potential enhancement of thermoelectric energy conversion in cobaltite superlattices
TRANSCRIPT
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
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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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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
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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.
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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
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