# Unconventional Thermoelectric Materials for Energy Harvesting and Sensing Applications

Matteo Massetti, Fei Jiao, Andrew J Ferguson, Dan Zhao, Kosala Wijeratne, Alois Würger, Jeffrey L Blackburn, Xavier Crispin, Simone Fabiano

► To cite this version:

Matteo Massetti, Fei Jiao, Andrew J Ferguson, Dan Zhao, Kosala Wijeratne, et al. Unconventional Thermoelectric Materials for Energy Harvesting and Sensing Applications. Chemical Reviews, 2021, 121 (20), pp.12465-12547. 10.1021/acs.chemrev.1c00218. [hal-03334482](https://hal.science/hal-03334482)

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# Unconventional Thermoelectric Materials for Energy Harvesting and Sensing Applications

Matteo Massetti, Fei Jiao, Andrew J. Ferguson, Dan Zhao, Kosala Wijeratne, Alois Würger, Jeffrey L. Blackburn,* Xavier Crispin,* and Simone Fabiano*

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Cite This: https://doi.org/10.1021/acs.chemrev.1c00218

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**ABSTRACT:** Heat is an abundant but often wasted source of energy. Thus, harvesting just a portion of this tremendous amount of energy holds significant promise for a more sustainable society. While traditional solid-state inorganic semiconductors have dominated the research stage on thermal-to-electrical energy conversion, carbon-based semiconductors have recently attracted a great deal of attention as potential thermoelectric materials for low-temperature energy harvesting, primarily driven by the high abundance of their atomic elements, ease of processing/manufacturing, and intrinsically low thermal conductivity. This quest for new materials has resulted in the discovery of several new kinds of thermoelectric materials and concepts capable of converting a heat flux into an electrical current by means of various types of particles transporting the electric charge: (i) electrons, (ii) ions, and (iii) redox molecules. This has contributed to expanding the applications envisaged for thermoelectric materials far beyond simple conversion of heat into electricity. This is the motivation behind this review. This work is divided in three sections. In the first section, we present the basic principle of the thermoelectric effects when the particles transporting the electric charge are electrons, ions, and redox molecules and describe the conceptual differences between the three thermodiffusion phenomena. In the second section, we review the efforts made on developing devices exploiting these three effects and give a thorough understanding of what limits their performance. In the third section, we review the state-of-the-art thermoelectric materials investigated so far and provide a comprehensive understanding of what limits charge and energy transport in each of these classes of materials.

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## CONTENTS

1. Introduction

2. Thermoelectric Effects in a Nutshell

2.1. Electron Conductors

2.2. Nonredox-active Electrolyte

2.3. Redox-active Electrolyte

3. Thermoelectric Device Concept

3.1. Thermoelectric Generators (TEGs)

3.2. Thermogalvanic Cells (TGCs)

3.3. Ionic Thermoelectric Supercapacitors

4. State-of-the-Art Materials

4.1. Organic TE Materials

4.1.1. p-Type TE Polymers

4.1.2. n-Type TE Polymers and Composites

4.1.3. Small Molecules

4.1.4. Thermal Transport in Organic Semiconductors

4.2. Carbon Nanomaterials

4.2.1. Synthesis, Processing, and Basic Properties of SWCNTs

4.2.2. Tunable Thermoelectric Performance of "Neat SWCNTs"

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4.2.3. TE Generators from CNT-based Fibers and Fabrics

4.2.4. Graphene

4.2.5. Nanocarbon Composites

4.3. Ionic Thermoelectric Materials

4.4. Materials for Thermogalvanic Cells (TGCs)

4.4.1. Redox Couples

4.4.2. Electrode Materials

5. Conclusions and Outlooks

Author Information

Corresponding Authors

Authors

Author Contributions

Notes

Biographies

Acknowledgments

AH

AK

AN

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BF

BF

BJ

BN

BQ

BQ

BQ

BQ

BQ

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BR

Received: March 14, 2021

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<table><tr><td>References</td><td>BR</td></tr></table>

# 1. INTRODUCTION

Currently, global electricity production largely relies on fossil fuels (67%), water (16%), and nuclear energy (11%), while a small but rising fraction of electricity is produced by wind (>4%) and solar energy (>2%).¹,² However, the conversion of primary energy sources into electricity is not always efficient (typically 35–50% for heat engines, 20% for solar thermal plants, 15–40% for solar cells), and thus a substantial amount of energy is lost as waste heat. While more than 50% of the available natural and waste heat is stored in large volumes of warm fluids (T < 250 °C), no viable technology exists to produce electricity from this low energy density heat. Efforts to adapt heat engines by replacing water steam with organic solvent steam (organic Rankine cycle) have enabled low-temperature operation.³ However, the conversion is less efficient, reaching approximately 15–20% with heat engines that work at T_cold = 25 °C and T_hot = 150 or 250 °C.⁴ Another technology that can utilize this heat resource is the thermoelectric generator (TEG). This solid-state electronic device converts the heat flow into an electric current without the need for maintenance because no moving mechanical parts or fluids are present. Like all heat engines, a thermoelectric device operates across a temperature gradient ΔT = T_hot - T_cold. Today, the efficiency of the best TEGs is 15% for inorganic thin films prepared in a laboratory setting and 6.4% for industrially scalable materials based on the ball milling and sintering of bismuth telluride alloy nanoparticles (T_cold = 25 °C and T_hot = 250 °C).⁵ On the basis of these materials, solar thermoelectric generators with an efficiency of 4.6% under AM1 light intensity have been fabricated.⁶

The first practical challenge for TEGs is the scaling up of material synthesis. The second limitation is the high material costs due to the scarcity of the atomic elements involved in those exotic materials.⁷ Because of these challenges, it is not straightforward for TEGs to become an economically viable solution, outperforming heat engines in large-area energy harvesting applications. Without breakthroughs, it is unlikely that TEGs for low-temperature waste heat harvesting (T < 250 °C) will significantly impact the problem of CO₂ emissions.⁸ Hence, research efforts in recent years have focused on thermoelectric materials composed of abundant atomic elements. In this context, widely available, although less efficient, carbon-based thermoelectric materials, i.e., carbon nanostructures⁹ and conducting polymers¹⁰ or composites of these materials,¹¹ have become a new subject of major interest in the scientific community. The interest in these materials is the motivation for this review.

Because one can find heat sources practically everywhere, TEGs are interesting power sources for many low-power technologies. TEGs have particular advantages over heat engines for small-scale implementations or applications, where maintenance is a major issue, such as powering autarkic sensor systems in nature or satellites in space. The miniaturization of heat engines to dimensions below 1 cm³ leads to high-speed rotation, which wears out the cylinder piston (>10 000 rpm). Microengines also face many challenges in combustion, among which the main problem is the gap necessary between the piston and cylinder. Therefore, the efficiency drops from 40% to 0.05% when scaling down a cylinder from 10 cm³ to 1 cm³.¹² In contrast, the efficiency of TEGs remains constant from centimeter-scale to microthermoelectric generators.¹³ In

addition, the efficiencies of thermomechanical engines decreases drastically to lower power levels when scaling down,⁸ such that a crossing point exists where thermoelectric devices become more efficient.

With these scaling considerations in mind, TEGs are attractive candidates for power sources or self-powered heat sensors for portable electronics and the Internet of Things (IoT) applications. For instance, in outdoor activities, TEGs are integrated with either cooking pots (PowerPot) or boots (Orange Power Wellies) to charge a mobile phone. Many low-power devices, such as radiator valves (Micropelt), sensor nodes,¹⁴ and industrial process monitoring,¹⁵ can be powered with TEGs. Another area of major impact is robotics and medical applications because many biosensors and medical implants¹⁶ can be powered by exploiting body heat.¹⁷ In these applications, the integration of TEGs on fabrics¹⁸ and flexible substrates¹⁹ or as electronic skin²⁰ might require various levels of mechanical flexibility and eventually some degree of stretchability. Unlike conventional inorganic thermoelectric materials, which are brittle and nonflexible, organic semiconductors can maintain mechanical integrity under various modes of stress and strain. Regarding the biomedical field, thermoelectric devices can also be integrated as self-powered temperature sensors for continuous monitoring, which has become exceedingly critical after the worldwide COVID-19 pandemic. These self-powered devices could directly transmit the electrical signal without any transduction, enabling easy integration of these temperature sensors in wearable electronics for real-time health monitoring.

Apart from wearable applications, TE-based devices can act as sensors in locations where solar-based energy harvesting/ sensing is not a viable solution, e.g., in critical or hazardous areas inside buildings, pipelines, mines, aircraft, ships, and cars. Therefore, TE sensors may potentially impact fire detection, homeland security, and many other aspects of human life.²¹,²² TE-based sensors can also be employed to monitor the environmental temperature along sections of a domestic or industrial plant for evaluating efficiency. For all of these applications, mechanical flexibility is a strict requirement, and organic semiconductors, thermoelectric fluids, and gels exploiting the thermogalvanic effect offer new opportunities in the field of stretchable and flexible electronics.²³ Finally, the exploration of the ionic Seebeck effect through the Soret effect in polymer electrolytes leads to a surprisingly high Seebeck coefficient (~10 mV K⁻¹), promoting a new research direction that uses this effect to charge supercapacitors with heat.²⁴

Figure 1 summarizes the conductivity-dependent Seebeck coefficients for various classes of thermoelectric materials. In contrast to conductivity (σ), which is directly proportional to the charge carrier concentration (n) in the material, the Seebeck coefficient (S) typically decreases with increasing charge carrier concentration. Therefore, Seebeck coefficients equal to or higher than 1 mV K⁻¹ are likely to be found in insulator materials having electronic structures with large energy gaps.²⁵ However, the low electrical conductivity of insulators (below 10⁻¹² S cm⁻¹) makes reliable thermovoltage measurements very challenging. For heat-to-electricity conversion, the thermoelectric power factor (S²σ) is the most critical parameter, which typically shows a maximum in the intermediate charge concentration region corresponding to semiconducting materials (10⁻⁵ S cm⁻¹ < σ < 10⁴ S cm⁻¹ and 500 μV K⁻¹ > S > 50 μV K⁻¹). This is because the thermal conductivity (k) of these materials is usually below unity, mainly due to the absence of long-range crystalline order. Thus, the phonon component is
<figure><img src="image_10.png" /><figcaption>**Figure 1.** Seebeck coefficient as a function of the conductivity for organic/hybrid thermoelectric materials (see relative section 4.1), carbon nanomaterials (see relative section 4.2), organic nanomaterial composites (see relative section 4.2.5), organic thermoionic materials (see relative section 4.3), thermogalvanic materials (see relative section 4.4), inorganic thermoelectric materials (green open symbols), and metals (gray open symbols). For the conductivity of redox electrolytes, only the transport of ions is considered.</figcaption></figure>

effectively reduced. Indeed, this behavior represents one of the greatest advantages of organic semiconductors over inorganic semiconductors, where most of the current research is focused on decreasing thermal conductivity. Hence, research and technology development on organic thermoelectric materials has mainly focused on semiconductors and semimetals. As shown in Figure 1, nonredox electrolytes (green symbols) and redox electrolytes (thermogalvanics, purple symbols) can generate larger Seebeck coefficients than electronic organic (blue symbols) or inorganic (yellow symbols) semiconductors, which makes them potentially suitable for thermoelectric applications despite their low conductivity.

In this article, we review three major classes of thermoelectric materials and devices working in the low-temperature range ($T < 250$ °C) and based on highly abundant elements: (i) thermoelectric generators based on conducting polymers, carbon nanostructures, and their composites, (ii) thermogalvanic cells, and (iii) ionic thermoelectric supercapacitors/batteries. In the first section, we present the basics of thermoelectric phenomena. The second section reviews the working principle behind the devices exploiting the different thermodiffusion phenomena, highlighting the key parameters necessary to improve their performances. Finally, in the third section, we provide an overview of the key materials investigated so far by the research community, and summarize the recent efforts to produce devices employing such materials.

## 2. THERMOELECTRIC EFFECTS IN A NUTSHELL

When mobile (uncharged) particles are subject to a temperature gradient, they undergo a steady drift toward the hot or cold side. Typically, mobile particles at the hot end move faster than those at the cold end. Thus, as shown in Figure 2a, the hot particles have a larger mean free path and diffuse further than the cold particles, leading to a high net particle density at the cold end. This effect, observed in gaseous, liquid, and even solid systems, is called thermophoresis (also thermodiffusion, or the Soret effect) and has been known for a long time.²⁶ Thermophoretic forces

<figure><img src="image_11.png" /><figcaption>**Figure 2.** Illustration of (a) the thermodiffusion ($S_T > 0$) of particles and (b) three different concepts of thermoelectric effects: (1) electronic conductors, (2) nonredox electrolytes, (3) redox electrolytes.</figcaption></figure>

have been used, for example, in the manufacturing of optical fibers in vacuum deposition processes²⁷ or to separate different polymer particles in field-flow fractionation.²⁸ Furthermore, thermophoresis has been demonstrated as a versatile technique for manipulating single biological macromolecules, such as DNA, in micro- and nanochannels by means of light-induced local heating.²⁹

When a colloidal suspension is placed in a temperature gradient $\nabla T$, the particle motion can no longer be described by Brownian motion. In the absence of an external force, particles then tend to accumulate at the cold or hot boundary. This is phenomenologically described by writing the particle current density ($j$) as:

$$j = -D\nabla c - D_T c \nabla T \quad (1)$$

where $c$ is the particle concentration, $D$ the diffusion coefficient of a colloidal suspension under Brownian motion, and $D_T$ the thermodiffusion coefficient. The two components of eq 1 represent the motion of particles under a concentration gradient ($\nabla c$, kinetic diffusion) and a thermal gradient ($\nabla T$, thermal diffusion).

The dispersed particles display a thermodiffusive drift velocity $\vec{v}$, which depends linearly on $\nabla T$ with a proportionality constant equal to $D_T$:

$$\vec{v} = -D_T \nabla T, \quad (2)$$

Depending on the sign of $D_T$, the particles will accumulate either at the cold or the hot side. In a closed system, the steady state implies zero current ($j = 0$), $c\nu + D\nabla c = 0$, leading to a stationary concentration gradient:

$$\nabla c = -cS_T \nabla T \quad (3)$$

where the Soret coefficient $S_T = D_T/D$ is provided by the ratio of $D_T$ and the ordinary diffusion coefficient $D$. These transport coefficients, $D_T$ and $D$, describe nonequilibrium properties.

Gaseous systems can be dealt with exactly in the framework of kinetic theory. In the 1910s, Enskog and Chapman found that in a gas mixture, heavier and larger atoms accumulate at the cold side.³⁰ These mass and size effects were confirmed experimentally with ammonium isotopes $N_{14}H_3/N_{15}H_3$ and dimer molecules $H_2, D_2, N_2$, and $O_2$.

A more complex problem is encountered in liquids because the transport properties are strongly influenced by molecular
interactions, such as electrostatic and dispersion forces, and the hydrodynamic flows induced by moving particles. In many instances, even the sign of the thermodiffusion coefficient $D_T$ cannot be determined from simple arguments.³¹ Moreover, the Soret coefficient shows an intricate size dependence. For very small molecules, the coefficients $D_T$ and $D$ carry the same hydrodynamic factor, and the Soret coefficient $S_T = -H/k_B T^2$ is provided by the molecular solvation enthalpy $H$, which is a purely thermodynamic quantity. However, for colloidal particles and polymers, $D_T$ is independent of the size or molecular weight, and instead, $D$ decreases with the hydrodynamic radius. As a consequence, the Soret coefficient $S_T$ increases with the particle size, and large particles and high polymers accumulate exponentially at the boundary.

If the particles subject to the thermal gradient are charged (Figure 2b), the accumulation of particles at the cold end will also produce a repulsive electrostatic force and therefore electric potential, which pushes the charges back to the hot end. The electric potential difference ($\Delta V$) produced by the temperature difference ($\Delta T$) is known as the Seebeck effect, and the proportionality constant is called the Seebeck coefficient ($S$):

$$ \Delta V = -S\Delta T \quad (4) $$

This Seebeck coefficient, or thermopower, is related to the entropy of charge carriers within a material with a temperature gradient. Carriers move from the hot side of the material (high entropy) to the cold side of the material (low entropy), establishing a steady-state voltage.

## 2.1. Electron Conductors

In electronic thermoelectric materials, charges are carried by free electrons and holes. At the hot end of the thermoelectric material, the number of free charge carriers is high, and accordingly, the entropy is larger than at the cold end. As a consequence, there is a steady flow of electrons and holes toward the cold end, where they recombine. From a thermodynamic point of view, the corresponding entropy flow from hot to cold drives the thermoelectric effect. In the first approximation, these electrons and holes behave like a gas of charged particles. Hence, their dynamics are described by kinetic theory, for example, in a quasi-classical approximation that keeps quantum statistics but relies on a classical equation of motion for the carriers.³² In a closed geometry, the Seebeck coefficient is readily obtained from the condition that thermodiffusion and thermoelectric currents cancel each other, $J = \sigma\Delta V - \sigma_T\Delta T = 0$, where the electronic transport is described by the conductivity $\sigma$ and the thermodiffusion coefficient $\sigma_T$. Analogous to the Soret effect discussed above, the Seebeck coefficient is defined as $S = \sigma_T/\sigma$. Whereas $\sigma$ is always positive, the sign of $\sigma_T$ depends on whether electrons or holes dominate thermodiffusion. For a more extensive discussion and explanation of the thermodynamic and kinetic models underlying the Seebeck effect, we refer the reader to the following review by Goupil et al.³³

In a p-doped semiconductor, holes are the majority mobile charge carriers; thus, a positive charge builds up at the cold end, generating a positive potential. Similarly, when electrons are the majority free carriers (n-type material), they will produce a negative potential at the cold end (see Figure 2b, example 1). By convention, the sign of the electronic Seebeck coefficient ($S_e$) is provided by the potential at the cold side with respect to the hot side, thus indicating the type of majority charge carriers, i.e., electrons or holes, with $S_e > 0$ for p-type semiconductors and $S_e < 0$ for n-type semiconductors.

In a metal, electronic transport properties are intimately related to the electronic band structure and to the density of states at the Fermi level $E_F$. To determine whether electrons or holes are the majority carriers in an electrical conductor, one needs to look at the density of states $N(E)$ in an energy window $k_B T$ centered at the Fermi level $E_F$. If $N(E)$ is constant in this range, the electron contribution to the thermodiffusion current cancels that of the holes, $J_e + J_h = 0$, as shown schematically in Figure 3a together with the Fermi-Dirac distribution $f(E) = 1/[e^{(E-E_F)/k_B T} + 1]$. If, on the other hand, the density of states increases with energy, $dN(E)/dE > 0$, there are more carriers with $E > E_F$ than “holes”, resulting in a net charge current $J$. In a closed geometry, this results in a polarization field and a thermoelectric potential with a negative Seebeck coefficient, $S_e < 0$ (Figure 3b). By the same token, if the density of states decreases with energy as in Figure 3c, there are more “hole” carriers below the Fermi level, resulting in a positive Seebeck coefficient $S_e > 0$.

These qualitative considerations can be made rigorous in terms of a Sommerfeld expansion,³² which results in the Mott formula³⁴,³⁵

$$ \begin{aligned} S_e(E_F) &= \frac{\pi^2 k_B}{3 e} (k_B T) \left[ \frac{\partial \ln \sigma}{\partial E} \right]_{E=E_F} \\ &= \frac{\pi^2 k_B}{3 e} (k_B T) \left\{ \frac{\partial[n(E)]}{n \partial E} + \frac{\partial[\mu(E)]}{\mu \partial E} \right\}_{E=E_F} \end{aligned} \quad (5) $$

This simple expression rationalizes two important features of the thermoelectric effect. First, noting that the last factor is of the order of the inverse Fermi level, one finds $S_e(E_F) \sim (k_B/e)(k_B T/E_F)$, which is smaller, by a factor of $k_B T/E_F$, than the result from classical Drude theory. Second, the Seebeck coefficient may be observed to have a negative sign, e.g., for bismuth, nickel, and alkali metals, whereas classical theory always predicts $S_e > 0$. These features highlight the quantum nature of the thermoelectric effect.

Organic semiconductors are generally crystalline or amorphous materials with weak intermolecular interactions. This weak coupling leads to narrow electronic bandwidths. Unlike nearly perfect crystals where the electronic states are delocalized, charge carriers in organic semiconductors are usually localized over spatially and energetically distributed transport sites. The widely accepted transport mechanism is that the charge carriers hop from one localized state to another within a lattice of molecular sites. In general, the Seebeck coefficient of conducting polymers varies broadly, ranging from <10 µV K⁻¹ for highly doped polymers up to 1000 µV K⁻¹ for lightly doped polymers. The dramatic reduction of $S_e$ at high doping levels arises from $E_F$ moving into the conduction band where the number of electronic states above and below $E_F$ are equivalent.³⁶

## 2.2. Nonredox-active Electrolyte

Similar to electrons, when ions in an electrolyte are subject to a thermal gradient $\nabla T$ (Figure 2b, example 2), they thermodiffuse from the hot side to the cold side, inducing a concentration gradient $\nabla c/c = -S_T \nabla T$ (Soret effect) and an internal electric field $E = -S_T \nabla T$ (Seebeck effect).

These effects result in entropy flow through a nonequilibrium system,²⁶ and the underlying thermodynamic forces on small ions are provided by gradients of the Massieu-Planck entropy potentials $-\mu_{\pm}/T$. Because the chemical potentials $\mu_{\pm}$ ($T, c_{\pm}$) depend on both temperature $T(r)$ and concentration $c_{\pm}(r)$, the thermodynamic forces comprise gradients of both quantities, $f_{\pm}$
<figure><img src="image_12.png" /><figcaption>Figure 3. (a) Case of a constant density of states N(E). The left and right sides correspond to the cold and hot ends of the material. The Fermi distribution (green curves) is plotted for the two temperatures (T = 0 K and T > 0 K). (b) Case of a linearly increasing N(E) with a net electron current and a negative thermopower. (c) Case of a linearly decreasing N(E) that is specific to conductors with a positive Seebeck coefficient.</figcaption></figure>

$= H_{\pm} \nabla T/T - k_B T \nabla c_{\pm}/c_{\pm}$, where the first term drives thermodiffusion with the solvation enthalpy $H_{\pm}$ and the second term drives the usual gradient diffusion. Imposing zero ion currents and Gauss' law, one readily obtains the steady-state Soret coefficient of the salt solution,³⁷

$$S_T = - \frac{H_+ + H_-}{k_B T^2} \qquad (6)$$

and the electrolyte Seebeck coefficient,

$$S_i = \frac{H_+ - H_-}{2eT} \qquad (7)$$

Measured ionic enthalpies are generally negative and on the order of $k_B T$,³⁸ such that the Soret coefficient is positive. The Seebeck coefficient may take either sign with absolute values on the order of $k_B/e \sim 100 \mu V K^{-1}$. The thermoelectric field often dominates thermophoresis in colloidal suspensions. Because of the different ion enthalpies, the Seebeck field is expected to change its sign upon varying the composition x of mixed electrolytes, such as $NaCl_xOH_{1-x}$.³⁹ This effect has been confirmed for suspensions of nanomicelles⁴⁰ and micrometer-scale colloidal particles,⁴¹ which is in quantitative agreement with the measured enthalpies.³⁸ As a caveat, note that eq 6 applies to small ions only, the size of which is small compared to the Debye screening length. In regard to ions larger than the Debye length, such as polyelectrolytes, the Seebeck coefficient takes a more intricate form.⁴²

Ionic enthalpies arise from complex water-ion interactions that are comprised of electrostatic, dispersion, and hydration forces. To date, there is no satisfactory theoretical framework for calculating enthalpies from first principles. More recently, nonequilibrium molecular dynamics simulations were used to study the Soret coefficient of alkali halide solutions and showed that the ordering of water around alkali cations impacts the Soret coefficient of the salts⁴³ and that the thermal polarization of free water affects the thermoelectric behavior of the solution.⁴⁴

Experimentally, the $S_i$ of a nonredox-active electrolyte sandwiched between two electrodes is defined as the ratio between the open-circuit potential ($\Delta V$) and the $\Delta T$ between the hot and cold sides (compared to the ionic Seebeck coefficient, the Seebeck coefficient of the metal electrodes is small ($<10 \mu V K^{-1}$) and can be neglected). A survey of nonredox-active electrolytes of interest for ionic thermoelectric applications is presented in section 4.

## 2.3. Redox-active Electrolyte

When the electrolyte subjected to the thermal gradient is also redox active, it undergoes an electrochemical reaction at the hot and cold electrodes (Figure 2b, example 3):

$$O_x + e^- \rightleftharpoons Red \qquad (8)$$

Upon electron transfer, the reaction releases or requires free energy ($\Delta G$) expressed as entropic ($\Delta S$) and enthalpic ($\Delta H$) contributions:

$$\Delta G = \Delta H - T\Delta S \qquad (9)$$

As a consequence, the concentrations satisfy the mass action law:

$$n_{ox} = e^{-\Delta G/k_B T} n_{red} = \varphi^{-1} n_{red} \qquad (10)$$

The temperature difference at the electrodes results in a spatial variation of the fugacity $\varphi$, which in turn imposes gradients of oxidant and reductant,

$$\nabla n_{ox} = \varphi^{-1} \nabla n_{red} + \varphi^{-1} n_{red} \frac{\Delta H}{k_B T} \frac{\nabla T}{T} \qquad (11)$$

In addition to eq 10, the stationary state satisfies the condition of zero current of each ion species $J_i = 0$, and Gauss's law. Proceeding as in ref 45, one finds a potential difference between hot and cold electrodes, $\Delta V_{redox} = -S_{redox} \Delta T$, where the Seebeck coefficient is given by the reaction enthalpy:
<figure><img src="image_13.png" /><figcaption>**Figure 4.** (a) Sketch of the different thermal energy fluxes passing through a thermoelectric leg subjected to a temperature gradient. Adapted with permission from ref 49. Copyright 2012 Royal Society of Chemistry. (b) Representation of a standard power output vs load resistance for TEG. (c) Classical thermoelectric modules, based on n- and p-legs, interconnects, and ceramic plates. Reproduced with permission from ref 49. Copyright 2012 Royal Society of Chemistry. (d) Woven-yarn structure to integrate TEG into textile for fibers. Reproduced with permission from ref 50. Copyright 2020 Nature. (e) Sketch of a planar thermoelectric module on a flexible substrate. Reproduced with permission from ref 51. Copyright 2011 IOP. (f) Unipolar thermoelectric generator, with a planar structure. Reproduced with permission from ref 47. Copyright 2017 Royal Society of Chemistry.</figcaption></figure>

$$S_{\text{redox}} = \frac{\Delta H}{eT} \qquad (12)$$

Taking the sum of the thermodiffusion and redox contribution, we have:

$$S = S_i + S_{\text{redox}} \qquad (13)$$

As a caveat, we note that the two terms are not additive in general.

# 3. THERMOELECTRIC DEVICE CONCEPT

## 3.1. Thermoelectric Generators (TEGs)

The fundamental element of a TEG is a thermocouple consisting of p-type and n-type TE material that are connected electrically in series and thermally in parallel, having an internal electrical resistance ($R_{pn} = R_p + R_n$) without considering the metallic electrical interconnections) and subjected to a temperature gradient $\Delta T = T_H - T_C$. Assuming that heat propagates from the hot to the cold side only through the thermocouple, then the losses through the metal contacts, due to their high thermal conductivity, can be neglected. This configuration is exploited to maximize the electrical output by creating a unidirectional charge flow inside the device by summing the thermovoltage

contributions from the two materials ($S = S_p - S_n$). The thermal energy flow consists of three contributions (Figure 4a): (i) the rate of Peltier heat absorbed from the hot side due to the thermogenerated current $I$ ($Q_{\text{Peltier}} = ST_H I$), (ii) the Joule heating emanating at the hot side due to the thermogenerated current $I$ ($Q_{\text{Joule}} = 1/2 I^2 R$), and (iii) the heat flux crossing the leg ($K\Delta T$) in the absence of an electrical current ($K$ is the thermal conductance of the thermocouple). At the hot junction, the thermal input power can be expressed as:

$$Q_H = ST_H I - \frac{1}{2}RI^2 + K(T_H - T_C) \qquad (14)$$

The electrical power output of the TEG when the current passes across a load resistance ($R_L$) is:

$$P = VI \qquad (15)$$

with an open-circuit voltage:

$$V = S\Delta T \qquad (16)$$

and a current:

$$I = S\Delta T/(R + R_L) \qquad (17)$$

The efficiency $\Phi$ of the TEG is the ratio between the electrical power output $P$ and the thermal power input $Q_H$:
$$ \Phi = \frac{S\Delta TI}{(ST_HI - \frac{1}{2}RI^2 + K\Delta T)} \quad (18) $$

By adjusting the dimensions of the thermocouple elements, it is possible to match the internal resistance $R_{pn}$ and the load resistance $R_L$. Thus, the maximum efficiency $\Phi_{max}$ can be written as$^{46}$

$$ \Phi_{max} = \frac{\Delta T}{T_H} \frac{\sqrt{1 + ZT} - 1}{\sqrt{1 + ZT} + T_C/T_H} \quad (19) $$

$$ Z = \frac{S^2}{K_{pn}R_{pn}} \quad (20) $$

where $ZT$ is the dimensionless figure-of-merit for the thermocouple and expresses the conversion efficiency of the thermocouple. From this parameter, another parameter ($zT$) has been derived, which is based on the thermoelectric properties of a single material ($S, k, \sigma$):

$$ zT = \frac{\sigma S^2}{k} T \quad (21) $$

The regimes of operation at maximum efficiency and maximum power output differ. At maximum power, the external load resistance equals the internal resistance of the device $R_L = R_{pn}$ (Figure 4b). The output voltage is half of the open-circuit voltage $V = V_{oc}/2$, so that the maximum power generated is:

$$ P_{max} = \frac{V_{oc}^2}{4R} = \frac{S^2(T_H - T_C)^2}{4R} \quad (22) $$

The Seebeck voltage produced by a single thermocouple is too small ($S \sim 200 \mu V K^{-1}$) to power any application; therefore, to reach suitable power outputs, the strategy is to connect multiple thermocouples electrically in series and thermally in parallel in a high-density configuration to maximize the power extraction per unit area. Finally, the thermoelectric module is sandwiched between ceramic plates to ensure good thermal exchange with the surroundings and to avoid short circuits between the metal contacts. The design of OTEGs varies in shape and size. In addition to the typical vertical geometry, which includes bulky pellets of active materials, OTEGs can also be manufactured as flat plates for thin films or woven-yarn textiles for fibers (Figure 4d,e). Using only one thermoelement, unipolar OTEG can sometimes be fabricated (Figure 4f).$^{47,48}$ Moreover, OTEG can be manufactured by low-cost, low-temperature, and large-area techniques such as printing and drop casting.

## 3.2. Thermogalvanic Cells (TGCs)

A thermogalvanic cell consists of two electrodes at different temperatures in contact with an electrolyte containing a redox couple. The redox electrolyte contains the reductant (red) and the oxidant (ox) of one electrochemical half-reaction. Importantly, the electron transfer of this half reaction at the electrode must be reversible. As illustrated in Figure 5, the thermogalvanic cell is an electrochemical thermoelectric generator that produces a difference in electric potential when a temperature gradient is set across the two electrodes. Because the redox electrolyte will transport an electric current through molecular shuttles by convection, diffusion, and migration between the two electrodes, a constant power can be extracted by connecting two electrodes to an external $R_L$.$^{52}$

<figure><img src="image_14.png" /><figcaption>Figure 5. (a) Schematic diagram and (b) equivalent circuit of a typical thermogalvanic cell. (c) $I-V$ and $P-V$ curves of a thermogalvanic cell. Reproduced with permission from ref 52. Copyright 2018 Linköping University Press.</figcaption></figure>

The equivalent electrical circuit of a TGC (Figure 5b) is a generator composed of an ideal potential source (Seebeck voltage) in series with an internal resistance $R$. The current-voltage characteristics follow that of a Thevenin generator:

$$ V = V_{oc} - RI \quad (23) $$

$$ V_{oc} = -S_{redox} \Delta T \quad (24) $$

The internal resistance is provided by the slope of the $I-V$ characteristic, while the intercept with the $V$-axis is the open-circuit voltage $V_{oc}$, and the intercept on the $I$-axis is the short-circuit current $I_{sc}$ (Figure 5c). The power varies quadratically with the thermogenerated current (Figure 5c):

$$ P = VI = V_{oc}I - RI^2 \quad (25) $$

The conditions for maximum power are found on the $I-V$ characteristics as the largest area under a rectangle determined by the point ($I_{sc}/2, V_{oc}/2$):

$$ P_{max} = \frac{V_{oc}I_{sc}}{4} \quad (26) $$

The maximum power corresponds to a choice of load resistance that is equal to the internal resistance (eq 24 and eq 16). Three types of resistive phenomena contribute to the internal resistance: (i) mass transfer resistance, (ii) ohmic resistance, and (iii) charge transfer resistance.$^{53-55}$ The ohmic resistance originates from the resistance due to the electrical connections to the external circuit, the bulk electrolyte resistance, and the electrode resistance. Often, the dominant phenomenon is the ionic resistance of the electrolyte ($R_i$):

$$ R_i = \frac{L}{A\sigma} \quad (27) $$

where $A$ is the cross-sectional area defined by the geometry of the two electrodes, and $L$ is the separation distance of the two electrodes. For noninteracting ions, the ionic conductivity $\sigma$ is approximated by:

$$ \sigma = \frac{F^2}{RT} \sum_i D_i c_i z_i^2 \quad (28) $$
<figure><img src="image_15.png" /><figcaption>Figure 6. Thermoelectric effects with (a) electronic charge carriers (electrons or holes) in a semiconductor and (b) ionic charge carriers in an electrolyte. Reproduced with permission from ref 64. Copyright 2017 Wiley-VCH.</figcaption></figure>

where $D_i$ is the diffusion coefficient of ion i, $c_i$ is the ionic concentration, and $z_i$ is its charge.

Hence, reducing the internal resistance is possible by increasing the concentration of the redox electrolyte and increasing the diffusion coefficient. In the Stokes–Einstein model, the ionic diffusion coefficient is a function of the viscosity $\gamma$ of the solvent and the hydrodynamic radius $r_H$ of the ions:

$$D = \frac{k_B T}{6\pi\gamma r_H} \qquad (29)$$

where $k_B$ is the Boltzmann constant and T is the absolute temperature. Hence, small, solvated ions in a low viscosity solvent are two factors minimizing the internal resistance of the TGCs.

The rate of the electron transfer reaction at the electrode/ electrolyte interface is related to the exchange current, $I_0$, which is characterized by the charge transfer resistance $R_{CT}$:

$$R_{CT} = \frac{RT}{AnF|I_0|} \qquad (30)$$

where A is the area of the electrode and R is the gas constant. The exchange current is provided by the Butler–Volmer expression of the electrochemical current (Tafel equation):

$$\ln(I) = \ln(I_o) - \frac{\alpha n F \eta}{RT} \qquad (31)$$

where $\eta$ is the overpotential and $\alpha$ is the charge transfer coefficient. $R_{CT}$ can be as small as 1 $\Omega$ in a good TGC, e.g., the electron transfer at the platinum electrode for a ferric/ ferrocyanide redox couple in salted water.⁵⁶

The mass transport resistance, i.e., the transport of reactants from the solution to the electrode is governed by three main processes (migration, diffusion, and convection) that limit the current $I_{lim}$:⁵⁵

$$R_{mt} = \frac{RT}{nF|I_{lim}|} \qquad (32)$$

If the steady-state current $|I_{lim}|$ is limited by the diffusion of reactants, it can be written as:

$$I_{lim} = \frac{nFADc}{\delta} \qquad (33)$$

where D is the diffusion coefficient of the electroactive species, A is the active electrode area, c is the concentration of the redox species, and $\delta$ is the thickness of the diffusion layer at the electrode, which theoretically is the maximum of the interelectrode distance. In practice, $\delta$ is smaller than the interelectrode distance because of convection.⁵⁵

The thermoelectric figure-of-merit ZT (eq 19) cannot provide a simple expression for thermogalvanic cells because several resistive contributions might limit the thermogenerated current. Because those contributions can be seen in series, depending on the geometry of the cell and its constituents (electrolyte, redox species, electrode), one resistive contribution can be dominant, and a simplified expression can be extracted. In many cases, the $R_{mt}$ is larger than (i) $R_i$ with high salt and redox couple concentrations and (ii) $R_{CT}$ using conventional metal electrodes. In that case, Abraham et al. proposed the following expression:⁵⁵

$$ZT^* = \left( \frac{z^2 F^2}{R} \right) S_{\text{redox}}^2 \frac{D_{\text{lim}} c}{\kappa} \qquad (34)$$

where z is the charge of the ion, F is the Faraday constant, $D_{lim}$ is the limiting diffusion coefficient, c is the concentration of the redox couple, $\kappa$ is the thermal conductivity, and R is the gas constant. The power conversion efficiency is the ratio between the output electric power and the input thermal power in the thermogalvanic cell:

$$\Phi = \frac{\text{electrical power}}{\text{thermal power}} \qquad (35)$$

The input thermal power depends on the geometry of the cell defining the thermal conductance but also the thermal contact between the various materials and the thermal conductivity of each material (ceramic, metal, electrolyte, separator). By neglecting the small thermal contributions, the rate of heat flow from the hot electrode to the cold electrode is dominated by the heat transport through the electrolyte and the rate of heat transfer induced by the electrochemical current.⁵⁷
<figure><img src="image_16.png" /></figure>

<figure><img src="image_17.png" /><figcaption>**Figure 7.** (a) Full charge and discharge cycle of the ionic thermoelectric supercapacitor (ITESC). (i) Establishing an ionic thermovoltage while a ΔT is applied to the device, (ii) thermoelectric charging of the supercapacitor, (iii) cooling to reach equilibrium at ΔT = 0, and (iv) discharging. (b) Electrical energy stored in a supercapacitor upon either electrical charging with an external applied V<sub>charging</sub> or thermoelectric charging (temperature gradient) through the electrolyte of the supercapacitor. Reproduced with permission from ref 24. Copyright 2016 The Royal Society of Chemistry.</figcaption></figure>

$$ \frac{dQ}{dT} = kA \frac{\Delta T}{d} + \frac{IT \Delta S}{nF} \qquad (36) $$

where k is the thermal conductivity of the electrolyte, A is the electrode cross-sectional area, ΔT is the temperature gradient between the two electrodes separated by a distance d, I is the cell current, and ΔS is the entropy of the cell reaction. In practice, the second term can be neglected because it is only a few percent of the total expression. Within these approximations, the total efficiency at maximum power is written as:<sup>54,58</sup>

$$ \Phi = \frac{V_{oc}I_{SC}}{4KA\frac{\Delta T}{d}} \qquad (37) $$

where I<sub>SC</sub> is the short-circuit current limited by the most resistive phenomena (R<sub>ET</sub>, R<sub>I</sub>, or R<sub>mt</sub>). To compare the efficiency Φ of different devices working at different temperatures (T<sub>H</sub> is the hot electrode temperature), the relative conversion efficiency Φ<sub>r</sub> is determined as the ratio of the total efficiency to the Carnot efficiency Φ<sub>C</sub>.

$$ \Phi_r = \frac{\Phi}{\Phi_C} \qquad (38) $$

$$ \Phi_c = \frac{\Delta T}{T_H} \qquad (39) $$

For the redox couple ferric/ferrocyanide and a flat platinum electrode, Φ<sub>r</sub> is approximately 0.50%.<sup>59,60</sup> Increasing the surface area of the electrodes is a strategy to increase efficiency. Indeed, multiwalled carbon nanotube electrodes<sup>61</sup> achieve Φ<sub>r</sub> = 1.4%, while electrodes made of composites of single-walled carbon nanotubes and reduced graphene oxide electrodes reach Φ<sub>r</sub> = 2.6%;<sup>62</sup> finally, platinum nanoparticle-decorated carbon aerogel electrodes reach a record high efficiency of 4.0%.<sup>63</sup>

### 3.3. Ionic Thermoelectric Supercapacitors

While ions can diffuse along the temperature gradient as electronic charge carriers, they cannot pass through a metal electrode to generate an external current (as illustrated in Figure 6). Instead, ions will accumulate at the interface between the metal electrodes and the electrolyte acting as an electric double

layer (EDL), which induces a transient current; additionally, the charge stored in EDL capacitors can be represented by integrating the current.<sup>64,65</sup> Therefore, the ionic thermoelectric effect is not suitable for typical thermoelectric generators under continuous operation. Alternatively, when using high capacitance electrode materials, such as carbon nanotubes and conductive polymers, the amount of accumulated charge can be greatly enhanced, and electrical energy can be harvested by charging supercapacitors (the so-called ionic thermoelectric supercapacitor ITESC)<sup>24,66,67</sup> or batteries. In this way, thermal energy is converted into stored electrical energy through the ionic thermoelectric effect and can be consumed upon discharging.

The basic operational protocol of an ITESC is illustrated in Figure 7a.<sup>24</sup> (i) First, a ΔT is established between the two electrodes separated by an electrolyte, and after a certain stabilization time (t<sub>st</sub>), a thermovoltage reaches V<sub>thermo</sub> = S<sub>i</sub>ΔT. (ii) The two electrodes are connected through a load resistance and V<sub>thermo</sub> charges the supercapacitor. The charge Q<sub>ch</sub> stored at the electrodes of the ITESC can be obtained by integrating the current during charging. (iii) When the device is fully charged, the external load resistance is disconnected (open circuit), and the heat source is turned off. The thermovoltage decays to zero without ΔT, and the open-circuit potential is mainly dependent on the stored charge at the electrode/electrolyte interface. (iv) Finally, the device is discharged by connecting the two electrodes through an external circuit. The amount of charge Q<sub>dis</sub> equals the integration of the discharge current. Ideally, the maximum stored electrical energy is described as:

$$ E_{ch} = \frac{1}{2}CV_{thermo}^2 = \frac{1}{2}C(\Delta TS_{i})^2 \qquad (40) $$

which quadratically increases with the Seebeck coefficient of the material. Figure 7b displays the electrical energy stored in a supercapacitor versus both voltage in the case of electrical charging and temperature in the case of thermoelectric charging. An ITESC is suitable for harvesting energy from intermittent heat sources, namely, the sun. During the daytime, ITESCs can be charged by being heated by the sun, and during the night, the stored electrical charge can be utilized by discharging the device.
Recently, Kim et al. developed an ITESC with a polystyrene sulfonic acid (PSSH) electrolyte and observed almost no leakage current, showing that the stored electric energy could be maintained for more than 24 h.⁶⁶ As the temperature changes, the equilibrium condition at the solid-liquid interface is disrupted, and the ion density varies, affecting the interfacial potential. The role of the electrode material work function⁶⁸ was systematically studied: by changing the temperature, the charge motion creates a varying potential difference across the double layer at the electrode-electrolyte interface such that a direct relationship between temperature sensitivity and work function of electrode potential could be shown. In addition, other studies were conducted to investigate the role of ion concentration⁶⁹ and ion size:⁷⁰ the results showed that the ions size is the main factor affecting the voltage generation (the bigger the ion, the lower the voltage generated for the same ΔT), while the ion concentration seemed to play a secondary role.

The efficiency of the heating-cooling cycle in ITESCs could be derived from the ratio of the produced electrical energy and the thermal energy passing through the device.⁶⁴ The maximum charging efficiency (Φ<sub>ΔTch</sub>) of an ITESC can then be expressed as:

$$ \Phi_{\Delta T_{\text{ch}}} = \frac{\Delta T}{2T_{\text{H}} + \frac{10k}{\sigma_{S_i}^2} - \frac{1}{2}\Delta T} \quad (41) $$

The maximum charging efficiency is closely related to the key material parameters, i.e., σ<sub>v</sub>, S<sub>v</sub> and k, which are gathered to define the ionic thermoelectric figure-of-merit of an electrolyte zT<sub>i</sub>:

$$ zT_i = \frac{\sigma_v S_i^2}{k} T \quad (42) $$

Similar to classic electronic materials, the efficiency will increase for electrolytes with high ionic conductivity and Seebeck coefficient and low thermal conductivity. They also directly derived Φ<sub>ch</sub> based on ZT<sub>i</sub>:

$$ \Phi_{\text{ch}} = \frac{\Delta T}{T_{\text{H}}} \frac{ZT_i}{2ZT_i + \frac{10T}{T_{\text{H}}} - \frac{1}{2}ZT_i \frac{\Delta T}{T_{\text{H}}}} \quad (43) $$

Compared with electronic thermoelectric materials, electrolytes have a higher Seebeck coefficient and lower thermal conductivity. However, because of their relatively low ionic conductivity, the zT<sub>i</sub> values and efficiencies are lower than those of good electronic thermoelectric materials (Figure 8a). Note that the efficiency evaluation here only considers the charging process and neglects the heat used to warm up the materials, which is proportional to the mass and heat capacitance of the electrodes and electrolyte. In a first approximation, this contribution to thermal energy is assumed to be similar for an ITESC and for a supercapacitor that is charged by a TEG. Importantly, the comparison of zT between an electronic conductor and an ionic conductor has meaning only if the electronic conductor is considered for a TEG coupled in series with a supercapacitor. Indeed, this circuit is an equivalent circuit of an ITESC. Despite the relatively low efficiency of ITESC, it shows one key advantage compared to a TEG coupled with a supercapacitor. At comparable temperature differences, electronic thermoelectric materials having good performances (e.g., Bi<sub>2</sub>Te<sub>3</sub> alloys) possess a lower Seebeck coefficient (~200 μV K<sup>-1</sup>) than ionic thermoelectrics (up to 10000 μV K<sup>-1</sup>). As a result, the stored energy in the same supercapacitor (C = 1 mF)

<figure><img src="image_18.png" /><figcaption>Figure 8. (a) Efficiency vs zT of typical electronic and ionic thermoelectric materials (ΔT = 10 K). (b) Stored energy in a supercapacitor with a capacitance of 1 mF as a function of the Seebeck coefficient of different thermoelectric materials at ΔT = 10 K. Adapted with permission from ref 64. Copyright 2017 Wiley-VCH.</figcaption></figure>

charged by ionic thermoelectric materials is many orders of magnitude larger than that charged with an electronic material, as illustrated in Figure 8b.

# 4. STATE-OF-THE-ART MATERIALS

## 4.1. Organic TE Materials

The versatile chemical synthesis, low-cost solution processability, and exclusive mechanical robustness endow organic conductors with broad appeal in thermoelectric applications. However, the limited solubility of organic conductive materials is typically a drawback with regard to building vertical legs (thickness of 10–100 μm) in the traditional architecture of a thermoelectric module. One must also consider the anisotropy of charge transport in organic materials, especially in the electrical conductivity where the difference between the in-plane/cross-plane σ can reach very high values (10<sup>5</sup> S cm<sup>-1</sup> in PEDOT:PSS).⁷¹ Processing powder into the shape of pellets, such as typically employed for inorganics, is of course a possible alternative.

Despite their low solubility, solution-processable organic conductors can form thin films, which provide new opportunities by enabling the design of OTEGs in various shapes and sizes, e.g., flat plates for thin films or woven-yarn textiles for fibers.⁷² Ultimately, OTEGs can be manufactured by low-cost, low-temperature, and large-area techniques, such as printing and drop casting. Although the power density is relatively low, OTEGs provide the possibility for a variety of novel low-power applications.⁷³ Because the main effort in OTEGs is to optimize the thermoelectric properties of materials, we review the milestones with regard to material development and understanding.

The molecular structures of some typical organic electronic (semi)conductors explored for thermoelectric applications are shown in Figure 9. The performance of materials with a zT value
<figure><img src="image_19.png" /><figcaption>Figure 9. Chemical structures of the typical conductors and dopant molecules used in OTEGs.</figcaption></figure>

larger than 0.1 or a PF larger than 100 µW m⁻¹ K⁻² are summarized in Table 1.

**4.1.1. p-Type TE Polymers.** Since high conductivity was discovered in I₂-doped polyacetylene (PA), many different conducting polymers, such as polypyrrole (PPy), polythiophene (PTh), polyaniline (PANi), and poly(3,4-ethylenedioxythiophene) (PEDOT), have been designed and synthesized, spurring the development of organic electronics.¹⁰⁰⁻¹¹¹ Typically, conducting polymers have relatively low thermal conductivity (0.1–0.7 W m⁻¹ K⁻¹). Thus, they are expected to be good candidates for organic thermoelectric applications if their Seebeck coefficient can be improved in their highly conducting state. Because conducting polymers are typically formed as thin films (<1 µm in thickness), it is experimentally challenging to accurately measure their out-of-plane thermal conductivity. For that reason, not all polymers are fully

characterized. However, the numerator of zT, which is the power factor (PF), enables the first comparison of the thermoelectric performance of various polymers:

$$PF = \sigma S^2 \quad (44)$$

The conductivity, Seebeck coefficient, and thermal conductivity of a conductive polymer are interrelated and each exhibit a different dependence on carrier density. Hence, the maximum zT is usually found for a specific carrier density. Other parameters also affect the transport properties, such as the regioregularity of the polymer chains¹¹² and the distribution of molecular weight,¹¹³ which depend on the chemical synthesis, molecular disorder, nanomorphology obtained upon film formation, and fabrication methods.⁸⁸

**4.1.1.1. Doping/Dedoping.** The use of chemical doping for the optimization of OTE materials dates back to as early as 1980,
<table><caption>Table 1. Summary of Reported Organic Thermoelectric Materials with PF Values above 50 µW m⁻¹ K⁻² or zT Values above 0.1.</caption><thead><tr><th>materials</th><th>doping/treatment</th><th>S (µV K<sup>-1</sup>)</th><th>σ (S cm<sup>-1</sup>)</th><th>PF (µW m<sup>-1</sup> K<sup>-2</sup>)</th><th>k (W m<sup>-1</sup> K<sup>-1</sup>)</th><th>zT<sup>a</sup></th><th>ref</th></tr></thead><tbody><tr><td colspan="8" style="text-align:center; font-weight:bold;">P-type</td></tr><tr><td>PA</td><td>I<sub>2</sub></td><td>28.4</td><td>11110</td><td>900</td><td></td><td></td><td>74</td></tr><tr><td>PA</td><td>FeCl<sub>3</sub></td><td>15.3</td><td>7530</td><td>180</td><td></td><td></td><td>74</td></tr><tr><td>PEDOT:Tos</td><td>TDAE</td><td>200</td><td>80</td><td>324</td><td>0.34</td><td>0.25</td><td>10</td></tr><tr><td>PEDOT:Tos</td><td>DMF</td><td>55</td><td>1500</td><td>454</td><td></td><td></td><td>75</td></tr><tr><td>PEDOT-Tos-PP</td><td></td><td>117</td><td>923</td><td>1270</td><td></td><td></td><td>76</td></tr><tr><td>PEDOT</td><td>BTFMSI</td><td>~37</td><td>~1070</td><td>147</td><td>0.19</td><td>0.22</td><td>77</td></tr><tr><td>PEDOS-C6</td><td></td><td>103</td><td>335</td><td>354.7</td><td></td><td></td><td>78</td></tr><tr><td>PEDOT:PSS</td><td>EG</td><td>62</td><td>900</td><td>350</td><td>0.37</td><td>0.28</td><td>79</td></tr><tr><td>PEDOT:PSS</td><td>EG/DMSO</td><td>73</td><td>880</td><td>469</td><td>0.33</td><td>0.42</td><td>79</td></tr><tr><td>PEDOT:PSS</td><td>TSA/DMSO N<sub>2</sub>H<sub>4</sub>/DMSO</td><td>49.3</td><td>1310</td><td>318.4</td><td>0.3</td><td>0.31</td><td>80</td></tr><tr><td>PEDOT:PSS/Te nanorod (NR)</td><td></td><td>163</td><td>19.3</td><td>70.9</td><td>0.22–0.30</td><td>0.1</td><td>81</td></tr><tr><td>PEDOT:PSS</td><td>HNO<sub>3</sub>/PEIE</td><td>32</td><td>2700</td><td>168</td><td></td><td></td><td>82</td></tr><tr><td>PEDOT/Bi<sub>2</sub>Te<sub>3</sub></td><td></td><td>~165</td><td>~500</td><td>1350</td><td>0.7</td><td>0.58</td><td>83</td></tr><tr><td>P3HT</td><td>TCB</td><td>~50</td><td>320</td><td>62.4</td><td>0.23</td><td>~0.1 (365 K)</td><td>84</td></tr><tr><td>P3HT</td><td></td><td>56</td><td>509</td><td>160</td><td></td><td></td><td>85</td></tr><tr><td>P(EtOPV-co-PV)</td><td>Mo(tfdCOCF<sub>3</sub>)<sub>3</sub></td><td>47</td><td>349</td><td>78</td><td></td><td></td><td>86</td></tr><tr><td>(BTBT)<sub>2</sub>AsF<sub>6</sub></td><td></td><td>15</td><td>4100</td><td>88</td><td></td><td></td><td>87</td></tr><tr><td>PBTTT</td><td>F<sub>4</sub>TCNQ</td><td>42</td><td>670</td><td>120</td><td></td><td></td><td>88</td></tr><tr><td>PBTTT</td><td>FeCl<sub>3</sub></td><td>9.4</td><td>2 10<sup>5</sup></td><td>~2000</td><td></td><td></td><td>89</td></tr><tr><td>PANi/Te NR</td><td></td><td>102</td><td>102</td><td>105</td><td>0.21</td><td>0.156</td><td>90</td></tr><tr><td>PDPP-g<sub>3</sub>2T<sub>0.3</sub></td><td>FeCl<sub>3</sub></td><td>56</td><td>360</td><td>110</td><td></td><td></td><td>91</td></tr><tr><td colspan="8" style="text-align:center; font-weight:bold;">N-Type</td></tr><tr><td>poly[K<sub>x</sub>(Ni-ett)]</td><td></td><td>~151.7</td><td>64</td><td>147</td><td>0.31</td><td>0.20 (440 K)</td><td>92</td></tr><tr><td>poly[K<sub>x</sub>(Ni-ett)]</td><td>electrochemical</td><td>~150</td><td>~310</td><td>~670</td><td>0.84</td><td>0.32 (400 K)</td><td>93</td></tr><tr><td>Ni NR/PVDF</td><td></td><td>~27</td><td>3000</td><td>220</td><td>0.55</td><td>0.15 (380 K)</td><td>94</td></tr><tr><td>PTEG-2</td><td>N-DMBI</td><td>~250</td><td>~13</td><td>~80</td><td>&lt;0.1</td><td>0.3 (393 K)</td><td>95</td></tr><tr><td>UFBDPPV</td><td>TAM</td><td>~198</td><td>22.5</td><td>80</td><td></td><td></td><td>96</td></tr><tr><td>FBDPPV</td><td>TAM</td><td>~32</td><td>21</td><td>51</td><td></td><td></td><td>97</td></tr><tr><td>TDPPQ</td><td>Bi</td><td>~585</td><td>3.3</td><td>113</td><td></td><td>0.23 (333 K)</td><td>98</td></tr><tr><td>A-DCV-DPPTT</td><td>N-DMBI</td><td>~666</td><td>4.9</td><td>217</td><td>0.34</td><td>0.23 (373 K)</td><td>99</td></tr></tbody></table>

when the Seebeck coefficient of undoped trans-PA was measured (equal to 900 ± 50 µV K⁻¹). Indeed, the possibility of reaching electrical conductivities exceeding 45000 S cm⁻¹ in heavily doped PA films represents an interesting opportunity for the development of a new family of thermoelectric materials.¹¹⁴,¹¹⁵ The reported maximum PF of PA is as high as 900 µW m⁻¹ K⁻² when doped with iodine and 180 µW m⁻¹ K⁻² when doped with FeCl₃.¹⁷⁴,¹¹⁶ However, the application of PA as a thermoelectric material has been limited by its insolubility and poor stability in air and even in an inert atmosphere.

A major breakthrough in organic thermoelectrics was made in 2011 by Crispin et al.,¹⁰ in which PEDOT-Tos was chemically polymerized by using iron(III) tris-*p*-toluenesulfonate as the oxidant; the resulting film had a conductivity as high as 300 S cm⁻¹ and a moderate Seebeck coefficient (40 µV K⁻¹) at room temperature. The oxidation state of the polymer was altered through reduction with strong amines such as tetrakis-(dimethylamino)ethylene (TDAE) in an inert atmosphere, yielding PF values that ranged from 38 µW m⁻¹ K⁻² for the pristine doped polymer to 324 µW m⁻¹ K⁻² at an optimum oxidation level of 22% (Figure 10a). With a lateral thermal conductivity of 0.37 ± 0.07 W m⁻¹ K⁻¹ measured through the 3ω technique, the zT value of PEDOT-Tos reached 0.25 (Figure 10b). Furthermore, the PEDOT-Tos films were stable in air at the optimum oxidation level. A vertically structured prototype of a completely organic TEG with 55 legs was made, in which the n-legs were composed of a blend of TTF-TCNQ in PVC, and the

p-legs were vapor-polymerized PEDOT-Tos. At a ΔT of 10 K, the extrapolated power density was 0.128 µW cm⁻². Recently, Villalva et al. investigated poly(ethyleneimine)-ethoxylate (PEIE) as work function modification agent for nitric acid (HNO₃) doped PEDOT:PSS as well as an encapsulation agent, preventing HNO₃ evaporation from the PEDOT:PSS film:⁸² indeed the presence of PEIE contributed to avoid the drop of electrical conductivity after air exposure by forming a barrier preventing the escape of nitric acid from the PEDOT film. In addition, PEIE ethoxy groups enhance charge carrier extraction through the modification of PEDOT work function, which, together with the lower ordering in the PEIE-coated films, might be responsible for the increase in S: the result is an increase in PF from 72 to 168 µW m⁻¹ K⁻². Other examples of doping/ dedoping processes, especially regarding PEDOT:PSS, to optimize both S and σ can be found in literature achieving high PF (>100 µW m⁻¹ K⁻²), showing that the tuning of the doping level is a crucial strategy to maximize the TE properties of organic semiconductors.⁷⁶⁻⁸⁰

In addition to chemical doping, the oxidation level of conducting polymers can be easily tuned by electrochemical doping/dedoping processes.¹¹⁷ This fine-tuning of the electronic properties of a polymer allows theoretical studies to be carried out to obtain a better understanding of the doping effects from a molecular level, which can guide the future development of OTE materials.¹¹⁸ We send the reader to other reviews covering all of this progress in detail.¹¹⁹⁻¹²³

<footnote>ªThe temperature for the summarized PF or zT values is room temperature unless otherwise stated.</footnote>
<figure><img src="image_20.png" /><figcaption>**Figure 10.** (a) Electrical conductivity, Seebeck coefficient, and *PF* and (b) *zT* of PEDOT:PSS at different oxidation levels after TDAE treatment. (a–b) Reproduced with permission from ref 10. Copyright 2011 Nature. (c) Seebeck coefficient vs electrical conductivity of different PEDOT derivatives and (d) Temperature dependence of electrical conductivity (normalized to 300 K) for the different PEDOT-Tos and PEDOT-PSS samples. (c–d) Reproduced with permission from ref 75. Copyright 2014 Nature. (e) Scheme of the morphology of TCB-treated P3HT films, and *PF* of TCB-treated P3HT films vs (f) the measurement direction and (g) temperature. (e–g) Reproduced with permission from ref 10, 75, and 84. Copyright 2016 Nature.</figcaption></figure>

**4.1.1.2. Morphology Control.** Upon solvent evaporation and film formation, conducting polymers do not form single crystal structures, but instead form amorphous, paracrystalline, or semicrystalline structures. Charge carrier transport, dominated

by interchain coupling, is therefore dramatically affected by the morphology of the film. Moreover, chemical defects, including twisted and kinked chains, are responsible for decreasing intrachain charge carrier mobility. Hence, the higher purity
and crystallinity are expected to improve the TE properties. Regarding PEDOT:PSS, for example, treatment with high boiling point solvents, such as dimethyl sulfoxide (DMSO) and ethylene glycol (EG), results in greatly improved electrical conductivities and $zT > 0.1$, with record values of $zT$ estimated to be 0.42 at room temperature (note: thermal conductivity was measured in the out-of-plane direction, while $\sigma$ and $S$ were measured in the in-plane direction).⁷⁹ The use of these solvents, in a process that is typically referred to as secondary doping, enhances the conductivity of the PEDOT:PSS film by removing the excess PSS and assisting the formation of an interconnected three-dimensional PEDOT network.

This morphology optimization of PEDOT-based conductors leads to the observation that the highly doped conducting polymer shows semimetallic character.⁷⁵ As shown in Figure 10c, different PEDOT samples with conductivities ranging from $10^{-2}$ to $10^3$ S cm$^{-1}$ and their corresponding Seebeck coefficients were investigated. While keeping the oxidation/doping level constant, the morphology is tuned from amorphous to semicrystalline, thus simultaneously enhancing the Seebeck coefficient and electrical conductivity. The PEDOT-Tos sample obtained by vapor-phase polymerization using a triblock copolymer of poly(ethylene glycol-propylene glycol-ethylene glycol) (PEG-PPG-PEG, 58000 Mw + DMF) as a template shows both the highest conductivity and Seebeck coefficient. The explanation for the improvement in both the Seebeck coefficient and conductivity is proposed as a transition from a Fermi glass to a semimetal; this theory is supported by the temperature dependence of conductivity (Figure 10d). The semimetallic character that originates from the formation of a network of bipolarons results in an empty bipolaron band that slightly overlaps with the valence band.

Morphology control has been used to enhance the thermoelectric properties of other doped conjugated polymers. In 2016, Chen et al. reported highly ordered poly(3-hexylthiophene-2,5-diyl) (P3HT) films prepared by the epitaxial growth method, in which the solvent, 1,3,5-trichlorobenzene (TCB), was used as the template.⁸⁴ The molecules in the quasi-one-dimensional (1D) P3HT film were highly oriented parallel to the fiber axis direction (as shown in the scheme in Figure 10e) with a decreased number of defects and enhanced carrier mobility; thus, the conductivity (parallel to the fiber axis) was markedly improved, reaching a maximum value of 320 S cm$^{-1}$. With this enhancement in electrical conductivity, maximum PF and $zT$ values of 62.4 $\mu$W m$^{-1}$ K$^{-2}$ and 0.1, respectively, were reached at 365 K in the direction parallel to the fiber axis (Figure 10f,g). This result is one of the highest values reported for pure P3HT materials. The same strategy of P3HT fiber alignment was also applied by Untilova et al. in 2020 by using high temperature rubbing,¹²⁴ a technique enabling the preparation of highly aligned polymer films with a thickness of a few tens of nanometers:⁸⁵ the alignment of the fibers, coupled with the molybdenum dithiolene complex Mo(tfdCOCF$_3$)$_3$ as doping specie, allowed reaching an impressive PF of 160 $\mu$W m$^{-1}$ K$^{-2}$.

The effect of morphology control on P3HT has been systematically studied by Lim et al. in 2019:¹²⁵ by vapor doping with $F_4TCNQ$, they were able to control the morphologies of the doped P3HT films. They observed that crystalline regions have higher doping efficiency compared to amorphous ones, which, coupled with higher charge mobility, leads to higher values of $\sigma$. On the other hand, $S$ appeared to be less affected by the morphology while instead being strongly influenced by structural ordering, in agreement with other research group

results,¹²⁶,¹²⁷ while also extending what Patel et al. previously observed with poly(2,5-bis(3-tetradecylthiophen-2-yl)thieno-[3,2-b]thiophene) (PBTTT) vapor-doped with $F_4TCNQ$, where a PF as high as 120 $\mu$W m$^{-1}$ K$^{-2}$ was achieved.¹²⁸

Highly ordered PANi chains have also been reported recently, with enhancements in electrical conductivity as well as the Seebeck coefficient by a factor of 3.3 and 1.2, respectively.¹²⁹ The use of a self-assembled supramolecule (SAS) like (3,6-dioctyl-decyloxy-1,4-benzenedicarboxylic acid) as a template has proven to be an effective way to obtain highly ordered PANi-SAS films, reaching a maximum PF of 31 $\mu$W m$^{-1}$ K$^{-2}$.¹²⁹ The importance of structural ordering has also been shown by Brinkmann et al., who recently reported a method to obtain highly oriented PBTTT films with high persistence length, capable of reaching exceptional electrical conductivities larger than $10^5$ S cm$^{-1}$ and corresponding PF of 2 mW m$^{-1}$ K$^{-2}$.⁸⁹

Thin-film nanostructuring has emerged as an alternative way to further enhance the thermoelectric properties of conducting polymers. For example, Sun et al. demonstrated a three-in-one improvement in the PF of PANi when the material is organized as nanotubes.¹³³ The power factor of the as-synthesized $\beta$-naphthalene sulfonic acid-doped polyaniline nanotubes was 2 orders of magnitude higher than that of pristine bulk PANi films. The reason for this improvement was ascribed to a sharper DOS, which leads to an increase in $S$ by a factor of 7, an increase in charge mobility alongside the nanotube axis (i.e., increase in $\sigma$) and higher phonon scattering (i.e., lower lattice component of the thermal conductivity). Subsequent studies on PPy nanotube films and the nanostructures of other conducting polymers have also been reported.¹³⁴ Conducting multilayered polymer structures with enhanced thermoelectric performance was published by Jo et al.¹³⁵ A 5-layer PEDOT:PSS/PANI-CSA film with a thickness of 20 nm was prepared through layer-by-layer deposition. The PF of the multilayer film reached 49 $\mu$W m$^{-1}$ K$^{-2}$, 2 times higher than that of the single PEDOT:PSS layer.

Another interesting strategy to control the morphology and mechanical properties of TE conducting polymer materials consists of realizing composites of conducting polymers and insulators, which can contribute to improving the material processability. For example, Yang et al. investigated the thermoelectric properties of poly(3-butylthiophene-2,5-diyl) (P3BT) nanowire interpenetrating networks in an insulating polystyrene (PS) matrix.¹³⁶ Because of the quasi-1D charge transport in this bulk interpenetrated network, the P3BT/PS blend (40-60% P3BT) showed a higher electrical conductivity (1 mS cm$^{-1}$) and lower thermal conductivity (0.2 W m$^{-1}$K$^{-1}$) than neat P3BT (0.1 mS cm$^{-1}$, 0.3 W m$^{-1}$K$^{-1}$, respectively), without sacrificing the Seebeck coefficient (600 $\mu$V K$^{-1}$). The PF of the blend was approximately 1 order of magnitude higher than that of the neat P3BT. Although the overall performance was not a record, it demonstrated the possibility of improving thermoelectric properties by blending conducting polymers with insulating matrices.

The use of an insulating polymer as a matrix is also a possible solution to improve the stability of the TE properties of organic semiconductors under mechanical stress. In 2019, Baran and coworkers demonstrated in two different works the possibility of achieving self-healing and stretchable PEDOT:PSS based composites with the addition of a polymeric surfactant (Triton X-100) and ionic liquids (1-ethyl-3-methylimidazolium cations with 4,5-dicyanoimidazolate and dicyanamide containing nitrogen anions), respectively.¹³⁰,¹³¹ The presence of Triton X-100
<figure><img src="image_21.png" /><figcaption>**Figure 11.** (a) SEM image of the cut on the PEDOT:PSS/Triton X-100 composite film and (b) Schematic illustration for the mechanical and thermoelectrical self-healing process of the composite film. (a–b) Reproduced with permission from ref 130. Copyright 2019 Wiley. (c) Schematic of the IL-induced morphological transition of PEDOT:PSS from the amorphous structure into crystalline nanofibrils inserted in IL-assisted soft domains. Reproduced with permission from ref 131. Copyright 2019 American Chemical Society. (d) Schematic of the stretchability of the PEDOT:PSS-polyurethane-IL composite. (e) Photograph of a stretchable TE module composed of 10 legs. (f) $V_{OC}$ and $R_{in}$ vs tensile strain ($ε$) applied to a module at $ΔT = 30$ K. (g) $P_{out}$ vs $R_{load}$ under different strains at $ΔT = 30$ K. (d–g) Reproduced with permission from ref 132. Copyright 2020 Nature.</figcaption></figure>

creates a viscoelastic composite in which the PEDOT:PSS is dispersed in the surfactant matrix which acts as a viscous fluid and thus after a cut, it can freely flow from both sides of the separated regions and form a physical contact, effectively mending the “wound” (Figure 11a,b), thus allowing the composite to recover its thermoelectric and electrical properties. On the other hand, the use of the ionic liquids (IL) allowed increasing at the same time both S and $σ$ of PEDOT:PSS, by modifying the doping level and nanostructure of PEDOT:PSS via an ion-exchange process. In addition, the presence of the IL resulted in the formation of PEDOT:PSS nanofibrils embedded in a soft matrix, improving its mechanical stretchability (Figure 11c).

Recently, Kim et al. combined the two previously discussed strategies by reporting an aqueous-processed poly(3,4-ethyl-enedioxythiophene)-polyurethane-ionic liquid composite, having both high conductivity (>140 S cm⁻¹) and stretchability (>600%),¹³² as shown in Figure 11d. The addition of water-borne polyurethane and of an ionic liquid (tricyanomethanide) not only increases tremendously the elasticity of the composite but also helps maintaining high values of electrical conductivity. The high conductivity of the composite is due to a smoother morphology of the film, possessing a roughness of around 42 nm because of the counterion exchange between the conducting

polymer and IL, which causes a morphological change in the PEDOT:PSS toward a crystalline nanofibrillar structure, which is beneficial not only to charge transport along the polymer chains but also for the formation of a percolating network within the polyurethane matrix. Finally, they realized a TEG composed of 10 thermocouples (Figure 11e), showing a constant open circuit voltage under different strain levels while keeping the variation in electrical resistance below 50%, even under high mechanical stresses (Figure 11f,g).

Apart from blending organic semiconductors with polymeric electrolytes/matrices or IL, other methods to obtain flexible organic thermoelectric have been reported: Kim et al. reported the wet-spinning into sulfuric acid of PEDOT:PSS, possessing a Young's modules of 1.9 GPa and a $σ > 800$ S cm⁻¹, that can be integrated into textiles.¹³⁷ Indeed, in a following work, they show the possibility of integrating a TEG with a vertical architecture into the sewing thread of a wool fabric, demonstrating the possibility of achieving a power of 1.2 µW with a $ΔT = 65$ K.¹³⁸

#### 4.1.1.3. *Hybrid Composites.*
Organic–inorganic composites present a unique way to optimize the $zT$ value of thermoelectric materials through proper molecular and interfacial designs. In 2010, a high-performance water-processable polymer–inorganic hybrid TE material with a room temperature $zT$ value of
<figure><img src="image_22.png" /><figcaption>**Figure 12.** (a) Cartoon depiction of Cu incorporation and the nucleation of Cu-Te alloy phases within PEDOT:PSS-Te NWs, and TEM images of (b) PEDOT:PSS-Te and (c) PEDOT:PSS-Te(Cux) NWs showing a transition from straight to bent wires. (a–c) Reproduced with permission from ref 140. Copyright 2016 American Chemical Society. (d) Scanning electron microscopy (SEM) image of a PEDOT/Bi₂Te₃(100) hybrid film. (e) Power factor of the PEDOT:PSS-Bi₂Te₃ composite as function of nanoparticle fraction, and (f) Change in the electrical resistance of the PEDOT:PSS-Bi₂Te₃ composite as a function of the number of bending cycles. (d–f) Reproduced with permission from ref 83. Copyright 2018 Nature.</figcaption></figure>

approximately 0.1 was reported.⁸¹ A smooth PEDOT:PSS/Te nanorod composite film, prepared by solution casting, had both high electrical conductivity and a high Seebeck coefficient at room temperature. The in-plane thermal conductivity of the composite film, characterized by means of the 3-ω method, ranged from 0.22 to 0.30 W m⁻¹ K⁻¹. This combination yielded excellent PF (~70 µW m⁻¹ K⁻²) and zT (~0.1) values at room temperature, demonstrating a record for organic composites. The same authors investigated the origin of the high performance of these hybrid materials.¹³⁹ By varying the nanowire:polymer ratio, a maximum conductivity (11 S cm⁻¹) was found at an intermediate content of Te nanowires (0.86 weight fraction). They also developed a series-connected model to explain the behavior of the composite, which emphasized the importance of understanding and controlling the interfacial volume in this organic–inorganic system to enhance carrier transport in hybrid materials and devices.

In 2016, Zaia et al. developed a method to control the growth of heterostructures within hybrid organic/inorganic nano-composites.¹⁴⁰ Nanoscale domains of a Cu₁.₇₅Te alloy at different doping concentrations were grown by adding different amounts of copper(II) nitrate hemipentahydrate to a dispersion of PEDOT:PSS-Te NWs (Figure 12a). A 22% enhancement in PF was achieved in the low-doping regime, with PF values reaching up to 84 µW m⁻¹ K⁻² (Figure 12c). This approach could be adapted to other systems for the rational design of solution-processable TE materials. Combining transport measurements with theoretical simulations, Kumar et al. further studied the thermoelectric properties of the PEDOT:PSS-Te(Cux) nanowire hybrid material system.⁸⁸ They proposed that the thermoelectric properties were dominated by the polymer matrix and that the positive effect of the inorganic NWs was due to the templating effect of the PEDOT:PSS morphology

at the interface with the Te(Cux) NWs as well as the charge redistribution between the two conductors.

Very impressive results have also been obtained by combining the two most iconic organic/inorganic thermoelectric materials: PEDOT:PSS and bismuth telluride. In 2018, flexible PEDOT/Bi₂Te₃ hybrid films with monodispersed and periodic Bi₂Te₃ nanophases (Figure 12d) were fabricated by Wang et al.⁸³ Periodic Bi₂Te₃ nanofillers could significantly scatter phonons, resulting in an ultralow in-plane thermal conductivity of ~0.7 W m⁻¹ K⁻¹, while the continuous PEDOT phase provided favored electronic transport with an ultrahigh power factor of ~1350 µW m⁻¹ K⁻² (Figure 12e). Consequently, a high zT value of 0.58 was obtained at room temperature. In addition, the as-fabricated hybrids showed superior mechanical flexibility, and their thermoelectric properties showed negligible changes after 100 bending cycles (Figure 12f). Katz et al. reported another PEDOT:PSS-Bi₂Te₃ composite with an enhanced PF.¹⁴¹ Both n- and p-type Bi₂Te₃ particles were ball-milled, dispersed in alcohol, and drop cast on a glass slide to form a thin layer. This step was followed by drop casting PEDOT:PSS on top of the Bi₂Te₃ particle layer. By controlling the load of PEDOT:PSS, maximum PFs of 70 and 40 µW m⁻¹ K⁻² were achieved for p- and n-type composite materials, respectively. This method facilitated the development of polymer–inorganic composite materials, and it overcame the organic–inorganic miscibility issues typical of these hybrid materials.

Similarly, P3HT–inorganic composites have also been studied as promising TE materials.¹⁴² Bi₂Te₃ was added to a solution of FeCl₃-doped P3HT and mixed via ultrasonication. A 15 µm thick composite film was then obtained by drop casting and vacuum drying the films at room temperature. The authors suggested an energy-filtering effect at the organic–inorganic semiconductor interface to explain why the P3HT-Bi₂Te₃
<figure><img src="image_23.png" /><figcaption>**Figure 13.** (a) Synthetic route for the synthesis of P(MeOPV-co-PV) and (b) Temperature dependence of the PF of camphorsulfonic acid-doped polyaniline (PANi) and iodine-doped P(MeOPV-co-PV) at a stretching ratio of 4.4. (a–b) Reproduced with permission from ref 146. Copyright 2006 Elsevier Ltd. (c) Illustration of DOS engineering to achieve a high Seebeck coefficient. Dependence of the (d) Seebeck coefficient and (e) PF on the HOMO-HOMO offset between the two polymer components and the offset between the HOMO of polymer A and the LUMO of the dopant. Closed symbols correspond to the P3HT:PTB7 (circles) and P3HT:TQ1 blends (squares). (c–e) Reproduced with permission from ref 149. Copyright 2018 Wiley.</figcaption></figure>

nanocomposite exhibited a PF (13.6 µW m⁻¹ K⁻²) more than 3 times higher than that of pure P3HT (3.9 µW m⁻¹ K⁻²). The authors suggested that this effect was due to lower energy carriers scattering at the interface between P3HT/Bi₂Te₃, contributing to increase S without significantly affecting σ.

4.1.1.4. *Chemical Structure Design.* In addition to the aforementioned strategies, modification of the chemical structures of the conducting polymer backbone or side chains can also affect the TE properties of organic semiconductors. For example, Leclerc et al. reported the thermoelectric properties of a series of polycarbazole derivatives.¹⁴³ By varying the side chains both on the carbazole cycle and on the nitrogen atom of the backbone unit, two poly(diindolocarbazole)s comprised of thiophene units were synthesized, showing the best combination of the Seebeck coefficient (71 µV K⁻¹) and conductivity (0.29 S cm⁻¹) but with a modest PF of only 0.15 µW m⁻¹ K⁻². Later, in 2009, three different carbazole-based copolymers, namely, poly[N-9'-heptadecanyl-2,7-carbazole-alt-5,5-(2,2'-bithiophene)] (PCDT), poly[N-9'-heptadecanyl-2,7-carbazole-alt-5,5'-(2,2'-(1,4-phenylene)dithiophene)] (PCDTB), and poly-[N-9'-heptadecanyl-2,7-carbazole-alt-5,5'-(4',7'-di-2-thienyl-

2',1',3'-benzothiadiazole] (PCDTBT), were synthesized through Suzuki coupling polymerization.¹⁴⁴ After optimization, the FeCl₃-doped PCDTBT films showed a relatively high electrical conductivity (160 S cm⁻¹) and Seebeck coefficient (34 µV K⁻¹), which resulted in a PF of 19 µW m⁻¹ K⁻².

Another interesting candidate for TE application is alkoxy-substituted poly(*p*-phenylene vinylene) (PPV), showing high conductivity values >200 S cm⁻¹ when doped with I₂.¹⁴⁵ In 2006, Hiroshige et al. first evaluated the thermoelectric properties of poly(2,5-dimethoxyphenylenevinylene) and its derivatives.¹⁴⁶ Stretched copolymers consisting of both 2,5-dimethoxy-substituted phenylenevinylene units and unsubstituted units (P(MeOPV-co-PV)) exhibited an electrical conductivity of 184 S cm⁻¹ and a Seebeck coefficient of 44 µV K⁻¹ at 313 K upon I₂ doping, resulting in a PF larger than 30 µW m⁻¹ K⁻², as shown in Figure 13b. The same authors also investigated the effect of different alkoxy substituents and stretching treatments on the TE properties of copolymers consisting of both 2,5-dialkoxy-substituted PPV units and unsubstituted units (P(ROPV-co-PV)) (RO = MeO, EtO, and BuO).⁸⁶ Among all the copolymers, I₂-doped P(EtOPV-co-PV)
<figure><img src="image_24.png" /><figcaption>**Figure 14.** (a) Molecular structures of P(gNDI-gT2) and AFM images of (b) pristine and (c) 30% doped σ p(gNDI-gT2). (a–c) Reproduced with permission from ref 158. Copyright 2018 American Chemical Society. Representation of coarse-grained molecular dynamics simulations of n-DMBI molecules dissolved in (d) pure N2200 side chain phase and (e) pure TEG-N2200 side chain phase. (f) Normalized number of contacts between dopant molecules in N2200 and TEG-N2200 at different concentrations. (d–f) Reproduced with permission from ref 159. Copyright 2018 Wiley. (g) Calculated spin density distribution for BBL and P(NDI2OD-T2). Reproduced with permission from ref 160. Copyright 2016 Wiley. (h) Effect of the annealing temperature on the secondary doping of BBL using N-DMBI. Reproduced with permission from ref 161. Copyright 2020 American Chemical Society.</figcaption></figure>

exhibited the best overall performance at a stretching ratio of 3.1, with a conductivity of 349.2 S cm⁻¹, Seebeck coefficient of 47.3 µV K⁻¹ and thermal conductivity of 0.25 W m⁻¹ K⁻¹ at 313 K. The corresponding PF of 78.1 µW m⁻¹ K⁻² and high zT value near 0.1 were recorded at the time for conducting polymers.

An interesting category for thermoelectric application is represented by random copolymers: they can be considered as irregular polymers with structural disorder along the backbone. This disordered structure can be favorable to adjust the energy level and the electrical properties to potentially enhance TE behaviors.<sup>147,148</sup> Therefore, Li et al. designed and synthesized a series of new random copolymers by copolymerizing D–A type and D–D type building blocks along polymer chains with the aim of improving the TE and electrical properties:<sup>91</sup> the tuning of the energy levels resulted in an increased doping level while strong D–A interchain interactions allowed obtaining a high degree of crystallinity for efficient charge transport. The result is an impressive PF of 110 µW m⁻¹ K⁻², due to an impressive charge mobility 1 cm² V⁻¹ s⁻¹ which leads to a σ > 300 S cm⁻¹ without affecting the S.

4.1.1.5. DOS Engineering. In Mahan and Sofo's pioneering work, a Δ-shaped transport distribution above the Fermi level was demonstrated to maximize the thermoelectric properties.<sup>150</sup> On the basis of this idea, Sun et al. used F4TCNQ as a p-type dopant in a mixture of P3HT and different fractions of poly(3-hexylthiophene-thiophene) (P3HTT). It was found that the Seebeck coefficient of P3HT mixed with 2% P3HTT increased from ≈580 µV K⁻¹ for pristine P3HT to 700 µV K⁻¹, while its conductivity remained at approximately 1.5 × 10⁻⁴ S cm⁻¹.<sup>151</sup> Li et al.<sup>152</sup> expanded on these results by offering a method to properly select the polymer blend composition to maximize PF. On the basis of these results, Kemerink et al.<sup>149</sup> reported a universal method to engineer the charge carrier DOS to maximize thermoelectric performance (explained in Figure 13d). A few conjugated polymers, namely, P3HT, poly({4,8-bis[(2-ethylhexyl) oxy] benzo[1,2-b:4,5-b'] dithiophene-2,6-diyl}{3-fluoro-2-[(2-ethylhexyl) carbonyl]thieno[3,4-b] thiophenediyl}) (PTB7), and poly[2,3-bis(3-octyloxyphenyl)-5,8-quinoxalinediyl]-2,5-thiophenediyl] (TQ1), were selected because they had different HOMO energy levels. P3HT was mixed with varying fractions of PTB7 and TQ1 and then doped
with F4TCNQ. They found that at the optimized mixing ratio, the Seebeck coefficient of the mixture largely exceeded that of any of the pure materials (1100 µV K⁻¹ at P3HT₀.₁:PTB7₀.₉₀ and 2000 µV K⁻¹ at P3HT₀.₀₅:TQ1₀.₉₅ vs 142 µV K⁻¹, 469 µV K⁻¹, and 1560 µV K⁻¹ for pure P3HT, PTB7, and TQ1, respectively), as shown in Figure 13d,e. These observed trends in the Seebeck coefficient were in good agreement with kinetic Monte Carlo simulations with parameters based on experimental data. In a following work, Kemerink et al.¹⁵³ reviewed the thermoelectric properties of conjugated blends, trying to review all the results obtained offering guidelines to design organic thermoelectric materials.

### 4.1.2. n-Type TE Polymers and Composites.
Regarding the fabrication of practical thermoelectric devices, both electron and hole conducting materials with large PF values are needed. Compared to their p-type counterparts, n-type conducting polymers have shown inferior thermoelectric properties, mainly due to their less effective and stable doping mechanisms. Conventional conducting polymers such as PA and PANi can behave as both p-type and n-type conductors when doped appropriately.¹⁵⁴,¹⁵⁵ However, because they are typically unstable in the n-type state under ambient conditions, their TE performance is insufficient. Although increasing examples of n-type conducting polymers with improved TE properties have been reported, more effort is needed to better understand the relationship between the chemical structure and doping of n-type TE materials to develop efficient and stable TE devices. In the following, we provide a few examples of state-of-the-art n-type TE materials.

In 2014, Chabinyc et al. reported the TE properties of the n-type polymer poly-N,N-bis(2-octyl-dodecyl)-1,4,5,8-naphthalenedicarboximide-2,6-diyl]-alt-5,5-(2,2-bithiophene) (P-(NDI2OD-T2)¹⁵⁶ doped in solution with an air-stable dopant (4-(1,3-dimethyl-2,3-dihydro-1H-benzoimidazol-2-yl)phenyl) (N-DMBI) before film formation.¹⁵⁷ The electrical conductivity of the P(NDI2OD-T2) films could be varied over several orders of magnitude by tuning the N-DMBI molar concentration. The electrical conductivity, however, reached an unsatisfactory maximum value of only 5 × 10⁻³ S cm⁻¹, which was attributed to the limited solubility/miscibility of the dopant in the polymeric matrix. The authors suggested that only ~1% of the DMBI molecules introduced in solution were active and contributed free carriers to the film, implying that only 1 in 1000 P(NDI2OD-T2) repeat units were doped.

This extrinsic limitation could be overcome by proper engineering of the polymer backbone and/or alkyl side chains. In two parallel studies, Müller et al.¹⁵⁸ and Koster et al.¹⁵⁹ showed that functionalizing the NDI-based polymer with polar side chains (Figure 14a), such as oligoethylene glycol-based side chains, led to a high electrical conductivity of more than 10⁻¹ S cm⁻¹ for an N-DMBI due to a low segregation of the dopant molecules (Figure 14b–f). Overall, the n-type polymers resulted in a larger thermoelectric PF of up to 0.4 µW m⁻¹ K⁻² compared to other NDI-based polymers functionalized with alkyl side chains. Caironi et al. reported that, compared to nonthionated P(NDI2OD-T2), the thionated P(NDI2OD-T2) gave rise to higher electrical conductivity and stable operation in air due to a deeper LUMO level, lying below the threshold for air stability at ~4 eV.¹⁶²,¹⁶³ Nevertheless, the thionated polymer showed a relatively low PF (0.04 µW m⁻¹ K⁻²). The same authors reported that the incorporation of “kinked” comonomers in NDI-based polymers enhanced dopant/polymer miscibility, allowing for a ~20-fold increase in doping efficiency.¹⁶⁴

Fabiano et al. argued that the low electron conductivity typically observed for n-type P(NDI2OD-T2) might be ascribed to the strong charge carrier intrachain localization,¹⁶⁰ which is a common characteristic of donor–acceptor polymers.¹⁶⁵ After replacing the donor bithiophene (T2) unit in P(NDI2OD-T2) with the more electron-deficient bithiazole (Tz2) unit (i.e., P(NDI2OD-Tz2)), they observed that the polymer donor–acceptor character decreased and its electron affinity increased, which resulted in an improved polymer doping efficiency. Moreover, because of the far lower intrachain steric demand of Tz2 than T2, the polymer backbone became more planar, thus enhancing the intermolecular π–π stacking interactions, which led to a higher electron conductivity (up to ~0.1 S cm⁻¹) and increased the thermoelectric response (PF up to 1.5 µW m⁻¹ K⁻²) compared to P(NDI2OD-T2).

Pei et al. synthesized a series of benzodifurandione-based PPV (BDPPV) samples with high electron affinities (LUMO levels deeper than −4.0 eV) and investigated their electrical conductivity, Seebeck coefficient, and PF values.¹⁶⁶ The LUMO levels of Cl-substituted BDPPV (CIBDPPV) and F-substituted BDPPV (FBDPPV) were 0.29 and 0.16 eV farther from vacuum than that of BDPPV, reaching values of −4.30 and −4.17 eV, respectively. The electrical conductivity of the N-DMBI-doped films increased dramatically with dopant concentrations lower than 10 wt %, reaching values as high as 14 S cm⁻¹. The electrical conductivity then decreased as the N-DMBI content increased above 10% in the blend. The Seebeck coefficient remained negative for all three N-DMBI doped polymers, but the absolute value of the thermopower continuously decreased with increasing dopant concentration. Overall, FBDPPV showed the highest PF (28 µW m⁻¹ K⁻²) at the optimal doping concentration of 5 wt % at room temperature. Later, Zhao et al. demonstrated that low-cost, air-stable, and solution-processable tetrabutylammonium fluoride (TBAF) could be used as an effective n-dopant for CIBDPPV to boost the electrical conductivity of this polymer up to 0.62 S cm⁻¹.¹⁶⁷ This n-doping mechanism relied on electron transfer from the F⁻ anions to the electron-deficient polymer CIBDPPV via an anion–π electronic interaction. A PF of 0.63 µW m⁻¹ K⁻² was obtained at a content of 25 mol % TBAF. This work provides an alternative air-stable dopant for high-performance solution-processable n-type thermoelectric materials. Recently, in 2021, they expanded their results by doping donor–acceptor (D–A) polymers, with dichloro substitutions, using N-DMBI: the Cl-substituted D–A due to the low lying LUMO level (<−4 eV) and to the dense packing of the hydrophobic groups, blocking O₂ and H₂O interactions, allowed to maintain 31% of the original σ after 222 days of air exposure.¹⁶⁸

In 2016, Fabiano et al. showed that solution-processable ladder-type conducting polymers, such as polybenzimidazobenzophenanthroline (BBL), were promising n-type materials for TE applications.¹⁶⁰ After exposure to tetrakis(dimethylamino)-ethylene (TDAE) vapors, the electrical conductivity of spin-coated BBL films could be dramatically increased, reaching values as high as 2.4 S cm⁻¹, while the Seebeck coefficient reached a value of −60 µV K⁻¹. Thus, by optimizing the TDAE exposure time, the PF of the doped BBL film could be optimized to a value of 0.43 µW m⁻¹ K⁻². Although the TE performance of BBL was not extraordinary, its high electrical conductivity, originating from its highly rigid and planar backbone (Figure 14g), provided a molecular-design guideline for new TE polymers. Recently, sequential doping has proven useful to
<figure><img src="image_25.png" /><figcaption>**Figure 15.** (a) Image of the poly[Kₓ(Ni-ett)]/PVDF composite on PET substrate. (b) *zT* value versus temperature. (c) Top-view and (d) cross-section SEM images of poly[Kₓ(Ni-ett)]/PVDF films. Reproduced with permission from ref 93. Copyright 2016 Wiley.</figcaption></figure>

improve the TE properties of BBL. First, the BBL film was spin-coated on a glass substrate and then annealed; subsequently, N-DMBI, dissolved in an orthogonal solvent (chloroform), was spin-coated on top of the polymer film. They reported a maximum PF of 1.5 µW m⁻¹K⁻² with an electrical conductivity of approximately 1 S cm⁻¹ (Figure 14h) and a Seebeck coefficient of −110 µV K⁻¹, showing an improvement by nearly a factor of 4 compared to the BBL/TDAE blend.¹⁶¹ The same authors recently reported an n-type all-polymeric ink based on a BBL:poly(ethyleneimine) (PEI) blend, having an electrical conductivity as high as 8 S cm⁻¹, and demonstrated high thermoelectric power output (56 nW per thermocouple with a temperature difference below 50K) when a BBL:PEI-based n-leg was coupled to a PEDOT:PSS-based p-leg.¹⁶⁹

Another growing interest in the field of organic thermo-electricity has been stimulated by metal-organic compounds. In 2012, Zhu et al. reported the n-type TE properties of the metal coordination polymer poly[Kₓ(Ni-ett)] (ett = ethylenetetra-thiolate) with a high zT value of 0.20 at 440 K.⁹² The authors synthesized a series of poly[Aₓ(Ni-ett)]s samples with different counterions (i.e., K⁺, Na⁺, and Bu₄N⁺). By air oxidation, poly[Aₓ(Ni-ett)] could be obtained as a precipitate. Among the different samples, poly[Kₓ(Ni-ett)] had the best thermoelectric performance, with an electrical conductivity of 64 S cm⁻¹, a Seebeck coefficient as high as −151.7 µV K⁻¹, and a thermal conductivity of only 0.31 W m⁻¹ K⁻¹ at 440 K. When the metal ligand was changed to Cu, the resulting metal-coordinated polymer showed p-type TE performance, allowing for the fabrication of modules composed of 35 p-n couples. The resulting TE module had an open-circuit voltage of 0.26 V at a temperature difference of 80 K, and its power output was estimated to be 2.8 µW cm⁻² at a temperature difference of 30 K.

Although these metal coordination polymers showed remarkably high thermoelectric performance, their insoluble and infusible characteristics hindered their processability and practical application in flexible thermoelectrics. To overcome this problem, the same authors used polymeric matrices to prepare solution processable composites.¹⁷⁰ The wet ball-milling method was employed to prepare thermoelectric composites of poly(M-ett). By optimizing the use of different polymer matrices, solvents, ball-milling times, weight ratios, and drying temperatures, maximum PF values of 0.43 µW m⁻¹ K⁻² and 0.86 µW m⁻¹ K⁻² were reached at 300 K for poly[Kₓ(Ni-ett)]/PVDF/NMP (n-type) and poly[Cuₓ(Cu-ett)]/PVDF/DMSO (p-type), respectively. To demonstrate their practical application in solution-processable thermoelectrics, a flexible device made of six inkjet-printed thermocouples was fabricated on a PET substrate.

In 2016, the same authors reported flexible, high-performance poly[Kₓ(Ni-ett)] films prepared by electrochemical deposition.⁹³ Thin-film samples, prepared by applying a potential of 0.6 V (vs Ag/AgCl), showed better performance than the aforementioned pellets. The Seebeck coefficient of the poly-[Kₓ(Ni-ett)] films was similar to that of the pellets, but its electrical conductivity (>300 S cm⁻¹) was several times higher. Therefore, even though the thermal conductivity was slightly higher than that of the pellets (0.84 W m⁻¹ K⁻¹ at 400 K for thin films vs 0.28 W m⁻¹ K⁻¹ at 400 K for pellets), the zT value of poly[Kₓ(Ni-ett)] films at 400 K could reach 0.32, a new record for n-type organic thermoelectric materials (Figure 15b). Later, Yee et al. systematically studied the reaction conditions of poly[Kₓ(Ni-ett)], such as the counterion, nickel equivalency, and oxidation extent, to reliably and reproducibly synthesize poly[Kₓ(Ni-ett)]/PVDF in large quantities of up to 15 g. After optimization, the PF of the poly[Kₓ(Ni-ett)]/PVDF composite
<figure><img src="image_26.png" /><figcaption>**Figure 16.** (a) Schematic of the crystal structure (BTBT)₂XF₆ and (b) Seebeck coefficient of (BTBT)₂XF₆ as a function of temperature. (a–b) Reproduced with permission from ref 87. Copyright 2016 American Chemical Society. (c) SEM image of TCNQ:TTT thin films with TCNQ:TTT molar ratio of 18.6 and (d) PF dependence TCNQ:TTT molar ratio. (c–d) Reproduced with permission from ref 180. Copyright 2017 Wiley. (e) Schematic of the measurement setup for the simultaneous measurements of S and σ of individual Sc₃N@C₈₀ molecules and (f) Histograms of S (green) at first contact for Sc3N@C80 compared to the values for C60 (in gray). (e–f) Reproduced with permission from ref 182. Copyright 2016 Nature. (g) σ, activation energy of conductivity (Eₐ), and S (a) of a C₆₀ thin film as a function of temperature. The sign of all plotted S values is negative. Reproduced with permission from ref 183. Copyright 2015 Japan Society of Applied Physics. (h) Schematic representation of PTEG-2 molecular packing, as atomistically resolved by molecular dynamics simulations, with the unit cell highlighted in blue. (i) zT of PTEG-2 as a function of temperature, showing the record for an n-type solution processable organic semiconductor. Reproduced with permission from ref 95. Copyright 2020 Nature.</figcaption></figure>

film improved to 23 µW m⁻¹ K⁻².¹⁷¹ Furthermore, thermal annealing was also investigated as an effective posttreatment process to simultaneously enhance the conductivity and thermopower of poly[Kₓ(Ni-ett)]/PVDF composite films.¹⁷²

Organic/inorganic composites have also been investigated for use as n-type thermoelectric materials, where the organic insulating polymers act as the polymeric matrices to effectively decrease the thermal conductivity and the inorganic semiconductors/metals function as the active thermoelectric material. For example, very recently, n-type composite thermoelectric materials with remarkable TE properties have been prepared by simply blending an insulating polymer together with metallic nanowires.⁹⁴

Drop-cast Ni-NW/PVDF films, manufactured with the Ni-NW content in the composites varying from 20 to 80 wt %, show

electrical conductivity and Seebeck coefficient increasing with the Ni-NW content and reaching maxima of 4701 S cm⁻¹ and -20.6 µV K⁻¹ at 80 wt % Ni, respectively. The temperature dependence of the conductivity of the Ni (80 wt%)/PVDF nanocomposite film showed typical metallic behavior, i.e., decreasing with increasing temperature, while the absolute value of the Seebeck coefficient showed the opposite trend. As a result, the PF of the composite film increased with temperature, from 200 µW m⁻¹ K⁻² at room temperature to 220 µW m⁻¹ K⁻² at 380 K. Combined with the low and nearly constant thermal conductivity of 0.55 W m⁻¹ K⁻¹ over the whole temperature range, a high zT value of 0.15 was achieved for the Ni (80 wt %)/PVDF composite film at 380 K.

A final mention must be made regarding the development of new dopant molecules to increase the efficiency of the doping
process in n-type organic semiconductors. Saglio et al. synthesized benzimidazoles with different linear and branched alkyl substituents and observed a clear dependence of the nature of the alkyl substituent on dopant intercalation into the semicrystalline morphology.173 It was found out that increasing the length or the steric hindrance of the alkyl substituents, the miscibility between dopant and copolymer is enhanced, leading to an increase by a factor of ≈7 in the electrical conductivity. Following this idea of optimizing the dopant specie, Yang et al. recently reported a DTF calculation-based method which led to the synthesis of a triaminomethane-type dopant (TAM), exhibiting high stability and strong hydride donating property due to its thermally activated doping mechanism. Thanks to this approach, they were able to efficiently dope fluorinated benzodifurandione-based poly(pphenylene vinylene) (FBDPPV), enabling it to reach an impressive value of n-type conductivity for an organic semiconductor (21 S cm⁻¹) and a PF of 51 µW m⁻¹ K⁻² even in films with thickness above 10 µm.⁹⁷ Finally, they reported one of the first examples of all-polymer TEG, with a power output of 77 nW at ΔT ≈ 50 K. On the basis of these impressive results, Lu et al. tried to investigate the role of molecular dopant cations on carrier transport by comparing two n-dopants, TAM and N-DMBI-H.⁹⁶ The weak π-backbone affinity of TAM cations, coupled with the spatial screening of the Coulomb potential thanks to its side chains, allows negligible perturbation of the polymer backbone conformation and microstructural ordering, thus resulting in a σ ≈ 22 S cm⁻¹ and a PF of 80 µW m⁻¹ K⁻². Apart from the impressive results, this work provided a successful strategy, applicable to other polymers, to increase the efficiency of the n-type doping process.

**4.1.3. Small Molecules.** In addition to conducting polymers, the thermoelectric properties of small molecules such as pentacene, fullerene, and phthalocyanine (Pc) have also been investigated. Theoretical studies have shown, for example, that Pc crystals can be excellent thermoelectric materials when n-doped, with the highest thermoelectric zT value of 2.5 in NiPc.174 Compared to conducting polymers, small molecules can be easily purified and crystallized, yielding well-defined structures that allow the establishment of structure–property relationships. In the following, we will examine some of the recent progress in small-molecule thermoelectrics.

**4.1.3.1. Charge Transfer Salts.** TTF-TCNQ is a well-known charge-transfer salt with a room temperature conductivity as high as 500 S cm⁻¹ in a single crystal. Its Seebeck coefficient and thermal conductivity have been reported to be −28 µV K⁻¹ and 1 W m⁻¹ K⁻¹, respectively. Although the calculated zT value of TTF-TCNQ is only on the order of 10⁻², it was used for the fabrication of prototype thermoelectric devices early on, mainly due to the lack of high-performance n-type materials at that time.¹⁰ Polycrystalline films of TTF-TCNQ show a much lower electrical conductivity (5–30 S cm⁻¹) than single crystals.¹⁷⁵⁻¹⁷⁷ Using N,N-dimethylacetamide (DMA) as the solvent, a drop-cast TTF-TCNQ polycrystalline film can be prepared with a high density of TTF-TCNQ nanofibers. The obtained film shows an improved electrical conductivity of up to 58 S cm⁻¹, while its Seebeck coefficient is 32.3 µV K⁻¹, similar to the reported values for typical TTF-TCNQ polycrystalline films. Thus, a maximum PF of 6.05 µW m⁻¹ K⁻² at room temperature is achieved.¹⁷⁸ In addition to TTF-TCNQ, the TE properties of other similar charge transfer complexes have also been evaluated. Evaporated films of TMTTF-TCNQ and HMTTF-TCNQ (TMTTF, tetramethyl-TTF; HMTTF, hexamethylene-TTF) and solution-processed nanoparticle films of (TTF)[Ni-

(dmit)$_2$]$_2$ (dmit: 1,3-dithiole-2-thione-4,5-dithiolato) were systematically investigated by Mori et al., with the (TTF)[Ni-(dmit)$_2$]$_2$ single crystals showing the highest PF of 29 µW m⁻¹ K⁻² at room temperature.¹⁷⁹ In 2016, Mori et al. reported the thermoelectric properties of another system, benzothienobenzothiophene (BTBT)-based molecular conductors.⁸⁷ Different charge transfer complexes of (BTBT)$_2$XF$_6$ (X = P, As, Sb, and Ta, structure in Figure 16a) were prepared, and their thermoelectric performances were characterized at room temperature. All the charge transfer complexes showed a similar Seebeck coefficient of approximately 15 µV K⁻¹, while (BTBT)$_2$AsF$_6$ possessed the highest electrical conductivity of 4100 S cm⁻¹. Thus, a PF as high as 88 µW m⁻¹ K⁻² was achieved for (BTBT)$_2$AsF$_6$, while all the other complexes had similar values (<50 µW m⁻¹ K⁻²) (Figure 16b).

Both p-doped and n-doped tetrathiotetracene (TTT) thin films were reported by Pudzs et al.¹⁸⁰ p-Type TTT$_2$I$_3$ thin films were obtained by vacuum deposition or the postdeposition of iodine (I$_2$), yielding an electrical conductivity of 1.3 S cm⁻¹, a Seebeck coefficient of 63 µV K⁻¹, and a PF of 0.52 µW m⁻¹ K⁻². For vacuum codeposited n-type TTT-TCNQ films, an electrical conductivity of 0.57 S cm⁻¹, a Seebeck coefficient of −75 µV K⁻¹, and a power factor of 0.33 µW m⁻¹ K⁻² were obtained at an optimized ratio of 1:2 (Figure 16d). A prototype thin-film TE generator based on both n- and p-type TTT complexes was also built as a proof of concept. Efforts have also been made on the development of metal–organic complexes such as Pc–metal complexes (CoPc, NiPc, CuPc, ZnPc, etc.), porphyrin-metal complexes (such as Co(tbp)I, tbp = tetrabenzoporphyrinato), and CuTCNQ with improved thermoelectric performance. We will not go into the details of the TE properties of these metal–organic complexes, but we refer to recently published reviews.¹⁸¹

**4.1.3.2. Small Molecule Semiconductors.** The thermoelectric properties of Na-, K-, and bidoped fullerenes were reported long ago.¹⁸⁴,¹⁸⁵ In 2013, Barbot et al. reported n-doped fullerene with cesium carbonate (Cs$_2$CO$_3$) via a cosublimation process.¹⁸⁶ The conductivity of the C$_60$:Cs$_2$CO$_3$ thin film sharply increased to 6.6 S cm⁻¹ at a molar ratio of 35.8% Cs$_2$CO$_3$, compared to undoped films with a conductivity of only 3.8 × 10⁻⁸ S cm⁻¹. The absolute value of the Seebeck coefficient decreased gradually with the dopant concentration, from 1715 µV K⁻¹ to less than 200 µV K⁻¹. After optimization, a high PF of 28.8 µW m⁻¹ K⁻² was achieved at a molar ratio of 15.2% Cs$_2$CO$_3$.

By using a modified scanning tunneling microscope to study the single-molecule junctions of an endohedral fullerene [a fullerene cage (C$_{80}$-Ih) encapsulating three scandium atoms joined to one nitrogen Sc$_3$N@C$_{80}$], Agrait et al. demonstrated that the sign and magnitude of the thermopower could be tuned by changing the orientation and pressure (Figure 16f).¹⁸² The origin of this bithermoelectric effect was also elucidated by means of density functional theory calculations, highlighting the importance of transport resonances at the molecular junctions. Later, Lambert et al. measured a positive Seebeck coefficient in exohedrally decorated C$_{50}$Cl$_{10}$ fullerene molecules and found that S could be increased by placing more C$_{50}$Cl$_{10}$ molecules in series.¹⁸⁷ Compared to a single C$_{50}$Cl$_{10}$, the Seebeck coefficient of the dimer increased from 8 µV K⁻¹ to 30 µV K⁻¹, while the PF increased from 5 µW m⁻¹ K⁻² to 60 µW m⁻¹ K⁻².

Recently, very large Seebeck coefficients of more than 150 mV K⁻¹ have been reported for pure C$_{60}$ thin films by in situ measurements under ultrahigh vacuum (Figure 16g).¹⁸³ Because of their high resistance, it is difficult to accurately
<figure><img src="image_27.png" /><figcaption>**Figure 17.** (a) Optical microscopy images of the fully printed vertical OTEG and (b) Photograph of the entire flexible monolithic OTEG. (c) Power density ($P_d$) vs output current ($I_{out}$). (a–c) Reproduced with permission from ref 190. Copyright 2020 Elsevier Ltd. (d) Mechanism of Bi interfacial doping of TDPPQ and molecular structure of TDPPQ. (e) Thermoelectric properties of TDPPQ doped with different thicknesses of Bi. (d–e) Reproduced with permission from ref 98. Copyright 2016 Wiley.</figcaption></figure>

measure the Seebeck coefficient of pure (undoped) $C_{60}$ films. However, Nakamura et al. developed a thermopower measurement system that solves this problem by using a custom high-input impedance differential amplifier. Although the electrical conductivity of undoped semiconductor films is relatively low ($<10^{-9}$ S cm$^{-1}$), thermoelectric materials with such high Seebeck coefficients could also be used for sensing, as they are able to react to small variations in temperature. Thus, the thermopower measurement system developed by Nakamura et al. will be useful for the future screening of high-impedance thermoelectric materials.

Improved thermoelectric properties of doped pentacene films were reported in 2011.<sup>188</sup> It was found that the electrical conductivity of an $I_2$-doped pentacene film decreased with increasing film thickness, and the obtained PF of 115 nm thick films was as high as 13 µW m<sup>-1</sup> K<sup>-2</sup> at room temperature. The stability of the $I_2$-doped pentacene film was also improved by encapsulating the doped films with a polyimide protective layer.

In addition to the development of n-type polymer thermoelectrics, very promising n-type thermoelectric performances have been reported in organic small molecules through the proper selection of dopants as well as an appropriate design of the molecular structure. Regarding molecular structure engineering and appropriate dopant selection, Koster et al. tailored the polarity of fullerene derivatives to enhance the miscibility between the host and dopant molecules.<sup>189</sup> They showed that functionalizing the fullerene with a hydrophilic triethylene glycol-type side chain led to a greatly improved doping efficiency of 18%, compared to <1% for the nonpolar side chain-functionalized reference fullerene using the same dopant concentration. The improved doping efficiency increased electrical conductivity, which reached a maximum of 2.05 S cm<sup>-1</sup> and enabled a PF of 16.7 µW m<sup>-1</sup> K<sup>-2</sup>.

They further improved their results by using as side chain a double ethylene glycol (PTEG-2) doped with 8 wt % N-DMBI, recording a high value of $zT > 0.3$ at 120 °C.<sup>95</sup> The obtained films showed an impressive PF of 80 µW m<sup>-1</sup> K<sup>-2</sup> coupled with an ultralow and constant thermal conductivity ($k < 0.1$ W m<sup>-1</sup> K<sup>-1</sup>) over the studied temperature range, exhibiting one of the highest PF values ever reported for a small-molecule semiconductor (Figure 16i).

Previously, Massetti et al. employed fullerene modified with a single ethylene glycol chain together with an inkjet printable formulation of PEDOT:PSS; this work was the first example of a monolithic and integrated µ-OTEG (Figure 17a), achieving a power density of 30 nW cm<sup>-2</sup> at an external temperature difference of 25 K (Figure 17b,c).<sup>190</sup> The 128 thermocouples were defined inside plastic and flexible PEN film, thus limiting the thermal losses that are generally linked to the use of plastic substrates by using a combination of laser micromachining and inkjet printing; this method provided the first example of a complete printed and integrated OTEG.

Zhu et al. demonstrated that Bi, an excellent dopant that has been widely used in inorganic TE materials, allows for the interfacial doping of small molecules (sketch in Figure 17e), resulting in high n-type performance.<sup>98</sup> The electrical conductivity of thiophene-diketo-pyrrolopyrrole based quinoidal (TDPPQ) films in contact with a 2 nm thick Bi layer dramatically increased from <0.01 S cm<sup>-1</sup> (undoped film) to 2.4 S cm<sup>-1</sup>, while the Seebeck coefficient was only slightly affected. With a thicker Bi film, the electrical conductivity became saturated, while the Seebeck coefficient began to drop. A maximum PF of 62 µW m<sup>-1</sup> K<sup>-2</sup> was reached at 20 °C for an
optimized Bi thickness of 2 nm (Figure 17f), and these values increased to 113 µW m⁻¹ K⁻² at a higher temperature (60 °C).

In contrast to the extrinsic doping approach, a solution- processable and self-doped small-molecule perylene diimide (PDI) was reported by Russ et al.¹⁹¹ Quaternary amine substituents containing different spacer groups were used to functionalize the PDI molecule. The electrical conductivity increased 100-fold when the alkyl spacer groups between quaternary amine substituents and the PDI core increased from two (PDI-1) to six (PDI-3) methylene units, reaching values as high as 0.5 S cm⁻¹. All three PDIs had an n-type Seebeck coefficient of ~ -200 µV K⁻¹, regardless of the alkyl spacer length. Thus, a PF as high as 1.4 µW m⁻¹ K⁻² was obtained for the PDI-3 film.

Zhu et al. demonstrated that small changes in the conjugated molecular structure dramatically affected the thermoelectric performance of small molecules. For example, by going from a quinoid to an aromatic structure, a 50 times higher PF with values up to 236 µW m⁻¹ K⁻² was achieved for chemically doped dihydropyrrolo[3,4-c]pyrrole-1,4-diylidene)bis(thieno[3,2-b]- thiophene) derivatives.⁹⁹ Experimental analysis and theoretical calculations revealed that the aromatic structure provided the molecule high electron mobility and appropriate energy levels for a more efficient doping reaction with N-DMBI. Because of the low in-plane thermal conductivity of 0.34 W m⁻¹ K⁻¹, a high zT value of 0.23 was measured at 373 K, which was the highest value reported to date for small-molecule organics. Kemerink et al. investigated the n-type thermoelectric properties of [6,6]- phenyl-C61-butyric acid methyl ester (PCBM) for different solution-based doping methods.¹⁹² The authors presented a novel inverse-sequential doping method where PCBM was deposited on a previously cast N-DMBI dopant film, achieving a PF as high as 35 µW m⁻¹ K⁻² with a conductivity of 0.40 S cm⁻¹. The authors also showed that adding a small fraction of higher electron affinity material (i.e., core-cyanated naphthalene diimide) into PCBM could lead to a higher S. This addition creates a second “trap” in the DOS, leading to a sharp increase in the Seebeck coefficient while at the same time decreasing the electrical conductivity. In addition, they presented a simple analytical model to unify the thermopower and conductivity of both p- and n-type TE materials and found good consistency between the experimental and calculated data. Remarkably, they found that most of the p-type data were described well by the model using a hop frequency of approximately 0.5–5 × 10¹¹ s⁻¹, whereas a higher frequency of about 0.5–5 × 10¹³ s⁻¹ was required for n-type data, implying that n-type organic semi- conductors seemed to have more potential to eventually achieve higher zT than p-type materials.

4.1.4. Thermal Transport in Organic Semiconductors.

So far, we mainly focused on showing different ways to enhance the PF of organic thermoelectric materials, trying to match the values of bismuth-telluride alloys. However, to achieve higher zT, thermal conductivity is another critical parameter that must be optimized. The thermal conductivity (k) of organic (semi-)conductors is generally considered to be lower than 1 W m⁻¹ K⁻¹.¹⁹³ However, similarly to the electrical conductivity, thermal conductivity in organic semiconductors is also an anisotropic property, as heat transport is generally more efficient along the chain direction than between chains. This is due to the weak nature of the van der Waals interactions between adjacent chains and/or the formation of lamella separated by the side chains, which limit the phonon and electron transport, thus reducing the lattice and the electronic components of the thermal

conductivity. Therefore, in many organic semiconductors, such as P3HT or PBTTT-C₁₄, it is possible to observe a significant anisotropy factor from the in-plane to the out-of-plane direction, showing larger phonon propagation distances along the π-conjugated backbone.¹⁹⁴ Stretch-aligning semicrystalline or amorphous polymer films/fibers also drastically increases the phonon propagation, with *k* values along the stretching direction being more than 1 order of magnitude higher than those perpendicular to the stretching direction.¹⁹⁵⁻¹⁹⁷ For example, *k* = 40 W m⁻¹K⁻¹ was recorded along the polyethylene axis in crystalline films, while the transverse value of *k* was around 100 times lower (around 0.3–0.4 W m⁻¹K⁻¹).¹⁹⁵ In addition, molecular dynamic (MD) simulations have shown a correlation between high thermal conductivities and rigidity of the polymer backbone due to the suppressed segmental rotations, leading to long-distance phonon propagation (that is, higher persistence length).¹⁹⁸ The increase of the persistence length is also linked to higher charge carrier mobility, which is also a critical parameter for σ and S, and thus further complicating independent optimization of the material parameters that determine zT. Therefore, to maximize the thermoelectric properties, the effect of polymer backbone rigidity on the lattice component of the thermal conductivity needs to be considered when designing new conducting polymers.

Furthermore, the process of doping organic semiconductors not only increases the electrical conductivity but also leads to a rise of the electronic component of *k* under the Wiedemann–Franz law (k ∝ Lσ, where *L* is the Lorenz number).¹⁹⁹ While noticeable for conductivities above 10 S cm⁻¹, this effect becomes predominant for σ exceeding 500 S cm⁻¹, as shown for PEDOT:PSS.¹⁹⁹,²⁰⁰ To further understand the impact of charge carriers on thermal conductivity, Kemerink et al.²⁰¹ analyzed the effect of energetic disorder and correlation length on the relationship between σ and *k*. Monte Carlo (MC) simulations showed that lowering the energetic disorder is a key strategy to reduce *L* (up to 5 times lower than the Sommerfield value). This decreases the contribution of charge carriers to the thermal conductivity and improves the thermoelectric figure of merit.

Finally, Campoy-Quilles et al. showed how, counterintuitively, the doping process (at least for low dopant concentrations) might be beneficial to reduce the thermal conductivity.²⁰² For the system PBTTT doped with F₄TCNQ, they observed that both annealing temperature (acting on the size of the polymeric crystalline domains) and doping process (at doping concentrations < 1 mol %) can be used to tune the thermal conductivity of PBTTT. In particular, they observed a 2-fold reduction in the value of *k* without significantly affecting the crystalline quality of the film. This process resembles the effect of phonon scattering observed in inorganic alloys.

Overall, while several theoretical studies have enhanced our understanding of thermal transport in organic semiconductors, more research needs to be done to enable a proper control of σ, S, and *k*. Here, we want to stress the critical role of anisotropy on the thermoelectric properties of organic semiconductors and the importance of measuring and correlating all thermoelectric parameters in the same direction. For a more detailed theoretical description of the physics behind heat transport in organic semiconductors, we refer the reader to the following seminal works.²⁰³⁻²⁰⁶
<figure><img src="image_28.png" /><figcaption>**Figure 18.** (a) Schematic of a section of a graphene sheet, showing the chiral indices (n,m) of SWCNTs formed by rolling the sheet from (0,0) to (n,m) along a roll-up vector (blue line). White hexagons correspond to (n,m) indices of semiconducting SWCNTs, while gray hexagons correspond to metallic SWCNTs. The green hexagons and orange hexagon correspond to the s-SWCNT species present in the polydisperse and monodisperse samples for which PL are displayed in (e) and (f). (b) Chemical structure of the (7,5) s-SWCNT, illustrating the nanotube diameter as well as the optical and electronic bandgap energies. (c) Photograph of raw SWCNT soot produced by the laser vaporization (LV) method. (d) Photograph of semiconducting SWCNT inks prepared by selective extraction with polyfluorene polymers. The color of each ink is determined by the diameter distribution of dispersed s-SWCNTs. (e) Photoluminescence excitation (PLE) map of polydisperse SWCNT sample with many s-SWCNTs present (green hexagons, a), each s-SWCNT peak labeled by its (n,m) indices. (f) The PLE map of a monodisperse sample is dominated by a single species, in this case the (7,5) s-SWCNT (orange hexagon, a). Phonon sidebands of the (7,5) excitation spectrum are labeled with asterisks.</figcaption></figure>

## 4.2. Carbon Nanomaterials

Carbon nanomaterials (CNMs) such as carbon nanotubes (CNTs) and graphene have received a great deal of attention over the past decades for energy conversion, storage devices, catalysts, and biosensors. Both CNTs and graphene are composed of an extended hexagonal lattice of sp²-bonded carbon atoms, which renders exceptional electrical, optical, thermal, chemical, and mechanical properties. CNTs and graphene high electrical conductivities, resulting from nearly identical and small effective masses for electrons and holes,²⁰⁷ make them interesting candidates for thermoelectric applications. At the level of an individual CNT or an isolated graphene sheet, this high electrical conductivity is offset by the strongly disadvantageous property of high intrinsic thermal conductivity for both materials.²⁰⁷ Additionally, the Seebeck coefficient is relatively low for both semimetallic graphene, multiwalled carbon nanotubes (which statistically are mostly metallic) and the ca. 33% of as-produced single-walled carbon nanotube (SWCNT) samples that have metallic electronic structure.²⁰⁸ Thick CNTs layers, often composed of unaligned “bundles and mats” can have much lower electrical and thermal conductivities than individual CNTs due to electrical and thermal contact resistances between bundles. Composites of CNTs/graphene with polymers (or inorganic materials) can potentially decouple electrical and thermal conductivity by introducing interfaces that provide more facile paths for electron transfer than phonon transfer. Additionally, the introduction of interfaces provides an opportunity to realize carrier energy filtering to decouple the thermopower and electrical conductivity. Molecular dopants offer a means to tune the dark carrier density and may also introduce sites that effectively scatter phonons. The complex interrelationship between the thermopower, electrical conductivity, and thermal conductivity in nanocarbon materials and

composites should remain an active area of fundamental research. Furthermore, recent strategies for separating SWCNTs by electronic structure and systematic tuning of the Fermi level by doping have enabled the realization of high thermoelectric power factors for both semiconducting and metallic SWCNT networks. Some of these results have been recently reviewed.²⁰⁸ Here we take specific remarkable examples of those achievements while also focusing on a number of new achievements that have emerged in this fast-moving field over the past few years since that review.

### 4.2.1. Synthesis, Processing, and Basic Properties of SWCNTs.
Single wall carbon nanotubes (SWCNT), discovered in 1991, can be thought of as a graphene sheet that has been wrapped up into a seamless cylinder (Figure 18a).²⁰⁹ Early studies utilizing scanning tunneling microscopy and spectroscopy found that both metallic and semiconducting carbon nanotubes existed, depending sensitively on the wrapping angle.²¹⁰ SWCNTs can be produced by a number of different methods and the resulting material properties can depend quite sensitively on the production method. Synthesis methods include chemical vapor deposition (fixed bed, fluid bed, floating catalyst, and plasma enhanced), laser vaporization, and arc discharge.²¹¹

Most synthetic routes to produce SWCNTs produce the statistically expected ratio between those two forms with 1/3 of metallic tubes (m-SWCNT) and 2/3 of semiconducting tubes (s-SWCNT). As such, most syntheses tend to produce raw SWCNT soot that is black (Figure 18b) because it contains a polydisperse mixture of many SWCNTs (Figure 18e) that absorb with different wavelengths. However, a number of effective strategies have been perfected over the past decade to separate s-SWCNTs and m-SWCNTs by electronic structure (Figure 18d,f).²¹² The optical bandgap of s-SWCNTs is
<figure><img src="image_29.png" /><figcaption>**Figure 19.** (a) Thermopower for five s-SWCNT species and one m-SWCNT species, calculated by density functional theory (DFT). Reproduced with permission from ref 9. Copyright 2016 Springer Nature. (b) Typical calculated results for (top panels) the electronic density of states and (bottom panels) Seebeck coefficient (dashed lines) and power function (solid lines) for (24,0) m-SWCNT and (26,0) s-SWCNT as a function of chemical potential ($\mu$). Reproduced with permission from ref 233. Copyright 2016 IOP.</figcaption></figure>

inversely dependent on diameter and varies from ca. 0.5 to 1.4 eV for a corresponding diameter range of 2.6–0.6 nm that is typical of most commonly used synthesis routes.²¹³ Given the relatively large diameter-dependent exciton binding energy,²¹⁴ this translates to an electronic bandgap (the relevant metric for thermoelectrics) of ca. 0.6–1.8 eV.

As-produced SWCNT samples typically consist of long (>1 µm) SWCNTs that are present in large bundles, consisting of, e.g., 10 or more SWCNTs, due to van der Waals attractions between nanotubes. A general strategy for incorporating SWCNTs into solid-state films required for applications such as thermoelectrics involves dispersing the SWCNTs into a solvent by ultrasonic processing. Coupled with a broad variety of surfactants, dispersing polymers, and solvents with varying Hansen solubility parameters, such processing routes can be used to produce a wide range of SWCNT dispersions in aqueous or organic solvents with or without a dispersing agent.²¹⁵⁻²¹⁸ These dispersions can then be deposited into films and structures of varying form factor by a number of different deposition methods, including vacuum filtration, spin-coating, ultrasonic, or pressure-driven spray deposition, ink-jet printing, blade coating, Gravure printing, and others.²¹⁹ Aqueous and organic solvent-based dispersions are also the starting points for most separations of SWCNTs by electronic structure, diameter, and chirality. The most relevant separation/enrichment strategy for TE applications is the separation of as-prepared SWCNT soot into populations of pure semiconducting, pure metallic, or tunable populations with controlled metal/semi ratios. Many separation techniques are applicable to aqueous SWCNT dispersions, including density gradient ultracentrifugation,²¹⁶,²²⁰ column chromatography separation,²²¹ and aqueous two-phase extraction.²²²,²²³ A predominant strategy for selective enrichment of semiconducting SWCNTs in organic solvents involves the use of polymers that preferentially wrap and disperse s-SWCNTs, leaving m-SWCNTs undispersed.²²⁴,²²⁵ These polymer-wrapped s-SWCNT inks tend to have vibrant tunable colors, due to the selection of a relatively narrow subset of s-SWCNTs (Figure 18d).

The cavitation energy in ultrasonic dispersion processes is used to débundle the SWCNTs, with an unavoidable side-effect being the progressive cutting and shortening of the SWCNTs with increasing sonication time. Because electrical conductivity

in an electronically coupled SWCNT network scales with increasing length of the individual nanotubes, the thermoelectric properties of such ultrasonic-processed SWCNTs is generally degraded as the sonication time is increased.²²⁶ Because of the ultrasonic-induced shortening of SWCNTs, new routes are being explored for less severe processing of SWCNTs for opto-electronic applications. Recent studies have demonstrated shear-force mixing as an attractive alternative to ultrasonic processing.²²⁷ Shear-force mixing does not feature any cavitation forces and has been shown to enable the dispersion of very long SWCNTs with retained electrical and optical properties. Another strategy that has emerged involves using as-produced SWCNTs that are never subjected to any ultrasonic or shear-force solution-phase processing.²²⁸⁻²³¹ Such samples tend to be highly fibrous interwoven SWCNT mats produced by floating catalyst chemical vapor deposition (FCCVD). Such mats can be gently processed by simple densification steps or spinning steps to produce high-density mats or fibers, respectively, that can be used directly in thermoelectric applications. We discuss these advances in section 4.2.3.

### 4.2.2. Tunable Thermoelectric Performance of “Neat SWCNTs”.
Before discussing the role(s) of CNTs for enhancing thermoelectric composites, we first consider the thermoelectric properties and performance of SWCNTs on their own. The understanding of this important model system is derived from the extensive research devoted to the control of the electronic structure with the goal to fine-tune the Seebeck coefficient and the electrical conductivity. The tuning of these properties has been achieved by controlling the ratio of semiconducting vs metallic, the presence of multiwalls possessing specific electronic properties, the doping by replacement of atoms in the tubes, or by redox reaction with another external chemical species. Here, we summarize some of the key advances and strategies employed for tuning, optimizing, and mechanistically understanding the thermoelectric performance of “neat” SWCNT samples. Several years ago, important computational work demonstrated that the Seebeck coefficient can be (1) particularly large in pure s-SWCNTs (Figure 19a) and (2) modulated with high sensitivity through doping (Figure 19b).⁹,²³²,²³³ Specifically, around 2015, a number of groups predicted exceptionally large peak thermopowers for s-SWCNTs that could exceed 1 mV K⁻¹ and that depended sensitively on Fermi energy.⁹,²³²,²³³ Other
<figure><img src="image_30.png" /><figcaption>**Figure 20.** (a) Example of molecular doping of SWCNTs, in this case, n-type doping by adsorption of benzyl viologen to the SWCNT surface. Reproduced with permission from ref 234. Copyright 2009 American Chemical Society. (b) Example of m-SWCNT (left) and s-SWCNT (right) doped n-type by encapsulation of ferrocene molecules. Reproduced with permission from ref 240. Copyright 2014 IOP Science. (c) Schematic of electrolyte gating experimental setup for measuring thermoelectric properties of carbon nanotube films. Reproduced with permission from ref 248. Copyright 2019 American Chemical Society.</figcaption></figure>

<figure><img src="image_31.png" /><figcaption>**Figure 21.** (a) Evolution of (6,5) s-SWCNT network optical absorption as a function of p-type dopant (OA) concentration that the film is soaked in, showing $S_{11}$ transition quenching. (b) Change in Fermi level ($\Delta E_f$) for (6,5) or LV SWCNTs as a function of the degree of $S_{11}$ quenching for progressively p-doped networks. (a,b) Reproduced with permission from ref 9. Copyright 2016 Springer Nature. (c) Plasma-torch s-SWCNTs progressively doped with AgTFSI, showing that the plasma resonance (PR) grows in as the $S_{11}$ is quenched. Reproduced with permission from ref 254. Copyright 2018 Taylor & Francis. (d) Evolution of absorption spectrum for arc-discharge m-SWCNTs with gate voltage ($V_g$). (e) Contour map of absorbance for same m-SWCNT sample as a function of absorbed photon energy and gate voltage. (f) Conductivity of the same m-SWCNT sample versus gate voltage. (g) Schematic diagram of density of state for the (15,6) s-SWCNT, chosen as representative of the arc-discharge SWCNTs. (d–g) Reproduced with permission from ref 243. Copyright 2014 Wiley.</figcaption></figure>

predictions also suggested multiple maxima in the Seebeck and power function due to the sharp van Hove singularities (vHS) of SWCNTs, especially in cases where the Fermi energy overlapped with a region of energy comprising multiple closely spaced vHS (e.g., Figure 19b, left).²³³

Even though CNTs had been studied in TE composites for many years by that point, these exciting predictions and associated experiments jump-started a renewed interest in the ultimate potential of neat SWCNT-based TE materials. By neat, we mean SWCNT thin films and samples where the SWCNTs can be considered the primary or sole transport medium, as opposed to SWCNT-based composites, where transport within both the SWCNT and matrix material(s) must be considered in concert. This is not to say that such films are comprised solely of SWCNTs; for example, many of these SWCNT networks contain residual “wrapping” polymer that was used to disperse the SWCNTs and most optimized networks include dopant molecules used to tune the Fermi energy and carrier density (in direct analogy to semiconducting polymer thermoelectrics). The important distinction employed in the following conversation is that the dominant charge transport pathways in “neat” SWCNT samples are the nanotubes and bundles themselves and not these secondary phases.

We break this section into subsections based on critical considerations for employing neat SWCNTs as thermoelectric materials. First, we discuss the noncovalent and covalent functionalization and doping strategies used to tune the thermoelectric properties of neat SWCNT networks. These can generally be applied to either mixed (semi/metal) samples or samples of highly pure s-SWCNTs or m-SWCNTs. Next, we discuss solution-processed SWCNTs that have been separated by electronic structure to explore the dependence of TE performance on the proportion of s-SWCNTs, m-SWCNTs, and the doping-controlled carrier density. Finally, we discuss the anisotropy of the TE transport properties within highly aligned SWCNT samples.

#### 4.2.2.1. Noncovalent Functionalization and Doping.
Because of the strong dependence of conductivity and thermopower on Fermi energy, a dominant route used to optimize the TE performance of CNTs involves tuning the electronic structure, carrier density, and Fermi level via noncovalent doping. The most typical strategy for noncovalent doping of SWCNTs is molecular redox doping, which involves exposing CNTs to organic molecules with strong electron donor or acceptor character to promote spontaneous electron transfer to/from the SWCNTs. Molecules can be adsorbed onto the surfaces of SWCNTs (Figure 20a)⁹,²³⁴⁻²³⁷ and/or may be
encapsulated within the endohedral volume of the SWCNTs (Figure 20b).²³⁸⁻²⁴² Donor molecules with HOMO values closer to vacuum than the first SWCNT conduction level ($c_1$) values serve as reductants, injecting electrons into the SWCNTs that can contribute to electrical conductance in a coupled SWCNT network. In contrast, acceptor molecules with LUMO values farther from vacuum then the first SWCNT valence level ($v_1$) serve as oxidants, removing electrons from SWCNTs to produce mobile holes in SWCNT networks. Another non-covalent method for doping SWCNT networks involves electrostatic gating in an electrochemical transistor-based (FET) geometry (Figure 20c). This method is particularly useful for systematically tuning the Fermi level with relatively fine control in both the p- and n-type regimes.²⁴³ Several studies have used this strategy to perform simultaneous measurements of conductivity and thermopower to probe the p- and n-type PF of a variety of SWCNT samples.²⁴³⁻²⁴⁸

Quantitative correlation of TE properties with carrier density and Fermi level is a difficult task with disordered organic semiconductors such as semiconducting polymers and SWCNTs. A significant amount of research is still devoted to understanding the appropriate analyses needed to make these correlations in SWCNTs, and we point the reader to a couple of recent book chapters²⁴⁹,²⁵⁰ and another recent study.²⁵¹ Briefly, optical and infrared spectroscopies are particularly useful for tracking the doping-induced modifications to the carrier density of SWCNTs. The strong excitonic optical transitions for s-SWCNTs in the visible and near-infrared are particularly sensitive to the carrier density, due to impact of phase space filling by charge carriers on the oscillator strength of excitonic transitions (Figure 21a,b). Because a direct estimation of areal or volumetric carrier density in SWCNTs can be quite challenging, the relative strengths of these optical transitions are routinely tracked as a function of doping as a proxy for the carrier density in a s-SWCNT network. X-ray photoelectron spectroscopy can also be used to probe the change in Fermi level, as a function of systematic changes in doping level because the kinetic energy of photoelectrons are referenced to the Fermi level (Figure 21b).⁹,²³⁵,²⁵²,²⁵³

Nonoguchi et al. demonstrated that the plasmon resonance of free charge carriers in the infrared can also be used to systematically track relative carrier density in s-SWCNTs (Figure 21c).²⁵⁴ While most molecular doping strategies do not modify the Fermi level strongly enough to quench the relatively high-energy excitonic transitions of m-SWCNTs,²⁵⁵ the plasmon resonance intensity can report on systematic changes to the carrier density in m-SWCNTs.²⁵⁶ However, electrostatic gating has been shown to be effective in modulating the Fermi level to a large enough degree to quench m-SWCNT excitonic transitions (Figure 21d,e), and the doping-induced tuning of the Fermi level increases the conductivity for both s- and m-SWCNTs, as shown in Figure 21f,g for electrostatic doping of m-SWCNTs.²⁴³

A wide variety of redox molecules can p/n-dope SWCNTs. Early doping work centered around aggressive acids and bases, many of which were known from decades of research to intercalate graphite. Bronsted acids such as nitric, sulfuric, and chlorosulfonic acids are widely known to produce p-type SWCNTs,²⁵⁷ as are halide gases with large electronegativities such as bromine and iodine.²⁵⁸ Alkali metals such as sodium, lithium, and cesium were shown in early SWCNT research to produce n-type SWCNTs.²⁵⁹ A wide variety of Lewis acids and bases has been explored for doping SWCNTs and other types of

CNTs.²⁶⁰ Some basic trends are as follows. Aromatic dopants functionalized with highly electron-withdrawing groups (e.g., F, CF₃, CN) tend to serve as good oxidants for producing p-type SWCNTs. Some examples include tetracyanoquinodimethane (TCNQ), tetracyanobenzene (TCNB), and many of their derivatives (e.g., F₄-TCNQ, etc.). Transition metal chlorides such as FeCl₃,²³⁰,²³¹,²⁶¹,²⁶² AuCl₃,²²⁸ and AgCl²⁴² also tend to be good p-type dopants for SWCNTs. The reactivity of very strong chloride-based oxidizing agents such as SbCl₅ can be difficult to control, leading to both oxidation and chlorination of aromatic species.²⁶³ As such, these acids can be “masked” with alkyloxonium cations (e.g., as triethyloxonium hexachloroantimonate [Et₃O⁺SbCl₆⁻]) to produce strong oxidants that exclusively form aromatic radical cations due to electronic doping (i.e., electron extraction from the HOMO of the aromatic system) without concomitant side-wall chlorination or alkylation.²⁶⁴ Triethyloxonium hexachloroantimonate (abbreviated here as OA) has been used to produce p-type SWCNTs with conductivity that is both highly tunable⁹,²⁰⁸,²⁵²,²⁵³,²⁶⁵ and relatively stable in ambient conditions.²⁶⁶

Bis(cyclopentadienyl)M(II) compounds with the formula M(C₅H₅)₂, where M is a transition metal (e.g., cobaltacene,²³⁸ ferrocene,²³⁹ nickelocene), tend to have small ionization potentials and serve as good reductants. Several studies have shown that these n-type dopants can actually be encapsulated within SWCNTs during the doping process,²³⁸,²⁴¹ which provides a unique strategy for improving the stability of n-type doping (vide infra). A number of amine-, imine-, and phosphine-based molecules also have ample reducing strength to produce n-type SWCNTs, as demonstrated by a few survey studies of such molecules.²³⁵,²³⁶ These include molecules such as hydrazine, ethylenediamine, methylamine, diethylamine, oleamine, diphenylphosphine, triphenylphosphine, and associated derivatives, as well as polymers such as polyethylenimine (PEI).⁵⁰,²³⁵,²³⁶ PEI has been especially popular for n-type doping of SWCNTs in TE applications,²²⁸,²³⁰,²⁶⁷⁻²⁷⁰ in large part for the relative stability imparted by the viscous polymer that tends to coat SWCNTs well and be somewhat impermeable to atmospheric ingress.²³⁵,²³⁶ Reduced viologens are also highly reductive and have been used widely to produce n-type SWCNTs.²³¹,²³⁴,²⁵²,²⁵³

Because of rapid compensation of electrons by oxygen and/or water molecules, n-type dopants tend to be significantly less stable than their p-type counterparts, so a significant amount of research has been devoted to developing means by which to improve n-type doping stability for TE applications of CNTs. Fukumaru et al. succeeded in intercalating cobaltocene (CoCp₂) inside the tubes to transform a p-type into a n-type SWCNT (Seebeck coefficient = -41.8 µV K⁻¹), with a superior air stability due to the molecular shielding effect.²⁴¹ Electron transfer from encapsulated CoCp₂ increased the carrier density in the s-SWCNTs and the electrical conductivity of the CoCp2@SWCNT film reached 432 S cm⁻¹, compared to 44.5 S cm⁻¹ for empty SWCNTs. The power factor of the CoCp2@ SWCNT film was 75.4 µW m⁻¹ K⁻², and combined with a low thermal conductivity of 0.15 W m⁻¹ K⁻¹, a zT value of 0.157 was achieved at 320 K (zT = 0.025 without CoCp₂ doping).²⁴¹

Nonoguchi et al. demonstrated air stability for mixed SWCNTs n-type doped with a variety of phosphine-, amine-, and imine-containing aromatic dopants²³⁶ and also demonstrated a strategy for combining ordinary salts such as sodium chloride (NaCl), sodium hydroxide (NaOH), and potassium hydroxide (KOH) with crown ethers for the production of
stable n-type SWCNTs (Figure 22a).<sup>237</sup> The stable complexes formed between crown ethers and metal-ions enhances the

<figure><img src="image_32.png" /><figcaption>**Figure 22.** (a) Benzo-crown ether-stabilized alkali metal n-type doping strategy developed by Nonoguchi et al. (b) Only the crown ether-stabilized metal salt succeeds in doping the SWCNTs n-type, whereas the isolated crown ether and sodium borohydride do not. (a,b) Reproduced with permission from ref 237. Copyright 2016 Wiley. (c) Adsorption isotherms of o-MeO-DMBI on SWCNT networks showing formation of multiple layers on the SWCNT surface. (d) n-Type stability is achieved only after the formation of the second layer of dopants. (c,d) Reproduced with permission from ref 271. Copyright 2019 American Chemical Society. (e) s-SWCNT networks doped n-type by potassium crown ether and BV are stable after encapsulation with 50 nm of Al₂O₃ via atomic-layer deposition. Reproduced with permission from ref 252. Copyright 2017 Royal Society of Chemistry.</figcaption></figure>

reducing power of the “naked” counteranions, which enhances the electron donating ability of anions that would not typically reduce SWCNTs, such as OH⁻ or Cl⁻, to enable n-type doping of SWCNTs (Figure 22b). A recent study from Nakashima et al. tracked adsorption isotherms of the n-type dopant 2-(2-methoxyphenyl)-1,3-dimethyl-2,3-dihydro-1H-benzo[d]imidazole (o-MeO-DMBI) and found that the formation of multilayers was critical to achieving long-term stability of the n-type thermopower and PF (Figure 22c,d).<sup>271</sup>

We note here that while several of these doping strategies have provided relatively air-stable mixed SWCNTs (containing both s- and m-SWCNTs), their success has been more limited in highly pure s-SWCNT samples. This has motivated other strategies, such as physical encapsulation, for retaining n-type

TE properties. As an example, MacLeod et al. encapsulated benzyl viologen-doped s-SWCNTs with Al₂O₃, deposited by atomic-layer deposition, to produce highly stable n-type s-SWCNTs (n-type PF > 400 µW m⁻¹ K⁻², Figure 17e).<sup>252</sup> In contrast, identically prepared unencapsulated s-SWCNT networks swiftly transitioned to intrinsic networks with low PF values.<sup>252</sup> It remains a challenge to the field to explore the roles of electronic structure, diameter, dopant redox potential, temperature dependence, and encapsulation strategy on the stability of n-type SWCNT networks for durable TE applications.

#### 4.2.2.2. Separated Semiconducting and Metallic SWCNTs.
Semiconductors remain the predominant materials explored for thermoelectrics due to the Seebeck coefficient being particularly small for metals. As such, a natural progression for the field of SWCNT thermoelectrics involves understanding the role of SWCNT electronic structure, and several studies over the past several years have explored SWCNT samples with finely tunable populations of s- and m-SWCNTs. Nakai et al. (Figure 23)

<figure><img src="image_33.png" /><figcaption>**Figure 23.** Power factor (top) and Seebeck coefficient (bottom) as a function of resistivity for s-SWCNT networks (filled symbols) and mixed or metallic SWCNT networks (open symbols). High resistivity data points for s-SWCNTs correspond to thermally annealed samples, while low-resistivity data points correspond to acid-doped p-type samples. Reproduced with permission from ref 272. Copyright 2014 IOP Science.</figcaption></figure>

produced an early study on SWCNTs separated by density gradient ultracentrifugation (DGU), demonstrating an inherently low Seebeck coefficient for m-SWCNTs (10–20 µV K⁻¹) and high Seebeck coefficients for s-SWCNTs (80–160 µV K⁻¹).<sup>272</sup> Theoretical calculations that followed closely thereafter captured the high Seebeck coefficient in s-SWCNTs (up to ca. 2000 µV K⁻¹) and also found that a decrease in the diameter (increase in the band gap) of the s-SWCNT leads to higher Seebeck coefficients.<sup>9,232</sup>

These early studies suggested that one of the general issues leading to limited PF in films of doped-SWCNTs is the presence
<figure><img src="image_34.png" /><figcaption>(c) Thermopower, $\alpha$ (μW/K) 100
Electrical conductivity, $\sigma$ (S m$^{-1}$)</figcaption></figure>

<figure><img src="image_35.png" /><figcaption>(d)</figcaption></figure>

<figure><img src="image_36.png" /><figcaption>(e)</figcaption></figure>

<figure><img src="image_37.png" /><figcaption>(b) Cleavable Polymers</figcaption></figure>

<figure><img src="image_38.png" /></figure>

**Figure 24.** (a) Three noncleavable and (b) two cleavable polymers used to selectively disperse semiconducting SWCNTs from raw soot. (c) Dependence of thermopower (top) and power factor (bottom) on conductivity for controllably doped laser vaporization (LV) s-SWCNT films prepared from either noncleavable PFO-BPy polymer (red squares) or SMP cleavable polymer (green diamonds and blue circles). Green diamonds represent removal of the SMP polymer in solution before spray-depositing s-SWCNT network. Blue circles represent removal of the SMP polymer in the solid-state after spray deposition of the s-SWCNT network. (d) Dependence of power factor on conductivity for controllably doped plasma torch s-SWCNT network prepared with cleavable PFPD polymer, where the polymer was removed in the solid state. Red squares show p-type PF resulting from doping with triethyloxonium hexachloroantimonate and gold diamonds show n-type PF resulting from doping with benzyl viologen. (e) Dependence of power factor on average bundle diameter (analyzed via AFM) in various s-SWCNT networks. (c) Reproduced with permission from ref 265. Copyright 2016 American Chemical Society. (d,e) Reproduced with permission from ref 252. Copyright 2016 Royal Society of Chemistry.

of metallic m-SWCNT in quantities large enough to create percolation paths dominating the electronic transport and providing low Seebeck coefficient. Research at the National Renewable Energy Lab has attempted to address this issue through studies utilizing selective polymer-based enrichment of s-SWCNTs,²²⁵ in conjunction with fine-tuning the Fermi level with molecular doping.⁹,²⁵²,²⁵³,²⁶⁵ With these strategies, Avery et al. reached a p-type PF = 340 µW m⁻¹ K⁻² at the optimal doping level for s-SWCNT networks composed of ca. 1.0 nm average diameter s-SWCNTs.⁹ The s-SWCNT networks produced in this study had some polyfluorene-based polymer remaining on s-SWCNT surfaces because typical “non-cleavable” wrapping polymers cannot be fully removed from the resulting s-SWCNT network (Figure 24a).

Follow-on work demonstrated that a polyfluorene wrapping polymer with easily cleavable hydrogen bonds between monomers (Figure 24b) enabled the improvement of the PF for a given SWCNT source material.²⁵²,²⁶⁵ The first study demonstrated that the p-type PF of s-SWCNT networks prepared from enriched laser vaporization SWCNTs could be improved from ca. 150 µW m⁻¹ K⁻² (noncleavable polymer) to 400 µW m⁻¹ K⁻² by complete removal of the cleavable polymer (Figure 24c).²⁶⁵ This result unambiguously confirms that the doped s-SWCNTs themselves are the sole source of the high TE performance in these networks. MacLeod et al. then expanded this approach by utilizing another cleavable polymer based on imine linkages between monomers.²⁵² The champion p-type PF of 700 µW m⁻¹ K⁻² in this study belonged to polymer-free s-SWCNT networks prepared from plasma torch synthesized

SWCNTs enriched by the imine-based cleavable polymer (Figure 24d). This study also demonstrated that, due to the equivalent electron and hole effective masses in SWCNTs, n-type s-SWCNT thin films with equally large PF could be produced by utilizing an amine-based n-type dopant (Figure 24d). The peak PF correlated inversely with the average diameter of s-SWCNT bundles in the network (Figure 24e). Huang et al. also demonstrated a flavin-based s-SWCNT extraction method, whereby the TE performance metrics were improved by 2 orders of magnitude when the flavin was fully removed from the s-SWCNT network.²⁷³

The thermal conductivity of mixed SWCNTs and enriched s-SWCNTs has also received significant attention recently, with several intriguing questions remaining to be answered. Avery et al. found that the thermal conductivity of undoped s-SWCNT networks was relatively high ($k = 16.5$ W m⁻¹ K⁻¹) and dominated by phonons because the network was not electrically conductive (Figure 25a).⁹ This value matches quite well with the values reported by Huang et al., ≈ 10 W m⁻¹ K⁻¹ for undoped electronically sorted SWCNT networks and ≈ 16 W m⁻¹ K⁻¹ for an undoped and unsorted SWCNT network (Figure 25b).²⁷⁴ Avery et al. found that doping the network dramatically lowered the thermal conductivity to $k = 2 - 4.5$ W m⁻¹ K⁻¹.⁹ For optimally doped polymer-free s-SWCNT networks, this relatively low thermal conductivity has led to the realization of $zT$ values in the range of ca. $zT = 0.01–0.12$.⁹,²⁵² This study and the subsequent follow-up demonstrated that there is a small, but real, electronic contribution to the thermal conductivity because $k$ increases weakly with carrier density but in a fashion that is
<figure><img src="image_39.png" /><figcaption>**Figure 25.** (a) Dependence of thermal conductivity on electrical conductivity for redox-doped s-SWCNTs at T = 298K. The trend line in the lower portion of the plot shows the dependence expected for the electrical component of thermal conductivity ($κ_e$) by the Wiedemann-Franz law. Reproduced with permission from ref 9. Copyright 2017 Springer Nature. (b) Dependence of thermal conductivity for SWCNT networks on the ratio of semiconducting and metallic SWCNTs (red circles). Also shown is the value measured for an unsorted network (black square). Reproduced with permission from ref 274. Copyright 2019 Taylor & Francis. (c) TEM images of unfilled SWCNTs and DWCNTs, along with CNTs filled with fullerenes. (d) Thermal conductivity ($κ$) and (e) thermopower ($S$) of fullerene-encapsulated CNT bundles measured at T = 300 K as a function of the CNT bundle diameter. The dashed lines show the average $κ$ and $S$ values of SWCNTs measured between different samples. (c–e) Reproduced with permission from ref 276. Copyright 2017 Springer Nature.</figcaption></figure>

<figure><img src="image_40.png" /><figcaption>**Figure 26.** Room temperature thermopower (a) and power factor (b) for polymer-free ~1.3 nm diameter laser vaporization (LV) s-SWCNT networks, as a function of room-temperature electrical conductivity. SWCNTs were originally dispersed with the cleavable SMP polymer, which was then removed in the spray-deposited film. Gray squares show data points for progressive doping of a single s-SWCNT network, while colored circles show data taken for samples that were measured with temperature-dependent thermopower and conductivity measurements. The blue shaded region shows the range where the Fermi level crosses into the first valence level ($v_1$) and the sample transitions from nondegenerate to degenerate. The red solid line shows the transport model used to simulate the Seebeck and $PF$ vs $σ$. Reproduced with permission from ref 253. Copyright 2019 Wiley.</figcaption></figure>

generally consistent with the Wiedemann-Franz law (Figure 25a).⁹,²,⁵,²⁶

Several other studies have explored the impact of molecular dopants, adsorbates, and encapsulants on the thermal conductivity of CNTs, and the effects seem to be mixed and may be complex.²⁷,⁵,²⁷ For example, doping was observed to reduce $k$ for acid-doped highly aligned SWCNT fibers, although iodine doping did not change $k$.²⁷⁵ Komada et al. found that fullerene encapsulation into SWCNTs and DWCNTs reduced $k$ by around 35%, while increasing $S$ by 40% (Figure 25c,e).²⁷⁶

Most recently, an interesting region was observed in the temperature-dependent thermal conductivity that suggests that the phonon contribution to $k$ ($k_{ph}$) can be completely eliminated within a narrow temperature range.²⁷⁷ Specifically, $k_{ph}$ reaches a minimum near 0 at T = 200 K for ~1.0 nm diameter HiPCO s-SWCNTs, whereas $k_{ph}$ steadily rises for larger diameter (~1.35 nm) plasma torch (PT) between 100 and 350 K. The source of this suppression of $k_{ph}$ is currently unclear but is likely due to a T-dependent phonon scattering mechanism. Such phonon scattering could arise from tube-tube interactions or
<figure><img src="image_41.png" /><figcaption>**Figure 27.** (a) Relationship between Seebeck coefficient and (b) power factor and conductivity of five samples with different semiconductor-metal ratios (table on right). The dashed lines in (a) are theoretical simulations. Mott's formula was used for sample 4 (Mott Cal.), whereas the linear response theory combined with the thermal Green's function method within the constant relaxation time approximation was used for samples 1–3 (LR Cal.). See reference for fitting details. Reproduced with permission from ref 248. Copyright 2019 American Chemical Society.</figcaption></figure>

interactions between the s-SWCNTs and exohedral or endohedral dopants and/or solvent molecules. Elucidation of this mechanism in future studies could lead to significant improvements in zT.

Mechanistically, the studies discussed above suggest a few ways forward for improving the performance of highly enriched s-SWCNT thermoelectrics. First, the morphology of the s-SWCNT appears to play a crucial role in transport. Removal of the insulating dispersion polymer improves carrier mobility, conductivity, power factor, and zT,²⁵²,²⁶⁵ and larger s-SWCNT bundles appear to reduce TE performance.²⁵² Molecules adsorbed on either exterior or interior surfaces of the SWCNTs dramatically impact conductivity and thermopower via doping and may modulate the thermal conductivity. Understanding the impacts on thermal conductivity may prove to be particularly fruitful, as it appears that the dominant phonon contribution can be dramatically reduced in some situations.

All of these observations point to a TE transport mechanism that depends heavily on the transport barriers present at the junctions between bundles, a conclusion that is in agreement with recent studies on polymer-free s-SWCNT networks²⁵³ that adapted the transport model originally used by Kang et al. to describe TE transport in semiconducting polymers.²⁷⁸ Specifically, the red lines in Figure 26 show the s = 1 relationship determined by this model that is expected for “intrinsic” transport properties of crystalline semiconductors with near-parabolic band dispersion near the band edge, where transport is limited by deformation potential scattering via acoustic phonons. This transport mode appears to describe the data on s-SWCNT networks well when the carrier density is high and the network is degenerately doped. The deviation from the “intrinsic” s = 1 behavior at low carrier density suggests that extrinsic factors limit transport and the ultimate PF achievable in the doped networks in this regime. It is quite possible that these “extrinsic” factors correspond to some of the morphological factors mentioned above that limit carrier transport from bundle-to-bundle in disordered s-SWCNT networks. Figure 26 demonstrates that identifying these extrinsic factors and mitigating them may provide a pathway toward the realization of substantially higher PF values in doped nanotube networks.

Second, it is critical to tune and optimize the doping level in s-SWCNTs to optimize the TE performance. This conclusion is reinforced by the recent study of Huang et al., who found a

relatively small difference between the TE performance of as-produced mixed metal/semi SWCNT thin films and highly enriched thin films in the absence of doping,²⁷⁴ reinforcing the need for carrier density control. Effective doping primarily requires finding dopants with the right balance between redox potential and environmental stability, but recent studies suggest morphology may play a role as well. Nonoguchi et al. found that the effectiveness of doping s-SWCNT networks depended on the SWCNT thickness and the dispersing polymer, and obtained a peak PF of 412 µW m⁻¹ K⁻² for ca. 200 nm thick films.²⁵⁴ Doping also appears to lower the thermal conductivity,⁹,²⁵² although further studies are needed to determine the precise mechanism at play. With respect to stability, s-SWCNT networks tend to be less stable to environmental degradation of n-type conductivity than their mixed and metallic counterparts. This sensitivity to ambient conditions appears to be an issue for both chemically doped s-SWCNTs²⁵² and s-SWCNTs doped via atomic substitution (e.g., N).²⁷⁹ Looking forward, it will be critical to further explore the mechanism(s) for this sensitivity and strategies for mitigating it. Encapsulation is certainly one strategy,²⁵² but it is possible that dopants with better inherent stability can also be developed.

While a number of research reports have now explored the potential TE benefits of enriched s-SWCNTs, very recent research suggests that the exploration of highly enriched m-SWCNTs may also prove fruitful. Ichinose et al. recently compared the carrier density-dependent TE power factor of SWCNT networks as a function of diameter, alignment, and metal/semi enrichment (Figure 27).²⁴⁸ They found that the typical behavior found for s-SWCNT networks, where S decreases continuously with increasing σ, does not hold for m-SWCNT networks. Instead, S increases continuously with increasing σ due to the increasing curvature in the DOS as E<sub>F</sub> is swept toward the first vHS via electrostatic doping (Figure 27a). The authors study samples with relatively small average diameter (ca. 0.7 nm) and large diameter (ca. 1.4 nm), in either unaligned or aligned films, and find the largest PF of ca. 300 µW m⁻¹ K⁻² for aligned networks of large diameter m-SWCNTs (Figure 27b).

It is encouraging to see that m-SWCNTs are not universally detrimental to TE performance and that enriched m-SWCNT networks can enable large TE PFs, exciting and somewhat unexpected findings that warrant more detailed fundamental
studies. Because alignment enhances the PF (measured in the alignment direction) by a factor of 4–5 compared to randomly aligned SWCNTs,²⁴⁶ the m-SWCNT sample is prepared from aligned, large-diameter arc discharge SWCNTs that were enriched by density gradient ultracentrifugation (DGU).²⁴⁸ Therefore, it is illuminating to compare aligned networks of enriched m-SWCNT and s-SWCNT prepared from similar large-diameter source material as a function of carefully controlled carrier density. Detailed comparisons of the thermal conductivity of m- and s-SWCNT networks will also be critical because little research exists for the former. It is also interesting to note that neither highly enriched m- or s-SWCNTs have been studied extensively within SWCNT-based TE composites. These studies should provide much-needed insight into the ultimate potential of highly enriched SWCNTs in organic thermoelectric devices and applications.

#### 4.2.2.3. Covalent Functionalization and Doping.
Beyond the noncovalent doping strategies discussed in section 4.2.2.1, a number of groups are also exploring covalent routes toward functionalizing and doping CNT networks to improve their TE performance. These strategies can be separated into two primary categories: (1) CNT sidewall functionalization with covalent defects and/or functional groups to modify carrier/phonon transport, and (2) substitutional doping of electron-deficient or electron-rich atoms into the CNT lattice to tune carrier density. In an early study belonging to the first category, Zhao et al. performed an Ar-plasma treatment to induce covalent defects within a free-standing bulky CNT paper.²⁸⁰ The authors reported that an optimized Ar-plasma treatment led to a significant enhancement of the Seebeck coefficient up to ~350 µV K⁻¹ and a reduction of the thermal conductivity to as low as 0.3 W m⁻¹ K⁻¹, leading to a zT value about 0.4 at 670 K (60 times higher than the pristine CNT paper).²⁸⁰ Recently, Nonoguchi et al. examined the effect of UV/ozone oxidation on the TE performance of s-SWCNT networks (Figure 28).²⁸¹ The authors suggest that the oxidative treatment introduced epoxy and carbonyl groups onto the s-SWCNT sidewall, which in turn increased the carrier concentration (Figure 28a). The enhanced Seebeck coefficient of the defective s-SWCNTs (Figure 28b) improved the power factor relative to similarly prepared s-SWCNT networks that were chemically doped with AgTFSI (Figure 28c).²⁸¹

A number of groups have also recently explored the potential of substitutional atomic doping for tuning the carrier density and providing more stable TE properties relative to those of chemically doped SWCNT networks. This strategy is analogous to the aliovalent doping strategies typically used in bulk covalent semiconductors such as silicon, where the incorporation of electron-deficient (e.g., boron) or electron-rich (e.g., phosphorus) leads to p- and n-type doping, respectively. There is a huge body of literature devoted to studying atomic substitutional doping in SWCNTs,²⁸²⁻²⁸⁴ but only recently has this approach been applied extensively to the TE properties of CNT networks. Liu et al. utilized an ammonia plasma to dope s-SWCNTs with nitrogen atoms and realize n-type TE transport.²⁷⁹ Before doping the s-SWCNT networks had a thermopower of 125 µV K⁻¹ and a power factor of 95 µW m⁻¹ K⁻². N-doped films showed a thermopower of -80 µV K⁻¹ in air at 110 °C, although they still behaved as p-type semiconductors under ambient conditions at room temperature.²⁷⁹ Dorling et al. use acetonitrile to introduce N atoms into a MWCNT synthesis to produce N-doped MWCNTs.²⁸⁵ The authors find n-type transport (negative thermopower) in the neat N-doped

<figure><img src="image_42.png" /><figcaption>**Figure 28.** (a) Schematic of UV/ozone treatment to induce oxygenic functional groups on sidewalls of s-SWCNTs. (b) Conductivity and Seebeck coefficient and (c) power factor of s-SWCNTs as a function of UV irradiation time. The dashed line in (c) references a chemically doped s-SWCNT network. Reproduced with permission from ref 281. Copyright 2019 AIP.</figcaption></figure>

MWCNTs, although composite films with P3HT have positive thermopower for MWCNT contents less than 40%. Interestingly, some p-type P3HT composites with low N-MWCNT content (e.g., 30%) can be converted to n-type by UV irradiation, an effect attributed to modification of the polymer microstructure and concomitant interactions with MWCNTs.²⁸⁶ The same group doped SWCNTs with N atoms via ammonolysis, but oxygenic functional groups induced by acid-based purification prevented the realization of a negative thermopower.²⁸⁵

Chiang et al. developed a wet chemistry-based strategy to produce boron-doped (mixed electronic structure) SWCNTs with boron content in the range of 0–3 atomic % (Figure 29).²⁸⁷ The authors found that boron doping improved both the conductivity and thermopower of the SWCNT networks, ultimately leading to a 6-fold improvement of the power factor up to ca. 60 µW m⁻¹ K⁻² at boron concentrations in the range of 2–3 atomic% (Figure 29f). Chemical doping of the B-doped SWCNTs led to an additional improvement in power factor up to ca. 100 µW m⁻¹ K⁻² (Figure 29g).²⁸⁷

#### 4.2.2.4. Transport Anisotropy in TE Properties.
SWCNTs are prototypical one-dimensional semiconductors with ultrahigh aspect ratios because a typical as-prepared individual nanotube has a diameter in the range of 0.6–2.0 nm and a length of >1 µm. As such, the structural, electronic, and optical properties of nanotubes are highly anisotropic. Recent studies have delved
<figure><img src="image_43.png" /><figcaption>**Figure 29.** (a–e) Possible boron configurations in a boron-doped SWCNT. (a) Borabenzene. (b) Borole. (c) Substituted boron. (d) C-BH-C. (e) C-BH₂. (f) Maximum power factors measured in the range of 37–200 °C, as a function of boron concentration. (g) Power factor of chemically doped B-SWCNTs (3.15 atomic % B) as a function of Fermi level shift (ΔE) at room temperature. Reproduced with permission from ref 287. Copyright 2019 American Chemical Society.</figcaption></figure>

into the implications of this anisotropy on TE properties at the macroscopic level for highly aligned thin films of SWCNTs.<sup>246,247,288</sup> Fukuhara et al. used slow vacuum filtration of unsorted arc discharge SWCNT dispersions to produce highly aligned SWCNT thin films and used electrolyte gating to tune the Fermi level in both the p- and n-type doping regimes (Figure 30).<sup>246</sup> They found that the conductivity along the alignment direction was several times larger than conductivity in the direction perpendicular to alignment (Figure 30b). In contrast, the Seebeck coefficient was isotropic, with a peak magnitude of ca. 50 µV K<sup>-1</sup> for appropriately doped p- and n-type SWCNTs for either parallel or perpendicular transport (Figure 30c). The conductivity anisotropy led to a parallel PF (ca. 160 µW m<sup>-1</sup> K<sup>-2</sup>), roughly four times that of the perpendicular PF (ca. 40 µW m<sup>-1</sup> K<sup>-2</sup>) (Figure 30d). Consistently, as discussed above, Ichonese et al. found that the TE power factor of highly aligned arc discharge m-SWCNT (>96% m-SWCNT purity) thin films was ca. five times higher than that of randomly aligned m-SWCNT (>99% m-SWCNT) thin films.<sup>248</sup>

Yamaguchi et al. studied the thermal conductivity of unsorted arc discharge SWCNT thin films that were prepared via vacuum filtration to be either highly aligned, poorly aligned, or completely random.<sup>288</sup> The thermal conductivity of the poorly aligned film was 28 ± 1.2 W m<sup>-1</sup> K<sup>-1</sup>, while that of the randomly aligned film was 14 ± 2.8 W m<sup>-1</sup> K<sup>-1</sup>, in a similar range to that found by a number of other groups for undoped randomly aligned SWCNT networks.<sup>9,274</sup> For the highly aligned film, they found a parallel thermal conductivity, k<sub>||</sub>, of 43 ± 2.2 W m<sup>-1</sup> K<sup>-1</sup> and a perpendicular thermal conductivity, k<sub>⊥</sub>, of 0.085 ± 0.017 W m<sup>-1</sup> K<sup>-1</sup>, for a thermal conductivity anisotropy of ca. 500. The large differences in thermal conductivity were explained by the different mechanisms controlling phonon transport (the dominant source of thermal conductance) in the different films. Specifically, the authors' analysis suggests that k<sub>⊥</sub> in the highly aligned film is dominated by the intertube thermal resistance, whereas k<sub>||</sub> is dominated by the internal thermal conductance of a SWCNT in the axial direction. In poorly aligned and random films, the thermal conductance is weighted by the contributions of both internal thermal conductance and intertube thermal resistance.

Nonoguchi compared the in-plane and through-plane TE performance of mixed SWCNTs doped with crown ether-stabilized alkali metal salts.<sup>237</sup> The KOH/18-crown-6-ether doped CNT film had an enhanced in-plane PF of 220 µW m<sup>-1</sup> K<sup>-2</sup> and through-plane PF of 27 µW m<sup>-1</sup> K<sup>-2</sup>, but because of the suppressed through-plane thermal conductivity (0.12 W m<sup>-1</sup> K<sup>-1</sup> compared to 39 W m<sup>-1</sup> K<sup>-1</sup>), the through-plane zT (0.07) value was more than 3 times higher than in-plane zT (0.02) at 310 K.

### 4.2.3. TE Generators from CNT-based Fibers and Fabrics.
The inherently threadlike nature of CNTs is an important advantage for thermoelectric energy harvesting, especially from the standpoint of wearable, flexible, and lightweight TE generators. As such, CNTs can be easily fabricated by a variety of means into fabric-like form factors. A particularly attractive aspect of these strategies arises from the relatively gentle processing methods that can be applied to as-grown CNT samples. Because these processing routes do not typically involve ultrasonication to produce surfactant- or polymer-based dispersions, they naturally leave the CNTs as long as possible, which dramatically improves network conductivity. These fabrics also often (but do not always) tend also to have well-aligned CNTs, improving the conductivity as discussed in section 4.2.2.4. Yarns and fibers can be directly spun from as-produced CNT “webs” or “socks”, often straight from the reactor in a continuous process and typically via pulling

<figure><img src="image_44.png" /><figcaption>**Figure 30.** (a) Schematic of highly aligned SWCNT film. (b–d) Comparison of electrical conductivity (b), Seebeck coefficient (c), and TE power factor (d) measured for electrostatically doped SWCNT samples in the direction parallel or perpendicular to SWCNT alignment. Reproduced with permission from ref 246. Copyright 2018 AIP.</figcaption></figure>
<figure><img src="image_45.png" /><figcaption>**Figure 31.** (a) Schematic of the fabrication of DWCNT yarn via direct spinning from FCCVD reactor onto a PDMS support. $FeCl_3$ and PEI were used to dope the frontside p-type and backside n-type, respectively. Reproduced with permission from ref 230. Copyright 2017 American Chemical Society. (b) Schematic of wet-spinning process from chlorosulfonic acid for fabrication of highly aligned CNT fibers. (c) Schematic of solvent-based process for forming either densified SWCNT films or low-density SWCNT foams. Reproduced with permission from ref 262. Copyright 2019 Wiley. (d) Schematic (left) and photograph (right) of $\pi$-type coil TEG formed by wrapping CNT yarn around PET spacer and brush-doping p- and n-type dopants. Reproduced with permission from ref 292. Copyright 2019 Royal Society of Chemistry. (e) Schematic (left) and photograph (right) of $\pi$-type sheet TEG formed by selective-area brush-doping of dopants onto continuous CNT sheet that is snaked through a PDMS support. Red and blue dots represent n-type and p-type dopants. (f) Schematic of out-of-plane Z-type sheet TEG, where each sheet is separated by a PDMS spacer (light blue) and junctions are formed by deposition of silver paste. Red and blue dots represent n-type and p-type dopants. (e–f) Reproduced with permission from ref 269. Copyright 2020 American Chemical Society.</figcaption></figure>

<figure><img src="image_46.png" /><figcaption>**Figure 32.** (a) Schematic of p/n CNT yarn (CNTY) doped p-type with PEDOT:PSS and n-type with PEI, after which it is twisted around PET filaments. (b) Schematic of out-of-plane thermoelectric textile (TET) fabricated by knitting CNT-twisted PET filaments ($\pi$-type knit). (c) Photographs of the flexible knitted TET. (a–c) Reproduced with permission from ref 270. Copyright Royal Society of Chemistry. (d) Schematic of fabrication process for CNT fibers doped alternately p- and n-type with PEDOT:PSS and oleanine, respectively, after which the CNT fibers are wrapped with acrylic fibers. (e) Schematic (left) and photographs (right) of out-of-plane knitted TET prepared from CNT/acrylic fibers ($\pi$-type knit). (d–e) Reproduced with permission from ref 50. Copyright 2020 Royal Society of Chemistry.</figcaption></figure>

through a liquid densification bath and rolling onto a roller (Figure 31a).²³⁰ Vertically aligned CNT arrays can also be dry-spun into fibers.²⁸⁹ Mild liquid-phase processing can be achieved by producing slurries in solvents (with no surfactants and no sonication) or by producing true solutions in strong acids such as chlorosulfonic acid.²⁴⁷,²⁷⁵ The latter technique can be paired with a coagulation-bath-assisted spinning process to produce highly aligned CNT fibers with exceptional alignment, strength, and conductivity (Figure 31b).²⁹⁰ As-produced web-like

material can also be densified into disordered or aligned fabric-like sheets that can be stacked, cut, and shaped by a variety of means and into a variety of form factors. Alternatively, web-like material can be expanded into flexible and moldable foams with very low density (Figure 31c).²⁶²

CNT fibers, yarns, and sheets can be fabricated into TE generators with a variety of shapes and form factors, but there are two predominant styles that have been explored for CNT TEGs. The “Z-type” or “accordion-style” TEGs place flat p- and n-type
<table><caption>Table 2. Summary of Fabrication Methods and Dopants for π-Type CNT-based TEGs</caption><thead><tr><th>CNT synthesis and (type)</th><th>fabric style</th><th>TEG style</th><th>p-type dopant</th><th>p-type dopant</th><th>peak p-type PF (μW m<sup>-1</sup> K<sup>-2</sup>)</th><th>peak n-type PF (μW m<sup>-1</sup> K<sup>-2</sup>)</th><th>key module metric(s)</th></tr></thead><tbody><tr><td>FCCVD (FWCNTs)</td><td>aligned sheet</td><td>π-type (sheet)</td><td>FeCl<sub>3</sub></td><td>PEI</td><td>250</td><td>255</td><td>23 mV<br>8 μW</td></tr><tr><td>FCCVD</td><td>aligned sheet</td><td>Z-type</td><td>FeCl<sub>3</sub></td><td>BV</td><td>2500</td><td>1100</td><td>0.66 μW cm<sup>-2</sup> (5.3 μW)</td></tr><tr><td>FCCVD</td><td>disordered sheet</td><td>Z-type</td><td></td><td>PEI</td><td>1800</td><td>1500</td><td>2.5 μW</td></tr><tr><td>FCCVD</td><td>yarn</td><td>π-type (knit)</td><td>PEDOT:PSS</td><td>PEI</td><td>513</td><td>668</td><td>500 V m<sup>-2</sup><br>50 mW m<sup>-2</sup> (383 μW)</td></tr><tr><td>DWCNT</td><td>wet-spun fibers</td><td>π-type (coil)</td><td>H-carbazole</td><td>Dpp</td><td>432</td><td>323</td><td>15.4 μW g<sup>-1</sup> (0.12 μW)<br>259 μW g<sup>-1</sup> (2.02 μW)</td></tr><tr><td>FCCVD</td><td>twisted fibers</td><td>π-type (knit)</td><td>PEDOT:PSS</td><td>oleamine</td><td>330</td><td>320</td><td>70 mW m<sup>-2</sup> (4.5 μW)</td></tr><tr><td>TUBALL (SWCNT)</td><td>solvent-prepped sheet</td><td>π-type (coil)</td><td></td><td>PEI</td><td>411</td><td>90</td><td>4.6 μW cm<sup>-2</sup> (10.3 μW)<br>2.0 μW cm<sup>-2</sup> (4.4 μW)<br>0.5 μW cm<sup>-2</sup> (1.1 μW)</td></tr><tr><td>FCCVD (DWCNT)</td><td>yarn</td><td>π-type (coil)</td><td>FeCl<sub>3</sub></td><td>PEI</td><td>2387</td><td>2456</td><td>1.2 V g<sup>-1</sup> (22.5 mV)<br>700 μW g<sup>-1</sup> (1.3 μW)<br>45 mV<br>4.2 μW</td></tr><tr><td>FCCVD</td><td>disordered sheet</td><td>π-type (in-plane)</td><td>Au/PANI</td><td>PEI</td><td>2454</td><td>1707</td><td>0.4 μW cm<sup>-2</sup><br>550 μW g<sup>-1</sup> (1.7 μW)</td></tr><tr><td>eDIPS (SWCNT)</td><td>porous foam</td><td>π-type (foam)</td><td>FeCl<sub>3</sub></td><td>BV</td><td>1.1</td><td>1.0</td><td>0.5 μW cm<sup>-2</sup><br>82 μW g<sup>-1</sup> (1.5 μW)</td></tr><tr><td>MWCNTs</td><td>coated PET yarn</td><td>π-type (in-plane)</td><td>PEDOT:PSS (separate leg)</td><td>PEI</td><td></td><td></td><td>143 mV<br>7.1 nW</td></tr></tbody></table>

CNT sheets in series with a flat insulating spacer in between each sheet (Figure 31f).<sup>269,291</sup> Conducting metal pastes may be used to “glue” isolated p-type and n-type sheets together and establish good contact at each p/n junction, or alternatively a continuous sheet may be alternately doped with different chemical dopants. The conductive nature of the CNTs allows for the fabrication of metal-free TEGs. When expanded, such TEGs resemble an accordion. Once collapsed, a thermal gradient can be established in the out-of-plane direction (with respect to the planes of individual p/n sheets and the planes of hot- or cold-side temperature sinks) to produce voltage and power via transport of carriers.

Another broad category of TEG commonly fabricated from CNTs is the π-type TEG, where p-n legs are arranged in such a fashion as to have their primary transport plane in the same direction as the thermal gradient (Figure 31d,e).<sup>269,292</sup> In most

cases, the SWCNT fiber/sheet transport plane is perpendicular to the planes of substrate or fabric into which it is integrated (out-of-plane geometry). This geometry of TEG is commonly used for “conventional” TEGs based on bulk inorganic semiconductors. However, in some cases, the SWCNT fiber/ sheet transport plane lies along the plane of a substrate or fabric into which it is integrated (in-plane geometry). The out-of-plane geometry tends to maximize the areal density of p/n legs for enhanced power density (i.e., W cm<sup>-2</sup>). Recent examples of π-type SWCNT-based TEGs based solely on doped SWCNTs utilize a continuous sheet geometry (π-type sheet, Figure 31e),<sup>269</sup> a coiled geometry in which yarns or sheets are coiled around an insulating spacer (π-type coil, Figure 31d),<sup>230,268,292</sup> in-plane substrate- or fabric-mounted geometries (π-type in-plane),<sup>231,293</sup> and a TEG based on PDMS-separated SWCNT foam blocks (π-type foam).<sup>262</sup> SWCNT fibers or yarns can also
<figure><img src="image_47.png" /><figcaption>**Figure 33.** (a) Scanning tunneling microscopy (STM) image of monolayer graphene sheet. Reproduced with permission from ref 297. Copyright 2012 IOP Publishing. (b) Real-space lattice, (c) Reciprocal-space lattice, and (d) Energy band structure of graphene showing Dirac cones meeting at K and K' points (white and black circles). (e) Zoom of K-point where linear electronic bands cross at the Dirac point. (d,e) Reproduced with permission from ref 297. Copyright 2012 IOP Publishing. (f) STM image of chemically grown graphene nanoribbons (GNRs). Blue and white overlays show cartoons of atomic structure for the GNRs. Reproduced with permission from ref 298. Copyright 2010 Springer Nature. (g) Experimental bandgaps for GNRs as a function of width. Reproduced with permission from ref 299. Copyright 2007 American Physical Society. (h) TEM image of monolayer graphene flake, exfoliated in NMP solvent. Scale bar is 500 nm. Reproduced with permission from ref 300. Copyright 2008 Springer Nature. (i) Cross-sectional SEM image of restacked reduced graphene oxide (rGO) thin film. Reproduced with permission from ref 301. Copyright 2007 Springer Nature.</figcaption></figure>

be combined with other conventional fibers (e.g., acrylics, PET, etc.) into integrated energy-harvesting textiles (π-type knit, Figure 32).⁵⁰,²⁷⁰ The p- and n-type doping in these π-type TEGs is typically achieved by soaking, spraying, or brushing chemical dopants onto targeted areas to form periodic repetitions of p-n junctions.

Table 2 summarizes the fabrication methodology and dopants used to fabricate a variety of Z-type and π-type CNT-based TEGs, along with the optimized performance metrics. Where reported, performance metrics such as peak open circuit voltage and peak power are provided as densities normalized to either area or mass of the TEG. Mass-normalized voltage and power density is one metric where CNTs compare especially well to their inorganic counterparts because carbon is typically substantially lighter than the elements used in compound semiconductors.

**4.2.4. Graphene.** Graphene is a Dirac semimetal with linear bands for electrons and holes that cross at the Dirac point (Figure 33). As such, it serves as a fascinating model system for studying the thermoelectric properties of massless Dirac fermions that result from linear band dispersion. The Fermi level or chemical potential of graphene can be modulated in a similar fashion to CNTs, by employing molecular doping, substitutional doping, and electrostatic doping. The reports of experimentally measured *zT* values for graphene are still rare, but many theoretical studies have been carried out to explore its thermoelectric properties and several recent experimental studies have begun publishing *PF* and *zT* values. A few recent review articles have been published on graphene-based thermoelectrics,²⁹⁴⁻²⁹⁶ so here we highlight many of the salient advances in the past decade.

It is important to point out that the term “graphene” is used to refer to a broad family of materials. At the purest level, monolayer graphene is a single sheet of hexagonally bonded sp²-

hybridized carbon atoms (Figure 33a). Monolayer graphene can be exfoliated from graphite or grown on copper foil by chemical vapor deposition (CVD). Monolayer graphene can be further nanostructured, most typically by the growth or cutting of a graphene sheet into nanoribbons (Figure 33f), where the narrow width of the graphene sheet confines charge carriers and opens up a band gap (Figure 33g). Graphene can also be exfoliated from graphite into solution as small flakes while retaining its identity as graphene (Figure 33h). Another distinct method of exfoliation involves the production of graphene oxide (GO) or reduced graphene oxide (rGO) by first oxidizing graphite, then exfoliating the GO, and finally reducing the oxide functional groups to produced rGO. Solution-exfoliated flakes of graphene, GO, or rGO can be restacked to form thick films (Figure 33i) that can be used directly in TE devices. Here we divide our discussion of graphene TE properties by focusing on these unique classifications accordingly.

#### 4.2.4.1. Monolayer Graphene.
Early work related to thermoelectric properties of monolayer graphene focused on thermal conductivity, which in theory should be exceptionally large.³⁰²,³⁰³ Balandin et al. introduced a scanning Raman technique (Figure 34a) that utilizes the temperature coefficient of the graphene G band to estimate a thermal conductivity of 4840–5300 W m⁻¹ K⁻¹.³⁰² Because phonons contribute significantly to this large thermal conductivity, subsequent studies demonstrated a dramatic reduction in *k* for isotopically engineered monolayer graphene sheets (Figure 34b).³⁰⁴

The dependence of conductivity and thermopower on Fermi level or chemical potential in graphene has been treated theoretically with a number of different approaches. Figure 35a displays an early study applying Mott's formula to this question.³⁰⁵ Conductivity increases linearly away from the charge-neutrality point (CNP), as expected for a Dirac semimetal. The Seebeck coefficient is zero at the CNP and has
<figure><img src="image_48.png" /><figcaption>**Figure 34.** (a) Schematic of Raman microscopy measurement introduced by Balandin et al. to measure thermal conductivity of monolayer graphene suspended over a trench. Reproduced with permission from ref 302. Copyright 2008 American Chemical Society. (b) Thermal conductivity of isotopically engineered monolayer graphene measured by similar technique by Chen et al. (closed diamonds) as a function of ¹³C concentration. Open orange squares show molecular dynamics simulations for comparison. Reproduced with permission from ref 304. Copyright 2012 Springer Nature.</figcaption></figure>

maxima for both p- and n-type graphene for small shifts of the chemical potential away from the CNP (Figure 35b).³⁰⁵ Zuev et al. probed the carrier density dependence of $\sigma$ (Figure 35c) and $S$ (Figure 35d) for exfoliated mesoscopic graphene samples where the chemical potential was swept in positive and negative directions via electrostatic gating.³⁰⁶ The Seebeck coefficient was carrier-density dependent (Figure 35d) and scaled linearly with temperature (inset), with a peak value of ~ -80 µV K⁻¹ at room temperature.³⁰⁶ Both this study³⁰⁶ and subsequent studies³⁰⁷ find divergent/anomalous behavior for thermopower near the Dirac point at low temperatures based on the characteristic behavior of massless Dirac quasiparticles. Later, a similar Seebeck coefficient value was reported together with a thermal conductivity of ~600 W m⁻¹ K⁻¹ at room temperature for monolayer graphene.³⁰⁸

Kanahashi et al. used a similar electrostatic gating approach to realize exceptionally large *PFs* in CVD-grown monolayer graphene (Figure 35e,f).³⁰⁹ The Seebeck coefficient scaled inversely with conductivity as $S \propto \sigma^{-1}$, as expected by the Mott equation, and the *PFs* reached as high as 6930 µW m⁻¹ K⁻² (p-type) and 3290 µW m⁻¹ K⁻² (n-type) at 300 K.

#### 4.2.4.2. Graphene Nanoribbons and Nanostructured Graphene.
Different groups have used the nonequilibrium Green's function formalism to calculate the thermoelectric properties of GNRs. GNRs with a width of 1.8 nm and a band

gap of $E_g = 667$ meV were predicted to have Seebeck coefficients ca. 500 times larger than monolayer graphene, reaching peak values of ±4000 µV K⁻¹ (Figure 36a).³¹⁰ Defects and edge imperfections served to increase the GNR thermopower and decrease thermal conductivity, but *zT* decreased for defective GNRs because these changes were outweighed by a large decrease in electrical conductivity.³¹⁰ This reduction in *zT* contrasts with the behavior calculated by Sevencli et al. for edge-disordered zigzag graphene nanoribbons (ZGNRs).³¹¹ In this case, the authors found a large reduction in thermal conductivity, whereas the electrical conductivity decreased to a much smaller degree at chemical potentials within the bandgap of the ZGNR and a very sharp derivative of electrical conductance at the onset of the ZGNR conduction and valence bands produced large *S*. As a result, a *zT* value of as high as 4 at room temperature was predicted (Figure 36b).³¹¹ Chen et al. predicted that periodic lattices of GNRs could improve *zT* due to the beneficial effect of interfaces between GNR sections with different widths.³¹²

Nanopatterning of graphene into nanoribbon sections are also predicted as a route to improving thermoelectric performance. Kim and Grossman used classical and quantum mechanical calculations to study the effect of H and C5 functionalization on the TE properties of patterned graphene "nanoroads", where functionalization introduced peaks in the density of states at the band edge (Figure 36c,d).³¹³ Unlike conventional 3D TE materials, the increase of *PF* and the reduction of thermal conductivity was achieved simultaneously, resulting in a high theoretical *zT* value of ~3 at room temperature (Figure 36d).

Recent experimental evidence provides support for the predicted enhancement of TE performance for GNRs and nanostructured graphene, relative to unstructured graphene. Graphene nanomeshes with GNR features having widths less than 10 nm were shown to have thermal conductivity reduced to as low as ~78 W m⁻¹ K⁻¹ and *S* as high as -520 µV K⁻¹, although the electrical conductivity was reduced as a result of the nanostructuring.³¹⁴ Recent studies on CVD-grown and suspended GNRs (Figure 37a) with 250 nm length and 40 nm width³¹⁵ demonstrated several-fold enhancement of *S* relative to unstructured graphene and a large improvement of the ratio of electrical to thermal conductivity ($\sigma/k$) relative to prior work.³⁰⁸,³¹⁴,³¹⁶ These enhancements ultimately resulted in a *zT* value of 0.12 at $120 < T < 220$ K (Figure 37b).³¹⁵

#### 4.2.4.3. Exfoliated Graphene Flakes and Reduced Graphene Oxide.
Graphene can also be solution-processed when exfoliated into small flakes in solvent. A wide variety of methods can be used, but sonication and centrifugation remains the most popular.³¹⁷ The primary advantages of these solution-dispersed graphene flakes is that the dispersions can be deposited by any number of solution deposition techniques to produce thick films of electrically percolated graphene flakes, and the dispersions allow for easy integration into composites with other organic or inorganic species. These thicker films and composites are more representative of the form factors needed for TE devices because it would be challenging to incorporate monolayer graphene and graphene NRs into macroscopic devices that can generate appreciable power densities. In the dispersion process, graphene is often oxidized in a modified Hummer's method to produce graphene oxide (GO), which is typically p-type because of the oxygenic groups. GO is typically chemically reduced (e.g., with hydrazine) to produce reduced graphene oxide (rGO), but the oxygenic functional groups are difficult to fully remove and the resulting rGO is still typically weakly p-type.³¹⁸⁻³²⁰ The oxygen
<figure><img src="image_49.png" /><figcaption>**Figure 35.** (a,b) Theoretical calculation of the low temperature (50 K) electrical conductivity (a) and Seebeck coefficient (b) of monolayer graphene as a function of chemical potential ($\mu$), where the chemical potential is set to zero at the CNP of graphene. Reproduced with permission from ref 305. Copyright 2007 American Physical Society. (c,d) Experimental measurement of temperature-dependent electrical conductivity (c) and Seebeck coefficient (d) of monolayer graphene as a function of applied gate voltage ($V_g$), where $V_g$ electrostatically modulates the chemical potential. Inset in (c) shows an SEM image of a typical device (2 µm scale bar), and inset in (d) shows temperature-dependent Seebeck coefficients taken at $V_g = -30$ V (square) and $-5$ V (circle). Dashed lines are linear fits to data. Reproduced with permission from ref 306. Copyright 2007 American Physical Society. (e) Seebeck coefficient as a function of electron and hole conductivities for large-area electrostatically doped monolayer graphene. Black trend lines show the $\sigma^{-1}$ dependence for $S$ expected from the Mott equation. (f) PF for electrostatically doped monolayer graphene. (e,f) Reproduced with permission from ref 309. Copyright 2019 Springer Nature.</figcaption></figure>

<figure><img src="image_50.png" /><figcaption>**Figure 36.** (a) Calculated thermopower of a perfect $n = 15$ armchair GNR as a function of chemical potential ($\mu$) at different temperatures: $T = 75$ K (light-blue squares), $T = 150$ K (green triangles), $T = 225$ K (red circles), $T = 300$ K (dark-blue line). The dashed-blue lines indicate the first conduction ($E_c$) and valence ($E_v$) sub-band edges. Inset shows a zoom for $\mu > E_c$. Reproduced with permission from ref 310. Copyright 2009 American Institute of Physics. (b) Calculated $zT$ versus chemical potential ($\mu$) at different temperatures for zigzag GNR with length of 4 µm. Zero $\mu$ is the charge-neutrality point, and the dotted vertical line represents the first conduction sub-band edge. Reproduced with permission from ref 311. Copyright 2010 American Physical Society. (c) Schematic of TE generator based on patterned graphene nanostructures with armchair boundary. Yellow shaded area indicates the patterned region. (d) H-functionalized graphene with a 6 Å pattern width (top) and C5-functionalized graphene (bottom). C and H atoms represented by gray and yellow spheres, respectively. Pink contour represents charge density of the conduction band. Green and blue spheres indicate C sites, where C5 chains are attached on the top and bottom surfaces of graphene, respectively. (e) Density of states per unit cell of H- and C5-functionalized graphene. (f) Calculated thermal conductivity (green), PF (orange), and $zT$ (red) of H- and C5-patterned graphene nanostructures at room temperature. Values are calculated along the pattern direction for $N_c = 8$ pattern width. (c–f) Reproduced with permission from ref 313. Copyright 2015 American Chemical Society.</figcaption></figure>
<figure><img src="image_51.png" /><figcaption>**Figure 37.** (a) Rapid-heating plasma-CVD method for synthesizing suspended GNRs along with 8-terminal Ni films. The 8-terminal test structure allows measurements of electrical and thermal conductivities and the Seebeck coefficient of GNRs on the same device. (b) Temperature-dependent *zT* of the GNRs produced by this rapid-heating plasma-CVD synthesis route (red spheres), as compared to other reports³⁰⁸,³¹⁴,³¹⁶ in the literature. Reproduced with permission from ref 315. Copyright 2019 American Chemical Society.</figcaption></figure>

functional groups also break up the sp² bonds and delocalized π electron density that is a highly desirable characteristic of graphene and typically limits the conductivity of rGO to ca. 10 S cm⁻¹.³²⁰ Li et al. demonstrated that dramatically increasing the reduction temperature to 3300 K substantially improves the

conductivity and TE performance of restacked rGO thin films.³²¹ These so-called high-temperature rGO (HT-rGO) thin films reached electrical conductivities in the range of 1500–4000 S cm⁻¹ (300 < *T* < 3000 K) and peak Seebeck coefficient of ca. 150 µV K⁻¹ around 1200 K, with a peak PF of 54.5 µW cm⁻¹ K⁻² (5450 µW m⁻¹ K⁻²) at 3000 K, making them attractive for high-temperature TE applications.³²¹

Novak et al. introduced a graphite-intercalation compound (GiC) strategy (Figure 38a) for producing nonoxidized graphene flakes (NOGF).³²³ The resulting restacked NOGF thin films were n-type, with Seebeck coefficient of −45.3 µV K⁻¹, electrical conductivity of 3280 S cm⁻¹, and PF of 673 µW m⁻¹ K⁻². Flexible devices were robust with respect to bending and produced 2.2 mV using body heat.³²³ Complementary n-type and p-type films could be prepared by changing the surfactant utilized in the intercalation-exfoliation process, with poly-vinylpyrrolidone (PVP) producing n-type NOGF and pyrene-butyric acid (PBA) producing p-type NOGF (Figure 38b).³²² Both types of NOGF films showed similar magnitudes of thermopower, conductivity, and PF (>600 µW m⁻¹ K⁻²), enabling an all-graphene TE device that produced 5.0 nW peak power at Δ*T* = 50 K (Figure 38c).³²²

### 4.2.5. Nanocarbon Composites.
Nanocarbons (i.e., carbon nanotubes, graphene, carbon quantum dots, etc.) have long been used as inclusions or “fillers” to improve the mechanical and electrical properties of composite materials.³²⁴,³²⁵ About a decade ago, attention turned to the exploitation of these unique nanoscale carbon systems in TE composites,³²⁶⁻³²⁹ where the primary focus has been on exploiting their impressive intrinsic carrier transport to enhance TE composites where the electrical conductivity of the host limits the overall TE performance (i.e., polymer-based hosts).

One of the primary motivations for the inclusion of nanoscale materials in TE composites is their potential to enhance the TE performance due to quantum effects that arise because of their size/shape and the nanoscale interfaces that are introduced with

<figure><img src="image_52.png" /><figcaption>**Figure 38.** (a) Schematic for intercalation/exfoliation procedure to produced nonoxidized graphene flakes (NOGF), where changing the surfactant modifies the majority carrier type. (b) Power factors achieved for restacked NOGF films comprised of p-type PBA-NOGF and n-type PVP-NOGF. (c) Voltage, current, and power for a device prepared from the n-type and p-type NOGF thin films. Reproduced with permission from ref 322. Copyright 2020 Wiley.</figcaption></figure>