Semiconducting large bandgap oxides as potential thermoelectric materials for high-temperature power...

INVITED PAPER

Semiconducting large bandgap oxides as potential thermoelectricmaterials for high-temperature power generation?

M. Backhaus-Ricoult • J. Rustad • L. Moore •

C. Smith • J. Brown

Received: 24 April 2014 / Accepted: 12 May 2014 / Published online: 25 June 2014

� Springer-Verlag Berlin Heidelberg 2014

Abstract Semiconducting large bandgap oxides are con-

sidered as interesting candidates for high-temperature

thermoelectric power generation (700–1,200 �C) due to

their stability, lack of toxicity and low cost, but so far they

have not reached sufficient performance for extended

application. In this review, we summarize recent progress on

thermoelectric oxides, analyze concepts for tuning semi-

conductor thermoelectric properties with view of their

applicability to oxides and determine key drivers and limi-

tations for electrical and thermal transport properties in

oxides based on our own experimental work and literature

results. For our experimental assessment, we have selected

representative multicomponent oxides that range from

materials with highly symmetric crystal structure (SrTiO3

perovskite) over oxides with large densities of planar crys-

tallographic defects (TinO2n-1 Magneli phases with a single

type of shear plane, NbOx block structures with intersecting

shear planes and WO3-x with more defective block and

channel structures) to layered superstructures (Ca3Co4O9

and double perovskites) and also include a wide range of

their composites with a variety of second phases. Crystal-

lographic or microstructural features of these oxides are in

0.3–2 nm size range, so that oxide phonons can efficiently

interact with them. We explore in our experiments the

effects of doping, grain size, crystallographic defects,

superstructures, second phases, texturing and (to a limited

extend) processing on electric conductivity, Seebeck coef-

ficient, thermal conductivity and figure of merit. Jonker and

lattice-versus-electrical conductivity plots are used to

compare specific materials and material families and extract

levers for future improvement of oxide thermoelectrics. We

show in our work that oxygen vacancy doping (reduction) is

a more powerful driver for improving the power factor for

SrTiO3, TiO2 and NbOx than heterovalent doping. Based on

our Seebeck-conductivity plots, we derived a set of highest

achievable power factors. We met these best values in our

own experiments for our titanium oxide- and niobium oxide-

based materials. For strontium titanate-based materials, the

estimated highest power factor was not reached; further

material improvement is possible and can be reached for

materials with higher carrier densities. Our results show that

periodic crystallographic defects and superstructures are

most efficient in reducing the lattice thermal conductivity in

oxides, followed by hetero- and homovalent doping. Due to

the small phonon mean free path in oxides, grain boundary

scattering in nanoceramics or materials with nanodisper-

sions is much less efficient. We investigated the impact of

texturing in Ca3Co4O9 ceramics on thermoelectric perfor-

mance; we did not find any improvement in the overall in-

plane performance of a textured ceramic compared to the

corresponding random ceramic.

1 Introduction

1.1 Potential value of oxides compared to other

thermoelectric materials: common advantages

and disadvantages

Thermoelectric power generators convert thermal energy

into electrical energy. Compared to other power generators,

they do not emit toxic gases, have long lifetime, low

operation and maintenance cost and can operate with waste

heat. Due to worldwide concerns about the increasing

demand of energy supply, limited availability of fossil fuel,

M. Backhaus-Ricoult (&) � J. Rustad � L. Moore � C. Smith �J. Brown

Corning Incorporated, Corning, NY 14831, USA

e-mail: [email protected]

123

Appl. Phys. A (2014) 116:433–470

DOI 10.1007/s00339-014-8515-z

impact of global warming, CO2 emission and human

health, green sources of energy become increasingly

attractive solutions. Thermoelectric power generation is

one option. Possible applications include recovery of waste

heat from industrial processes, power plants, incinerators,

melting furnaces, gas heaters, engine exhaust streams and

many others.

Let us consider the example of heat extraction from the

automotive exhaust gas: Currently, only roughly one-third

of the fuel energy is used for the translation of the car; one-

third leaves the car in the form of exhaust gas heat and one-

third heats the radiator fluid. Thermoelectric generators in

the EGR loop, in the exhaust stream after the exhaust gas

after-treatment system (DOC, DPF) or, closer to the

engine, directly at the exhaust manifold or on the engine

surface itself, are considered as a promising technology

that can contribute to lower fuel consumption and decrease

CO2 emission.

The conversion of thermal into electrical energy in

thermoelectric generators is based on the Seebeck

effect: If a semiconductor is exposed to a temperature

gradient, the temperature dependency of its carrier

concentration produces a potential difference across the

material that is proportional to the temperature differ-

ence. Suitable thermoelectric materials produce a large

thermopower (potential difference across the sample)

when exposed to a temperature gradient. They typically

show a strong dependency of their carrier concentration

on temperature and have high carrier density, high

carrier mobility and a low thermal conductivity. Pure

p-type (n-type) materials have only positive (negative)

mobile charge carriers, electron holes (electrons) and a

positive (negative) Seebeck coefficient. Most materials

have both, positive and negative charge carriers, and

may also have in addition ionic charge carriers, so that

the sign of the Seebeck coefficient depends on the

predominant carrier.

The thermoelectric conversion efficiency of a thermo-

electric generator depends on the figure of merit ZT of its

material

ZT ¼ TS2r=j ¼ TPF=j; ð1Þ

with T being temperature, S Seebeck coefficient or ther-

mopower, r electric conductivity, j thermal conductivity

and PF the power factor [1]. Materials for efficient ther-

moelectric generation should simultaneously exhibit large

S, large r and small j and behave as ‘‘an electron crystal

and a phonon glass.’’ Property optimization for a specific

material is difficult due to the coupling of the properties.

The key properties all depend on the carrier concentration;

electrical and thermal conductivity typically increase with

carrier density, while the Seebeck coefficient decreases

with carrier density. Best thermoelectric performance is

expected for heavily doped semiconductors with carrier

concentrations in the range of 1019–1021 mol-1 [1].

Low thermal conductivity in conjunction with high

electrical conductivity is difficult to achieve since heat is

not only carried via lattice vibrations, but also by the

electronic carriers:

j ¼ je þ jlat; ð2Þ

where the lattice conductivity jlat represents the heat por-

tion of the overall thermal conductivity j that is transported

by phonons and the electrical thermal conductivity je the

fraction associated to electronic carrier transport. Carrier

heat transport is related to electrical conductivity and

temperature by the Wiedemann–Franz law

je ¼ LrT ; ð3Þ

with L being the Lorenz number (L = 2.4 9 10-8 V2/K2

for free electrons) [1].

A thermoelectric device typically includes two types

of semiconducting materials, p- and n-type conductors

that form the n- and p-legs of a device. The equilibrium

carrier concentration in a semiconductor is dependent on

temperature. If a device with its p- and n-legs is placed

in a temperature gradient, the carrier concentrations

differ at the hot and cold side of the legs and lead to a

flow of n- and p-carriers, respectively; for the device

with its combination of n- and p-legs, the heat provided

to the hot side of the device then causes an overall flow

of electrical current through the device and delivers

electrical power. Typical thermoelectric modules are

composed of an assembly of alternating n- and p-legs,

thus requiring both n-type and p-type thermoelectric

materials.

To recover about 20 % of the heat energy in some of the

above mentioned applications, n- and p-materials with a

figure of merit of around 2 are needed [1].

1.2 General summary on thermoelectric material

performance

In the past years, extended research efforts have focused on

the development of thermoelectric materials. Presently

only tellurides have reached broad commercialization.

Bi2Te3-based modules are used as Peltier cooling elements,

in refrigerators or cooled car seats. Bi2Te3/PbTe2 thermo-

electric generators (TEG) are used in prototypes and small

serial productions in cars to recover exhaust gas heat from

the automotive tailpipe. CoSb3-based skutterudites are in

the development for the same application, with the goal to

also handle the higher exhaust temperatures that have to be

filtered out by an over-temperature valve for the telluride

modules. Si–Ge alloys found application in outer space

probes, where nuclear energy is used as heat source, so that

434 M. Backhaus-Ricoult et al.

123

electric energy can be produced over extremely long

periods of time (contacts are spring-loaded and modules

operate in constant temperature gradients).

Several low-temperature thermoelectric materials have

been developed. (Pb, Bi,…)(Te,Se,S) tellurides have been

optimized to ZT = 1–1.2 over a broad temperature range

for both, n- and p-type materials and, with silver or thal-

lium substitution, reach even ZT [1.5. Due to chemical

stability restrictions and low-melting point, applications of

Bi2Te3-based materials are limited to low temperature

(\450 �C) and, in addition, require protective surface

coating. Lead telluride has higher-temperature stability, but

is considered as environmentally unfriendly (regulated in

several countries). Limited chemical stability also restricts

applications of clathrates, skutterudites and silicides. Sur-

face protection against oxidation extends application tem-

peratures of skutterudites and silicides to 600–650 �C.

Some high-performance thermoelectric materials have

limited implementation due to environmentally unfriendly

component, such as lead, thallium,… or rare/expensive raw

materials, such as germanium, indium, ruthenium and

others.

Needs for low- and intermediate-temperature applica-

tions are in principle covered by tellurides, skutterudites,

silicides and silicon-based alloys [2, 3], even though fur-

ther material improvement is required to reach high con-

verter efficiencies. However, efficient, low cost,

environmental-friendly thermoelectric materials that can

operate at high temperatures (800–1,000 K) in air are still

lacking. Oxide materials have been considered as promis-

ing candidates based on their larger chemical and tem-

perature stability, their low cost and lack of toxic elements

[4, 5]. The best thermoelectric oxides include Na-interca-

lated Co oxide [6], layered misfit cobaltites, perovskites

[7], ZnO [8], In2O3(ZnO)n [9]. According to the literature,

single crystals of p-type sodium cobaltite excelled in ZT

and reached for optimized composition and orientation ZT

on the order of 1 [6], while ZT remains lower for known

n-materials. Significantly lower ZT values are reported for

polycrystalline materials. Ca3Co4O9 [10], doped SrTiO3

[11] and doped ZnO [8] show promising properties, but

stable polycrystalline materials with reproducible high

performance suitable for large-scale production have not

yet been demonstrated. Compared to other higher-perfor-

mance thermoelectric materials, most semiconducting

oxides suffer low carrier mobility at high carrier concen-

tration. They typically exhibit rather small power factors,

but offer relatively low thermal conductivity. Although

general concepts such as donor doping for enhancing the

carrier concentration are applicable to oxides, independent

tuning of the electrical properties seems to be more difficult

for oxides than for other small bandgap semiconductors. A

similar statement can be made for enhanced phonon

scattering through nanostructure engineering. Despite the

numerous efforts of introducing secondary-phase nano-

particles into oxides, no (reproducible) breakthrough has

been achieved.

2 Scope of this review

Due to their high-temperature stability, lack of toxicity and

low cost, semiconducting oxides are considered as poten-

tial candidates for thermoelectric power generation. Com-

pared to other high-performance thermoelectrics, they offer

low thermal conductivity, but reach only small power

factors as result of low carrier mobility and localized

charge carriers. The aim of this review was to assess the

potential of large bandgap semiconducting oxides by

evaluating existing approaches for tuning their electronic

and thermal properties. A key challenge for all thermo-

electric materials is the strong coupling of the thermo-

electric properties that has to be overcome to reach high

performance. The question is raised if it is more difficult to

overcome this coupling for oxides with their high carrier

density and low carrier mobility.

Doping is a standard procedure for semiconductors to

increase their carrier concentration and can also be applied

to oxides. However, the high carrier densities in semicon-

ducting oxides make effective doping more difficult; we

wonder if effective tuning and band engineering through

doping is possible for oxides and to what extent.

A further challenge in oxide materials becomes visible

when the widely used approach of nanostructured, engi-

neered microstructures is considered for enhancing phonon

scattering and decreasing the thermal conductivity. Struc-

turation at the nanoscale is efficient for materials with

mean free phonon path in the range of 10 nm and more and

can strongly reduce their thermal conductivity. It is more

difficult to match the mean free phonon path in oxide

materials (range of 1–2 nm or less) by microstructural

features of suited size for efficient scattering. We will

analyze and compare in this review the efficiency of pho-

non scattering at point defects, hetero- and homovalent

dopants (alloy scattering), grain boundaries, nanoprecipi-

tates and other nanoscale structures and especially analyze

the efficiency of crystallographic defects, such as stacking

faults, dislocations and point defects and superlattice

structures compared to nanograin size and nanosize sec-

ond-phase precipitates.

For our systematic experimental assessment, we select a

number of representative multicomponent oxides, that

range from materials with simple crystal structure (perov-

skite) to oxides that form planar crystallographic shear

defects (single sets of crystallographic shear planes in

Magneli phases [12] or intersecting shear planes in block

Semiconducting large bandgap oxides 435

123

structures or more randomized channel structures), adopt

layered superstructures and/or actively partner with a wide

range of secondary phases in composites. Many of the

selected oxides were chosen because their crystallographic

or microstructure features are within the size range of

0.3–20 nm and thus suited to interact with oxide phonons

in materials. In addition, for many of these oxides, doping

can be used to introduce intergap states so that some tuning

of the materials to higher electric performance seems fea-

sible. We explore the potential of these oxide materials by

studying the effects of doping, grain size, defect engi-

neering, addition of second phases, considering some

aspects of the impact of processing, providing a general

analysis of the figure of merit of these materials and dis-

cussing the expected device performance. We provide a

brief comparison of oxides to other thermoelectrics and

extract the key differences.

3 Concepts for managing thermoelectric properties

and their applicability for oxides

Enhancing the figure of merit of thermoelectric materials is

a key goal in the development of efficient thermoelectric

generators that led to the search for materials with phonon

glass, electron-crystal behavior. Several general approa-

ches have been described and implemented to reach this

aim. They can be roughly divided into measures with focus

on improving the power factor, while preserving (or even

decreasing) the thermal conductivity and approaches for

decreasing the thermal conductivity, while preserving or

even increasing the power factor. Even though the ther-

moelectric properties are highly coupled and no full sepa-

ration of both approaches is possible, this rough

classification helps to organize and review the approaches

implemented so far.

3.1 Concepts for improving the power factor

Doping has turned out to be a very efficient path to higher

power factors in many different semiconducting materials

[1]. Despite a strong coupling of properties, an overall

benefit in the thermoelectric power factor is often achieved

for semiconductors from doping, which leads to an increase

in ZT, since the charge carrier densities and their contri-

bution to the thermal conductivity remain low. The same

concept of increasing the carrier concentration through

doping can be applied to oxides. However, unlike semi-

conductors with their limitation by the carrier concentra-

tion, oxides have relatively high carrier concentrations and

are limited by their low carrier mobility. The question then

has to be asked whether doping is an efficient approach for

oxides. The benefit of an increase in power factor can

easily be annihilated by a substantial increase in thermal

carrier conductivity. High-conductivity oxides would suf-

fer most by increase in thermal conductivity.

Band engineering can be used to reach large power

factors. Based on Vining’s [13] expressions for conduc-

tivity and Seebeck coefficient from the density of states, a

special role can be attributed to states close to the Fermi

level, especially high densities of state in the bandgap just

below the Fermi level. Based on statistical thermodynamics

relationships between the Seebeck coefficient and the

derivative of the DOS at the Fermi level, an anomaly high

Seebeck coefficient is expected for high DOS and large

value for its slope close to the Fermi level. Tuning of the

bandstructure by doping is more efficient for semicon-

ductors than for oxides, since the densities of dopant-

induced intergap states remain very small in oxides.

In case of anisotropic distribution of carriers in a structure,

carrier pockets can be targeted and engineered for best ther-

moelectric property combination. The concept was described

for Si/SiGe and GaAs/AlAs superlattices, where the carrier

distribution was rendered anisotropic in the superlattice

structure [14], but also for anisotropic bulk materials.

Additional states can be introduced in the DOS of a

material through interaction with a dopant or second-phase

particles by a so-called resonance effect and yield to an

increase in Seebeck coefficient. Heremans [15] reported

doubling of ZT of PbTe by doping with thallium; thallium

provided additional energy levels (resonance levels) and

increased the density of states close to the Fermi level,

which then drove a higher Seebeck coefficient. Another

example is the interaction of ErAs nanoprecipitates with a

GaInAs matrix [16].

In nanostructured materials with the correct nanodi-

mensions, low-energy electrons are scattered at the inter-

face barriers, while higher-energy electrons pass unaffected

so that an energy filtering of the electrons takes place. For

the correct combination of nanodimension and electron

energy, the electron density distribution is narrowed by

such scattering in proximity of the Fermi level, so that a

beneficial effect on the Seebeck coefficient can be obtained

[14, 17].

Significant increase in ZT was achieved for low-

dimensional materials, when not only the thermal con-

ductivity was decreased, but also the Seebeck coefficient

increased by quantum confinement and strong localization

of states close to the Fermi level. Introduction of nanosize

alters the density of electronic states by quantum confine-

ment. The discontinuity of the electric properties can lead

to a decoupling of Seebeck coefficient, electrical conduc-

tivity and thermal conductivity and in special configura-

tions increase the figure of merit [14]. Examples are the

layered quantum dot structures of doped PbTe with (PbS-

eTe) dots that reached ZT = 3.5 (570 K) [18].


123

3.2 Concepts for reducing thermal conductivity

Several approaches focus directly on enhanced phonon

scattering to decrease the thermal conductivity.

Amorphous materials or glasses offer lowest thermal

conductivity; however, they do not meet the requirement of

high electrical conductivity for good thermoelectric per-

formance. Quasi-crystalline alloys and conducting glasses

so far have not reached promising performance.

Introduction of homovalent and heterovalent dopants in

the crystal lattice yields enhanced scattering of phonons at

the perturbed lattice sites. The volume fraction of the alloy

scattering centers is limited by the maximum solubility of

the various atoms/ions in the structure. Thus, its contribu-

tion remains often of limited importance for oxide mate-

rials with small solubilities and high carrier concentrations.

Heavy-ion species with large vibrational amplitude

(rattlers) at partially filled structural sites provide efficient

phonon scattering, as can be seen in skutterudites, clath-

rates and other cage structures [19]. In addition, a mixture

of different rattler atoms can be used to reach efficient

scattering in a broader range of the phonon spectrum

(demonstrated in multi-filled skutterudites [20]).

Nanostructured monolithic materials, composites and

superlattices offer a large number of grain boundaries and/

or interfaces that can be designed to reduce the thermal

conductivity more than the electrical conductivity [21].

The mean free path of electrons in solid matter is in general

much shorter than the mean free paths of phonons. In

addition, phonons show a very broad spectrum that for

materials such as silicon can range up to tens of micro-

meters, so that structural and mass perturbations within a

certain length scale range can be created in a crystal to

produce strong scattering of the phonons, but not of the

electrons. In some cases, only the low-energy electrons are

scattered at an interface barrier so that, in the best con-

figuration, energy filtering [14, 17] of the electrons takes

place and a narrower electron density distribution in

proximity of the Fermi level is obtained that then provides

in addition a beneficial effect on the Seebeck coefficient.

Structuration of a material at the nanoscale cannot only

include introduction of nanograins or fine second-phase

dispersions, but also nanopores. Extended efforts on con-

trolling nanovoid size and distribution together with

assessing related changes in thermoelectric performance

have been undertaken for ZnO [22].

Minimization of jlat through efficient nanostructuration

has been proposed as a path to high ZT. Phonon scattering

in nanostructured monolithic and composite materials has

been suggested and successfully implemented for example

for Si-based materials. Si–Ge alloys with grain sizes from

micrometer down to 10 nm have been systematically

studied [23]. Decreasing the grain size into the nanometer

range, adding second-phase particles and amorphous grain

boundary layers have been used to enhance phonon scat-

tering in Si–Ge and shown significant improvement in the

figure of merit due to enhanced phonon scattering at the

grain boundaries in the nanomaterials. The success for Si–

Ge alloys cannot easily be reproduced for other materials

because silicon has a very high lattice conductivity

([10 W/mK at 1,000 K), a huge phonon mean free path of

200–300 nm and an extremely wide distribution in its

phonon energies, ranging from free phonon paths of 3 nm

to a broad tail with more than 10 or even 50 lm. Phonons

of those energies can be efficiently scattered by nanograins

of 100, 200 nm and even larger. Broad grain size distri-

bution in the nanomaterials offers enhanced scattering for

broad phonon distributions.

For the evident reason that most materials have not

similarly large mean free path and broad distribution as

silicon, the nanostructuration approach for enhanced pho-

non scattering was applied with less success other materi-

als, such as tellurides, silicides and skutterudites.

3.3 Potential of renowned concepts of thermoelectric

performance tailoring for oxide materials

Compared to metallic semiconductors, oxides typically

have a larger bandgap, higher carrier density, more local-

ized electrons and much lower carrier mobility. While for

metallic semiconductors the power factor can typically be

improved by increasing the carrier concentration through

doping and this is sufficient to raise the figure of merit, the

high carrier concentrations in conductive oxides make this

approach more difficult and produce a stronger coupling of

electrical conductivity, Seebeck coefficient and thermal

conductivity. Therefore, traditional approaches are less

efficient to increase the figure of merit in large bandgap

semiconducting oxides.

Phonons in oxide materials have a typical mean free path

in the range of 0.2–2 nm, which is about a factor of hundred

smaller than in silicon. In silicon, the phonon distribution

shows a mean free path of 200–300 nm and an extremely

long tail that extends up to tens of micrometers. Therefore,

structural and mass perturbations at length scales from

10 nm to micrometers are able to produce strong phonon

scattering in silicon without having any major impact on the

electrons (at least for large wavelengths). Besides silicon,

very few materials have such large phonon mean free path

and broad distribution. Due to the small phonon mean free

path in oxides, structuration at the scale of 100 nm or even

10 nm, as typically found in nanoceramics or materials with

second-phase dispersions, is not efficient for phonon scat-

tering in oxides. Features of 0.5–2 nm size are needed to

introduce efficient phonon scattering. Features that have

multiple characteristic distances or cover a range are


123

preferred to a unique size feature, so that a broader frequency

range of phonons can be scattered.

Mass and size fluctuations are introduced in the oxide

crystal lattice by point defects, such as vacancies, homo-

valent and heterovalent dopants in both, cation or anion

sublattices; they offer a suited length scale for scattering

the oxide phonons. However, the volume fraction of such

defects is typically small (\1 % or even\10-3), so that the

number of scattering centers remains relatively small and

the overall benefit limited. Crystallographic defects, such

as dislocations, stacking faults or any other planar defects

also have suited sizes, but again their density has to be high

enough to introduce efficient scattering. Dislocation den-

sities in ceramics remain rather small, even in plastically

heavily deformed materials. Planar defects can occur at

high density in some oxide structures, they include twin-

ning, crystallographic shear, antiphase boundaries. It was

pointed out by Backhaus-Ricoult et al. [24] that crystallo-

graphic shear planes and twin boundaries contribute effi-

ciently to phonon scattering in titania and niobium oxide

ceramics and produce low lattice thermal conductivity.

Layered superstructures offer another opportunity for

enhancing phonon scattering in oxides. It is known that

misfit cobaltites Ca3Co4O9 with their alternating layers of

conducting CoO2- and insulating rock salt units exhibit

anisotropic properties with high electrical conduction in the

plane and low conductivity perpendicular to the plane.

Electrons propagating in the direction of the conducting

layers are less scattered than phonons propagating in this

direction, so that an overall in-plane advantage in the figure

of merit is expected in textured materials.

A comparison of high-temperature lattice thermal con-

ductivities of various oxide materials shows that typical

values range from 3 to 20 W/mK (e.g., 3.5 W/mK for

TiO2, 4–5 W/mK for SrTiO3, 7 W/mK for ZnO, and 20 W/

mK for alumina), while Nb oxide block structures, double

perovskites [25] and misfit cobaltites can show lattice

conductivity below 2 W/mK.

4 Crystallographic shear defects (CS) in large bandgap

oxides

Since crystallographic shear defects play an important role

in several materials of this review, their nature is described

in the following short excursion.

Crystallographic shear occurs in several transition metal

oxides WO3, MoO3, Nb2O5 and the rutile form of titania or

its combination with vanadium, chromium or other oxides

[26]. A crystallographic shear plane is a planar crystallo-

graphic defect that is associated with a change in the anion

to cation ratio without any change in coordination or sig-

nificant deformation of the anion polyhedra surrounding

the metal atoms. The metal coordination polyhedron is

usually an octahedron of oxygen atoms. Oxygen ions are

linked by corners or edges and corners to form a rather

open structure. Upon reduction, either by direct removal of

oxygen or by reaction with lower valence compounds, the

open structure collapses along specific crystallographic

planes, the crystallographic shear planes, to provide lower

energy structures, from which an entire plane of oxygen is

removed. The crystallographic shear vector is equal to an

oxygen–oxygen distance. With the frequency of this so-

called Wadsley defect (or CS defect) in the structure, the

non-stoichiometry can be varied. A homologous series of

defined compounds with different shear plane distances can

be formed. Simple structures exhibit only one set of shear

planes, such as titanium oxide.

Other materials can undergo simultaneously shear on

different planes and form different types of CS defects,

which intersect and result in a block structure. The CS

planes intersect to divide the materials into columns of

corner-linked octahedra. Local rearrangement in block size

and changes in their packing can be introduced by com-

position changes (cation substitution). Defects consisting

of clusters or walls of inappropriate blocks are very com-

mon in these compounds; their occurrence is strongly

controlled by processing.

5 Literature review on large bandgap oxide materials

5.1 SrTiO3

Doped, oxygen-deficient SrTiO3 materials have received

particular attention because of their high Seebeck coeffi-

cients and electrical conductivity [7]. SrTiO3 crystallizes in

the cubic perovskite ABO3 structure with the smaller Ti4?

ions at the cube corners being surrounded by 6O2-,

forming the TiO6 octahedral building units, and the large

Sr2? ions occupying the cube center. Stoichiometric

SrTiO3 is an insulator with a bandgap of 3.2 eV. The

perovskite structure offers some flexibility in composition,

such as oxidation, ultimately leading to the formation of

A2B2O7 pyrochlore, reduction toward A2B2O5 Brown-

millerite, incorporation of AO excess to form Ruddlesden

Popper phases Srn?1TinO3n?1 or integration of BO2 excess

that would ultimately lead to AB2O5 pseudobrookite for-

mation. Intrinsic transport properties of SrTiO3 can be

described by considering Sr-vacancies, O-vacancies and

electronic carriers (Ti-vacancies only have to be considered

at very high temperatures) [27]. Further on, homovalent

and heterovalent substitutions on both, A- and B-sites, as

well as introduction of oxygen vacancies via reduction are

possible. These simple point defects were sufficient to

describe the behavior of SrTiO3 [28]. Substitution,


123

deficiency or excess of the ions in the structure were typ-

ically found to lead not only to chemical and carrier density

changes, but also to structural distortions that lead to sig-

nificant modifications of the band structure. Similar other

oxides, doped n-type SrTiO3 has carrier concentrations in

the range of 1021/mol and low Hall mobility (0.3–0.4 cm2/

Vs), but excels by its exceptionally large effective carrier

mass (m* = 5–10 m0 in Nb-doped SrTiO3, while

m* = 2 m0 in Nb-doped titania and m* = m0 in spinel)

[29]. The large Seebeck coefficient of doped (reduced)

SrTiO3 is a direct result of the large effective mass, the

highly symmetric cubic structure with its triple-degener-

ated t2g state and the high density of states close to the

Fermi level. It results in power factors that are larger than

those of other oxides. The high lattice conductivity of

undoped and doped strontium titanate (5–12 W/mK) is a

disadvantage that is difficult to overcome by nanostruc-

turing for promoted phonon scattering since the phonon

mean free path is in the order of 1–2 nm at 1,000 K. The

amorphous limit of the lattice conductivity is estimated

around 1.8 W/mK (1,000 K).

The layered Ruddlesden Popper phases of the homolo-

gous series (SrO)n?1TinO3n?1 [30] with their superstruc-

tures of n perovskite layers per inserted distorted rock salt-

type SrO layer (also described by removal of TiO2 layers

along (100) and shear of the remaining perovskite block by

1/2[111]) were considered as particularly attractive since

their internal planar defect interspacing promises enhanced

phonon scattering and resulting low lattice conductivity.

Effectively, various Ruddlesden Popper phases were

intensively investigated [31, 32].

5.2 Titanium oxide

We chose titanium oxide as a second model material

family in this review, since it includes not only materials

with a wide range of point defect concentrations (oxygen

vacancies, cation substitution, anion substitution) in

undoped and doped rutile or anatase, but also different

types and densities of planar crystallographic shear

defects and microtwins in its Magneli phases. As a con-

sequence, the titania material family allows us to study

and compare the impact of point defects, planar defects

and grain size on the thermoelectric properties. Scattering

of phonons occurs at the point defects in a crystal and

scales with the concentration of those defects, which can

be impurities, dopant or substitution atoms and randomly

or ordered intrinsic defects, such as oxygen vacancies or

interstitial titanium ions. The TinO2n-x Magneli phases

are particularly interesting for thermoelectric applications

since their crystallographic shear plane distances are in

the same range as the mean free phonon path in these

oxides.

Based on the literature, the Ti–O phase diagram [33, 34]

presents a number of different oxide phases, including

titanium dioxide which crystallizes in three different

structures, rutile, anatase and brookite, the homologous

series of Magneli phases TinO2n-1 with n [ 4, Ti3O4,

Ti2O3, TiO1-x, a number of low-temperature suboxides and

a metallic solution.

At low oxygen activity, titania is an n-type conductor

with oxygen vacancies VO:: and Ti3? interstitials as dom-

inant point defects; the defects are more or less randomly

distributed. Rutile is the thermodynamically stable high-

temperature structure; anatase forms at low temperature

and for small particle size. Rutile has an extended stoi-

chiometry range and is able to host a large number of more

or less randomly distributed oxygen vacancies together

with charge-compensating cations [35]. Point defect mod-

els for rutile have been derived by several authors as

function of oxygen chemical potential, temperature and

doping level [36, 37]. The models indicate that the elec-

trical conductivity of titanium dioxide increases with

decreasing oxygen partial pressure. While donor dopants

improve the conductivity in the donor-controlled interme-

diate oxygen partial pressure range, the highest electrical

conductivities are reached in the extremely reducing

intrinsic regime, where the conductivity is controlled by

Ti3? interstitials and electrons [38]. Random oxygen

vacancies can only accommodate non-stoichiometries up to

x = 10-4 in TiO2-x. Once the oxygen vacancy concen-

tration in the rutile structure exceeds this threshold, the

randomly distributed oxygen vacancies collapse into a

planar shear defect; edge- and face-shared octahedral

replace the corner-shared TiO6-octahedral, and the cation

to anion ratio changes, but the cation coordination and

TiO6 octahedra structure remain unchanged. Crystallo-

graphic shear defects form on (1–21), (12–3),… rutile

planes with shear in the [110] direction. Shear plane type

and density vary, giving rise to a large number of rutile

Magneli phases [39].

The oxygen-deficient Magneli phases TinO2n-1 that

form at very low oxygen activity show CS superstructures

that accommodate the oxygen deficiency through oxygen-

deficient Ti2O3-type crystal slabs that are periodically

inserted every 1, 2,…, n-2 TiO2 layers to form Ti4O7,

Ti5O9,… TinO2n-1 [12]. For the higher Magneli phases

with n = 18–20, the CS densities remain low; the CS

planes are no longer periodically arranged and are strongly

associated with intertwinning. Lower Magneli phases are

reported to exhibit a periodic insertion of their crystallo-

graphic shear planes. Thus, Ti4O7 is built of rutile-like

slabs of TiO6-octahedra that extend infinitely in the a- and

b-directions and that are 4 octahedra thick in the c direction

inbetween the (110)-type shear planes. Ti5O9 shows a

periodicity of five octahedra inbetween the shear planes. In


123

higher Magneli phases, the energetically most favorable

shear planes typically form first. However, Magneli phases

with n = 4, 5, 6 require very high shear plane densities,

which cannot be provided by the lowest energy planes, so

that energetically less-favorable shear planes with higher-

defect energy form. In the Ti–O system, simultaneous

formation of CS defects on different planes was not

observed.

Shear defects in Magneli phases are separated by 4, 5,

… x d121rutile planes, with interdefect spacings of 0.7–2 nm,

which matches extremely well the mean free path of pho-

nons in titania (about 1 nm), so that a strong phonon

scattering can be expected.

We used literature data [40, 41] and carried out our own

studies [24] for doped rutile, anatase and various undoped

and doped Magneli-phase materials. Improved ZT values

compared to our own results were reported in the literature

for anion-doped Magneli-phase ceramics with nitrogen

Tin(O,N)2n-2 [42] for n [ 15 that showed similar sub-

stantial reduction in thermal conductivity as our own

research on undoped Magneli ceramics, but reached a

slightly higher ZT = 0.25 (1,000 K).

In addition, we compared in our own work the impact of

micrometer and nanometer grain size relative to the impact

of point or planar defects on the thermoelectric properties.

5.3 Niobium oxide

Niobium oxide block structures were included in this

review in the goal of enhancing the CS defect density and

introducing multiple coexisting interdefect distances. Nio-

bium oxides with compositions between NbO2.5 and

NbO2.42 adopt block structures. Two sets of crystallo-

graphic shear planes intersect to divide the materials into

columns of corner-linked octahedral. In projection, the

columns look like rectangular blocks. The phase range of

these structures can be extended by partial substitution of

Nb2O5 by TiO2, which drives the stoichiometry toward

MO2 and introduces additional smaller size blocks, or

WO3, which drives the stoichiometry toward MO3 and

introduces larger size block. Compositional changes occur

at the block periphery. Changes in local composition can

be implemented by local rearrangements in block size and

packing. Defects consisting of clusters or walls of inap-

propriate blocks are very common in these compounds and

are often processing specific. Phase diagram and crystal-

lographic-phase information are provided in [5, 44–46].

Based on the combined information, the niobium–oxygen

phase diagram reveals several oxides, NbO, NbO2, NbO2.5,

the homologous series of structurally related niobium oxide

phases with general formula Nb3n?1O8n-2, n = 5, 6, 7, 8

(Nb16O38, Nb19O46, Nb22O54, Nb25O62), Nb12O29

(12Nb2O5–2O) and Nb94O232 (47Nb2O5–3O).

NbO has a defective NaCl-type structure with both,

ordered cation and anion vacancies, and extends only over

a narrow homogeneity range from NbO0.982 to NbO1.008.

NbO2 crystallizes at room temperature in a deformed

rutile structure with chains of edge-shared NbO-octahedral

along the c axis, while neighboring chains are connected by

shared oxygen atoms. The Nb4? ions are shifted and form

metal–metal pairs within the chains. The bonded Nb–Nb

atoms are separated by a distance of 0.28 nm, whereas the

Nb–Nb distance along the chain is 0.312 nm. Bonding

within chains is reported to be homopolar and expected to

trap all available electrons in these bonds. NbO2 transforms

at about 1,125 K into a basic rutile (The transformation

temperature decreases with decreasing pO2.). NbO2 is a

small polaron conductor. Simple point defect models with

interstitial niobium and niobium vacancies as predominant

defects allow us to describe the intrinsic semiconductor and

metallic conductor behavior.

Nb2O5 is an insulator with a bandgap of 3 eV between

the oxygen 2p-valence band and the Nb 5d� band. It either

adopts a large number of different structures or was

described in the form of slightly different symmetries

[47]. The Nb2O5 structures can be roughly described by

very similar low- and medium-temperature phases that

transform at 1,273 K into a high-temperature block

structure of NbO6 octahedral (3 9 4 and 3 9 5) that share

corners with octahedral in their own block and edges with

octahedral in other blocks. One of the 28 niobium atoms

in each unit cell occupies a tetrahedral site at some block

junction. The low-/medium-temperature unit cell has a

more asymmetric pattern with 42 oxygen atoms forming

large open circles and 8 Nb ions being located in distorted

octahedral, while another 8 Nb ions occupy pentagonal

bipyramids. The remaining 0.8 Nb ions per unit cell are

located in interstitial 9-coordinated sites in the unit cell.

Both stoichiometric phases contain already several types

of niobium ions in the structure; electron densities and

carrier localization are expected to differ for those dif-

ferent niobium sites.

In addition to the stoichiometric phases, numerous

Nb2O5-x phases were reported. Even though no full

agreement on different phases and homologous series was

achieved by the different authors (some due to use of dif-

ferent crystallographic reference systems), most findings

can be summarized as homologous series of structurally

related niobium oxide phases with general formula

Nb3n?1O8n-2, n = 5, 6, 7, 8 (Nb16O38, Nb19O46, Nb22O54,

Nb25O62) and additional oxides of formula Nb12O29

(12Nb2O5–2O) and Nb94O232 (47Nb2O5–3O) [5, 26]. At

least the two latter phases seem to be stable, while others

are described as transient phases that form upon oxidation

of NbO or NbO2. Some of the metastable phases can be

considered as mixtures of different compounds of the


123

homologous series or more complex mixtures of those with

stable compounds. Some structural details are reported

based on HRTEM [30].

Based on high conductivity, tunable structural defects,

wide redox range and extended possibilities for hetero- and

homovalent substitution, niobium oxide-based materials

promise potential as n-type materials for thermoelectric

power generation at high temperature. Since only limited

data on thermoelectric properties are available in the lit-

erature, we investigated the properties of the various nio-

bium oxide phases and the effect of doping, substitution

and second-phase addition [48].

5.4 Tungsten oxide

Oxygen-deficient tungsten oxides were included in our

review, since they form highly defective structures ranging

from defective perovskites over Magneli-type phases to

chain structures. Details of the W–O phase diagram are

reported in [49]. The fully oxidized WO3 phase can be

considered as a modified perovskite-type (ABO3) structure,

where the large A atom is missing and the B atom is W6?.

The W6? atoms occupy the center of the oxygen octahedra.

At room temperature, WO3 is monoclinic rather than cubic

due to tilting of the octahedra. With increasing tempera-

ture, WO3 transforms from monoclinic to orthorhombic

(*603 K) and from orthorhombic to tetragonal

(*1,013 K) with transitions to other tetragonal structures

at higher temperatures. With increasing oxygen deficiency,

two series of Magneli phases are formed, WnO3n-1 and

WnO3n-2. In the WnO3n-1 series, point defects condense to

form crystal shear (CS) planes on {102}, where the lowest

n value is *14. These phases are referred to as {102} CS

phases. In the WnO3n-2 series, the CS planes lie on {103},

n ranges from about 12–28, and they are referred to as

{103} CS phases. For the W oxide Magneli phases, it was

found that each individual Magneli phase WnO2n-1

exhibits only a single n-specific type of shear plane and that

Magneli phase and shear plane change to another type,

when the shear planes get too close and the shear plane

interaction energy becomes too high. At lower oxygen

levels, two additional stable WO3-x phases were found,

W12O34 (WO2.83) and W18O49 (WO2.72). These are not

Magneli phases, rather the octahedra align to form pen-

tagonal columns. They are referred to as PC phases. Below

W18O49, WO2 is the only stable oxide. WO2 has a structure

similar to VO2 and is related to the rutile structure.

5.5 Misfit cobaltites

Even though misfit cobaltites are p-type oxides and do not

fit in the category of the so far considered substoichio-

metric n-type oxide materials of this review, we have

included them to illustrate both, the impact of their

incommensurate layered crystal structure on the thermo-

electric properties and the effect of texturing in case of

anisotropy in the transport properties.

NaxCoO2 [6] was the first metallic oxide with large

Seebeck coefficient that revealed excellent thermoelectric

properties (extrapolated ZT(1,000 K) = 0.7), but unfortu-

nately had a low high-temperature stability (sodium loss at

high temperature and hygroscopic in air). NaxCoO2 is built

of CoO2 layers with sodium atoms inserted inbetween the

layers. Numerous higher misfit cobaltites have been

derived with larger inserted oxide building blocks that are

more stable, such as Ca3Co4O9, BiCa2Co5/3Ox and others.

Ca3Co4O9 [50–52] adopts a CdI2-type crystal structure,

where layers of CoO2 with edge-sharing CoO6 octahedra

alternate with a triple-rock salt layers (Ca,Ca,CoO3). The

two layers have common lattice parameters in the a-

direction of the plane, but different parameters in the b-

direction. They form an incommensurate structure. It is

commonly stated that the CoO2 layers provides directed

planar electrical conduction, whereas the rock salt layers

can be considered as a reservoir for charge injection into

the CoO2 layers. The length scale of the superstructure and

charge transfer can be modified with dopant and substi-

tuting atoms of different sizes.

6 Modeling approaches for predicting thermoelectric

performance of oxides

First principles methods are widely used in the screening

and discovery of new thermoelectric materials, giving a

detailed picture of the band structure and suggesting new

ideas for how to adjust it to optimize thermoelectric per-

formance [53]. A major problem in such applications is the

strongly correlated nature of the electrons in many oxide

materials [54]. In such systems, one-electron theories often

yield limited insight and the many-body aspects of the

electronic structure, ostensibly, have to be explicitly con-

sidered [55]. The electronic structure of strongly correlated

oxide systems, even very simple ones such as TiO2,

remains an active area of research with important impli-

cations in the theory and understanding of oxide thermo-

electric materials. A major question is the extent of

localization of the electrons. If the electrons are delocal-

ized, the transport properties can be estimated from the

band structure, with the band curvature giving the effective

mass. The Boltzmann transport equations can be conve-

niently solved with codes such as BoltzTrap [56] to make

estimates of thermoelectric properties providing informa-

tion such as the optimal doping level for high thermopower

[57]. On the other hand, if the electrons are localized, and

electron transport is dominated by polaron hopping, the


123

calculation of conductivities requires knowledge of the

hopping matrix elements and reorganization energies. In

such cases, the thermopower can be estimated with simple

statistical mechanical principles.

Computational methods have also been widely applied to

the calculation of thermal conductivities. In comparison with

other simple transport properties, accurate calculation of

thermal conductivities is notoriously difficult and remains an

active area of research, even for simple materials such as

GaN [58]. Because of the long integration times required to

converge the heat–current correlations, even using non-

equilibrium techniques, the methods have only been applied

to systems with empirical potentials and thus have limited

generality. There have not really been sufficient studies

carried out to know how to best construct potential functions

that give accurate conductivities; much of the research effort

still goes into simply computing converged conductivities

(i.e., precision) without worrying too much about accuracy.

The particularly recalcitrant combination of strongly corre-

lated electronic structure and the difficulty of computing

reliable, reproducible thermal conductivities make explora-

tion of thermoelectric materials through molecular modeling

approaches highly challenging.

The TiO2 system illustrates well the inherent difficulties

in searching for optimal dielectric materials using atomistic

modeling approaches. First, the electronic structure of

excess electrons in TiO2 has been a subject of considerable

controversy with contradictory theoretical and experimen-

tal evidence for both localized and delocalized electronic

states. For example, X-ray structural studies in the Ti4O7

(2TiO2–Ti2O3) system have shown evidence for a Verwey-

type transition from localized electronic states at low

temperature to delocalized states at room temperature [59].

This has been recently shown to be consistent with

LDA ? U electronic structure calculations [60]. Recent

calculations [61] have suggested the existence of two

coexisting populations of localized and delocalized elec-

tronic states. In this work, it was shown that, because of the

extremely low mobilities expected for the localized elec-

trons, transport properties should be controlled by the de-

localized electrons.

It is known that there is strong conduction anisotropy in

reduced TiO2 with excess electrons, with high conductivity

along the [001] direction connecting edge-sharing TiO2

octahedra [62]. However, GGA ? U calculations [24] find

negligible anisotropy in the conductivities in contradiction

to experiments. This is an important check on the method;

while absolute conductivities cannot be computed within

the Boltzmann equation without knowledge of the relaxa-

tion time, the extent of anisotropy can still be predicted, at

least under the assumption that the anisotropy in the

relaxation time itself is small. This suggests that the de-

localized picture may not be correct, as it is obvious that a

hopping mechanism will give the expected anisotropy due

to the close Ti–Ti distances along the [001] direction (see

calculations in [63]). On the other hand, Janotti et al. [64]

showed in their Fig. 1 that the effective mass in the [110]

(M) direction is higher than along [001] (Z), a finding that

appears to differ qualitatively from the more complete

analysis provided by BoltzTraP. Further, studies of the

thermopower over the series TiO(1.75–1.91) in Backhaus-

Ricoult et al. [24] give excellent agreement with the Cha-

ikin–Beni expressions for the thermopower expected from

a U0 � kT Hubbard model in the localized-carrier regime

[65]. Presented with such contradictory evidence for a

system-like TiO2-x, it is difficult to see how electronic

structure calculations, at the present time, can serve as a

reliable guide for screening thermoelectric materials in

oxide systems. Clearly the community can look forward to

exciting improvements in computational screening of

thermoelectric oxide materials as research on electronic

structure calculations in strongly correlated oxide systems

moves forward.

Interesting preliminary insights into thermal conductiv-

ity in the TiO2-x system have been provided by molecular

dynamics calculations of thermal conductivity using

empirical potentials [66]. The calculations showed the

expected decrease in conductivity with interplanar spacing

in the Magneli phases; however, they showed an unex-

pected anisotropy with retarded heat conduction along the

Magneli shear planes. The reported effect is small but in

the right sense for improvement of thermoelectric proper-

ties as it is the same direction as the fast direction for

electrical conductivity in Ti4O7 [67]. This finding illus-

trates the promise of atomistic simulation approaches in

designing materials for optimal thermal scattering.

7 Oxide-specific problems and difficulties

for measuring thermoelectric properties at high

temperature

High-temperature electrical conductivity and Seebeck

coefficient have been frequently measured simultaneously

by using a ZEM (Ulvac Technologies) or similar equip-

ment. The thermal conductivity was then measured sepa-

rately. While the Laserflash method was broadly applied

and provided sufficient quality in its heat diffusion mea-

surements, the assessment of heat capacity in the same

equipment was often of low quality and specialized DSC

equipment produced higher-quality heat capacity mea-

surements. The van der Pauw approach allows us to mea-

sure simultaneously electrical and thermal properties on the

same sample; unfortunately, this approach is limited to low

temperatures due to the difficulty of making reliable high-

temperature contacts [1].


123

We used in our work a ZEM3 with platinum electrodes,

carbon contact foil and helium gas with a residual oxygen

content of 5–10 ppm oxygen (0.1 bar total pressure) and

acquired data in the temperature range from 400 to 1,100 K

on plan-parallel polished bar-shaped samples of size

2–3 mm 9 2–3 mm 9 12–15 mm. We restricted our

study to temperatures where the samples were stable during

the measurement time in the ZEM at low oxygen level and

did not change their stoichiometry (reproducible runs).

Stoichiometry changes were observed for some oxide

materials at temperatures [1,000 K or when exposed to

higher (lower) oxygen content than the processing or post-

annealing condition. In our work, we measured thermal

conductivity on carbon-coated 10 mm 9 10 mm 9

2–3 mm samples between 400 and 1,050 K in Argon

(residual oxygen level about 10 ppm) by the laser flash

method in an ANTAR 3. The thermal conductivity j was

derived as product of thermal diffusivity d, specific heat cp

and bulk density q of the sample, j = d � cp � q. The

density of the samples was determined at room temperature

from sample weight and dimensions; thermal expansion

coefficients from the literature were used. Heat capacity

was measured either in a Netzsch DSC 400 or a Laserflash

(Anter), comparing changes to a known standard of alu-

mina (Netzsch) or Pyroceram� (Anter).

In order to extract levers for low thermal conductivity

and compare the measured properties of different materials,

it is useful to derive the lattice thermal conductivities of the

dense materials by extrapolating measured values to full

material density and subtracting the electron contribution

from the thermal conductivity k using the Wiedemann–

Franz law, which is typically well obeyed at high tem-

perature. For lack of better data, we used in our work the

Lorenz number for free electrons 2.45 9 10-8 WX/K2 to

derive our lattice conductivities as function of temperature.

There are some particular challenges for property mea-

surements of oxide materials.

7.1 Effect of oxygen partial pressure and material

relaxation during measurements

Even though it is often stated that thermoelectric oxide

materials offer the advantage of being stable in air, this

statement has to be considered very carefully. Oxides and

especially those which are known for their n- or p-ther-

moelectric performance undergo changes in their oxygen

stoichiometry in their thermodynamic stability range.

Under thermodynamic equilibrium conditions, an oxide

equilibrates with its surrounding oxygen-containing envi-

ronment under oxygen ex- or incorporation. At low tem-

peratures, the kinetics may be sluggish, and a sample may

not be able to respond within the measurement time to a

change in pO2. At high temperature, kinetics may be fast

enough so that a bulk sample may respond rapidly to the

changes in the surrounding environment, particularly to the

oxygen partial pressure, so that a complete bulk sample or

only a surface layer changes its composition. The change in

oxygen stoichiometry corresponds to ‘‘oxygen vacancy

doping’’ and is often used as a powerful driver during the

optimization of oxide thermoelectrics. Therefore, care has

to be taken when thermoelectric properties are measured

for oxide materials, and it has to be evaluated whether the

Fig. 1 Schematics of the organization of MO6 octahedral in a perfect

rutile structure (TiO2, NbO2, WO3), b Magneli structure with

crystallographic shear defects on one type of shear plane (homologous

series of TinO2n-1, n = 4–9, …), c block structure with crystallo-

graphic shear defects on intersecting planes. The yellow and blue

squares represent full MO6 octahedral with a central M-atom

symmetrically surrounded by six oxygen atoms. Bonding between

octahaedra occurs via corner-sharing (represented as corner-

connected squares), edge-sharing (represented in the sketch as the

partial overlap of blue and yellow squares which corresponds to the

projected view of edge-sharing octahedral). Additional isolated M

ions are in the block structure (black dots) filling s interblock gaps.

The sketch is obtained by shear on different crystallographic planes

and represents a typical niobium oxide block structure with 3 9 4,

3 9 5, 3 9 3 blocks of NbO6 octahedral


123

material is ‘‘modified’’ (reduced or oxidized) during the

measurement. It is recommended to run measurements

during heat up and cool down or over long time to identify

material modifications. Often different piece of equipment

are also operated under different atmosphere; for example,

the inert gas ‘‘nitrogen’’ at one site may contain a different

residual oxygen content than the inert nitrogen or argon gas

used at a different site (rough levels of oxygen in bottled

nitrogen or argon are in the ppm range, but depend on the

gas quality).

The example of heavily reduced Zn(Al,Ce)O is used

as illustration. If doped zinc oxide is fabricated at

1,200 �C under heavily reducing environment, and its

electrical properties are measured in argon (5 ppm O2)

environment in a ZEM, during heating, at temperature

around 800 K, the electrical conductivity typically

shows a discontinuity in its slope due to oxidation of the

sample. Conductivity data acquired during sample cool

down will follow the lower conductivity branch of the

oxidized material.

7.2 Effect of material anisotropy in measurements

Many thermoelectric materials have complex (non-cubic)

crystal symmetry and show substantial anisotropy in their

properties. The anisotropy of a material can be further

enhanced by processing; uniaxial pressing, forging, tem-

plating, extrusion, etc., and may introduce anisotropy in

the microstructure and in the pore structure of the mate-

rial so that the electrical and thermal properties of a

sample differ in different sample directions. To assess the

thermoelectric properties of such anisotropic samples or

materials, it is important that electric conductivity, See-

beck coefficient, and thermal conductivity are measured

in the same direction of the sample (and preferentially on

the same sample).

The literature on oxide thermoelectric materials is

unfortunately full of data sets that were not measured in

same direction and show an exaggerated figure of merit.

We will use two examples for illustration, a hot-pressed

porous material with flattened pores in the compression

axis and a layered crystal structure, misfit cobaltite

Ca3Co4O9 with high conductivity in the CoO2 plane and

low conductivity perpendicular to it.

For small hot-pressed disk samples, bar samples are

typically cut from the disk for assessing electrical proper-

ties by ZEM in the direction perpendicular to the pressing

direction, which contains the smaller pore projections and

consequently has higher conductivity. If the disk itself or

part of it is used to assess the thermal transport by the

Laserflash method, the thermal conductivity is measured in

the direction of the compression axis, which means through

sections with large pore projections and thus lower thermal

(and electrical) conductivity. If those two sets of measured

data are combined, the resulting figure of merit is overes-

timated and false:

ZTfalse ¼ T rradSradial=vaxial� �

[ T rradSradial=vradial� �

¼ ZT

By precaution for a new process or material, it is rec-

ommended that measurements are always done in the same

direction or for both directions to assess the anisotropy of

the samples.

The crystalline anisotropy in layered crystal structures,

as for example misfit cobaltites, is praised for its low

thermal conductivity in the direction perpendicular to the

layering and the high electrical conductivity in the direc-

tion of the layer. A clear benefit from the anisotropy in

thermal and electric conduction is not evident. In textured

cobaltite ceramic, the grains are no longer randomly ori-

ented, but preferentially aligned in one direction. The

texturing yields anisotropic properties and makes it nec-

essary to assess thermoelectric material performance from

property measurements in the same directions, see Figs. 24

and 25.

8 Extraction of drivers for oxide thermoelectric

properties from experimental data

The strong coupling of the thermoelectric properties

makes it of little use to evaluate the effect of material

changes on a single thermoelectric property. Instead, the

impact on the overall thermoelectric performance has to

be compared. In order to compare materials and extract

trends in their thermoelectric properties as function of

dopant nature and concentration, grade of reduction,

second-phase addition, etc., we plotted both the Seebeck

coefficient and the lattice thermal conductivity as a

function of the logarithm of the electrical conductivity. In

these plots, our measured properties were extrapolated to

fully dense ceramics. The first type of plot is commonly

known as Jonker plot. Jonker had shown that for non-

degenerate semiconductors with band conduction, the

electrical conductivity (r) and Seebeck coefficient (S) are

related according to the equation:

S ¼ � k

eðln r� ln r0Þ ¼ �2:3026

k

elog r� log r0ð Þ; ð4Þ

where ? is for n-type and - is for p-type semiconductors

[68]. Similar relationships have been formulated for small

polaron conductors [69, 70].

Jonker plot and lattice-versus-electrical conductivity

plot allow us to identify drivers (inhibitors) for thermo-

electric performance. Drivers are found in the high Seebeck

coefficient–high conductivity corner of the Jonker plot with

largest possible deviation from the line representing the


123

ideal semiconductor behavior and in the lattice-versus-

electrical conductivity plot in the corner of high electric

conductivity and small lattice conductivity. If the experi-

mental data of a material family can be described by a

common relationship between S and r, then PF and ZT of

that material family or group of compositions can be cal-

culated from that relationship.

9 Material results

9.1 SrTiO3

9.1.1 Literature results on SrTiO3 thermoelectric

properties

In the field of oxide thermoelectrics, doped, oxygen-defi-

cient SrTiO3 ceramics have received particular attention

because of their high Seebeck coefficient and high electrical

conductivity, although their high thermal conductivity was

an obstacle for the realization of high ZT bulk materials.

Many studies explored the effects of reduction and doping

on the thermoelectric properties of SrTiO3 [4]. In the lit-

erature, results for oxygen vacancy doping [71], heterova-

lent doping [72] and homovalent A-site doping [29] were

reported. Some of the highest ZT values for n-type oxide

bulk materials were reported for Nb-doped strontium tita-

nate: ZT(1,000 K) = 0.35 for SrNb0.2Ti0.8O3 [73]. Efforts

on doping, grain size decrease and second-phase disper-

sions were implemented for enhanced phonon scattering

and led to a variety of materials with ZT(1,000 K) =

0.3–0.35. We demonstrated ZT (1,050 K) = 0.33 in poly-

crystalline bulk materials SrNb0.2Ti0.8O3-d and (La,Y)0.2-

Sr0.8TiO3-d with TiC addition [74]. We also investigated

the effects of combined heterovalent and homovalent A-site

doping (La0.15Y0.05CaxSr0.8-xTiO3-d), A ? B-site co-dop-

ing ((La,Y)xSr1-xNbyTi1-yO3-d), titania excess ((La0.15-

Y0.05Sr0.8)1-yTi1?yO3-d), and different second-phase

additions (TiB2, TiN, NbC, LaB6, LaAlO3, Y3Al5O12) on

the thermoelectric performance of strontium titanate [75,

76]. Although no bulk materials with ZT(1,000 K) values

higher than 0.33 were discovered, we found different paths

for reaching this value and gained insight into drivers and

limitations for further material development.

High thermal conductivity and limited dopant solubility

were key challenges for the successful development of

strontium titanate bulk materials with higher ZT. For thin

films and particularly constrained films on substrates,

higher ZT was reported (ZT = 0.37 for pulsed laser

deposited expitaxial films) [77]; particularly high ZT

(ZT = 2.4) was reported for the two-dimensional interfa-

cial layers between doped SrTiO3 film and LaAlO3 sub-

strate [78].

The layered crystal structures of the Ruddlesden Popper

phases offered another approach for reducing the thermal

conductivity. Effectively, a decrease in lattice conductivity

(about 40–50 % at room temperature and 80–90 % at

1,000 K compared to that of the symmetric perovskite

structure) was found, but it was unfortunately accompanied

by low electrical conductivity. The latter was caused by the

insulating nature of the SrO layers and a small effective

mass, only half of the value in the Nb-doped perovskite

structure. ZT typically ranged only around 0.15–0.2 at

1,000 K.

Several attempts were reported on modeling lattice

vibrations and their impact on thermal and carrier con-

duction [79].

9.1.2 Results from own work

In our own work, we have investigated the effects of

reduction, A-site, B-site and co-doping, as well as second-

phase addition on the high-temperature thermoelectric

properties of strontium titanate materials.

Processing We derived strontium titanium oxide materi-

als from mixtures of SrCO3, La2O3, Y2O3, TiO2, Ti2O3

and, in some cases, additional CaCO3 or Nb2O5 powders

that were turbula-mixed, cold-pressed into pellets, calcined

at 1,200 �C for 12 h in air, milled, then recalcined under

the same conditions and re-milled. For some materials, fine

second-phase particles of TiC, TiN, TiB2, LaB6, Al2O3 or

zirconia were added. Cold-pressed pellets were reduced in

a graphite bed at 1,400 �C for 6 h. The reduced materials

were milled and sieved to -325 mesh. Powders were

densified in a graphite die by current-assisted rapid sin-

tering (SPS) with hold at top temperature 1,500 �C for 15 s

under 35 MPa. Reduction of the materials was induced by

heat treating the powders in a graphite bed and, in some

cases, was enforced by adding reducing second phases.

9.1.2.1 Oxygen vacancy doping Reduction, also known

as oxygen vacancy doping, is a key driver for the ther-

moelectric properties of strontium titanate; high electrical

conductivity is only achieved in highly reduced materials.

Intrinsic (unreduced) SrTiO3 ceramics have very low

electrical conductivity and large Seebeck coefficient,

associated with a thermal conductivity in the range of

6–8 W/mK, so that their figure of merit remains below

ZT(1,000 K) = 0.01. Donor doping with La or Nb leads to

unreduced materials with Seebeck coefficients of -250 to

-500 lV/K and electrical conductivities that roughly scale

with the level of dopant and range from 2,000 to 8,000 S/m

at 1,000 K. Even highest concentrations of donor dopant,

for example 27 % of niobium, still do not provide prom-

ising thermoelectric properties; the figure of merit is only


123

ZT(1,000 K) = 0.05. The electrical conductivity is

expected to scale with the dopant level; 1 % of M3? donor

dopant on an A-site can potentially contribute 3.8 9 1.020/

cm3 carriers. However, the conductivities of unreduced

materials suggest that these carriers remain partially

localized and do not fully contribute to the conductivity.

Literature provides Hall mobilities for unreduced La- and

Nb-doped materials in the range of 1–4 cm2/Vs. Electronic

structure calculations showed a strong increase in the

effective mass m* with niobium doping compared to the

intrinsic material.

Reduction or oxygen vacancy doping is the most suc-

cessful driver to higher thermoelectric performance. It

leads to a large increase in the carrier mobility (Hall

mobility in the range of 10 cm2/Vs), increase in the

effective mass up to m* = 8 and to additional increase in

the carrier concentration. After reduction, niobium-doped

strontium titanate can reach an electrical conductivity of

15,000 S/m at 1,000 K and preserve a Seebeck coefficient

of -250 lV/K, so that, at similar thermal conductivity, the

figure of merit is dramatically increased to ZT = 0.25

(1,000 K). At the same level of reduction, donor dopants

play an important role in reduced material and drive the

electrical conductivity to higher values, while the Seebeck

coefficient is hardly altered, so that the figure of merit is

continuously improved with increasing donor dopant level.

While the above described trends were reflected in the

literature data, they were not clearly attributed. Therefore,

we systematically studied the effects of doping and

reduction in our own work and confirmed the behavior, see

Fig. 2.

9.1.2.2 Co-doping From the initial discovery of high

thermoelectric performance in reduced, n-doped SrTiO3,

further material optimization was included a systematic

exploitation of heterovalent and homovalent doping and

co-doping on A- and B-sites of the perovskite. Heterova-

lent A-site doping had been reported as a successful path to

higher ZT in the literature. We studied the effect of het-

erovalent A-site co-doping with La and Y in (La,Y)xSr1-

xTiO3-d materials with x = 0.05–0.25. Depending on raw

materials and processing conditions, our materials had

grain sizes in the range of 1–10 lm. With increasing sub-

stitution of strontium by trivalent yttrium and/or lantha-

num, an improvement of the thermoelectric properties was

observed. An optimum was achieved at 20 % strontium

substitution, before the ZT started to decline at higher

doping levels (either due to material inhomogeneity and/or

limit of dopant solubility). Figure 3 summarizes the evo-

lution of the figure of merit of reduced SrTiO3 (using our

reduction process) as function of dopant content. The

summary plot includes results from niobium B-site doping

and lanthanum and/or yttrium A-site doping and co-doping

in single-phase ceramics and composites with 5 wt% TiC

nanoparticles. Independent of the nature of dopants, the

figure of merit of the materials increased with the total

dopant content from 0.03 to a maximum ZT of 0.25–0.33 at

dopant levels of about 0.2 before decreasing. The data

revealed a solubility limit for A- and B-site dopants in a

range between 0.27 and 0.37 at 1,050 K, in broad agree-

ment with rigid-band Boltzmann transport calculations

[80]. The nature of the dopant and substitution site show a

slight impact on the ZT reached at a fixed dopant con-

centration. Niobium is the most efficient dopant, followed

by lanthanum, followed by yttrium. Mixed (La, Y) A-site

doping combined with addition of reducing second phase

(5 %TiC) produced a similar ZT evolution as simple nio-

bium B-site doping. Based on the experimental data, it is

not clear if the decrease in ZT at higher dopant levels was

caused by the limit in dopant solubility or reflected the

Fig. 2 Schematic presentation of the impact of donor doping and reduction on electrical conductivity and Seebeck coefficient of strontium

titanate


123

decline in power factor that was predicted by modeling for

a carrier density larger than 1 9 10-21 cm-3 (larger than

the dopant-associated carrier concentration in our

experiments).

The ratio between the two donor dopants Y and La in

(La0.2-xYx,Sr0.8)TiO3 was varied over a range of x = 0–0.2

to study the impact of A-site co-dopant ratio on the proper-

ties. The overall figure of merit was found to increase for

addition of lanthanum and also for addition of yttrium; best

improvement was demonstrated for co-doping with

0.025–0.05 yttrium and 0.025–0.15 lanthanum. The best

values, ZT (1,050 K) = 0.07 and 0.15, were reached for

unreduced and reduced La0.15Y0.05Sr0.8TiO3, respectively

[74].

Composite materials were made with addition of TiC

nanoparticles that were added prior to material densifica-

tion. TEM analysis revealed the presence of grain boundary

and triple-point decoration by a carbon-rich phase, possibly

a solid solution of TiC. The electrical conductivity of the

composite increased by addition of TiC compared to

material that was simply reduced in a carbon bed, thereby

suggesting that the addition of TiC induced a further

reduction of strontium titanate. The second-phase addition

was varied over a range of 0–30 %; best ZT was achieved

for about 5 wt% of TiC nanopowder addition. At 1,050 K,

the electrical conductivity of the best composite material

with TiC was twice that of TiC-free reduced material.

Seebeck coefficients were negative for all materials and the

absolute values increased with increasing temperature.

Even though the Seebeck coefficient preserved some cou-

pling with the electrical conductivity, an overall benefit in

the power factor of reduced material was visible. The

reduced material achieved power factors in the range of

5 9 10-4 W/mK2 (1,050 K). Addition of nanoTiC

produced an additional large increase in electrical con-

ductivity by a factor of 2 at 1,050 K (factor 6 at 550 K!),

while the Seebeck coefficient did not change. As a result,

the power factor reached 8 9 10-4 W/mK2 at high tem-

perature (1,000 K). The thermal conductivity was observed

to increase with reduction, but the thermal conductivity of

the reduced material without and with TiC addition

remained very similar. As shown in Fig. 4, the lattice

thermal conductivity of TiC-containing reduced material

was actually lower than the lattice conductivity of both,

calcined and reduced materials without TiC addition. This

suggests that fine TiC particles in the final composite

contribute to the phonon scattering. The effect produced a

decrease in lattice conductivity from 3.0 to 2.5 W/mK over

a wide temperature range. The figure of merit of both

reduced materials was higher than that of unreduced

material and increased with temperature. A ZT of 0.3 at

1,050 K was obtained for reduced composite material with

5 % TiC addition. A comparison of the thermoelectric

properties for calcined and reduced La0.15,Y0.05Sr0.8TiO3-d

and its composite with 5 wt% TiC as a function of tem-

perature is presented in Fig. 4. Further details were

reported in [75].

With the aim of further reducing the lattice thermal

conductivity, we combined homovalent and heterovalent

A-site doping and introduced Ca in addition to Y, La on the

A-site in monolithic materials (La0.15Y0.05)CaxSr0.8-x-

TiO3-d

, x = 0.1, 0.2, 0.3 and their composites with 5 wt%

TiC. In the monolithic materials (without TiC), Ca-doping

produced a large decrease in electrical conductivity with

little change in Seebeck coefficient and therefore very low

power factors. In materials with TiC addition, the power

factor increased with Ca-doping, although the highest

achieved power factors were similar to the Ca-free material

Fig. 3 Impact of dopant nature

and concentration in reduced

monolithic SrTiO3 and its

composites with TiC on ZT at

1,050 K; an approximate

solubility limit is indicated by a

dotted vertical line


123

with TiC. At low temperatures (\750 K), the thermal

conductivity effectively decreased with increasing Ca due

to lower lattice conductivity. At higher temperatures

([950 K), lattice conductivity and total thermal conduc-

tivity converged for all Ca-levels. The lowest lattice con-

ductivity was 2.5 W/mK at 1,050 K. The highest ZT at

1,050 K (0.27) was reached for a composite with 0.3Ca and

5 % TiC. Further details were reported in [81].

To fully explore all cation doping options, we partially

substituted titanium by the heavy donor atom niobium on

the B-site, (La,Y)xSr1-xNbyTi1-yO3-d, in the hope to fur-

ther increase the power factor and decrease the lattice

conductivity.

For a total dopant level (x ? y) = 0.2, no improve-

ments in properties were found for different A-/B-site

doping ratios (x:y) over material without any B-site

doping. Materials were also prepared with different total

dopant concentrations (x ? y) = 0.1, 0.175, 0.2, 0.25 at

fixed Sr/Nb ratio x:y = 1 and fixed ratio [La]/[Y] = 1

with and without TiC additions. Figure 5 shows the

results. For materials without TiC second phase, the

electrical conductivity was low (\5,000 S/m), and power

factor and thermal conductivity decreased with increasing

total dopant level. ZT = 0.1 (1,050 K) was the best

value for reduced samples with x ? y = 0.125 without

any second-phase addition of TiC. For materials with

TiC addition, the electrical conductivity increased with

the total dopant level (x ? y), while the absolute value of

the Seebeck coefficient decreased. The thermal conduc-

tivity was similar for different total dopant levels

(x ? y). Thus, the ZT curves for different TiC levels

were similar with maximum values around ZT = 0.3 at

1,050 K. The figure of merit for A ? B-site doping was

comparable to La ? Y co-doping in (La0.15Y0.05)0.2-

Sr0.8TiO3-d [81].

In order to better understand the impact of A/B stoi-

chiometry, (La,Y)-co-doped strontium titanates were pro-

duced with titanium-excess y = 0.04–0.11 in composition

(La0.15Y0.05Sr0.8)1-yTi1?yO3-d. The excess titanium oxide

was added as Ti2O3 to the starting powder; calcining,

reduction and sintering were not changed compared to the

previously described process. No TiC was added. Best

results were obtained for y = 0.078–0.09. The electrical

conductivity of the material with excess Ti was signifi-

cantly increased over that of the stoichiometric composi-

tion (y = 0), but remained lower than conductivities of

reduced stoichiometric composites with TiC. The Seebeck

coefficient of the stoichiometric composite material and the

excess Ti material were the same. The lattice thermal

conductivity, however, was higher in the material with Ti-

excess. ZT(1,050 K) = 0.25 was achieved with Ti-excess

and reduction.

Fig. 4 Comparison of thermoelectric properties as a function of temperature for (La,Y)0.2Sr0.8TiO3-d materials as-made, reduced and reduced

with 5 wt% TiC addition


123

9.1.2.3 SrTiO3 composites Since addition of TiC as sec-

ond phase had a strong impact on the thermoelectric prop-

erties of reduced co-doped strontium titanates, the effect of

other second-phase additions was investigated. TiC, TiB2,

TiN, NbC, LaB6 LaAlO3, Y3Al5O12 were added to a base

composition of La0.15Y0.05Sr0.8TiO3-d. TiC, TiN, TiB2 and

LaB6 efficiently increased the electrical conductivity and

the power factor and produced a decrease in lattice thermal

conductivity, see Fig. 6. TiC, TiB2 and TiN were most

effective and gave ZT = 0.3–0.33 at 1,050 K [81].

Further on, the addition of insulating second phases,

such as alumina, zirconia, was explored at levels between 2

and 4 wt%. No improvement in the overall ZT was

achieved compared to the monolithic ceramics

(ZT(1,050 K) = 0.34). The result was not affected by the

size of the second-phase particles, supporting once more

that second-phase dispersions need size below 10 nm for

efficient interaction with the phonons.

9.1.2.4 Comparison of strontium titanate-based materi-

als The Jonker plot in Fig. 7 includes doped SrTiO3

materials and composites. The dotted line in Fig. 7

describes the behavior of an ideal single-parabolic band

semiconductor with slope k/e; the line was fitted to

materials with low conductivity. Figure 7 demonstrates

that the family of SrTiO3 materials cannot be described by

a single line in the Jonker plot. Many materials with low

electrical conductivity deviate from the line, showing

lower absolute Seebeck coefficient and thus smaller power

factors. High-conductivity materials align on a common

line with a slope in the order of 0.5 k/e; they distinguish

from the low conductivity materials by the benefit of higher

Seebeck coefficient compared to ideal semiconductor

behavior and represent a group of materials with improved

power factor.

Figure 7b shows the thermal conductivity as function

of the logarithm of the electrical conductivity. Reduced

SrTiO3 with homo- and heterovalent doping and/or sec-

ond-phase addition shows lattice conductivities in the

range of 2.5–3.5 W/mK. The plot suggests an overall

slight decrease in lattice conductivity with increasing

electrical conductivity. This suggests that the penalty of a

higher carrier contribution to the thermal conductivity for

high electrical conductivity materials is partially or fully

compensated by lower lattice conductivity. Particularly

low lattice conductivity at high electrical conductivity

was found for reduced doped materials made with second-

phase addition of TiC, TiN, LaB6, TiB2.

Fig. 5 Thermoelectric properties of A- and B-site co-doped SrTiO3, (La,Y)xSr1-xNbyTi1-yO3-d with La and Y on the A-site and Nb on the

B-site for reduced monolithic material and composite obtained with 5 % TiC addition


123

9.2 Titanium oxide-based materials

9.2.1 Thermoelectric performance from the literature

View the interest over the past decade in titania for solar

applications, photocatalysis, photo-electrochemical water

splitting, sensing applications and others, many of the

physical properties of rutile, anatase and brookite have

been reported in the literature, including electrical con-

ductivity and to a lesser extent thermal conductivity and

thermopower [36, 82–85]. Undoped titanium dioxide

(rutile) is an n-type semiconductor with a bandgap of about

3.3 eV. The n-type properties are promoted by donor-type

intrinsic defects including oxygen vacancies and interstitial

titanium cations. Defects that support p-type conduction,

titanium vacancies, form only at high oxygen activity and,

Fig. 6 Impact of various second-phase additions on the thermoelectric properties of reduced La0.15Y0.05Sr0.8TiO32d

Fig. 7 Presentation of a Seebeck coefficient and b lattice thermal conductivity as function of the logarithm of the electrical conductivity for

SrTiO3 materials with various dopants, reduction levels and second-phase additions


123

in addition, are rather immobile and require very high

temperatures for equilibration. Therefore, the intrinsic

defect disorder in titania can be described by defect for-

mation and equilibration equations for oxygen vacancies,

titanium interstitials Ti3?, Ti4?, electrons e0, electron holes

h, the charge neutrality condition and the site conservation

in the crystal.

Since the electrical conductivity is the weighted product

of the concentration of the charge carriers and their charge,

and the ionic conduction is small compared to electronic

conduction, the overall conductivity can be approximated

by r = Ne e be ? Nh 9 h 9 bh (N—carrier concentration,

b—mobility).

Based on the defect chemistry, the intrinsic behavior of

titania is controlled by oxygen vacancies in an interme-

diate to low oxygen activity regime, while at very low

oxygen activity, Ti3? interstitials become the dominant

defects with concentrations that increase with decreasing

oxygen activity. As a consequence, the electrical con-

ductivity is expected to follow V-shape as function of

oxygen activity.

Experimental data for electrical conductivity, Seebeck

coefficient and Hall mobility for titanium dioxide are

reported in the literature as function of temperature and to a

limited extent also as function of oxygen partial pressure.

Data for pure titanium dioxide are summarized in [86–90].

The electrical conductivity follows the expected V-shaped

curve as function of oxygen activity, and the V-curves shift

to lower values with decreasing temperature. Only few of

the reported data were measured under controlled oxygen

activity and temperature and for clean samples. High-

quality data include those by Tani and Baumard [37] and

the multiple studies by Novotny. Novotny measured both,

electrical conductivity and Seebeck coefficient, for single

crystals and for polycrystalline materials with different

dopants. The anisotropy in conductivity was scarcely

explored [62, 91]; results suggest an eight times higher

conductivity in the \001[ direction than in the \110[direction.

Reduced titania Magneli phases TinO2n-1 have been

object of numerous studies since their discovery by Mag-

neli himself [40]. Early work on Magneli phases

TiO1.94–1.86 (fabricated from anatase powder by reduction

in H2 at high temperature and subsequent cold pressing)

advanced the hypothesis of ZT(350 K) = 0.2 and, based

on experimental Seebeck coefficient of -0.5 mV/K at

70 �C and extremely optimistic assumptions (not mea-

sured) on the thermal conductivity, possibly much better

high-temperature values [92]. Those assumptions were

proven wrong by later work. In [24, 93], various reduced

rutile and Magneli phases were processed and their elec-

trical and thermal properties measured; ZT(1,000 K) was

not found to exceed 0.2.

Heavily reduced titania TiO1.1–1.2 (processed by con-

trolled combustion synthesis from metallic titanium with

sodium perchlorate, etc., and densified by spark plasma

sintering) was reported to reach exceptional figures of

merit of 1.2–1.64 at temperature of 700–1,200 �C [94], but

these results were not confirmed by the scientific commu-

nity, including our own work.

Thermal properties can be derived from electronic

structure calculations, but require intensive computation

and typically only provide properties of perfect structures,

while thermal properties of real material are significantly

affected by their point defects, impurities, planar defects

and grain boundaries (in polycrystals), which produce

phonon scattering and lead to a decrease in thermal lattice

conductivity. Few computational results are available on

thermal properties. In [95], strong anisotropy in thermal

conduction was shown in the rutile single-crystal phonon

DOS.

Experimental high-temperature measurements are

scarce, especially under defined thermodynamic conditions

or as function of crystal orientations. There is a general

consent on significant anisotropy in the thermal conduc-

tivity of rutile with values perpendicular to the c direction

being larger than in the c direction with 13 W/mK parallel

and 9 W/mK perpendicular to the c axis at 273 K [95].

Experimental results also reflected an impact of the rutile

non-stoichiometry on the thermal conductivity.

Experimental observations of lowered lattice conduc-

tivity in Magneli phases compared to nanoceramics of

rutile or anatase confirmed this expectation. Ceramics with

the highest shear defect density, Ti4O7, showed a 25 %

decrease in lattice conductivity compared to rutile ceram-

ics. While similar findings were reported in the literature

[93], no comparison was provided for the effectiveness of

phonon scattering by point defects, grain boundaries or

nanodispersions in such titania ceramics.

9.2.2 Thermoelectric performance results from own work

With the aim of comparing the impact of point defects

from hetero- and homovalent cation doping or oxygen

vacancy ‘‘doping’’ via reduction, planar crystallographic

defects and grain boundaries on the thermoelectric prop-

erties, we conducted own research, fabricated a wide range

of undoped and doped titanium oxides and their composites

with micro- and nanograin size and determined their

properties, covering a wide range of reduction from stoi-

chiometric titanium dioxide to TiO.

Processing Our own titanium oxide-based materials were

prepared via partial reduction at increased temperature

from unreacted or reacted powders or densified ceramics of

undoped or doped rutile or anatase and their composites.


123

Some materials were made by reduction in gas mixtures

(O2/N2, CO/CO2, H2/H2O), in graphite containment or by

mixing them with TiO, NbO, niobium metal, graphite, TiC,

TiN, NbC, WC, W, Mo. Materials were fabricated by

natural sintering in air or controlled environment and by

current-assisted rapid densification in uncoated and/or

alumina-coated graphite dies. Details on processing were

reported in [24]. Our assessment also included rutile single

crystals with different orientations and various levels of

reduction.

Our titania ceramics covered a wide range of grain sizes

(single crystals to 30 nm grain size nanoceramics) and

electrical conductivity (insulator to metallic behavior).

Most materials had densities of 95 % and more. In the

presented own data, we extrapolated all material properties

to 100 % density to enable best comparison.

9.2.2.1 Electrical properties of undoped titanium oxides

Rutile single crystals (MTI) in their as-received state

exhibited the typical high Seebeck coefficient and low

electrical conductivity of an insulator, 700 lV/K and 100

or 200 S/m in air at 1,000 K, respectively, perpendicular to

(001) and (100). For polycrystals, an average of all single-

crystal directions would be expected (in case of ideal

behavior). However, our measured electrical properties of

polycrystalline (micrometer grain size) and nanocrystalline

rutile and anatase showed lower power factors with con-

ductivities around 100 S/m (lowest values for smallest

grain size) and 500 lV/K as Seebeck coefficient in air at

1,000 K. The results were attributed to the imperfect nature

of the poly- and nanocrystalline materials, small levels of

acceptor impurities (Fe, alkaline earth and alkali) in the

bulk and segregation at grain boundaries. At similar grain

size and same oxygen partial pressure, anatase slightly

exceeded in electrical conductivity, but had a lower See-

beck coefficient (-400 to -500 lV/K) than nanorutile

(-600 lV/K). In response to equilibration in reducing

environment, the absolute value of the Seebeck coefficient

decreased and the electrical conductivity increased; (100)

single crystals, poly- and nanocrystalline materials were

found to follow a common evolution upon reduction,

reaching Seebeck coefficients of about -200 lV/K at

about 10,000 S/m. The conductivity increased with

increasing temperature. The evolution of the electrical

properties with reduction was visualized in a Jonker plot,

Fig. 9, where the Seebeck coefficient is plotted as a func-

tion of the logarithm of the electrical conductivity. For

ideal single-parabolic band semiconductors, a straight line

with slope k/e is expected, since the electrical conductivity

is proportional to the carrier density r = N e b (N carrier

density, e carrier charge and b carrier mobility), while the

Seebeck coefficient exhibits inverse proportionality to the

carrier density S = k/e [ln(2pm* kT/h2N)] ? const = k/

e lnr ? k/e ln(2pm*kT/h2)0b0 const. Surprisingly, a com-

mon straight line with slope k/e was found for all rutile

ceramics and single crystals at different states of reduction;

anatase ceramics closely followed the same line.

The Seebeck coefficients of the Magneli phases spanned a

wide range. Magneli phases with low non-stoichiometry

exhibited semiconductor behavior with increasing electrical

conductivity for increasing temperature, while Magneli

phases Ti4O7 and Ti5O9 reflected metal behavior with

decreasing conductivity for increasing temperature, Fig. 8.

At 1,000 K, Ti4O7 reached electrical conductivities in the

range of 60,000 S/m at a Seebeck coefficient of -100 lV/K.

The electrical conductivity was found to increase with

increasing oxygen non-stoichiometry, and power factors

were dramatically improved with reduction (Figs. 8 and 9).

Even though a strong coupling between Seebeck coef-

ficient and electrical conductivity was expected, it was

surprising that the Magneli phases followed the unique line

of reduced rutile in the Jonker plot and that the line

extended over three orders of magnitude in conductivity

from the metallic Magneli phases to the slightly reduced

rutile single crystal, Fig. 9. Properties of Ti2O3 and TO1?x

with their vanishing Seebeck coefficients slightly deviated

from the common behavior of the other titanium oxides.

The observation that semiconducting and metallic Magneli

phases follow the same dependency as insulating rutile

reflects a major lack of differentiation in the band structure

features of their contributing carriers. We had expected that

structural CS defects would introduce special localized

states with differing degeneracy or additional delocalized

states, but apparently the TiO6 octahedral environment

remains the determining feature in all structures. While this

may explain the behavior of the semiconducting higher

Magneli phases, an explanation of highly localized elec-

trons in the metallic Magneli phases Ti4O7 and Ti5O9

cannot be used. It was even more surprising that the See-

beck coefficients of all Magneli phases fitted the Heike’s

formula for localized-carrier transport

S ¼ � kB

ej j ln2Sn þ 1

2Snþ1 þ 1

� �1� x

x

� ��

The predicted Seebeck coefficients perfectly matched

the experimental data of the Magneli phases at high tem-

perature. Based on our observations, we had to conclude

that models for band and polaron conduction predicted the

same Seebeck coefficient and matched the experimental

data for the titanium oxide Magneli phases.

9.2.2.2 Impact of grain boundaries and crystallographic

shear defects on thermal conductivity Rutile single

crystals revealed anisotropy in their lattice thermal con-

ductivity, as shown by the values of 4 and 6.5 W/mK at

550 K and 3.5 and 5 W/mK at 1,050 K that we measured


123

in the [100] and [001] directions, respectively. Polycrys-

talline rutile with micrometer grain size was found within

those bounds, close to the lower value, thus demonstrating

some contribution of phonon scattering at grain boundaries

in the polycrystalline ceramic. Compared to large grain

size ceramic, nanorutile and anatase ceramics showed

slightly lower lattice conductivities of 3.3 and 3.0 W/mK

(1,000 K), respectively. Anatase has a structural advantage

and therefore exhibited lower thermal conductivity than

rutile. Significantly lower lattice thermal conductivity was

observed for the Magneli phases. For these ceramics, no

advantage was achieved by decreasing the grain size. The

lattice conductivity was found to decrease with increasing

shear plane density (or non-stoichiometry), reaching a

minimum of about 2.5 W/mK at 1,000 K for Ti4O7,

Fig. 10. The lattice conductivity of a given Magneli-phase

ceramic was found to be the same for micro- and nano-

crystalline materials and thus independent of grain size.

The differences in thermal conductivity were more pro-

nounced for lower temperatures, since the carrier contri-

butions to the thermal transport remained low (scaled with

temperature).

The combination of electrical and thermal perfor-

mances provided a clear increase in ZT from single

crystalline to polycrystalline to nanocrystalline rutile or

anatase, followed by a large leap in performance to the

Magneli phases with their crystalline shear defects,

Fig. 8d. In our studies, Ti8O15–Ti10O19 showed the

largest ZT; they were closely followed by the other

semiconducting Magneli phases. Ti5O9 had clearly

lower ZT and Ti4O7 again substantially lower ZT, even

though the values remained larger than those of unre-

duced rutile. Even though Ti4O7 had the lowest thermal

lattice conductivity and the highest electrical conduc-

tivity, its ceramic did not provide the highest ZT. Its

Fig. 8 Electrical conductivity, Seebeck coefficient, thermal lattice conductivity and ZT as function of temperature for as-processed and reduced

single-crystalline rutile, micro- and nanorutile, anatase, various Magneli phases and Ti2O3

Fig. 9 Jonker plot for pure titanium oxide phases with different

degrees of reduction


123

high electrical conductivity yielded large contributions

of the carriers to the thermal conductivity.

Since each Magneli phase provides only one particular

CS defect spacing, the question was asked whether a

variety of CS spacings in a mixture of different Magneli

phases would enhance the phonon scattering and result in

even lower thermal conductivity and higher figure of merit.

To test this hypothesis, we fabricated various mixtures of

2–5 different Magneli phases. We found that the electrical

properties and lattice conductivity matched the phase

average without any particular benefit in the lattice con-

ductivity or the overall performance.

Based on the above results, we identified three dif-

ferent drivers for TE performance of pure titanium

oxide phases, oxygen vacancy doping (reduction) that

improves ZT mainly through improvement of the elec-

trical performance, phonon scattering at grain boundary

that provides a slight improvement in thermal lattice

conductivity and phonon scattering at crystallographic

shear defects as strong driver for decreasing the thermal

conductivity. Phonon scattering at the crystallographic

shear defects was more efficient than phonon scattering

at grain boundaries in nanomaterials. The key role of

CS defects was explained by the very high densities of

crystallographic shear planes and the particularly good

match of the shear defect spacing with the mean free

phonon path.

9.2.2.3 Cation doping and substitution In the literature,

the effect of dopants on the defect chemistry, the electrical

conductivity and the Seebeck coefficient of titania was

extensively studied for n-dopants, such as Ta, Nb. It was

shown that Nb-doping (0.65 %) increased the carrier con-

centration and improved the electronic conductivity by

several orders of magnitude. Nb-doping shifted the n–p

transition to higher oxygen activity. Metallic-like conduc-

tion was obtained at low oxygen activity and semicon-

ductor behavior prevailed in the high oxygen activity

range. The defect chemistry was analyzed in detail in [96].

We explored a wide range of donor dopants, including

Nb, Mo, Ta, V, W, homovalent substitution by heavy

cations Zr and Sn and low levels of acceptor doping with

Fe, Sr, Al and some combinations.

Niobium can be hosted in the rutile lattice up to

levels beyond 10 %. In our work, niobium-doped rutile

effectively showed a substantial increase in electrical

conductivity in air and reached much higher power

factors than undoped rutile ceramics, Fig. 11. As a result,

the figure of merit reached 0.05 at 1,000 K, and thus

doubled the ZT of undoped rutile. However, the power

factors remained below 1 9 10-4 W/mK2 (1,000 K) for

highest doping levels and did not come close to the

values of the Magneli phases that were about an order of

magnitude higher.

Reduction (oxygen vacancy doping) of niobium-doped

rutile produced a significant increase in electrical conduc-

tivity (20,000–25,000 S/m at 1,000 K) and Seebeck coef-

ficient, so that the power factors (5 9 10-4 W/mK2 at

1,000 K) exceeded those of the undoped Magneli phases,

Fig. 11. Unfortunately the possibility of forming CS defects

is suppressed by doping with pentavalent niobium, so that

the lattice thermal conductivity of niobium-doped rutile

remained much higher (2.9–4 W/mK at 1,000 K) than that

of Magneli phases with their high CS densities, Fig. 11. As

a consequence, the overall ZT of Nb-doped, reduced rutile

ranged only in between the best Magneli phases. In sum-

mary, niobium was found to act as an efficient electron

donor, but stabilized the rutile structure so that benefits for

the thermal conductivity were restricted to scattering by the

heavier niobium atoms and could not be combined with the

more efficient phonon scattering at CS defects.

We fabricated several mixed composites of Nb-doped

rutile and Magneli phases in the hope of producing phonon

scattering at the CS defects and achieve an overall benefit,

but, even in case of fine nanocomposites, did not observe

any advantage in ZT.

We further tried to enable the CS defect formation in

niobium-doped titanium oxide ceramics through targeted

co-doping with trivalent cornerstone atoms, which we

expected to introduce and fix the CS defects, but not to

contribute to the electrical conductivity. The concept was

explored through experiments for Al, Ga or Y addition.

Unfortunately, the optimization of the carrier concentration

at highest possible shear defect density required, based on

Fig. 10 Plot of lattice thermal conductivity as function of the

logarithm of the electrical conductivity for rutile single-crystal,

polycrystalline rutile, nanocrystalline rutile, nanocrystalline anatase,

several Magneli phases and doped rutile


123

our estimates, about 10 % trivalent co-dopant cations, a

concentration that exceeded the solubility and caused phase

transformation or second-phase formation.

We explored the effect of other dopants at levels

between 1 and 3 % for an exemplary composition of

approximately Ti(D)O1.94 (D standing for dopant). We

fabricated homogeneous powders from precursors and

densified them by current-assisted sintering. Plots of See-

beck coefficients and lattice thermal conductivity as func-

tion of the (log) electrical conductivity provided insight in

the response of the different dopants. Figure 12 shows that

the doped materials all followed the straight line of the

reduced pure titanium oxide phases in the Jonker plot,

showing thereby that doping and substitution did not break

the strong coupling between the electrical properties. Ta

and Mo produced compositions with high electrical con-

ductivity, but showed a similar disadvantage of very high

lattice thermal conductivity as niobium. Vanadium doping

produced fewer changes. Surprisingly, homovalent doping

with zirconium also yielded high conductivity and high

thermal conductivity. Most dopants, such as Fe, V, Sn, Al,

produced only small conductivity modifications. Co-dop-

ing with Nb ?Al, Li ? Ta, Sr ? Nb substantially

decreased the electrical conductivity, but preserved the

high lattice thermal conductivity of simple Nb or Ta

doping. The overall benefits from doping for the figure of

merit can be illustrated by plotting the power factors as

function of the lattice thermal conductivity. Drivers for the

figure of merit, such as Sn, are found in the right bottom

corner of the plot; while materials in the left upper corner

of the plot that includes Nb, Ta, W, lower ZT.

Fig. 11 Thermoelectric properties of unreduced and reduced undoped and Nb-doped rutile and Magneli phases

Fig. 12 a Lattice thermal conductivity (green) and Seebeck coeffi-

cient (blue) as function of the logarithm of the electrical conductivity

for doped TiOx; x was kept approximately constant, dopant nature and

levels are indicated for each data point


123

Within the range of dopants that we explored, we did

not achieve an increase in ZT compared to the best

undoped Magneli phases, but we were able to reach

similar ZT through niobium doping of rutile. The donor-

or acceptor-dopant induced changes in carrier concentra-

tion brought no advantages due to the strong coupling of

electrical conductivity, Seebeck coefficient and carrier

thermal transport.

Introduction of mass fluctuations in the oxide lattice

through substitution of Ti by heavier homovalent atoms,

such as Sn, Zr, proved as efficient. For 1–3 % Sn substi-

tution, a 20 % reduction in lattice thermal conductivity was

reached at high temperature, while zirconium substitution

provided only a 10–15 % decrease.

9.2.2.4 Composite approach Decreasing the lattice

thermal conductivity through enhanced phonon scatter-

ing at fine second-phase particles is a common approach

used in semiconductor thermoelectrics. In order to study

the effect of second-phase dispersions in Magneli

composites, we fabricated a number of composites with

second phases of TiC, TiN, WC, SrTiO3. For compos-

ites with TiC and TiN, their nanopowders were mixed

with rutile nanopowder, while all other composites with

oxide second phases, WC or metal were made via

hydrolysis of mixed precursors to ensure homogeneity

and small grain size.

Composites of TiO2-x with TiC, TiN and WC were

reactive-phase mixtures that were not brought to thermo-

dynamic equilibrium, but only partially reacted and then

quenched to a non-reactive state. Those composites cov-

ered a wide range in conductivity and Seebeck coefficient

depending on phase ratio and processing conditions, but

exhibited the same coupling of electrical properties as the

pure Magneli phases. For small volume fraction of finer

WC nanoparticles, a minor improvement in lattice con-

ductivity was observed. No advantages for the lattice

conductivity were observed for TiC and TiN second pha-

ses. This was attributed to an increase in the composite

lattice thermal conductivity due to the highly conducting

TiC or TiN fraction, so that any benefit of the second-phase

dispersion was lost.

It was expected that dispersions of insulating oxide

particles were more efficient. We observed that composites

of Magneli phases with small level of alumina or SrTiO3

preserved the electrical properties of the dispersion-free

materials up to secondary-phase fractions of 3 %; for

higher dispersion levels, degradation in electrical conduc-

tivity was observed. The introduction of oxide second-

phase dispersion in Magneli composites produced in some

cases a further decrease in lattice conductivity. Thus,

additional SrTiO3 nanoparticles provided 10–15 % further

decrease in lattice conductivity.

Even for a simple system, such as silicon with sec-

ond-phase dispersions, a clear ranking of drivers for

lattice conductivity reduction by heavy mass, interface

structure, second-phase particle size or particle distri-

bution has not been fully determined. In the case of

oxides, such consideration may be more complex, but

also much simpler, when considering the mean free

phonon path of few nanometers. A mean free phonon

path of 1–2 nm suggests that even nanoprecipitates are

too large in size to cause any substantial scattering. A

further disadvantage is the low-volume fraction of sec-

ond-phase dispersions. It is then evident that the CS

defects in the Magneli phases with their high densities

and interspacing of 1–4 nm are highly suited for phonon

scattering and, as demonstrated, produce lowest lattice

thermal conductivities.

Based on the above result (with ZT optimized com-

posites being only slightly better than PF optimized com-

posites), it is not obvious that a ZT optimization of doped

and second-phase Magneli-phase composites can provide

any significant improvement in ZT beyond the perfor-

mance of pure Magneli phases.

9.2.2.5 Comparison of titanium oxide-based materi-

als The Jonker plot of rutile, anatase and pure Magneli

phases was extended by including doped materials and

composites, Fig. 13a. Materials with dopants such as Sn, V,

W, Zr, Mo, Nb, Fe or second phases such as TiC, TiN, alu-

mina, SrTiO3 all match the Seebeck coefficient–electrical

conductivity relationship of the pure Magneli phases. We

conclude therefore that dopants or nanoparticles provide no

benefit for the electrical properties; they only drive the

properties along the electrical conductivity scale. While for

most dopants such shifts were small, they were very large for

niobium and provided very high electrical conductivity.

Extension of the comparative thermal conductivity

analysis, Fig. 13b, showed that the addition of dopants or

second phases affected the lattice thermal conductivity and

that, compared to the pure Magneli phases, certain dopants

provided substantial advantages or disadvantages at simi-

lar electrical properties. The enhancement in electrical

conductivity and Seebeck coefficient in the presence of

Nb, Ta, W, Mo was coupled with a stabilization of the

rutile phase and the impossibility of forming CS defects,

so that these heavy n-type dopants did not provide the

expected advantage. Addition of homovalent Zr and

especially Sn revealed the expected effect of heavier mass

and an associated decrease in lattice conductivity, espe-

cially for tin.


123

9.3 Niobium oxide

9.3.1 Thermoelectric properties reported in the literature

Electrical conductivity classifications of Nb2O5-x block

structures were derived from the filling level of the con-

duction band (2 electrons per missing oxygen) by Ruescher

[97]. Ruescher proposed band-like conduction for high

carrier concentrations and polaron conduction for NbO2.5-x

with x \ 0.05 (carrier concentrations \1.7 9 1021 cm-3

and 1 nm average distance between adjacent polaron cen-

ters). It was further suggested that bipolarons and single

polarons coexist in a metal–insulator transition regime with

0.01 \ x \ 0.05, while NbO2.417 was suggested to possess

localized electrons and quasi-free electrons and the end

member of the block-type series NbO2.4 was expected to

show metallic conductivity. Based on studies of the

anisotropy in properties, Ruescher advanced the hypothesis

of a quasi one-dimensional charge carrier system with

insulator-like behavior perpendicular to the block columns

and metallic behavior along the block columns (b-direc-

tion, needle axis) as a result of polaron overlap in this

crystallographic direction [97]. These considerations sug-

gested excellent thermoelectric properties for Nb2O5-x and

motivated us to explore the performance in this system.

Literature included electrical conductivity and Seebeck

coefficient measurements for a number of niobium oxides,

but no complete evaluation for the figure of merit.

Experimental data for NbO (made from sintered, oxi-

dized niobium powder) showed [98] very small thermo-

electric power (10–15 lV/K at 1,000 K) and electrical

conductivity below 1 S/m (1,000 K). Seebeck coefficients

of -100 to -200 lV/K (widely independent of stoichi-

ometry) at electrical conductivity of 100–20,000 S/m were

reported for NbO2 [98, 99]; the best power factors

(2 9 10-4 W/mK2 at 1,000 K) did not enable

ZT(1,000 K) [ 0.1. The same authors [100, 101] reported

Seebeck coefficients ranging from -100 to -500 lV/K

(1,000 K) for Nb2O5 with highest values close to the 2:5

stoichiometry. A similar set of data [102] projected slightly

lower electrical conductivities. Most of the reported

materials were made from coarse powders by natural sin-

tering; even though information on grain size and density is

lacking, the large differences in electrical conductivities

suggest that the grain boundaries in those materials acted as

strong barriers for transport (impurity segregation, inter-

granular glass, space charge layers or…).

Few results are published on doped niobium oxides.

Reported data include doping of NbO2 rutile with boron

[103] and vanadium [104]. While an increase in electrical

conductivity was achieved, the reports disagreed on the

impact on the high-temperature Seebeck coefficient.

9.3.2 Thermoelectric performance results from own work

In order to fully evaluate the potential of niobium oxide—

based materials, we conducted own research, fabricated a

wide range of undoped and doped niobium oxides and their

composites with micro- and nanograin size and determined

their properties.

Processing details We fabricated different niobium

oxide-based n-type materials that contained a single pure or

doped/substituted niobium oxide phase or a mixture or

composites with second phases such as NbO or metals

such as Nb, W, Mo, or carbide such as TiC, NbC, WC, or

nitride such as TiN or oxide such as TiO2. We used

micrometer (1–10 lm) or nanometer-size niobium oxide

starting powders that were mixed with dopant or second-

phase powders, partially reduced by carbothermal reduc-

tion (Nb2O5 ? C ? Nb2O5-x ? CO), disproportionation

(Nb2O5 ? x NbO (Nb) ? Nb2O5-x) or reaction with TiC,

TiN, NbC (Nb2O5 ? x TiC ? Nb(Ti)2O5-x ? CO) and

cold-pressed to pellets. The pellets were sintered in sealed

Fig. 13 a Jonker plot and b lattice-versus-electrical conductivity plot of pure and doped titanium oxide-based materials and their composites,

temperature 1,000 K


123

ampoules or rapidly densified by current-assisted sintering

under pressure in a Spark Plasma Sintering machine under

3–20 MPa at 1,000–1,300 �C. Some samples were post-

annealed under different oxygen partial pressure–temper-

ature conditions. Most materials had 95 % and more den-

sity. In the presented own data, we extrapolated all material

properties to 100 % density to enable best comparison.

9.3.2.1 Thermoelectric properties of pure niobium oxi-

des The thermoelectric properties of niobium oxide

ceramics were systematically explored. All niobium oxides

were n-type conductors. NbO was an excellent conductor

with more than 200,000 S/m, a Seebeck coefficient of

25 lV/K and high lattice conductivity of 8 W/mK

(1,000 K). NbO2 ceramics showed lower electrical con-

ductivity, 15,000 S/m, Seebeck coefficient of 150 lV/K

and lattice conductivity of 3 W/mK (1,000 K). Pure Nb2O5

materials exhibited the low electrical conductivity of an

insulator (2,600 S/m at 1,000 K) with Seebeck coefficient

(220 lV/K), so that reduction was needed to achieve good

thermoelectric performance. The beneficial effect of CS

defects in Nb2O5 was immediately apparent in its low

lattice conductivity, which was without further optimiza-

tion in the order of 2 W/mK.

Sub-stoichiometric Nb2O5-x, NbO2 and Nb12O29 pure-

phase materials were obtained by disproportion reaction

followed by slow natural sintering or rapid current-assisted

reactive sintering; the Nb47O116 composition showed an

Nb2O5-x-type-related structure. Figure 14 shows typical

microstructures of the materials. The thermoelectric per-

formance of the sub-stoichiometric phases and their mix-

tures are shown in Fig. 15. Compositions with Nb2O5-type

structures clustered in one group with low electrical con-

ductivity (\5,000 S/m at 1,000 K) with negative temper-

ature dependency and a moderate Seebeck coefficient

(-150 to -300 lV/K at 1,000 K) that improved with

temperature. Compositions in the range from Nb12O29 to

NbO2?x formed a second group of materials with semi-

conductor-type electrical conductivity and reached about

30,000 S/m at 1,000 K with smaller Seebeck coefficients

around -100 lV/K at 1,000 K. Both material groups

showed the same low lattice thermal conductivity of about

2 W/mK2 at high temperature, but differed at low tem-

perature. The Nb12O29 group adopted higher lattice con-

ductivity, while Nb2O5-type materials demonstrated

temperature-independent lattice conductivity. The figure of

merit of the Nb12O29-based materials was higher and

reached ZT(1,000 K) = 0.16, Fig. 15.

9.3.2.2 Thermoelectric properties of niobium oxide-based

composites Nb2O5-x composites with NbC, WC, TiC, TiN

underwent some reaction during rapid current-assisted den-

sification and formed additional reaction products; some of

the latter arranged as shell around the second-phase particles,

Fig. 14b. For a wide range of composite types, phase ratios

and their associated widely varying electrical conductivity, a

rather fixed high-temperature Seebeck coefficient of about

-100 lV/K was found. The lattice thermal conductivity of

the composites was increased compared to that of pure

substoichiometric Nb2O5. This was attributed to a volume–

fraction-based contribution of the higher lattice conductivity

of the second phases. The second-phase dispersions appar-

ently did not induce any visible advantage for the lattice

conductivity through enhanced phonon scattering. We come

back to our earlier findings for titania Magneli phases; the

dimensions of the second-phase nanodispersions were again

too large compared to the phonon mean free path.

ZT(1,000 K) = 0.2 was reached for several composites,

again a trade-off between electrical conductivity and See-

beck coefficient had to be found. Results for various com-

posites with TiC, TiN, SiC and WC are summarized in

Fig. 16.

9.3.2.3 Comparison of niobium oxide-based materials In

the Jonker plot, Fig. 17, NbOx-based materials aligned on a

single line with a slope in the order of �(k/e) and thus

differentiated from the ideal semiconductor behavior that

was also adopted by TiO2 materials. The deviation from the

Fig. 14 a Nb12O29 with minor amount of Nb2O5 second phase and b NbO2–TiN composite


123

ideal semiconductor behavior suggested some level of

degenerate semiconductor behavior. Best niobium oxide-

based materials covered a similar range in electrical

properties as TiO2-derived materials, but their maximum

power factor occurred at higher electrical conductivity and

significantly lower lattice conductivity. The low lattice

conductivities are related to the block structure with its two

sets of crystallographic shear planes at interdistances as

small as 0.7 nm, which offered highly effective scattering

sites for the oxide phonons. The impact of stoichiometry,

W, Mo and Ti dopants and second-phase particles TiC,

TiN, NbC, WC and metallic W were investigated. Com-

posites with TiN or W second-phase addition reached the

best performance a maximum ZT = 0.23 (1,000 K).

While titania-based materials respected the strong cou-

pling between electrical conductivity and Seebeck coeffi-

cient with a penalty in Seebeck coefficient for every

increase in electrical conductivity, in the niobium oxide-

based materials, a certain degree of decoupling was

achieved at high temperature. Composites showed a rela-

tively constant Seebeck coefficient over a wide range in

electrical conductivity. Such decoupling was typically

observed for other degenerate semiconductors, such as

SrTiO3 and misfit cobaltites.

9.4 Tungsten oxides

Since the literature reports included only limited information

on high-temperature conductivity and Seebeck coefficients

[105], we synthesized various WOx polycrystalline ceramics

by solid-state reaction and spark plasma sintering and deter-

mined their high-temperature thermoelectric properties.

Processing details WO3-x materials were made by

mixing WO2 and WO3 powders, reacting them under

vacuum and densifying them by spark plasma sintering.

The starting powders were tungsten (IV) oxide (Alfa

Aesar 40367, WO2, 99.9 %, -100 mesh) and tungsten

Fig. 15 Electrical conductivity, Seebeck coefficient, lattice thermal conductivity and ZT as function of temperature for niobium oxides with

different reduction levels


123

(VI) oxide (Alfa Aesar 13398, WO3, 99.998 %). Stoi-

chiometric batches of the two powders were weighed

and mixed to make WO3-x compositions: W18O49

(WO2.722), W12O34 (WO2.833), W20O58 (WO2.9), W25O73

(WO2.92), W25O74 (WO2.96). The mixtures were balled-

milled for 1 h in a planetary ball mill. The milled

powder mixtures were cold-pressed into pellets and

vacuum-sealed in quartz. The pellets were heat-treated

at 1,000 �C for 50 h, then re-milled and sieved to

–325 mesh. Powders were densified under rapid heating

in a graphite die in a spark plasma sintering machine.

Samples of each of the five reacted powders, the jet

milled WO3 powder and the ball milled WO2 powders

were sintered in the SPS. The maximum temperature

and hold time were 1,200 �C for 4 min with and applied

force of 35 MPa.

The transport properties of reduced tungsten oxides with

stoichiometries WO3-x, x = 0.04–0.28 and WO2 were

measured over a temperature range from 460 to 1,050 K.

Electrical conductivity was found to increase with

increasing reduction, with the exception of WO2, which

situated between W18O49 (WO2.72) and W12O46 (WO3.83).

A metal–semiconductor transition was observed at

x * 0.1. Seebeck coefficients were negative for all phases,

with absolute values increasing with increasing tempera-

ture and decreasing x (Fig. 18). The thermal conductivity

increased with increasing x, with WO2 exhibiting an

exceptionally high lattice thermal conductivity (6.8 W/

mK) compared to the other phases (Fig. 18). The extremely

low lattice conductivity values at high x that had been

derived from the total thermal conductivity was not correct

and indicated that our extrapolation with use of the

Fig. 16 Electrical conductivity, Seebeck coefficient, lattice thermal conductivity and ZT as function of temperature for niobium oxide

composites


123

Wiedemann–Franz law did not apply. Over the temperature

and composition range of the study, Seebeck coefficients

varied from -23 to -94 lV/K and did not fit the Jonker

model for non-degenerate semiconductors, Fig. 19. The

figure of merit changed with the tungsten oxide stoichi-

ometry and evolved through a maximum; WO2.9 exhibited

the highest thermoelectric figure of merit, ZT(1,050 K) *0.10 (Fig. 18).

Fig. 17 a Jonker plot and b lattice-versus-electrical conductivity plot including a large number of undoped and doped niobium oxide and their

composites

Fig. 18 Thermoelectric properties, electrical conductivity, Seebeck coefficient, lattice conductivity (extrapolated from the lattice conductivity

by using the Wiedemann–Franz law) and figure of merit ZT for tungsten oxide ceramics WO32x at 1,050 K


123

9.5 Niobium-based double perovskites

Double perovskites adopt a layered crystal structure of

slabs of perovskite (SrNbO3) and additional SrO layers that

are stacked with a shear-type defect. The combination of

layered crystal structure, heavy elements Sr and Nb and

even heavier potential dopants such as La or W trigger for

these compounds expectations of extremely small lattice

conductivity. In the literature, effectively, very low con-

ductivity at low temperatures was reported perpendicular to

the stacked layers in single crystals [25]. We hoped to

achieve a benefit in high-temperature properties for poly-

crystalline ceramics and processed undoped and doped Nb-

based double perovskites of type SrNb2O6, Sr2Nb2O7,

Sr4Nb4O14 and Sr5Nb5O17 with heavy element doping (La

and W) and various reduction states. Representative

microstructures of the double perovskites are shown in

Fig. 20; they illustrate the platy grain shape of the layered

crystal structure. XRD confirmed that the as-processed

densified ceramics adopted a double perovskite structure.

Seebeck coefficients of the as-processed materials ranged

from 200 to 400 lV/K, but the electrical conductivity

remained very small.

View that both, SrTiO3 perovskite and Nb oxides,

benefitted from partial reduction, we tried to reduce the

double perovskites during processing. Sr2Nb2O6 showed

significant improvement in electrical conductivity, but the

Seebeck coefficients dropped below 100 lV/K. Higher

members of the series decomposed upon reduction under

formation of insulating SrNbO3 perovskite.

The thermal conductivity of undoped and doped dou-

ble-perovskite ceramics met the expectations of low lat-

tice conductivity, Fig. 21. The lattice conductivity was

found to be in the range of 1.5 W/mK (constant from RT

to 1,000 K). An increase in lattice conductivity was

noticed upon reduction due to formation of normal

perovskite.

Double perovskites with their extremely strong anisot-

ropy can almost be considered as one-dimensional con-

ductors. They would have been ideal in the scope of this

review to illustrate the effect of anisotropy and texturing on

ZT, but unfortunately we were not able to identify a

composition–reduction window in our work that provided

acceptable power factor of at least 1 9 10-4 W/mK2.

For that reason, we used Ca3Co4O9 materials for illus-

trating effects of anisotropy and texturing.

Fig. 19 Jonker plot and lattice-versus-electrical conductivity for tungsten oxide ceramics WO32x at 1,050 K

Fig. 20 Representative microstructures of double perovskites


123

9.6 Ca3Co4O9

9.6.1 Literature reports on thermoelectric properties

A large number of experimental and modeling results were

reported in the literature on misfit cobaltites. Both, first

principles band structure calculations in combination with

standard Boltzmann transport theory, where correlation

effects were neglected, and application of the Hubbard

model (paradigmatic model for strongly correlated sys-

tems) predicted large power factors. Exceptionally high ZT

(extrapolated ZT = 1 at 1,000 K) was measured for small

single crystals [10]. Polycrystalline materials were so far

not successful to reach that performance. The reason is a

comparably much lower electrical conductivity that is

caused by the combination of lattice anisotropy of electric

and thermal transport properties, distribution in grain ori-

entation and high grain boundary resistance. Numerous

attempts were made to reach the single-crystal performance

by texturing. Different approaches were used, including

template growth from aligned Co(OH)2 precursor platelets

[106], magnetic field alignment [107] and sinter forging

[108]. Texturing was reported to improve the material

performance (ZT(1,000 K) = 0.25 [106]).

Cation substitution for both elements and in both layers

was widely studied, but did not indicate clear trends. The

large scattering and even contradiction of reported results

may be related to some extend to an assessment of elec-

trical and thermal properties in different sample directions,

which resulted for anisotropic materials in highly overes-

timated ZT. From the large number of published results, we

identified processing as a dominant driver for properties.

While we do not want to discuss the impact of doping or

analyze in detail the different processing approaches, we

want to use some examples from our own work to illustrate

the importance of the Co-oxidation state throughout pro-

cessing and contribute to the understanding on the impact

of texturing on the properties.

9.6.2 Results of own work

9.6.2.1 Impact of the Co-oxidation state throughout pro-

cessing on final material performance We used in our

work the solid-state templating reaction that was first

introduced by Tani [109]. Ca3Co4O9 starting powder was

made by solid-state reaction from CaCO3 powder and

Co(OH)2 precursor platelets. Reaction temperature and time

and especially the quality of the Co(OH)2 precursor plate-

lets had a strong impact on the quality of the Co3Co4O9

powder. Cobalt hydroxide was fabricated in a hydrothermal

process under nitrogen flow from cobalt chloride solution

by slowly adding sodium hydroxide solution and forming a

fine blue–green hydrated precipitated that, during aging,

turned under water loss into a fine pink precipitate that was

carefully dried and stored under exclusion of air.

Co H2Oð Þ6� 2þþ2OH�

blue�green

! Co H2Oð Þ4 OHð Þ2�

þ 2H2Opink

! b-Co OHð Þ2#pink

Co(OH)2 and CaCO3 powders were mixed in stoichi-

ometric ratio and reacted at temperatures between 800 and

950 �C in air; best results were obtained for reaction at

850 �C for 8 h. The reacted powders were densified by

direct current sintering (SPS). Best results were obtained

for a hold of 10 min at top temperature 800 �C under

70 MPa.

Even small traces of oxidation of the Co(OH)2 precursor

at any stage of the process produced lower-quality cobaltite

samples with substantially lowered electrical conductivity

(at same Seebeck coefficient and lattice conductivity),

Fig. 22. The Co(OH)2 templates were very sensitive to air

and moisture and readily degraded into cobalt(III)

hydroxide in the presence of both air and moisture (visible

as brownish coloration of cobalt III hydroxide). Traces of

cobalt(III) hydroxide produced undesirable by-products

during the solid reaction with calcium carbonate or other

Fig. 21 Seebeck coefficient and lattice conductivity of various double perovskites with different degrees of reduction


123

Ca sources, such as Ca3Co2O6 (chain structure with low

conductivity), lowered the electrical conductivity of the

final materials and resulted in low ZT.

The impact of various other processing parameters, such

as platelet size, solid-state reaction environment, raw

material sources, densification parameters, etc., was stud-

ied; no other parameter had as much impact as traces of

oxidation of the precursor platelets.

9.6.2.2 Impact of texturing We explored two different

approaches to process textured ceramics from Ca3Co4O9

powders that were obtained by solid-state reaction from

Co(OH)2 precursor platelets.

In the first approach, Ca3Co4O9 containing slip was

tape-casted; the thin tape was stacked and calendared into a

thick film; thick films were stacked in a die and densified

by high-temperature pressing in the SPS. Materials that

were obtained by natural sintering or hot-pressing were

found to exhibit an unexpected random orientation of their

platy grains (texture factor 0). While the grain size could be

modified by sintering time, temperature, oxygen partial

pressure and applied pressure, such process changes did

help to produce any texturing of the final ceramic. SEM

studies revealed the presence of agglomerates in the slurry

that inhibited the alignment of the individual plates in the

tape casting process and provided a random distribution of

the spherical agglomerates and their plates, Fig. 23b. The

resulting tape-cast ceramic had an in-plane (in-tape) elec-

trical conductivity that was slight smaller than in sintered

powder pellets (9,000 S/m at 173 lV/K instead of

13,000 S/m at 150 lV/K, 1,050 K) and had a similar figure

of merit, 0.14 instead of 0.15, as material made from

pressed powder.

In the second approach, Ca3Co4O9 powder was sinter-

forged. In a first step, powder was pressed into a pellet

under low force (30–60 MPa) in the SPS (1,073–1,173 K,

5–10 min). In the second step, the pellet was placed into a

larger die and repressed under higher force (decreasing

Fig. 22 Electrical conductivity Seebeck coefficient, lattice thermal

conductivity and figure of merit ZT as function of temperature for

Ca3Co4O9 dense ceramics that were made with exactly the same

processing parameters, but from Co(OH)2 precursor with different

oxidation levels (Co3? content)


123

force over the pressing process from 120 MPa to 60 MPa)

at 1,123–1,173 K. Strongly aligned microstructures with

texture factors in the range of 0.6–0.9 (texture factor 1 for

perfect alignment) were obtained, Fig. 24c. The alignment

improved with forging temperature and force and benefit-

ted from slow heating rates. We measured the thermal

transport in pressure-forged samples parallel and perpen-

dicular to the forging direction, see Fig. 24. Compared to

untextured materials with random plate distribution, all

forged samples showed improved in-plane electrical con-

ductivity perpendicular to the forging direction,

15,000–18,000 S/m at relatively constant Seebeck coeffi-

cient, 158–168 lV/K. The lattice conductivity in samples

with random orientation ranged from 1.5 to 1.8 W/mK

(1,000 K); in the textured samples, it was lower in the

direction perpendicular to the layers, 1.0–1.2 W/mK and

higher in-plane, 2.7 W/mK. Combination of electrical and

thermal transport data provided exactly the same figure of

merit of ZT(1,000 K) = 0.15 for random materials and for

textured material in the in-plane direction, Fig. 25. The

random and textured materials of Fig. 25 were all made

from the same raw materials and by the same process and

had similar densities. The results strongly suggest that

texturing does not provide any advantage, since not only

the electrical conductivity was increased, but also the

thermal conductivity. Our finding did not meet the expec-

tations of a performance benefit due to energy filtering in

the layered cobaltite structure. It also did not match liter-

ature results on textured ceramics that reported an

improvement of in-plane ZT through texturing [109, 110].

However, it cannot be excluded that the reported results in

[109, 110] were affected by use of different raw materials

for random and textured ceramics. Many published results

on ‘‘random’’ and textured Ca3Co4O9 materials used a

combination of in-plane electrical conductivity and out-of-

plane thermal conductivity and thus obtained a highly

overestimated figure of merit. Our Fig. 25 includes such

‘‘overestimated ZT’’ from in-plane conductivity and per-

pendicular to plane thermal conductivity, which is by more

than a factor 2 larger than the correctly measured in-plane

ZT and has to be considered not only as overestimated, but

false!

Our best undoped (Co3?-free) materials with random

microstructure reached a figure of merit ZT(1,050 K) =

0.32. Dopants were not found to provide significant further

Fig. 23 SEM micrographs of polished cross-sections of a pressure-sintered powder pellet, b pressure-sintered tape-casted film stack and

c pressure-forged pellet

Fig. 24 Sample preparation from a pressed disk for evaluation of

electrical and thermal conductivity; a and b show typical cuts that

lead easily to a mixed parallel and perpendicular assessment of

electrical and thermal properties and leads to overestimated ZT; e–

f cut and orientation of sample for in-plane thermal conductivity

measurement


123

improvement. Some initial work on further optimization of

processing and densification conditions indicated further

possible paths for improvement. Our textured materials did

not show any advantage compared to the same random

materials and were far away from the single-crystal per-

formance reported in the literature. The main reason must

be the high resistance of the grain boundaries. Detailed

analysis of the grain boundary chemistry and grain

boundary electrical properties would advance the under-

standing of causes for high resistance and provide ideas for

efficient grain boundary engineering.

10 Conclusions

Even though large bandgap semiconducting oxides are

promising candidates for high-temperature thermoelectric

power generation (700–1,200 �C) due to their high-tem-

perature stability, lack of toxicity and low cost, they typi-

cally reach only much smaller power factors than Bi/Pb

tellurides, skutterudites or silicon–germanium alloys.

Compared to other thermoelectrics, oxides have rather high

carrier concentrations, but low carrier mobility with more

localized charge carriers. Compared to intermetallics or

silicon alloys, the thermal conductivity of oxides is rather

small and, in some cases, can be even extremely small.

This general picture roughly describes many large bandgap

semiconducting oxides; it does not include any principle

obstacle for high thermoelectric performance. However,

experience from the literature and our own work showed

that the strong coupling of the transport properties makes a

thermoelectric performance optimization difficult.

In order to assess how such shortcoming can be best

overcome, we reviewed the concepts and approaches that

were successfully applied for other thermoelectric semi-

conductors and examined their value for oxide materials.

Performance of those other semiconductors was mainly

limited by their carrier concentration; the carriers had high

mobility. Therefore, most approaches focused on increas-

ing the carrier concentration through doping and were

rather successful, since contributions of carriers to the

thermal transport remained low for small carrier concen-

trations. In oxides, the situation is different. Doping oxides

with their already high carrier densities leads only to a

minor relative increase in carrier concentration and also

produces immediately a significant carrier-induced

Fig. 25 Thermoelectric properties of simply pressed, random and dynamically forged, textured Ca3Co4O9 with indication of the measurement

directions (in-plane and perpendicular to the conduction plane of the platelets)


123

increase in thermal conductivity. In addition, solubility

limits restrict doping in many oxides to small concentra-

tions; only in exceptional cases, levels above 1 % can be

reached. For the n-type conductors of this review, SrTiO3,

TiO2-x, Nb2O5-x, WO3-x, reduction or oxygen vacancy

doping turned out to be much more efficient than cation

doping and provided significant improvement in the power

factor. Even for materials with a small stoichiometry range,

such as SrTiO3 or rutile, substantial improvement was

achieved. This suggests that oxygen vacancy ‘‘doping’’

does not only increase the carrier concentration, but also

affects the carrier mobility. Our experimental results

demonstrated a power factor improvement for doped,

reduced SrTiO3 by addition of second reducing phases,

such as TiC or TiN, which effectively promote further

reduction. Literature also reported increases in the effective

mass for reduced doped SrTiO3-x and TiO2-x and illus-

trated that the increase in effective mass was much larger

in reduced SrTiO3 than in La- or Nb-doped SrTiO3. Based

on these findings, oxygen vacancy doping was identified as

a key driver for the electrical performance that allowed to

reach some decoupling of the electric transport properties.

The Jonker plots of SrTiO3 and NbOx effectively showed

the largest deviation from the ideal semiconductor behavior

for highly reduced materials. A comparison of different

material families in the Jonker plot showed that all titanium

oxide-based materials followed a single straight line over a

wide range of conductivity with the characteristic slope k/e

of a single-parabolic band semiconductor, while niobium

oxide-based materials aligned on a common line, but with

much smaller slope (about 0.5 k/e). SrTiO3-based com-

posites showed less alignment on a common line and

showed significant deviation from the k/e slope for high

electrical conductivity. The misfit cobaltites finally illus-

trated the behavior of a fully degenerate semiconductor

with almost constant Seebeck coefficient for a (extended)

range of electric conductivities. Compared to the n-type

materials of this review, the conductivity range that was

covered by the cobaltites was too limited (5,000–20,000 S/

m) to derive a meaningful Jonker plot.

We concluded that the more degenerate semiconductors

offered the highest potential for tuning their electrical

performance and thus achieving higher figures of merit.

We fitted the behavior of each material family (in its

most promising range in the Jonker plot) by a straight line

and computed the evolution of the power factor as function

of the carrier concentration, Fig. 26. The resulting curves

suggested a maximum power factor at 1,000 K of about

5 9 10-4 W/mK2 at an electrical conductivity around

20,000 S/m for titanium oxide-based materials and

5 9 10-4 W/mK2 at comparatively higher electrical con-

ductivity, about 100,000 S/m, for niobium oxide-based

materials, 3 9 10-3 W/mK2 at an electrical conductivity

around 50,000 S/m for SrTiO3-based materials, respec-

tively. We reached the predicted maximum power factors

for titanium and niobium oxides in our own experiments for

heavily reduced material (Magneli phases and block struc-

tures). From this, we concluded that no further improve-

ment in electric properties is expected from simple carrier

density optimization. For SrTiO3-based materials, neither

literature results for bulk materials, nor our own experi-

ments reached the predicted maximum power factor. All

our results were obtained for lower electrical conductivity

than the predicted optimum conductivity. Therefore, further

improvement of the material can be expected from an

increase in carrier density. While this may be very difficult

for bulk materials and traditional processing, thin film

technology offers opportunities to move far away from the

thermodynamic equilibrium. Effectively, results by Ohta

et al. on extremely high ZT (2.4) of an interfacial layer of

doped, reduced SrTiO3 encourage pursuing such path.

Our literature review and own work also had the goal to

identify drivers for low thermal conductivity in oxide

materials. In our experimental assessment, we studied

multicomponent oxides with a simple crystal structure

(SrTiO3 perovskite), high densities of planar crystallo-

graphic defects (single sets of crystallographic shear planes

in TinO2n-1 Magneli phases, intersecting shear planes in

NbOx block structures or more irregular or channel struc-

tures in WO3-x) and layered superstructures (Ca3Co4O9).

We investigated monolithic materials with different grain

sizes (including nanoceramics) and composites. Many of

these oxides had crystallographic or microstructure fea-

tures in the size range of 0.3–20 nm. Our experimental

results allowed to compare the impact of doping, grain size,

crystallographic defects, superstructures, second-phase

addition, texturing and (to a limited extend) processing on

the thermal conductivity and figure of merit.

We found that a high density of crystallographic defects,

such as crystallographic shear planes or superlattice

structures, were by far the best drivers for low lattice

conductivity. They were followed by high levels of do-

pants. Nanograin size and second-phase dispersions played

only a secondary role for phonon scattering in oxides due

to the difficulty of matching their dimension to the phonon

mean free path (\2 nm). For layered structures, we con-

firmed the benefit of superlattices for the thermal conduc-

tivity. Ca3Co4O9 and double perovskites reached very low

lattice conductivity in the range of 1.5 W/mK, which is

close to the conductivity of their corresponding random

compounds.

Literature suggests that texturing of layered oxides

allows us to reach higher figures of merit, especially for

incommensurate crystal structures like Ca3Co4O9. Effec-

tively, many efforts were made to texture misfit cobaltites;

and some benefits were reported in the literature. However,


123

for many of the reported results either material processing,

raw materials or property, measurements were changed

between the assessment of random and textured materials.

We have shown that processing is a key driver for the

thermoelectric properties of Ca3Co4O9, so that small vari-

ations in the fabrication can easily drive the changes that

were reported in the literature. In addition, many literature

reports presented a highly overestimated figure of merit by

assessing electrical and thermal properties in different

texturing direction (high in-plane electrical conductivity

combined with low perpendicular thermal conductivity).

We have shown in our work that such false assessment can

overestimate the figure of merit by a factor 2! We prepared

random and textured materials using exactly the same

process and showed that random materials and textured

materials (in-plane) had exactly the same figure of merit.

The higher in-plane thermal conductivity compensated the

benefit of higher in-plane electrical conductivity in the

textured material.

In conclusion, we identified a number of key drivers for

electrical and thermal properties in oxide thermoelectrics

and provide guidelines for future material research.

Acknowledgments The authors want to thank Kim Work, Michelle

Wallen, Robert Fretz, Indrajit Dutta, Ron Davis, Mike Carson, Ron

Parysek, Bryan Wheaton, Erica Stapleton, Teresa McDermott and

Andrew Russell for materials characterization and materials pro-

cessing and acknowledge Deenamma Vargheese’s and Todd St.

Clair’s contributions to the larger area of this research and mention

discussions and support by many other colleagues from Corning

Incorporated.

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