Master’s Thesis

Potential enhancement of thermoelectric energy conversion in cobaltite superlattices

Tasos Englezos - S1463144 Enschede 03/09/2015 Nanotechnology University of Twente Faculty of Science and Technology Inorganic Materials Science

Graduation committee: Prof. dr. ing. Guus Rijnders Prof. dr. ir H.J.W. Zandvliet Dr. ir Mark Huijben

I

II

Abstract

Interest in the research on thermoelectric materials was reenacted in the past fifteen years

due to developments in the field of nanotechnology. Through novel methods, tools and techniques,

researchers gained a better understanding of nanoscale physical and chemical properties as well as

the capability to tune them. Through various types of nanoscale manipulation of mater such as carrier

doping, epitaxial growth, defect and interface structural engineering, attributes such as carrier density

and mobility, band gap and conductive channels, phonon conductivity etc. can be tuned in such a way

to improve the efficiency of thermoelectric materials.

In this research we investigate the combination of two single crystalline cobaltite oxides,

NaxCoO3 and Ca3Co4O9 into one epitaxially grown thin film superlattice. Both materials have proven to

be interesting and promising for thermoelectric technology, with NaxCoO3 being more electrically

conductive but less chemically stable while Ca3Co4O9 having a higher Seebeck coefficient and being

chemically inert under ambient conditions and high temperatures. The in-plane crystallographic

similarities of the “rock salt” CoO2 layers which are common in the structure of both cobaltites, makes

it possible to grow the two materials on top of each other stacking them along the c-axis.

The superlattice films were grown at 430oC and partial O2 pressure in a common Pulsed Laser

Deposition setup. The Substrate used was (La0.3Sr0.7)(Al0.65 Ta0.35)O3, commonly referred to as LSAT.

Due to its cubic unit cell, the mismatch between film and substrate should be quite large however

under certain orientation a 12-fold symmetry with the in-plane parameters NaxCoO3 is possible.

The thin films grown this way have been scanned in the θ/2θ direction and from the diffraction

pattern it is evident that their crystalline plane coherency is maintained while growth is oriented along

the c-axis. The d-spacing of the superlattice samples was found to be in between the values of the

individual films.

Two different types of superlattice growth where investigated in this work: a) Keeping a fixed

period thickness at 10:10nm for the two materials while increasing the total film thickness and b)

Keeping the total film thickness constant at 140nm while varying the number of periods and hence

the period thickness. Both methods gave useful insight regarding the grain morphology, the

crystallinity and the electronic qualities of the thin films.

For the fixed period thickness the best performing sample was the 60nm with sheet resistivity

and Seebeck values at 11.35mOhm*cm and 121.9μV/K respectively and a combined power factor of

1.3x10-4 W/m*K2. For the films with variable number of periods the best performance was for the 7-

period sample with sheet resistivity at 7.85 mOhm*cm, Seebeck of 82.6 μV/K and power factor of

0.8x10-4 W/m*K2.

According to the results acquired in this work it can be concluded that epitaxial growth of

NaxCoO3 and Ca3Co4O9 in a superlattice structure is possible in both combinations and that the

structural and electronic properties are maintained at a good level. Even though complete stability for

the NaxCoO3 could not be achieved, the samples are expected to show reduced thermal conductivity

in measurements that will be conducted in future work. If the thermal conductivity is indeed reduced,

then the superlattice approach for the cobaltite oxides could be proven to be a significant step

towards improving the efficiency of thermoelectric cobaltite oxides.

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IV

Contents:

1. Introduction 1

2. Theory on thermoelectricity 3

2.1. Thermoelectricity and the Seebeck effect 3

2.2. Electrical conductivity 5

2.3. Thermal conductivity 6

2.4. Figure of Merit 7

2.5. The thermoelectric module concept 9

3. Oxide materials, Thin films and Superlattices 11

3.1. Thermoelectric oxides 11

3.2. NaxCoO2 11 3.3. Ca3Co4O9 13

3.4. Thin films 14

3.5. Epitaxial cobaltite oxide superlattices 15

4. Fabrication and characterization Methods 16

4.1. Pulsed laser deposition 16

4.2. Sample preparation and deposition 17

4.3. Atomic Force Microscopy (AFM) Characterization 18

4.4. X-Ray Diffraction (XRD) Characterization 19

4.5. Reactive Ion Etching and deposition of gold contacts 21

4.6. Room temperature resistivity measurements 21

4.7. Room temperature Seebeck Coefficient measurements 22

4.8. Physical properties measurement setup (PPMS) , mobility and carrier concentration 23

5. Results and discussion 25

5.1. Work outline 25

5.2. Calibration samples 26

5.3. Ion etching profile calibration 28

5.4. CCO:NCO Superlattice: Constant layer thickness, variable total thickness 29

5.5. Single films 70 nm 33

5.6. CCO:NCO superlattice samples: Constant total thickness, variable number of periods 35

6. Conclusion and Recommendation 43

6.1. Concluding Overview 43

6.2. Experimental considerations 44

6.3. Recommendation for future research 45

Acknowledgements 48

Bibliography 50

V

1

1. Introduction

Motivation The global human population is steadily increasing and with it the demand for energy

resources is escalating. Moreover, recent reports on the global warming phenomena have

dramatically outlined the fact that there is imminent need to limit the consumption of fossil fuels for

our energy needs since the greenhouse gasses and pollution produced in the process, irreversibly

harm the earth’s environment and contribute to the global warming.

Evolution of nuclear power generation has contributed towards limiting the use of fossil fuels.

However high maintenance and equipment costs, issues related to the safe disposal of the toxic waste

by-products as well as the risks of nuclear meltdown, limit the use of nuclear power generation as an

alternative power source.

Other promising renewable power generation alternatives such as wind turbines and solar

panel power plants are becoming more and more efficient. However a large scale implementation is

usually needed in the form of “wind parks” or “solar power farms” in order to achieve adequate power

output towards the main power grid. This fact renders the implementation costly and environmentally

challenging.

One factor that is common in almost every energy production and energy consumption

method is heat losses. Commonly known as Waste heat, it is the dominant energy loss factor in the

majority of industrial applications nowadays. Loss of heat caused by friction, hot exhaust gasses,

resistances etc. can more than 60% (figure 1). This energy could potentially be harvested and recycled

into electricity directly by using the capability of thermoelectric energy conversion. This way the

output of the current power generation technologies can be improved and energy lost into heat during

consumption can be partially recovered into the main power grid. Moreover, standalone

thermoelectric power generators could be used wherever sufficient temperature gradients are

possible providing another alternative sustainable power source which can decrease the consumption

rates of fossil fuels.

Figure 1: Waste heat and potential recovery into usable energy by various fields in industry through the use of thermoelectrics.1

2

Thermoelectric (TE) power generation has been successfully used in some niche applications

and cheap low output TE modules are nowadays commercially available. Large scale application

however has been prohibited by factors such as poor output efficiency, material complications, high

temperature incompatibility, and costs related to the rarity and treatment required for thermoelectric

materials. The fact that thermoelectric modules do not involve any moving parts significantly lowers

the maintenance costs due to the increased reliability while also permits for scalability and makes

implementation much easier. To date, promising application of TE energy recovery has been in

automobiles, where a lot of waste heat is produced in the engine coolant or exhaust gas, which could

be recycled directly into energy for the car. TE power generation has also been widely used in space

technology where energy recovery is of the outmost importance.

The best performing thermoelectric materials, however, are either scarce and therefore

expensive, or contain toxic elements such as tellurium or antimony, which degrade when exposed to

high temperature air. Oxides suitable for thermoelectric research, are abundant in nature, nontoxic,

high temperature tolerant and offer tunable properties rendering them appropriate for a wide range

of possible applications (figure 2). Although their inherent thermoelectric properties are worse than

that of the previously mentioned elements, novel fabrication methods from a nanoscopic perspective

can be utilized to improve their thermoelectric performance.

Figure 2: Schematic comparison of various thermoelectric materials for waste heat recovery and refrigeration applications with respect to (a) the operational temperature and ecological friendliness and (b) in terms of abundancy. Adapted from2

3

2. Theory on thermoelectricity

2.1. Thermoelectricity and the Seebeck effect

The Seebeck effect is the direct conversion of temperature differences into a voltage

differential and hence into electricity. In 1821, Thomas Seebeck, a German physicist, realized that

when two different metallic elements which are joined in two places forming a closed circuit, while at

the same time they are held in different temperatures (ΔΤ) , a compass needle would be deflected.

The phenomenon was attributed to the different response of the metals. Due to their compositional

difference, to the temperature gradient formed between them, generating a current loop and a

magnetic field. The effect was termed "thermoelectricity" and it can be described as the way a

material responds to the temperature gradient applied to it in order to maintain its electronic balance.

Figure 3: Schematic representation of a one level semiconductor device the energy difference between the chemical potential μ and the energy of the conduction level is of the order of a few kT.

A relatively easy way to explain this effect is by using the bottom up approach for a small scale

one level device (elastic resistor) where the free electron approximation is valid (figure 3). The two

sides of the device in this example is the same material with its two sides held at different

temperature. The energy level (E) of the conduction band can be approximated by the parabolic

dispersion relation with respect to the electron wave number (κi) and the directional effective mass

of the electron (𝑚𝑖∗) with (i=x,y,z):

𝐸3𝑑(𝜿) =ħ2

2(∑

𝜅𝑖2

(𝑚𝑖∗)

2

𝑖

)

And the density of the electronic states corresponding to that energy, D(E), is given by:

𝐷(𝐸) =1

2𝜋2 (2⟨𝑚∗⟩

ħ2 )

3/2

𝛦1/2

Where:

⟨𝑚∗⟩ = √𝑚𝑥𝑚𝑦𝑚𝑧3

And (ħ) is the reduced Plank constant.

4

Using expression for the Fermi distribution function f(E,μ,T), where (E) is the energy level of the

conduction band, (μ) is the chemical potential (T) is the temperature and (k) the Boltzmann constant;

𝑓(𝐸, 𝜇, 𝑇) =1

𝑒(𝐸−𝜇)

𝑘𝑇 + 1

and the overall electronic conductance (G):

𝐺 = ∫ 𝑑𝐸 (−𝜕𝑓

𝜕𝐸) 𝐷(𝐸)

We can estimate the current (I) running through the device:

𝐼 =1

𝑞∫ 𝑑𝐸 𝐷(𝐸)(𝑓1 − 𝑓2)

With (q) being the charge of the carrier.

For the hypothetical device with one conduction level and for very small variations in temperature

and chemical potential between the two contacts, the current can be approximated through Taylor

expansion as:

𝐼 ⋍ 𝐺 (𝜇1 − 𝜇2

𝑞) + 𝐺𝑠(𝑇1 − 𝑇2) = 𝐺𝛥𝑉 + 𝐺𝑠𝛥𝛵

With the indicators 1,2 referring to contact 1 and 2 respectively and with (Gs) being the conductance

attributed to the temperature gradient:

𝐺𝑠 = ∫ 𝑑𝐸 (−𝜕𝑓

𝜕𝐸) 𝐷(𝐸)

𝐸 − 𝜇

𝑞𝑇

Expressed in terms of voltage difference we get:

𝛥𝑉 =1

𝐺𝐼 −

𝐺𝑠

𝐺𝛥𝛵

Where the ratio between (Gs) and (G) is the Seebeck coefficient:

𝑆 =𝐸 − 𝜇

𝑞𝑇

It is clear from the derivation above that due to the nature of the Fermi function, current flow

can be obtained not only by differing the chemical potential, but also by different temperature without

application of external voltage. The physical aspects of this derivation are explained in figure 4.

5

Figure 4: Schematic representation of the interaction between the Fermi function and the density of states for a hypothetical one level n-type semiconductor device.

At higher temperatures the Fermi distribution function is changing gradually over a range of

a few kbT, from zero to one; figure 4.a. At very low temperatures close to zero Kelvin the Fermi

distribution changes abruptly from zero to one along the chemical potential; figure 4.b. When these

two states are in contact, figure 4.c, the electrochemical potential is initially at the same level but the

difference of the Fermi function due to the temperature gradient, enables current flow from contact

1 to contact 2 (hot to cold in this example), for conducting states above the chemical potential and

from contact 2 to contact 1 (cold to hot), for conducting states below the chemical potential.

For a typical semiconductor the chemical potential lies roughly in the middle of the forbidden

energy band between the valence and conduction band which lies at energy (E). Due to the nature of

the density of states D(E) in semiconductors usually resembling a parabola, the interaction of the

Fermi distribution function with D(E) allows for conduction in the hot contact while it prohibits

conduction in the cold contact figure 4.d and when two sides are connected, current is allowed to flow

between them and the conduction electron population is the product of (f1-f2) and the Density of

states D(E) figure 4.e.

Therefore semiconductor materials with the conduction band way above the chemical

potential (large bandgap) are found to have a high Seebeck coefficient. However a very large bandgap

would significantly hinder electron conductance. This is the reason why usually the materials chosen

for thermoelectric research have a bandgap such that it allows for high Seebeck without limiting the

conductivity. I.e. materials with increased carrier mobility.

2.2. Electrical conductivity

The electrical conductivity (σ) is used to measure the freedom of charge carriers to move

through a material. For a crystal lattice it is given as the interaction between the electron charge (e),

the relaxation time between electron collisions (τ), the electronic carrier density (n) and the electron

effective mass (m*), by the Drude equation:

6

𝜎 =𝑒2𝜏𝑛

𝑚∗=

1

𝜌

Where ρ is the resistivity of the material. A relation between the charge, the collision time and the

effective mass is also expressed as the carrier mobility (μ):

𝜇 =𝑒𝜏

𝑚∗

By combining the two equations the conductivity can be expressed as a function of the carrier density

and carrier mobility:

𝜎 = 𝑛𝑒𝜇

2.3. Thermal conductivity

Thermal conductivity is a measure of the ability of a material to allow the flow of heat from

its warmer surface through the material to its colder surface. Understanding the mechanism of

thermal conductivity is a major factor in the research on thermoelectric materials. Thermal

conductivity is the parameter that affects the time under which the induced thermal gradient can be

maintained throughout a sample’s geometry as well as the magnitude of the temperature difference

that can be achieved. According to the theory by Debye and Peierls for a crystal, at the lowest

temperatures the thermal conductivity depends on the size and shape of the crystal and increases

with temperature in relation to the specific heat. The maximum thermal conductivity is limited by the

scattering of phonons and is characteristic of the material. Near the maximum, the thermal

conductivity is sensitive to the imperfections and impurities in the crystal lattice3.

Like electrical conductivity where the associated charge carriers are electrons or holes, the

parameter attributed to thermal conductivity is (k), and it has contribution from the electronic charge

carriers (ke) as well as the lattice vibration modes (phonons) (kL).

𝑘 = 𝑘𝑒 + 𝑘𝐿

Where Ke can be related to the Electrical conductivity through the Lorentz factor (L) and the

temperature (T). This relation is given by the Wiedermann-Franz law4:

𝑘𝑒 = 𝐿𝜎𝑇 = 𝑛𝑒𝜇𝐿𝑇

The Lorentz factor for free electrons is:

𝐿 =𝜋2

3(

𝑘𝐵

𝑒

2

) = 2.45 ∗ 10−8 𝑊𝛺𝛫2

7

Since the lattice contribution kL cannot directly be measured, it is calculated as the difference between

the measured k and the electronic contribution. Hence there is need for accurate estimation of ke.

Electronic thermal conductivity

The electronic contribution to the thermal conductivity of a material is given by:

𝑘𝑒 =1

3𝐶𝑒𝑣𝑓𝑙𝑒 =

𝜋2𝑛𝑘𝛣2 𝑇𝜏𝑒

3𝑚𝑒∗

Where (Ce) is the electron specific heat, (νf) is the Fermi velocity, (le) is the electron mean free path

and (τe) is the average collision time of electrons.

Lattice thermal conductivity

Lattice thermal conductivity of a crystal is attributed to phonons and is determined by three

contributions: The frequency dependent specific heat of phonons (Cph), the phonon group velocity

(vph) and the mean free path of phonons (lph). It can be modeled by:

𝑘𝐿 = 𝑘𝑝ℎ =1

3𝐶𝑝ℎ𝑙𝑝ℎ𝑣𝑝ℎ

The mean free path of phonons is determined by two factors: the rate of scattering with other

phonons at high temperatures and by scattering with static impurities or boundaries in the crystal

lattice at lower temperatures. The transition between the two contributions is dependent on the

Debye Temperature (TD) of the material which can vary between 100-1000K. At high temperatures

lph is decreasing with 1/T.

The phonon specific heat at temperatures exceeding the Debye limit is given in its classical

form from the Dulong Petit law:

𝐶𝑝ℎ = 3𝑁𝑘𝛽

With 3N being the number of normal phonon modes, and is independent of temperature

2.4. Figure of Merit

The efficiency of energy conversion of a thermoelectric material is determined primarily by three

properties: (S) Seebeck coefficient, (σ) the electrical conductivity and (κ) the thermal conductivity of

the material. A simplified way to quantify this efficiency through the connection of all the properties

and the temperature of application (T) in the dimensionless figure of merit (zT).

𝑧𝑇 =𝑆2𝜎

𝜅𝛵

Where the nominator is also called the Power factor and it is indicative of how well a thermoelectric

material performs with respect to its electronic properties:

8

𝑃 = 𝑆2𝜎

In more detail the electrical conductivity is given by the equation:

𝜎 = 𝑛𝑒𝜇 =1

𝜌

And the Seebeck coefficient expressed as function of the effective mass and carrier density:

𝑆 =8𝜋2𝜅𝛽

2

3𝑒ℎ2𝑚∗𝑇 (

𝜋

3𝑛)

23⁄

The rest of the factors (ρ) is the resistivity, (n) is the carrier density and (μ) the carrier mobility of the

material, (m*) is the electron effective mass and (kβ), (h) the Boltzmann constant and Plank’s

constant respectively.

It can be easily understood that in order to enhance the thermoelectric efficiency there are

three different paths to follow. Enhancing the Seebeck coefficient, enhancing the electrical

conductivity or lowering the thermal conductivity. However it has been proven that this is not a minor

task. The three properties are not by default decupled from each other and most of the time,

optimizing one factor comes at the cost of diminishing another. This fact can also be depicted in figure

5, where it is noted that the zT factor is optimal at a different carrier concentration value than the

power factor.

Another approach for enhancing the Seebeck coefficient is by increasing the effective mass

m* of the carriers i.e. by narrowing the bands via designing the density of states5 or via nanostructure

engineering6. However this approach may significantly reduce the mobility of the carriers, while there

are also studies supporting that higher performance can be achieved through an effective mass

reduction7.

Figure 5: Optimizing zT through carrier concentration tuning. The value range for the other parameters plotted against the y-axis are: S (0-500μVK-1), σ(0-5000Ω-1cm-1),k(0-10Wm-1K-1),adapted from4

9

Thermoelectric materials are usually heavily doped semiconductors and with carrier

concentrations in the range of 1019-1021 per cm3. A reduction in the lattice thermal conductivity can

significantly increase the figure of merit zT. Figure 6 denotes this fact.

Figure 6: Lower lattice thermal conductivity directly increases the zT and increases the Seebeck coefficient due to lower electronic thermal concentration ke. Plot is based on a model system (Bi2Te3), adapted from4

Through the above analysis of the thermoelectric properties, three paths leading to a potential

increase of the zT factor and thus the energy conversion efficiency of devices made of nanostructures

can be derived. 1) Introduce Interfaces and boundaries of nanostructures to constrain the electron

and phonon waves, which lead to a change in their energy states and correspondingly, their density

of states and group velocity. 2) Use of quantum size effects and classical interface effects to influence

the symmetry of the differential conductivity with respect to the Fermi level. 3) Utilize interface

scattering and induce variations of the phonon spectrum in low-dimensional structures in order to

reduce the phonon thermal conductivity.

2.5. The thermoelectric module concept

A typical thermoelectric module for power generation consists of both n-type and p-type

thermoelectric materials connected in series with a conductive material. A temperature gradient

applied across the module causes the charge carriers to diffuse towards the cold side, generating a

thermoelectric voltage. This way the electron and hole transport from the n-type and p-type materials

respectively, is additive and leads to the generated current. A Schematic diagram representing the

concept of such a module is presented in figure 7.

10

Figure 7: Schematic diagram of a typical thermoelectric module for electrical power generation. Components of n-type (red) and p-type (blue) materials are connected in series and then contained between ceramic substrates. Heat is applied to one side of the module, causing the charge carriers to diffuse across the module and generating an electrical current8.

It has to be clarified that the figure of merit zT described in the previous chapter, is only

referring to the material’s thermoelectric performance and not to the overall efficiency of the

thermoelectric module which contains those materials. This is primarily because the material

properties (S, κ, σ) are also dependent on temperature and due to factors related to the

interconnectivity and cumulative performance of all the parts co-existing in TEG module. The overall

performance of a thermoelectric generator is best described by the Carnot efficiency (η):

𝜂 =𝛥𝛵

𝛵ℎ

√1 + 𝑧𝑇 − 1

√1 + 𝑧𝑇 +𝑇𝑐𝑇ℎ

Where (Tc) and (Th) are the temperatures in the cold side and the hot side of the module respectively

and ΔΤ is the difference between them4.

11

3. Oxide materials, thin films and Superlattices

3.1. Thermoelectric oxides

Metal oxides are ionic compounds consisting of metal cations and oxygen anions alternately

paced and held together via attractive Coulombic interaction between them. In such ionic compounds,

the charge carriers (electrons or holes) polarize the surrounding crystal lattice by strongly interacting

with it, localizing themselves on the lattice points while inducing lattice distortion and limiting the

overlap of the atomic orbitals. Transport of such localized carriers known also as small polarons, is

done by a hopping mechanism accompanied by the surrounding lattice distortion. Due to this

transport mechanism results in carrier mobility values much lower than that for the band conduction

in the range of 1 – 0.1 cm2/Vs. These attributes result in a stronger coupling of the three factors

(electrical conductivity, thermal conductivity and Seebeck coefficient). The mean free path of phonons

in oxides ranges between 0.2 – 2nm and thus, for effective phonon scattering to be achieved,

patterning and nano features induced in the materials should be of comparative length scales9.

Initially oxides where believed to be inadequate as thermoelectric materials due to low

motility values. However they have other inherent properties that render them a good candidate for

thermoelectric material research. They are non-toxic and environmentally friendly while attributes

such as large thermal and chemical stability allow for their application over a wide temperature

gradient in air environment. Due to the large temperate gradient tolerance not only a high Carnot

efficiency and can be achieved but also, nonlinear, nonlocal TE effects (such as the benedicks effect)10

may me induced, playing also a positive role towards the thermoelectric potential.

Moreover oxides can be chemically adaptive and structurally complex which makes them

suitable for nanoscale material engineering both in aspects of composition and structure. Finally they

can be found in abundancy in nature thus radically decreasing the cost of raw material. Although the

evaluated zT values of the researched oxides are still lower that of state of the art thermoelectric

materials, the positive factors mentioned previously indicate that research on oxides from the

thermoelectric point of view is certainly worthwhile2.

Cobaltite-oxides and more specifically NaxCoO2 and Ca3Co4O9 as well as variations with

different dopant elements11-13 have been reported to yield significant thermoelectric properties

reaching zT values that exceed unity2. If those performance values can be confirmed through further

testing and reproducibility these two materials can be probable candidates for the next generation

thermoelectric applications.

3.2. NaxCoO2

One of the most important steps towards the recognition of oxides as promising

thermoelectric materials, was done by Tersaki et al.14 in 1997. Measurements performed on single-

crystal Sodium cobaltite oxide NaCo2O4 or 2x(Na0.5CoO2) metallic transition-metal, pronounced the

low resistivity values and the high in-plane thermoelectric power at 300k, contrary to what was

expected at that time. NaxCoO2 is a layered complex metal oxide compound consisting of two CdI2

type CoO2 layers holding the sodium atoms in between. The CoO2 layers directly contribute to the

electrical conductivity and the large Seebeck coefficient while the Na atoms act as a structural unit,

influencing the electronic concentration in CoO2 layers and decreasing the thermal conductivity along

12

Figure 9: Crystal structures of NaxCoO2 a) two layer structure (related to γ-phase) b) three layer structure related to α- and β-phases. The red spheres represent the O atoms the blue spheres the Co atoms and the yellow spheres represent the Na atoms. Adapted from 17.

the stacking direction c-axis15. Figure 8, shows the computed lattice thermal conductivity of NaxCoO2

In the in plane and out of plane direction.

Figure 8: computed Lattice (phonon) thermal conductivity of NaxCoO2 in-plane (black κ‖ ) and out of plane (blue κz )16.

Different stoichiometry of the Na atoms which can vary from 0.25<x<0.9, leads to significant

variation of the crystal structure and hence of the related properties17. Due to this reason the lattice

parameters of the NaxCoO2 need to be evaluated for every growth technique and substrate. Based on

comparison between the measured properties for PLD growth on LSAT substrate and literature

reported values, the stoichiometry of the films used in this work, has been estimated to be Na0.67CoO2

with very limited variation over increasing the thickness of the film18-20. The lattice parameters

corresponding to this stoichiometry are a=2.82 Å, c=11.05 Å 17. A schematic representation of the

crystal structure of both structural phases is shown in figure 9.

The NaxCoO2 single crystals show resistivity values of 200 mOhm*cm and a large room

temperature Seebeck coefficient of 100 mV/K. The zT exceeds unity above 800 K for bulk crystal.

However NaxCoO2 is challenging as a material since it decomposes into insulating Co(OH)2 in high

humidity environment because the Na+ ions dissolve easily in H2O.

13

3.3. Ca3Co4O9

Similarly to NaxCoO2, Calcium cobaltite oxide Ca3Co4O9 has a complex crystal structure

consisting of alternating stacks of triple rock salt-type Ca2CoO3 layers and single CdI2 -type hexagonal

CoO2 layers. The lattice parameters a=4.834 Å, c=10.835 Å and β (tilt) =98.14o, are the same for

both types of layers which have monoclinic symmetry and stack on top of each other along the c-axis.

However the b-parameter differs in the two layers creating a mismatched structure, with b1=2.824 Å

for the CoO2 layers and b2= 4.558 Å for the Ca2CoO3 layers. The CoO2 layers are conductive and the

Ca2CoO3 block is insulating and considered as a charge reservoir12,21,22. The structural characteristics

of the Ca3Co4O9 are shown in figure 10.

Literature reports indicate the high thermal and chemical stability of Ca3Co4O9 up to 1000+

degrees in air as well as high room temperature Seebeck coefficient of 120-140μV/K and metallic-

like electronic conductivity23. Ca3Co4O9 is reported to grow more crystalline at higher deposition

temperatures 750oC and low deposition rates. At the initial stages of growth the material forms a less

crystalline buffer layer of a few nanometers which is believed to be connected to the high

thermoelectric performance 24,25 .

Due to the structural complexity, Ca3Co4O9 shows reduced thermal conductivity especially in

the out of plane direction. This observation can be explained by the enhanced interlayer scattering of

the phonons (which are calculated to contribute to 90% of the thermal conductance). These features

render the material very promising for further investigation as a thermoelectric element. The

computed phonon thermal conductivity of Ca3Co4O9 in both out of plane and in plane direction26 is

shown in figure 11.

Figure 10: Crystal structure of Ca3Co4O9 unit cell along [010] and [100] axes adapted from25

14

3.4. Thin films

Thin films grown from these materials on different substrates, generally show differences on

their atomic structure compared to bulk samples either single crystal or polycrystalline. Due to the

growth mechanics of PLD as discussed earlier, crystallinity can be maintained, increased or decreased,

grain formation and sizes can vary, defects can be formed or induced and stain and relaxation

parameters play a significant role. Epitaxially grown thin films can form a strongly oriented structure

which can preserve the intrinsic electrical properties of the bulk crystals. At the same time, the small

dimensions can induce quantum confinement effects and further reduction of the thermal

conductivity due to lattice mismatches, grain size variation and scattering of phonons at surface and

interface boundaries. Boundaries and defects engineering in conventional approaches are reported to

play a significant role in enhancing the zT factor.

The substrate material is also very important in order for the epitaxially grown thin films to

maintain a high quality single crystal structure, planar coherency and orientation. This way electronic

conduction can stay close to the bulk value while thermal conductivity is limited due to phonon

confinement. The crystal structure of (LaAlO3)0.3(Sr2AlTaO3)0.7 (LSAT) used as the substrate in this

work is cubic with a-lattice parameter of 3.868 Å. The in-plane lattice mismatch between the cobaltite

oxides and the LSAT is expected to induce some strain effects in the growth of the film. The possible

in-plane ordering of NaxCoO2 on the LSAT substrate is a 12-fold symmetry at 15o orientation

mismatch and is shown in figure 12.

Figure 12: Representation of the in-plane ordering of NCO (red mesh) on LSAT (black mesh). Adapted from27

Figure 11: Anisotropy of total thermal conductivity in [Ca2CoO3]0.60CoO2 as a function of temperature. Out-of-plane thermal conductivity is far lower than in-plane thermal conductivity26.

15

The structure of Ca3Co4O9 and NaxCoO2 misfit-layers as well as the stacking faults observed

in the first layers of grown thin films work as an inherent barriers which impose a diminishing effect

on the thermal conductivity attributed to phonon-phonon interactions and phonon scattering.

3.5. Epitaxial cobaltite oxide superlattices

Superlattices are periodic, nanocomposite structures consisting of alternating material layers

with thicknesses as small as a few nanometers. Such structures have been commonly designed in

order to control electron transport. Phonon-related thermal transport, however, can also be affected

through the proper engineering of the superlattice unit cell. This approach is in line with the Phonon

Glass Electron Crystal theory28,29 which suggests that an ideal thermoelectric material should behave

as a conductive crystal for electrons but block phonons like an amorphous glass. This way superlattices

consisting of materials with similar lattice parameters, have the potential to show high zT values by

reducing the thermal conductivity due to interface density and epitaxial strain effects, while

maintaining good electron transport properties due to the crystalline similarities of the materials.

The superlattice approach on oxides with a focus on enhancing the thermoelectric potential

by reducing the thermal conductivity has been studied for a variety of materials in literature30,31

including Ca3Co4O9 and NaxCoO2 and other oxides27,32,33. For most of the samples effective reduction

of the thermal conductivity is reported with a clarified specification that the out-of plane parameter

is reduced significantly more than the in-plane parameter. This outcome is attributed to the

contribution of the interfaces to the phonon scattering processes. There is a lot of research on how

the period thickness and density influence this effect, however a conclusive model is not yet defined

mainly because of the complexity of the phonon scattering mechanics that involve transitions from

coherent to incoherent transport modes. It is clear however that for effective thermal conductivity

reduction the layer thickness has to correspond to the phonon mean free path related to the materials

that constitute the superlattice.

In this work the superlattice approach is revisited for the Ca3Co4O9 and NaxCoO2 oxides,

focusing mainly on structural and electronic properties of the superlattice samples. Unfortunately

thermal conductivity measurements were not possible during the course of this work, however if the

samples are proven retain good levels of the measured properties, such measurements may be

conducted as future work. Figure 13, shows an illustration of a superlattice film.

Figure 13: Illustration of a superlattice thin film with layers of Ca3Co4O9 and NaxCoO2 oxides grown on a crystal substrate. Optimally the electronic transport is minimally affected along the oxide layers while the phonons transport is hindered due to scattering at the interfaces formed between the two oxides. Adapted from27.

16

4. Fabrication and characterization methods

In this chapter all the sample fabrication, preparation and characterization methods that were

used throughout the course of this work are listed together with a short explanation of their working

principle. All the sample thin films have been deposited by the Pulsed Laser Deposition technique.

Atomic Force Microscopy (AFM) and X-Ray Diffraction (XRD) are then used to observe the growth

quality. Subsequently the samples were prepared for the electronic measurements via an etching and

sputtering procedure and lastly resistivity and Seebeck measurements were conducted.

4.1. Pulsed lased deposition

Pulsed Laser Deposition (PLD), is a proven method for growing thin films at variable deposition

rates while the stoichiometry of the target material is preserved at a good level. In its common format,

figure 14, it consists of a high vacuum chamber which contains the target holder with the target

material(s) and the substrate(s), a heater stage where the substrate is placed, a load lock with a

loading stick to handle the target stage and the heater and a Rheed electron gun with a phosphor

screen. The atmosphere inside the chamber can be varied with different gasses. Commonly for the

deposition of oxide films, O2 is used as background atmosphere to fully oxygenate the deposited

samples.

Figure 14: Schematic illustration of a typical PLD setup and the relevant components. Adapted from34

A high energy UV pulsating laser is directed via optical components into the chamber and

focused on the desired target. Mirrors are used for alignment of the laser beam while focus lenses and

metallic masks are utilized to control the energy density, the homogeneity and the laser spot size and

fluency on the material target. Across the target stands the heater stage where the substrate material

is held. The distance between the target material and the heater stage as well as the alignment (x, z,

θ) can be adjusted. The heater stage regulates the temperature of the substrate.

17

For every laser pulse on the target, energy is first converted to electronic excitation and then

into thermal, chemical and mechanical energy resulting in evaporation, ablation, plasma formation

and even exfoliation. The outcome is that atomic species from the target material are ablated in the

form of a plasma plume and directed towards the substrate via a pressure gradient. When reaching

the substrate which is usually a single crystal, they crystalize in accordance to its crystal lattice and the

energy conditions which are tuned by changing the temperature of the heater in combination with

the pressure/gas ambiance in the chamber. In between the pulses or pulse bursts the atomic species

have the time to thermally diffuse on the “hot” substrate surface and find the energetically optimal

positions to crystalize, until the next wave of atomic species arrives. This way, layer by layer, thin films

are grown epitaxially on the substrate.

The target holder can contain up to 5 different material targets which can be independently

used for a single film deposition. Optimally the angle at which the laser hits the target is kept at 450.

During the deposition the target is moving with respect to the beam, in direction parallel to the heater

surface, so that a scan area is formed, this way preserving the homogeneity of the plasma plume while

limiting the depth of penetration into the target material.

4.2. Sample preparation and deposition

As mentioned earlier the samples were deposited using PLD. The growth parameters have

been calibrated according to the most recent working conditions of the PLD system. First the LSAT

substrate was chosen over LaAlO3 and Sapphire due to the higher room temperature Seebeck

coefficient values acquired previously from test deposited films while the crystallinity of the films is

still at high level27. The substrates are treated to an annealing procedure for 10 hours at 1050oC and

scanned with AFM to ensure cleanness and surface quality.

Before each deposition the PLD chamber is pumped down to high vacuum conditions ~10-7

mBar which is an indication of pure ambiance in the chamber. I.e. Absence of contamination species

that could interfere with the quality of the deposited film. The temperature of the heater is then risen

to 430oC and pressure is set at 0.4 mbar O2. These growth conditions are set for optimized growth of

the NaxCoO2 crystalline film19. Since it is the more volatile of the two elements deposited and it cannot

withstand the higher (optimally) deposition temperatures of 750oC for the CCO23. Oxygen

environment in used during the process to ensure that the final film will be a saturated oxide. The

laser repetition rate is kept at 1 Hz for the film deposition to guarantee enough time for the species in

the plasma to acquire good crystalline ordering on the substrate as well as for saturation of the oxygen

atoms.

After the deposition of the film, the sample is treated to an automated oxygen flush up to

1000 mbar. The cooldown procedure is then initiated and set to drop from 430oC to room

temperature at a rate of 10oC per minute. 1000mbar O2 pressure is maintained during the cooldown.

After the sample has reached a temperature <50oC, the deposition of the amorphous Al2O3 capping

layer is initiated. This layers is to prevent or at least limit the degradation over time of the volatile

NaxCoO2 layers. The capping layer thickness is always tuned to be at least equal or greater than

100nm.

18

4.3. Atomic Force Microscopy (AFM) characterization35

AFM is a non-intrusive characterization method suitable for investigating the surface

characteristics of a sample and was used to analyze the surface quality of the substrates before the

deposition as well as the surface of the films after the deposition. A sharp tip of atomic dimensions,

follows the topography of the surface and through intensity, phase, frequency variations, we can get

qualitative and quantitative information of the scanned area.

Bruker Dimension Icon AFM and a Multimode SPM have been used for all measurements in

this thesis. All measurements have been conducted at ambient temperature and ex-situ using tapping

mode (TM). In this mode the oscillating AFM tip with a frequency between 100 to 400 kHz, is brought

close to the surface (<10nm) where it “feels” a repelling force from the sample. Since the force

induced at the tip is kept stable, the tip is forced to move up and down following the variations in the

topography, through changing the vibrational amplitude. These variations are then translated to

different color contrast in the program, enabling visual illustration of the samples’ surface as well as

qualitative and quantitative analysis of the scanned area. An illustration of the working principle of a

typical AFM setup is given in figure 15.

Figure 15: Illustration and working principle of a typical AFM setup35.

Initially the AFM is utilized to ensure the quality of the substrates that were used for each film.

After the film (and the A2lO3 capping layer) is deposited, the AFM is used again first to evaluate the

surface roughness of the film. Low surface roughness of the amorphous capping layer is an indication

of better coverage and thus protection of the underlying film. Moreover the AFM images are used to

extract a rough grain size profile. This is possible since the amorphous capping layer is expected to

follow the pattern of the underlying crystalline film. Typical scans for observing the quality of the

substrate are shown in figure 16.

19

Observed LSAT steps are roughly 0.5nm high. The LSAT substrates were being treated to an

annealing procedure for 10 hours at 1050oC and in O2 atmosphere. Difference in the terraces on the

substrates are an indication of the different miscut angle over different batches of substrates.

4.4. X-Ray Diffraction (XRD) Characterization36,37

Three types of X-ray characterization measurements were performed on the samples, under

ambient temperature and humidity conditions.

1) A low angle 2θ/θ (0<2θ<8) optical reflectivity measurement to determine the thickness

and roughness of the grown films.

2) A wider range 2θ/θ (10<θ<110) out-of-plane diffraction scan to determine the film’s

crystal orientation with respect to the substrate as well as a more detailed 2θ/θ focused on

the (002) plane peak of the film (12<θ<20).

3) A rocking curve measurement (ω-scan) along the film’s most intense peak is also taken, to

indicate the preferred orientation of the film.

An illustration of the symmetrical-reflection θ/2θ measurement and the rocking curve

measurement principles is given in figure 17.

Figure 16: Typical AFM images of the LSAT 001 substrate used throughout this research. The miscut angle is higher for the substrate in the left scan, with more terraces exposed over similar surface area. From the line profile it is evident that the terrace step size is roughly 0.5 nm in both images.

20

Figure 17: Left: illustration of symmetrical-reflection θ/2θ measurement. Right: Rocking curve measurement. Adapted from36.

The incoming X-ray beam is initially aligned on the center of the sample (X, Y, and Z) positions

and then the X-ray source as well as the detector can be rotated around the sample. With this setup

crystallographic measurements can be performed as well as optical measurements, e.g. to determine

the thickness of the thin film.

For the most common measurement, symmetrical reflection, the diffracted X-rays from the

crystal lattice planes parallel to the sample surface are collected, providing crystallographic

information along the sample surface normal vector. The intensity and the sharpness, low Full Width

at Half Maximum (FWHM) of the peak are an indication of how well the film/superlattice has

grown/stacked along the c-axis. The information obtained by this method can be interpreted by

utilizing Bragg’s law:

𝑛𝜆 = 2𝑑𝑠𝑖𝑛𝜃

Where (λ) is the X-ray wavelength, (θ) the angle between the sample and the incoming X-ray beam,

n is the number of crystal planes and (d) is the distance between the crystal panes. According to this

formula, when the wavelength of the X-rays is equal to twice the distance between two crystal planes,

peaks appear in the 2θ/θ scan. Diffraction peaks appearing at fixed intervals indicate the periodic

structure of the film as well as that of the substrate.

For the low angle reflectivity measurements the 2θ angle is scanned while ω angle is kept

fixed at half of the scanning angle, ω=θ. The beam is aligned on the optical surface of the sample

(2θ=0.4o, ω=0.2o) to maximize the optical reflection intensity. The film is then scanned between

(0o<2θ<6o). The fringes that appear are a result of the difference between the refractive indices of

the thin film and the substrate, which is still optically accessible due to the films’ low thickness. The

thickness of the film can be derived either by direct analysis of two consecutive fringes, or by Fourier

analysis of the whole plot up to the point where (data) noise becomes too noticeable.

For the Rocking curve (RC) measurement the orientation axis, which has been previously

determined by the symmetrical reflection measurement, is taken into account and the 2θ-angle is

kept fixed. The diffraction intensities from the lattice planes along the favored direction are then

noted. Finally, the degree of alignment to this preferred orientation is estimated from the FWHM

profile of the rocking curve peak. Defects like dislocations, film curvature and grain mosaicity, create

disruptions in the perfect parallelism of atomic planes are observed as broadening of the rocking

curve. The center of the rocking curve is determined by the d-spacing of the peaks. Very sharp RC

peaks are observed only when the crystal is properly tilted so that the crystallographic direction is

parallel to the diffraction vectors however due to instrument broadening and the intrinsic width of

the crystal material a certain width will also be observed.

All XRD measurements presented in this thesis have been done on the PANalytical X-pert pro MRD

21

4.5. Reactive Ion Etching and deposition of gold contacts

After the film’s surface and crystalline growth have been analyzed, usually in this order, the

following step is to prepare the sample for the electronic measurements by depositing conductive

contacts on the four corners of the film. To achieve this, a custom made-cross like, metallic mask was

manually mounted on the film leaving only the four corners of the sample exposed. The sample is then

placed in the argon ion etcher and the exposed corners are subjected to etching time estimated

according to the film and capping layer thickness. The goal is to etch through the Al2O3 capping layer

and through (most of) the deposited film thus exposing enough active film layers attach with the

conductive contacts.

After etching is complete the sample is transferred in a Perkin Elmer sputtering machine for

the deposition of the contacts. The transfer time is kept under a minute to keep the sample exposure

to the humidity of the ambiance to a minimum. The material chosen for the contacts is gold due to its

chemical inertia and low resistance. During sputtering, argon ions are accelerated towards the surface

of metal target, when in contact atoms from the target are extracted and propelled due to a voltage

gradient, towards the sample thus forming a metal layer. The Argon ions are accelerated due to a bias

applied to the target. A constant Argon gas pressure of 10-2 mbar is maintained in the vacuum

chamber on their way to the target. This way the argon ions ionize even more argon atoms resulting

in a retaining a constant deposition rate metal atoms onto the sample34.

The sample’s exposed corners are firstly subjected to deposition of a thin titanium layer to act

as a precursor for the gold contacts which are deposited right after. The thickness of the gold contacts

is calculated to exceed the overall depth of the etched areas thus covering completely the exposed

corners and ensuring optimal contact with most of the layers in the film.

After the sputtering, the sample is taken out and the metallic mask is carefully removed. The

contact pads of the sample are connected to a PPMS puck using a soft wire in the wire bonder. Van

der Pauw resistivity measurements as well as transport measurements are then conducted. Ideally

during the same day of the resistivity measurements, the Seebeck coefficient is also measured in a

custom measurement setup.

4.6. Room temperature resistivity measurements

The resistivity of the samples was measured at room temperature and in ambient conditions

using the Van der Pauw (VdP) four point measurement method. An illustration of the common wiring

setup for conducting VdP resistivity measurements on the film is shown in figure 18. The resistance is

measured both in horizontal and vertical direction across the four contact points in the corners of the

tetragonal samples as seen in figure. Measurements where performed in - HP34401a - digital

multimeter and the values R1=VDC/iAB , R2= VBC/iAD are being noted.

22

Figure 18: Illustration of sample connectivity for Van der Pauw resistance measurement.

Then using the VdP formula,

𝑒−𝜋

𝑅1𝑅𝑆

+ 𝑒

−𝜋𝑅2𝑅𝑆

= 1

the sheet resistance (Rs) in (Ohm/□) is calculated using a matlab script. When the thickness of the

sample is known and it is relatively homogeneous the resistivity (ρ) can be calculated using ρ=RS*t

Where (t) is the total thickness of the conducting (active) layer of the film.

4.7. Room temperature Seebeck Coefficient measurements

Measurement of the room temperature Seebeck coefficient for the thin films is done in a

custom made setup. It consists of Three HP34401a digital multimeters to measure the three voltages

(V1, V2 and VSeebeck), two Peltier elements for temperature control and a Keithley 2400 source-

meter to power them. The sample (5x5mm) is placed on top and in between of two Peltier elements

separated by air gap of 2.9 mm. When placing the sample roughly in the middle, approximately (1.6x5

mm) of each side of the substrate bottom surface is in contact with the Peltier elements. In figure 19

an illustration of the measurement setup and the measuring points between the thermocouples are

being shown.

Figure 19: Illustration of the custom setup used to measure the room temperature Seebeck coefficient. The wire connection is also shown. The elements in the picture are not in scale.

The measurement cycle is run by a LabView program which is set to apply current at six

different stages in order to heat up one Peltier element while the other is cooled down and vice versa.

This way it generates a temperature gradient between the two Peltier elements and thus in the

23

sample. From the top side two thermocouples are in contact with the relevant contact points on the

sample. Two thermocouples consisting of Alumel and Chromel, measure the temperature gradient

that arises across the sample and translate it into an induced Seebeck voltage simultaneously. The

temperature gradients measured, vary usually between 1-3 degrees oC. A reference point for this

voltage is acquired by keeping a junction at 0oC. The formula used to estimate the Seebeck coefficient

of the sample is:

𝑉𝑆𝑒𝑒𝑏𝑒𝑐𝑘 = − ∫ 𝑆𝑎𝑙𝑢𝑚𝑒𝑙𝑑𝑇 −𝑇1

𝑇𝑟𝑒𝑓

∫ 𝑆𝑠𝑎𝑚𝑝𝑙𝑒𝑑𝑇 −𝑇2

𝑇1

∫ 𝑆𝑎𝑙𝑢𝑚𝑒𝑙𝑑𝑇𝑇𝑟𝑒𝑓

𝑇2

= ∫ 𝑆𝑎𝑙𝑢𝑚𝑒𝑙𝑑𝑇 −𝑇2

𝑇1

∫ 𝑆𝑠𝑎𝑚𝑝𝑙𝑒𝑑𝑇𝑇1

𝑇2

Where (T1) and (T2) are the temperatures measured at the contacts of the sample and (Tref) is the

temperature at the junction set at 0oC. The induced Seebeck voltage is then plotted against the

temperature gradient and from a linear fitting the Seebeck coefficient is estimated. The measurement

is performed along all the 4 sides of the sample and across the diagonals and the average Seebeck

value is finally calculated. The contribution of the Chromel wires is canceled out, however since the

Alumel material of the thermocouples also generate a Seebeck voltage due to the temperature

difference, to determine the actual temperature difference between the two thermocouples out of

the voltage difference between V1 and V2 a value of (+18.5 μV/K) has to be subtracted from every

measurement.

With this setup both ΔΤ and ΔV are done simultaneously and at the same point, enhancing our

control over the measurement conditions and without the need to wait for some stabilization time

that the system is in thermal equilibrium. The error value has been determined by measuring the

Seebeck coefficient of a Bi2Te3 calibration sample. Comparison of the Seebeck values measured here,

with values from the literature resulted in 5% difference and this percentage is used as the error in

the measurements.

4.8. Physical properties measurement setup (PPMS) , mobility and carrier concentration

The measurement principle of the PPMS is the Hall Effect. Hall Effect measurement can be

primarily utilized to determine the Hall voltage (VH). From that point on , other Important parameters

such as carrier mobility, carrier concentration (n), Hall coefficient (RH), resistivity, magneto-resistance

(R), and the conductivity type (n or p) can all be derived from the Hall voltage measurement.

For a more detailed measurement of the Hall resistance and its response to variation in

magnetic field the film is placed in the PPMS under helium atmosphere and in 300 oK. The sample is

wired in Hall configuration. A magnetic sweep is then run from 0 to -9 Tesla to +9 Tesla and back to

0 Tesla, and a plotting of over the Hall resistance over magnetic field B is then extracted. From the

slope of the curve we calculate the carrier density and in turn the carrier mobility. Since the samples

are very sensitive to thermal effects the fitting was done only for the last phase of the swipe (0T to

9T and back 0T) and a resistivity correction due to thermal drift over time, has been carried out.

24

The related measured and calculated factors are:

The Hall voltage, given in relation to the Current (I), the magnetic field (B) the carrier density (n), the

carrier charge (q) and the thickness of the film (t):

𝑉𝐻 = −𝐼𝐵

𝑛𝑞𝑡

The Hall resistance RH

𝑅𝐻 =𝑉𝐻

𝐼

The relation between the Hall resistance and the slope

𝑅𝐻 = −𝐵 (1

𝑛𝑞𝑡) = 𝐵 ∗ (𝑠𝑙𝑜𝑝𝑒)

Estimation of the carrier density

𝑛 =1

𝑠𝑙𝑜𝑝𝑒 ∗ 𝑞𝑡

Estimation of mobility (μ) from carrier density

𝜇 =1

𝑞𝑛𝑅𝑠

Where (Rs) is the sheet resistance.

With these equations in mind we estimate the carrier density and mobility for the superlattice samples

related to the main part of this work.

25

5. Results and discussion

5.1. Work outline

Plan: The plan for this research was to utilize an already optimized thin film growth method

of stacked Ca3Co4O6 and NaxCoO2 layers and to further investigate the possible enhancement of the

thermoelectric qualities the deposited films. The background idea behind the combination of these

two materials into an epitaxialy grown thin film superlattice, was firstly the structural similarity

(compatibility) and secondly the fact that one could improve on the properties which the other was

lacking, as discussed in section 2. More specifically thin films of Ca3Co4O6 show a high Seebeck

coefficient but also relatively high resistivity values in the order of 10-20 mOhm*cm while thin films

of NaxCoO2 show resistivity values in the order of 5-10 mOhm*cm while maintaining a Seebeck value

of roughly 70% compared to Ca3Co4O6 ,considering films of same thickness. By maintaining a constant

film thickness while changing the number of interlayers (periods) the aim was to investigate whether

Seebeck and Resistivity values are maintained within good limits and also to examine in future

experiments, the effect on the thermal conductivity of the films.

Calibration: In the beginning of this research the PLD system had been relocated and the laser

optical path was significantly elongated. Therefore there was need for a significant part of the

procedure to be re-evaluated and re-optimized. For this reason the first quarter of the timeframe of

the research was utilized to calibrate the growth of all the materials involved under the new conditions

and evaluate the deposited reference films. Due to the limited timeframe and since I had to proceed

to the main research questions, only partial re-optimization of the process was achieved. This involves

the thickness homogeneity of the deposited films due to alignment issues and the effectiveness of the

deposited Al2O3 amorphous capping layer. All of the thin films which contained NaxCoO2 layers with

the capping layer presented improved stability over time as compared to those without, however not

absolute.

Main aim: For the main research question, superlattice thin films of Ca3Co4O6 : NaxCoO2

combination where grown up to 140nm total thickness using optimized parameters for NaxCoO2. The

films where grown on insulating 5x5mm2 LSAT 001 substrates, since previous reports have shown

sufficient crystallinity and higher Seebeck values compared to other substrates. The individual layers

thickness was varied between approximately 1nm and 17.5nm resulting in superlattice films with

periods of (4, 7, 14, 28, 70) inter-layers. Growth quality indicators such as crystallinity, surface

roughness and grain size have been monitored and electronic conductivity, and the Seebeck

coefficient where measured. Electron carrier densities and nobilities were also calculated.

For all the grown samples, characterization and measurements where conducted in ambient

humidity and temperature conditions. From here on in the main text shortcuts of the material names

are going to be used. Example: ALO/(CCO:NCO)n/LSAT stands for n-periods of Ca3Co4O6 : NaxCoO2

on LSAT substrate covered with an amorphous layer of Al3O2.

26

5.2. Calibration samples

The growth rate the materials involved in this work had to be calibrated three times during

the course of this research. Initially in the very beginning due to the relocation of the PLD system later

on once more due to the modification of the laser alignment procedure and lastly due to the change

of the laser intensity meter equipment. Initially three calibration samples where grown on Al2O3

substrates because films grown on this substrate before have shown reduced surface roughness which

is a parameter that can influence the accuracy of the thickness evaluation from the XRD-low angle

reflectivity measurement.

The first two samples where CCO/ALO deposited at two different pressure and temperature

combinations in order to understand the difference in structure and crystallinity as well as to measure

the growth rate. Seebeck coefficient and resistivity measurements were also performed.

Figure 20: AFM scans (3x3 μm) of 85nm CCO film grown at 750 oC (left) and 27.5nm CCO film grown at 430οC (right), on ALO substrate .

Figure 21: Diffraction pattern on the (002) plane for both films grown at different temperatures.

It has to be noted that the films are of different thicknesses, however from the AFM images

(figure 20) a clear difference in the grain growth between the two can be observed and from the

Diffraction pattern shown in figure 21, the (002) plane peak for the CCO can only be seen for the high

temperature growth. This was also the only observable peak for a plain CCO film throughout this

research and it is going to be used as a reference for the general preferred d-spacing for CCO.

27

Figure 22: Typical low angle X-ray reflectivity scan in order to measure the film thickness and evaluate the growth rate.

The film thicknesses have been measured by X-ray low angle reflectivity scans like the one

shown in figure 22. For the CCO the growth rates were evaluated to 54 pulses/nm at 750 oC, 0.01

mbar O2 pressure and 73 pulses/nm at 430 oC, 0.4 mbar O2 pressure. For NCO the growth rate was

67 pulses/nm at 430 oC, 0.4 mbar O2 pressure. From here after and up to the point where new

calibration was needed, the film samples were deposited with these growth rates as reference and all

of the films where grown at 430 oC, 0.4 mbar O2 pressure due to the nature of NCO as earlier

discussed.

Figure 23: X-ray diffraction pattern of both films on ALO. The inset image shows the different (002)-plane 2θ peak value for CCO and ALO which is indicative of their d-spacing difference.

The different value for the 2θ diffraction peak of the (002) plane for CCO and NCO crystalline

film is shown in figure 23. From Bragg’s law as explained in section 4.4, we calculate the d-spacing for

the two crystalline films. For NCO (θ≈8.035±0,05ο), d≈11.02 Å and for CCO (θ≈8.26±0,05ο),

d≈10.72 Å which are consistent with values reported in literature17,18,38.

After the new railway and alignment of the laser the growth rates of CCO, NCO were evaluated

again this time on LSAT Substrate. From the XRD-reflectivity profiles the growth rate was recalculated

at 72 pulses/nm for CCO and 64 pulses/nm for NCO. Finally after the implementation of the new

measuring device for the laser intensity, only the growth rate of CCO was re-measured for a third time

and the was confirmed at 72 pulses / nm.

28

After consultation and since precise control of the layer growth was not the focus of this

research, this value was kept as the growth rate for both materials. The relative growth ratio would

then be roughly 1:1.125 (CCO: NCO) resulting in slightly more NCO being deposited in every period

for the superlattice structures.

Calibration of ALO: For the growth rate amorphous layer of ALO a piece of silicon wafer was

used partially treated with a photoresist mask. After the deposition, the photoresist layer was

removed in acetone along with the ALO deposited on top. Therefore ALO remained only in the

exposed part of the silicon. AFM height profile was taken subsequently to measure the height of

deposited layer and later calculate the rate of growth. The AFM image and the growth profile are

shown in figure 24.

Figure 24: Left. AFM image of a step profile of ALO on Silicon. The lower area has been normalized by a 3-point leveling technique. Right. Line scan of the height profile.

From the step profile the height difference was measured at 145 nm and a value of ~76

pulses/nm for the growth rate was extracted. The small kink at the step-edge is a result of more

material species getting accumulated at the interface between the substrate and the edge of the

photoresist layer and it is commonly observed.

After the new laser alignment conditions and the replacement of the intensity meter however,

the growth rate of the ALO capping layer had to be re-evaluated. An LSAT substrate was used and a

droplet of photoresist approximately 3mm in diameter was applied on the center of the substrate.

After the ALO deposition the photoresist was thoroughly washed away in an acetone bath. For the

7200 pulses/5Hz used in this deposition, it was estimated by the previous growth rate calibration to

result in an ALO layer of 103nm. The sample was subsequently scanned in the AFM and height profiles

were extracted in the four (N,W,S,E) quarter-circles of the approximately circular area. The four

measurements varied between 127nm and 168nm and an average height profile was estimated

149nm. From this new value the deposition rate was estimated at 48 pulses/nm.

The significant inconsistency of the four line scans may indicate incorrect alignment of the

substrate towards the material target. Overall however the growth rate was found to be considerably

higher than previously calculated, which may be due to a more accurate calibration of the laser

fluency. There is a concern regarding the validity of this growth rate measurement, however due to

shortage of time no further optimization could be made. This new value was applied for all the later

samples that followed and commonly depositions of 5500 pulses were used resulting in an

approximately 120nm thick capping layer.

29

5.3. Ion etching profile calibration.

As a final calibration step, the etch rate of the ALO plus the material needed to be estimated.

Two previously deposited samples of ALO/NCO/LSAT were used. For both samples a layer of

photoresist mask was placed and fixed on top of each sample leaving only a certain area in the center

of the surface exposed. The samples were etched in two separate processes. One sample was etched

for 6 mins while the second was etched for 6+6 minutes, with a pause of 5 mins in between to prevent

sample overheating, using the same conditions and parameters. Both samples were later scanned in

the AFM. For the 6 min etched sample the height profile was c.a 68nm while for the 6+6 mins etching

the profile was c.a 145 nm. The corresponding etch rates were calculated at 11.3nm/min and 12.1

nm/min respectively, suggesting a probable increase of the etch rate of NCO after the ALO layer has

been breached. Unfortunately due to the limited time frame we did not evaluate separately the etch

rates for all the relative materials. The 12.1nm/min etch rate was applied for the calculations

regarding the contact deposition procedure. Since this procedure was not accurately calibrated for

most of the samples etching times between 13-14.30 mins were used

5.4. CCO:NCO Superlattice samples: Constant layer thickness, variable total thickness.

A total of three CCO:NCO superlattice samples were deposited keeping the individual layer

thickness fixed at 10nm and varying the total thickness of the active film from 60 to 140nm, to

investigate how the relative properties change. One sample of 50nm total thickness, but ending with

a layer of NCO thus breaking the periodicity was also prepared and investigated. AFM, XRD, Seebeck

and resistivity measurements have been conducted. Seebeck and resistivity values regarding a sample

of 200nm total thickness with related characteristics are also reported in this section for comparison,

however it has to be noted that this sample was deposited under the previous operating conditions

of the PLD system.

AFM characterization

The surface characteristics of the three samples were analyzed with the AFM. Scans of

3x3μm2 from the 50nm ALO/NCO:(CCO:NCO)2/LSAT sample and a 60nm ALO/(CCO:NCO)3/LSAT

are shown in figure 25. The surface roughness and grain analysis statistics are noted in table 1. The

AFM data for the 140nm thick sample are reported later in chapter 5.6.

Figure 25: 3x3μm2 scans of the 60nm (CCO:NCO)3 (left) and the aperiodic 50nm NCO:(CCO:NCO)2 (right) grown on LSAT. Color scale is the same for both images.

30

Table 1: Roughness and grain statistics for the 60nm and 50nm samples.

The difference between the two samples is one CCO layer of approximately 10nm. An

interesting observation here is that the grain size and shape vary significantly between the two films.

According to the FFT analysis, smaller and more homogeneous grains have formed for the 50nm

sample ending with NCO layer while for 60nm which has an extra 10nm CCO layer the estimated

grains have double the size and show broader inhomogeneity in the in-plane dimensions. A smooth

surface is maintained for both films with roughness <7%, the roughness range however is broader

for the 60nm film.

XRD analysis

The four samples where scanned in the XRD. As explained previously the scans are aligned

with respect to the substrate peak, which is single crystal LSAT (001). The error in alignment was

roughly evaluated at ±0.05o taking into account device accuracy and deviations between the

measurements. Figure 26 shows cumulatively scans over the (002) plane peak of the films.

Figure 26: Scan on the (002) plane peak for samples with variable total thickness. Higher intensity is consistent with higher thickness and also indicator of good crystallographic plane coherence.

From the intensity of the observed (002) peak for each sample, increased intensity with

increasing the sample thickness is evident and indicative that planar coherence is maintained as the

thickness of the superlattice increases. The peak position is maintained at roughly (θ≈8.185o±0.05o),

and the corresponding d-spacing value is d≈10.82Å which is in between the values measured for the

individual NCO, CCO films. The FWHM profile of all the peaks is 0.31±0.02 degrees.

Total thickness (nm)

surface roughness (Rms) in (nm)

Roughness range

(Rms) in (nm) Grain size (nm)

Grains min and max

dimension (nm)

60 3.3 ±1.2 285 135-436

50 3.5 ±0.2 144 114-174

31

Figure 27: Rocking curve of 60nm (CCO:NCO) 10:10 period superlattice film, on the 002 plane peak.

An interesting observation was also done for the omega rocking curve scan of the 60nm

sample shown in figure 27 where individual peaks of both materials are still observable. Peak positions

at 7.985o±0.05o and 8.285o±0.05o are noted, which closely correspond to the separate NCO and CCO

peaks previously observed on ALO substrate. The d-spacing matching these values are 11.09Å and

10.69Å respectively. LSAT has a lattice parameter of 0.387nm which is smaller than the 0.476nm

for ALO. This can explain the small variation from the positions of the (002) peaks observed earlier

on ALO substrate. Since the crystal structure of NCO and CCO will be more constricted in the a, b-axes

on the LSAT substrate in comparison to ALO and this will force a stretch in the c-axis. At higher

sample thicknesses these two peaks seem to merge into one at an intermediate position, probably

Indicating that the superlattice film is obtaining a d-spacing at relaxed position between the two.

Seebeck and resistivity analysis

After the surface and crystalline analysis, the samples were prepared for the electronic

measurements. A summary of the Seebeck and resistivity values is given in table 2. In Figures 28 and

29 the Seebeck, the Sheet resistivity and the combined power factor are plotted against the sample

thickness.

Table 2: Summary of the Seebeck and resistivity values. A value for the room temperature power factor is also given. Each layer is approximately 10nm thick.

Sample Thickness

(nm)

Sheet resistivity

(mOhm*cm)

Seebeck coefficient

(μV/K) Power factor RT

(W/m*K2) NCO:(CCO:NCO)x2 50 9.04 109.3 13.2151Ε-05

(CCO:NCO)x3 60 11.35 121.9 13.0922Ε-05 (CCO:NCO)x5 100 9.78 80.4 6.6096E-05 (CCO:NCO)x7 140 7.85 82.6 8.6914E-05

(CCO:NCO)x1027 200 6.2 95 14.5565Ε-05

32

Figure 28: Seebeck and sheet resistance over total thickness (constant period thickness 10nm/10nm). The data points not connected to the lines belong to the 50 nm sample which ends with an NCO layer and is (non-periodic). Error values are 7% on the resistivity and 5% on the Seebeck. The lines are guides to the eye. The data points for the 200 nm sample were taken from literature27.

Figure 29: Calculated power factor plotted against the sample thickness. Error bars are a product of error propagation from the two factors S, R.

From the values obtained by these measurements a clear trend was evident only for the

resistivity which is consistently decreasing with increasing thickness. This behavior is expected since

use of more material would increase the number of active electronic carriers. Furthermore the sample

of 50nm (2x 10:10 NCO/CCO +10nm NCO) has significantly lower resistivity of 9.04 mOhm*cm

compared to the 60nm sample which was measured to 11.35 mOhm*cm, while maintaining high

Seebeck value. The lower resistivity value can be anticipated since the extra NCO layer increases the

ratio of NCO: CCO to 3:2 (in 10nm layers) making the low resistivity material dominant.

For the Seebeck it can be noted that the values obtained in this work for the superlattice

samples are very close to an average between single NCO and CCO films of same total thickness,

deposited on LSAT as reported in literature*27,34. By reducing the overall film thickness the Seebeck is

increasing at least up to the 60nm thick film which was the minimum thickness periodic superlattice

film measured in this work. The Seebeck increase can be attributed to size effects and confinement

which enlarges the spacing between the energies at which the conduction levels lie and the chemical

potential.

*Comparison made between the single NCO film and CCO film values reported in Ch.4: P71, Ch.5: P104 of ref 27 and Ch.5:

P41 of ref34.

0

20

40

60

80

100

120

140

4

6

8

10

12

14

16

18

20

40 60 80 100 120 140 160 180 200 Seeb

eck

co

effi

cien

t (μ

V/K

)

Shee

t re

sist

ivit

y

(mO

hm

*cm

)

Sample thickness (nm)

0.0E+00

2.0E-05

4.0E-05

6.0E-05

8.0E-05

1.0E-04

1.2E-04

1.4E-04

1.6E-04

1.8E-04

40 60 80 100 120 140 160 180 200

Po

wer

fac

tor

(W

/m*K

2)

Sample thickness (nm)

33

5.5. Single films 70 nm

To inspect whether the 140nm thick (CCO:NCO) superlattice film was performing as an

average between the combined stoichiometry of 70nm NCO + 70nm CCO. Two more reference

samples were deposited. The active material thickness of these samples was 70nm CCO on 10nm of

NCO, and 70nm NCO respectively. Both samples were deposited on LSAT substrates. The 10nm

buffer-transition layer for the CCO film was to induce better crystallization of the CCO at the low

deposition temperature as explained earlier in the thesis. Capping layer was only used on the pure

NCO film. AFM, XRD, Seebeck and resistivity measurements have been performed.

AFM Characterization

Both films where scanned with the AFM Images of 3x3μm scans for both films are also

presented, together with line-profile taken to indicate the height profile of the individual grains in

figure 30 and a summary of the calculated values is given in table 3.

Figure 30: Left. CCO:NCO 70:10nm. Right. NCO 70nm. Both scans are 3x3μm2 wide. The profiles below the images correspond to the line scans.

Table 3: Grain and roughness statistics.

Sample reference

Total thickness

(nm)

surface roughness

(Rms) in (nm)

Roughness

error

(Rms) in (nm)

Grains min and

max dimension

(nm)

Grain

size

(nm)

CCO:NCO 70:10 14.0 ±1.5 72-91 82 NCO 70 16.1 ±1.4 44-93 68

34

From the AFM scan on the (CCO:NCO 70:10 nm) sample it seems that the CCO layer grows

on top of the NCO layer, forming pillar like structures separated from each other rather than forming

a coherent film layer. The more uniform grains of NCO buffer layer are visible in between these

structures. For the 70nm NCO sample a more uniform film layer can be observed with less spacing in

between the grains. The surface roughness of both films is quite high at 17.5% and 23% of the total

thickness respectively. The average grain size from the FFT analysis is calculated at 82 and 68 nm and

with a range of dimensions below 100nm for both samples. These grain sizes is significantly lower

compared to the superlattice samples.

XRD Analysis

For both samples, θ/2θ diffraction and rocking curve scans where performed along the

diffraction peak of the (002) plain. The different d-spacing for NCO and CCO is once more noticeable

in consistency with our previous observations. Moreover from the rocking curve scan the two distinct

peaks for the CCO:NCO 70:10nm sample again suggest that two planar orientations are maintained,

likely corresponding to the substrate NCO buffer layer and to the CCO layer. The high intensity of both

peaks however may suggest that the two different planar orientations may belong to the “CCO pillars”

Figure 31: Left. 2θ diffraction pattern on the (002) plane peak for both samples. Right. Rocking curve on the (002) plane peak for both samples. Color code is the same for both figures.

There is significant difference in the relative intensity between the two peaks despite their

small difference in thickness. The reason for this may be the pillar-like formation observed also in the

AFM scans. The height of these structures may exceed the average film thickness which was calibrated

in the deposition parameters. Therefore more (002) planes are exposed to the X-ray beam, drastically

enhancing the number of counts in the diffraction pattern.

Seebeck and Resistivity measurements

The room temperature sheet resistivity was measured at 10.92±0.76 mOhm*cm and

3.37±0.24 mOhm*cm and the Seebeck at 90.3±4.5 μV/K and 21.1±1.05 μV/K for the CCO and NCO

samples respectively. Averaging these two numbers we get a resistivity of 7.14±0.5 mOhm*cm and

a Seebeck of 55.7±2.8 μV/K.

35

5.6. CCO:NCO Superlattice samples: Constant total thickness, variable number of periods.

Varying the amount of periods and interlayers in a sample of constant thickness is expected

to influence the structural and electronic properties of the cobaltite superlattices.

In this chapter a total of five Superlattice samples of constant total thickness 140nm and

variable period density of NCO/CCO (4, 7, 14, 28, and 70) layers were deposited and studied. The

choice of thickness at 140nm was made with the rationale that it has to be enough for a sufficient

number of periods to fit inside, while it retains good properties as investigated previously in chapter

2.1. AFM, XRD, Seebeck, resistivity and PPMS analysis has been performed on all of the samples. Since

complete stability was not achieved by the ALO capping layer, priority was given towards the Seebeck

and resistivity measurements which were usually undertaken one or two days after the deposition.

What follows is an analytical overview of these samples as well as comparison with the corresponding

properties of the individual reference films.

Surface and grain analysis

All the samples where scanned with the AFM. Larger scan areas of 10x10 μm2 or 20x20 μm2

were used for the statistical analysis regarding the roughness and grain size estimation while scans of

3x3 μm2 and smaller were used for individual grain observations and further qualitative analysis and

discussion. It has to be reminded that the surface layer of all the samples is amorphous ALO roughly

thicker than 100 nm and hence the scan is not performed directly on top of the materials of interest.

Therefore the AFM characterization can only be accounted for these specific samples and only for

rough estimations. The numbers are far from being absolute but they can be used to show a relative

trend in the roughness and grain sizes. Images of 3x3 μm2 scans of the samples are shown as well as

a typical FFT transform image chosen to indicate the possible variation in grain sizes/dimensions are

shown in figure 32 (a-e).

The grain size estimation has been performed by transforming the scanned area using 2D Fast

Fourier Transform as shown in figure 32-f and then taking a line scan across the transformed pattern.

Line scans were taken in both horizontal and vertical directions and the average value 1/ (2*FWHM)

of the peaks is noted as a rough estimation of the grain size. The variation of the grain shape in

minimum and maximum dimensions is evaluated from the ratio between the two line scans.

The surface roughness value calculations were performed by taking the average between the

two Rms Rq values which were the result of the aggregate value of vertical and horizontal line scans

over the height profile image of the scanned area. The error margins are the average between the two

calculated ranges. Table 4 contains a summary of the statistical estimation of the quantities. The

roughness and in-plane grain size values are also plotted against the number of periods in the film for

qualitative analysis figures 33, 34.

36

a.

b.

c.

d.

e.

f.

Figure 32: (a. - e.) 3x3μm2 scans of the all the periodic (CCO: NCO) samples in order of decreasing amount of periods (image a. is for 70 periods). f. image example of a typical 2D FFT transform pattern. The shape of the pattern indicates that the grains are elongated towards one dimension rather than being symmetrical.

37

Table 4: Summary of Statistical quantities for the superlattice (CCO:NCO) films of same total thickness (140nm). Color coding is kept for easier comparison with the XRD- analysis reported in the next subchapter.

Color

reference

Number of

periods

Period

thickness (nm)

Surface roughness

(Rms) in (nm)

Roughness range

(Rms) in (nm)

Grains

(nm)

Grains min and max

dimension (nm)

█ 4 35 6.7 ±0.9 189 98-366

█ 7 20 13 ±4.6 277 135-418

█ 14 10 6.6 ±1.6 143 93-281

█ 28 5 15.4 ±5.2 130 124-136

█ 70 2 25.6 ±7.2 191 172-210

Figure 33: Grain size over number of periods. The connected marks indicate the average grain size while the -, + symbols indicate the minimum and maximum calculated dimension according to the FFT analysis. The line is a guide to the eye.

Figure 34: Rms average surface roughness over number of periods. Error bars indicate the minimum and maximum rms value of the roughness. The line is a guide to the eye.

From this analysis it is evident that the surface roughness is increasing significantly with

increasing the amount of periods in the superlattice. Values up to 18.2% of the main film thickness

(140nm*) for the 1:1 nm 70 periods film and 11% for the 2.5:2.5nm, 28 period film. These high

values could be subjected to the following explanation.

Considering that the growth of NCO and CCO crystals is oriented along the c-axis and that the

c-axis parameter is reported to be roughly 1.1 nm, then the individual layer growth of these films is

roughly 0.9 monolayers and 2.3 monolayers respectively. At this thickness value and considering

diffusion mechanics, it is rather probable that the initial island formation of the deposited material is

not enough to form a well-structured and oriented layer.

*The Al2O3 amorphous capping layer is assumed to not contribute to the roughness factor since if it is optimally grown, it is

considered to follow the pattern of the underlying structure.

0

50

100

150

200

250

300

350

400

450

0 10 20 30 40 50 60 70 80

Gra

in s

ize

(nm

)

Number of periods

0

5

10

15

20

25

30

35

0 10 20 30 40 50 60 70 80

Av

g. R

ou

ghn

ess

(nm

)

Number of periods

38

This assumption can be backed by literature observations reporting the formation of an

unstructured buffer layer and stacking faults in the initial stages of growth of the CCO material when

grown directly on the substrate39, and could also be happening on a thin NCO layer . As the deposition

continues, the incoming material species could be latching onto the already formed islands with grater

probability rather than proceeding to the space in between. Therefore, as the growth of the film is

progressing with more species deposited on the formed islands and less in between, the roughness is

maintained at a high value. This qualitative analysis is a speculation and further quantitative analysis

would be needed for a conclusive report, i.e. TEM cross section images could give a more definite

explanation.

Other parameters which could influence the roughness is the initial structural coherency of

the substrate, thermal gradients on the substrate, impurity species in the deposition chamber,

fluctuations of the laser flux and texture of the target material surface. The roughness for the 14 and

4 period samples is notably lower both in value and in range while for the 7 period sample the

increased roughness could be related to misalignment between target stage and substrate.

Regarding the grain size the conclusion that can be drawn with respect to these five

superlattice samples, is that by increasing the period density and thus decreasing the individual layer

thickness there is significantly less variation in the grain dimension ratio. Therefore thinner interlayers

lead to more uniform grain formation. The exact composition of the grains however is not determined

since the NCO, CCO layers may be intermixing instead of stacking on top of each other. A trend in the

grain size cannot be clearly decided.

Crystalline analysis

The superlattice samples where scanned in the XRD. The analysis presented here took place

within two consecutive weeks to limit the variation of the device parameters and ambiance

conditions. A cumulative plotting of the full range scan for all the samples, as well as a closer θ/2θ

scan and a rocking curve scan over the (002) diffraction plane peak are given in figures 35 and 36.

Figure 37 shows the FWHM value of the spread of the rocking curve peak.

Figure 35: Cumulative plotting of the full range diffraction pattern for all the samples. The square (▪) symbols indicate the (002) planes of the film and the circles (○) indicate the LSAT substrate (001) planes. The graphs are normalized and verticaly shifted for clarity, the scale is logarithmic.

39

Figure 36: 2theta Scan of the (002) plane diffraction peak of the film. Rocking curve (Ω-scan) along the (002) peak, vertically shifted for clarity. FWHM values of the peaks indicate the planar deviation from the preferred orientation of growth.

Figure 37: Low FWHM value indicates the peak sharpness is grater and thus the planar orientation of the films is retained.

The cumulative plotting indicates the optimal alignment with respect to the LSAT substrate

(001) peak. A closer scan on the (002) peak of the film points out the fact that there is a merged

situation between the NCO and CCO preferred d-spacing. With limited variation for all five samples

the (002) peak is located at (θ≈8.185o±0.05o), and the corresponding d-spacing value is d≈10.82Å.

Since all the measurements in this chapter where conducted using the same intensity

attenuation factor, the relative intensity is a competent evidence of the degree of crystallinity of the

films with respect to the coherency of the observed planes. It is evident that a maximum crystallinity

is achieved for the 14-period sample hence individual layer thickness of 5 nm which accounts for 4-

5 unit cells for each material layer (NCO, CCO). Increasing or decreasing the periodicity relative to

that value, the influence on the crystallinity is negative.

In agreement to this observation, Full Width at Half Maximum (FWHM) values of the rocking

curve scans, demonstrate that the 14-period superlattice sample retains the preferred orientation

with minimum deviation in comparison to the other four films.

Seebeck and Resistivity analysis

The effect of the different number of periods on the Seebeck coefficient and sheet resistivity

have been measured for the (CCO:NCO)n superlattices. The measurements are reported in table 5.

Both attributes and the corresponding power-factor are plotted against the number of periods in

figures 38 and 39. The error values on the measurements are 5% and 7% over the measurement for

0

0.5

1

1.5

2

2.5

3

3.5

0 20 40 60 80

FWH

M v

alu

e (o

)

Number of periods

40

the Seebeck and the resistivity respectively while the error margin for the power factor is calculated

by error propagation theory from text books.

Table 5: Resistivity, Seebeck and combined power factor values.

Color

reference Number of

periods

Period

thickness

(nm)

Sheet resistivity

(mOhm*cm)

Seebeck coefficient

(μV/K) Power factor RT (W/m*K2)

█ 4 35 5.8 50.8 4.44938E-05 █ 7 20 7.85 82.6 8.69141E-05 █ 14 10 8.13 73.98 6.73191E-05 █ 28 5 6.29 68.27 7.40985E-05 █ 70 2 4.87 34.33 2.42002E-05

Figure 38: Sheet resistivity and Seebeck value over the number of periods. Error bars are 7% for the resistivity and 5% for the Seebeck. The line is a guide to the eye.

Figure 39: Calculated power factor over the number of periods. The error values are estimated by error propagation from the Seebeck and resistivity.

We were able to observe a trend in the effect of the different periodicity towards the sheet

resistivity and Seebeck values. Regarding the sheet resistivity the first important observation was that

it remained at low values despite the introduction of more electron scattering interfaces and possibly

more point defects. Increasing the amount of periods seems to have a negative effect on the Seebeck

coefficient which decreases by up to a factor of two however the resistivity values are also decreasing

0

20

40

60

80

100

120

140

0

1

2

3

4

5

6

7

8

9

10

0 10 20 30 40 50 60 70 80

Seeb

eck

co

effi

cien

t (μ

V/K

)

Shee

t re

sist

ivit

y

(mO

hm

*cm

)

Number of periods

0.E+00

2.E-05

4.E-05

6.E-05

8.E-05

1.E-04

1.E-04

0 10 20 30 40 50 60 70 80

Po

wer

fac

tor

(W/m

*K2

)

Number of periods

41

by roughly the same factor. The color reference of the samples is kept the same as in the XRD analysis

to mitigate linked observations.

The sheet resistivity values present a maximum at 7.85 and 8.13 mOhm*cm for the samples

with 7 and 14 periods respectively. These samples also show the maximum planar coherency as

concluded from the XRD analysis. Correspondingly these two samples show the maximum Seebeck

values at 82.6 and 74 μV/K. For the samples with 7, 14 and 28 periods the power factor was similar

within the error margins with the maximum for the 7 period sample at ~0.9 (μWK-2cm-1).

The sample with 70 periods showed the minimum sheet resistivity, accompanied however

with the lowest Seebeck coefficient. The films with 4 periods of 35nm showed an intermediate

behavior.

Electronic analysis

In the Van der Pauw resistivity measurements as well as the Seebeck measurements, most of

the superlattice samples have shown a significant variation in measurements along the two directions.

In order to further investigate the source of this variation, as a final measurement all the superlattice

films with variable period density were analyzed with the physical property measuring system (PPMS).

Hall resistivity measurements where conducted at room temperature and through a magnetic sweep

from -9 to 9 tesla. From the slope of the plotting of the Hall resistances over the applied magnetic

field, a rough estimation of the carrier density was achieved. The estimated carrier density together

with the resistivity values where subsequently used to calculate the carrier mobility.

For reasons related to thermal stability, only the last part of the magnetic swipe was taken

into account for the calculations in all the measurements. Thermal drift effects where corrected prior

to the analysis by calculating the change of resistance over time with respect to the value at zero tesla

and subtracting this value from the rest of the measurements. As previously explained in the thesis,

all of the samples containing NCO where not completely stable over time. Therefore for the PPMS

measurements the resistivity was determined again for every superlattice sample at the moment of

the measurement and this “new” value was for taken into account for the carrier density and mobility

calculations. All the calculations summarized In Table 6.

Table 6: Summary of the carrier density and carrier mobility calculations. The n1, μ1 and n2, μ2 values correspond to the two

different slopes for RH1, RΗ2 over B in every sample. The average n-value is the average between the two and it is used to

estimate the μ-average.

Color

reference

Number

of

periods

n1

(*1022)

(cm-3 )

n2

(*1022)

(cm-3 )

navg

(*1022)

(cm-3)

μ1

(cm2/Vs)

μ2

(cm2/Vs)

μavg

(cm2/Vs)

resistivity

(mOhm*cm)

█ 4 0.677 0.203 0.440 0.153 0.511 0.236 6.02

█ 7 0.419 0.078 0.248 0.144 0.776 0.243 10.35

█ 14 0.216 0.093 0.154 0.372 0.867 0.520 7.78

█ 28 0.178 0.158 0.168 0.526 0.595 0.558 6.65

█ 70 0.966 0.414 0.690 0.128 0.299 0.179 5.05

42

Figure 40: Average mobility (blue) and carrier density (red) plotted against the number of periods. The lines are guide to the

eye.

As expected from the previous observations the difference in the two Hall resistances were

again confirmed, and through the analysis it can be attributed to different carrier densities and hence

different carrier motilities, along the two directions of the Hall measurements.

The reported values n1, n2 are given in 1022*cm-3 and correspond to the slopes of the linear

fitting of the two different RH1 and RH2 over the magnetic field B observed in all the samples, navg is

the average value between the two. The μ1, μ2 values are in (cm2/Vs) and are calculated from n1, n2,

respectively while the μavg value is calculated from navg. The relation of the calculated average values

for carrier density and mobility to the period density of the superlattices is also plotted in figure 40.

The error from the fitting of the line in terms of the fitting parameter R2 was very small to be

considered significant.

The directional inhomogeneity of the carrier density and mobility values can be directly

related to the variation in the shape of the grains observed earlier in the AFM analysis. For the

superlattices with 28 and 70 period density, where the grain statistics showed relative homogeneity,

the estimated mobility also shows limited dispersion. For the superlattices with 4, 7 and 14 periods

however, for which the shape of the grains shows significant variation, the calculated values n1, μ1 and

n2, μ2 differ from each other in correspondence.

There are other factors that could induce a difference in the carrier density and mobility values

along the two directions of measurement for the superlattice films. One could be Inhomogeneity in

the film’s thickness and another factor could be the instability of the NCO layers of the superlattice.

Regarding the later, as indicated previously in the thesis, NCO interacts with H2O in the atmosphere

and such an interaction initiates at the borders of the 5x5mm2 sample and proceeds inwards at a

random pattern. By doing so it passivates the ionic species thus altering the carrier density and the

mobility in an unpredictable way.

0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

0.9

0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

0.9

0 10 20 30 40 50 60 70 80

(n1+

n2)/

2 (

cm-3

) (*

10

22)

μ a

vg

(cm

2/V

s)

Number of periods

43

6. Conclusion and Recommendation

6.1. Concluding overview

Epitaxial thin films of NaxCoO3 and Ca3Co4O9 both in single film composition and in superlattice

combinations were successfully deposited and studied during the course of this research. Structural

properties of the superlattices and their dependency on the total thickness and period density have

been evaluated. Preliminary correlation between structural and electronic properties has been

achieved.

An overview of the Seebeck, resistivity and power factor values for all of the samples that

were studied during this work is given in table 7. With (cyan and green), we denote the pest

performing CCO, NCO single films in accordance and with (gold and purple) the best performing

superlattice films in the two superlattice categories studied: variable total thickness and variable

period density over constant thickness correspondingly. The comparison is done between samples

grown on LSAT 001 substrates The samples in (red) are those for which instability in the O2 pressure

during cooldown phase was observed and the deposition was repeated with the non-affected sample

giving better results.

Table 7: Overview of the Seebeck, Resistivity and power factor values for all the samples studied during the thesis.

Composition Ratio Total Film Thickness

(nm)

Seebeck coefficient

(μV/K)

Sheet resistivity

(mOhm*cm)

Power factor RT (W/m*K2)

x10-4 CCO/ALO single film 85 108.4 5.6 2.1 CCO/ALO single film 27.5 101.8 18.9 0.55

CCO/ALO at 750oC single film 60 135.4 21.4 0.86 CCO/LSAT single film 91.2 140.3 19.5 1 CCO/LSAT single film 53 141.9 24.22 0.83 NCO/LSAT single film 100 87.2 1.61 4.7 NCO/LSAT single film 70 21.1 3.37 0.13 NCO/LSAT single film 50 22 4.92 0.1

CCO:NCO/LSAT 70:10 80 90.3 10.92 0.74 NCO:(CCO:NCO)x2 10:(10:10) 50 109.3 9.04 1.3

(CCO:NCO)x3 (10:10) 60 121.9 11.35 1.3 (CCO:NCO)x5 (10:10) 100 54.6 10.27 0.29 (CCO:NCO)x5 (10:10) 100 80.4 9.78 0.66 (CCO:NCO)x4 (17.5:17.5) 140 50.1 5.8 0.43 (CCO:NCO)x7 (10:10) 140 82.6 7.85 0.87

(CCO:NCO)x14 (5:5) 140 74 8.13 0.67 (CCO:NCO)x28 (2.5:2.5) 140 42.3 7.6 0.24 (CCO:NCO)x28 (2.5:2.5) 140 68.3 6.29 0.74 (CCO:NCO)x70 (1:1) 140 34.3 4.87 0.24

44

From the XRD analysis of the films, It was clear that the crystalline planar coherency is affected

by the period density, with the 5nm:5nm (CCO/NCO)x14 giving the sharpest diffraction peak and the

least spread in the rocking curve scan. This is suggestive that 4-5 monolayers of each material are

required in one period to achieve the optimal crystal structure in the orientation of growth of such a

superlattice and before planar coherence is affected by strain effects. Planar coherence however is

not the dominant factor for the measured in-plane electronic conductivity to be optimal. Correlation

between AFM and electronic results indicate that electronic conductivity it is rather dependent on the

size and homogeneity of the grain formation. However a certain level of variation in the grain shape

is required for a high Seebeck value. Further evaluation of the in-plane crystal coherency and

Transmission Electron Microscopy measurements of the samples could shed more light in the

correlation between structural an electronic properties.

Overall we have observed that the thermoelectric properties of the superlattice NCO:CCO

films grown on LSAT are maintained in good levels with respect to the overall film thickness. The best

performing superlattice films in terms of power factor were 1.3 (W/m*K2) x10-4 for the 60nm thick

(CCO:NCO) x3 with 10nm:10nm period thickness ratio and 0.87(W/m*K2) x10-4 for the 140nm thick

(CCO:NCO) x7 again with 10nm:10nm period thickness ratio. The power factor values are comparable

in the order of magnitude with results reported for topotactic synthesis of alternatively stacked

CCO:NCO composites from ref21 taking into account that the reported values are at 1000 oK.

The effect of the period density in superlattice CCO/NCO thin films on the thermal

conductivity remains to be seen in future work.

6.2. Experimental considerations

PLD: Laser alignment and fluctuations. After the elongation of the laser path there were

significantly more complications in aligning the laser beam and keeping it stable on the exact same

spot. Since every small vibration can greatly destabilize the linear laser path affecting both the fluency

variation on the material target and the angle of incidence which in turn influences the directionality

of the plasma plum. These factors can affect the overall quality of the deposited film especially for the

3-4 hour long depositions of the thicker samples. Moreover, although the distance between the heater

stage and the target stage (y-value) is kept stable, assuming optimal system calibration, the final

parallel (z, x, θ) alignment of the substrate towards the target stage is always done by eye. After a lot

of trial and error depositions, where the misalignment affect was obvious, a certain optimization was

achieved. However deviations are always to be expected and accounted for.

Unstable Oxygen pressure during the cooldown phase of the deposition. This situation has

been observed in several occasions during the sample depositions. For some of these samples the

deposition has been repeated. And the corresponding samples indicated higher values of surface

roughness, lower degree of crystallinity and performed worse in measurements compared to the re-

deposited ones that had no cooldown complications. This instability was accidentally noticed and was

not always sought after. Therefore there may have been more “affected” samples. Oxygen defects

and vacancies play a crucial role in the structural coherency of the deposited films affecting the

crystallinity and in turn the electronic properties of the sample.

Capping layer. In order to improve the coverage of the amorphous ALO layer on the sample,

after the initial deposition of a certain layer thickness directly on the sample’s surface (at 0-degrees

orientation towards normal surface vector), the heater stage can be set into two more positions at 45

45

and -45 degrees of tilt and rotation and for both positions a few nm of ALO material can be deposited.

This way the sides of the sample would also be better protected

Ion etching and deposition of gold contacts. Due to the volatility of the NaxCoO2 the final

samples are prohibited to come in contact with acetone or ethanol. Therefore the more precise

photoresist dry etching masks could not be used in this research. Instead we used the hard cross-like

metallic mask which has to be manually placed on each sample and held on by kapton tape, before

the etching and sputtering treatment, leaving a lot of room for error. Two of the most common

problems are: scratching the surface of the sample which can affect the quality and performance of

the capping layer and inhomogeneous exposed surface (in the corners) which leads to unequal etching

and variable contact sizes. The quality of the contacts in turn, affects the electronic, and Seebeck

measurements.

An optimized solution could be to use metallic “open box” like cover masks of predefined

dimensions and with precisely defined exposed features in the corners or in 4 symmetrical points

further inside the sample area, if we want to avoid the sides where the NCO layers are more likely to

be exposed to H2O in the environment. In this case, taking into account the substrates’ dimensions, a

metallic open-box of (5.1x5.1x1.2) mm3 with precisely defined exposed corners i.e. (500x500) μm2

could be used. This could improve the quality of the etching in terms of accuracy as well as the

homogeneity of the gold contacts both in respect to positioning and in respect to the amount of gold

sputtered in each contact. Moreover such a mask should induce improved protection over the sides

of the sample during the dry ion etching procedure, preventing undesired etching of the capping layer.

A precise estimation of the etching rate of both the capping layer and the combined stack of CCO/NCO

layer is also essential for the optimization of the procedure of the contact deposition.

Custom Seebeck measurement setup. Ideally the custom setup could be placed under a hood

with a more controlled ambiance and less fluctuations in the air flow that can influence the

measurements. A mechanism able to apply soft pressure on the sample surface or air suction from

the bottom of the substrate could help optimize a consistent position and contact of the sample with

respect to the two Peltier elements. This suggestion could help increase the induced temperature

gradient on the sample due to the better contact and also improve the stability and accuracy of the

Seebeck measurements.

6.3. Suggestions for future research

As future work in continuation to the research on the CCO:NCO superlattices firstly the

thermal conductivity of chosen samples should be measured and compared with single material films

of the same thickness, along with the Seebeck and conductivity values. Possible methods to do this

measurements by some further treatment of the samples is the Harman method40 , the time domain

thermo-reflectance41 and the 3-ω technique42,43. The samples are expected to show reduced out-of-

plane thermal conductivity, however the in-plane thermal conductivity measurement is more relevant

and can be used in relation to our measurements to calculate the zT figure of merit.

Furthermore the variation in the period thickness within a superlattice sample i.e. intermixing

periods of CCO:NCO with different thicknesses or inducing some period thickness gradient could be

tested. Such structures could potentially reduce even further the lattice thermal conductivity by

blocking an extended number of phonon modes due to the increased structural complexity of the film.

46

This suggestion involves taking into account the optimal period growth, in terms of planar coherence,

shown in this work i.e. keeping each layer thickness between 2.5-10nm.

An alternative approach regarding the CCO:NCO superlattices is the pre-patterning of several

substrates i.e. LSAT as used in this work, with photoresist in a desired way, such as stripes of different

nanometer thickness along the substrate surface. Then a layer of CCO material could be deposited at

high temperature 750oC, which is the optimal scenario for CCO, in the unmasked area. The substrates

with the CCO are then treated to remove the photoresist (and the material on top), leaving only the

desired stripes of CCO material on the substrate. The samples should then be used as substrates for

the growth of a superlattice structure as was done in this work. Starting with an NCO layer which will

be deposited directly on the substrate material in between the CCO stripes as well as on top of the

CCO, maintaining the initial patterning of the film. The growth of the superlattice should then continue

up to a desired thickness. This method should be possible due to the structural compatibility of the

cobaltites and could be useful in order to induce in-plane interfaces to the superlattice and further

the benefits from limiting the thermal conductance.

Finally a slightly different direction could be to avoid the volatile and unstable NaxCoO3 and

the capping layer and use another doped cobaltite with structural similarities in favor of reducing the

resistivity. Some suggestions could be elements such as Bi, Fe, B and Ga, for which reports show

already an improvement in the thermoelectric properties44-46.

47

48

Acknowledgements

It has been 12 months since I have undertaken this master assignment. Starting from nearly

zero level regarding lab experience and only partial knowledge in materials science, there was a lot of

catching up to do and it was not always easy . Through a procedure of trainings, experiments, studying

and writing this productive and self-developing period has reached a final point, the outline of which

is included in this report. But it cannot all be captured in written word as the whole experience has

been much more than what a report can show; and it is the people, the attitude and the group

environment that make the story whole.

Taking this opportunity I would like to express my gratitude to the whole Inorganic Material

Science Group where I spent this exciting year, starting from the research chair Prof. Guus for being

an excellent host. A big thanks to Peter who was my daily supervisor during his final months in the

group and gave his best effort to train me up to a level where I could be self-sufficient, also keeping

his cool during the “oups, what the …. happened?” moments. A special thanks to Mark for trusting in

me when I approached him and asked for a project in renewable energy and materials while declaring

my “inexperience” in the topic and for his motivating guidance throughout the project.

Thanks to Alim, David, Tom and Ron for their help, useful insight and recommendations during

the times when I experienced problems with the PLD depositions, as well as for their efficient co-

operative planning of the use of the laser. Special thanks to Kurt for his XRD-training and for shedding

light on the measurements analysis but also for all the other totally irrelevant but highly interesting

and fun discussions. Big thanks of course to my fellow master students “fishbowl buddies” Jaap,

Jasper, Jeffrei, Thomas and Roi for their help in specific topics but mostly for all the fun experience

and the “warm climate” in and out of the group, during this period.

I would also like to thank the technical personnel Dominic, Henk, Laura, Jose and Karin for

being available when specific tools, laser refills and other maintenance issues where required as well

as for keeping the equipment running as smoothly as possible. Thanks also to Frank for the training

on the Perkin Elmer and the Argon Ion Etcher. Appreciation also to Gertjan and Andre as well as all

the rest of the IMS people, physicists and chemists alike for all the interesting conversations during

the coffee brakes and lunches, all of you played a significant part in the overall experience.

Lastly I would like to thank Harrold for being my study advisor throughout my MSc and for

always being encouraging and motivating. Here I would also like to appreciate the enormous support

of my Family back In Greece, as well as all of my friends who played their part in maintaining the

balance between my study and social life.

49

50

Bibliography:

1 Writers, S. Waste heat to electricity, <http://www.energy-daily.com/reports/World_record_holder_999.html> (last visited 20/08/15)

2 He, J., Liu, Y. & Funahashi, R. Oxide thermoelectrics: The challenges, progress, and outlook. Journal of Materials Research 26, 1762-1772, doi:10.1557/jmr.2011.108 (2011).

3 Holland, M. G. Analysis of lattice thermal conductivity. Physical Review 132, 2461-2471, doi:10.1103/PhysRev.132.2461 (1963).

4 Snyder, G. J. & Toberer, E. S. Complex thermoelectric materials. Nature Materials 7, 105-114, doi:10.1038/nmat2090 (2008).

5 Heremans, J. P. et al. Enhancement of thermoelectric efficiency in PbTe by distortion of the electronic density of states. Science 321, 554-557, doi:10.1126/science.1159725 (2008).

6 Humphrey, T. E. & Linke, H. Reversible thermoelectric nanomaterials. Physical Review Letters 94, doi:10.1103/PhysRevLett.94.096601 (2005).

7 Pei, Y., Lalonde, A. D., Wang, H. & Snyder, G. J. Low effective mass leading to high thermoelectric performance. Energy and Environmental Science 5, 7963-7969, doi:10.1039/c2ee21536e (2012).

8 Sigma-Aldrich. Materials for Advanced Thermoelectrics, <http://www.sigmaaldrich.com/materials-science/metal-and-ceramic-science/thermoelectrics.html> (last visited 20/08/15)

9 Backhaus-Ricoult, M., Rustad, J., Moore, L., Smith, C. & Brown, J. Semiconducting large bandgap oxides as potential thermoelectric materials for high-temperature power generation? Applied Physics A: Materials Science and Processing 116, 433-470, doi:10.1007/s00339-014-8515-z (2014).

10 Piotrowski, T., Jung, W. & Sikorski, S. Application of non-linear theory to analysis of Benedicks effect in semiconductors. Physica Status Solidi (A) Applications and Materials Science 204, 1063-1067, doi:10.1002/pssa.200674126 (2007).

11 Huang, Y., Zhao, B., Lin, S., Ang, R. & Sun, Y. Enhanced thermoelectric performance induced by Cr doping at Ca-Sites in Ca3Co4O9 system. Journal of the American Ceramic Society 97, 3589-3596, doi:10.1111/jace.13144 (2014).

12 Wu, T. et al. On the origin of enhanced thermoelectricity in Fe doped Ca 3Co4O9. Journal of Materials Chemistry C 1, 4114-4121, doi:10.1039/c3tc30481g (2013).

13 Obata, K. et al. Thermoelectric properties of Ca3Co4O 9/[Ca2(Co0.65Cu0.35) 2O4]0.624CoO2 composites. Journal of Electronic Materials 43, 2425-2429, doi:10.1007/s11664-014-3109-2 (2014).

14 Terasaki, I., Sasago, Y. & Uchinokura, K. Large thermoelectric power in ${\mathrm{NaCo}}_{2}{\mathrm{O}}_{4}$ single crystals. Physical Review B 56, R12685-R12687 (1997).

15 Liu, P. et al. High temperature electrical conductivity and thermoelectric power of NaxCoO2. Solid State Ionics 179, 2308-2312, doi:10.1016/j.ssi.2008.08.010 (2008).

16 Demchenko, D. O. & Ameen, D. B. Lattice thermal conductivity in bulk and nanosheet NaxCoO 2. Computational Materials Science 82, 219-225, doi:10.1016/j.commatsci.2013.09.049 (2014).

17 Viciu, L. et al. Crystal structure and elementary properties of Nax Co O2 (x=0.32, 0.51, 0.6, 0.75, and 0.92) in the three-layer NaCo O2 family. Physical Review B - Condensed Matter and Materials Physics 73, doi:10.1103/PhysRevB.73.174104 (2006).

18 Brinks, P., Heijmerikx, H., Hendriks, T. A., Rijnders, G. & Huijben, M. Achieving chemical stability in thermoelectric Na xCoO 2 thin films. RSC Advances 2, 6023-6027, doi:10.1039/c2ra20734f (2012).

51

19 Brinks, P. et al. Enhanced thermoelectric power factor of NaxCoO2 thin films by structural engineering. Advanced Energy Materials 4, doi:10.1002/aenm.201301927 (2014).

20 Brinks, P., Rijnders, G. & Huijben, M. Size effects on thermoelectric behavior of ultrathin NaxCoO2 films. Applied Physics Letters 105, doi:10.1063/1.4901447 (2015).

21 Liu, J., Huang, X., Sun, Z., Liu, R. & Chen, L. Topotactic synthesis of alternately stacked Ca3Co 4O9/γ-Na0.66CoO2 composite with nanoscale layer structure. CrystEngComm 12, 4080-4083, doi:10.1039/c0ce00430h (2010).

22 Wu, T. et al. A structural change in Ca 3Co 4O 9 associated with enhanced thermoelectric properties. Journal of Physics Condensed Matter 24, doi:10.1088/0953-8984/24/45/455602 (2012).

23 Brinks, P., Van Nong, N., Pryds, N., Rijnders, G. & Huijben, M. High-temperature stability of thermoelectric Ca<inf>3</inf>Co<inf>4</inf>O<inf>9</inf> thin films. Applied Physics Letters 106, doi:10.1063/1.4917275 (2015).

24 Klie, T. P. a. R. F. Characterization of Thermoelectric Ca3Co4O9 Microstructure Using Transmission Electron Microscopy. Journal of Undergraduate Research 5 (2011).

25 Sun, T., Hng, H. H., Yan, Q. & Ma, J. Effects of pulsed laser deposition conditions on the microstructure of Ca 3Co 4O 9 thin films. Journal of Electronic Materials 39, 1611-1615, doi:10.1007/s11664-010-1261-x (2010).

26 Fujii, S. et al. Impact of dynamic interlayer interactions on thermal conductivity of Ca3Co4O9. Journal of Electronic Materials 43, 1905-1915, doi:10.1007/s11664-013-2902-7 (2014).

27 Brinks, P. Sise Effects in Thermoelectric Cobaltate Heterostructures, University of Twente, (2014).

28 Takabatake, T., Suekuni, K., Nakayama, T. & Kaneshita, E. Phonon-glass electron-crystal thermoelectric clathrates: Experiments and theory. Reviews of Modern Physics 86, 669-716, doi:10.1103/RevModPhys.86.669 (2014).

29 Sales, B. C. Electron crystals and phonon glasses: A new path to improved thermoelectric materials. MRS Bulletin 23, 15-21 (1998).

30 Koga, T., Sun, X., Cronin, S. B. & Dresselhaus, M. S. Carrier pocket engineering to design superior thermoelectric materials using GaAs/AlAs superlattices. Materials Research Society Symposium - Proceedings 545, 375-380 (1999).

31 Peranio, N., Eibl, O. & Nurnus, J. Structural and thermoelectric properties of epitaxially grown Bi 2Te3 thin films and superlattices. Journal of Applied Physics 100, doi:10.1063/1.2375016 (2006).

32 Ravichandran, J. et al. Crossover from incoherent to coherent phonon scattering in epitaxial oxide superlattices. Nature Materials 13, 168-172, doi:10.1038/nmat3826 (2014).

33 Landry, E. S., Hussein, M. I. & McGaughey, A. J. H. Complex superlattice unit cell designs for reduced thermal conductivity. Physical Review B - Condensed Matter and Materials Physics 77, doi:10.1103/PhysRevB.77.184302 (2008).

34 M., I. Structural engineering of Ca3Co4O9 thermoelectric thin films, Univercity of Twente, (2013).

35 Atomic Force Microscopy, <http://www.teachnano.com/education/AFM.html> (last visited:20/08/15)

36 Mitsunaga, T. X-ray thin-film measurement techniques. 37 Inaba, K. X-Ray Thin-Film Measurement Techniques I. Overview. The Rigaku Journal 24, 10-

15 (2008). 38 Masashi, M. et al. Thermoelectric Properties of Two Na x CoO 2 Crystallographic Phases.

Japanese Journal of Applied Physics 42, 7383 (2003). 39 Qiao, Q. et al. Effect of substrate on the atomic structure and physical properties of

thermoelectric Ca3Co4O9 thin films. Journal of Physics Condensed Matter 23, doi:10.1088/0953-8984/23/30/305005 (2011).

52

40 Satake, A., Tanaka, H., Ohkawa, T., Fujii, T. & Terasaki, I. Thermal conductivity of the thermoelectric layered cobalt oxides measured by the Harman method. Journal of Applied Physics 96, 931-933, doi:10.1063/1.1753070 (2004).

41 Belliard, L. et al. Determination of the thermal diffusivity of bulk and layered samples by time domain thermoreflectance: Interest of lateral heat diffusion investigation in nanoscale time range. Journal of Applied Physics 117, doi:10.1063/1.4908068 (2015).

42 Liu, Y. Q. et al. in IOP Conference Series: Materials Science and Engineering.1 edn. 43 Maize, K. et al. in Annual IEEE Semiconductor Thermal Measurement and Management

Symposium. 185-190. 44 Tian, R. et al. Enhancement of high temperature thermoelectric performance in Bi, Fe co-

doped layered oxide-based material Ca3Co4O 9+δ. Journal of Alloys and Compounds 615, 311-315, doi:10.1016/j.jallcom.2014.06.190 (2014).

45 Demirel, S., Avci, S., Altin, E., Altin, S. & Yakinci, M. E. Enhanced thermoelectric properties induced by chemical pressure in Ca 3Co4O9. Ceramics International 40, 5217-5222, doi:10.1016/j.ceramint.2013.10.090 (2014).

46 Constantinescu, G. et al. Effect of Ga addition on Ca-deficient Ca3Co4O y thermoelectric materials. Ceramics International 40, 6255-6260, doi:10.1016/j.ceramint.2013.11.083 (2014).

53