Non-metal Alkylsilyl Compounds as
Precursors in Atomic Layer Deposition
of Chalcogenides and Pnictides
Tiina Sarnet
Laboratory of Inorganic Chemistry
Department of Chemistry
Faculty of Science
University of Helsinki
Helsinki, Finland
Academic Dissertation
To be presented, with the permission of the Faculty of Science of the University of Helsinki,
for public criticism in Auditorium A110 of the Department of Chemistry (Chemicum) on
December 21st 2015 at 12 o’clock noon.
Helsinki 2015
Supervisors
Professor Mikko Ritala
and
Professor Markku Leskela¨
Laboratory of Inorganic Chemistry
Department of Chemistry
University of Helsinki
Helsinki, Finland
Reviewers
Professor Harri Lipsanen
Department of Micro- and Nanosciences
School of Electrical Engineering
Aalto University
Espoo, Finland
Doctor Claudia Wiemer
Laboratorio MDM
Institute for Microelectronics and Microsystems
National Research Council (CNR)
Agrate Brianza, Italy
Opponent
Professor Kornelius Nielsch
Institute for Metallic Materials
Leibniz Institute of Solid State and Materials Research Dresden
Dresden, Germany
ISBN 978-951-51-1821-9 (paperback)
ISBN 978-951-51-1822-6 (PDF version)
http://ethesis.helsinki.fi/
Unigrafia
Helsinki 2015
“Inspiration is the momentary cessation of stupidity.”
Abstract
Materials are crucial to the technological advances of society. The never ending need for data
storage and new energy sources pushes research towards clear goals. Perhaps some of today’s
solutions can in the future be replaced or augmented with phase change memories and thermo-
electric materials. Phase change materials store data in their amorphous and crystalline phases
that have great differences in their electrical and optical properties. Thermoelectric materials
can utilize waste heat and produce electricity from temperature differences. They can also be
utilized in temperature control as they can be used to create a temperature difference by using
electricity.
Shrinking device sizes and increasing device complexity require that deposition methods such
as atomic layer deposition (ALD) are used. ALD is based on sequential, saturative surface
reactions. Precursors are brought to the surface one at a time, separated by purges. Because
of the saturative reactions, each ALD cycle deposits a constant amount of material up to a
monolayer, making film thickness control very simple.
ALD of chalcogenides has focused mainly on sulfides, and the chemistries for selenide and tel-
luride depositions have been limited. Pnictides have a similar situation. The ALD chemistries
for arsenides include only a few combinations of precursors, and antimonides are barely demon-
strated. This is why a new group of precursors was needed. The alkylsilyl non-metal precursors
react very efficiently with metal halides in a dehalosilylation reaction. These types of reactions
have now been utilized in both chalcogenide and pnictide thin film growth.
In this thesis, several chalcogenide and pnictide ALD processes were studied in detail by uti-
lizing the appropriate alkylsilyl non-metal precursors. In general, typical ALD characteristics
were found. Growth rates saturated with respect to precursor pulse lengths; film thicknesses
increased linearly with the number of deposition cycles; and the films were stoichiometric with
low impurity contents. Application wise, the ALD chalcogenide and pnictide films had the re-
quired properties. The phase of the phase change materials could be repeatably and quickly
changed, and the thermoelectric films showed a proper response to a temperature gradient.
iv
Preface
The natural sciences, and especially chemistry, have always been my favorite subjects at school.
I must thank my teachers Mrs. Anja Peitola for bringing out the joy of chemistry and teaching
the basic skills needed later on, and Dr. Elina Na¨sa¨kka¨la¨ for for letting us roam free in the lab.
During my studies at the university, it became abundantly clear that inorganic chemistry was
the only way to go. This feeling became stronger after a couple of summer jobs at the laboratory.
In the end, I never left.
My supervisors Prof. Mikko Ritala and Prof. Markku Leskela¨ are thanked for sharing their
seemingly endless knowledge about the elements, and ALD in particular. It has been a privilege
to learn from you during these years.
I am deeply grateful for the insightful and constructive comments of the official reviewers of
this thesis, Prof. Harri Lipsanen and Dr. Claudia Wiemer. My work was greatly improved due
to your suggestions.
In research, no one is an island. There have been so many people who had a large part in
this work. The following co-authors and collaborators are thanked for their involvement in
this work. I would like to thank Dr. Viljami Pore for initially introducing me to the world of
ALD, Mr Timo Hatanpa¨a¨ for his synthetic genius and preparing all the alkylsilyl precursors
used during this research, and Dr. Marianna Kemell for her vast knowledge relating to the
SEM and EDX. In addition, I am grateful to Dr. Marko Vehkama¨ki for his expertise with the
TEM and electrical characterizations; Dr. Esa Puukilainen and Mr. Miika Mattinen for atomic
force microscopy; Dr. Jani Ha¨ma¨la¨inen for sharing his knowledge of the inner workings of our
reactors; Dr. Kjell Knapas for knowledge on all things chemistry; and Mr. Robert Huggare for
breathing life into my ailing reactor. Special thanks go to Mr. Mikko J. Heikkila¨, for his tireless
help with everything relating to x-rays.
Many people outside our laboratory have greatly contributed to this work. Prof. Timo Sajavaara
and Dr. Mikko Laitinen from the University of Jyva¨skyla¨ as well as Dr. Kenichiro Mizohata from
the Material Physics Division are thanked for film composition analysis. I would like to thank
Prof. Jouni Ahopelto and Mr. Timo Flyktman from VTT for thermoelectic characterizations.
The help with LEIS and its intricacies from Mr. Rik ter Veen of Tascon is greatly appreciated.
Dr. Alejandro Schrott, Dr. Yu Zhu and Dr. Huai-Yu Cheng from IBM and Macronix are thanked
for their assistance with phase change related measurements. My heartfelt thanks go also to
Dr. Simone Raoux from the Helmholtz-Zentrum Berlin for not only the research collaboration,
v
but especially for the knowledge shared and all the fruitful conversations along the years. Mr.
Howard McKee is thanked for his expertise in the English language, for proofreading this thesis
and some of the articles it consists of.
The long-term financial support from ASM Microchemistry, and the support from the Finnish
Centre of Excellence in Atomic Layer Deposition is gratefully acknowledged. Furthermore, I am
most thankful of the Kemian Pa¨ivien Sa¨a¨tio¨ for awarding me a personal grant for my studies.
Of course I would like to thank all present and former colleagues at the Laboratory of Inorganic
Chemistry for fruitful work-related conversations, for the completely absurd discussions and
especially for the silly jokes. In addition, thanks are due to everyone playing floorball on Fridays.
What better way to end a stressful week at work than to tackle a co-worker or a supervisor
during a game.
The effect of peer support can never be underestimated. I have met some very special people
along the years. Timo, you have been there for literally most of my life, and I hope we will
have even more to reminisce about in the coming years. Thanks to Tatu, Reeta, Sonja and all
the others involved with Esitisle and HYK. It was a welcome distraction from all the studying.
Special thanks go to my partner in crime, Miia. We have been together since we were still young
and innocent first year chemistry students, and still share an office. We have vented frustration,
laughed at various shenanigans, and eaten enough chocolate throughout the years.
I owe my deepest gratitude to my parents, who have encouraged me in everything I do. I can
only say I have tried to make the best of the chances I have gotten. The bragging rights are
truly yours.
In the end, I would like to express my most sincere appreciation and love to my fiance´ Jason.
You kept me remotely sane and functional in the darkest depths of thesis desperation, made
me laugh every day, and took me out of the house when I was going stir crazy. It is time to get
a life now.
Helsinki, November 2015
Tiina Sarnet
vi
List of original publications
This doctoral dissertation consists of a summary and of the following publications, which are
referred to in the text by their Roman numerals. The author’s contribution to the publications
is described in the indentations.
I Atomic Layer Deposition of Antimony and its Compounds Using
Dechlorosilylation Reactions of Tris(triethylsilyl)antimony
Viljami Pore, Kjell Knapas, Timo Hatanpa¨a¨, Tiina Sarnet, Marianna Kemell, Mikko
Ritala, Markku Leskela¨ and Kenichiro Mizohata Chem. Mater. 23, 247 (2011)
The author planned and conducted the deposition experiments with V.P.,
and analyzed most of the thin films with M.K.. T.H. synthesized the alkylsilyl
precursor. K.K. perfomed the in-situ reaction mechanism studies. V.P. and
K.K wrote the paper, with other authors contributing.
II Atomic Layer Deposition and Characterization of GeTe Thin Films
Tiina Sarnet, Viljami Pore, Timo Hatanpa¨a¨, Mikko Ritala, Markku Leskela¨, Alejandro
Schrott, Yu Zhu, Simone Raoux and Huai-Yu Cheng J. Electrochem. Soc. 152, D694
(2011)
The author planned and conducted the deposition experiments and per-
formed most of the analysis on the films. T.H. synthesized the alkylsilyl
precursor. Laser testing and resistivity experiments were performed by A.S.,
Y.Z., S.R. and H-Y.C. The author wrote the first draft, and finalized the
paper together with other co-authors.
vii
III Atomic Layer Deposition and Characterization of Bi2Te3 Thin Films
Tiina Sarnet, Timo Hatanpa¨a¨, Esa Puukilainen, Miika Mattinen, Marko Vehkama¨ki,
Kenichiro Mizohata, Mikko Ritala and Markku Leskela¨ J. Phys. Chem. A 119, 2298
(2015)
The author planned and conducted the deposition experiments and per-
formed most of the analysis on the films. T.H. synthesized the alkylsilyl
precursor. AFM analysis was performed by E.P. and M.M., TEM analysis
by M.V. and TOF-ERDA by K.M.. The author wrote the first draft, and
finalized the paper together with M.R.
IV (Et3Si)2Se as a Precursor for Atomic Layer Deposition: Growth Analysis of
Thermoelectric Bi2Se3
Tiina Sarnet, Timo Hatanpa¨a¨, Marko Vehkama¨ki, Timo Flyktman, Jouni Ahopelto,
Kenichiro Mizohata, Mikko Ritala and Markku Leskela¨ J. Mater. Chem. C, 3, 4820
(2015)
The author planned conducted the deposition experiments and performed
most of the analysis on the films. T.H. synthesized the alkylsilyl precursor.
TEM analysis was performed by M.V., TOF-ERDA by K.M., and electrical
characterization by T.F. and J.A.. The author wrote the first draft, and
finalized the paper together with M.R.
V Alkylsilyl Compounds as Enablers of Atomic Layer Deposition: Analysis of
(Et3Si)3As Through the GaAs Process
Tiina Sarnet, Timo Hatanpa¨a¨, Mikko Laitinen, Timo Sajavaara, Kenichiro Mizohata,
Mikko Ritala and Markku Leskela¨ Submitted, (2015)
The author planned and conducted the deposition experiments and per-
formed most of the analysis on the films. T.H. synthesized the alkylsilyl
precursor. TOF-ERDA and RBS analysis was performed by M.L., T.S. and
K.M.. The author wrote the first draft, and finalized the paper together with
M.R. and M.L..
viii
List of abbreviations
AFM atomic force microscopy
AIST Ag−In−Sb−Te alloys
ALD atomic layer deposition
ALE atomic layer epitaxy
CD compact disc
CVD chemical vapor deposition
DRAM dynamic random access memory
DVD digital versatile disk
EDX energy dispersive x-ray analysis
FESEM field-emission scanning electron microscopy
FIB/SEM focused ion beam/scanning electron microscope
GIXRD grazing-incidence x-ray diffraction
GST Ge−Sb−Te alloys
HTXRD high-temperature x-ray diffraction
LED light-emitting diode
LEIS low-energy ion-scattering
MBE molecular beam epitaxy
ML monolayer
MOCVD metal-organic chemical vapor deposition
MOSFET metal-oxide-semiconductor field-effect transistor
MOVPE metal-organic vapor phase epitaxy
P power
PEALD plasma-enhanced atomic layer deposition
PGEC phonon glass, electron crystal
PLA pulsed laser ablation
R reflectivity
S Seebeck coefficient
T temperature
TEM transmission electron microscopy
TFEL thin film electroluminescent display
thd 2,2,6,6-tetramethyl-3,5-heptanedione
TMDC transition metal dichalcogenide
TOF-ERDA time-of-flight elastic recoil detection analysis
UHV ultra high vacuum
VLS vapor-liquid-solid
XRD x-ray diffraction
XRR x-ray reflection
ZT thermoelectric figure of merit
κ thermal conductivity
κe electronic component of thermal conductivity
κp phononic component of thermal conductivity
ρ electrical resistivity
σ electrical conductivity
ix
Contents
Abstract iv
Preface v
List of original publications vii
List of abbreviations ix
Contents x
1 Introduction 1
2 Introduction to chalcogenide and pnictide materials 3
2.1 Phase change materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3
2.2 Thermoelectric materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 10
2.3 Topological insulators . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17
2.4 III-V semiconductors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17
3 Atomic layer deposition of chalcogenides and pnictides 19
3.1 Principles of atomic layer deposition . . . . . . . . . . . . . . . . . . . . . . . . 20
3.2 Atomic layer deposition of chalcogenides and pnictides . . . . . . . . . . . . . . 23
x
3.2.1 Chalcogenides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 23
3.2.2 Pnictides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 24
3.3 Alkylsilyl compounds as precursors . . . . . . . . . . . . . . . . . . . . . . . . . 25
4 Experimental 29
4.1 Atomic layer deposition of chalcogenide and pnictide thin films . . . . . . . . . . 29
4.2 Chalcogenide and pnictide thin film characterization . . . . . . . . . . . . . . . . 30
5 Results and discussion 32
5.1 Properties of chalcogenide and pnictide ALD processes . . . . . . . . . . . . . . 33
5.1.1 Growth rate saturation . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33
5.1.2 Purge effects . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 35
5.1.3 Temperature dependence . . . . . . . . . . . . . . . . . . . . . . . . . . . 36
5.1.4 Film growth . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 37
5.1.5 Conformality . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 38
5.2 Properties of ALD chalcogenide and pnictide thin films . . . . . . . . . . . . . . 39
5.2.1 Composition . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 39
5.2.2 Crystallinity and morphology . . . . . . . . . . . . . . . . . . . . . . . . 41
5.2.3 Phase change properties . . . . . . . . . . . . . . . . . . . . . . . . . . . 44
5.2.4 Thermoelectric properties . . . . . . . . . . . . . . . . . . . . . . . . . . 45
6 Conclusions and outlook 47
References 49
xi
xii
Chapter 1
Introduction
Materials engineering or materials development has been an integral part in pushing mankind
forward. Learning how to make bronze and iron has furthered the human race by giant leaps.
In the modern era, many of the great technological improvements of the 20th century have
centered around silicon-based electronics. Our lives have been greatly transformed in just a few
decades. As Moore’s law is still valid to describe the development in the electronics industry,
and green thinking permeates all aspects of life, new and innovative solutions and materials are
needed to fulfill increasingly difficult demands.
As society becomes progressively digitalized, more and more storage capacity is needed to store
the great variety of data that are constantly produced. Memory these days needs to be fast,
non-volatile, durable, rewritable and cheap. One candidate for a new memory concept is phase
change memory. Phase change materials were discovered in the 1960s and entered the consumer
market in CDs and DVDs a few decades later. A more recent application is the phase change
random access memory. The data are stored as zeros and ones by the crystalline and amorphous
phase of the material. Phase change materials have a large contrast in electrical and optical
properties between the phases, making readout very simple. Rewriting is also easy by changing
the phase. Phase change materials are mostly chalcogenide and pnictide materials.
Energy is an important supporter of our way of life. Our society does not function without
electricity, and neither do smartphones and air-conditioners. Energy needs to be supplied to
the most remote corners of the Earth, and at the same time, excess heat is released from factories
and electronic devices without further utilization. Thermoelectric materials create electricity
from a temperature difference and can also use electricity to create a temperature difference.
Due to simple and robust device structures, thermoelectrics are used in various applications
1
from Voyager space probes to car-seat coolers. Many of the most used thermoelectric materials
belong to the class of chalcogenide and pnictide materials.
Although the material properties utilized in the aforementioned applications are different, the
compositions are not. Phase change materials and thermoelectrics are mainly chalcogenides and
pnictides, compounds formed by the elements of groups 15 and 16 such as selenides, tellurides,
antimonides and arsenides.
The motivation for scientific research can come from various sources; the researcher can, for
example focus on interesting scientific questions using curiosity as the main driver. Motivation
can also come from striving towards a desired end result. ALD research fits well into the middle
of these two extremes. Thin film research stems from fundamental chemical research in finding
new chemistries to deposit materials. In addition, clear goals and specifications can be inferred
from the applications.
Thin film research has become increasingly important for the microelectronics industry due to
the decrease in feature sizes and transformations towards three dimensional substrates. If the
developmental pace of the industry continues, new technologies, such as ALD, are needed to
face the challenges as much as new materials are required for the changing environment.
This thesis describes a process that started from a clear application, phase change memories,
and progressed to basic research in trying to determine the correct chemistry to fulfill the expec-
tations. Phase change materials, such as Ge2Sb2Te5, contain elements that were not common
in ALD, especially germanium and antimony. Additionally, telluride chemistries, and those of
chalcogenides in general, were quite simple and rare.
From this background, the alkylsilyl compounds were found, along with their efficient dehalosi-
lylation reactions [1]. This approach was extended to many non-metals, mainly the chalcogens
and pnictogens. Further research could then be extended to other industrially relevant materi-
als, such as thermoelectrics and III-Vs, with the same chemistry.
The structure of this thesis follows the progress of the research from the desired end results
to the chemistry that was needed to achieve them and to the results that followed. Chapter 2
describes in detail the materials classes that were the desired outcome, starting with phase
change materials that started this whole avenue of research. Chapter 3 describes the main
scientific method for this thesis, ALD, and the chemistry behind it. A brief summary of the
experimental procedures for depositing and analyzing thin films is given in Chapter 4. Selected
results of the alkylsilyl-enabled ALD processes are presented in Chapter 5. Conclusions from
the research and some further outlooks are discussed in Chapter 6.
2
Chapter 2
Introduction to chalcogenide and
pnictide materials
Chalcogens are elements from group 16 (6A), comprising sulfur, selenium and tellurium. When
discussing chalcogens, oxygen and polonium are usually disregarded. Pnictogens are elements
from group 15 (5A), and include nitrogen, phosphorus, arsenic and antimony. Chalcogenides
and pnictides are thus compounds formed by these elements.
Chalcogenide and pnictide compounds form an interesting whole, consisting of several indus-
trially relevant material categories. These include phase change materials, thermoelectric ma-
terials, topological insulators, along with II-VI and III-V semiconductors. What makes them
intriguing is the fact that the same compound can have properties from more than one material
group, and often does so. For example, Bi2Te3 and Sb2Te3 have phase change and thermoelectric
properties, and can also be characterized as topological insulators.
2.1 Phase change materials
The study of phase change materials began in the 1960s. In 1968, Stanford Ovshinsky proposed
a group of materials that could be repeatably converted from a high resistivity state to a low
resistivity state and back with the aid of an electric field. According to him, these materials
would mainly be amorphous intrinsic semiconductors, such as oxygen- and boron-based glasses,
and materials combining tellurium and arsenic with the elements from groups 13, 14 and 16.
Ovshinsky also purported that with some compositions, the aforementioned low resistivity
state would remain, even when the electric field is removed [2]. Even though the current phase
change material compositions are discussed later, the last condition stated by Ovshinsky is
3
vitally important. Phase change materials are indeed non-volatile memory materials, meaning
their information is not lost when electricity is no longer supplied.
There have been discussions on a universal memory that would be able to replace most currently
used memory technologies [3]. In addition to the already mentioned non-volatility, this memory
would need to pack large amounts of information into small spaces and store it for a long
time. The read, write and erase operations should be very fast, repeatable and managed with
little power. Finally, the whole construction should be inexpensive and compatible with silicon
technology [3, 4]. Although these requirements may sometimes seem contradictory, they serve
as a good basis, against which new accomplishments can be compared. The requirements also
present clear goals to strive towards.
As with any memory, distinct states are required to keep the information. In the case of phase
change materials, the zeros and ones for the memory devices are created by the amorphous and
crystalline phases of the materials, and the changes from one phase to the other. Therefore, the
material properties between the phases need to be different.
In general, although the amorphous phase does not have the long-range order of a crystal, a
short-range order can exist. This means that some order can be detected at the nearest-neighbor
and next-nearest-neighbor distances, like similar bond lengths, coordination numbers and bond
angles [5].
Due to these structural differences, there are two distinct material properties of the phase change
materials, which differ significantly from the amorphous to the crystalline state. Reflectance
values are higher for the crystalline than the amorphous materials. The difference can be in the
order of 20 % [6], depending on the alloy and wavelength [4]. The difference in resistivity between
the phases is even greater, a phenomenon already presented by Ovshinsky. This difference is
often several orders of magnitude.
Both of these differences in material properties are utilized in the memory devices. The reflec-
tivity difference has been utilized in rewritable CDs and DVDs already for decades [4]. The
more recent application of phase change random access memory uses the large difference in
resistivity, possibly enabling even multilevel storage [7] and logic operations [8]. The phase
change random access memories are already in production, in the 45 nm node [9], and used in
mobile phones, such as the Nokia Asha [10].
4
The changes in phase are brought about by laser or electrical pulses. The detection of the
amorphous or crystalline state is performed by a low intensity pulse that does not affect the
material. The amorphous material is crystallized by a moderate power pulse, while amorphiza-
tion is achieved by a shorter, more powerful pulse [11]. This is demonstrated in Figure 2.1. In
general, the change in phase is caused by laser or current induced joule heating [9], though
electronic excitations [12] might also have an influence. Phase change materials also exhibit a
phenomenon called threshold switching [13], which allows for less power being used for crystal-
lization [4]. When the phase change material is subjected to the threshold voltage or electric
field, a fast electronic transition enables a larger current to pass through the still amorphous
material which, in turn, heats the material and leads to its crystallization [4].
Figure 2.1: Differences in phase change material properties during crystallization and amor-
phization. P is the power of the laser/electrical pulse, R the reflectivity and ρ the resistivity of
the material.
The crystallization in itself is a thermodynamically favored process. The kinetic hindrance to
the crystallization, on the other hand, enables data-retention [9], which is a crucial property
for memory applications. The required storage time is 10 years at 100 or even at 150 ◦C, if
automotive applications are considered [11].
The quick crystallization is caused by the poor glass-forming properties of the phase change
materials. Glass-formers are inorganic substances that do not crystallize during melt-quenching,
but solidify as an amorphous phase. The glassy states have higher energy, entropy and volume
than their crystalline counterparts. Moreover, since good glasses do not crystallize easily [12],
poor glasses are thus needed.
Interestingly, the time needed for the crystallization is not constant. The initial crystallization
of the amorphous material is the slowest process [11]. When this material, or parts of it, are first
melt-quenched back to the amorphous phase and then recrystallized, this second and further
5
crystallizations progress much faster than the initial crystallization [14]. This can be explained,
for example, with the recrystallization starting from the interface with the surrounding crys-
talline material. Crystallization can thus occur without additional nucleation [15]. In addition,
the melt-quenched amorphous material can have much more medium range order than the
as-deposited amorphous material, making the recrystallization faster [16]. The recrystallization
process can occur even in less than 1 ns [17].
The faster recrystallization is especially true for the growth-dominated materials, where most
of the crystallization occurs via growing nuclei. The additional order also helps with nucleation-
dominant materials, where the crystalline volume grows mainly by newly forming nuclei [18, 19].
Some examples of phase change memory devices are presented in Figure 2.2. The first example
is an optical memory device in Figure 2.2a. The phase change material is only one of the
layers in the multilayer stack consisting, for example, of dielectric ZnS−SiO2, reflective Al and
the polycarbonate substrate [20]. The most simple electrical device structure is the so-called
mushroom cell in Figure 2.2b, where the phase change material is sandwiched between the
top and bottom electrodes and heaters. The crystallized region forms into the horizontal phase
change structure right above the heater. While this type of structure is good for testing, it does
not enable sufficient scaling. With scaling, also the reduction of the reset current is desired [21].
This could be achieved by decreasing the contact area between the phase change material and
the heater [22]. Overall, in more complex structures, the phase change material is no longer as a
two dimensional layer, but confined into a trench [22] or a pore [23] such as in Figure 2.2c. These
confined structures also result in better thermal insulation from neighboring structures. Such
complicated structures and shrinking devices require specialized thin film deposition methods.
(a) (b) (c)
Figure 2.2: Examples of typical phase change memory device structures. Single layer rewritable
DVD (a) and phase change random access memory configurations (b), (c). (a) Reprinted with
permission from [15]. Copyright (2010) American Chemical Society. (b) and (c) Reprinted
(adapted) from [24] with permission from Elsevier, Copyright (2006).
6
Once the device is in operation, it needs to endure the rewriting process and be reversibly
switchable. To compete with dynamic random access memory (DRAM), 1016 − 1018 cycles are
required. Currently, at least 1011 cycles have been proven [9].
Most of the materials studied today are in the ternary composition range of Ge-Sb-Te (GST),
demonstrated in Figure 2.3, along with their common applications. The binaries GeTe and
Sb2Te3 mix well and form a GST alloy. Numerous studies have been made on the compositions
that lie in the tie-line between GeTe and Sb2Te3, making Ge2Sb2Te5 the “classical” phase change
material. The material composition can be varied since the GST alloys have stable compositions
along the whole tie-line. Another important family of phase change materials is the Ag and
In doped Sb2Te, or AIST, family. In addition, Ge-doped Sb has proven to have phase change
properties. Overall compositions in these alloys are tailorable to specific applications [4].
Figure 2.3: Phase diagram of phase change materials.
When choosing the right composition, also the crystallization properties need to be taken into
account. The AIST family crystallization is growth dominated, while that of the GST family
is nucleation controlled [18, 19].
Good reflectance properties at desired wavelengths are needed from materials in optical appli-
cations. Therefore, as an example, with changes in the laser wavelengths from IR to red and
further to blue, compositions in the GeTe-Sb2Te3 tie line moved towards GeTe, as they have
a higher contrast in the shorter wavelength, i.e. blue, region of the visible spectrum [4]. The
changes in optical properties enable compositional tailoring for a specific application.
7
There have been great difficulties in determining the exact difference in structure between the
amorphous and crystalline phases. The crystalline structure is naturally easier to determine,
for example with x-ray diffraction (XRD). An amorphous material, by definition, does not have
long range order. In the case of phase change materials, however, there have been indications of
short range order with neighboring atoms, similar to crystalline materials. Pathways, such as the
“umbrella-flip” of the germanium atom between a tetrahedral coordination in the amorphous
state and an octahedral coordination in the crystalline state, have been suggested for the phase
transition [25]. Nevertheless, it seems that Ge atoms in amorphous form also have mostly
octahedral coordination. Therefore, it could be stated that similar bonding conditions exist in
both forms, and the transitions within the structure are quite minor [26].
There are some factors that the crystalline structures of phase change materials have in common.
These could be used as a basis for finding new materials that fit into this category. Phase change
materials exhibit resonant bonding. As an example, Sb has three valence p-electrons to form
orthogonal bonds. In its crystalline structure, Sb has six roughly equivalent nearest-neighbors to
form bonds with, giving rise to two limiting possibilities of bond structures. Simplified versions
of this are presented on the left and right sides of Figure 2.4. In reality, the hybrid form of these
exists (Fig. 2.4, middle), making the bond structure resonant, even with some Peierls distortions
in the crystal structure [27]. Overall, the hybridized structures are fairly symmetric, or close to
the octahedral-like coordination [28]. The resonance bonding is not present in the amorphous
phase, thus presenting one reason for the large contrast in electrical properties between the
phases [26].
Figure 2.4: Simplified resonance structure of Sb, with the extremes on the sides and the hybrid
structure in the middle. Reprinted by permission from Macmillan Publishers Ltd [27], Copyright
(2008).
In many of the phase change materials, these same three p-electrons per atom are present.
When we also take into account the ionicity of the bonds, denoted by the difference in valence
radii of the p-orbitals [28], a plot such as Figure 2.5 [26] can be made. It is abundantly clear that
all known phase change materials are located in a very small area of the diagram, indicating
that a very specific bonding character is needed. A curious outcome of this type of plot of
8
chalcogenide materials is the overlap of the material classes. The same combination of ionicity
and hybridization of phase change materials is also found in known thermoelectic materials
and topological insulators. These types of theoretical studies are important as they open whole
new possibilities in using the materials at hand, and provides a good tool for discovering new
materials within the desired property ranges.
Figure 2.5: Map of materials with, on average, three p-electrons per site. Reprinted with per-
mission from [26], Copyright (2012) John Wiley and Sons.
In general, phase change materials can be made by both physical vapor deposition (PVD) and
chemical vapor deposition (CVD) methods. In addition, chemical synthesis of nanoparticles is
also an option [29, 30]. Thin films of change materials are mostly made by sputtering, against
which all the other methods, and the material properties produced by them, can be compared.
All PVD methods use high or ultra high vacuum (UHV) conditions. In PVD, material is removed
from a solid source or target by supplying energy in the form of heat, molecular bombardment or
a laser beam. Atoms or molecules traverse in direct paths through the vacuum to the substrate.
These direct paths of the atoms make PVD a line-of-sight method, meaning that only the top
surface of the substrates can be coated. In addition to sputtering, phase change materials have
been made by simple evaporation [31, 32], pulsed laser ablation (PLA) [33, 34], molecular beam
epitaxy (MBE) [35, 36]. Nanowires [37] have been made using the vapor-liquid-solid (VLS) [38]
method, along with nanoparticles created by PLA [39, 40].
9
CVD methods utilize chemical reactions to form the products. In a typical process, gaseous
precursors are brought into the reaction chamber and to the substrate surface, where they
adsorb and react, forming the desired end product along with a number of volatile side prod-
ucts [41]. Often these precursors are metal-organic, making MOCVD (metal-organic CVD) the
most common type of CVD. Unlike in PVD, CVD processes enable depositions on three dimen-
sional substrates. Phase change materials have been deposited by both thermal CVD [42–45]
and by plasma-assisted CVD [46, 47].
2.2 Thermoelectric materials
The thermoelectric phenomenon is mainly based on two effects. According to the Seebeck effect,
a temperature gradient across a material or a device produces a voltage. The complementing
Peltier effect causes a temperature difference across a material or a device when a current is
passed through it [48, 49]. Examples of the devices are depicted in Figure 2.6. The devices
are connected electrically in series with conducting strips. Thermal conduction is ensured with
electrically insulating but thermally conducting plates to connect the separate p/n-pairs, i.e.
the legs, into a module [49].
(a) (b)
Figure 2.6: Schematics of a Seebeck device (a) and a Peltier device (b).
In general, thermoelectric materials have a high specific power (watts per volume or mass),
which enables scaling to low power levels. They are very useful in applications where low
weight and small size are more essential than the production costs. In addition, thermoelectric
devices have no moving parts and thus need no maintenance, which makes decade-long operation
possible.
10
The Seebeck effect is used in thermoelectric power generation, such as on the Apollo Lunar
missions, Voyager space probes, and in many hard-to-access locations on Earth where constant
power and low maintenance are required [49]. The same effect can also be utilized in waste
heat utilization, in factories for example. The Peltier effect is used in cooling applications, be
it a simple USB-plugged beverage container [50], a car-seat temperature-control system [51], or
more sophisticated devices such as lasers and infra-red sensors [52, 53].
The efficiency of thermoelectric materials is usually determined by a dimensionless figure of
merit, ZT (Equation 2.1), where S is the Seebeck coefficient, σ is the electrical conductivity, κ
is the thermal conductivity and T is the absolute temperature. Therefore, a good thermoelectric
material or device needs to have a high Seebeck coefficient to achieve a larger voltage, a high
electrical conductivity to minimize the parasitic or joule heating of the cold side, and a low
thermal conductivity to maintain the temperature difference as high as possible [54].
ZT =
S2σ
κ
T (2.1)
The difficulty lies in fulfilling these seemingly simple terms in the same material. Although
there is no fundamental upper limit for ZT [49], there is interdependency between the factors
affecting ZT through the electronic properties of the material. When the Seebeck coefficient
increases, the electrical conductivity decreases due to carrier density effects. Moreover, when the
electrical conductivity increases, thermal conductivity increases as well due to the Wiedemann-
Franz law [55]. This can make it difficult to achieve ZT values greater than 1 in bulk or 3D
materials [55, 56]. In addition, the thermal conductivity κ consists of electronic and phononic
components, which is demonstrated in Equation 2.2.
κ = κe + κp (2.2)
The focus is often at decreasing the phononic component of the thermal conductivity as the the
electronic component is related to the electric conductivity according to the Wiedemann-Franz
law [55, 57]. The more the lattice vibrations can be disturbed, the better. An all-encompassing
approach has been proposed to achieve all this. The aim is a phonon-glass electron-crystal
(PGEC) structure: the electronic properties of a crystalline solid (high S and σ), where the
electron mean free paths are long, and the thermal properties of a glass (low κp), where the
11
phonon mean free path is short. These materials could be found amongst heavy-element com-
pound semiconductors that have small band gaps and complex crystal structures [58].
Up until the 1990s, thermoelectric materials were made mainly in bulk form, and their ZT values
were at the highest close to unity [52, 59]. A new possible approach was discovered in 1993 [60,
61], suggesting that significant increases in ZT values were possible by using nanostructured
materials. This enabled engineering thermoelectric materials from the start, not just using
the natural materials available. The so-called second generation materials were already able
to achieve ZT = 1.3 − 1.7 by decreasing the lattice thermal conductivity by using mixed
compositions and nanoscale precipitates. The current third generation materials are reaching
values of ZT = 1.8− 2.2 by optimizing band structures and suppressing phonons on all length
scales [62]. Nanostructuring the devices gives an opening for thin film deposition methods, such
as ALD. General trends in ZT values for different materials can be found in Figure 2.7.
(a) (b)
Figure 2.7: ZT values of p-type (a) and n-type (b) materials in terms of absolute temperature T.
Reprinted (adapted) with permission from [63], Copyright (2009) Cambridge University Press.
The nanostructuring can be done in different length scales, since the lattice vibrations, i.e.
phonons, have different wavelengths or mean free paths. Complex crystal structures increase
the number of phonon modes. Each of the phonon modes can then be handled separately. The
shortest wavelength phonons are scattered by the shortest length scale defects, such as point
defects and dopant atoms on lattice sites. These include both vacancies and mass contrast
caused by heavier or lighter dopant atoms. For example, heavy atoms within the structure, such
as within voids, can vibrate independently causing disturbances in the phonon propagation. The
medium wavelength phonons can be scattered by basic nanostructuring, such as the interfaces
between the matrix and the precipitate. Mass contrast can also be created between the different
12
phases. The long wavelength phonons are accordingly scattered by the largest scale defects in
the material. This can be achieved via grain boundaries by breaking down the crystals to
the mesoscopic scale (100 nm - 5 µm). Finally, the material can be sintered back into a bulk
object [62, 64].
Nanostructuring can be combined simultaneously with hierarchical architecture, band align-
ment and other electrical structure engineering. Zhao et al. [62] called this all-encompassing
engineering the “panoscopic approach”. The band engineering is performed by having the va-
lence and conduction band energies of the two or more phases close to each other to ensure
good transport properties.
Thermoelectric materials include a myriad of materials, only some of which are mentioned here.
Those known for decades are Bi2Te3, PbTe and SiGe alloys. They are used at low, intermedi-
ate and high temperatures, respectively, due to the varying temperature dependence of their
ZT values. It needs to be taken into account, however, that all the mentioned thermoelectric
materials are doped to optimize their properties, even if not always explicitly noted.
The bismuth telluride alloys have their maximum ZT levels a little above room temperature,
at about 350− 450 K. Bi2Te3 and its alloys are therefore the main sources for refridgeration
and also energy generation at low temperatures. Sb2Te3 is often alloyed with Bi2Te3 to increase
ZT by decrasing thermal conductivity and leaving the electronic properties mostly unchanged.
This type of an alloy leads to a p-type composition, often near (Sb0.8Bi0.2)2Te3. In a similar
fashion, an n-type alloy is formed by alloying Bi2Te3 with Bi2Se3, to compositions such as
Bi2(Te0.8Se0.2)3 [49, 57, 65].
Lead telluride and its alloys were one of the first thermoelectric materials discovered. They have
a cubic NaCl-type crystal structure and excellent thermoelectric properties. They are generally
used in the temperature range of 500− 900 K, where intrinsic ZT values as high as 1.4 can
be reached. A PbTe−PbSe has even achieved ZT = 1.8. Lead tellurides, or selenides, are also
doped, often with halides and alkali metals. Further ZT improvement can probably be achieved
through nanostructuring [66]. Good performance has also been found in materials similar to
lead tellurides, such as the TAGS, (AgSbTe2)1-x(GeTe)x, or LAST, (AgSbTe2)x(PbTe)1-x, where
values of ZT > 1.5 can be reached [49, 57, 67].
The complete miscibility of Si and Ge ensures a large variety of possible alloys such as the
common Si0.8Ge0.2. Although the miscibility is good, care needs to be taken during preparation,
as Si- and Ge-rich regions form easily. The SiGe alloys work best at very high temperatures,
around 900− 1300 K, although they can be operated at much lower temperatures as well. A
13
wide temperature range can mean very efficient operation. For thermoelectric devices, the SiGe
alloys are doped with boron for p-type and phosphorous for n-type materials [68].
The demands for higher ZT values and the onset of multileveled approaches to nanostructuring
have brought along several new materials classes into the family of thermoelectrics. Many of
these are also chalcogenides and pnictides. In general, the recent additions to thermoelectrics
strive for PGEC properties with complex crystalline structures.
Half-Heuslers are a group of semiconducting materials made of metallic elements. The XYZ
composition is made of an electropositive metal X (Ti, Zr, Hf), a late transition metal Y (Co, Ni)
and a main group element Z (Sn, Sb). The crystal structure can be thought as an fcc Z lattice,
with metal X in the octahedral holes and metal Y in half of the tetrahedral holes, leading to an
XZ rocksalt sublattice and a YZ zincblende sublattice. In general, these half-Heusler materials
have large power factors (S2σ), although combined with a large κ. By including several of the
possible metals in the same structure, point defects can be formed, leading to a decrease in κ.
Therefore, ZT values of ≥ 1 have been achieved at temperatures above 750 K, when practically
all the aforementioned metals were used together in an n-type composition [56].
Skutterudites are a class of compounds with a CoAs3 type structure. They can be generalized
with the formula MX3, where M is a transition metal (Co, Rh, Ir) and X is the pnictogen (P,
As, Sb). Skutterudites have an open body-central-cubic crystal structure, with a void in the
center of the unit cell. The pnictogens form planar σ-bonded rings within the structure. The
binary skutterudites have high mobilities but also a high thermal conductivity, which makes
them unsuitable for thermoelectrics as such. When the structural void is filled with atoms with
a large mass and a small diameter, such as rare-earth or alkaline-earth metals, the thermal
conductivity is reduced and thermoelectric applicability is achieved. Not all voids need to be
filled and can be filled with different elements at the same time. Depending on the inserted
atoms, both p- and n-type materials can be achieved. The inserted elements cause a rattling
effect, where the heavy atom vibrates independently from the rest of the crystal structure,
which significantly decreases thermal conductivity, enabling ZT values of over 1.5 in the range
of 700− 1000 K [53, 69, 70].
Zintl compounds are represented with the nominal formula AM2X2, where A is an alkaline-earth
or rare-earth metal, M is a transition metal or a main group metal, and X is an element from
groups 14 and 15. They are usually air sensitive, with a complex crystalline structure such as
ThCr2Si2. The large unit cell and the combination of covalent and ionic bonds fulfill the PGEC
criteria of low thermal and high electrical conductivity. [53, 57]
14
There are still several other material groups that have thermoelectric properties and do not
contain chalcogen or pnictogen elements. Oxides, such as cobaltites and perovskites are studied
because they are chemically and thermally stable even in high temperature gradient conditions
and in air. In addition, they can be easily chemically modified and compositionally tailored.
Low costs and environmental friendliness are also desirable properties. At best, oxides have
reached ZT values of 0.65 [71]. Another group of inorganic materials that has demonstrated
thermoelectric properties is metal silicides, which can be used at high temperatures. The best
ZT values around unity have been achieved with magnesium and manganese silicides [49, 68].
All the above described materials are inorganic. Nevertheless, organic materials have been
proven to have thermoelectric properties as well. They would be an alternative to the inorganic
materials due to elemental abundance, low cost, chemical modifiability and mechanical flexi-
bility. Organic thermoelectric materials are mainly conducting pi-conjugated polymers such as
polyaniline and polyacetylene. Their low intrinsic ZT values have been sometimes compensated
with hybrid structures with more familiar inorganic thermoelectric materials. Overall, a good
deal of work is still to be done here. Especially the thermal stability is an issue, as polymers
do not withstand high temperatures [72].
Thermoelectric materials are a varied group and cannot be categorized simply. They contain
many elements in increasingly complex structures. Material structure engineering has led to
various compounds having ZT values in excess of their bulk values presented previously, as can
be seen in Figure 2.8. Nevertheless, high ZT values do not lead to an instant success in real
world applications. To achieve more than a niche market, mass production is needed, where cost
is always an issue. Therefore, with elements like antimony, the noble metals and the rare earths,
there can be a market risk of increasing costs due to scarcity and a small number of production
sites. Thus, thermoelectric materials need to be evaluated not only on their properties but also
on their applicability to production [73].
Actual thermoelectric devices, especially the legs (see Figure 2.6) made of the n- and p-type
thermoelectric materials, are still macroscopic in size to accommodate large temperature differ-
ences across the devices. Therefore, also the methods to manufacture these materials are bulk
methods.
The initial thermoelectric material is often made with powder metallurgical tehniques or by
mechanical alloying [49]. Other possibilities include melt-spinning [64] or melting from ele-
ments [74].
15
Figure 2.8: State-of-the-art ZT values for different materials. Reproduced from [62] with per-
mission from The Royal Society of Chemistry, Copyright (2014).
Since nanostructuring is practically a requirement these days, the starting material can also be
made in the form of nanoparticles. These can be made via wet chemical methods, hydrothermal
growth or inert-gas condensation [53]. If the thermoelectric material is made in bulk form, it is
usually ball-milled into an appropriate grain size.
When the thermoelectric material is in a powder form, it needs to be compacted into the final
form of a leg in a device, or part thereof. The compacting of the powders is often achieved
via sintering methods, such as spark plasma sintering and current-assisted sintering [64]. Other
compacting methods include wet chemical methods, hydrothermal growth, and inert-gas con-
densation [53]
As all of the thermoelectric materials have their optimal temperature ranges where they perform
the best, segmenting has become a tool. Several thermoelectric materials can be combined in the
segments of a device leg so that each one is functioning in its optimum temperature range [57].
As these segments need to be bonded together, the preparation methods need to be well chosen
to achieve a good contact between the materials.
16
2.3 Topological insulators
Topological insulators are a curious group within the chalcogenide materials. Their existence
was theorized a number of years ago [75], after which they gained great scientific interest.
The theoretical predictions were confirmed with experimentally measured results very soon
afterwards [76, 77].
As the name implies, topological insulators are insulating in the bulk. The intriguing properties
arise from the surface, where metallic conducting states can be found. This is made possible
by the time-reversal symmetry of the conducting states [78–80], which is typical for heavier
elements [79–81].
It appears that topological insulators are mostly chalcogenides with layered crystal structures.
For example, several of the proven compounds, such as Sb2Te3 [78], Bi2Te3 [82] and Bi2Se3 [77]
all have the same crystal structure.
Spintronics and quantum computing have been envisioned as possible applications for the topo-
logical insulators [82, 83]. This again implies that the materials need to made in non-bulk form.
A good way to make a surface-reliant material would be a surface-controlled method such as
ALD.
2.4 III-V semiconductors
As the name implies, III-V semiconductors are comprised from the elements of groups 13 (III)
and 15 (V). Examples of these compounds are GaAs and InP. Not naturally occurring, this
category of materials became an intense subject of study already in the 1950s [84, 85]. After
this, their interesting electric properties were enough to fuel the research efforts.
Many III-V compounds have their band gaps close to the visible range, making them ideal
candidates for optical applications. As an example, one of the most studied III-V compounds,
GaAs, has a direct band gap of 1.4 eV, which puts it in the infrared region, while the band
gaps of others such as AlN are in the ultraviolet region [86]. Many industrially important III-
V semiconductors are ternaries or even quaternaries, such as AlxGa1-xN and AlxGayIn1-x-yAs,
which means their properties, like band gaps, can be adjusted by compositional variations.
Consequently, tailor-made materials can be made for the applications [87].
17
It is the direct band gap that enables an efficient radiative recombination of electrons and holes,
which leads to photon emission. A very simple LED can be achieved with a p-n junction in
GaAs [88, 89]. When the alloys mentioned above are combined in multilayer stacks or quantum
wells and other complicated structures, a wide range of wavelengths can be covered [87]. The
blue LEDs based on GaN [90] have been the latest addition to LED wavelengths, thus giving
rise to significantly more research on white LEDs, for example. The same III-V direct band gaps
and optical properties derived thereof, also enable lasers [91–93] with similar heterostructures
and quantum wells. Another optoelectronic application of III-V compounds is solar cells [94].
Due to the high carrier mobility, compared to Si in particular, GaAs has been an attractive
candidate for microelectronics, as transistors could be made to work faster this way [95]. InAs
has an even higher electron velocity [87]. The special band structure of GaAs enables negative
resistance effects [96, 97], which can be utilized in microwave production such as in a Gunn
diode [87, 98]. Electronic components usually require heterostructures, which can be easily
formed as GaAs has a small lattice mismatch with many of its alloys [87].
As III-V compounds are needed for their electrical properties, low impurity levels are crucial.
Therefore, single crystal methods have been popular for making them. Single crystals can be
grown, for example, with the liquid encapsulated Czochralski [99, 100] and different Bridgman
and gradient freeze methods [101–103].
In optoelectronics applications, thin films are also needed. These can be made, for example, with
liquid phase epitaxy [87, 104]. In addition, there are physical vapor deposition methods such
as evaporation [105] and molecular beam epitaxy [104, 106, 107]. Chemical vapor deposition
(CVD) methods include metal-organic CVD (MOCVD) [108], also known as metal-organic
vapor phase epitaxy (MOVPE), and atomic layer epitaxy (ALE) [109–111].
Generally, III-V thin films have been deposited on III-V substrates, either on the same material
or on something with a small lattice mismatch. Compositionally different materials, such as
sapphire and silicon, have also been used as the interface layer is important for high crystal
quality. Overall, the deposition temperatures have been relatively high, in the 400− 700 ◦C
range. Some materials, such as GaN, require even higher temperatures of over 1000 ◦C [87].
18
Chapter 3
Atomic layer deposition of
chalcogenides and pnictides
Atomic layer deposition was initially developed for depositing thin, pinhole-free, conformal films
for thin film electroluminescent displays (TFEL). The first process was also a chalcogenide,
the phosphor material in TFELs, ZnS:Mn [112]. Soon thereafter, another ALD process for a
chalcogenide, this time ZnTe, was presented [113]. Since that, ALD has proven itself in many
different nanotechnological applications such as microelectronics [114].
Although TFELs are still in production, the main applications of ALD these days are in the
semiconductor industry. High-k materials are needed in MOSFET and DRAM device structures
as well as in double-patterning lithography. Other applications include the magnetic read/write
heads for hard disks and protective coatings for silver jewelry to prevent tarnishing during
storage [115].
The reason for taking ALD into use in microelectronics and many other applications is the
miniaturization of devices, which increased the requirements for deposition controllability.
Miniaturization has also led to increasingly complex three-dimensional devices such as tri-
gate and FinFET transistors. For most PVD methods, thin film deposition outside line-of-sight
is difficult. Therefore, chemical methods, such as ALD, are needed for conformal, repeatable,
and good quality film deposition.
A large number of materials have been deposited by ALD during the 40 years the method
has been used [116]. Both chalcogenides and pnictides have been the focus of intense study,
and thus several chemistries have been used to deposit these thin films, both epitaxially and
non-epitaxially.
19
3.1 Principles of atomic layer deposition
Atomic layer deposition (ALD) is a gas phase thin film deposition method, a modification of
chemical vapor deposition (CVD). Unlike in CVD where precursors can react already in the
gas phase before reaching the substrate, in ALD the precursors are brought to the substrate
one at a time. A simple ALD cycle is demonstrated in Figure 3.1.
Figure 3.1: Representation of an ALD cycle.
The precursor molecules react with surface groups until all available sites are occupied, making
ALD a saturative method. It is thus self-limiting, as growth stops when all the surface groups
have reacted. A purge separates the precursors from each other in space or in time, removing
excess precursor molecules and reaction byproducts. The second precursor can then be brought
to the substrate surface where a new set of saturative reactions occur. Another purge completes
the basic ALD cycle.
When each ALD cycle deposits the same amount of material - a fraction of or a full monolayer of
film - the desired film thickness can be achieved simply by controlling the number of deposition
cycles. The saturative reactions enable uniform thin film deposition on even the most demanding
three-dimensional and large area substrates [117, 118].
Generally, ALD processes are not strongly temperature dependent. There can be a wide “ALD
window”, where the growth rate stays fairly constant [116]. Some temperature variations in
growth are shown in Figure 3.2. Part 3 represents the stable ALD window, which does or does
20
not exist for a process. A process can still have all the typical ALD characteristics, even if there
is no ALD window. In part 3A, the growth rate is temperature independent, while in 3B the
growth rate is temperature dependent. This difference can be caused by a possible temperature
dependence of the density of the reactive sites [117].
Figure 3.2: ALD process properties with respect to the deposition temperature.
The increase in growth rate in part 1 can be due to multilayer adsorption of the precursors
or condensation of low vapor-pressure precursors. The effect becomes more prominent with
decreasing temperatures, because the volatility continues to decrease. In part 2, the growth
rate decreases with decreasing temperature. This might be due to kinetics slowing down the
reactions so that they are not completed. With the temperature increase, the system has more
energy, which can speed up the reactions towards completion [117].
The growth rate increase in part 4 means that the self-limiting conditions no longer apply due
to precursor decomposition. This causes more reacting species to be involved in the growth
reactions. These unwanted species might interfere with the intended process by competing or
additional reactions. The self-limiting conditions can be lost also due to precursor desorption,
which causes the growth rate decrease in part 5. If the precursor no longer sticks to the surface,
less reactions occur and less material is deposited [117].
This basic form of ALD is also called thermal ALD, since the activation of the chemical reactions
is achieved through heat. Other forms, usually with lower deposition temperatures, utilize
plasma (plasma-enhanced ALD, PE-ALD) [119] or light (photo-assisted ALD) [120] as sources
of energy. The energetic species created by the plasma, such as ions or radicals, add energy to
the system [121]. Recent developments in ALD, and especially in reactor design, have added
some old and new variations to the method. Spatial ALD, where the precursors are separated
in space and the substrate moved between them, has made a return into the ALD toolbox,
21
having been the initial concept in Suntola’s work [112]. Spatial ALD has been developed so
far now that some robust ALD processes are used to deposit thin films on very long substrate
surfaces traveling from roll-to-roll [115].
For any ALD process to work properly, certain important properties are required from the
precursors. The most obvious property is sufficient volatility, as ALD is a gas phase method.
The precursor should have a vapor pressure of 0.1− 1 mbar [115] at a reasonable temperature
below the deposition temperature so that condensation before reaching the substrate is avoided.
Next, the precursor needs to be highly reactive with its counterpart, yet not reactive enough to
etch or otherwise destroy the growing film and the reactor parts. The byproducts also need to
be volatile to be easily removed by the purge. In addition, as thermal stability is needed, the
precursor should not decompose at the deposition temperature. [115–117]
Several other factors should also be taken into account when considering precursors. Toxicity
to the user is one issue. Many precursors are relatively safe to handle with minor precautions,
yet some require special attention [115, 117]. For example, H2O is perfectly safe, and H2S can
be used rather safely from an external gas bottle even in a regular laboratory. Nevertheless,
any gaseous pressurized precursors need to be handled carefully. When progressing down in
the periodic table, H2Se and H2Te are very toxic already in small concentrations where they
cannot be smelled, unlike H2S. Therefore many more precautions would be needed to use these
precursors [122].
Environmental friendliness is another issue, as waste is always created during an ALD process,
particularly through the unreacted precursors and byproducts that are purged away during a
deposition cycle. One consideration is cost. While in the scale of research laboratories, proof-
of-concept studies can be made with very little precursor, industrial applications might not be
feasible due to costs. [115–117]
ALD precursors encompass a wide variety of different types of compounds. Metal precursors
are usually inorganic, metalorganic or organometallic. Inorganic precursors are simple - either
elements themselves or halides. Common metalorganic precursors can be organometallic and
have direct carbon-metal bonds with ligands, such as alkyls and cyclopentadienyls, or they
can be plain metalorganic and have ligands bonded through a non-carbon such as alkoxides,
β-diketonates, amides, imides, amidinates and phosphines. The most common non-metal pre-
cursors are hydrides such as H2O and NH3 [115–117].
22
3.2 Atomic layer deposition of chalcogenides and pnic-
tides
The assortment of ALD processes has been growing ever since the early days of Suntola. Both
scientific curiosity and the needs coming from industrial applications have driven the research.
Nevertheless, a number of potentially industrially relevant chalcogenide and pnictide materials
did not have ALD processes as little as a decade ago. This group included the already mentioned
material categories of phase change materials and thermoelectric materials.
3.2.1 Chalcogenides
Zinc sulfide was one of the early ALD success stories, and there is a substantial list of sulfides
already made by ALD [116]. Most of the reported sulfide materials can be made by com-
bining a metal halide [123–126] or a metal-thd (thd = 2,2,6,6-tetramethyl-3,5-heptanedione)
complex [123, 127, 128], with H2S, the most common sulfur precursor. In some cases, such
as ZnS [112, 129] and CdS [130, 131], the elements can be used for the metal and non-metal
precursors.
Only two selenide materials, ZnSe and CdSe, have been reported prior to introducing alkylsilyls
selenides as precursors. Nevertheless, there are a number of different processes reported for ZnSe
and CdSe. As with the sulfides, elements can be used to form both selenides [131–134].
There were few reports of telluride ALD until the recent alkylsilyl-enabled processes. The
simplest of these processes used the respective elements as precursors, for example in MgTe [135,
136], MnTe [136–139], ZnTe [132, 138] and CdTe [131, 136–139].
When elements cannot be used for depositing selenides, H2Se becomes the precursor of choice.
It is most often combined with a zinc alkyl, either dimethyl [140, 141] or diethyl zinc [142],
although ZnSe can also be deposited from elemental Zn and H2Se [143]. Both dimethyl zinc -
H2Se processes used low deposition temperatures, with the ALD window at around 150
◦C, and
growth rates saturating to 1 ML/cycle. The diethyl zinc - H2Se process, or a silylamido zinc
- H2Se process, needed deposition temperatures a few hundred degrees higher [142]. Similarly,
the use of in situ created ZnCl2 as the zinc precursor led to high deposition temperatures, with
the ALD window as high as 400− 500 ◦C. All of these processes used H2 as the carrier and
purge gas. There has been only one instance where CdSe was mentioned being prepared from
non-elemental sources, by combining dimethyl cadmium with H2Se [144].
23
Differently from the selenides, alkyl tellurides have been utilized in most of the reported ALD
telluride processes. On each instance these have been combined with alkyl compounds of the
metals. Also with tellurides, the carrier and purge gas used was H2. The most common telluride
material to be produced is CdTe, while HgTe [145, 146] and ZnTe [147] are barely mentioned.
To deposit CdTe, Me2Cd is combined with
iPr2Te [148, 149], Et2Te [147, 150] or MeAyTe [145,
147, 150], where Ay is the allyl ligand. Most CdTe processes have their optimum deposition
temperatures around 300 ◦C, while the growth rates saturate to 1 ML/cycle.
3.2.2 Pnictides
Pnictides, and especially the III-V semiconductors have been industrially important materials
for several decades already. The nitrides are the most abundant group of ALD pnictide pro-
cesses. In nitride thermal ALD, metal halides [151, 152], alkylamides [152] and alkyls [116] have
been most commonly combined with NH3 [116, 152] for film formation.
The next group is the phosphides, of which only AlP, GaP and InP have been presented. These
processes mostly utilize toxic PH3 as a precursor [153–155]. The metal precursors have also
been fairly simple, including mainly alkyls [154, 156] and halides [153].
The only non-alkylsilyl-enabled ALD process for an antimonide, InSb, used the elements as
precursors [157]. The ALE InSb served as an interface layer for further MBE growth. The
best overall results were achieved when this interfacial layer was deposited at 300 ◦C with 85
atomic layers. This temperature is low for an ALE process [110], making the crystalline quality
somewhat doubtful after only the ALE growth. The temperature was raised for the subsequent
MBE, which should improve the crystallinity. The InSb growth was initially three dimensional,
although the islands coalesced leading to two dimensional growth already after 40 cycles. The
ALE InSb interfacial layer improved the crystallinity, optical and electrical properties of the
5 µm InSb film deposited by MBE [157].
As III-V semiconductors have industrially important properties, arsenides, in particular, have
been the focus of intense study. In general, arsenides have been formed by ALE/ALD using
arsine as the arsenic precursor. Compounds were generally used for group III precursors, as the
elements have a low vapor pressure [110]. Thus the metal precursors have been fairly varied,
including halides [158, 159], alkyls [160–162], amines [163] and cyclopentadienyls [164]. Other
arsenic sources have been aminoarsine [164, 165] and tert-butylarsine [166, 167].
24
A good example is the quintessential III-V compound GaAs, which became the focus of intense
study in the 1980s [160, 168] when the deposition method was still called ALE, atomic layer
epitaxy. Though there were numerous studies on the subject [169], the epitaxial films were de-
posited with a limited number of precursor combinations, of which the combination of trimethyl
gallium and arsine was the most popular. Hydrides in general have proven themselves as very
good precursors, although in this case, precautions need to be taken. Arsine is very toxic and
as a gas requires good safety measures.
Other alkyl gallium precursors [170, 171] were also used, along with gallium halides [158, 172,
173]. In addition, as mentioned earlier, alkyl compounds of arsine [174, 175] were also used
as arsenic precursors. Practically all of these processes had growth rates that saturated to 1
ML/cycle, making the processes very stable.
The consistent problem with ALE GaAs processes was carbon contamination. Both trimethyl
and triethyl gallium are thermally unstable and can crack before reaching the substrate or on the
substrate [110]. Therefore, the gallium arsenide films were consistently p-type, if preventative
measures were not taken to prevent this carbon contamination. Arsine doses needed to be
sufficiently large and the gallium alkyl pulses short to achieve higher purity [111]
These days, most carrier gases used in ALD are inert. Curiously, in the GaAs studies, hydrogen
was most often used as either the transport gas or even as a third precursor [176, 177]. Laser-
assisted ALE was also used to ensure complete reactions on the surface. [178] Furthermore,
the temperature range of these processes was limited to high temperatures, in the order of
300− 500 ◦C. This temperature could be too high for some applications and more sensitive
substrates.
3.3 Alkylsilyl compounds as precursors
Alkylsilyl compounds are a recent addition to the ALD precursor toolbox [1]. They can be
described with a general formula (R3Si)xNm, where R is an alkyl group and Nm is a chalcogen
or pnictogen. The number x is determined by the non-metal as chalcogens require two and pnic-
togens three ligands. An example of a crystalline structure of an alkylsilyl tellurium compound,
(tBuMe2Si)2Te, can be found in Figure 3.3.
25
Figure 3.3: Molecular structure of (tBuMe2Si)2Te. Reprinted with permission from [1]. Copy-
right (2009) American Chemical Society.
Alkylsilyl related chemistry is not new as such. For example, (Me3Si)3As was combined with
gallium and indium halides to form GaAs and InAs in the liquid phase already in 1989 [179]. The
dehalosilylation reaction has also been utilized in the synthesis of other III-V materials [180],
such as GaSb nanoparticles [181]. Even single-source precursors have been synthesized [182,
183].
The initial applicability of a precursor for ALD can be determined by thermo-gravimetric
analysis (TGA). As the temperature of the sample is raised, the precursor weight loss should
proceed in a single step. This would mean a clean evaporation and thus good applicability as a
precursor. If the final residue is also close to 0, the thermal stability of the compound, another
ALD requirement, is good. However, as the TGA method is based on the weight loss of the
solid or liquid form, no further information about the gaseous product can be deduced [122].
TGA curves for several alkylsilyl chalcogenides and pnictides can be found in Figure 3.4.
All of the presented compounds show a single step evaporation, which indicates good volatil-
ity. Most of them also had very low residues, indicating that they do not decompose while
evaporating. Some differences, however, could still be found. The compounds with methylsilyl
ligands evaporated at lower temperatures than their counterparts with larger alkyl groups. The
telluride compounds are a clear example of this behavior. (Me3Si)2Te shows evaporation already
below 50 ◦C, while (tBuMe2Si)2Te and (Et3Si)2Te start to evaporate around 100
◦C. There was
also molecular weight dependency in the volatility, as group 15 alkylsilyls evaporated at higher
temperatures than group 16 alkylsilyls. In addition, there were some residues for group 15,
increasing from arsenic to antimony to bismuth, indicating some decomposition.
26
Figure 3.4: TGA curves for various alkylsilyl compounds of group 15 and 16 elements [184].
The use of alkylsilyl compounds with metal chlorides in ALD is based on very specific reactions,
namely dehalosilylation reactions. The reactions can be understood in terms of the Lewis hard-
soft acid-base (HSAB) theory, which states that hard acids and bases are small and only slightly
polarizable species, while soft acids and bases are larger and more polarizable [185]. In the
alkylsilyl compounds, silicon is a hard acid that is bonded to a large soft base such as selenium
or tellurium. In the second precursor, the chloride is a fairly hard base, bonded to a possibly soft
and large metal cation such as antimony(III). When the ligand exchange occurs, the results are
more favorable hard-hard and soft-soft pairs of silicon-chloride and metal-chalcogen -pairs [122].
Examples of these dechlorosilylation reactions are presented in Equations (3.1)-(3.3).
2 SbCl3(g) + 3 (Et3Si)2Te(g) −−→ Sb2Te3(s) + 6 Et3SiCl(g) (3.1)
GeCl2 ·C4H8O2(g) + (Et3Si)2Te(g) −−→ GeTe(s) + 2 Et3SiCl(g) + C4H8O2(g) (3.2)
SbCl3(g) + (Et3Si)3Sb(g) −−→ 2 Sb(s) + 3 Et3SiCl(g) (3.3)
27
Equation (3.1) is representative of all the dechlorosilylation reactions. The overall reactions
are identical, whether the alkylsilyl compound is a chalcogenide or a pnictide. Consequently,
making selenides or arsenides does not differ from depositing tellurides. Moreover, the oxidation
state of the cation is not significant for the reaction. GeTe2 can deposited as efficiently as the
more common GeTe [186, 187].
The dehalosilylation reaction is not disturbed by an additional neutral ligand such as 1,4-
dioxane in Equation (3.2). According to mechanistic studies, the dioxane ligand is completely
released during the GeCl2 ·C4H8O2 pulse [188]. A special example of the utilization of the
dechlorosilylation reaction is the formation of elemental antimony (Eq. (3.3)), which is enabled
by the disproportionation reaction of antimony [I].
The similar reactions, and the self-limiting nature of ALD, give a good way to deposit ternary
materials by combining binary processes. When Equations (3.1) and (3.2) are combined, the
result is a GST process. When GeTe and Sb2Te3 are pulsed alternately, Ge2Sb2Te5 can be
formed [1]. When all three of the above ALD processes are combined in various deposition
cycle ratios, the compositions can be controlled over a wide range.
All alkylsilyl compounds mentioned here are air and moisture sensitive, although somewhat
less so when the alkylsilyl groups are large. This means that they need to be handled carefully
while loading them into a reactor. Ethyl- and methylsilyl compounds are liquids and tert-
butylsilyls are low-melting solids [122]. The former are thus good ALD precursors also for
industrial applications since the surface area of a liquid does not change, keeping the precursor
dose constant. The vapor pressure of the methyl silyl compounds is relatively high, making
them easy to deliver in large doses but increasing the safety risk to some extent.
28
Chapter 4
Experimental
4.1 Atomic layer deposition of chalcogenide and pnictide
thin films
All thin films in the studies included in this thesis were deposited in a hot-wall, flow-type F120
ALD reactor (ASM Microchemistry Oy, Helsinki, Finland). Nitrogen was used as the carrier
and purge gas, with the operating pressure being below 10 mbar. The nitrogen was made on-site
with a Nitrox UHPN 3000 generator (rated purity 99.9995 %)
The reactor design limits the size of the substrates, all of which were 5× 5 cm2. The most
commonly used substrate materials were native SiO2 -terminated silicon and soda-lime glass.
Silicon substrates were used for most of the results that follow, while soda-lime glass was used
for reflectivity and some of the resistivity measurements.
Each precursor was housed separately inside the reactor, evaporated from open glass vessels
and heated individually to the required temperatures. GeCl2 ·C4H8O2 is a white needle-like
solid, evaporated at 65 ◦C. GaCl3, SbCl3 and BiCl3 are white powders, evaporated at room
temperature, 30 and 140 ◦C. GaCl3 has such a high volatility and reactivity with air that
more careful handling is required than with the other chlorides as HCl is easily released. All
alkylsilyl compounds used were liquids. (Et3Si)2Se, (Et3Si)2Te, (Et3Si)3As and (Et3Si)3Sb were
evaporated at 35, 40, 80 and 85 ◦C, respectively. Extra precautions, such as a gas mask and
gas detector were used when necessary, especially in the case of (Et3Si)3As. Precursor pulsing
was performed with inert gas valving.
29
4.2 Chalcogenide and pnictide thin film characterization
Scanning electron microscope images were acquired using a Hitachi S-4800 field emission scan-
ning electron microscope (FESEM). Film thicknesses and compositions were determined from
energy dispersive X-ray (EDX) spectra, measured using an INCA Energy 350 EDX spectrom-
eter connected to the S-4800. The measured k-ratios were transformed into thickness values by
using a GMR electron probe thin film microanalysis program [189].
Light impurities were analyzed using time-of-flight elastic recoil detection analysis (TOF-
ERDA), with a 5 MV tandem accelerator EGP-10-II, using 45 MeV 127I or 48 MeV 79Br as
a primary beam. In another setup, also 8 MeV 35Cl incident ions were used.
Films were also analyzed with Rutherford backscattering spectrometrry (RBS) using a 4.8 MeV
He beam and a silicon detector located at 169 degrees.
Transmission electron microscopy (TEM) was performed with a FEI Tecnai F20 200 kV trans-
mission electron microscope. Samples were prepared with a FEI Quanta 3D 200i DualBeam
focused ion beam/scanning electron microscope (FIB/SEM).
When applicable, thicknesses were also determined by X-ray reflectivity (XRR) measurements
with a Bruker D8 Advance X-ray diffractometer. Film crystallinity was analyzed with grazing
incidence X-ray diffraction (GIXRD) and θ−2θ measurements using a PANalytical X’Pert Pro
MPD X-ray diffractometer. For in situ high temperature XRD (HTXRD) measurements, an
Anton-Paar HTK1200N oven was used.
Atomic force microscopy was performed with a Veeco Instruments Multimode V with a Nano-
scope V controller. Samples were measured in tapping mode in air using phosphorus-doped
silicon probes (RTESP, Veeco Instruments) or silicon probes (RFESP, Bruker).
Crystallization times of phase change materials were measured using a static laser tester. For
crystallization studies, an amorphous sample was exposed to laser pulses of variable length
and power. For recrystallization studies, the sample was first crystallized by annealing and
then melt-quenched by laser pulses. Recrystallization was then attempted by applying pulses
of variable power and time to the same location.
30
The electrical properties were measured with an Ecopia HMS-5000 Hall Effect Measurement
system. Measurements were made in a 0.55 T magnetic field using the van der Pauw configura-
tion in a temperature range from 80 to 350 K. Small indium droplets were used to contact the
thin films. Routine resistivity measurements were made with a four point probe (CPS Probe
Station, Cascade Microtech) connected to a Keithley 2400 source meter.
Resistivity as a function of temperature was measured from phase change materials using a
custom-made set-up. Two large aluminum pads with a small and well defined gap were deposited
prior to the measurements and were contacted by two contact pins. The resistance between the
pins was measured during heating with a rate of 1 Ks−1.
Thermoelectric measurements were performed on glass substrates heated from one end by a heat
plate. Voltages were measured with a Keithley 2400 source meter and substrate temperatures
with a simple K-type probe attached to a multimeter.
Surface analysis by low-energy ion-scattering (LEIS) was performed at Tascon GmbH, using
an ION-TOF Qtac100 dedicated LEIS instrument. An overview of all elements heavier than B
were achieved with 3 keV 4He+ ions; 5 keV 20Ne+ ions were used for a quantitative analysis of
Bi and Te. The primary ion doses for the analyses were 1.4× 1014 He ions/cm2 and 3× 1013 Ne
ions/cm2. Quantification was performed by means of statistical calibration. The samples were
analysed ’as received’ after shipping, and after in situ cleaning with O atoms for 10 minutes to
remove hydrocarbon contamination. The statistical calibration was performed on the cleaned,
and thereby well-defined samples.
31
Chapter 5
Results and discussion
Alkylsilyl chalcogenides and pnictides have already enabled a large number of ALD processes
through the dehalosilylation reactions. When combined with chloride precursors, many previ-
ously unseen materials, such as GeTe, Sb and Bi2Se3, can now be made by ALD. In addition,
some new routes are provided for materials like ZnTe and GaAs that had also been deposited by
ALD earlier. All of the ALD processes demonstrated so far with alkylsilyl non-metal precursors
are presented in Figure 5.1. In this chapter, aspects of some of these ALD processes, mainly
those of Sb [I], GeTe [II], Bi2Te3 [III], Bi2Se3 [IV] and GaAs [V] are studied more closely.
Figure 5.1: ALD processes enabled by non-metal alkylsilyl precursors.
32
5.1 Properties of chalcogenide and pnictide ALD pro-
cesses
Although a large number of ALD processes with different precursors are covered in this thesis,
they are remarkably similar in many respects. The underlying reason must be the chemistry
that these processes share. The dehalosilylation reactions presented in Section 3.3 ensure that
the end results must be similar. This section delves closely into the characteristics of the ALD
processes and their similarities.
5.1.1 Growth rate saturation
Growth rate saturation is the key property in defining a process to be true ALD. The growth
rate of thin films should saturate after a sufficiently large precursor dose is received, which
ensures reactions with all available surface groups. Figures 5.2 and 5.3 show that this condition
is fulfilled in all processes that use alkylsilyl precursors.
0 0 . 5 1 . 0 1 . 5 2 . 0 2 . 5 3 . 0 3 . 500 . 10 . 20 . 30 . 40 . 5
0 . 6   Growth rate (Å/cycle) G e C l 2  C 4 H 8 O 2  p u l s e  ( s ). 0 0 . 5 1 . 0 1 . 5 2 . 0 2 . 5 3 . 0 3 . 500 . 10 . 20 . 30 . 40 . 50 . 6  
 Growth rate (Å/cycle) ( E t 3 S i ) 2 T e - p u l s e  ( s )
(a)
0 1 2 3 4 5 600 . 20 . 40 . 60 . 81 . 01 . 21 . 41 . 61 . 8
2 . 0   Growth rate (Å/cycle) B i C l 3  p u l s e  l e n g t h  ( s ) 0 1 2 3 400 . 20 . 40 . 60 . 81 . 01 . 21 . 41 . 6  
 Growth rate (Å/cycle) ( E t 3 S i ) 2 T e  p u l s e  l e n g t h  ( s )
(b)
0 1 2 3 4 5 600 . 51 . 01 . 5
2 . 0   Growth rate (Å/cycle) B i C l 3  p u l s e  l e n g t h  ( s ) 0 1 2 3 400 . 51 . 01 . 52 . 0  
 Growth rate (Å/cycle) ( E t 3 S i ) 2 S e  p u l s e  l e n g t h  ( s )
(c)
Figure 5.2: Effect of precursor pulse lengths on the growth rate of GeTe [II] (a), Bi2Te3 [III] (b)
and Bi2Se3 [IV] (c) thin films.
33
Both telluride processes, GeTe and Bi2Te3, show very similar characteristics with respect to
(Et3Si)2Te. There is saturation from early on with 1 s pulses. After that, there might even be a
decline in growth rates when precursor pulse lengths are increased. The reasons for this might
be similar to those for the purge effects, discussed in Section 5.1.2.
The only selenide process Bi2Se3, differs somewhat from the telluride processes. There appears
to be a slight increase in the growth rate after the stable growth region is achieved. Either
the (Et3Si)2Se is slower to adsorb and react on the surface than (Et3Si)2Te, or the slight drift
might just be due to a measurement uncertainty from the EDX or due to minor differences
in deposition temperature. There was a strong temperature dependence of the growth rate,
discussed in detail in Section 5.1.3.
0 0 . 5 1 . 0 1 . 5 2 . 0 2 . 500 . 10 . 20 . 30 . 40 . 50 . 6   Growth rate (Å/cycle) ( E t 3 S i ) 3 S b  p u l s e  l e n g t h  ( s )
(a)
0 0 . 5 1 . 0 1 . 5 2 . 000 . 10 . 20 . 30 . 40 . 50 . 6
0 . 7   Growth rate (Å/cycle) G a C l 3  p u l s e  l e n g t h  ( s )
(b)
Figure 5.3: Effect of precursor pulse lengths on the growth rate of Sb [I] (a) and GaAs [V] (b)
thin films.
It appears that the pnictide processes GaAs and Sb behave similarly to each other. Saturation
with (Et3Si)3As and (Et3Si)3Sb was achieved with 1 s pulses, after which the changes in growth
rates were minor.
As for the counterparts of the alkylsilyls, the metal or semimetal precursors, their saturation
behavior was very straightforward. Most saturated almost immediately, only BiCl3 needed
longer pulses to reach the growth rate saturation. The films deposited with shorter pulses
showed no differences in properties when compared to the films deposited with pulses further
in the saturation region.
Overall, proper ALD-like growth rate saturation behavior can be found in all of the alkylsilyl-
enabled processes. This saturation leads to repeatability and process stability, which enables
parameter tweaking for the desired reactor types and applications.
34
5.1.2 Purge effects
In general, purges should not have an effect on ALD processes, as long as they are sufficient
in removing all the excess precursors and by-products. Therefore, it was somewhat surprising
that the lengths of purges affected the growth of Bi2Te3 and Bi2Se3 so drastically. An example
with Bi2Se3 is given in Figure 5.4.
Figure 5.4: Effect of purge lengths on the growth rate of Bi2Se3 thin films [IV].
When only the purge after one precursor was increased from 1 to 8 s, the growth rate halved.
The overall decrease in growth rate was up to two thirds, when both purges were increased
at the same time. The reasons for this can be multitudinous. The adsorption might not be
complete immediately, and there could be some physisorption involved. This would mean that
the unreacted groups have more time to desorb from the surface, pushed along with the purge
gas flow.
Similar behavior has been seen in a GST process using alkoxides and alkylsilyl telluride where
precursor adsorption and desorption were balanced, allowing for ALD-like saturation behavior
in combination with the significant decrease in growth rate with respect to increasing purge
lengths [186]. Reaction mechanism studies on the Sb2Te3 process [188] have shown, however,
that a ligand exchange reaction occurs during both pulses. After such a reaction, the desorption
should be more difficult than in the case of chemisorption alone. Partial desorption is still
possible, possibly through ligand rearrangement.
35
5.1.3 Temperature dependence
The most striking property of the alkylsilyl ALD processes is their strong temperature depen-
dence. As discussed in Section 3.1, many ALD processes have growth rates that stay nearly
constant within a range of temperatures. This was definitely not the case for the chalcogenide
processes shown in Figure 5.5a.
All the processes showed a similar trend of a pronounced decrease in the growth rate with
increasing deposition temperatures. This was not connected to the absolute deposition temper-
atures, since Bi2Te3 and Bi2Se3 could be deposited at much higher temperatures than GeTe or
Sb2Te3. In fact, Bi2Te3 and Bi2Se3 deposition occurs at temperatures, where GeTe or Sb2Te3
no longer grow at all. In addition, ZnTe and ZnSe have been deposited at much higher temper-
atures (400 ◦C) than Bi2Te3 and Bi2Se3. Nevertheless, all of the chalcogenide processes follow
the trend of decreasing growth rates within their own temperature range.
The pnictide processes (Figure 5.5b) mostly behave in the same way as the chalcogenides. The
only difference is that both Sb and GaAs processes show a plateau of sorts in the middle of their
deposition temperature range. The usual explanation for this overall growth rate decrease would
be precursor desorption or a decrease in the density of the reactive sites. This type of behavior
is not restricted to alkylsilyl-enabled ALD processes - a similar temperature dependence was
also seen in ALD of AlF3 [190], for example.
1 0 0 1 5 0 2 0 0 2 5 000 . 2 50 . 5 00 . 7 5
1 . 0 01 . 2 51 . 5 0Growth rate (Å/cycle) D e p o s i t i o n  t e m p e r a t u r e  ( o C )
 G e T e B i 2 T e 3 B i 2 S e 3
(a)
1 0 0 1 5 0 2 0 0 2 5 0 3 0 000 . 20 . 4
0 . 60 . 81 . 0Growth rate (Å/cycle) D e p o s i t i o n  t e m p e r a t u r e  ( o C )
 S b G a A s
(b)
Figure 5.5: Effect of deposition temperature on the growth rate of chalcogenide (a) and pnictide
(b) thin films.
36
Similar temperature effects have been seen with MBE [35] and sputtering [191] of tellurides. A
suggested reason for this kind of behavior has been the high vapor pressure of tellurium, which
as a continuous flux can even etch a whole film away [35, 36].
5.1.4 Film growth
As can be expected of proper ALD processes, all alkylsilyl processes enable efficient film thick-
ness control. Regardless of the temperature dependence or purge related behavior, film thickness
was always determined by the number of deposition cycles. There appeared to have been no
incubation periods in any of the processes, as the relation between the film thickness and the
number of deposition cycles was linear from the very beginning.
Nevertheless, this did not mean that growth progressed in the quintessential layer-by-layer
mode. It could have been the case with the processes that produced amorphous films, but LEIS
has indicated that growth may very well follow the island-growth mode in the early stages of
processes producing crystalline films [III].
The initial stages of Bi2Te3 film growth were followed with LEIS, the results of which can be
seen in Figure 5.6. In Figure 5.6a, the silicon signal slowly disappeared as the bismuth and
tellurium signals increased, indicating that the film was closing over the substrate surface. Full
closure occurred after about 70-90 deposition cycles, which would mean a film thickness of
about 10 nm according to EDX. LEIS revealed another interesting aspect of the early Bi2Te3
film growth, namely the changing composition of the outermost layer of the growing film. In
Figure 5.6b, it can be seen that mostly bismuth is deposited during the first few cycles, even
though (Et3Si)2Te was pulsed last in the cycle. Close to stoichiometric, and stabilized, ratios
of Bi and Te could be found on the surface after eight deposition cycles [III].
37
0 5 0 1 0 0 1 5 0 2 0 0051 01 52 0
2 53 03 54 0 Surface fraction (%)  
  B i T e S i O
C y c l e  n u m b e r
(a)
0 5 0 1 0 0 1 5 0 2 0 00 . 00 . 20 . 4
0 . 60 . 81 . 0  B i + T e / ( B i + T e + S i ) T e / ( B i + T e )Surface fraction rati
o
C y c l e  n u m b e r
(b)
Figure 5.6: Surface fraction of the elements (a) and their ratios with respect to the number of
deposition cycles (b), measured with LEIS [III].
5.1.5 Conformality
The self-limiting saturative growth that defines ALD enables conformal deposition onto high
aspect ratio substrates. As the precursors are transported to the substrate in the gas phase,
they can react with all available surfaces, assuming the precursor molecule is small enough to
physically fit into the possible crevices. A good example can be seen in Figure 5.7.
(a) (b)
Figure 5.7: FESEM image of Sb (a) and a TEM image of GeTe [II] (b) thin films deposited
onto demanding three-dimensional substrates.
38
In Figure 5.7a, the Sb film follows the surface of the trench structure closely. The film thickness
appears equal on both the top surface and at the bottom of the trench, as well as on the side
walls. Similar behavior is demonstrated by the GeTe film in Figure 5.7b. The GeTe film was
deposited into a 31 × 84 nm via at the bottom of the trench. The via is completely filled, the
film follows the substrate shape with little thickness variation, without a visible seam. These
two processes are a good example of ALD conformality, as the Sb produces crystalline films
and the GeTe process amorphous films. This crystallinity difference has no meaning for the
inherent properties of ALD, and good conformality is the end result. Conformality is especially
important for applications like phase change memories, where the memory element is located in
a tight space like in Figure 2.2c. Voids or unevenness in the film can cause device malfunction.
For complete via filling, an amorphous structure is more suitable because a polycrystalline film
will produce a grain boundary in the middle of the via.
5.2 Properties of ALD chalcogenide and pnictide thin
films
In addition to being notable additions to the materials selections of ALD processes, the chalco-
genide and pnictide films should also be of good quality, and have the right properties to be
industrially applicable. The ALD chalcogenide and pnictide thin films should present the prop-
erties needed for their most typical and intended applications. The following sections present
studies on the properties of these thin films.
5.2.1 Composition
The chemistry of the dehalosilylation reaction has proven to be highly efficient in producing
stoichiometric films in most deposition conditions. The correct composition is one of the most
important properties a thin film must have. The chemistry should generally lead to stoichiomet-
ric compound formation and not just alloying, which maybe in some cases the desired result.
These compositions should also be stable, and not affected by minor parameter changes.
The versatility of ALD can be seen when different binary processes are mixed to form ternar-
ies. This was definitely the case for phase change materials, where the binary processes were
highly compatible. By combining GeTe, Sb2Te3, GeSb and Sb processes in the same deposition,
numerous compositions within the phase diagram in Figure 5.8 were achieved.
39
Figure 5.8: Compositions of ALD phase change materials. Filled circles represent examples of
measured EDX compositions of thin films. Reprinted from [122] with permission from Elsevier,
Copyright (2013).
In the instance of thermoelectric materials Bi2Te3 and Bi2Se3, the deposition processes were
also in this aspect practically identical. Both processes produced n-type chalcogen-rich films
in all deposition conditions. The compositional stability is excellent for repeatability, although
thermoelectric materials are usually tailored alloys, where both p- and n-types are needed.
In the Sb process, the disproportionation reaction between the two Sb-containing precursors
can only lead to elemental Sb, which was also the case. Even such a reaction still leaves a
great deal to be studied composition-wise. For example, impurities are very important factors
in determining other film properties.
The only greater variations in compositions were seen in the GaAs process. According to EDX,
stoichiometric films in terms of the Ga/As ratio were produced only at the ALD window
temperature of 200 ◦C. This was also verified with RBS, meaning that the chemistry was not
as stable as it was with all the other alkylsilyl processes, where the temperature did not affect
the composition.
For the most part, alkylsilyl-enabled ALD processes produced very pure films. This is especially
noteworthy when the deposition temperatures are considered. They vary between 80 and 160 ◦C,
and are very low for typical ALD processes. In general, higher deposition temperatures lead to
less impurities in the films. As shown for the Sb process in Figure 5.9, in the films discussed
in this chapter, only a few atomic percent or fractions thereof of the most typical impurity
elements, such as carbon, nitrogen, hydrogen, oxygen and chlorine, were found. In addition,
40
most of the oxygen was found on the surface of the films - the impurity contents within the
films were even lower.
Figure 5.9: TOF-ERDA depth profile of an Sb thin film [I].
Gallium arsenide was the exception here as well. At low temperatures and long GaCl3 pulses,
the chlorine content was much higher, even about 20 at.% according to EDX. TOF-ERDA
measurements from a film deposited at 200 ◦C indicated that most impurities were at the level
of a single atomic per cent or less. Only the amount of oxygen was higher, and the exact location
of the oxygen could not be determined. The most probable location is on the surface of the film,
such as in Figure 5.9. Some of the impurities were most likely due to some incomplete reactions
of GaCl3. Overall, it seems that the GaAs process is the most unstable of the alkylsilyl-enabled
ones.
5.2.2 Crystallinity and morphology
Crystallinity and morphology are very important thin film properties. Many applications require
the film to be either amorphous or crystalline, and exist in the correct crystalline phase if
there are many. It cannot be known beforehand whether a particular ALD process produces
amorphous or crystalline films; and if crystalline, in what phase. The phase is influenced by
deposition parameters, which need to be optimized. Undesirable phases are one reason to search
for new precursors and chemistries. The ALD processes discussed in this thesis had crystallinity
and morphology that were very much process dependent, and to some extent temperature
dependent.
41
The only amorphous films were produced by the GeTe and GaAs processes. Also, most ternary
phase change materials, with a sufficiently low content of either Sb or Sb2Te3, were amorphous
as-deposited at the lowest deposition temperatures. This is good since amorphous materials are
preferred in phase change applications due to smoother film surfaces.
For example, GeTe thin films could be crystallized only by annealing and not during the de-
position. An example HTXRD is provided in Figure 5.10. The film stayed amorphous past
the possible deposition temperature range and crystallized only at 165 ◦C. The initial crystal-
lization formed a rhombohedral phase, which then transformed to the high-temperature cubic
phase at 385 ◦C. All the peaks show a slow drift towards smaller 2θ values, more than only a
thermal expansion would cause. In addition, the two peaks at about 43◦ combine into a single
peak with the cubic transformation. The high-temperature cubic phase transformed back into
the rhombohedral phase when cooled, and the double peaks could again be resolved [II]. This
crystallization and diffraction behavior of ALD GeTe is very similar compared to literature
values for sputtered GeTe [192].
Figure 5.10: HTXRD of a GeTe thin film deposited at 90 ◦C [II].
The rhombohedral phase was the common denominator for Bi2Te3 and Bi2Se3. They both
belong to the same space group, and thus have almost identical diffraction patterns. The slight
differences in the diffraction patterns arise from the differences in the lattice parameters caused
by the telluride ion being bigger than the selenide ion. The θ− 2θ diffractograms of Bi2Te3 and
Bi2Se3 were indeed very similar, as can be seen from Figure 5.11a. Both clearly show diffraction
peaks arising from the same crystal planes, namely the (003) and its multiples. This is an
indication of preferred orientation of growth in the (001), or the c-axis, direction.
42
The similarities were also seen in TEM, where the initial stages of growth for both Bi2Te3
and Bi2Se3 included oriented small grains. The TEM images in Figure 5.11b demonstrate the
clear crystalline structure of the grains. The distance between the fringes in both images, about
1 nm, corresponded to the distance between the (003) planes in the crystal structure, which is
the physical length of one quintuple layer, (Te/Se)-Bi-(Te/Se)-Bi-(Te/Se), of the material. The
preferred orientation seen in XRD is thus confirmed by TEM. The growth of Bi2Te3 and Bi2Se3
proceeds mainly in the (001) direction.
(a) (b)
Figure 5.11: XRD (a) and TEM (b) of Bi2Te3 and Bi2Se3 thin films deposited at 160
◦C [III, IV].
The surface morphologies have been found to be a direct consequence of the crystalline structure
of the growing films. Amorphous substances are naturally featureless on the surface, while in
crystalline materials the differences in the original crystalline orientation and the growth rates
of the different crystalline facets will cause roughness on the surface [193–195]. Examples are
presented in Figure 5.12. There is little to be seen in the image of GeTe (Fig. 5.12a) as it was
amorphous. Sb2Te3 (Fig. 5.12b) and Bi2Te3 (Fig. 5.12c), on the other hand, show practically
identical surface structures of flakes. This stems from the crystalline structures, which are both
rhombohedral and belong to the space group 166. This type of flaky morphology exists at the
very surface of the films, while the parts deeper in the film show a dense structure and firm
attachment to the substrate, as seen in the TEM images above. The flaky surface structure
is therefore likely to be caused by the cohesive failure within the crystallites of the film as it
grows thicker.
43
(a) (b) (c)
Figure 5.12: FESEM images of GeTe [II] (a), Sb2Te3 (b) and Bi2Te3 [III] (c) thin films. The
scale bar in all images is 1 µm.
5.2.3 Phase change properties
Phase change memories rely on the phase change between the amorphous and crystalline phases.
This change needs to be easy and repeatable. There also needs to be a clear difference in the
optical and electrical properties between the phases. All of these requirements were fulfilled by
ALD GeTe, and earlier also for ALD GST [196].
The optical changes are demonstrated in Figure 5.13, where a laser has been used to crystallize
the amorphous as-deposited thin film (Fig. 5.13a), or to first melt-quench an annealed film
and then recrystallize it (Fig. 5.13b). The results from these two images are in line with the
literature, as recrystallization occurs much faster than the initial crystallization. This has been
explained with small crystallites being left behind after melt-quenching [16]. In this case, the
crystallization took place in 300 ns, which is a comparable value to those measured from sput-
tered films. The recrystallization occurred in only 50 ns, which is faster than that in sputtered
films [II].
The electrical resistivity was also compared to the sputtered films. The behavior was again
comparable, as can be seen from Figure 5.14. The only difference is the slightly higher crystal-
lization temperature of the ALD GeTe, 185 ◦C. The difference in resistivity between amorphous
and crystalline phases is more than enough for readout purposes or multilevel storage, six orders
of magnitude [II].
44
(a) (b)
Figure 5.13: Laser crystallization (a) and recrystallization of melt-quenched (b) GeTe [II].
Figure 5.14: Resistivity of ALD (squares) and sputtered (triangles) GeTe [II].
It can be concluded that thin films made by ALD work at least just as well as the customary
sputtered films. This has been shown here for GeTe, and previously also for GST [196].
5.2.4 Thermoelectric properties
The properties used for determining the thermoelectric applicability of a material are all found
in the figure of merit ZT (Eq. 2.1). Electrical measurements bring out the differences between
Bi2Te3 and Bi2Se3, as can be seen from Figure 5.15. As Bi2Te3 has much better conductivity
than Bi2Se3, it is also fairly safe to assume that from these two ALD materials, Bi2Te3 would
perform better as a thermoelectric material. Additionally, Bi2Se3 is most commonly alloyed in
general and not so much used on its own as a thermoelectric.
45
5 0 1 0 0 1 5 0 2 0 0 2 5 0 3 0 0 3 5 01 . 0 x 1 0 2 01 . 5 x 1 0 2 02 . 0 x 1 0 2 02 . 5 x 1 0 2 03 . 0 x 1 0
2 03 . 5 x 1 0 2 0  B i 2 T e 3 B i 2 S e 3   Carrier concentration (cm-3 ) T e m p e r a t u r e  ( K )
(a)
5 0 1 0 0 1 5 0 2 0 0 2 5 0 3 0 0 3 5 0051 01 52 02 5
3 0  B i 2 T e 3 B i 2 S e 3Mobility (cm2 /Vs) T e m p e r a t u r e  ( K )
(b)
5 0 1 0 0 1 5 0 2 0 0 2 5 0 3 0 0 3 5 002 0 04 0 06 0 08 0 0
1 0 0 0Conductivity (S/cm) T e m p e r a t u r e  ( K )  B i 2 T e 3 B i 2 S e 3
(c)
Figure 5.15: Temperature dependence of carrier concentration (a), mobility (b) and conductivity
(c) of Bi2Te3 (squares) and Bi2Se3 (circles) thin films deposited at 160
◦C [III].
The thermoelectric performance was studied with the setup from Figure 5.16. The results
demonstrate clear thermoelectric behavior for both Bi2Te3 and Bi2Se3. Surprisingly, they be-
haved almost identically. Both measurements lead to slopes of −180 µV K−1, which agree very
well with the literature values for n-type Bi2Te3 [49] and Bi2Se3 [77, 197–199].
(a) (b)
Figure 5.16: Thermoelectric measurement setup (a) and the thermoelectric response (b) of
Bi2Te3 (triangles) and Bi2Se3 (circles) thin films deposited at 160
◦C [IV].
Although ALD thin films have not been used in thermoelectric devices, the previous results
prove that this could be possible. When the substantial thremoelectric response is combined
with the good quality of the films and deposition processes themselves, ALD should be consid-
ered as a method to deposit thermoelectric materials for applications.
46
Chapter 6
Conclusions and outlook
The importance of materials and device engineering becomes increasingly evident as time goes
by. The progress in the microelectronics industry, fueled by Moore’s law, leads to smaller and
smaller devices. The laws of physics will put an end to the shrinkage at some point. Eventually,
new materials are needed and device structures become increasingly complicated and three-
dimensional. It is this complexity that gives more opportunities for ALD to be used. In the end,
some thin film structures can only be made with ALD due to the conformality of the method.
This has led to a great increase in ALD research as can be seen from the increasing number
of scientific papers published in the field. As a result, there is also a need for more application
driven research. Applications need materials with specific compositions and properties.
Overall, the alkylsilyl compounds have proven to be excellent ALD precursors. They evaporate
cleanly and have shown good temperature stability. With the deceptively simple dehalosilylation
reaction, good quality thin films have been made. The processes themselves have presented all
the necessary requirements of ALD. Growth rates saturate with respect to precursor pulse
lengths; film thicknesses are easily controlled by the number of deposition cycles; and the films
are conformally deposited on complex substrate surfaces.
Moreover, the ALD thin films themselves present excellent properties. Generally, films have been
stoichiometric, with low impurity contents. Desired crystalline phases were observed, and even
oriented growth was detected. In addition, the phase change and thermoelectric properties of the
appropriate materials have been proven, which is essential for using ALD in those applications.
Further research ventures in to this area should delve into the reasons behind the untypical
purge-related behavior seen with Bi2Te3 and Bi2Se3. The cause behind this behavior should be
determined and whether it is more common or only limited to these particular processes.
47
One interesting avenue to follow would be making thermoelectric devices from ALD thin films.
Since most devices are macroscopic in size, and the thermoelectric material made to bulk form,
it presents an obvious challenge. Two questions to be addressed are: How thin can the devices be
made? What kinds of structural changes in the device would be needed with a vapor deposition
method?
In the phase change materials, there is still much of the Ge-Sb-Te phase diagram to be explored.
Even though ALD of many compositions has been proven, only a few have been tested for
repeatable phase change. In addition, as doping has always been an important part of phase
change material tuning, other elements, such as gallium, could be introduced to the processes.
This should not lead to any major compatibility issues, providing the dehalosilylation reaction
is not disturbed.
There are still a number of chalcogenide and pnictide processes that have only been demon-
strated to work. Further studies could be made on indium chalcogenides and aluminum pnic-
tides, for example. It is evident that the dehalosilylation reaction will work for more precursor
combinations than presented in this thesis.
The next possible materials to be explored with this alkylsilyl chemistry are the transition
metal dichalcogenides (TMDC). They have become the “hot topic” in the research community
due to their possibilities in the 2D scale. In general, 2D TMDCs have band gaps, which makes
them usable in electronic applications where graphene cannot be used.
The processes demonstrated here cover a variety of technologically important materials. With
the growing ALD market, it will be interesting to see whether some of these processes make it
into production and onwards to the consumers. Nevertheless, there will definitely be a need for
more ALD research, also with alkylsilyl non-metal precursors.
48
References
1. Pore, V., Hatanpa¨a¨, T., Ritala, M., and Leskela¨, M. J. Am. Chem. Soc. 131(10), 3478–
3480 (2009).
2. Ovshinsky, S. R. Phys. Rev. Lett. 21(20), 1450–1453 (1968).
3. Wuttig, M. Nat. Mater. 4(4), 265–266 (2005).
4. Wuttig, M. and Yamada, N. Nat. Mater. 6(11), 824–832 (2007).
5. Elliott, S. R. In Phase Change Materials: Science and Applications, Raoux, S. and Wuttig,
M., editors, 63–80. Springer Science & Business Media (2009).
6. Yamada, N., Ohno, E., Nishiuchi, K., Akahira, N., and Takao, M. J. Appl. Phys. 69(5),
2849–2856 (1991).
7. Nirschl, T., Philipp, J., Happ, T., Burr, G., Rajendran, B., Lee, M., Schrott, A., Yang,
M., Breitwisch, M., Chen, C., Joseph, E., Lamorey, M., Cheek, R., Chen, S.-H., Zaidi,
S., Raoux, S., Chen, Y., Zhu, Y., Bergmann, R., Lung, H., and Lam, C. In Int. Electron
Devices Meet., 461–464, (2007).
8. Cassinerio, M., Ciocchini, N., and Ielmini, D. Adv. Mater. 25(41), 5975–5980 (2013).
9. Lacaita, A. L. and Redaelli, A. Microelectron. Eng. 109, 351–356 (2013).
10. Elliott, S. R. Int. J. Appl. Glass Sci. 6(1), 15–18 (2015).
11. Raoux, S. and Wuttig, M., editors. Phase Change Materials: Science and Applications.
Springer Science & Business Media, (2009).
12. Veiko, V., Ageev, E., Kolobov, A. V., and Tominaga, J. J. Optoelectron. Adv. M. 15(5–6),
371–394 (2013).
13. Adler, D., Henisch, H. K., and Mott, S. N. Rev. Mod. Phys. 50(2), 209–220 (1978).
14. Raoux, S., Shelby, R., Mun˜oz, B., Hitzbleck, M., Krebs, D., Salinga, M., Woda, M., Aust-
gen, M., Chung, K., and Wuttig, M. In E*PCOS 2008.
15. Raoux, S., Welnic, W., and Ielmini, D. Chem. Rev. 110(1), 240–267 (2010).
16. Cheng, H.-Y., Raoux, S., and Chen, Y.-C. J. Appl. Phys. 107(7), 074308 (2010).
49
17. Loke, D., Lee, T. H., Wang, W. J., Shi, L. P., Zhao, R., Yeo, Y. C., Chong, T. C., and
Elliott, S. R. Science 336(6088), 1566–1569 (2012).
18. Terao, M., Morikawa, T., and Ohta, T. Jpn. J. Appl. Phys. 48(8), 080001 (2009).
19. Bruns, G., Merkelbach, P., Schlockermann, C., Salinga, M., Wuttig, M., Happ, T. D.,
Philipp, J. B., and Kund, M. Appl. Phys. Lett. 95(4), 043108 (2009).
20. Shi, L. In Phase Change Materials: Science and Applications, Raoux, S. and Wuttig, M.,
editors, 253–284. Springer Science & Business Media (2009).
21. Lee, S.-H., Hwang, Y., Lee, S., Ryoo, K., Ahn, S., Koo, H., Jeong, C., Kim, K., Koh, G.,
Jeong, G., Jeong, H., and Kim, K. In Symp. on VLSI Tech., 20–21, (2004).
22. Pellizzer, F., Pirovano, A., Ottogalli, F., Magistretti, M., Scaravaggi, M., Zuliani, P., Tosi,
M., Benvenuti, A., Besana, P., Cadeo, S., Marangon, T., Morandi, R., Piva, R., Spandre,
A., Zonca, R., Modelli, A., Varesi, E., Lowrey, T., Lacaita, A., Casagrande, G., Cappelletti,
P., and Bez, R. In Symp. on VLSI Tech., 18–19, (2004).
23. Breitwisch, M., Nirschl, T., Chen, C., Zhu, Y., Lee, M., Lamorey, M., Burr, G., Joseph,
E., Schrott, A., Philipp, J., Cheek, R., Happ, T., Chen, S.-H., Zaidi, S., Flaitz, P., Bruley,
J., Dasaka, R., Rajendran, B., Rossnage, S., Yang, M., Chen, Y., Bergmann, R., Lung,
H., and Lam, C. In Symp. on VLSI Tech., 100–101, (2007).
24. Lacaita, A. Solid-State Electron. 50(1), 24–31 (2006).
25. Kolobov, A. V., Fons, P., Frenkel, A. I., Ankudinov, A. L., Tominaga, J., and Uruga, T.
Nat. Mater. 3(10), 703–708 (2004).
26. Wuttig, M. and Raoux, S. Z. Anorg. Allg. Chem. 638(15), 2455–2465 (2012).
27. Shportko, K., Kremers, S., Woda, M., Lencer, D., Robertson, J., and Wuttig, M. Nat.
Mater. 7(8), 653–658 (2008).
28. Lencer, D., Salinga, M., Grabowski, B., Hickel, T., Neugebauer, J., and Wuttig, M. Nat.
Mater. 7(12), 972–977 (2008).
29. Baldwin, R. A., Foos, E. E., and Wells, R. L. Mater. Res. Bull. 32(2), 159–163 (1997).
30. Caldwell, M. A., Raoux, S., Wang, R. Y., Wong, H.-S. P., and Milliron, D. J. J. Mater.
Chem. 20(7), 1285–1291 (2010).
31. Ie, S. Y., Bea, B. T., Ahn, Y.-k., Chang, M. Y., You, D. G., Cho, M. H., Jeong, K., Oh,
J.-H., Koh, G.-H., and Jeong, H. Appl. Phys. Lett. 90(25), 251917 (2007).
32. Wagner, T., Orava, J., Prikryl, J., Kohoutek, T., Bartos, M., and Frumar, M. Thin Solid
Films 517(16), 4694–4697 (2009).
33. Hu, D., Zhu, J., and Lee, J. Integr. Ferroelectr. 84(1), 233–238 (2006).
50
34. Neˇmec, P., Moreac, A., Nazabal, V., Pavliˇsta, M., Prikrˇyl, J., and Frumar, M. J. Appl.
Phys. 106(10), 103509 (2009).
35. Braun, W., Shayduk, R., Flissikowski, T., Ramsteiner, M., Grahn, H. T., Riechert, H.,
Fons, P., and Kolobov, A. Appl. Phys. Lett. 94(4), 041902 (2009).
36. Shayduk, R. and Braun, W. J. Cryst. Growth 311(7), 2215–2219 (2009). The 15th
International Conference on Molecular Beam Epitaxy (MBE-XV).
37. Jung, Y., Lee, S.-H., Ko, D.-K., and Agarwal, R. J. Am. Chem. Soc. 128(43), 14026–14027
(2006).
38. Wagner, R. S. and Ellis, W. C. Appl. Phys. Lett. 4(5), 89–90 (1964).
39. Yoon, H., Jo, W., Lee, E., Lee, J., Kim, M., Lee, K., and Khang, Y. J. Non-Cryst. Solids
351(43–45), 3430–3434 (2005).
40. Park, G.-S., Kwon, J.-H., Kim, M., Yoon, H. R., Jo, W., Kim, T. K., Zuo, J.-M., and
Khang, Y. J. Appl. Phys. 102(1), 013524 (2007).
41. Ohring, M. Materials Science of Thin Films: Deposition and Structure. Academic Press,
second edition, (2002).
42. Kim, R.-Y., Kim, H.-G., and Yoon, S.-G. Integr. Ferroelectr. 90(1), 80–87 (2007).
43. Longo, M., Salicio, O., Wiemer, C., Fallica, R., Molle, A., Fanciulli, M., Giesen, C.,
Seitzinger, B., Baumann, P., Heuken, M., and Rushworth, S. J. Cryst. Growth 310(23),
5053–5057 (2008). The Fourteenth International Conference on Metalorganic Vapor Phase
Epitaxy.
44. Abrutis, A., Plausinaitiene, V., Skapas, M., Wiemer, C., Salicio, O., Pirovano, A., Varesi,
E., Rushworth, S., Gawelda, W., and Siegel, J. Chem. Mater. 20(11), 3557–3559 (2008).
45. Kim, R.-Y., Kim, H.-G., Park, K.-W., Ahn, J.-K., and Yoon, S.-G. Chem. Vap. Deposition
15(10–12), 296–299 (2009).
46. Choi, B. J., Choi, S., Shin, Y. C., Hwang, C. S., Lee, J. W., Jeong, J., Kim, Y. J., Hwang,
S.-Y., and Hong, S. K. J. Electrochem. Soc. 154(4), H318–H324 (2007).
47. Choi, B. J., Choi, S., Eom, T., Ryu, S. W., Cho, D.-Y., Heo, J., Kim, H. J., Hwang, C. S.,
Kim, Y. J., and Hong, S. K. Chem. Mater. 21(12), 2386–2396 (2009).
48. Cadoff, I. and Miller, E. Thermoelectric Materials and Devices. Reinhold Publishing
Corporation, New York, (1960).
49. Rowe, D., editor. CRC Handbook of Thermoelectrics. CRC, Boca Raton, Florida, USA,
(1995).
50. ThinkGeek. USB Thermoelectric Cooler & Warmer - http://www.thinkgeek.com/
product/96b3/.
51
51. Bell, L. E. Science 321(5895), 1457–1461 (2008).
52. Heremans, J. P., Dresselhaus, M. S., Bell, L. E., and Morelli, D. T. Nat. Nanotechnol.
8(7), 471–473 (2013).
53. Mart´ın-Gonza´lez, M., Caballero-Calero, O., and Dı´az-Chao, P. Renew. Sus. Energ. Rev.
24, 288–305 (2013).
54. Dresselhaus, M., Chen, G., Ren, Z., Dresselhaus, G., Henry, A., and Fleurial, J. JOM J.
Min. Met. Mat. S. 61(4), 86–90 (2009).
55. Dresselhaus, M., Lin, Y., Cronin, S., Black, M., Rabin, O., and Dresselhaus, G. In Pro-
ceedings of the Nonlithographic and Lithographic Methods for Nanofabrication MRS Sym-
posium, (2001).
56. Bos, J.-W. G. and Downie, R. A. J. Phys.: Condens. Matter 26(43), 433201 (2014).
57. Snyder, G. J. and Toberer, E. S. Nat. Mater. 7(2), 105–114 (2008).
58. Slack, G. In CRC Handbook of Thermoelectrics, Rowe, D., editor, 407–440. CRC (1995).
59. Kanatzidis, M. G. Chem. Mater. 22(3), 648–659 (2010).
60. Hicks, L. D. and Dresselhaus, M. S. Phys. Rev. B 47(19), 12727–12731 (1993).
61. Hicks, L. D. and Dresselhaus, M. S. Phys. Rev. B 47(24), 16631–16634 (1993).
62. Zhao, L.-D., Dravid, V. P., and Kanatzidis, M. G. Energy Environ. Sci. 7(1), 251–268
(2014).
63. Godart, C., Gonc¸alves, A., Lopes, E., and Villeroy, B. In Mater. Res. Soc. Symp. Proc.,
volume 1166-N08-01, 183–193, (2009).
64. Gonc¸alves, A. P. and Godart, C. Eur. Phys. J. B 87(2), 42 (2014).
65. Goldsmid, H. J. Materials 7(4), 2577 (2014).
66. LaLonde, A. D., Pei, Y., Wang, H., and Snyder, G. J. Mater. Today 14(11), 526–532
(2011).
67. Alam, H. and Ramakrishna, S. Nano Energy 2(2), 190–212 (2013).
68. Vining, C. In CRC Handbook of Thermoelectrics, Rowe, D., editor, 329–338. CRC (1995).
69. Uher, C. In Proc. 22nd Int. Conf. Thermoelect., 42–47, (2003).
70. Mu¨chler, L., Casper, F., Yan, B., Chadov, S., and Felser, C. Phys. Status Solidi RRL
7(1–2), 91–100 (2013).
71. He, J., Liu, Y., and Funahashi, R. J. Mater. Res. 26(15), 1762–1772 (2011).
72. Chen, Y., Zhao, Y., and Liang, Z. Energy Environ. Sci. 8(2), 401–422 (2015).
52
73. Gaultois, M. W., Sparks, T. D., Borg, C. K., Seshadri, R., Bonificio, W. D., and Clarke,
D. R. Chem. Mater. 25(15), 2911–2920 (2013).
74. Fano, V. In CRC Handbook of Thermoelectrics, Rowe, D., editor, 257–266. CRC (1995).
75. Zhang, H., Liu, C.-X., Qi, X.-L., Dai, X., Fang, Z., and Zhang, S.-C. Nat. Phys. 5(6),
438–442 (2009).
76. Hsieh, D., Xia, Y., Qian, D., Wray, L., Dil, J. H., Meier, F., Osterwalder, J., Patthey, L.,
Checkelsky, J. G., Ong, N. P., Fedorov, A. V., Lin, H., Bansil, A., Grauer, D., Hor, Y. S.,
Cava, R. J., and Hasan, M. Z. Nature 460(7259), 1101–1105 (2009).
77. Hor, Y. S., Richardella, A., Roushan, P., Xia, Y., Checkelsky, J. G., Yazdani, A., Hasan,
M. Z., Ong, N. P., and Cava, R. J. Phys. Rev. B 79(19), 195208 (2009).
78. Hsieh, D., Xia, Y., Qian, D., Wray, L., Meier, F., Dil, J. H., Osterwalder, J., Patthey, L.,
Fedorov, A. V., Lin, H., Bansil, A., Grauer, D., Hor, Y. S., Cava, R. J., and Hasan, M. Z.
Phys. Rev. Lett. 103(14), 146401 (2009).
79. Hasan, M. Z. and Kane, C. L. Rev. Mod. Phys. 82(4), 3045–3067 (2010).
80. Moore, J. E. Nature 464(7286), 194–198 (2010).
81. Qi, X.-L. and Zhang, S.-C. Rev. Mod. Phys. 83(4), 1057–1110 (2011).
82. Chen, Y. L., Analytis, J. G., Chu, J.-H., Liu, Z. K., Mo, S.-K., Qi, X. L., Zhang, H. J., Lu,
D. H., Dai, X., Fang, Z., Zhang, S. C., Fisher, I. R., Hussain, Z., and Shen, Z.-X. Science
325(5937), 178–181 (2009).
83. Moore, J. Nat. Phys. 5(6), 378–380 (2009).
84. Welker, H. Z. Naturforsch. A 7, 744–749 (1952).
85. Welker, H. Z. Naturforsch. A 8, 248–251 (1953).
86. Haynes, W. M., editor. CRC Handbook of Chemistry and Physics. CRC Press/Taylor and
Francis, Boca Raton, Florida, USA, 95th edition, (2014).
87. Kasap, S. and Capper, P., editors. Springer Handbook of Electronic and Photonic Mate-
rials. Springer, New York, (2006).
88. Pankove, J. I. and Berkeyheiser, J. E. Proc. IRE 50(9), 1976–1977 (1962).
89. Black, J., Lockwood, H., and Mayburg, S. J. Appl. Phys. 34(1), 178–180 (1963).
90. Schubert, E. F. Light-emitting Diodes. Cambridge University Press, (2003).
91. Nathan, M. I., Dumke, W. P., Burns, G., Dill, F. H., and Lasher, G. Appl. Phys. Lett.
1(3), 62–64 (1962).
92. Hall, R. N., Fenner, G. E., Kingsley, J. D., Soltys, T. J., and Carlson, R. O. Phys. Rev.
Lett. 9(9), 366–368 (1962).
53
93. Quist, T. M., Rediker, R. H., Keyes, R. J., Krag, W. E., Lax, B., McWhorter, A. L., and
Zeigler, H. J. Appl. Phys. Lett. 1(4), 91–92 (1962).
94. Green, M. A., Emery, K., Hishikawa, Y., and Warta, W. Prog. Photovolt.: Res. Appl.
17(5), 320–326 (2009).
95. Mimura, T. IEEE Trans. Microwave Theory Tech. 50(3), 780–782 (2002).
96. Ridley, B. K. and Watkins, T. B. Proc. Phys. Soc. 78(2), 293 (1961).
97. Hilsum, C. Proc. IRE 50(2), 185–189 (1962).
98. Gunn, J. Solid State Commun. 1(4), 88–91 (1963).
99. Oliver, J., Fairman, R., Chen, R., and Yu, P. Electron. Lett. 17(22), 839–841 (1981).
100. Holmes, D. E., Chen, R. T., Elliott, K. R., and Kirkpatrick, C. G. Appl. Phys. Lett. 40(1),
46–48 (1982).
101. Hoshikawa, K., Nakanishi, H., Kohda, H., and Sasaura, M. J. Cryst. Growth 94(3), 643–
650 (1989).
102. Chen, T., Huang, T., Chen, L., and Guo, Y. J. Cryst. Growth 106(23), 367–376 (1990).
103. Bourret, E. and Merk, E. J. Cryst. Growth 110(3), 395–404 (1991).
104. Moss, S. and Ledwith, A., editors. The Chemistry of the Semiconductor Industry. Blackie
& Son Ltd., Glasgow, (1987).
105. Campomanes, R. R., da Silva, J. H. D., Vilcarromero, J., and Cardoso, L. P. J. Non-Cryst.
Solids 299-302(Part 2), 788–792 (2002).
106. Ku¨nzel, H. Physica B + C 129(1-3), 66–80 (1985).
107. Panish, M. B. and Temkin, H. Annu. Rev. Mater. Sci. 19(1), 209–229 (1989).
108. Stringfellow, G. B. Organometallic vapor phase epitaxy - theory and practice. Academic
Press, San Diego, (1999).
109. Goodman, C. H. L. and Pessa, M. V. J. Appl. Phys. 60(3), R65–R82 (1986).
110. Usui, A. and Watanabe, H. Annu. Rev. Mater. Sci. 21(1), 185–219 (1991).
111. Ozeki, M. Mater. Sci. Rep. 8(3), 97–146 (1992).
112. Suntola, T. and Antson, J. (1977). US Patent 4,058,430.
113. Ahonen, M., Pessa, M., and Suntola, T. Thin Solid Films 65, 301–307 (1980).
114. Knez, M., Nielsch, K., and Niinisto¨, L. Adv. Mater. 19(21), 3425–3438 (2007).
54
115. Leskela¨, M., Niinisto¨, J., and Ritala, M. In Comprehensive Materials Processing, Hashmi,
S., Batalha, G. F., Tyne, C. J. V., and Yilbas, B., editors, 101–123. Elsevier, Oxford
(2014).
116. Miikkulainen, V., Leskela¨, M., Ritala, M., and Puurunen, R. L. J. Appl. Phys. 113(2),
021301 (2013).
117. Ritala, M. and Leskela¨, M. In Deposition and Processing of Thin Films, Nalwa, H., editor,
volume 1 of Handbook of Thin Film Materials, chapter 2, 103–159. Academic Press (2003).
118. Leskela¨, M. and Ritala, M. Angew. Chem. Int. Ed. 42(45), 5548–5554 (2003).
119. Knoops, H. C. M., Langereis, E., van de Sanden, M. C. M., and Kessels, W. M. M. J.
Electrochem. Soc. 157(12), G241–G249 (2010).
120. Henke, T., Knaut, M., Hossbach, C., Geidel, M., Rebohle, L., Albert, M., Skorupa, W.,
and Bartha, J. W. ECS J. Solid State Sci. Technol. 4(7), P277–P287 (2015).
121. George, S. M. Chem. Rev. 110(1), 111–131 (2010).
122. Hatanpa¨a¨, T., Ritala, M., and Leskela¨, M. Coord. Chem. Rev. 257(23–24), 3297–3322
(2013).
123. Leskela¨, M., Niinisto¨, L., Niemela¨, P., Nyka¨nen, E., Soininen, P., Tiitta, M., and Va¨ha¨kan-
gas, J. Vacuum 41(4-6), 1457–1459 (1990).
124. Asikainen, T., Ritala, M., and Leskela¨, M. Appl. Surf. Sci. 82/83, 122–125 (1994).
125. Pore, V., Ritala, M., and Leskela¨, M. Chem. Vap. Deposition 13(4), 163–168 (2007).
126. Szczerbakow, A., Godlewski, M., Dynowska, E., Ivanov, V. Y., S´wiatek, K., Goldys, E. M.,
and Phillips, M. R. Acta Phys. Pol. A 97(3), 579–582 (1998).
127. Tammenmaa, M., Antson, H., Asplund, M., Hiltunen, L., Leskela¨, M., Niinisto¨, L., and
Ristolainen, E. J. Cryst. Growth 84, 151–154 (1987).
128. Johansson, J., Kostamo, J., Karppinen, M., and Niinisto¨, L. J. Mater. Chem. 12, 1022–
1026 (2002).
129. Suntola, T., Antson, J., Pakkala, A., and Lindfors, S. In SID International Symposium in
San Diego, California, 29 April-1 May 1980, Digest of Technical Papers, 108–109 (SID,
Los Angeles, California, 1980).
130. Tadokoro, T., Ohta, S., Ishiguro, T., Ichinose, Y., Kobayashi, S., and Yamamoto, N. J.
Cryst. Growth 130, 29–36 (1993).
131. Ezhovskii, Y. K. and Klyuikov, A. I. Zr. Prikl. Khim. 73, 881–884 (2000). [Russ. J. Appl.
Chem. 73, 933 (2000)].
132. Yao, T. and Takeda, T. Appl. Phys. Lett. 48(2), 160–162 (1986).
55
133. Guziewicz, E., Godlewski, M., Kopalko, K.,  Lusakowska, E., Dynowska, E., Guziewicz,
M., Godlewski, M. M., and Phillips, M. Thin Solid Films 446, 172–177 (2004).
134. Feng, Y. Z. and Guo, S. P. J. Mater. Sci. Lett. 15, 1824–1827 (1996).
135. Hartmann, J. M., Cibert, J., Kany, F., Mariette, H., Charleux, M., Alleysson, P., Langer,
R., and Feuillet, G. J. Appl. Phys. 80, 6257–6265 (1996).
136. Hartmann, J. M., Charleux, M., Mariette, H., and Rouvie`re, J. L. Appl. Surf. Sci. 112,
142–147 (1997).
137. Herman, M. A., Jylha¨, O., and Pessa, M. J. Cryst. Growth 66, 480–483 (1984).
138. Faschinger, W., Hauzenberger, F., Juza, P., Pesek, A., and Sitter, H. J. Electron. Mater.
22, 497–500 (1993).
139. Hartmann, J. M., Feuillet, G., Charleux, M., and Mariette, H. J. Appl. Phys. 79, 3035–
3041 (1996).
140. Hsu, C.-T. Jpn. J. Appl. Phys., Part 1 35, 4476–4479 (1996).
141. Yokoyama, M., Chen, N. T., and Ueng, H. Y. J. Cryst. Growth 212, 97–102 (2000).
142. Rees, W. S., Green, D. M., Anderson, T. J., Bretschneider, E., Pathangey, B., Park, C.,
and Kim, J. J. Electron. Mater. 21, 361–366 (1992).
143. Koukitu, A., Saegusa, A., Kitho, M., Ikeda, H., and Seki, H. Jpn. J. Appl. Phys., Part 2
31, L2165–L2168 (1992).
144. Bisengaliev, R. A., Novikov, B. V., Aleskovskii, V. B., Drozd, V. E., Ageev, D. A.,
Gubaidullin, V. I., and Savchenko, A. P. Fiz. Tverd. Tela 40, 820–821 (1998). [Phys.
Solid State 40, 754 (1998)].
145. Bhat, I., Ehsani, H., Wang, W. S., Ghandhi, S. K., and Karam, N. H. J. Vac. Sci. Technol.
B 10, 1376–1379 (1992).
146. Karam, N. H., Wolfson, R. G., Bhat, I. B., Ehsani, H., and Gandhi, S. K. Thin Solid
Films 225, 261–264 (1993).
147. Wang, W.-S., Ehsani, H., and Bhat, I. J. Electron. Mater. 22, 873–878 (1993).
148. Emerson, R. M., Hoyt, J. L., and Gibbons, J. F. Appl. Phys. Lett. 65, 1103–1105 (1994).
149. Dakshinamurthy, S. and Bhat, I. J. Electron. Mater. 27(6), 521–526 (1998).
150. Wang, W.-S., Ehsani, H. E., and Bhat, I. B. J. Cryst. Growth 124, 670–675 (1992).
151. Hiltunen, L., Leskela¨, M., Ma¨kela¨, M., Niinisto¨, L., Nyka¨nen, E., and Soininen, P. Thin
Solid Films 166, 149–154 (1988).
152. Kim, H. J. Vac. Sci. Technol. B 21, 2231–2261 (2003).
56
153. Usui, A., Sunakawa, H., Stu¨tzler, F. J., and Ishida, K. Appl. Phys. Lett. 56(3), 289–291
(1990).
154. McDermott, B. T., Reid, K. G., El-Masry, N. A., Bedair, S. M., Duncan, W. M., Yin, X.,
and Pollak, F. H. Appl. Phys. Lett. 56, 1172–1174 (1990).
155. Hirose, S., Yamaura, M., and Munekata, H. Appl. Surf. Sci. 150, 89–94 (1999).
156. Ishii, M., Iwai, S., Kawata, H., Ueki, T., and Aoyagi, Y. J. Cryst. Growth 180, 15–21
(1997).
157. Thompson, P. E., Davis, J. L., Waterman, J., Wagner, R. J., Gammon, D., Gaskill, D. K.,
and Stahlbush, R. J. Appl. Phys. 69, 7166–7172 (1991).
158. Koukitu, A., Nakai, H., Saegusa, A., Suzuki, T., Nomura, O., and Seki, H. Jpn. J. Appl.
Phys. 27(5A), L744 (1988).
159. Ahopelto, J., Kattelus, H. P., Saarilahti, J., and Suni, I. J. Cryst. Growth 99, 550–555
(1990).
160. Bedair, S. M., Tischler, M. A., Katsuyama, T., and El-Masry, N. A. Appl. Phys. Lett.
47(1), 51–53 (1985).
161. Dapkus, P. D., DenBaars, S. P., Chen, Q., Jeong, W. G., and Maa, B. Y. Prog. Cryst.
Growth. Charact. Mater. 19, 137–147 (1989).
162. Ishizaki, M., Kano, N., Yoshino, J., and Kukimoto, H. Jpn. J. Appl. Phys., Part 2 30,
L435–L436 (1991).
163. Kitahara, K., Ohtsuka, N., Ashino, T., Ozeki, M., and Nakajima, K. Jpn. J. Appl. Phys.,
Part 2 32, L236–L238 (1993).
164. Ozeki, M., Haraguchi, T., and Fujita, A. J. Cryst. Growth 298, 90–93 (2007).
165. Fujii, K., Suemune, I., Koui, T., and Yamanishi, M. Appl. Phys. Lett. 60, 1498–1500
(1992).
166. Tran, C. A., Masut, R. A., Brebner, J. L., and Leonelli, R. Appl. Phys. Lett. 62, 2375–2377
(1993).
167. Are`s, R., Hu, J., Yeo, P., and Watkins, S. P. J. Cryst. Growth 195, 234–241 (1998).
168. Nishizawa, J., Abe, H., and Kurabayashi, T. J. Electrochem. Soc. 132(5), 1197–1200
(1985).
169. Usui, A. Proc. IEEE 80(10), 1641–1653 (1992).
170. Sakuma, Y., Ozeki, M., Ohtsuka, N., and Kodama, K. J. Appl. Phys. 68(11), 5660–5664
(1990).
171. Are`s, R., Watkins, S. P., Yeo, P., Horley, G. A., O’Brien, P., and Jones, A. C. J. Appl.
Phys. 83(6), 3390–3397 (1998).
57
172. Mori, K., Yoshida, M., Usui, A., and Terao, H. Appl. Phys. Lett. 52(1), 27–29 (1988).
173. Kobayashi, R., Jin, Y., and Hasegawa, F. J. Cryst. Growth 113, 491–498 (1991).
174. Maa, B. Y. and Dapkus, P. D. Appl. Phys. Lett. 58(20), 2261–2263 (1991).
175. Chen, Q., Beyler, C. A., Dapkus, P. D., Alwan, J. J., and Coleman, J. J. Appl. Phys. Lett.
60(19), 2418–2420 (1992).
176. de Keijser, M. and van Opdorp, C. Appl. Phys. Lett. 58(11), 1187–1189 (1991).
177. Kobayashi, R., Ishikawa, K., Narahara, S., and Hasegawa, F. Jpn. J. Appl. Phys., Part 2
31, L1730–L1732 (1992).
178. Doi, A., Aoyagi, Y., and Namba, S. Appl. Phys. Lett. 49(13), 785–787 (1986).
179. Wells, R. L., Pitt, C. G., McPhail, A. T., Purdy, A. P., Shafieezad, S., and Hallock, R. B.
Chem. Mater. 1(1), 4–6 (1989).
180. Wells, R. L. and Gladfelter, W. L. Journal of Cluster Science 8(2), 217–238 (1997).
181. Schulz, S., Martinez, L., and Ross, J. L. Adv. Mater. Opt. Electron. 6(4), 185–189 (1996).
182. Janik, J. F., Wells, R. L., Jr., V. G. Y., Rheingold, A. L., and Guzei, I. A. J. Am. Chem.
Soc. 120(3), 532–537 (1998).
183. Foos, E. E., Jouet, R. J., Wells, R. L., Rheingold, A. L., and Liable-Sands, L. M. Journal
of Organometallic Chemistry 582(1), 45 – 52 (1999).
184. Hatanpa¨a¨, T., Pore, V., Sarnet, T., Ritala, M., and Leskela¨, M. In ALD2010. Proceedings
of the 10th International Conference on Atomic Layer Deposition, M8, (2010).
185. Huheey, J. E., Keiter, E. A., and Keiter, R. L. Inorganic Chemistry: Principles of Structure
and Reactivity. Harper Collins College Publishers, New York, 4th edition, (1993).
186. Eom, T., Gwon, T., Yoo, S., Choi, B. J., Kim, M.-S., Buchanan, I., Xiao, M., and Hwang,
C. S. Chem. Mater. 26(4), 1583–1591 (2014).
187. Eom, T., Gwon, T., Yoo, S., Choi, B. J., Kim, M.-S., Buchanan, I., Ivanov, S., Xiao, M.,
and Hwang, C. S. Chem. Mater. 27(10), 3707–3713 (2015).
188. Knapas, K., Hatanpa¨a¨, T., Ritala, M., and Leskela¨, M. Chem. Mater 22(4), 1386–1391
(2010).
189. Waldo, R. In Microbeam Analysis. Proceedings of the 23rd Annual Conference of the
Microbeam Analysis Society, Newbury, D., editor, 310. San Francisco Press, (1988).
190. Ma¨ntyma¨ki, M., Heikkila¨, M. J., Puukilainen, E., Mizohata, K., Marchand, B., Ra¨isa¨nen,
J., Ritala, M., and Leskela¨, M. Chem. Mater. 27(2), 604–611 (2015).
191. Raoux, S., (2012). personal communication.
58
192. Raoux, S., Mun˜oz, B., Cheng, H.-Y., and Jordan-Sweet, J. L. Appl. Phys. Lett. 95(14),
143118 (2009).
193. Nilsen, O., Karlsen, O. B., Kjekshus, A., and Fjellv˚ag, H. Thin Solid Films 515(11),
4527–4537 (2007).
194. Nilsen, O., Karlsen, O. B., Kjekshus, A., and Fjellv˚ag, H. Thin Solid Films 515(11),
4538–4549 (2007).
195. Nilsen, O., Karlsen, O. B., Kjekshus, A., and Fjellv˚ag, H. Thin Solid Films 515(11),
4550–4558 (2007).
196. Ritala, M., Pore, V., Hatanpa¨a¨, T., Heikkila¨, M., Leskela¨, M., Mizohata, K., Schrott, A.,
Raoux, S., and Rossnagel, S. M. Microelectron. Eng. 86(7–9), 1946–1949 (2009).
197. Kadel, K., Kumari, L., Li, W., Huang, J. Y., and Provencio, P. P. Nanoscale Res. Lett.
6, 57 (2011).
198. Qiu, X., Austin, L. N., Muscarella, P. A., Dyck, J. S., and Burda, C. Angew. Chem. Int.
Ed. 45(34), 5656–5659 (2006).
199. Giani, A., Bayaz, A. A., Foucaran, A., Pascal-Delannoy, F., and Boyer, A. J. Cryst.
Growth 236(1-3), 217–220 (2002).
59