Chimie des Interactions Plasma-Surface (ChIPS), Faculty of Sciences
Ph.D. Thesis
Oxygen vacancy stabilized zirconia;
synthesis and properties
Mohsin Raza
Chimie des Interactions Plasma-Surface (ChIPS)
Chimie des Interactions Plasma-Surface (ChIPS), Faculty of Sciences
University of Mons
BE-7000 Mons, Belgium
Mohsin Raza
Printed by UMons, March 2017
Mons, Belgium
Chimie des Interactions Plasma-Surface (ChIPS), Faculty of Sciences
Ph.D. Thesis
Oxygen vacancy stabilized zirconia;
synthesis and properties
Mohsin Raza
Chimie des Interactions Plasma-Surface (ChIPS)
16 March 2017
Promoter and co-promoter Stéphanos Konstantinidis and Rony Snyders
Jury Members
Prof. Dr. Jochen M. Schneider (RWTH Aachen University, Germany)
Prof. Dr. Jean F. Pierson (University of Lorraine, France)
Prof. Dr. Roberto Lazzaroni (University of Mons, Belgium)
Dr. Thomas Godfroid (Materia Nova, Belgium)
Dr. Jerome Cornil (University of Mons, Belgium)
Chimie des Interactions Plasma-Surface (ChIPS), Faculty of Sciences
University of Mons
BE-7000 Mons, Belgium
i
Abstract
The crystal structure of any material is a decisive factor for controlling its properties1,2. In
this respect, zirconia (ZrO2) is a material which exists in three crystallographic phases: the
monoclinic phase which is stable up to 1205 °C; the tetragonal phase appears from 1205 °C
to 2377 °C; and finally the cubic phase is thermodynamically stable from 2377 °C to 2710
°C3. Because of the superior chemical stability, high hardness, high dielectric constant and
prominent optical properties, zirconia can be exploited for a broad range of applications such
as medical implants, oxygen detectors and as wear resistant or thermal barrier coatings
(TBCs). However, for pure zirconia, it is not possible to exploit most of the above-mentioned
applications as this is restricted by the change in volume of the zirconia-based components
due to the phase transformation upon heating and cooling, which ultimately leads to the
deterioration of the device components.
Therefore, for decades and whatever the zirconia is synthesized as a bulk material or as a thin
film coating, the cubic phase is stabilized by doping zirconia with cations such as Y (yttrium).
1 A. Van De Walle, A complete representation of structure-property relationships in crystals., Nat. Mater. 7
(2008) 455–458. doi:10.1038/nmat2200
2 G.B. Olson, Computational Design of Hierarchically Structured Materials, Science (80-. ). 277 (1997) 1237–
1242. doi:10.1126/science.277.5330.1237.
3 J.P. Abriata, J. Garcés, R. Versaci, The O-Zr (Oxygen-Zirconium) system, Bull. Alloy Phase Diagrams. 7
(1986) 116–124. doi:10.1007/BF02881546.
ii
Doing so, not only stabilizes the cubic phase but also leads to the generation of oxygen
vacancies in the zirconia lattice. This situation makes Yttria Stabilized Zirconia (YSZ) useful
for TBCs but also as an electrolyte with high ionic conductivity to be used in Solid Oxide
Fuel Cells (SOFCs). However, such stabilization strategies lead to the perturbation of the
periodic potential of the oxide-ion array, which results in higher energy barrier for O2- ions
during their diffusion to a vacant site as compared to intrinsic vacancy-doped oxides 4 .
Therefore, an intense research has been developed during the last fifteen years to promote
stabilization without incorporating aliovalent cations. Stabilization phenomena for
nanometer thick zirconia films have been related to the grain size, energy input during
growth, stresses in the zirconia film and oxygen vacancy or nitrogen atom incorporation. So
far, a consensus over what drives the phase formation and stabilization of cubic zirconia has
not been reached.
In the present study, according to quantum-chemistry based calculations, It is shown that the
cubic phase is the most stable phase if more than 3 at.% of oxygen vacancies are incorporated
in the ZrO2 lattice. On the other hand, carefully designed cold plasma-based reactive
magnetron deposition experiments allowed to control the amount of O vacancies incorporated
inside the zirconia lattice. X-Ray diffraction analysis of these oxygen vacancy doped zirconia
thin films are in remarkable agreement with theoretical predictions, hence emphasizing that
the incorporation of oxygen vacancies is the sole responsible mechanism for the stabilization
of the zirconia cubic phase. However, it is also observed that any deviation from the
optimized synthesis conditions leads to the change in film phase constitution.
4 J.B. Goodenough, Ceramic technology: Oxide-ion conductors by design, Nature. 404 (2000) 821–823.
doi:10.1038/35009177.
iii
In a second step, the thermal stability of the oxygen vacancy stabilized zirconia thin films is
addressed. X-Ray diffraction experiments, show that these films are stable up to 750 °C. The
ionic and electronic conductivity measurements highlight that the ionic conductivity is in the
range of YSZ however a colossal increase in the ionic conductivity is observed when the film
thickness is ≈ 10 nm i.e. 7.4 S.cm-1 at 725 °C. However, these coatings are electronically
non-conductive. These results highlight the possibility to use oxygen vacancy doped zirconia
films in SOFC devices. Finally, the photoluminescence analysis of OVSZ revealed the
presence of 2 emission peaks centered at 388 nm and 488 nm originating from the states
present in the band gap. This result validates the theoretical calculations data predicting the
appearance of energy states in the band gap upon O vacancy incorporation.
In conclusion, it is demonstrated that the incorporation of oxygen vacancies, and therefore
material defect chemistry plays a very important role in the phase formation and properties
of reactively sputtered thin films, particularly in zirconia. This observation could pave the
way to the development of new thin film growth strategies or to the synthesis of functional
thin films with new or enhanced properties.
iv
v
Acknowledgements
I am grateful to a large number of people who have contributed, helped and supported me in
this entire thesis work.
Stéphanos Konstantinidis and Rony Snyders, thank you very much for giving me the
opportunity to perform this thesis work, for your guidance at every step, encouragement, time
and support. I have learnt a lot from you in these last 4 years. I am very fortunate to have you
as my promoter. Thank you so much for all of this!
All the collaborators, I am very thankful to all of you for helping me out and performing the
requested measurements and calculations.
All the members of Chimie des Interactions Plasma-Surface (ChIPS), thank you all for
providing me such a nice working environment. Mattia, Claudia and Nikolay many thanks
for providing good company in and off the lab. Sabine and Dany thank you for taking care of
administrative stuff each and every time.
My friends, thank you all for being part of my life, for your support and encouragement.
My Siblings, thank you all for making my life wonderful. I can never forget the time which
we had together. I love you all, a lot!
My Parents, who pray for my success day and night, I could not have achieved this without
your support, love, and encouragement. I am very thankful for your every effort, you put to
give me a good life and to make me a good person. I love you very much!
vi
vii
Table of Contents Abstract i
Acknowledgements v
Table of Contents vii
1. Introduction 1
1.1. Background 1
1.2. Aim of the research work and strategy 5
References 7
2. Zirconia 9
2.1. The Crystal structure of Zirconia (ZrO2) 9
2.2. Stabilization strategies for zirconia c-phase 10
2.2.1. Stabilization of Zirconia c-phase by doping 11
2.2.2. Stabilization of high temperature phase of Zirconia without doping 13
2.3 Properties and applications of zirconia 22
2.3.1 Mechanical properties of zirconia 23
2.3.2 Thermal properties of zirconia 24
2.3.3 Ionic conduction of stabilized zirconia 24
2.3.4 Optical properties of Zirconia 27
References 28
3. Thin film growth 31
3.1. Sputtering 32
3.1.1. Glow discharge 34
3.1.2. Magnetron sputtering 38
3.1.3. Reactive magnetron sputtering 41
3.2. Thin film growth 44
3.2.1. Early stages of thin film formation 45
3.2.2. 3D thin film growth 48
viii
3.2.3. Microstructure of thin film 50
3.3. Influence of energetic species on thin film properties during sputtering 53
3.4. Formation and evolution of stresses in thin films 56
References 60
4. Zirconia thin film deposition and characterization 63
4.1 Modelling and computational details 64
4.2 Thin film deposition and process monitoring 65
4.2.1 Voltage feedback control unit 68
4.3 Film characterization tools 70
4.3.1 X-ray diffraction (XRD) 70
4.3.1.1 Bragg-Brentano (θ-2θ) mode 71
4.3.1.2 Grazing incidence XRD (GIXRD) mode 72
4.3.2 Transmission electron microscopy (TEM) 72
4.3.3 Secondary electron microscopy (SEM) 73
4.3.4 Chemical composition of zirconia thin films 74
4.3.4.1 Rutherford backscattering spectroscopy (RBS) 74
4.3.4.2 Nuclear reaction analysis 75
4.3.5 Electrochemical Impedance spectroscopy (EIS) 76
4.3.6 Photoluminescence (PL) 77
4.3.7 Heat flux microsensor 79
References 81
5. Influence of oxygen vacancies on the phase constitution of zirconia thin films 83
5.1 Phase stability of oxygen deficient zirconia; quantum chemistry based DFT calculations84
5.2 Synthesis of ZrO2-x by reactive magnetron sputtering 86
5.3 Conclusion 90
References 91
6. Influence of the deposition parameters on the phase constitution of oxygen vacancy doped
thin films 93
6.1 Influence of pressure and discharge current 94
6.1.1 Experimental details 94
6.1.2 Results and discussion 95
6.1.2.1 Evolution of the phase constitution as a function of pressure and discharge current
95
6.1.2.2 Normalized energy flux measurements 97
6.2 Influence of film thickness on the crystal structure of OVSZ films 102
ix
6.2.1 Experimental details 102
6.2.2 Results and discussion 102
6.2.2.1 Evolution of XRD diffractograms as a function of film thickness 102
6.2.2.2 Analysis of film cross-section by SEM and HRTEM 104
6.3 Conclusions 108
References 110
7. Thermal stability of OVSZ thin films 111
7.1 Experimental details 112
7.2 Results and discussion 113
7.3 Conclusion 124
References 126
8. Ionic conductivity of OVSZ thin films 127
8.1 Experimental details 128
8.2 Results and discussion 128
8.3 Conclusion 137
References 138
9. Optical properties of OVSZ thin films 139
9.1 Experimental details 140
9.2 Results and discussion 140
9.3 Conclusion 146
References 148
Outlook 149
Appendix 154
1
1. Introduction 1.1. Background
Thin films are layers of materials whose thickness ranges from a few monolayers to few
micrometers. These films sometimes exhibit unique properties that cannot be observed in
bulk materials. The history of making thin films dates back to the metal ages and especially
to the ancient craft of gold beating. In ancient times, i.e. more or less 5000 years ago, the
Egyptians were the first to practice this art [1] to make decorative gold leafs. Today thin films
are still being used for many decorative proposes but also to enhance the surface properties
of materials as well as to build functional devices in various application fields such as optics,
mechanics, microelectronics, sensors, energy production such as solar energy, solid oxide
fuel cells.
Obviously, the properties of a film and thereby its area of application is mainly determined
by the elemental composition of the material. But, besides the elemental composition of the
coating, the spatial arrangement of the film forming atoms, i.e. the crystal structure, is also a
decisive factor [2,3]. One example of such is zirconia (ZrO2). Zirconia is a polymorphous
material which exists in three crystallographic phases at atmospheric pressure: (i) the
monoclinic phase (m, space group P21/c) stable up to ~ 1205 °C; (ii) the tetragonal phase (t,
space group P42/nmc) appears from ~ 1205 °C to 2377 °C; and finally (iii) the cubic phase
(c, space group Fm-3m) from 2377 °C to 2710 °C (melting temperature)[4]. Since ZrO2
2
exhibits high chemical stability, high hardness[5], high dielectric constant[6] and prominent
optical properties[7], ZrO2 films have been exploited for a broad range of applications e.g.,
medical applications[8,9], wear resistant coatings[10] and for thermal barrier coatings
(TBCs) [11,12]. However, in the case of pure zirconia, it is not possible to exploit most of
the above mentioned applications as this is restricted by the change in volume of the zirconia-
based components (~5 vol.%) due to the phase transformation upon heating and cooling of
the device, which ultimately leads to the deterioration of the device components [13,14]. To
overcome this situation, the understanding of the synthesis process and the origin of the
properties of interest must be unraveled. An example of the latter is the stabilization of the
high temperature c-phase of zirconia at room temperature with the help of aliovalent dopants
(e.g. Y3+ or Mg2+)[15]. By adding around 12 mol. % of yttria (Y2O3), the c-phase of zirconia
is found to be stabilized at room temperature. This material is known as yttria-stabilized
zirconia (YSZ)[15]. In YSZ, some zirconium (Zr4+) ions are replaced by yttrium (Y3+) so that
to maintain the charge neutrality, for two substituting yttrium cations, one oxygen vacancy is
created. This makes YSZ not only useful for TBCs but also as an electrolyte membrane in
solid oxide fuel cells (SOFC)[16–18] and in oxygen sensors[19] because of its very good
ionic conductivity in between 600 °C – 1000 °C [20].
However, it has also been found that the doping by aliovalent cations leads to the perturbation
of the periodic potential of the oxide-ion array, which results in higher energy barrier for O2-
ions during their diffusion to a vacant site in the solid as compared to intrinsic vacancy-doped
oxides[21]. Therefore, to stabilize high temperature c-phases of zirconia at room temperature
without any doping of yttria, an intense research has been developed during the last 15 years
using various thin film deposition techniques. Among the various thin films growth methods,
3
the most common are electroplating[22], chemical vapor deposition (CVD)[23] and physical
vapor deposition (PVD)[24].
Electroplating is one of the simplest thin film deposition techniques and is very common in
the industry due to its cost effectiveness. In such technique, the object to be coated (the
substrate) is immersed into the solution in which the metal is dissolved. Then negative
potential is applied to the object. In this way, the positive metal ions are attracted to the
surface of the object and a coating is formed. Unfortunately, it’s hard to produce high quality,
dense and defect-free, coatings by such method. This technology also produces a
considerable amount of hazardous wastes.
Chemical vapor deposition (CVD) is a family of processes in which volatile gasses
(precursors) are let into the deposition chamber and the coating is grown on a heated substrate
by a chemical reaction(s) occurring on or in its vicinity. This technique is also widespread in
the industry because of its high deposition rates. It also provides homogeneous films on
complex shaped surfaces and it offers the possibility to deposit a great variety of films, from
metals to organic compounds, with a high control on film composition. The major drawback
of CVD is that most of the time very high temperatures are needed to make the different
constituents react in the gas phase or on the substrate. This means that CVD cannot be applied
to temperature sensitive substrates (e.g. plastic) and also in areas where the thermal expansion
coefficients of films and substrates are different e.g. on steel and electrical components. In
some cases, such limitations can be (partially) overcome by using a plasma to activate the
process. In this case, the process is known as plasma-enhanced CVD.
Physical vapor deposition (PVD) relates to a family of processes in which the film is
deposited on a substrate by the condensation of a vaporized material generated from a solid
4
or liquid source (called target). One such method of PVD is DC sputter deposition. In DC
sputter deposition, the target is bombarded by ions generated in a plasma. The bombardment
of the target induces the ejection (sputtering) of the atoms from the target surface. These
atoms are then transported to the substrate where they condensate and form a film. This
process is generally based on the diode configuration with facing electrodes constituting the
anode and cathode, the plasma being located in between. In such an arrangement, the cathode
plays the role of the target to be sputtered. The disadvantage of DC sputter deposition is that
the plasma is not confined efficiently close to the sputtering region[24] and require sputtering
pressure from several tenths of mTorr to 100 mTorr. One way to confine plasma near the
target and to lower the sputtering pressure is by placing a pair of concentric, permanent,
magnets behind the target. This method is known as DC magnetron sputtering. The
advantages of this configuration are the reduced background gas pressure, the very good
quality of the films, and a better control over the deposition parameters. In such configuration
one can deposit elemental, alloyed, or compound films depending on the target composition.
Compound film materials such as Metal-Nitrides or Metal-Oxides are also synthesized by
adding a reactive gas such as N2 or O2 to the argon carrier gas, in the deposition chamber. In
this case, the technique is called DC reactive magnetron sputtering. The deposition of zirconia
(ZrO2) thin films can be achieved using such kind of process.
In the field of thin film synthesis, in particular in the case of PVD and cold-plasma based
magnetron sputtering methods, several mechanisms have been proposed for the stabilization
of the cubic phase of zirconia without the use of any aliovalent ions. The stabilization is either
related to the grain size[13,25], the energy input during film growth[26], to stresses in the
film[27,28] or to the presence of O vacancies (Vo) or N atoms in the lattice[29]. But it appears
5
that a consensus over what is the mechanism(s) responsible for the stabilization in c-phase
formation has not been reached so far.
1.2. Aim of the research work and strategy
The aim of this thesis work is to understand what drives the synthesis and stabilization of the
high temperature c-phase of zirconia at room temperature. For this purpose, the impact of
oxygen vacancy doping on the phase formation will be first investigated using quantum
chemistry based calculations. Then, a well optimized deposition procedure, based on reactive
magnetron sputtering, will be utilized to synthesize vacancy doped zirconia films. The
deposition parameters (i.e. pressure, discharge current leading to the change in energy
deposited during the film growth) will be varied systematically in order to determine their
influence on the c-phase formation, without doping of aliovalent atoms. In the last step,
thermal stability, the ionic conduction and optical characteristics of the developed material
will be measured and discussed.
This manuscript is organized as follows. Following the introductory chapter, a general
overview of zirconia crystal structure, c-phase stabilization mechanisms, zirconia properties
and applications is given in the chapter 2. In Chapter 3, an overview of the reactive magnetron
sputtering thin film deposition technique is given. In chapter 4, details about the theoretical
calculations and the deposition and film characterization experiments are given. Moreover,
the way that the oxygen vacancy doped zirconia thin films are obtained, are presented in
details. The chapter 5 describes the influence of O vacancies on the phase constitution of
zirconia and the chapter 6 describes influence of deposition parameters on the phase
formation of zirconia thin films and the influence of film thickness. Chapter 7 consists on the
thermal stability of oxygen vacancy stabilized cubic zirconia (OVSZ) thin films. Finally, the
6
properties of the OVSZ films are investigated. Ionic conductivity and photoluminescence
measurements are presented in chapter 8 and 9, respectively.
7
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[2] A. Van De Walle, A complete representation of structure-property relationships in crystals., Nat. Mater. 7 (2008)
455–458. doi:10.1038/nmat2200.
[3] G.B. Olson, Computational Design of Hierarchically Structured Materials, Science (80-. ). 277 (1997) 1237–
1242. doi:10.1126/science.277.5330.1237.
[4] J.P. Abriata, J. Garcés, R. Versaci, The O-Zr (Oxygen-Zirconium) system, Bull. Alloy Phase Diagrams. 7 (1986)
116–124. doi:10.1007/BF02881546.
[5] R.C. Garvie, R.H. Hannink, R.T. Pascoe, Ceramic steels?, Nature. 258 (1975) 703.
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considerations, J. Appl. Phys. 89 (2001) 5243. http://dx.doi.org/10.1063/1.1361065.
[7] Q. Zhang, J. Shen, J. Wang, G. Wu, L. Chen, Sol-gel derived ZrO 2 -SiO 2 highly reflective coatings, Int. J.
Inorg. Mater. 2 (2000) 319. http://dx.doi.org/10.1016/S1466-6049(00)00037-4.
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applications., J. Dent. 35 (2007) 819–26. doi:10.1016/j.jdent.2007.07.008.
[10] X. Zhou, I. Balachov, D.D. Macdonald, The effect of dielectric coatings on IGSCC in sensitized type 304 SS in
high temperature dilute sodium sulfate solution, Corros. Sci. 40 (1998) 1349. http://dx.doi.org/10.1016/S0010-
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Mater. Rev. 50 (2005) 20.
[14] J. Chevalier, L. Gremillard, A. V. Virkar, D.R. Clarke, The tetragonal-monoclinic transformation in zirconia:
Lessons learned and future trends, J. Am. Ceram. Soc. 92 (2009) 1901–1920. doi:10.1111/j.1551-
2916.2009.03278.x.
[15] H.G. Scott, Phase relationships in the zirconia-yttria system, J. Mater. Sci. 10 (1975) 1527–1535.
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3.00007-7.
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Gel Sci. Technol. 27 (2003) 119–136.
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magnetron sputtering of transition metal oxide films, Appl. Phys. Lett. 85 (2004) 748. doi:10.1063/1.1777412.
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of ZrO 2 coatings deposited by reactive pulsed magnetron sputtering (PMS), Thin Solid Films. 377–378 (2000)
37–42.
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8
9
2. Zirconia 2.1. The Crystal structure of Zirconia (ZrO2)
Zirconia (ZrO2) is a polymorphous material. At atmospheric pressure it can be found in three
fluorite related crystallographic phases; (i) the monoclinic (m), occurs naturally as a mineral
Baddeleyite and is stable up to ~ 1205 °C; (ii) the tetragonal phase (t) appears from ~ 1205
°C to 2377 °C; and finally (iii) the cubic phase (c) from 2377 °C to 2710 °C (melting
temperature)[1]. The unit cells related to m-, t-, and c-zirconia are presented in Fig.2.1 and
their characteristic crystallographic properties on Table 2.1.
Fig. 2. 1: Crystal structure of Zirconia (a) monoclinic, (b) tetragonal, (c) cubic
Table 2. 1 Crystal structure of pure zirconia (ICDD PDF # 013 0307; 010 070 6627; 00 049 1642).
Temperature
(°C) Phase
Space
group
Space
group number
Cation
coordination number
Cell parameters
a(Å) b(Å) c(Å) α(°) β(°) γ(°)
0-1205 Monoclinic P21/c 14 7 5.1477 5.2030 5.3156 90.00 99.38 90.00
1205-2377 Tetragonal P42/nmc 137 8 3.5948 3.5948 5.1824 90 90 90
2377-2710 Cubic Fm-3m 225 8 5.1280 5.1280 5.1280 90 90 90
10
Zirconia high temperature c-phase in general has superior properties over the room
temperature stable m- and t-phase because of highly symmetric atomic arrangement.
Therefore, for most of the engineering applications, zirconia c-phase is preferred. However,
this is limited by the instability of cubic zirconia at room temperature. Further limitation in
the use of zirconia for high temperature applications comes from the change in its unit cell
volume as a function of temperature [2], as shown in Fig 2.2. It could be observed the unit
cell volume of zirconia first increases up-to 1400 K and then, due to the phase transformation,
the unit cell volume decreases abruptly. This change in volume during the increase in
temperature as well during the cooling results in the formation of cracks in the coating, thus
making it un-useful in such circumstances.
Fig. 2. 2: Variation of the Zirconia unit cell volume of m, t, and c-phase as a function of temperature,
adopted from [2]. The t-phase here is described in terms of face-centered tetragonal cell of unit cell
parameter á. The usual crystallographic has a = a/√2 [2].
2.2. Stabilization strategies for zirconia c-phase
Two strategies are found in the literature regarding the stabilization of the high temperature
cubic phase of zirconia at room temperature. One is to introduce a foreign element (doping)
11
in the zirconia lattice. The second is related to a mechanism that does not involve any doping
process Here (in section 2.2.1 and 2.2.2) both approaches are presented.
2.2.1. Stabilization of Zirconia c-phase by doping
Fluorite oxides (MO2) are the oxides which exhibit cubic crystal structure consisting of a
cubic oxygen lattice with alternate body centers occupied by eight coordinated cations. The
cations are arranged into a face centered cubic structure with the anions occupying the
tetrahedral sites. Usually the size of these tetravalent cation (M) is big enough to sustain the
cubic (fluorite) structure. However, in case of zirconia (ZrO2), the size of Zr4+ is too small to
sustain the cubic structure at low temperatures. Therefore, to stabilize the cubic structure, it
has to be partly substituted with a larger cation than Zr4+. For this purpose, since more than
ninety years, the stabilization of the high temperature c-phase of zirconia at room temperature
is achieved by doping zirconia (ZrO2) by larger cations of lower valence than Zr4+, e.g., by
incorporating Y3+, Ca2+, or Mg2+ in the ZrO2 lattice. This strategy was first reported by Ruff
et al.[3] in 1929. In the beginning, due to the lack of advanced measuring equipment, the
dopant values reported for the stabilization of c-phase were a bit higher than the recently
published values. It has been reported by Gaudon, et al.[4] that by doping around and above
7 mol% of yttria (Y2O3), the c-phase of zirconia can be stabilized at room temperature, as
shown in Fig.2.3. This material is known as yttria stabilized zirconia (YSZ). On the other
hand, doping by 2-7 mol% of yttria, leads to the stabilization of the tetragonal phase also
known as partially stabilized zirconia. Below 2 mol% of yttria doping, the monoclinic phase
is present.
12
Fig. 2. 3: Phase diagram of Zirconia-yttria system. m, t and c represent the monoclinic, tetragonal and
cubic phase respectivley[4].
In YSZ, on addition of Y3+ in the zirconia lattice, Zr4+ cations are replaced by Y3+. To
maintain the charge neutrality, for each two substituting yttrium cations, one oxygen vacancy
is created. This process is summarized by the eq. (1 and 2) using Kröger-Vink notation and
is schematically illustrated in Fig. 2.4.
𝑌2𝑂3 𝑍𝑟𝑂2→ 2𝑌′𝑍𝑟 + 𝑉𝑂
.. + 3𝑂𝑂𝑥 (1)
𝑂𝑂𝑥 ↔ 𝑉𝑂
.. + 2𝑒 +1
2𝑂2 (2)
𝑌′𝑍𝑟 represent the Y in the Zr site with the apparent negative charge, 𝑉𝑂.. is the vacancy in the
oxygen site with double positive charge, 𝑂𝑂𝑥 is the lattice oxygen, i.e., oxygen in the oxygen
site with net charge of zero.
13
Fig. 2. 4: Schematic representation of the insertion of Yttria (Y2O3) in the zirconia (ZrO2) lattice and
the creation of oxygen vacancies. Image adopted from [5].
2.2.2. Stabilization of high temperature phase of Zirconia without doping
In order to stabilize t- and c-phase of zirconia at room temperature, especially in case of
zirconia thin films, several mechanisms have been proposed so far. The stabilization has been
attributed to the grain size [6–8], energy input during the film growth [9–11], O vacancies
and/or incorporation of N atoms in the lattice [12–14] and has been also related to the stresses
in the film [8,9,11]. These mechanisms are briefly presented here.
Grain-size effect
The effect of grain size on stabilizing the zirconia t-phase at room temperature without any
dopants was first reported by Gravie [15] in 1965. In their study they reported that 100% t-
ZrO2 can be achieved at room temperature by having a grain size in the range of 11-17 nm.
However, if the grain size ranges from 17 nm to 30 nm, a mixture of tetragonal and
monoclinic phases was found. It is concluded in their study that a critical size exists for the
14
stabilization of the metastable tetragonal phase, which is found to be 30 nm. Further it has
been also reported by Chen et al. [16] that the tetragonal phase of zirconia can also be
stabilized for temperature in the order of 700 °C (as shown in Fig. 2.5). This is a much lower
temperature range than what is reported on the phase diagram (see Fig. 2. 3) where the
temperature ranges from 1205 °C to 2377 °C. According to these authors, such a lowering of
the temperature at which the t-ZrO2, but also the metastable cubic phase, is obtained is
achieved by applying an external hydrostatic pressure on powdered zirconia. This is due to
the densities of tetragonal, orthorhombic, and cubic crystal structures which are higher than
that of the monoclinic structure. The application of an external hydrostatic pressure converts
the monoclinic structure, which is stable at lower temperatures, into a denser structure.
Further a relation between the grain size and phase transformation temperature was also
found by Garvie et al. [17,18]. Garvie et al. shows the phase transformations temperature
Fig. 2. 5: Schematic of pressure-temperature phase diagram of bulk ZrO2. Image adopted from [19].
15
can be lowered by reducing the grain size[17,18], therefore it’s possible to obtain the
tetragonal and cubic phases at room temperature instead of Orthorhombic-I, -II by applying
an external pressure. Shukla et al.[7] have reported in their study that an external hydrostatic
pressure is required to stabilize the tetragonal/cubic phase at room temperature for zirconia
grain size greater than 10 nm. However, in the case of a grain size less than 10 nm, there is
no need to apply external hydrostatic pressure. Therefore, there must be sufficient hydrostatic
pressure acting inside the grain to stabilize the tetragonal and cubic phases at room
temperature in this situation. In the case of liquid particle of radius (r); the magnitude of the
hydrostatic pressure (∆P) can be calculated from Gibbs-Thompson equation
∆𝑃 =2𝛾
𝑟(3)
where γ is the surface tension. In the case of solids, the surface tension has to be replaced by
the surface stress f, which is given by the expression (4) where ε is the strain.
𝑓 = 𝛾 +𝑑𝛾
𝑑𝜀(4)
Since in solids, internal hydrostatic pressure is due to the surface stress which is of the order
of surface energy. Therefore, for spherical isotropic particles dγ/dε can be neglected. Thus
the surface tension γ can be substituted by surface stress f in in equation (3)
∆𝑃 =2𝑓
𝑟(5)
The magnitude of f has been estimated by Winterer et al. for nano-crystallite zirconia powder
with a grain size ranging from 5-30 nm and is equal to 5 Nm-1[20]. Substituting this value in
equation 5 and calculating the internal hydrostatic pressure for grains of 10 nm and below
reveals that the internal hydrostatic pressure increases with the decrease in grain size and is
as high as 25 GPa for a grain-size of 1 nm, (Fig. 2.6). Further, Winterer et al.[20] estimated
16
the grain size for the stabilization of tetragonal phase. This critical grain size is equal to 8-12
nm. The later value indicates that the internal hydrostatic pressure should be about 2.5 GPa.
On the other hand, to stabilize cubic phase of zirconia, a slightly larger internal hydrostatic
pressure (~30 GPa) is required. Nitsche et al.[21,22] also synthesized nanosized zirconia
powder characterized by grain size from 5-30 nm using a gas condensation technique.
However, from HRTEM analysis, they found that the size of the zirconia grains lies in the
range of 7-32 nm and the grains exhibit a core-shell morphology with the tetragonal phase as
the core and the monoclinic phase as the shell.
Fig. 2. 6: Variation in internal hydrostatic pressure as a function of Nano-crystallite (grain) size,
calculated using equation (5), image adopted from [7].
Influence of energy deposition and stresses on phase stabilization
In a plasma based process, a film grown on the substrate surface is subjected to intense
particle bombardment. The impinging species can be among others, plasma ions (e.g. Ar+,
O2-, Zr+), electrons, photons, condensing metal atoms, neutral particles [23]. The flux of each
of these particles will influence to amount of energy delivered to the film during growth. The
17
influence of energy input on stabilizing high temperature cubic phase of zirconia was first
noticed by Goedicke et al.[9] in 2000, for films of 100-150 nm thickness synthesized by
reactive pulsed magnetron sputtering (PMS) at various deposition pressure. Goedicke et al.
varied the working pressure from 0.3 Pa to 3.5 Pa as well as the target to substrate distance.
During the deposition, they did not supply any intentional heating to the substrate. However,
they observed an increase of maximum 30 °C in the substrate temperature during the
deposition by using thermos strips. In their study, when the films were deposited at low
pressure i.e. 0.3 Pa, films exhibited pure monoclinic phase. On the other hand, when films
were deposited at 3.5 Pa, they exhibited pure cubic phase as shown in Fig. 2.7. The reason
Goedicke et al. propose for such a behavior is related to the energy of the condensing
particles. They suggested that higher sputtering pressure reduces the mean free path of the
sputtered particles and therefore the mobility of the depositing species on the substrate
surface. They further relate lower mobility to microporosity of the film and therefore reduced
hardness and lower compressive stresses in the film. Goedicke et al. also measured residual
stresses in their study and found that, indeed, films deposited at low sputtering pressure (0.3
Pa) exhibit high compressive stresses (~ 1800 MPa) while the films deposited at higher
sputtering pressure (3.5 Pa) exhibit low tensile stresses (~139 MPa).
18
Fig. 2. 7: Diffractograms of ZrO2 thin films deposited by PMS at sputtering pressure of (a) 0.3 Pa,
(b) 3.5 Pa. Image adopted from [9].
In reactive magnetron sputtering having O2 as a reactive gas, an important feature is the
emission of negatively charged oxygen ions O- from the oxidized part of the target
surface[11,24–26]. These O- ions are accelerated in the cathode sheath. They bombard the
growing film with energies higher than the other depositing plasma species, in the range of
several hundreds of eV [26]. The energy of these O- ions is correlated to the magnitude of the
target voltage, while the number of these O- ions is determined by portion of the target
covered with the oxide layer. In several studies, the influence of these O- ions on the
19
formation of crystalline structures has been reported. For example, in 2006, Mraz et al.[27]
studied the influence of the energy brought by these oxygen ions emitted from the oxidized
target during the growth of transition metal oxides (Nb, Ta, Zr, and Hf) films by reactive
magnetron. These authors proposed that the evolution of the crystalline structure of transition
metal oxide thin films may depend on the presence of O− ion bombardment induced adatom
mobility. Further, Ngaruiya et al. [11] also studied the structure formation of various
transition metal oxides of group 4 (Ti, Zr, Hf), 5 (V, Nb, Ta) and 6 (Mo, W) deposited by
reactive magnetron sputtering at 6 mTorr. Ngaruiya et al. found that Zr and Hf-based
sputtering processes allow for the formation of the monoclinic phases of their respective
oxides. While the other transition metals from group 5 (V, Nb, Ta) and 6 (Mo, W) form
amorphous films. For the zirconium target, Ngaruiya et al. observed that, with the increase
in oxygen flow, the films show a decrease in compressive stresses up to a critical limit which
marks the onset of complete target oxidation. At complete target oxidation, film stresses
increased abruptly (up to (-1500 MPa) and these fully oxidized films were found to be in the
monoclinic phase. They attributed the generated compressive stresses in the film to the
energetic bombardment of oxygen negative ions as result of target oxidation. The latter is
monitored by observing the target voltage as a function of the oxygen flow introduced in the
chamber (Fig. 2. 8). For group 4 oxides (Ti, Zr, Hf), it is observed that the cathode potentials
are moderate as compared to group 5 and 6, which result in providing moderate O- ion energy
for film crystallization and atomic arrangement that causes stresses build up. Finally, it is
concluded that the high flux of low energy oxygen negative ions emitted from the Zr target
accounts for the crystallization and the stress build up in the zirconium oxide films. On the
other hand, for group 5 (V, Nb, Ta) and 6 (Mo, W) oxides, it is proposed that the oxygen ions
get accelerated by the high cathode potential and bombard the growing film. To the contrary
20
of the group 4 transition metals, in this case, the swift bombardment results in the relaxation
of stresses via thermal spikes and in the production of an amorphous structure.
Fig. 2. 8: Variation in Zr target voltage and film stresses as a function of oxygen flow [11].
In 2006, Severin et al.[28] controlled the target oxidation state and therefore the emission of
the fast O- ions by adding N2 to the sputtering process. Later in 2008, Severin et al.[10]
reported the stabilization of the high temperature cubic phase of zirconia films at room
temperature by using such N2 addition strategy (Fig. 2.9). When nitrogen is not added in the
sputtering process and zirconium is sputtered in an argon/oxygen atmosphere, the deposited
films are mainly amorphous and monoclinic. There is only a very small window around 1.7
sccm of oxygen flow where the cubic zirconia films are obtained. On addition of 0.75 sccm
nitrogen flow, the cubic phase zirconia thin films are deposited at lower partial pressure of
21
O2 and the operating window is broadened. This window is further enlarged by increasing
the nitrogen flow to 1.5 sccm and lowering the oxygen flow.
Fig. 2. 9: Phase diagram of reactively sputtered ZrOxNy films for various oxygen and nitrogen flows
as reported by Severin et al.[10].
In 2010 Sarakinos et al.[13] demonstrated for HfO2 (which is isostructural to ZrO2) that it’s
not the bombardment of O- ions which governs the cubic or tetragonal phase formation but
the incorporation of O vacancies and/or the substitution of O atoms by N atoms in the
nonmetal sub-lattice. Sarakinos et al. used the same dc-reactive magnetron sputtering and
sputtering ambient as the one reported by Severin et al. [28]. To prevent the fast O- ions
emitted from the target (i.e. the racetrack) to reach the film, they mounted a Cu ring above
the sputtering target, as presented in Fig. 2.10. It should be noted here that Sarakinos et al.
used a high power pulse magnetron sputtering discharge (HPPMS) to grow the films. This
technique is known to enhance ion bombardment during film growth.
22
Fig. 2. 10: Schematic of the strategy employed by Sarakinos et al to block the bombardment of O-
ions on the growing film. Image adopted from [13].
Influence of oxygen vacancies on zirconia c-phase stabilization
The role of oxygen vacancies in the stabilization of zirconia tetragonal and cubic phase has
been also studied theoretically by Fabris et al. [12]. for a cell of 96 atoms (32 Zr, 64 O) these
authors demonstrated that by incorporating 1 (equivalent to incorporating ~3.2 % mol Y2O3)
and 4 (equivalent to incorporating ~14.4 % mol Y2O3) oxygen vacancies in the zirconia lattice
i.e. by removing oxygen atom(s) from the cell, that the presence of oxygen vacancies is
responsible for the stabilization of the high temperature tetragonal and cubic phase of
zirconia. However, they also suggested that a such stabilization procedure i.e. by doping the
zirconia crystal with oxygen vacancies, may be achieved theoretically only.
2.3 Properties and applications of zirconia
Zirconia is an extremely versatile material finding its applications in medical field[29–32],
as decorative coatings [33–35], cutting blades [7], thermal barrier coatings [36–38], in
23
oxygen sensors [39–41], solid oxide fuel cells [41–45] and in other high temperature
applications because of its superior chemical stability, optical, mechanical, thermal and ionic
properties. Therefore, a brief overview of these mechanical, thermal, ionic, and optical
properties is given in the next sections.
2.3.1 Mechanical properties of zirconia
The elastic properties of pure bulk zirconia have been studied by Chan et al. [46] as a function
of the crystal structure of the material. It is reported that the bulk moduli of monoclinic and
tetragonal zirconia hovers around 150-200 GPa while cubic zirconia has a bulk modulus
around 171-288 GPa. On the other hand, high pressure phases i.e. orthorhombic-I and
orthorhombic-II phases, have values around 224-273 GPa and 254-444 GPa, respectively.
Further, the monoclinic phase of zirconia has hardness about 9.2 GPa[47] for samples with a
density >98% and 4.1-5.2 GPa[48] for samples with a density >95%, whereas hardness of
amorphous zirconia vary between 5 and 25 GPa[49]. The hardness values have been observed
to increase with the addition of yttria in the zirconia lattice, e.g. the hardness approached 11
GPa with a doping of 1.5 mol% of yttria [48]. The addition of larger amounts of yttria i.e. 8
mol.% of yttria leads to a further increase in the hardness which reaches a value of 15 GPa
[50].
As can be seen from the above mentioned values of elastic modulus and hardness, tetragonal
and cubic zirconia have superior mechanical properties than monoclinic. This is the reason
why most of the mechanical engineering applications make use of tetragonal and/or cubic
phases by stabilizing them at room temperature with the help of dopants.
24
2.3.2 Thermal properties of zirconia
Zirconia in its monoclinic phase exhibits a thermal conductivity of 7.2 W/m.K [51] at room
temperature. This value decreases to 2.5 W/m.K at 1100 °C (see Fig. 2.11). This makes this
material very interesting for thermal barrier coatings (TBCs) [36]. However, an increase in
temperature leads to the phase transformation (shown in Fig. 2.2) and thus causes
delamination and cracks in the coating. This is avoided by doping the zirconia lattice with
yttria [36], i.e. by stabilizing the high temperature cubic phase of zirconia. Such stabilization
not only help to avoid phase transformation but also lowers the thermal conductivity of YSZ,
e.g., from 1.42 W/m.K (at room temperature) to 1.35 W/m.K at 1200 °C as c-phase of zirconia
has lower thermal conductivity as compared to m-zirconia.
Fig. 2 11: Thermal conductivity of pure zirconia (ZrO2) [51].
2.3.3 Ionic conduction of stabilized zirconia
Stabilization of zirconia at room temperature with the help of dopants not only make it useful
for mechanical and thermal barrier applications but also as an ionic conductor (electrolyte)
to be incorporated in solid oxide fuel cells (SOFCs) and oxygen sensors. It is because, on
25
addition of aliovalent ions in the zirconia lattice, oxygen vacancies are created to maintain
the charge neutrality.
It was first expected that the ionic conductivity will increase with the increase in O vacancy
concentration i.e. by increasing dopant content (Y2O3). However, later it was observed the
maximum ionic conductivity in yttria-stabilized zirconia (YSZ) occur at 7 – 9 mol% Y2O3 at
327 – 1227 °C [52,53]. On the other hand, higher amount of Y2O3 was found to lower the
mobility of O vacancy by increasing the diffusion energy across an Y–Y common edge as
compared to the diffusion across one with a Zr–Y common edge [52]. The formation of
oxygen vacancies allow oxygen ion O2- migration[43,54], as schematically represented in
Fig. 2.12. YSZ is one of the most frequently used electrolyte materials in SOFCs and oxygen
sensors because of its great availability, cost effectiveness, stability, and high ionic
conductivity.
Fig. 2. 12: Schematic of oxygen vacancy migration in YSZ, Image adopted from [5].
Stabilized zirconia and fuel cells for energy production
A Fuel cell produces electricity from the electrochemical combination of a fuel with an
oxidant. A fuel cell consists of an anode and a cathode separated by an electrolyte. The fuel
(e.g., hydrogen) is fed to the anode where it is oxidized and electrons are released to the
external (outer) circuit. While the oxidant (e.g., oxygen) is fed to the cathode where it is
26
reduced and electrons are accepted from the external circuit. So, electrons flow from the
anode to the cathode through an external circuit and produce an electric current. The
electrolyte conducts ions between the anode and cathode. The key feature of fuel cells is their
high energy conversion efficiency because the fuel cell directly converts chemical energy
into electrical energy. Moreover, fuel cells also offer many advantages over traditional energy
conversion methods e.g., significant higher conversion efficiency, modular construction, high
efficiency at part load, minimal siting restriction, potential for cogeneration, and the most
important, much lower production of pollutants [43].
The principle of fuel cell was first reported in 1839 by Sir William Grove [55]. His fuel cell
used diluted sulfuric acid as an electrolyte and operated at room temperature. The discovery
of solid-oxide electrolytes came much later, in 1899, by Nernst [56] and the operation of the
first fuel cell containing solid-oxide electrolyte (also called solid oxide fuel cell) at 1000 °C
was reported by Baur and Preis in 1937 [57]. Since then, SOFC technology has developed a
lot. Mutlikilowatt fuel cells based on stabilized zirconia electrolyte have been operated for
thousands of hours and have shown good performance[43]. However, the conductivity
requirement for the electrolyte determines the operating temperature of the SOFC e.g., the
ionic conductivity of 8 mol% yttria stabilized zirconia equals 1x10−5 S/cm at 500 °C but
ramps up to 4x10−3 S/cm at 900 °C [43,44,54,58]. This situation requires to the use of high
operating temperatures for SOFCs, typically about 1000 °C [43]. For the device to operate at
such high temperatures, the device components must need to meet certain requirements. For
example, each component of device must have the proper stability (chemical, phase,
morphological, and dimensional) in oxidizing and/or reducing environments, chemical
compatibility with other device components, and proper conductivity. In addition to these,
the components of device must have similar coefficients of thermal expansion to avoid
27
separation or cracking during fabrication and operation. Beside the above mentioned
requirements, the electrolyte and interconnect must be dense to prevent the gas mixing. The
electrolyte used in the SOFCs are normally of thickness ranging several micrometers (thick
films). However, the use of thicker films leads to the high ohmic losses. Therefore, to
overcome such high ohmic losses high operating temperatures are used. An alternate way to
overcome this limitation is the use of thin films (tens of nm thick) e.g., superionic
conductivity i.e. 1S/cm has been reported for 58 nm thick YSZ films in the temperature range
150-500 °C [59] This will not only help to lower the operating temperature and ohmic losses
but will also help to lower the manufacturing cost of the SOFCs by allowing the use of low
cost device components.
2.3.4 Optical properties of Zirconia
Zirconia has also very good optical properties. The refractive index typically equals 2.1-2.2
[9,60,61]. Zirconia also presents a large optical band gap of 5.1-6 eV[62–66] and is therefore
transparent in the visible range. Zirconia is also a high K dielectric (K=25) [67] material,
making zirconia a good candidate in the silicon microelectronics element base applications.
Further, zirconia is a prospective candidate for the role of an active medium in the next
generation Resistive Random Access Memory (ReRAM) [68]. The optical and transport
properties of zirconia, as for other high-K dielectrics, are determined by the presence of
defects in the structure. According to the literature, a blue luminescence band with an energy
of about 2.5-2.8eV is observed in zirconia [67,69–73]. However, the blue luminescence is
attributed to the presence of defects and impurities in the zirconia lattice but the origin of
these luminescence is unclear.
28
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31
3. Thin film growth
Thin films can be deposited in a number of ways. In this regard, methods commonly used to
synthesize tetragonal and cubic zirconia thin films are presented in Fig. 3. 1. As discussed in
section 1.1, zirconia thin film deposition can be divided in two major disciplines, i) chemical
deposition methods and ii) physical vapor deposition methods (PVD). Chemical deposition
method further can be divided into chemical solution based synthesis and chemical vapor
deposition (CVD) method. In chemical solution based method, precursor solution is
employed on the substrate and later processed to have a thin film, it includes Sol-gel method
[1–4] and Spray pyrolysis [5–7] methods. On the other hand, in chemical vapor deposition
method, volatile gasses (precursors) are let into the deposition chamber, which is under
vacuum and the zirconia coating is grown on a heated substrate by a chemical reaction
occurring on or in its vicinity. CVD itself is the parent of a family of film growth processes.
The most common sub-CVD methods for zirconia thin film deposition are atomic layer
deposition [8–12] and metal organic chemical vapor deposition (MOCVD) [13–17]. Like
CVD, the term PVD also includes a family of techniques. During physical vapor deposition,
a vapor of the film forming species is created from a solid, or sometimes, from a liquid source,
by physical means. The source of material is usually named the target and the vapor of
depositing species can be obtained e.g. through heating or as a result of the interaction of fast
ions or a LASER beam with the target surface. The vapor is then transported through the gas
32
phase to the substrate, where it condensate and form the film. The common methods among
PVD to deposit zirconia thin films include cathodic arc evaporation [18], pulsed laser
deposition [19–22], ion beam assisted deposition [23–25] and sputtering [26–36].
Fig. 3. 1: Flowchart summarizing the most common deposition methods used to deposit zirconia
thin films.
3.1. Sputtering
Sputtering is a widely used method not only to deposit binary compound materials such as
Metal-Oxides or Metal-Nitrides thin films but also to deposit elemental, alloy, mixture or a
compound film depending on the target composition. Further the choice of sputtering comes
from its conceptual simplicity as well as from its easy scalability. Sputtering is by definition
the ejection (removal) of atoms from the target material which is bombarded by energetic
species. The bombarding species are ions or fast neutrals extracted from the plasma but
(positive) ions are usually considered as the most important particle. Depending on the energy
of the incident/bombarding ion, the sputtering process can be divided into three regimes; i)
single knock-on (low energy, 10-30 eV), ii) linear cascade (moderate energy, 100 eV-10 keV)
33
and iii) spike (high energy, >10 keV ), as shown in Fig. 3.2. In a single knock-on event, in
the low energy regime, a small fraction of the target atoms is set into motion by interaction
of incident ion with the surface atoms. The bombarding ion provides enough energy to the
surface atoms to overcome the surface binding energy and to sputter out from the target
surface. While in linear cascade regime, the incident particle goes under a series of collisions
with several target atoms. This situation triggers the displacement of several target atoms
from their sites, which go in collision with other surface atoms (linear cascade) and thus
sputtering them out. In the spike regime, the incident ion carries a substantially high amount
of energy which provide enough energy to all the atoms along its path to overcome their
binding energy, causing a higher spatial density of moving atoms as compared to linear
cascade regime. In the spike regime the density of recoil atoms is so high that the majority of
atoms with in the spike volume are in motion and the region of collisions become so dense
that multiple collisions occur simultaneously. In this case the process becomes a complicated
process of many-body interactions between hundreds and tens of thousands of atoms, which
cannot be treated with the binary collisional process (as in the case of linear cascade).
The number of atoms ejected out from the target per incident ion is referred as the sputtering
yield. The latter depends on the incident ion energy, the angle of incidence, the surface
binding energy of the target material, and on the mass of the incident ion and the target atom
[37]. Beside the ejection of atoms from the target surface during sputtering, some additional
phenomena are also observed. Such phenomena include backscattering of the incident
species, change in the surface structure and morphology, as well as the emission of electrons
from the target surface and the production of phonons[38]. The electrons emitted from the
target by the ion bombardment are known as secondary electrons. The emission of secondary
electrons is of particularly important in the case of plasma-based sputtering methods. Indeed,
34
the secondary electron emission results in the generation of new ions through inelastic
collisions with atoms in the surrounding gas phase, these newly formed ions will in turn
bombard the target. This situation enables the continuous emission of secondary electrons
and gas atom ionization which allow the formation of a quasi self-sustained glow discharge.
Glow discharges are discussed in the next section.
Fig. 3. 2: Sputtering regimes (a) Single knock-on (low energy) (b) Linear cascade (c) Spike (high
energy) [37].
3.1.1. Glow discharge
The sputtering process is generally based on a diode configuration with facing electrodes,
where the cathode is the target to be sputtered. To realize sputtering, high-purity noble gas
(typically Ar) is introduced in the continuously pumped chamber. The working pressure
ranges from a few tenths of mTorr5 up to 100’s of mTorr. Then a constant negative potential
is applied to the target using an electric power supply. This result in the acceleration of the
few electrons already present in the low pressure gas. These initial electrons-ions pairs are
generated by external sources such as X-Rays or UV photons coming from the surroundings
(e.g. cosmic rays). The primary electrons, accelerated by the electric field of the cathode
(typically hundreds of volts), promote a wave of ionization in the noble. The positive ions
5 1 mTorr = 0.133 Pa
35
(Ar+) produced during this ionization cascade can be in turn accelerated towards the
negatively biased sputter target. The ion bombardment results in the removal of target atoms
(sputtering) as well as in the emission of secondary electrons as described in the previous
section. These fast secondary electrons allow to sustain the ionization of the gas phase. As
the steady state is reached, a partially ionized gas consisting of energetic particles i.e. ions,
electrons, photons and a much larger number of gas neutrals is obtained. This particular
medium is called glow discharge or plasma[39]. The term plasma was first used by Irving
Langmuir in order to describe an ionized gas in 1927[40].
The generation and stabilization of the plasma is the core of the sputtering processes. For this
purpose, collisions in the gas phase are essential. Since the gas phase is a combination of
ions, electrons, neutrals and molecules, therefore ideally one should consider the interaction
between all possible pair of permutations. These collisions can be divided into elastic and
inelastic collisions. In elastic collisions, the kinetic energy of the particle is transferred to the
target particle and the target particle is accelerated i.e. the internal energy of the target particle
remains unaffected. In inelastic collisions, part of the kinetic energy of the missile is
transferred to the target particle but in this case, the internal energy of the target particle
increases. Hence, the target particle can be excited, ionized, or dissociated (if the target
particle is a molecule).
In low pressure plasmas such as those utilized for sputtering, the collisions involving
electrons usually dominate the ionization processes[39]. Ionization by electron impact is
allowed if the electron has at least a high enough kinetic energy i.e. above the ionization
potential of the target atom or molecule. Actually, every collisions process occurring in the
gas phase is characterized by a certain probability which is determined by the so-called
36
collision cross-section. Cross sections are energy dependent quantities[39]. Other ionization
processes observed in the plasma phase include ion-neutral and metastable-neutral collisions.
The counter part of the ionization is electron-ion recombination. Due to the recombination
processes which mainly happens at the chamber walls, an external source is always needed
for the generation of new electrons to sustain the plasma. This is the role of the negative
potential applied to the target as described in the beginning of this section. Further, a less
dramatic transfer of the energy would cause excitation, where the bound electrons of the
target atom/molecule jump to a higher energy level. As the lifetime of the excited state is
finite, the excited electron decays to a lower energy state after that finite time by releasing
the excess energy in the form of a photon. The emitted photon may have energies
corresponding to the visible spectral range. Consequently, the plasma glows.
Formation of sheath
Plasma in general is a quasi-neutral i.e. there is an equal amount of negative and positive
particles (ne≈ni)[39]. This changes dramatically when the plasma is locally disturbed, e.g.
when an electrically isolated (floating) surface is inserted into the plasma. In this situation,
the substrate gets exposed to a flux of electrons, ions and neutral atoms. The flux of impinging
species is, in a first approximation, equal to 𝑛𝑐
4 [39] where n is the density of the
corresponding particles and 𝑐 their average velocity. In the very beginning, since the
electrons are much faster than the other plasma species, because of their much lower mass,
the floating substrate experiences a much higher electron current density as compared to the
ion current density. This result in the buildup of negative net charge on the floating substrate
and hence of a negative potential with respect to the plasma. Consequently, newly arriving
electrons get repelled by this negative potential while positive ions are attracted. This
37
negative potential appearing on the electrically insulated surface is known as the floating
potential (Vf) and the potential of the undisturbed bulk plasma is the plasma potential (Vp).
The plasma potential has the most positive value in the glow discharge. Once electron and
ion fluxes equilibrate, a potential difference equal to |Vp|-|Vf| appears between the plasma
bulk and the substrate surface. As a consequence, only electrons with energy higher than the
potential difference |Vp|-|Vf| can overcome the barrier and a net positive charge appears
around the surface. Locally the neutrality relation (ne≈ni) doesn’t hold anymore and this
space charge region is named sheath [39].
Now let’s assume that the kinetic energy of the electrons obeys the Maxwell-Boltzmann
distribution, then the plasma density in the sheath ��𝑒 can be written as[39]
��𝑒𝑛𝑒= 𝑒𝑥𝑝 − (
𝑒(𝑉𝑝 − 𝑉𝑓)
𝑘𝐵𝑇𝑒) (3.1)
Here ne is the plasma density of undisturbed plasma and Te is the electron temperature. Since
the plasma potential difference |Vp|-|Vf| result in the spatial variation of the potential
(∆𝑉(𝑥))within the sheath and is given by the equation[39],
∆𝑉(𝑥) = (𝑉𝑝 − 𝑉𝑓) exp (−𝑥
𝜆𝐷) (3.2)
where 𝜆𝐷 = (𝑘𝐵𝑇𝑒𝜀0
𝑛𝑒𝑒2)1
2 and is called the Debye length and 𝜀0 is the permittivity of free
space[39]. Debye length is the characteristic plasma dimension on which the charge densities
can exist. From equation 3.2, it can be understood that over a distance 𝜆𝐷 , any plasma
perturbation is reduced to 0.37(1/e) of its initial value. Therefore, the plasma species located
at 2-3 Debye length from the point of perturbation will practically remain unaffected.
38
Assuming 𝑛𝑒 = 1016𝑚−3, and 𝑘𝐵𝑇𝑒 = 2𝑒𝑉 [41],typical value of the Debye length for a cold
low pressure plasma is in the order of 10-4 m [41].
Fig. 3. 3: Plasma potential distribution in a dc glow discharge[39].
Fig. 3. 3 represents the potential distribution in a glow discharge generated by applying a DC
(direct current) continuous voltage at the cathode. The area between the cathode (target) and
anode (chamber walls) is known as the plasma potential (Vp) and has the most positive
potential. In the vicinity of the cathode and anode sheaths are formed, so that the perturbation
of plasma due to the presence of electric is restricted in this region.
3.1.2. Magnetron sputtering
The dc diode base sputtering, discussed in previous section, has several drawbacks. the most
important is that the process requires relatively high sputtering pressure (typically….), which
result in increased scattering of the sputtered species in the gas phase and thus low deposition
rates as well as poor film quality. Further the use of high cathode voltages has a consequence
on the bombardment of anode by fast electrons, which are repelled from the target (cathode)
because of their negative charge. This in return cause substantial heating of the growing film.
39
It is very undesired phenomena which affect the resulting film properties. To avoid such
drawbacks and to be able to maintain the glow discharge at lower sputtering pressures and
voltages, an increase in the ionization of plasma species is essential. This can be achieved by
combining the existing electric field of the cathode with a magnetic field, i.e. by placing a a
set of permanent magnets behind the cathode. In such configuration, the force F experienced
by an electron can be written as
�� = −𝑞(�� + ��×��) (3.3)
where �� is the electric and �� magnetic field vectors, �� is the velocity of the electron and q its
charge. In such scenario, the electron repelled by the negative potential of the cathode will
move helicoidally around the magnetic field lines due to the superposition of the electric field
force �� = 𝑞�� and the Lorentz force 𝐹𝐿 = −𝑞��×��
Fig. 3. 4: the principle of magnetron sputtering. Electron are trapped by Lorentz force in an
inhomogeneous magnetic field and result in increased ionization[42].
The electrons following the helicoidal trajectories are trapped in the vicinity of the target and
result in the increased ionization probability of the gas particles in that region, as shown in
40
Fig. 3. 4. Such configuration in sputtering is known as magnetron sputtering and is the most
popular technique in physical vapor deposition techniques[42,43].
One of the disadvantage of the early developed magnetron sources was that the plasma was
too effectively confined near the target surface due to the magnetic field emerging and
reentering the target in a closed loop type of pattern, also known as balanced magnetic field
configuration[43] (see Fig. 3. 5(a)). In 1986 the issue was resolved by Window and Savvides
[44] by introducing an unbalanced magnetron, Fig. 3. 5(b). In such unbalanced configuration,
stronger outer magnetic ring is used which cannot be compensated by the inner weak
magnetic ring. Such arrangement allows some electrons to escape from the confining
electromagnetic field and create plasma away from the target surface area. This consequently
leads to the better transport of charged particles towards the substrate. Thus enhancing the
ionization in the substrate vicinity allowing for more intense bombardment on the growing
film.
Fig. 3. 5: Schematic illustration of the cross section of (a) balanced and (b) unbalanced magnetron.
41
3.1.3. Reactive magnetron sputtering
Single or multi-component thin films can be deposited by magnetron sputtering only when
the elemental or compound target is used, respectively. In this way, the target has to be made
of the same constituent elements of the material that has to be deposited. This result in the
same stoichiometry of the films as of the target or some time in slight variations, due to
preferential sputtering of the target atoms and/or by the way the sputtered material is
transported to the growing film. Further, this limits the possibility of using magnetron
sputtering where the films composition has to be different than the target. The alternate of
such is the use of reactive gas along with the noble gas (Ar) in the magnetron sputtering
process, and the principle technique is known as reactive magnetron sputtering. Using such
strategy, one can deposit metal oxide, nitride and carbide films by adding O2, N2 or
CH4,/C2H2 in the magnetron sputtering process, respectively. Moreover, it also allows the
thin film scientist to deposit films with various compositions by tuning the flow of the
reactive gas into the deposition chamber.
When a reactive gas is added in the sputtering process, it interacts with the target surface and
with the collecting areas of the chambers, i.e. walls of the chamber and substrate. The
injection of reactive gas has tremendous implications on the sputtering process, as show in
Fig. 3. 6. In the Fig. 3. 6, the reactive gas partial pressure, target voltage and the deposition
rate are plotted during the sputtering of a zirconium target at constant current in a Ar/O2
ambient, as a function of reactive gas (O2) flow[45]. In the beginning, i.e. at very low flow
of O2, the reactive gas partial pressure remains very low (Regime I, Fig. 3. 6(a)) as the O2
flow is increased. This can be attributed to the gettering (the removal of reactive gas atoms
by a getter e.g., Zr) of the reactive gas molecules by the Zr atoms that are freshly deposited
on the substrate surface and the chamber walls [46]. This low O2 partial pressure in regime I
42
results in the very low target coverage (the target surface is not oxidized) with high deposition
rate. As a consequence, the films deposited in this regime contain high metallic content and
is the reason why this regime I is also called metallic mode. The oxygen partial pressure in
the chamber increases abruptly when the oxygen flow rate exceeds the gettering rate, regime
III. This results in the full coverage of the target leading to a ZrO2 compound formation on
the target surface and the decrease in the deposition rate at the substrate. This decrease in
deposition rate is a common feature of reactive magnetron sputtering as the sputtering yield
of formed compound is always lower than the one of the corresponding metal[47]. This could
be explained by the presence of high ionic/covalent bond i.e. high binding energy as
compared to metallic target[48]. This regime III is also called poisoned or compound mode
due to the full coverage/oxidation of the target. In this situation stoichiometric oxide films
are grown. The regime II, lying between the metallic and the poisoned mode, also known as
transition zone, is frequently abrupt and is often accompanied by a hysteresis effect. The
abrupt change in the target coverage and the hysteresis is the result of a nonlinear relationship
between the reactive gas flow and its partial pressure in the chamber as a result of gettering
phenomena. Trying to operate in the regime II with a normal reactive flow control of the
reactive gas is very difficult if not impossible for most reactive systems. It could be also seen
from Fig. 3. 6(b) that such transition from metallic to compound mode has implications on
the target voltage. The increase of target voltage in the case of Zr-O2/Ar indicates the increase
of plasma impedance as the target current was kept constant in this case. Such relation
between the target current-voltage has been studied by Depla et al. [49] for various oxides
and the change in voltage has been attributed to the variation of the secondary electron
emission coefficient. They also found that there is an inverse relationship between the
43
secondary electron emission coefficient and the target voltage, i.e. the voltage drops for
materials where the secondary electron emission increases and vice versa.
As discussed, in the metallic mode films with a very high metallic content are synthesized.
On the other hand, in poisoned mode, only stoichiometric films with very low deposition rate
can be deposited. The only way to deposit films with high deposition rates and with variation
in film chemical composition (under-stoichiometric) is to work inside the transition zone
[30,33,36,47]. The difficulty here, is that this regime is usually unstable and extremely
sensitive to the reactive gas partial pressure. One can overcome this drawback and work
inside the transition zone by using a plasma emission monitoring device or by using a voltage
feedback control unit. In the present thesis work, reactive magnetron sputtering was chosen
for film deposition and the film composition was varied by working inside the transition with
the help of voltage feedback control unit.
Fig. 3. 6: Effect of oxygen flow on the (a) oxygen partial pressure, (b) target voltage and on (c)
deposition rate during the Zr-O2/Ar reactive magnetron sputtering at constant target current [45].
44
3.2. Thin film growth
Once the atoms are sputtered from the target surface, they are transported through the vapor
phase and reach the substrate surface where they condensate and form a film. The
mechanisms involving different processes of thin film formation could be roughly divided
into two stages; i) early stages i.e. the arrival of sputtered atoms onto the substrate surface
(deposition rate, R) and accommodation, their migration along the substrate surface (surface
diffusion) and finally their incorporation into stable clusters (nucleation) and thus 2-
dimensinal (2D) film grow. After initially forming one or more 2D monolayers, further layer
growth becomes energetically unfavorable and 3-dimensional (3D) islands form. The
schematic of these mechanisms is presented in Fig. 3. 7.
Fig. 3. 7: Schematic of processes leading to film growth [50]
45
3.2.1. Early stages of thin film formation
The atoms arriving the substrate from the vapor phase are known as adatoms. When an
adatom lands on the substrate, it loses most of its momentum and kinetic energy through
dissipation into vibrations of the substrate lattice (phonons). In fact, the adatom starts to
interact with the substrate surface when they are at a distance of several Å[41] i.e. adatom
experience an attractive potential of the substrate surface atoms as they approach the
substrate. Thus they gain energy which is in the order of cohesive energy of the substrate
material. Since for the adatom to stick to the surface i.e. not to bounce back to the vapor
phase, it has to reduce its total kinetic energy through dissipative mechanisms to a value lower
than the adsorption energy (Ead). The ratio of the number of deposited adatoms to the number
of impinging adatoms is known as sticking coefficient. Theoretical calculations[51] for
nearly the same masses of adatoms and substrate surface atoms shows that the sticking
coefficient is unity if the kinetic energies are up to 25 times the adsorption energy. Sticking
coefficient are significant lower than unity when high substrate temperature or if atoms
lighter that the substrate atoms are considered. The residence time (𝜏𝑠) of an adatom before
it re-evaporates in to the gas phase is given by equation (3.4)[43].
𝜏𝑠 =1
𝑣𝑜exp (
𝐸𝑎𝑑
𝑘𝐵𝑇) (3.4)
where 𝑣𝑜is the adatom-surface vibrational frequency which depends on the adatom-surface
combination and 𝑘𝐵 is the Boltzmann constant. So an adatom will combine with a surface if
large surface residence time and high sticking coefficients are reached, this criterion is met
only when 𝐸𝑎𝑑 ≫ 𝑘𝐵𝑇.
Once the adatom has landed the substrate surface, the next step is surface diffusion. Surface
diffusion is a very important mechanism for migration of adatoms to the site which allows
46
nucleation and growth. Surface diffusion can be understood as a 2 dimensional random walk
during which an adatom jumps from one potential well to another.
After surface diffusion, next is the nucleation, i.e. the adatoms form aggregates (nuclei) which
grow in size or dissociate. For a relatively low density of nuclei the growth take place via
surface diffusion as shown in Fig. 3. 8. In a first approximation the process depicted in Fig.
3. 8 can be treated thermodynamically and thus the change in free energy on formation of
cap-shaped nuclei with a mean size r can be written as
∆𝐺 = 𝑎3𝑟3∆𝐺𝑉 + 𝑎1𝑟
2𝛾𝑣𝑓 + 𝑎2𝑟2𝛾𝑓𝑠 − 𝑎2𝑟
2𝛾𝑠𝑣 (3.5)
where γ are interfacial energies and v, f and s stands for vapor, film, and substrate,
respectively. The pre-factors, a1, a2 and a3 are geometrical constants which are determined by
the nucleus-substrate wetting angle θ, as shown in Fig. 3. 8. ∆𝐺𝑉 is the change in Gibbs free
energy due to gas-solid transformation, supersaturation in the vapor phase leads to the
negative Gibbs free energy which is the driving force for nucleation.
Fig. 3. 8: Atomistic mechanisms during nucleation[41].
The dependence of ∆𝐺𝑉 on the nuclei mean size r is plotted in Fig. 3. 9. As can be seen from
Fig. 3. 9, there is a critical nuclei size 𝑟∗ on which the ∆𝐺𝑉 depend and thus the nuclei growth.
. The nuclei whose size 𝑟 > 𝑟∗ will grow otherwise, in case of 𝑟 < 𝑟∗ , the nuclei will
dissociate.
47
Fig. 3. 9: Dependence of nuclei free energy on the mean nuclei size r [41]
For the development of a continuous film and its microstructure, the growth of 2-dimensional
nuclei is of paramount importance. Based on interfacial energies, in the case of epitaxial
growth, three growth modes have been established[50], as shown in Fig. 3. 10. The term
epitaxy is used here to describe the growth of a crystalline films on a crystalline substrate.
Depending on the interfacial energies i.e. if 𝛾𝑠𝑣 < 𝛾𝑓𝑠 + 𝛾𝑣𝑓, then the energy balance requires
the minimization of the area covered by the nuclei. This will lead the nuclei to grow in the
form of three dimensional islands, also called (a) Volmer-Weber growth mode (Fig. 3. 10(a)).
In case of ideal homoepitaxy (the case when the film and substrate have perfect lattice match,
i.e. are of same material), 𝛾𝑠𝑣 = 𝛾𝑓𝑠 + 𝛾𝑣𝑓 , will lead to the uninterrupted layer by layer
growth, such growth mode is known as (b) Frank-van-der-Merwe growth mode (Fig. 3.
10(b)). If 𝛾𝑠𝑣 > 𝛾𝑓𝑠 + 𝛾𝑣𝑓, then the area covered by the nuclei should be maximized and will
lead to the formation of one layer at a time, but due to the film-substrate lattice constant
mismatch, strain will develop in the layers. After the growth of few layers, the interfacial
strain energy 𝛾𝑓𝑠 will break the condition for the layer by layer growth mode and will lead to
48
the island growth. Such growth mode is known as (c) Stranski-Krastanov growth mode and
is depicted in Fig. 3. 10(c).
Fig. 3. 10: Schematic of basic growth modes [50].
One has to keep in mind that in sputter deposition, the films are normally grown non-
epitaxially and far from equilibrium. Therefore, the above-mentioned growth mode
classifications are hardly observed. 3-dimensional growth features are more frequently
observed which are predominantly kinetically controlled and result in the production of
polycrystalline films.
3.2.2. 3D thin film growth
The growth stages of a polycrystalline film formation are shown in Fig. 3. 11 [52]. This
evolution of the structure starts by nucleation (Fig. 3.11a) leading to grain growth (Fig.
3.11b), which further leads to the coalescence of the grains either by complete liquid-like
coalescence resulting in the formation of single crystals (Fig. 3. 11c) or by an incomplete
coalescence resulting in polycrystalline islands and channels (Fig. 3. 11d), and finally to a
continuous film (Fig. 3.11e).
49
Fig. 3. 11: Stages of polycrystalline films growth; (a) nucleation, (b) crystal growth, (c) island
coalescence, (d) growth by filling the channels and (e) formation of a continuous film [52].
In such continuous film growth, the same mechanisms are followed; the so-called
fundamental structure forming phenomena. Theses phenomena involve; i) nucleation, ii)
crystal growth and iii) the grain growth[52].
Nucleation involves the same mechanisms as discussed in previous section 3.2.1 i.e. the
condensation of adatoms on the substrate.
Crystal growth involves the incorporation of the deposited material into the condensed phase.
In polycrystalline films two kinds of crystal growth processes can be observed; i) the growth
of discrete crystals which form on the substrate surface, and ii) the growth of crystals on
already formed crystals.
The grain growth in polycrystalline films can also be divided into two categories depending
on the mobility of the grain boundary; i) grain growth by coalescence of the islands, ii) grain
growth by repeated nucleation, also called abnormal grain growth.
50
3.2.3. Microstructure of thin film
Depending on the fundamental structure forming phenomena and the film deposition
conditions, the microstructure of thin film can be described using structure zone model[53].
For physical vapor deposited pure elemental films, structure zone model has been developed
by considering the substrate temperature (Ts) and melting temperature (Tm) of the material
being deposited. Based on the homologous temperature (Ts/Tm), microstructure and growth
evolution of thin film is classified into 3 categories; a) Zone I (0 < Ts/Tm < 0.2), b) Zone T
(0.2 < Ts/Tm < 0.4), c) Zone II (Ts/Tm > 0.4)[53], shown in Fig. 3. 12.
Fig. 3. 12: Structure zone model of pure elemental films as a function of homologous temperature
and film thickness [53].
In Zone I (0 < Ts/Tm < 0.2), the film is composed of very thin fibers of diameters in the range
of 1-10 nm and determined by the nucleation density and statically fluctuations. With
deposition time, the crystalline fibers grow out of the primary nuclei and proceed to the top
of the film while keeping the orientation of the nuclei. However, along the substrate surface,
51
as the Ts/Tm increase, the diameter of the fiber increases. Since in this zone, Ts/Tm <0.2, bulk
and surface diffusion have no remarkable value, the film presents an under-dense structure.
In Zone T (0.2<Ts/Tm <0.4), because of the relatively high temperature, the self-diffusion of
the adatom is remarkable but still the grain boundary mobility is very limited. This result in
inhomogeneous structure along the film thickness, i.e. very fine crystalline at the substrate
but composed of V-shaped grains in the next thickness range.
Zone II (Ts/Tm >0.4), because of the high homologous temperature in this zone, the bulk
diffusion is significant in this regime. This leads to grain boundary mobility not only during
the coalescence but also during the film thickening. The grain boundaries being mobile, the
grain boundary energy minimization also takes place and result in the formation of grain
boundaries perpendicular to the film plane. The films grown in this zone exhibit
homogeneous structure in the growth direction as well as the columnar structure.
Thornton added another parameter in the structure zone model, namely the pressure of the
sputtering gas[54]. This parameter was added to investigate the role of energetic
bombardment on the film microstructure, a basic feature of sputtering process. The model
presented by Thornton predicts that, for films grown at low pressures, the shift of Zone I to
Zone T, or of Zone T to Zone II, to lower homologous temperature (Ts/Tm) due to the
increased energy transfer to the film by the energetic bombardment of plasma species.
As the structure zone model presented above provides the microstructure and growth
evolution of pure elemental thin films grown by magnetron sputtering method. Therefore, the
extension of structure zone model towards the multicomponent and/or multiphase thin films
is needed, especially for the films grown by reactive magnetron sputtering. For the
development of such structure zone model, the determination of the dependence of the
52
fundamental structure forming phenomena on the growth conditions is required. In this
regards, beside the substrate homologous temperature (Ts/Tm), an additional parameter
describing the concentration of contaminant species is considered[53]. Here it should be
noted that the term “contaminant species” does not only describe the impurities which
unintentionally contaminate the films but also the atoms of reactive gas or a second element,
which are intentionally incorporated in the host material lattice for the formation of
compound films. The contaminant species impinging on the growing film can adsorb and
segregate on the faces of the growing crystal or can be dissolved in the crystal lattice[55].
Experimental results show that the contaminant species can either promote or hinder the
operation of fundamental structure forming phenomena[56]. One particular example of such
case is the growth of O containing Al films. Fig. 3. 13 shows the influence of O concentration
on the growing film microstructure. In general oxygen has a low solubility in Al and
segregates on the surface and at the grain boundaries. This lead to the 2-dimensional layer
formation at these sites and hinders the surface diffusion and grain boundary mobility. At
low O/Al flux ratio (Jo/JAl ~ 10-3), O accumulates at the grain boundary and hinders the grain
boundary mobility. This result in zone II films with a poor texture compared to pure Al thin
films (Fig. 3. 13 a & b). At higher O concentration (Jo/JAl ~ 10-2), coarsening during the
coalescence is partially suppressed which lead to the formation of randomly oriented grains
i.e. zone T films (Fig. 3. 13 c). On further increase of O concentration (Jo/JAl ~ 0.1-1), an
oxide layer form on all islands of all possible orientation, leading to the periodic interruption
of grain growth and the start of secondary nucleation. This result in nano-size grains (Fig. 3.
13d) and is also referred as to zone III. This zone III is also predicted by structure zone models
when high deposition temperatures close to the melting point are used[53]. Finally, when the
53
O atoms are in majority as compared to Al atoms i.e. Jo/JAl > 1-5, then the film consists
mainly of Aluminum oxide which is an amorphous phase at room temperature (Fig. 3. 13e).
Fig. 3. 13: Influence of oxygen concentration on Al thin film microstructure[56].
3.3. Influence of energetic species on thin film properties during
sputtering
An interesting feature of plasma-based sputtering is the bombardment of the growing film by
energetic plasma species. As discussed previously, plasma species can be ions, electron,
photons and as well as neutrals [41]. The interaction of these plasma species especially of
ions with the growing film is determined by their kinetic energy and flux. [38]. Such
54
interaction of bombarding ions with the growing film depending on their energy is shown in
Fig. 3. 14. Of course the exact energy values depend on the growth conditions. It could be
seen in Fig. 3. 14, energy range span from few eVs to 1000 eV. The lower limit of the scale
i.e. below 0.1 eV describes the thermalized species which do not affect the growing film. The
activation of surface process starts from 0.1 eV, where physisorption of impinging species is
possible. An order of magnitude higher, i.e. at 1 eV, chemisorption processes start where the
chemical reactions are involved and last up to 10 eV. The surface characteristics of growing
film surface starts to get affected when these energies reach tenths of electron volts. Finally,
when the energies reach 100th of eV, the impinging atoms start to influence the bulk of the
growing film. The upper limit of the energy is defined where the kinetic energy and
momentum of the impinging species is so high that the deposited material is again vaporized
(re-sputtering). This effect starts about 1000 eV, depending on the material.
Fig. 3. 14: Schematic of effect of energetic bombardment at different energy levels, after [38].
In the energy range of chemisorption, the depositing ions are largely ineffective to influence
the film physical properties. On the other hand, the energy range of surface and bulk effects
are the most important as they provide additional means to tailor the film properties. For
55
example, in the energy range of surface effects (tenth of eVs), depositing ions are only able
to influence the atoms which are in their vicinity of impinging position. This will not cause
any damage but will support the annealing of defects frozen in as a result of limited adatom
mobility. On the other hand, in the energy range of bulk effects i.e. 100th of eVs, the
impinging ions get implanted in the subsurface of the growing film and cause subsurface
effects. These implantation leads to the lattice defects e.g., displacement of atoms from their
lattice site and creation of vacancy. These phenomena lead to the generation of residual
compressive stresses, which are discussed in next section. Sometime these implantations can
also cause the collision cascade, which result in the re-sputtering of the deposited material.
Moreover, these high energetic impinging ions can also penetrate deep in the growing film,
known as channeling mechanism. In this case, the ballistic damage is significantly reduced
and the energy of the impinging ions is minimized by electronic excitations. Not always all
the high energetic bombarding ions are implanted in the growing film, sometime they are
also backscattered. Finally, the interaction of bombarding ions with growing film surface can
also cause electron and photo emission.
The influence of bombarding ions has also shown implications on the crystal structure and
microstructure of the zirconia films prepared by reactive magnetron sputtering [28,31,32,57],
as already discussed in section 2.2.2. It has been shown that at low sputtering pressure or at
less target-substrate distances or at high target powers (i.e. at high energy) the deposited
zirconia films exhibit dense microstructure, which turn porous on the increase of sputtering
pressure. Increasing the target-substrate distances or at lowering the target power (i.e.
decreasing the ion density) have an identical effect on the zirconia film microstructure. One
also has to note this the data presented here deal with the energy brought to the growing film
by the ions only. But in reality the total energy brought to the growing film is a combination
56
of heat radiations (photons) towards the substrate, power transferred by electrons and ions,
neutral depositing atoms and as well as of background gas [58–61]. Taking into account these
contribution, the total energy flux per deposited metal atom can reach several keV [61].
Further, it has also been observed that, with the increase in impinging atoms energies, films
exhibit an increase in compressive stresses. Not only this, with the change in imping ion
energy a change in film crystal structure is also observed i.e. pure monoclinic or mixed
tetragonal and monoclinic or pure cubic can be obtained.
3.4. Formation and evolution of stresses in thin films
Films grown by physical vapor deposition method experience a continuous bombardment of
energetic plasma species, as discussed in the previous section. This leads to the enhanced
grain boundary mobility as well as in the formation of defects in the growing films.
Fig. 3. 15: Schematic of (a) tensile and (b) compressive stresses in the deposited film [62].
As a result, mechanical stresses are generated in the thin films and is called intrinsic stresses.
Such generated intrinsic stresses affect the films crystallinity by causing lattice distortion and
57
sometime these stresses reach so high value that they cause the phase transformation [28].
Moreover, this also affect the mechanical, thermal and electrical properties of the deposited
films. In general film intrinsic stresses can be divided into two parts; a) tensile stress and b)
compressive stress, shown in Fig. 3. 15. In polycrystalline thin films, in the post coalescence
growth stages, grains interact in order to minimize their energy and thus close any existing
gap. This leads the grain boundaries to shrinkage, which in turn causes tensile stresses [63].
Magnitude of such tensile stresses can be calculated using the equation (3.6)[41].
𝜎𝐺𝐵 =𝐸𝑓𝐷
𝐿𝑔(1 − 𝑣𝑓)(3.6)
Here 𝐸𝑓 is the elastic modulus, 𝑣𝑓 is the Poisson’s ratio of the film, 𝐿𝑔 the grain size and D
is the grain boundary relaxation distance. Similar to grain boundary shrinkage, shrinkage of
film volume due to the recrystallization, phase transformation and defect annihilation cause
tensile stresses[64]. Contrary to the tensile stresses, compressive stresses are generated due
to the expansion of the crystal lattice. In PVD, the lattice expansion is caused by the
implantation of energetic plasma species, as discussed earlier. However, compressive stress
relaxation beside the compressive stress generation at high energies is also observed
frequently[65].
Fig. 3. 16: Stress dependence on the energy of bombarding species [66].
58
The rationale behind it is, when the energy of the bombarding specie is too high, this cause
the intense local heating of the film (known as thermal spike). This provides enough energy
to the implanted and displaced atoms to move from their metastable position to film surface
(out-diffusion). Since these tensile and compressive stresses in the film are dependent on the
adatom energy, three energy regimes related to the residual stresses can be considered as
shown in Fig. 3. 16[66]; (i) low energy regime (0.1-1 eV), in this regime tensile stresses due
to the grain boundary shrinkage are dominant. (ii) moderate energy regime (1-20 eV), in this
energy regime compressive stresses are dominant due to the plasma species implantation and
atom displacement. (iii) high energy regime (>25 eV), in this energy regime stresses
relaxation occur due to the thermal spikes.
The intrinsic stresses in the film deposited on a substrate lead to the change in the curvature
of the substrate (δ), as shown in Fig. 3. 15. Using Stoney’s equation and change in substrate
curvature and film, substrate mechanical properties, one can calculate the values of those
intrinsic stresses generated in the film [67].
Beside intrinsic stresses caused by bombardment of energetic species, stresses at the film-
substrate interface can also be generated, when the film is deposited at a different temperature
than the surrounding temperature. For example, deposition of film at high temperatures. Such
stresses are called thermal stresses and originate because of the difference in film and
substrate coefficient of thermal expansion. Thermal stresses in a film can be calculated using
equation (3.7) [68]
𝜎𝑡ℎ =𝐸𝑓
1 − 𝑣𝑓(𝑎𝑓 − 𝑎𝑠)(𝑇𝑠 − 𝑇𝑎) (3.7)
59
here 𝐸𝑓 is the elastic modulus, Ts is the substrate temperature during deposition, Ta is the
temperature during measurement, 𝑣𝑓 is the Poisson ratio, 𝑎𝑠 and 𝑎𝑓 are thermal expansion
coefficient of substrate and film respectively. Thermal stresses can be compressive or tensile
depending on the sign of 𝜎𝑡ℎ. If 𝜎𝑡ℎis positive, then the stresses are tensile and if 𝜎𝑡ℎnegative,
then the stresses are compressive.
Beside stress generated by the film depositing species or by the difference in film-substrate
thermal expansion coefficient, evolution of film stresses has also been observed as a function
of film thickness [67,69]. In this regard Doerner et al. has discussed two hypothetical cases,
i) the stress increase continuously as the film grows i.e. a linear relation between the stress
and the film thickness and ii) the stress grow as function of film thickness up-to a certain
value and then stabilize. Nouveau et al. [69] have observed a similar case to (ii) for CrN films
deposited by magnetron sputtering as a function of film thickness. and observed three growth
regimes: Regime A) (up to 150nm) where the stress keeps increasing, B) (150-500nm) the film
changes its structure to release stress, C) (above 500nm) the stress value is stable and relaxation is
induced by the growth of less dense planes.
60
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4. Zirconia thin film deposition and
characterization
Today several theoretical and experimental methods exist study the growth of thin films (as
discussed in chapter 3) and their properties. In this thesis work, Density Functional Theory
(DFT) calculations are used to better understand the influence of oxygen vacancy
incorporation on zirconia phase formation. This theoretical investigation is then compared to
the properties of oxygen vacancy doped zirconia thin films synthesized by cold plasma-based
dc reactive magnetron sputtering (dc-RMS). The choice of reactive magnetron sputtering is
not only based on its simplicity of use but also because it allows controlling the film elemental
composition e.g. with the help of a voltage feedback control unit. To characterize the
deposited zirconia thin films, several analysis techniques are used. Obviously, the choice of
the analysis methods depends on the properties one wishes to investigate, the resources
available, and on the precision required.
In this chapter, the technical details related to the DFT calculations and thin film deposition
are summarized. Moreover, a description of the thin film characterization methods used to
address the film phase constitution and microstructure as well as the ionic conductivity and
photoluminescence properties is also provided.
64
4.1 Modelling and computational details
Theoretical calculations are performed at the DFT level using the 3.2 version of the SIESTA
code [1] including periodic boundary conditions. The calculations are made in the
Generalized Gradient Approximation (GGA) using the Perdew-Burke-Ernzerhof (PBE)
functional [2]. Trouillier-Martins pseudopotentials were used to simulate the nucleus and
core electrons. We used 5s2 4d2 as valence configuration for Zr with 3.04, 3.19 and 2.68 Bohr
for s, p, and d channels, respectively. For oxygen, we have used 1.14 Bohr for all channels.
The atomic basis set is described using a double zeta basis plus polarization orbitals (DZP).
Supercells of cubic, tetragonal, and monoclinic polymorphs of ZrO2 are built using 96 atoms
i.e. 32 Zr and 64 O by a duplication of the experimental unit cell in the three directions. The
final lattice vectors of the three polymorphs are: a = 10.14 Å for cubic; a = b = 10.10 Å, c =
10.36 Å for tetragonal; a = 10.29 Å, b = 10.42 Å, c = 10.62 Å, β = 99.23° for monoclinic.
Oxygen vacancies are inserted in the zirconia lattice by removing O atoms from the pristine
ZrO2 cells. Three configurations were tested. The vacancies are distributed randomly, apart
from each other, and as clusters, to assess the influence of O vacancies on the phase
constitution. In case of random vacancy generation, O atoms were first assigned numbers and
then from those total assigned numbers some number were randomly picked by the software
and O atoms corresponding to those numbers were removed. In case of apart O vacancies
were created manually by removing O atoms apart from each other. In case of cluster 2-4
vacancy were clustered together.
The optimization of the structure and the calculation of the electronic structure are performed
at 0 K using a mesh cut-off of 190 Ry and a Monhkorst-Pack grid of (2x2x2) k-points. The
atomic positions were relaxed until the forces on the atom were less than 0.04 eV/Å, while
65
the lattice vectors of the unit cell were fixed during the calculation. Moreover, the density of
states (DOS) of cubic structure were also extracted from DFT calculations data, for various
O vacancy concentrations.
4.2 Thin film deposition and process monitoring
To assess experimentally how the incorporation of O vacancies influences the zirconia phase
formation/stabilization, experiments were designed using a conventional dc reactive
magnetron sputtering (dc-RMS) setup where a Zr target is sputtered in a reactive atmosphere
containing argon and oxygen. The sputtering setup is built in such a way that it not only
helped to incorporate O vacancies in the zirconia lattice as the film grows, but also to vary
their concentration.
For this purpose, a voltage feedback control unit6 [3–6] (Speedflo mini from Gencoa, UK)
was used. The voltage feedback control unit is an auxiliary device which provides a rapid
control over the oxygen partial pressure using the target voltage as a feedback signal. This
way, the reactive sputtering process was allowed to work inside the so-called metal-to-
compound transition zone [7] of the sputtered zirconium target. The transition zone is an
experimental working window in between the metallic and oxidized modes of the sputtered
target where under-stoichiometric films are grown, as shown in Fig. 4. 1. Using this well-
optimized synthesis setup, 100 ± 10 nm thick films were deposited on Si (100) single crystal
substrates at 50, 65, 75, 80, and at 100% signal set point values. The 0 % represents the
metallic mode set point while 100 % represents the peak of the transition zone (See Fig 4.1).
Film were also grown in the oxidized mode. In order to reach a precise control over the sputter
target and therefore of the film-chemistry, O2 gas was injected at the target during the film
6 Discussed in details in section 4.2.1
66
growth and the length of the tubing system, between the O2 mass flow and the vacuum vessel,
was shortened as much as possible. It should be noted that the films used to perform the
elemental analysis, 18O2 (purity 97.1 %) was used instead of conventional 16O2. The Ar gas
was injected away from the target by using a conventional mass flow controller. A schematic
representation of the zirconia thin film deposition process is shown in Fig. 4. 2 and a summary
of deposition conditions is given in Table 4. 1.
Fig. 4. 1: Target voltage curve of Zr target as a function of O2 flow, shows the transition zone and
working points inside the transition zone as well as in the poisoned zone where the ZrO2-x and ZrO2
films were deposited, respectively.
Table 4. 1: Experimental details of the film deposition by reactive magnetron sputtering.
Base pressure <2×10-6 mTorr
Working pressure 5, 10, 20 mTorr
Target to substrate distance 6.5 cm
Discharge current (I) 0.2, 0.3, 0.4 A
Reactive gas 18O2 or O2
Sputtering gas Ar
Substrate Si (100); 525 ±20 μm
thick
Film thickness 100±10 nm
67
Every deposition run was carried out in a high vacuum chamber (length 78 cm, diameter 42
cm). The base pressure < 2×10-6 mTorr (4×10-4 Pa) was reached with help of Turbo pump
backed by primary pump (pumping speed 12 m3/h). During the depositions, the working
pressure was kept constant at e.g. 10 mTorr (1.33 Pa) in each case using a throttle valve. To
introduce sample in the deposition chamber, a load lock system was used and before
depositing each film, the target was first sputtered clean. During the depositions, non-
intentionally heated substrates were placed at a distance of 6.5 cm from a 5 cm in diameter
purity (99.97%) Zr target. The DC current applied to the sputter target (i.e., the cathode of
the system) was ranging from 200 mA up to 400 mA. The power supply was an Advanced
Energy MDX 500 dc power supply equipped with an arc suppressor (Sparkle from Advanced
Energy).
Fig. 4. 2: Schematic of Zirconia thin film deposition used in this work. The target voltage is read by
the feedback control unit (right hand side) and allows controlling the oxygen flow during film
growth.
68
4.2.1 Voltage feedback control unit
The voltage feedback control unit (Speedflo mini from Gencoa, UK) is a device which
actively monitors reactive magnetron sputtering processes and controls them by continuously
monitoring the target voltage or other variables. The principle of the voltage feedback control
unit is similar to a basic automatic control system which is described by a block diagrams in
Fig. 4. 3. When the user of the automatic control system inputs a reference value (r) which
has to be achieved (e.g. the temperature of a room), then the controller having an algorithm
will send some manipulated variable (m1) to the control elements (for example a heater). The
final control elements will then give a new manipulated variable (m2) (for example heat flow).
That variable will be (eventually) combined with other loads (l) from the system (e.g. sun or
human heat) which are not manipulated, ultimately giving an overall manipulated variable of
the reaction m3. By having a way of monitoring what the reaction was (for example measuring
the temperature in the room) one can get what is called the controlled variable value c. This
value c is fed back to the controller which compares it with the reference value (r). Based on
the error ‘e’, controller will act consequently in the loop control.
Fig. 4. 3: The basic automatic control system [3].
In voltage feedback control units, the target voltage is monitored continuously and is used as
input. Input is processed by the embedded software and an actuator output signal is produced,
69
based on certain algorithm parameters. The actuator output signal is sent to the mass flow
controller which responds according to the signal and regulates the reactive gas flow to
maintain the set point value condition. For a reliable voltage feedback control, the system
requires effective sensor information. For this purpose, voltage feedback system generates a
series of commands to the actuator via a processing algorithm which should be appropriate
in relevance, magnitude and speed of change.
The speed of data transfer between the voltage feedback control used here and the PC
(Windows) is called Update Speed. This parameter can be varied between a minimum of 0.5
to 5 sec depending on the working situation. In this thesis work, the highest update speed 0.5
seconds is used. There is also the Sampling/Actuation rate of the voltage feedback unit that
should not be confused with the Update speed above. The sampling/actuation is related to the
Input/Output communications with sensors and actuators that are in the millisecond range.
The voltage feedback controller used in this thesis work can handle large numbers of data
sampling with typical rates between 1-10 ms depending on the number of mass flow
controllers connected and the type of action on each process. In the thesis work, only one
mass flow was connected to the process and the sampling response type during the process
was selected as AUTO, allowing the voltage feedback unit to operate as fast as the processing
time allows (around 1 ms or less).
Beside using voltage feedback control unit to work inside the transition zone, there are other
options available as well e.g., plasma emission monitoring (PEM) and lambda probe. In case
of PEM, the intensity of an emission line of the element of interest present in the plasma is
monitored and used to work inside the transition zone. To use PEM, one needs to collect the
70
plasma emission efficiently and in a robust way as any variation in the line of sight or
deposition of sputtered material on the collecting window can lead to misleading conditions.
The lambda probe or lambda sensor gives a measure of the partial pressure of oxygen with
respect to the atmosphere. Lambda probe is directly installed in the process chamber. It is
based on the Nerst differential voltage potential created when the process side of the probe is
exposed to a different partial pressure of oxygen than the atmospheric side of the probe. The
probe needs O- diffusion in order to generate the potential and achieves this via a platinum
catalyst and a zirconia membrane. Since the zirconia and platinum require high temperatures
in order to be operational, therefore the sensor is also fed with an independent electrical
power. The choice of Voltage feedback unit in this thesis work is based on its simplicity to
use and the geometrical restrictions of the chamber.
4.3 Film characterization tools
To characterize the deposited zirconia thin films, various characterization techniques are used
in this thesis work and are briefly presented in this section.
4.3.1 X-ray diffraction (XRD)
X-ray diffraction (XRD) [8] is a powerful technique vastly used in material science division
to identify the crystalline phases present in materials and to measure their structural properties
(grain size, phase composition, stress). The sample to be analyzed is irradiated by x-rays
(having wavelength in the range of interatomic spacing in the crystals). The X-rays interact
with the matter, get diffracted, and are collected by a detector. Diffraction of x-rays by
periodic arrays of atoms in the crystals result in constructive and destructive interferences
which give rise to the diffraction pattern. This phenomenon is described by Bragg’s Law [9]
2𝑑𝑆𝑖𝑛𝜃 = 𝑛𝜆 (4.1)
71
where, d is the lattice spacing, 𝜃 is the scattering angle, 𝜆 is the wavelength of x-rays and n
is an integer number. During XRD measurements, the sample is scanned for a wide range of
angles and the diffracted intensities are recorded as a function of the latter. Such recorded x-
ray pattern is called diffractogram. Recorded diffractogram of the crystalline material is
viewed as a fingerprint of the crystal structure and is compared to the diffractograms in the
database to identify the material crystal structure. Beside identifying the crystal structure, one
can also extract the information regarding the grain size, crystal orientation, phase
composition and stresses present in the material by using the peak position, peak intensity,
peak width, and the change in peak position
4.3.1.1 Bragg-Brentano (θ-2θ) mode
The diffractograms are typically recorded by using the measurement setup in the so-called
Bragg-Brentano (θ-2θ) [8] geometry, where θ is the incidence angle of the x-rays measured
with respect to the surface of the sample and 2θ is the diffracted angle. In this configuration,
the source of the x-rays beam and the detector are scanned synchronously such that the
incidence angle and the diffracted angle remain the same throughout the scanning. Therefore,
in such a geometry only the crystal planes which are parallel to the sample surface are probed.
A schematic of such Bragg-Brentano (θ-2θ) geometry is shown in Fig. 4. 4(a).
Fig. 4. 4: Schematic of (a) Bragg-Brentano (θ-2θ) and (b) GIXRD geometry.
72
4.3.1.2 Grazing incidence XRD (GIXRD) mode
Normally in Bragg-Brentano (θ-2θ) mode, during the x-ray diffraction of thin films
(thickness in the range of tens of nm), the intensity of the diffracted peaks is very low as well
as the diffractograms are dominated by diffractions from the substrate due to the larger
penetration depths of x-rays (of the order of µm). This is particularly the case when highly
ordered substrates such as Si single crystals are used. In order to avoid such kind of issues,
the method of grazing incidence x-ray diffraction (GIXRD) [8] is used. During GIXRD, the
x-rays enter the sample at very small fixed incidence angles (few degrees or less) of incidence
thereby increasing the path travelled by the x-rays significantly (see the schematic shown in
Fig. 4. 4(b)). Since the incidence angle (𝜔 in Fig 4.4.) is fixed in GIXRD mode, therefore,
the diffractograms are obtained by varying the detector position. This allows probing crystals
whatever their orientation in the film.
In the present thesis work, crystallinity of the deposited films was analyzed by BB method
and GIXRD using PANalytical Empyrean with a Cu-Kα radiation (λ = 1.54 Å) source. The
diffractograms were recorded with a step size of 0.07° using an incidence angle of 0.5° at 40
mA, 45 kV of generator settings. The resulting diffractograms were compared to ICDD PDF
cards representing patterns of three polymorphs of ZrO2: monoclinic (PDF# 37-1484),
tetragonal (PDF# 81-1544) and cubic (PDF# 49-1642) in order to know the crystal structure
of the deposited films.
4.3.2 Transmission electron microscopy (TEM)
Transmission electron microscopy (TEM) [8] is a powerful technique for examining the
crystal structure and the microstructure of solid materials e.g., metals, ceramics,
semiconductors, polymers, and composites. In TEM, a focused electron beam with energy
(100-300 kV) is incident on a thin (typically less than 200 nm) sample. The signal in TEM is
73
obtained from both un-deflected and deflected electrons that penetrate the sample. A series
of magnetic lenses at and below the sample position deliver the signal to a detector, usually
a fluorescent screen, a film plate, or a video camera. Accompanying this signal transmission
is a magnification of the spatial information in the signal by as little as 50 times to as much
as a factor of 106. One considering factor regarding the TEM lenses are the diaphragms or
apertures of these lenses employed at certain positions in TEM. The purpose of these
apertures is to filter either the source or the transmitted signal. The most important diaphragm
in TEM is called the objective aperture, lying at the back focal plane of the objective lens. In
this plane the scattered electron waves recombine to form a diffraction pattern. The use of a
small objective aperture lenses while operating in the image mode, blocks all diffracted
beams, and serve to enhance the image contrast significantly. The use of a large objective
aperture allows the passage of many diffracted beams thus enhanced diffraction pattern, is
the modus operandi for the technique and is referred to high-resolution transmission electron
microscopy (HRTEM).
In the present thesis work (HR)TEM is used to study the crystal structure of zirconia films at
various positions from the substrate-film interface.
4.3.3 Secondary electron microscopy (SEM)
Scanning electron microscopy (SEM) [8] is one of the most widely employed techniques to
the study the surface topography in three dimensions of a material. In SEM, a beam of
electrons, like in TEM, with energy ranging from few keV to 50 keV is emitted from a W-
filament or LaB6-crystal.The beam is focused on the surface of the sample with the help of
magnetic lenses. As the electrons impinge on the surface, interactions occur. As a result of
this, the emanating secondary and backscattered electrons are collected by suitable detectors
74
and are used to extract the microstructure and surface topography information in the form of
images of the sample.
In the present thesis work SEM is used to study the microstructure of the deposited samples
by having their cross-sectional images.
4.3.4 Chemical composition of zirconia thin films
The chemical composition of zirconia thin films i.e. concentration of Zr and O in this thesis
work was realized by combining Rutherford backscattering spectrometry (RBS) [8] and
Nuclear reaction analysis (NRA) [8].
4.3.4.1 Rutherford backscattering spectroscopy (RBS)
Rutherford Backscattering Spectroscopy (RBS) is one of the commonly used nondestructive
technique for quantitative depth-profiling of thin films. In RBS, sample is bombarded with a
beam of high-energy (Eo) particles (MeV range) with mass M1, which undergoes an elastic
collision with the sample stationary atoms to be investigated (having mass M2). On collision,
the energy of backscattered M1 particle is detected at a given angle θ. By using laws of
conservation of energy and momentum, the mass of the target particles M2 is calculated.
Moreover, since the probability of scattering in a certain angle is known by the so-called
Rutherford cross section, this makes it possible to estimate the abundance of M2 by counting
the yield of scattered particles M1 in a certain solid angle, covered by the detector.
Results obtained by RBS are insensitive to sample matrix and typically do not require the use
of standards, which makes RBS the analysis of choice for depth profiling of major
constituents in thin films. Detection limits range from a few parts per million (ppm) for heavy
elements to a few percent for light elements.
75
In the thesis work, for elemental analysis, zirconia thin films were deposited on graphite foil.
The concentration of Zr in the samples was probed by RBS. The incident energy of the alpha
particles was 2 MeV, and the beam impinged the sample surface at normal incidence.
Backscattered particles were collected at 165° in a passivated implanted planar silicon (PIPS)
detector. Spectra were analyzed with the SIMNRA software assuming Rutherford
backscattering cross-section.
4.3.4.2 Nuclear reaction analysis
While RBS is more suitable for heavier elements, Nuclear reaction analysis (NRA) is used to
determine the concentration and depth distribution of lighter elements in solids but like RBS,
NRA is also a nondestructive technique. NRA is also isotope specific, making it the ideal tool
for isotopic tracer experiments. This characteristic also makes NRA less vulnerable than
some other ion scattering methods to interference effects that can overwhelm signals from
low abundance elements. In NRA, an ion beam with an energy ranging from a few hundred
keV to several MeV is produced in an accelerator and bombards the sample. Nuclear
reactions with low-Z nuclei in the sample are induced by this ion beam. Products of these
reactions (typically p, d, t, He, α particles, and γ rays) are detected, hence producing a
spectrum of particle yield versus energy. That accumulated spectrum is then compared to a
body of data accumulated through research in low-energy nuclear physics to determine
concentrations and distributions of specific elements or isotopes in the material under
investigation. Depending on the reaction type, NRA can be divided in to two different
categories: The Resonant Nuclear Reaction Analysis (RNRA) and the non-resonant Nuclear
Reaction Analysis. RNRA uses beam energies near narrow isolated resonances of relevant
nuclear reactions to determine the depth distribution of elements in a sample. On the other
hand, when reaction cross sections are sufficiently large over an extended energy range, the
76
entire depth profile may be obtained using a single incident beam energy. This is known as
non-resonant NRA.
In this thesis work for elemental analysis, zirconia thin films were deposited on graphite foil.
Specific 18O depth profiles were determined using the resonant nuclear reaction 18O(p,α)15N
at 151 keV. In this case, in order to allow NRA of the zirconia films, isotopic 18O was
introduced in the vacuum chamber instead of conventional high-purity oxygen. The latter is
being 16O. This way, only the oxygen introduced in the film during deposition is probed and
oxygen pollution from venting is discarded. Samples were tilted at 30° with respect to the
incident beam and the alpha particles were collected in large area passivated implanted planar
silicon (PIPS) detectors facing the sample surface and parallel to it. The incident energy was
varied from 145 to 200 keV. Depth profiles were deconvoluted in order to take into account
the energy straggling of the beam and then quantified with the help of a Si18O2 standard
produced by thermal oxidation in a pure 18O atmosphere. These RBS and NRA measurements
were performed in the Laboratoire d’Analyse par Réaction Nucléaire (LARN, Pr. S. Lucas)
at the University of Namur.
4.3.5 Electrochemical Impedance spectroscopy (EIS)
Electrochemical Impedance spectroscopy (EIS) [10] is a well-known method for
characterizing many of the electrical properties of materials with electronically conducting
electrodes. For example, it can be used to investigate the dynamics of a bound or mobile
charge in the bulk or interfacial regions of a solid or liquid material (e.g., electrolyte,
semiconducting, mixed electronic–ionic). EIS involves the measurement of impedance by
applying a single-frequency voltage to the cell (i.e. the electrolyte equipped with the
electrodes) and measuring the phase shift and amplitude, or real and imaginary parts, of the
resulting current at that frequency using either an analog circuit or fast Fourier transform
77
(FFT) analysis of the response. Later the measured impedance of a cell is compared to an
equivalent RC model circuit, leading to the electrolyte resistance i.e., impedance R. Using
this measured value of resistance and the geometry of the cell, one can calculate the ionic
conductivity using equation 4. 2.
𝜎 = 1
𝑅
𝑎
𝑏. 𝑑(4.2)
Where, 𝜎, is the ionic conductivity of the electrolyte, a, is the distance between the two
electrodes, b, is the electrode width normal to the current flow and d is the film thickness, as
shown in Fig. 4. 5.
Fig. 4. 5: Schematic of film on a substrate of which impedance is measured.
In the present thesis work, EIS was performed on the deposited samples to measure their
lateral ionic conductivity in air. Measurements were carried out in a two-electrode
configuration geometry using silver paste to ensure the connectivity onto the film surface.
EIS measurements were performed in the 475 – 725 °C temperature range with 25 °C steps.
During the measurement, the frequency was varied from 42 Hz to 1 MHz, with a 0.2 V
alternating voltage signal.
4.3.6 Photoluminescence (PL)
Luminescence is referred to the emission of light by a material through any process other
than blackbody radiation. The term Photoluminescence (PL) [8] narrows this down to any
78
emission of light that results from the optical stimulation of the material. In our everyday life,
PL finds its uses from food to materials research area. For example, many inorganic materials
including, semiconductors, crystalline ceramics, and glasses are studied by PL to probe their
optical properties and to have a look at impurities and defects.
In PL, when a material is irradiated it gains energy by absorbing the photons at some
wavelength and an electron is excited from a low to a higher energy level. Such a process can
be described as a transition from the ground state to an excited state of an atom or molecule,
or from the valence band to the conduction band of a semiconductor crystal (electron-hole
pair creation). After excitation, the system may undergo a non-radiative internal relaxation
involving interaction with crystalline, and then the excited electron moves to a more stable
excited level, such as the bottom of the conduction band. After a system-dependent
characteristic lifetime in the excited state (which can last from picoseconds to many seconds)
the electron returns to the ground state by emitting the excess energy in the form of light
(known as radiative transition). This emitted light is detected as photoluminescence, and the
spectral dependence of its intensity is analyzed to deduce the information about the properties
of the energy states of the material. The time dependent emission can also be monitored to
get informations about energy levels coupling and lifetimes. The light used in PL excitation
and emission usually fall in the range of 0.6-6 eV (roughly 2000-200 nm). Many electronic
transitions of interest lie in this range, and efficient sources and detectors for these
wavelengths are available easily.
In this thesis work, photoluminescence (PL) measurements were carried out at room
temperature in order to investigate optical properties of OVSZ samples. For this purpose, a
Xenon lamp was used as excitation source and an edge filter was used as to cut the excitation
79
wavelength above 325 nm. During the characterization, excitation wavelengths used are 270
nm (4.6 eV), 280 nm (4.4 eV), 290 nm (4.3 eV), 300 nm (4.1 eV), and 310 nm (4 eV).
Detection was made possible using a CCD camera and, in order to compare the PL spectra,
PL intensities were normalized to the excitation power.
4.3.7 Heat flux microsensor
In order to investigate the role of energetics on the growing film, a Heat flux microsensor
(HFM) developed by the GREMI research group from the University of Orléans (Dr. A.-L.
Thoomann [11–15]) was used. The probe was mounted at the substrate position, 6.5 cm above
the sputter target. The heat flux sensor converts heat into an electrical signal [11–15] and is
recorded as an energy density (W/m2) delivered to the substrate – film system during
deposition. Such measured energy actually represents the “global energy” delivered during
the deposition process, it detects simultaneously the contributions of i.e. Ar+ and other plasma
ions bombarding the film, fast particles such as O- or secondary electrons emitted from the
target, electrons, photons, IR photons emitted by the hot sputtered target [13], condensing
sputtered atoms, the enthalpy of formation of metal-oxygen bonds. In this study, the energy
density recorded by the HFM was normalized with respect to the deposition rate in order to
allow the comparison between various deposition conditions.
The heat flux instrument used is built around a commercial sensor from Vattel (Vattel-HFM-
7) and is composed of a thermopile for the energy flux measurements and a resistance
temperature detector (Pt100) for the temperature control. The thermopile used contained
1600 thermocouple junctions per cm. The small diagnostic active area (approx. 17mm) and
the ms time resolution provided a very sensitive probe for energy flux measurements. The
device was calibrated according to NIST protocols using a cylindrical black body. During the
measurement, the temperature of the HFM was maintained around 5 °C using a proper
80
cooling system in order to avoid the heating of the sensor as this can lead to energy loss by
IR emission. Before starting the measurement, a 100 μm thick copper foil used as substrate
was pasted on the HFM sensitive surface using a vacuum-compatible thermal paste (JELT-
6017-COMPOUND SILICONE 20G). The presence of the copper foil only influences the
response time of the system. In order to avoid the growth of thicker film on the copper foil,
the latter was changed periodically.
81
References [1] J.M. Soler, E. Artacho, J.D. Gale, A. Garcìa, J. Junquera, P. Ordejòn, D. Sànchez-Portal, The SIESTA Method
For Ab Initio Order-N Materials Simulation, J. Phys. Condens. Matter. 14 (2002) 2745–2779. doi:10.1088/0953-
8984/14/11/302.
[2] J. Perdew, K. Burke, Y. Wang, Generalized gradient approximation for the exchange-correlation hole of a many-
electron system, Phys. Rev. B. 54 (1996) 16533–16539. doi:10.1103/PhysRevB.54.16533.
[3] V. Bellido-González, B. Daniel, J. Counsell, D. Monaghan, Reactive gas control of non-stable plasma conditions,
Thin Solid Films. 502 (2006) 34–39. doi:10.1016/j.tsf.2005.07.230.
[4] W.D. Sproul, D.J. Christie, D.C. Carter, Control of reactive sputtering processes, Thin Solid Films. 491 (2005) 1–
17. doi:10.1016/j.tsf.2005.05.022.
[5] D. Depla, G. Buyle, J. Haemers, R. De Gryse, Discharge voltage measurements during magnetron sputtering,
Surf. Coatings Technol. 200 (2006) 4329–4338. doi:10.1016/j.surfcoat.2005.02.166.
[6] I. Safi, Recent aspects concerning DC reactive magnetron sputtering of thin films: a review, Surf. Coatings
Technol. 127 (2000) 203–218. doi:10.1016/S0257-8972(00)00566-1.
[7] J. Musil, P. Baroch, J. Vlček, K.H. Nam, J.G. Han, Reactive magnetron sputtering of thin films: Present status
and trends, Thin Solid Films. 475 (2005) 208–218. doi:10.1016/j.tsf.2004.07.041.
[8] C. Richard Brundle, C.A. Evans, S. Wilson, eds., Encylopedia of Materials Characterization, 1st ed., Manning
Publications Co., 1992.
[9] W.H. Bragg, M. A, F.R. S, The Reflection of X-rays by Crystals, R. Soc. 17 (1913) 43.
doi:10.1098/rspa.1913.0040.
[10] E. Barsoukov, J.R. Macdonald, eds., Impedance Spectroscopy; Theory, Experiments and Applications, 2nd ed.,
John Wiley & Sons, Inc., USA, 2005. doi:10.1002/0471716243.
[11] A.L. Thomann, N. Semmar, R. Dussart, J. Mathias, V. Lang, Diagnostic system for plasma/surface energy
transfer characterization, Rev. Sci. Instrum. 77 (2006) 33501. doi:10.1063/1.2166467.
[12] A.L. Thomann, P.A. Cormier, V. Dolique, N. Semmar, R. Dussart, T. Lecas, B. Courtois, P. Brault, Energy
transferred to the substrate surface during reactive magnetron sputtering of aluminum in Ar/O2 atmosphere, Thin
Solid Films. 539 (2013) 88–95. doi:10.1016/j.tsf.2013.05.075.
[13] P.A. Cormier, A.L. Thomann, V. Dolique, A. Balhamri, R. Dussart, N. Semmar, T. Lecas, P. Brault, R. Snyders,
S. Konstantinidis, IR emission from the target during plasma magnetron sputter deposition, Thin Solid Films.
(2013). doi:10.1016/j.tsf.2013.07.025.
[14] P.A. Cormier, A. Balhamri, A.L. Thomann, R. Dussart, N. Semmar, J. Mathias, R. Snyders, S. Konstantinidis,
Measuring the energy flux at the substrate position during magnetron sputter deposition processes, J. Appl. Phys.
113 (2013) 13305. doi:10.1063/1.4773103.
[15] H. Kersten, H. Deutsch, H. Steffen, G.M.W. Kroesen, R. Hippler, The energy balance at substrate surfaces during
plasma processing, Vacuum. 63 (2001) 385–431. doi:10.1016/S0042-207X(01)00350-5.
82
83
5. Influence of oxygen vacancies on
the phase constitution of zirconia
thin films Zirconia (ZrO2) is a polymorphous material which exists in three crystallographic phases
under atmospheric pressure: (i) the monoclinic phase stable up to ~ 1205 °C; (ii) the
tetragonal phase appears from ~ 1205 °C to 2377 °C; and finally (iii) the cubic phase from
2377 °C to 2710 °C (melting temperature) [1]. As cubic zirconia exhibit superior mechanical
and thermal properties over monoclinic zirconia (discussed in chapter 2), the stabilization of
cubic phase of zirconia at room temperature has always been of great importance. For decades
the stabilization has been achieved by the doping the zirconia lattice with cations of lower
valence than Zr (e.g. Y) [2,3] and the resulting material is known as yttria stabilized zirconia
(YSZ) [3]. As doping has its own disadvantages, therefore, to stabilize the high temperature
c-phases of zirconia at room temperature without any doping of yttria, an intense research
has been developed during the last one and a half decade using various synthesis techniques.
The stabilization procedure has been related to the grain size, energy input during growth,
stresses in the film and O vacancies/N atom incorporation in the zirconia lattice [4–14].
However, a consensus over what drives the phase formation in zirconia has not been reached
so far.
84
In this Chapter, we demonstrate that O vacancy incorporation is the sole responsible for the
stabilization of the high temperature c-phase of zirconia, at room temperature. To achieve
this, we coupled cold plasma-based reactive magnetron sputtering experiments to quantum
chemistry based Density Functional Theory (DFT) calculations, performed in collaboration
with Dr. D. Cornil and Prof. J. Cornil, Service de Chimie des Matériaux Nouveaux,
University of Mons.
5.1 Phase stability of oxygen deficient zirconia; quantum chemistry based
DFT calculations
To investigate the influence of oxygen vacancy incorporation on the zirconia phase stability,
O vacancies were introduced in the zirconia lattice randomly, apart from each other and in
the form of clusters (details in chapter 4). Based on DFT calculations, the impact of oxygen
vacancy concentration and their distribution in the zirconia super cells is quantified in terms
of energetics of the ZrO2-x t- and c-phases7. In figure Fig. 5. 1, It is observed that in case of
randomly and apart introduced O vacancies, the c-phase is thermodynamically the most stable
phase if more than 3 at.% of O vacancies are incorporated in the lattice. Similar theoretical
results have also been reported, via a self-consistent tight-binding model, by Fabris et al. [15].
In their publication, these authors suggest that the stabilization of the t- and c-phase of
zirconia can be achieved (only in theory) solely by incorporating O vacancies in the zirconia
lattice. In case of clustered O vacancies, the trend is not clear and is due to the clustering of
O vacancies which result in big holes in the built cell. Finally, we found a very good
agreement from our DFT calculations for the transition enthalpies of ZrO2. The enthalpy for
the m to c transition, equals 14.45 kJ/mol in our case while the one measured by X. Luo et al.
7 Data in numerical form is provided in the Annex I at the end of thesis as well as the energy (eV/atom) is also
presented in the form of graphs for sake of comparison.
85
equals 14.26 kJ/mol [16].
Fig. 5. 1: Influence of O vacancies on zirconia phase constitution. The energies of each phase are
compared as a function of the concentration of oxygen vacancies. Three ways to distribute the
vacancies in the lattice are presented: (a) random, (b) apart from each other, and (c) clustered.
86
5.2 Synthesis of ZrO2-x by reactive magnetron sputtering
To assess experimentally how the incorporation of O vacancies influence the zirconia phase
formation/stabilization 100 ± 10 nm thick ZrO2-x films were deposited inside the transition
zone as well as in the oxidized mode, according to the procedure described in chapter 4. To
deposit the films, Zr target was fed by 200 mA of dc current and the pressure was kept
constant to 10 mTorr. Such deposition parameters were chosen after systematically varying
the deposition parameters.
The chemical composition of the films deposited in such condition is shown in Fig. 5. 2b. It
is observed that the films deposited inside the transition zone at 50, 65, 75, 80 and 100 %
signal (Fig. 5. 2a) are indeed under-stoichiometric and thus contain O vacancies (32, 20, 16,
6, and 3 at.%, respectively). In such under-stoichiometric zirconia, the local Zr charge state
may vary from Zr+ to Zr4+ depending on the surrounding atoms and density of vacancies [17–
21]. Interestingly, the quantum-chemical calculations show that the net charge on zirconium
cation varies from 2.54 |e| for ZrO2 down to 1.87 |e| for 15% of O vacancy. These values are
in good agreement with those reported in [18], which highlight a variation from 2.57 |e| for
ZrO2 down to 2.02 |e| for Zr2O3. According to the DFT calculations presented in Fig. 5. 1 a
and b, such amount of O vacancies should induce the formation of the c-phase. In contrast,
the film deposited in the oxidized mode is stoichiometric. The observed decrease in O
vacancy concentration with the increase in signal (i.e., increase in the target voltage as the
18O2 partial pressure increases) is due to the increased oxide compound formation on the
target surface.
To check the influence of film chemical composition on film crystallinity, grazing incidence
x-ray diffraction (GIXRD) was performed and the resulting diffractograms were compared
87
to ICDD PDF cards representing patterns of the three polymorphs of ZrO2: monoclinic (PDF#
37-1484), tetragonal (PDF# 81-1544) and cubic (PDF# 49-1642). The peak positions of the
films deposited inside the transition zone matched very well with those corresponding to the
tetragonal and/or cubic phases and were hard to distinguish. To unambiguously identify the
phase, we compared the theoretical diffractograms of the t- and c-phases (obtained from the
DFT optimized structures) with the diffractogram of one of our film grown at 100% sensor
signal i.e. film with the lowest concentration of vacancies, as shown in Fig. 5. 3. Furthermore,
Lamas et al [22] have pointed out the splitting of the (400) peak of the c-phase of ZrO2 into
the (004) and (400) peaks in the t-phase of ZrO2, with more than one degree of separation in
their XRD spectra. In Fig. 5. 3, the same peak splitting can be clearly seen for the theoretical
diffractogram of the t-phase while we did not observe any splitting of the (400) peak of the
film grown at 100% sensor signal, thus confirming that the oxygen vacancy-doped films (i.e.
deposited inside the transition zone) belong to the cubic zirconia system.
Fig. 5. 2: (a) Target voltage curve of Zr target as a function of O2 flow shows the transition zone and
working points inside the transition zone as well as in the poisoned zone where (b) chemical analysis
data reveal that ZrO2-x and ZrO2 films were deposited, respectively.
88
The diffractograms of the DFT structures obtained for the various O vacancy concentrations
are shown in Fig. 5. 4(a) together with the diffractograms of the as-deposited films recorded
using an incidence angle of 0.5° (Fig. 5. 4(b)). The remarkable agreement between theory
and experiments demonstrates that the oxygen-vacancy doped films (synthesized inside the
transition zone) are of pure c-phase while the film changes to phase pure monoclinic when
deposited in the oxidized mode, i.e. when containing no O vacancies. However, in DFT based
diffractograms Fig. 5. 4(a), it is observed that some new peaks start to appear with the
increase in O vacancies. The latter could be an artifact coming from the large number of
vacancies in the cell. On the other hand, it is also possible that these peaks are not detected
in the experimental diffractograms because of a too low signal-to-noise ratio.
Fig. 5. 3: Comparison of the diffractograms of the cubic and tetragonal phases obtained at the DFT
level with the experimental diffractogram of the film grown at 100% sensor signal inside the transition
zone (vacancy concentration = 3 % at.). The inset data shows that there is no splitting of the
experimental c(400) peak, as it is observed for the tetragonal (004) and (400) peak.
It has been also reported that incorporating O vacancies induces lattice distortions and result
in contraction of the zirconia lattice, which may result in similar zirconia lattice as in cubic
zirconia due to the coulomb forces between O vacancy-Zr and O vacancy-O atoms [23,24].
89
Furthermore, Fabris et al. [15] have also shown in their work that having a low content of O
vacancies, i.e. 1 at.% (similar to 3.2 mol.% Y2O3), leads to the tetragonal distortion, while
having a higher content of O vacancies, i.e. 4 at.% (=14.4 mol.% Y2O3) results in every
Fig. 5. 4: GIXRD diffractograms of the (a) structures resulting from the DFT calculations
and (b) deposited films, as a function of concentration of O vacancies.
oxygen atom to be a neighbor of a vacant site or at least four of its six neighboring oxygen
atoms, leads to the cubic structure. Moreover, as we do not observe any peak shift in the XRD
90
spectra of the deposited films, is an indication that the deposited films are not stressed. For
these reasons, we believe that the only possible atomistic mechanism behind stabilization of
c-phase of zirconia is lattice distortion caused by O vacancy incorporation in the zirconia
lattice, ultimately forcing the O and Zr atoms to arrange themselves in the high symmetry
cubic crystal.
5.3 Conclusion
In conclusion, based on our DFT calculations and their remarkable agreement with our
experimental data, we conclude that incorporating (as low as 3 at.%) O vacancies is the sole
mechanism responsible of promoting zirconia to the high-temperature c-phase, at room
temperature. In this case, the film deposition was performed with a discharge current of 200
mA, and total pressure of 10 mTorr. However, we observed that deviation from this
“optimized” growth condition (200 mA, 10 mTorr) influence the zirconia phase constitution.
Studying the influence of the working conditions is the subject of chapter 6.
91
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6. Influence of the deposition
parameters on the phase
constitution of oxygen vacancy
doped thin films
In the previous chapter 5, the stabilization of the high temperature c-phase of zirconia at room
temperature is attributed to the incorporation of O vacancies in the zirconia lattice and is
demonstrated for 200 mA and 10 mTorr sputtering condition. A deviation from the
experimental condition (200 mA and 10 mTorr) for which the diffractogram only exhibits the
cubic reflections, lead to the change in zirconia phase constitution. Moreover, the results
presented in the previous chapter 5 are related to films of thickness ≈ 100 nm. In the
literature, it is shown that the film thickness may also influence the film crystallinity,
microstructure, mechanical and optical properties [1–5]. In similar way one can wonder what
would happen in the case of OVSZ thin films if their thickness is increased. Therefore, to
investigate the influence of such deposition parameters, this chapter is divided into two
sections; i) the evolution of the X-Ray diffractograms is presented as a function of the
deposition pressure and discharge current. The evolution of the XRD data is discussed in
94
terms of energy transferred to the substrate – film system during growth. ii) influence of film
thickness on OVSZ films is investigated by varying the films thickness but the film
deposition conditions were kept constant during the whole deposition process.
6.1 Influence of pressure and discharge current
6.1.1 Experimental details
The deviation from the established sputtering condition is carried out by systematically
varying the deposition condition i.e. by working in the transition zone with the help of the
voltage feedback control unit at 5, 10 and 20 mTorr using 200, 300 and 400 mA of discharge
current, fed to the 5 cm in diameter Zr cathode. In this section, the film characterized by 16
at% O vacancy concentration for all above mentioned working condition and 3 at.% O
vacancy concentration, deposited only at 200 mA, 5, and 10 mTorr are discussed. During the
deposition, the substrates were placed at a distance of 6.5 cm from the target surface.
As the pressure and discharge current can have a direct influence on the energetic of the
plasma species, the global energy flux is measured using the heat flux microsensor developed
by the research group GREMI of the University of Orléans [6–8]. The HFM measures the
contribution emanating from the bombardment of Ar+ and other ions, fast particles such as
O- or secondary electrons emitted from the target, plasma e-, hν, IR emitted by the hot
sputtered target, condensing atoms, etc. During each of these deposition conditions, the heat
flux probe was placed at the substrate position. These measurements were performed at
ChIPS lab, in the frame of a one-week exchange program between ChIPS and GREMI
laboratories. Once the global energy flux was recorded by the heat flux probe, then it was
normalized with respect to the deposition rate i.e. the number of film-forming Zr atoms
reaching the surface of the sample as determined from Rutherford Backscattering
95
Spectroscopy (RBS). More technical details on thin film deposition, GIXRD, energy flux
measurement and on RBS can be found in chapter 4.
6.1.2 Results and discussion
6.1.2.1 Evolution of the phase constitution as a function of pressure and
discharge current
GIXRD performed on the 16 at.% O vacancy films deposited for each of the above mentioned
deposition conditions revealed that working at 10, 20 mTorr using a discharge current 200,
300 and 400 mA lead to the formation of c-phase (Fig. 6.1a). On the other hand, using a
discharge current of 300 and 400 mA and working at low pressure of 5 mTorr, lead to the
mixed c- and m-phase formation. However, working at 200 mA with a pressure of 5 mTorr
resulted in a grey area of phase formation i.e. no clear evidence of pure cubic or mixed c- and
m-phase. All the films were deposited at the same position inside the transition zone i.e. at
the same target coverage. Therefore it is believed the films possess the same film chemistry
i.e. the same O vacancy concentration when they are grown using the same set point value
[9]. It should be noted here that the O vacancy concentration was determined for the case of
200 mA, 5 and 10 mTorr using 18O and NRA/RBS analysis. The vacancy concentration was
found to be identical i.e. 16 at.%, corresponding to a set-point value for the voltage feedback
control unit equal to 75 %. Similarly, NRA/RBS analysis of the films deposited using a set-
point 100% at 200 mA, and pressure 5 and 10 mTorr exhibited the identical concentration of
O vacancy i.e. 3 at.%. GIXRD of 3 at.% O vacancy film revealed film deposited at 200 mA,
10 mTorr is pure cubic, while film deposited at 200 mA, 5 mTorr exhibit a mixed c- and m-
phase (Fig. 6. 1b).
96
Fig. 6. 1: GI-XRD (Cu Kɑ1) diffractograms of zirconia thin films containing (a) 16 at.% O
vacancies, deposited at various discharge currents and pressures. (b) 3 at.% O vacancies.
97
According to the above mentioned observations, one could argue that decreasing the working
pressure and/or increasing the discharge current would increase the energy deposited on the
substrate surface during film growth. Moreover working at higher set point (100% set point
in order to synthesize 3 at.% O vacancy concentration films) inside the transition zone for the
same pressure and discharge current, lead to the higher target voltage as compared to the
lower set-point, 75% signal (16 at.%), see Fig. 5. 2. Increasing the current leads to the
production of more numerous ions and electrons, decreasing the pressure results in an
increase of the mean free path and increase in set-point i.e. voltage will lead to the increase
in energy of the depositing ions. Therefore, to study the influence of deposition parameters
on the global energy flux arriving at the substrate, and ultimately on the film phase
constitution, the energy flux was measured using the heat flux probe. To allow the
comparison between the various deposition conditions, the energy flux was normalized by
dividing the flux value by the flux of deposited Zr atoms, as measured by RBS.
6.1.2.2 Normalized energy flux measurements
The normalized energy flux data are presented in Table 6. 1 for each working condition.
Table 6. 1: Global energy flux per deposited Zr atom (eV/Zr atom) at the substrate position is
reported for various dc-RMS working conditions for the deposition of ZrO2-x thin films containing
16 and 3 at.% O vacancies. It has to be noted the values for 300 mA for 5, 10, 20 and for 400 mA,
10 mTorr are extrapolated values.
16 at.% O vacancy concertation
5 mTorr (eV/Zr) 10 mTorr (eV/Zr) 20 mTorr (eV/Zr)
200 mA 1031 598 395
300 mA 1314 762 504
400 mA 1598 927 613
3 at.% O vacancy concentration
200 mA 1313 823 n.a
98
Normalized energy flux data reveal that low energetic values are obtained for a given
discharge current at 20 mTorr working pressure. The normalized energy flux increases
significantly with the decrease in working pressure and reached the keV range on lowest
working pressure i.e. at 5 mTorr. This increase in normalized energy with the decrease in
pressure is primarily related to the increase of the flux and kinetic energy of film-forming
species as the mean free path becomes longer. Other phenomena could also contribute in the
increase of normalized energy flux in this condition, such as the production of fast negative
ions emitted from the oxidized target [10,11]. The later are accelerated in the target sheath
and, in the case of a lower pressure condition, as the target voltage is increased significantly
i.e. 260 V for 5 mTorr and 210 V for 20 mTorr (and thus the kinetic energy of the negative
oxygen ions emitted from the target too) to keep the current constant.
Fig. 6. 2: Influence of energetics on the phase constitution of zirconia thin films containing 16 at.%
O vacancies.
Plotting the resulting film phase as a function of the normalized energy (shown in Fig. 6. 2)
reveal that critical energy level (i.e. ≺ 1000 eV/Zr atom) under which the c-phase is formed
99
could be determined (white zone in Fig. 6. 2). Above 1050 eV/Zr atom (red zone in Fig. 6.
2), the formation of m-phase occurs as well although the film is synthesized in the transition
zone. In between 1000-1050 eV/Zr atom is the grey area with no clarity of phase pure cubic
or mixed c- and m-phase formation. It should be noted that the presence of a critical value of
the normalized energy flux has also been reported by Mraz et al [12] in the case of reactively
sputtered titanium dioxide films. In their study they show that below 540 eV/Ti atom only
the reflections from the anatase phase are obtained. While above 1000 eV/Ti atom, only the
rutile phase appears. In between these two values, a mixture of anatase – rutile phases is
obtained. Cormier et al. [13] have also reported normalized energy flux data in the case of
reactively sputtered titanium dioxide films. In their case it is found that, having ≈7 keV/Ti
atom, the films are composed of anatase phase while for an energy range of 7 – 13 keV/Ti
atom, a mixture of anatase – rutile phase is obtained. Similarly, the influence of energy on
zirconia phases has also been reported by Goedicke et al.[14] in 2000, for films of thickness
equal to 100-150 nm and synthesized by reactive pulsed magnetron sputtering (PMS) at
various deposition pressure. Goedicke et al. also varied the working pressure from 2.2 mTorr
to 26.2 mTorr as well as the target to substrate distance. The sputter power was kept constant.
In their study, when the films were deposited at low pressure i.e. at 2.2 mTorr, the
diffractograms only exhibited the monoclinic lines. On the other hand, when films were
deposited at 26.2mTorr, they exhibited only the cubic peaks. These results are similar to ours,
however, these authors do not specify the elemental composition of their films. They
controlled the oxygen gas inlet by an optical emission monitor but it is not specified if they
work in the transition zone, and they didn’t analyze the XRD data in depth to verify if they
synthesized the tetragonal or the cubic phase. The reason Goedicke et al. provide for such
behavior is related to the energy of the condensing particles. They suggested that the use of
100
higher sputtering pressure reduces the mean free path of the sputtered particles and therefore
the mobility of the depositing species on the substrate. They further relate this lower mobility
to the microporosity of the film and therefore to the reduced hardness and lower compressive
stresses in the film. Goedicke et al. also measured residual stresses and found that, indeed,
films deposited at low sputtering pressure (2.2 mTorr) exhibit high compressive stresses (~ -
1800 MPa) while the films deposited at higher sputtering pressure (26.2mTorr) exhibit low
tensile stresses (~139 MPa). The generation of compressive stresses at low sputtering
pressure could be explained by e.g. the energetic species which get implanted in the growing
film and cause subsurface effects [15]. This implantation leads to lattice defects e.g.,
displacement of atoms from their lattice site. In reactive magnetron sputtering, these high
energetic species could be negatively charged oxygen ions (O-) formed at the oxidized
fraction of the target surface [10,16–18]. As already specified above, these O- ions are
accelerated in the cathode sheath and bombard the growing film with energies in the range
of several hundreds of eV [10]. The energy of these O- ions depend on the magnitude of the
target voltage, while the number of these O- ions is determined by the target coverage by the
oxide layer. In several studies, the influence of these O- ions on the structure formation has
been reported. For example, in 2006, Mraz et al.[11] studied the influence of the these oxygen
ions emitted from the oxidized target during the growth of Nb, Ta, Zr, and Hf oxide films by
reactive magnetron. These authors proposed that the evolution of the crystalline structure of
the transition metal oxide may depend on the presence of O− ion bombardment induced
adatom mobility. Ngaruiya et al. [18] also studied the structure formation of various transition
metal oxides of group 4, 5 and 6 by reactive magnetron sputtering at lower pressure (6
mTorr). Ngaruiya et al. found that Zr and Hf-based sputtering processes allow the formation
of the monoclinic phases of their respective oxides. To the contrary, the other transition metal
101
oxides from group 5 and 6 (Nb, Ta, W, V, and Mo) form amorphous films in the same
deposition conditions. For the zirconium target, Ngaruiya et al. observed that monoclinic
zirconia films were deposited with high deposition stresses (-1500 MPa) in the case of a fully
oxidized target. They attributed the generated compressive stresses in the film to the flux of
oxygen ions emitted from the oxidized target. In the work of Severin [19], nitrogen is added
to the Ar/O2 atmosphere to reduce the flux of fast negative O- emitted from the target. If the
flow of nitrogen is not high enough, the film phase constitution is monoclinic. On the other
hand, when enough nitrogen is added, the phase is cubic. Actually, by adding N2, the reactive
discharge can be stabilized in the transition zone. Later on, Sarakinos et al [20] discussed the
phase formation of hafnium oxide (which is isostructural with ZrO2) by reactive sputtering.
They confirmed that the suppression of the O- bombardment and the incorporation of oxygen
vacancies favors the generation of cubic crystals but they incorporated the vacancies by using
a High-Power Impulse Magnetron Sputtering (HiPIMS) discharge which is known to increase
the ion bombardment of the film during growth [21].
From the above mentioned results, it appears that the presence of energetic species (such as
fast O- ions) increases the value of the energy deposited during growth and the stress levels
in the coatings. This variation in the growth conditions may induce some modifications in
the film phase constitution. From the data displayed in Fig. 6. 1, and 2, it could be speculated
that in our condition, and despite the incorporation of vacancies, the appearance of the
monoclinic phase is the result of the increased energy flux above a certain threshold, which
leads to the build-up of large compressive stress. However, to unambiguously pinpoint the
effect of such energetic species a more detailed investigation must be carried out. Moreover,
as the presented normalized energy flux is a combination of various plasma species
102
contribution therefore, it would be interesting to develop a strategy to figure out the role of
each specie in the phase formation.
6.2 Influence of film thickness on the crystal structure of OVSZ films
6.2.1 Experimental details
OVSZ films with 16 at.% O vacancies were deposited by working inside the transition zone
at 10 mTorr and 200 mA (i.e. the optimized conditions) of discharge current using dc reactive
magnetron sputtering. In order to investigate the influence of film thickness on film
crystallinity, film thickness was varied from 150 nm to 1300 nm. To investigate the influence
of the film thickness on the film crystallinity, GIXRD as well as Bragg-Brentano XRD was
performed on the deposited samples. High resolution transmission electron microscopy (HR-
TEM) was also performed on the deposited samples and for HR-TEM analysis, 2000 nm
thick sample with 16 at.% O vacancies was used and the sample cross section preparation
(sandwich gluing, mechanical polishing and ion polishing) was carried out at -60 °C to avoid
any sample degradation. TEM analysis were carried out at the Institut des Materiaux de
Nantes (Dr. A.A. El Mel and E. Gautron). To investigate the film microstructure, scanning
electron microscopy (SEM) was performed on the cross sections of the 2000 nm thick films.
More technical details regarding the film deposition, HR-TEM, XRD can be found in chapter
4.
6.2.2 Results and discussion
6.2.2.1 Evolution of XRD diffractograms as a function of film thickness
GIXRD and Bragg-Brentano XRD diffractograms of OVSZ films whose thickness equals
150, 275, 615 and 1300 nm are shown in Fig. 6. 3 (a) and (c), respectively. It has to be noted,
as the thicker (615 and 1300 nm) films have more material volume as compared to thinner
103
(150 and 275 nm) films, giving higher intensity of the XRD peaks as compared to thinner
ones. Therefore, to avoid any peak suppression, a narrow scan in the range of 26 - 33° was
also performed on 150 and 275 nm thin films with a 10 time more acquisition time as
compared to 1300 nm thick film and no additional peaks were observed in this 26 - 33° range
Fig. 6. 3: (a) GIXRD, (b) GIXRD narrow scan with 10 time more acquisition time, (c) Bragg-
Brentano diffractograms of 16 at.% O vacancy containing zirconia films of various thickness.
104
(diffractograms shown in Fig. 6. 3b). The first feature that can be noticed with the increase
in film thickness from the GIXRD as well as Bragg-Brentano spectra (shown in Fig. 6. 3) is
the competition between the c(111) and the c(200) peak intensities. It is observed that the
c(111) orientation dominates at lower film thickness i.e. at 150 nm, while with the increase
in film thickness i.e. around 600 nm, the c(200) orientation overtakes the c(111) peak
intensity. Similar feature has also been observed by Nouveau et al. [5] in Bragg-Brentano
XRD diffractograms for CrN films which have a similar cubic crystal structure as cubic
zirconia. Nouveau et al. showed that the compressive stress varies with respect to the film thickness
[5,22] and they observed three growth regimes : Regime A) (up to 150nm) where the compressive
stress keeps increasing, B) (150-500nm) the film changes its structure i.e. preferred plane orientation
to release compressive stress, C) (above 500nm) the compressive stress value is stable and relaxation
is induced by the growth of less dense planes. Another feature can be observed in Fig. 6. 3: the
film with the lowest thickness (150 nm) exhibit pure c-phase, however with the increase in
film thickness a small bump appears around 28°. This is a reflection related to the monoclinic
phase of zirconia. At 1300 nm, this peak associated with the appearance of the monoclinic
phase is significant.
6.2.2.2 Analysis of film cross-section by SEM and HRTEM
To investigate the appearance of monoclinic peak, SEM cross-sectional images and HR-TEM
analysis were performed on 2000 nm thick OVSZ film. Fig. 6. 4 shows the cross-section
images taken at (a) middle, (b) near the top, (c) middle of film cross-section and (d) film
substrate interface reveal that, near the interface the microstructure of the film indicates a
dense nucleation zone (Fig. 6. 4d) followed by columnar feather-like growth (Fig. 6. 4 (b)
and (c)).
105
Fig. 6. 4: SEM cross-sectional images of 2000 nm thick zirconia film having 16 at.% O vacancies (a)
middle, (b) near the top, (c) middle of film cross-section and (d) film-substrate interface. For all these
cross-sectional images small white bar at the right bottom of images correspond to the scale bar which
is 100 nm.
Fig. 6. 5: Cross sectional image of the film and the substrate as obtained by TEM. The substrate is
Si (100). Colored circles indicate the area where selected area electron diffraction analysis was
carried out.
106
Fig. 6. 5 shows the cross-section of the 16 at.% O vacancy film of thickness 2000 nm. To
perform the SAED (Selected Area Electron Diffraction) analysis along the substrate – to –
film surface axis, different areas were selected from the film-substrate interface, as shown in
Fig. 6. 5. SAED analysis, close to the substrate, revealed that the film is textured and belongs
to cubic structure. This area of film could be related to the 150 nm and 275 nm-thick films
from which the XRD analysis which revealed the same cubic structure (Fig. 6. 3). SAED
analysis performed while moving away from the film-substrate towards the middle of the
film also revealed the presence of cubic crystals. No monoclinic crystals were detected as
observed in XRD diffractograms (Fig. 6. 3). The film away from the interface has feather-
like columnar structure as shown in Fig. 6. 4. To locate the monoclinic crystals, a much more
careful investigation was carried out at the intersection of those feathers, as shown by the
yellow rectangle in Fig. 6. 6.
107
Fig. 6. 6: HR-TEM images of 16 at.% OVSZ film at around 850 nm from substrate.
At the intersection of two feathers, due to the superposition of feathers, it was hard to
distinguish if the phase is amorphous or not. To avoid this, Fast Fourier Transformation (FFT)
was used. FFT analysis revealed there exists crystals which could only be indexed by
monoclinic phase (Fig. 6. 7). Furthermore, the monoclinic crystals were localized around 850
nm and 2000 nm away from the film-substrate interface hence confirming the initial XRD
analysis (Fig. 6. 3). The above mentioned results highlight that, despite the film growth
conditions are kept constant during the whole deposition process, the film microstructure and
phase constitution may evolve.
108
Fig. 6. 7: Electron diffraction of 16 at.% OVSZ film obtained at around 850 nm from the substrate-
film interface. The diffraction spots are indexed with monoclinic phase.
6.3 Conclusions
In the section 6.1 we have presented that oxygen vacancy doped cubic zirconia coatings are
actually obtained in some specific working conditions i.e. low current, moderate to high
pressure. As the current is increased and/or the pressure decreased, monoclinic peaks appear
on the diffractograms, although the films are deposited inside the transition zone. The
evolution of the film phase constitution with the deposition parameters is interpreted in terms
of energy flux at the substrate surface. Pure cubic oxygen vacancy doped zirconia thin films
are obtained by working in conditions for which the normalized energy flux lies below ~1000
eV/Zr atom. Above this threshold, up-to 1050 eV/Zr atom the diffractograms exhibit the grey
area with no clarity of phase pure cubic or mixed c- and m-phase formation but, above 1050
109
eV/Zr atom both cubic and monoclinic peaks. The increase in film stress as a result of the
increased normalized energy flux is proposed to explain this behavior.
In the section 6.2, 16 at. % OVSZ film with various thicknesses, ranging from 150 to 2000
nm, have been deposited. GIXRD diffractograms reveal the appearance of monoclinic peaks
with the increase in film thickness, while the films are deposited in the transition zone. In the
case of thicker films, SEM cross-sectional images show dense nucleation near the substrate-
film interface followed by columnar feather-like growth of the film away from the substrate-
film interface. TEM analysis of these thick films allows to determine that the monoclinic
crystals indeed appear in the upper part of the film, e.g. above 850 nm. The reason for the
appearance of the monoclinic crystals close to the film surface is not yet clarified.
Nevertheless, an important information can be extracted from these results. One should pay
attention when synthesizing (thick) films because, even though the deposition conditions are
kept constant throughout the deposition process, the film microstructure and its phase
constitution may vary.
110
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7. Thermal stability of OVSZ thin
films
Zirconia is one of the materials which exhibit very low thermal conductivity (1.2 - 2.6
Wm−1K−1[1]) making it a good candidate for thermal barrier coatings (TBCs)[2]. However,
the use of zirconia in high temperature applications is restricted by the change in its volume
(~5 Vol.%)[3] due to phase transformation at elevated temperatures as well as while cooling
down the device. This phase transformation and change in volume result in the crumbling of
the zirconia-based components. Since last couple of decades, phase transformation of
zirconia upon thermal cycling is avoided by stabilizing the high temperature c-phase of
zirconia. Typically, stabilization is achieved by doping zirconia with cations of lower valence
than Zr e.g., Y (as discussed in chapter 2). In such stabilization process some Zr4+ cations are
replaced by Y3+. in this situation, to maintain the charge neutrality, one oxygen vacancy is
created for each two substituting Y3+ cations. This makes YSZ not only useful for TBCs but
the presence of vacancies also makes it a good ionic conductor used as an electrolyte in solid
oxide fuel cells (SOFC)[4–6] and in oxygen sensors[7].
In chapter 5, it is demonstrated that the zirconia high temperature c-phase can be stabilized
at room temperature by incorporating the right amount of oxygen vacancies in the lattice [9].
This material was then called oxygen vacancy stabilized zirconia (OVSZ). By definition,
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OVSZ contains vacancies and, like YSZ, it could be useful as ion conductor in Solid Oxide
Fuel Cell.
In this chapter, the thermal stability and structural transformation at elevated temperature is
investigated using temperature-resolved XRD measurements in order to verify if, indeed,
OVSZ thin films can be used in high-temperature applications.
7.1 Experimental details
To investigate the thermal stability of the OVSZ coatings, two identical sets of 100±10 nm
thick c-ZrO2-x films containing ~16 and 3 at.% O vacancies were deposited according to the
procedure described in chapter 4. The samples were deposited on Si(100) substrates whose
thickness is 525 ±20 μm (Siegert Wafer).
To analyze the film crystallinity of the as-deposited films, grazing incidence x-ray diffraction
(GIXRD) is performed using PANalytical Empyrean with a Cu-Kα1 radiation source (λ =
0.15406 nm). The diffractograms were recorded with a step size of 0.07° with a duration of
15 seconds using an incidence angle of 0.5° at 40 mA, 45 kV of generator settings.
Temperature-resolved XRD measurements are performed on a Bruker D8 Discover
diffractometer, using a Co-K⍺ radiation (λ = 0.17902 nm) and a parallel beam with a diameter
of 1mm. The measurements were carried out at the Institut Jean Lamour, in Nancy (Pr. J.F.
Pierson and Dr. P Boulet). Samples were heated either in air or in a N2 ambient with a Domed
Hot Stage (DHS1100 from Anton Paar), installed on the goniometer. The heating speed was
8 °C/s. During annealing in the N2 ambient, a flow of 1-2 L/h of N2 was blown at the sample
surface and the pressure inside the dome was same as outside the dome, i.e. 1 atm. These
measurements were carried out in grazing incidence geometry as well but at a grazing angle
of 3° with a step size of 0.025° and duration of 4 s. During the scan, for each working
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temperature, the annealing stage temperature was kept constant for 2 hours to allow the scan
to complete at that specific temperature. Further, the detection was assured by a scintillator
counter before which long Sollers slits are installed to select the diffracted X-rays.
7.2 Results and discussion
XRD diffractograms of the as deposited films containing 16 and 3 at.% O vacancies acquired
at an incidence angle of 0.5° are shown in Fig. 7. 1. The diffraction peaks from the cubic
phase of zirconia are identified.
Fig. 7. 1: GIXRD spectra of the as-deposited films recorded using Cu x-ray source at an incidence
angle of 0.5° at room temperature.
TR-XRD diffractograms acquired at an incidence angle of 3° for the same films containing
16 and 3 at.% O vacancies annealed in air and N2 from 200 °C to 900 °C and then cooled to
100 °C are shown in Fig. 7. 2. It is observed, for low temperatures (i.e. from 200-300 °C),
that the diffractograms do not show any peaks. Let’s recall that, for these specific
experiments, the films are scanned at an incidence angle of 3° which is a higher value than
for the diffractograms presented in Fig. 7. 1, where the same films where scanned at an
incidence angle of 0.5°. It is known that in GIXRD, the intensity of the peak is very much
dependent on the incidence angle i.e. at low incidence angles, close to the critical angle of
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115
Fig. 7. 2: TR-GIXRD spectra of the films containing approximately 16 at.% O vacancies annealed in
(a) air and (b) N2. Films containing 3 at.% O vacancies are also annealed in (c) air and (d) N2. X-ray
source Co and x-ray incidence angle 3° was used.
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total reflection, the beam path inside the probed material is increased (beam path ≈ t/Sinθ, t
is film thickness and θ is beam incidence angle[10]). As a result, an increase in the intensity
of the reflected peak is observed when films are scanned at lower incidence angles. Further
annealing of films at 350 °C shows that the cubic diffraction peaks become more intense,
meaning the films get better crystallized. This could be explained in terms of diffusion of
atoms inside the film material at higher temperature. This phenomenon is supported by the
observed increase in the intensity of c(111) peak as a function of temperature, shown in Fig.
7. 3. This situation allows already existing small cubic crystals to grow in size by gathering
diffusing atoms. At 350 °C those crystals are big enough (average grain size 12-15 nm at 350
°C while average grain size of as deposited films is 8-9 nm) to be detected in the used TR-
XRD setup. The evolution of the crystallite size as a function of the temperature is shown in
Fig. 7. 4. The crystallite size is calculated using the Scherrer Equation [11] and taking the
c(111) peak into consideration. Data show for 16 and 3 at.% Vo films that the crystallite size
increases monotonously with the increase in temperature, up to 700 °C. After passing 700 °C
a relatively fast increase is observed. Further, on cooling down the films after reaching 900
°C, it is observed that, in both cases, the average crystallite size remains around 20 nm.
Fig. 7. 3: Evolution of c(111) and m(111) as function of temperature of film containing 16% O
vacancies.
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Fig. 7. 4: Evolution of cubic crystallite size as a function of temperature, calculated using C(111)
peak of films containing (a) 16 at. % of O vacancies and (b) 3 at. % of O vacancies.
The analysis of the annealed samples shows that, after passing 750 °C, m(111) peak appears
and gets more and more prominent with the increase in temperature. It is also observed that
nor the annealing ambient (i.e., air or N2) nor the Vo concentration has any influence on the
appearance of the m-peak. However, with the increase in annealing temperature a significant
shift to lower angles in all peak positions is observed on the TR-XRD spectra (however this
peak shift is different for each peak, not shown here), whatever the oxygen vacancy
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Fig. 7. 5: Shift of C(111) peak as a function of temperature of the films containing (a) 16 at.% of O
vacancies (b) 3 at.% of O vacancies to lower angles during annealing in Air and N2 and recovery
while cooling down.
concentration (16 or 3 at.%) or the annealing ambient is (air or N2). At 350 °C, the c(111)
peak exhibits a shift of less than 0.2°. With the increase in annealing temperature, the c(111)
peak shifts further to lower angles due to the thermal expansion. At 900 °C, a maximum shift
to lower angle of about 0.5° is reached as shown in Fig. 7. 5. While on cooling down the films
it is observed that the c(111) peak tries to return to its original position. One should also note
that the m(111) peak also did not vanish upon cooling the films to room temperature (shown
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in Fig. 7. 2). It is also noted from Fig. 7. 2 and Fig. 7. 5, that the films containing 16 and 3
at.% Vo and annealed in air and N2 show a similar evolution.
The more detailed analysis of the temperature–resolved diffraction data of 16 and 3 at.% Vo
reveals that, after passing 750 °C, the area under the c(111) start to decrease (Fig. 7. 6). On
the other hand, at this precise temperature, the m(111) peak appears as shown in Fig. 7. 2.
The variation of the area under the m(111) peak as a function of temperature is almost
identical (in absolute value) to the decrease in area under the c(111) peak, as shown in Fig.
7. 6 for 16 at.% and 3 at.% Vo. Such a decrease in area under the c(111) peak after passing
750 °C would highlight the destruction of cubic zirconia crystals. To the contrary, monoclinic
crystals are formed. From the respective variation in intensities, one could, at first sight,
believe that all the monoclinic crystals are coming from the transformation of the cubic
crystals after 750 °C. Here it should be highlighted that the diffraction intensities related to
the c(111) and m(111) reflections are not identical when looking at the synthetic XRD
spectrum constructed from monoclinic and cubic unit cells, built from quantum chemistry–
based calculations. These calculations and the ZrO2 unit cell utilized to generate the synthetic
XRD spectra are presented in chapter 5. From the synthetic spectra it is found that the
relationship between the diffracted intensities, for a single diffracting unit cell, is Im(111) = 0.35
Ic(111). So having almost the same decrease and gain in area by the cubic and monoclinic peaks
respectively would point out there are more monoclinic crystals formed than cubic crystals
destroyed.
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Fig. 7. 6: Area under the c(111) and m(111) peak as a function of temperature of film containing 16
at.% O vacancies. After passing the 750 °C a drop in c(111) and increase in m(111) peak area is
observed. Similar evolution is observed for film containing 3 at.% O vacancies.
According to these observations, the appearance of m-phase could eventually be explained
by two mechanisms; i) the oxygen uptake by a small fraction of the annealed OVSZ thin film
from the surroundings after passing 750 °C. That part of the film would therefore turn into
stoichiometric m-ZrO2 and/or ii) the martensitic transformation of existing cubic crystals, as
in case of YSZ. It should be also noted that monoclinic structure is the thermodynamically
stable phase below ~1200°C, for fully stoichiometric ZrO2.
Since the OVSZ thin films are, by definition, under-stoichiometric, the possibility of oxygen
uptake and the annihilation of the oxygen vacancies by the incorporation of O from the
surrounding ambient cannot be ignored. To this this purpose, RBS and NRA measurements
were carried out on the OVSZ thin films (deposited using 18O and a graphite sheet as
substrate) at room temperature and after annealing at 750 °C for two hours. It was observed,
due to the annealing at 750 °C, 18O was completely replaced by 16O. Here, it should be
mentioned that the top view of the films grown on the graphite substrate shows the presence
of cracks already before the annealing process, Fig. 7. 7. The
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Fig. 7. 7: SEM top view images of as deposited films showing cracks. Films are deposited on
graphite substrate for RBS and NRA analysis.
presence of cracks most likely change dramatically the way the film interacts with the
surrounding atmosphere, and the amount of oxygen it takes up upon annealing. However,
because of some instrumental reasons it was hard to calculate the 16O uptake concentration
from the surrounding ambient. Furthermore, as it has been demonstrated in the previous
chapter 5, that oxygen vacancy incorporation promotes the formation of the c-phase at low
temperature and, as the films still retain some of the c-phase up to 900 °C and even on cooling
down to 100 °C (Fig. 7. 2), it seems not appropriate to claim that oxygen vacancies are
completely annihilated and the films fully turns to stoichiometric ZrO2 at such higher
temperatures (when the film are grown on silicon single crystals). However, it could be that
only a small volume fraction of the film turns to ZrO2. The monoclinic crystals could nucleate
at these particular places. The formation of a stoichiometric monoclinic (passivation) layer
on the film surface would impede the further oxidation of the under-stoichiometric zirconia
left below. In this condition, the diffraction peaks related to the presence of the cubic phase
still remain intense at 900°C and as well as when cooled down to 100 °C, as observed in Fig.
7. 2 and Fig. 7. 6. On the other hand, here one should also note that the evolution of the film
phase constitution as a function of temperature is identical whatever the annealing
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experiments are performed in air or in the nitrogen ambient. It could be that the latter contains
enough oxygen or water vapor to allow the partial oxidation of the ZrO2-x film.
The 2nd possible mechanism is related to the martensitic transformation of existing cubic
crystals into monoclinic crystals. It is similar to c/t to m phase transformation in sol-gel based
ZrO2 films [12] or stress/pressure induced t to m phase transformation in YSZ [13–15].
Mehner et al. [12] deposited c/t, 1000 – 1200 nm thick, zirconia films on steel substrates
using sol-gel method and studied their thermal stability. Mehner et al. zirconia films were
stable up-to 600 °C and after that temperature, monoclinic peaks appeared on their Bragg-
Brentano diffractograms. They also observed a small shift in the diffraction peak position to
lower angles with the increase in temperature, similar to the shift monitored in the present
study (shown in Fig. 7. 5). Mehner et al [12]attributed the shift to the generation of
compressive stresses in film during annealing. As the buildup of compressive stresses parallel
to the surface causes a vertical expansion of the deposited film and thus results in the increase
of the distance between the planes which are parallel to the substrate surface, the observed
shift in peak to lower angles is an indication of the presence of in-plane compressive stresses
in the films upon annealing. To calculate such annealing generated stresses, Mehner et al.
performed Sin2ψ method on his samples using the c(333) peak. They found below 625 °C
films did not possess high stresses but above 625 °C films exhibit ≈ −750 MPa compressive
stress. They correlated it to the onset of monoclinic phase formation. In the present study,
due to the film thickness (≈ 100 nm), the grazing incidence XRD geometry was used, and in
such geometry the planes probed are not all parallel to the sample surface. Therefore, it’s
hard from the present measurements to calculate the residual stresses from such peak shifts
(Fig. 7. 5). However, since stresses are also generated due to the difference in film-substrate
thermal expansion coefficients (TEC), therefore, a theoretical calculation can be made to
123
estimate the thermal stresses. As Si has a lower TEC (3.6×10-6/K[16]) as compared to c-
zirconia (8.8-10.6×10-6/K[1]) and as the deposited film is bound to the substrate, at elevated
temperature, the film will experience compressive stresses. Such thermal stresses arising
from the mismatch between the film and substrate TEC can be calculated using Eq. 6. 1 where
𝐸𝑓 is the Young’s modulus, 𝑣𝑓 is the Poisson ratio, 𝛼𝑓 and 𝛼𝑠 are the respective TECs of the
film and substrate, 𝑇𝑠 is the substrate temperature before starting the annealing (25°C), and
𝑇𝑎 is the temperature during the measurement. A negative value of 𝜎𝑡ℎ corresponds to
compressive stress.
𝜎𝑡ℎ =𝐸𝑓
1 − 𝑣𝑓(𝛼𝑓 − 𝛼𝑠)(𝑇𝑠 − 𝑇𝑎) (6.1)
To assess the theoretical thermal stress values in such deposited films, the Young’s modulus
and Poisson ratio was taken from the literature (E =200 GPa [17,18], , 𝑣𝑓 = 0.25 [12]) and
we assume that the zirconia Young’s modulus variation with the temperature is negligible.
The calculated compressive stress was found to increase with the increase in temperature as
shown in Fig. 7. 8. Calculation indicates that, at 750 °C, the film stress reaches a value of -
856 MPa. At 900 °C, the value reaches -1032 MPa. Comparing these results with Mehner’s
et al. report (c/t to m phase transformation occurs at ≈ −750 MPa, at) it appears that OVSZ
film are more resilient, i.e. stable up to ≈ −856 MPa, than the sol-gel based zirconia films of
these authors. Moreover, it can be assumed that in such circumstances, where the
compressive stresses are building up inside the film, another probable reason for the
appearance of the m-peak (beside oxygen uptake from the annealing atmosphere) is that some
cubic crystals turn into monoclinic crystals in order to cope with the mechanical stress, i.e.
through a martensitic, diffusion-less, transformation. So, per these arguments, it seems
reasonable to assume that our OVSZ films reached a critical threshold of stress after passing
124
750 °C, and such stresses might force the high symmetric cubic structure to transform in to
the low symmetric monoclinic structure. From our data, this critical threshold value of stress
appears to be in the range of (≈ −856 MPa), as shown in Fig. 7. 8.
As a final word, it should be added that when thermal cycling was performed (6 times in the
present study) below 700 °C, the diffractograms of the OVSZ films only exhibit the c-phase
reflections (data not shown).
Fig. 7. 8: Theoretical residual stress evolution as a function of temperature. Vertical dashed lines
show the temperature where monoclinic peak appeared.
7.3 Conclusion
From the above presented data, we conclude that oxygen vacancy stabilized cubic zirconia
(OVSZ) films deposited on Si wafers are stable up-to 750 °C. It is suggested that a
combination of two mechanisms is responsible for the appearance of m-phase above the
critical temperature of 750 °C; i) oxygen uptake by a fraction of OVSZ thin film from the
surroundings which leads to the formation of stoichiometric ZrO2. The later crystallizes under
the monoclinic structure, thermodynamically stable in such circumstances. ii) The martensitic
transformation of existing cubic crystals, due to the compressive stress forcing the cubic
125
crystals to rearrange themselves in to the low symmetric m-phase. However, to
unambiguously separate these two mechanism a more detailed study is needed.
The development of such high temperature stable c-zirconia film containing oxygen
vacancies might not only be helpful in thermal barrier coating applications but also in SOFC
applications, as ionic conductive electrolyte membranes. This thin film material would allow
to lower the operating temperature of such devices while offering high ionic conductivity.
The analysis of the ion conductivity of the as-deposited OVSZ films is the subject of the next
chapter of this thesis work.
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8. Ionic conductivity of OVSZ thin
films
Use of stabilized zirconia as an electrolyte in SOFCs has always been a top choice due to its
stability at operating temperatures, good ionic conductivity and because of its desirable
chemical stability in both oxidizing and reducing atmospheres [1]. The ionic conductivity in
stabilized zirconia is attributed to the mobility of oxygen vacancies, which are created by the
addition of aliovalent dopants, as discussed in chapter 2. It was first expected that the ionic
conductivity will increase with the increase in O vacancy concentration i.e. by increasing the
dopant content. However, later it was observed the maximum ionic conductivity in yttria-
stabilized zirconia (YSZ) occur at 7 – 9 mol% Y2O3 at 327 – 1227 °C [2,3]. On the other
hand, the incorporation of a higher amount of Y2O3 was found to lower the mobility of O
vacancy by increasing the diffusion energy across an Y–Y common edge as compared to the
diffusion across one with a Zr–Y common edge [2]. In the present chapter, we will measure
the ionic conductivity of Oxygen Vacancy Stabilized Zirconia (OVSZ) thin films in order to
establish if such material can also be utilized as an electrolyte membrane e.g. in SOFC
applications. In the chapter 7, the stability of the magnetron sputtered OVSZ films is already
demonstrated and OVSZ can withstand annealing procedure up to 750 °C.
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In this section, we measure the ionic conductivity and investigate the influence of O vacancy
concentration, film thickness, and the nature of the substrate on the ionic conductivity of
OVSZ thin films.
8.1 Experimental details
In order to perform this study, c-phase OVSZ thin films containing 3 or 16 at.% of O
vacancies and characterized by a thickness of 10, 20, 50 or 100 nm were prepared using dc-
reactive magnetron sputtering, as described in chapter 4. The films were deposited on
polished NdGaO3 (100) substrates (size 5 x 5 x0.5 mm3) purchased from Crystal Germany.
To study the influence of substrate, films containing 16 at.% O vacancies of thickness 10, 20
and 50 nm were also deposited using the same deposition conditions on polished MgO
substrates (size 5 x 5 x0.5 mm3) purchased again from Crystal Germany. To determine the
ionic conductivity, electrochemical Impedance Spectroscopy measurements were carried out,
at the Department of Energy Conversion and Storage, Technical University of Denmark with
the help of Dr. S. Sanna and Dr. V. Esposito), to measure the ionic conductivity at various
temperatures. Four-probe method was also used on the samples containing 3 and 16 at.% O
vacancies of 100 nm thickness deposited on Si (100) and borosilicate glass in order to
measure the film resistivity.
8.2 Results and discussion
The Arrhenius plot, i.e. the plot of the logarithm of the ionic conductivity σ as a function of
the temperature T, for 10 nm thick OVSZ thin film containing 3 at.% O vacancies is shown
in Fig. 8. 1. According to this figure, the oxygen vacancy doped zirconia thin film is
exhibiting a measurable ionic conductivity in a temperature range which is comparable to
SOFC systems.
129
For the sake of comparison, the logarithm of the ionic conductivity of an YSZ film (thickness
= 15 nm) containing a similar amount of O vacancies (1 at.% O vacancy = 3.2 mol% Y2O3)
and reported by Kosacki et al. [1] is also presented as a function of the temperature on Fig.
8. 1.
Fig. 8. 1: The Arrhenius plot of as deposited OVSZ thin film containing 3 at.% Vo (blue data set) is
compared to an YSZ thin film (red data set, from Kosaki et al [1]) presenting similar physical
characteristics.
The logarithm of the ionic conductivity is proportional to the temperature [4,5]. In both cases,
an increase in ionic conductivity with the increase in temperature is observed (Fig. 8.1). Data
shows that the 3 at.% OVSZ thin film exhibits ~1 S.cm-1 ionic conductivity at 650 °C. The
conductivity of the film rises to ~7.4 S.cm-1 at 725 °C. On the other hand, the YSZ film
exhibits ~ 0.3 S.cm-1 at 725 °C [1]. This is almost 24-fold increase in the ionic conductivity
exhibited by OVSZ as compared to YSZ films.
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The ionic conductivity 𝜎 can be expressed as 𝜎 = 𝑛𝑞𝜇 , where n is charge carrier
concentration (cm-3), q is the electric charge of the moving ion (here in our case O2- ions) and
μ is the mobility of charge carriers (meaning mobility of O ions) (cm2 s-1 V-1) [6]. As the
charge carrier and their number i.e. O vacancy concentration, in 3 at.% OVSZ thin film and
in a 10 mol% YSZ is almost the same, therefore, this lead to the conclusion that the mobility
of O ions is higher in OVSZ film as compared to YSZ coating because of the high ionic
conductivity measured (~7.4 S.cm-1) at 725 °C. This is also consistent with Pornprasertsuk et
al. [2] who suggested that the ionic conductivity in YSZ is lowered due to the bigger ionic
radius of Y3+ (0.90 Å) as compared to Zr4+ (0.72 Å) resulting in smaller space for the O ions
to move. Since the ionic conductivity also depends on the number ‘n’ of charge carriers
present in the material, therefore based on this, one can postulate that by further increasing
O vacancy concentration in OVSZ, most probably the ionic conductivity will further increase,
as there are no dopants.
Fig. 8. 2: Arrhenius plot of OVSZ thin film containing 16 and 3 at.% Vo. the data reported by
Kosaki et al for 15nm thick YSZ film is also plotted for comparison.
131
This is not the case as shown in Fig. 8. 2. Data reveals that, when the O vacancy concentration
is increased from 3 at.% to 16 at.%, the ionic conductivity is not significantly modified but
its higher than YSZ. The ionic conductivity in case of 16 at.% O vacancy is found to be equal
to ~0.9 S.cm-1 at 650 °C and rises to ~3 S.cm-1 at 725 °C. This behavior in ionic conductivity
for the OVSZ thin film containing 16 at.% O vacancies as compared to 3 at.% O vacancy
concentration film could be explained in terms; i) lower film density of 16 at.% O vacancy
film as compared to 3 at.% O vacancy film as the number of Zr atoms for 16at ,% was found
lower (5.42 E+17/cm2) as compared to 3 at.% (7.5 E+17/cm2) by RBS measurements. ii) by
lattice distortion, induced by the larger number of O vacancies. Such lattice distortion around
O vacancy in pure zirconia has already been reported in several studies [7–10]. Our quantum
– chemistry based calculations results also show that the cubic cell is slightly distorted as the
number of oxygen vacancies is increased as shown in Fig. 8. 3.
Fig. 8. 3: Zirconia cell after relaxation with (a) 0 at.% O vacancy (b) 10 at.% O vacancy, blue atoms
represent Zr atoms and red atoms represent O atoms.
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Since for high ionic mobility, the lattice has to be highly symmetric [5,6,11], the
incorporation of a larger number of O vacancies (e.g. 16 at.%) in the crystal most probably
leads to a pronounced alteration of the lattice. This situation ultimately results to a relatively
lower ionic mobility of the charge carriers and thus, to a lower ionic conductivity compared
to what is expected for such heavily doped zirconia films. Similar finding was also reported
by Kimpton et. al. [12] who suggested that excessive distortion of the oxygen ion conduction
path by large size ionic dopants is responsible for the decline in ionic conductivity. Further,
the activation energies (Ea) of 10 nm films containing 16 and 3 at.% O vacancies were also
calculated from the slope of the linear fit and were found to be 1.28 and 1.59 eV respectively
(Fig. 8. 2). This puzzling result could be (partially) explained by the fact that having higher
amount of O vacancies means smaller jump to diffuse from position A to B. While in case of
lower amount of O vacancies, the diffusion path from A to B can be longer, so ion would
need more energy.
The influence of the thickness on the ionic conductivity of OVSZ thin films containing 16
and 3 at.% O vacancies is shown in Fig. 8. 4 (a) and (b). It is observed in Fig. 8. 4 (a) and (b)
that the ionic conductivity of OVSZ films increases with the decrease in film thickness. At
725 °C, the ionic conductivity of 100, 50, 20 and 10 nm thick films containing 3 at.% O
vacancies was found to equal ~ 0.04, 0.05, 0.1 and 7.4 S.cm-1, respectively. The ionic
conductivity of 100, 50, 20 and 10 nm thick films containing 16 at.% O vacancies follows
the same trend and was found to be respectively equal to ~ 0.02, 0.06, 0.2 and 3 S.cm-1 at 725
°C. Such an impact of the film thickness on the ionic conductivity is known for oxide ion
conductors and is attributed to the enhancement of ionic mobility due to interfacial
effects[1,6,11,13]. These effects are divided into two categories; (i) the formation of space
charge layer (SCL) at the interface. The SCL is a region of continuum electric charge
133
occurring in dielectrics, where the mobile charge carrier diffuses away by the influence of an
electric field leaving a depletion region of, for example, ionized species such as oxygen
vacancies. Such depletion is electrostatically balanced by accumulation of opposite sign
carriers [11,14]. (ii) The Lattice Strain, induced by slightly or largely mismatched interface
between the substrate and the ion conducting film itself [11]. The degree of disorder at the
interface plays an important role in determining which of these factor is dominant because it
directly controls the formation of defects at the interface, e.g., at the simplest level of analysis,
higher the disorder higher the concentration of the segregated defects at the interface [11].
This higher disorder at the interface caused by lattice strain or by space charge region will
lead to several orders of magnitude greater diffusivity as compared to that of the lattice
[1,6,11,13]. Indeed, it has been shown by Kosacki et al. [1] that such interfacial effects get
more pronounced for films having a thickness lesser than 60 nm. However, we observe a
pronounced increase in the ionic conductivity when thickness of the OVSZ films is
134
Fig. 8.4: Influence of film thickness on the ionic conductivity for a) the 3 at.% vacancy
concentration and b) the 16 at.% vacancy concentration.
135
Fig. 8. 5: Logarithm of the ionic conductivity as function of thickness of films containing 16 and 3
at.% O vacancies, at 725 °C.
decreased from 20 nm down to 10 nm, as shown in Fig. 8. 5. The increase in ionic
conductivity for the 16 at.% O vacancy film is 2 fold as the thickness is reduced from 50 to
20 nm but a dramatic increase (14 fold) is observed when the thickness decreases further
from 20 nm down to 10 nm. Similarly, for the 3 at.% O vacancy film, a colossal (73 fold)
increase is observed when the film thickness is decreased from 20 nm to 10 nm. This increase
in ionic conductivity indicates that the interfacial effects are pronounced for the 10 nm-thick
films.
136
Fig. 8. 6: Influence of interfacial strain on the ionic conductivity of films containing 16 at.% O
vacancies and having various thickness.
To separate the effect related to the space charge layer and the interfacial lattice strain, films
of thickness equal to 50, 20, and 10 nm containing 16 at.% O vacancies were deposited on
MgO (Zirconia - MgO lattice mismatch equals 18 %) and NGO substrates (Zirconia - NGO
lattice mismatch equals 25 %). Ion conductivity data are displayed in Fig. 8. 6 reveal that
films characterized by a thickness of 50 and 20 nm deposited on NGO exhibit a lower
conductivity as compared to the same films deposited on MgO. However, for the 10 nm thick
films, it is observed that the films deposited on NGO (25% lattice mismatch) exhibit a higher
ionic conductivity as compared to one deposited on MgO (18% lattice mismatch). This leads
to the conclusion that the strong enhancement of the ionic conductivity for the 10 nm thick
137
films is more marked by interfacial lattice strain. Above 10 nm, most probably, the lattice
strain relaxes and lead to more bulk (lattice) like diffusion.
Finally, the four-probe analysis performed on the as deposited samples at room temperature
revealed films do not exhibit any electronic conductivity at all i.e. are highly insulating (a
critical requirement for the electrolyte to be used in solid oxide fuel cells).
8.3 Conclusion
Ionic conductivity measurements of OVSZ thin films containing 16 and 3 at.% O vacancies
(1 at.% of O Vac = 3.2 mol.% of Yttrium in YSZ) have been carried out and reveal that OVSZ
thin films are good ion conductors. Ionic conductivity is in the range of reported values for
YSZ. Then increase of O vacancy concentration from 3 at% to 16 at% does not improve the
ion conductivity as first expected. The incorporation of larger amounts of vacancy leads to
lattice distortions which hinders ion mobility. However, a colossal increase in the ionic
conductivity is observed when the film thickness is decreased to 10 nm. This nanoscale effect
seems to be mainly originated from lattice strain, induced by depositing the OVSZ film onto
the right substrate (polished NdGaO3 (100) single crystals). Such an enhancement in the ionic
conductivity of the OVSZ thin films might offer new opportunities for ion conductors.
Moreover, it should be noted that the OVSZ films have a high resistivity as measured by
four-point probe method. This combined with the high ionic conductivity exhibited by OVSZ
makes this material a promising alternate of YSZ to be used in solid oxide fuel cells.
138
References [1] I. Kosacki, C.M. Rouleau, P.F. Becher, J. Bentley, D.H. Lowndes, Nanoscale effects on the ionic conductivity in
highly textured YSZ thin films, Solid State Ionics. 176 (2005) 1319–1326. doi:10.1016/j.ssi.2005.02.021.
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the systems Zr0.75Ce0.08M0.17O1.92 (M=Nd, Sm, Gd, Dy, Ho, Y, Er, Yb, Sc), Solid State Ionics. 149 (2002)
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139
9. Optical properties of OVSZ thin
films
Zirconia is a wide band gap material (bang gap energy = 5.1 - 6 eV[1–5] ) characterized by a
high refractive index (n = 2.1 - 2.2[6–8] ). These properties make zirconia, and zirconia –
based materials, useful candidates for high k dielectric applications, [9–11]. Most of the
zirconia properties are exploited by stabilizing its high temperature tetragonal and cubic
phases with the help of doping, e.g. with Y3+ cations, as previously shown (see chapter 2).
Beside the use of dopants, one can stabilize high temperature c-phase of zirconia at room
temperature by incorporating O vacancies in the zirconia lattice i.e. by growing under-
stoichiometric zirconia (ZrO2-x) thin films, as demonstrated in the chapter 5. In such oxygen
vacancy stabilized zirconia (OVSZ) films, not all the Zr-d electrons are accommodated with
O atoms. This situation leads to the promotion of Zr-labeled states into the bandgap of
zirconia. It is therefore expected that the optical properties of the materials might be modified,
such as in allowing new optical transitions in the visible range. In this last section, we present
the first results on the optical characterization of Oxygen Vacancy Stabilized cubic Zirconia
thin films. In particular, we study experimentally and theoretically how oxygen vacancy
incorporation influences the bandgap and thus optical properties of OVSZ thin films.
140
9.1 Experimental details
In order to investigate the optical properties of zirconia thin films, two sets of samples having
thickness 100±10 nm were prepared on Si (100) and on borosilicate glass substrates. Thin
film samples were synthesized at 10 mTorr and with a discharge current of 200 mA. Samples
were prepared by working inside the transition zone as well as out of it, using dc-RMS (details
can be found in chapter 4). Each of these sets is comprised of 3 samples, 2 of which are c-
ZrO2-x samples with oxygen vacancy concentration equal to either 3 at.% or 16 at.% and one
samples is pure ZrO2 (i.e. no vacancies incorporated). In this last case, as previously reported,
the X-ray diffractograms shows only reflection emanating from the monoclinic phase.
For the optical characterization of such deposited samples, spectrophotometry and
photoluminescence (PL) measurements were carried out at room temperature at the Institut
Jean Lamour, in Nancy (Dr. David Horwat and Dr. Hervé Rinnert). In case of
photoluminescence, a Xe lamp was used as excitation source and an edge filter was used as
to cut the excitation wavelength above 325 nm. During the characterization, excitation
wavelengths used are 270 nm (4.6 eV), 280 nm (4.4 eV), 290 nm (4.3 eV), 300 nm (4.1 eV)
and 310 nm (4 eV). Detection was made possible using a CCD camera and, in order to
compare the PL spectra, PL intensities are normalized to the excitation power.
Beside these experiments, data were extracted from the DFT calculations previously
performed to gain a better insight on the experimental results. Details of such calculation can
be found in Chapter 4.
9.2 Results and discussion
The impact of various amounts of O vacancy incorporated in the zirconia unit cell on the
bandgap, as derived from the DFT calculations, is shown in Fig. 9. 1. At first sight it can be
141
seen the unit cell with 0 O vacancy shows no state in the bandgap. While the DOS of 1, 2, 3
and 10 at.% O vacancy show that, as soon as we incorporate O vacancies in the unit cell, new
states appear in the bandgap. The number of states increases with the increase in O vacancy
concentration. It is also observed that the Fermi level shifts towards the conduction band with
the incorporation of O vacancies. These results underline that these states, located in the
bandgap, are occupied.
Fig. 9. 1: DOS of zirconia unit cells containing 0, 1, 2 3 and 10 at.% O vacancies. Dashed red line
shows the fermi level.
Further analysis of the DOS of the 3 at.% OVSZ film is shown in Fig. 9. 2. It reveals that
these band gap states are not coming from the O atoms but from the Zr atoms. Moreover, it
is also observed that the states appearing in the bandgap are localized near the O vacancy and
belong to Zr atoms (see Fig. 9. 2a, DOS of Selected Zr atoms located close to the O vacancy).
DFT calculations suggest that the formation of an isolated O vacancy in pure zirconia lead to
142
a situation where the O-p band can no longer hold all the Zr-d electrons. As a result, Zr states
accommodating the excess electron appear in the bandgap.
One thing regarding the calculated bandgap has to be noted; its energy is about 3.7 eV.
Experimental values reported for the zirconia bandgap range between 5.1 and 6 eV [1–5].
This discrepancy in the bandgap calculated theoretically and experimentally is well
known[12] and arises due to the PBE functional used in our calculations.
Fig. 9. 2: DOS of cubic zirconia containing 3 at.% O vacancy, calculated using DFT. DOS of (a) Zr
atoms located close to O vacancy, (b) all Zr atoms in the cell, (c) all O atoms, (d) of all atoms.
PL spectra of the deposited films containing 0, 3, and 16 at.% of oxygen vacancies are
reported in Fig. 9 .3 (a), (b) and (c), respectively. PL spectra of 0 at.% oxygen vacancy film
show no emission peak (the very small bump at ~320 nm is due to the parasitic light
emanating from the setup), while the spectra related to the films with 3 and 16 at.% oxygen
vacancy concentrations exhibit two emission peaks. The first peak is centered at 388 nm (3.2
eV) and the second is centered at 488 nm (2.5 eV). The presence of these two emission peaks
143
in the visible range for these two OVSZ films highlight that photoluminescence results from
O vacancy incorporation in the zirconia lattice.
The PL spectra of 3 at.% O vacancy film shows that at low excitation energy (4 eV), emission
intensity of 388 peak is high while the intensity of the 488 nm peak is low. Further, an
evolution of reverse trend is observed with the increase in excitation energy for these two
peaks (Fig. 9. 3b). Since the excitation energy in the present study is smaller than the zirconia
bandgap, therefore, there is no possibility to excite electrons from the valence band (VB)
directly to the conduction band (CB). However, as we discussed, DFT calculations show that
the incorporation of O vacancies leads to the formation of energy states in the bandgap to
accommodate the excess electrons (Fig. 9. 1. Electron from those new states could be excited
to the levels above the CB, leaving holes behind. Those excited electrons decay non-
radiatively to the CB from where they decay back into the holes and lead to the emission of
388 nm (3.2) and 488 nm (2.5) peaks, Fig. 9. 4. However, the DOS of the film doped with 3
at.% O vacancy (Fig. 9. 1) shows that the new states appearing in the bandgap lie at energies
of 1.8, 1.4, and 1.2 eV from the CB. These energies do not correspond to the observed
fluorescence signal wavelengths. As it was discussed earlier, the bandgap calculated here,
using PBE functional, is lower than the reported experimental bandgap energies. To correct
such calculated 3.7 eV bandgap, one way is to lift the CB to the experimental value i.e. to 5.1
eV. By doing so, new values corresponding to the states appearing in the bandgap from CB
become 3.2, 2.8, and 2.6 eV. After this adjustment, one can find the 3.2 and 2.6 eV energy
values. These values fit well with the emission wavelengths hence confirming that the origin
of fluorescence lie in these states appearing in the bandgap as a result of incorporation of O
vacancies. A tentative schematic representation of electron excitation upon light absorption
during the PL experiments, and involving the energy states located in the band
144
Fig. 9. 3: PL spectra of films containing 0 (a), 3 (b) and 16 (c) at.% O vacancy.
145
Fig. 9. 4: Schematic representation of the possible transitions.
gap, is presented in Fig. 9. 4. Further, the decrease in the emission intensity of 388 nm peak
and increase in the emission intensity of 488 nm peak with the increase in excitation energy
for 3 at.% O vacancy film suggest that most probably there is a transfer of energy between
these two peaks.
Fig. 9. 3c shows the PL spectra of the 16 at.% oxygen vacancy film. Those spectra reveal that
the intensity of emission peak at 488 nm increases with the increase in excitation energy,
while the intensity of peak at 388 nm is very weak and excitation energy has no influence on
it. The origin of both peaks is the same i.e. the states appearing in the bandgap as a result of
O vacancy incorporation, as discussed earlier. However the weaker intensity of 388 nm peak
in case of 16 at.% O vacancy and as compared to 3 at.% O vacancy film is probably due to
the quenching of the PL phenomenon by the larger amount of O vacancies present in the
zirconia lattice. It is known for O vacancies that they can trap the excitation energy through
non-radiative transitions and quench the emission [13]. We believe that it’s the reason why
we observe a decrease in the intensity of the peak at 388nm with the increase in O vacancy
concentration.
146
Fig. 9. 5: Transmittance spectra of films containing 16, 3 and 0 at.% oxygen vacancy.
The transmittance spectra of films containing 0, 3 and 16 at.% O vacancy films are shown in
Fig. 9. 5. It is observed, below 270 nm, the samples do not transmit any light due to the glass
substrate which is not transparent to the UV. While in the visible and IR range, films show a
high transmittance above 80 %. Spectra also show small bumps around 900 nm which
correspond to the change of grating and detector. The other small bump around 1050 nm,
probably also coming from the glass substrate. In general, the transmittance data show that
films are transparent in visible-IR range.
9.3 Conclusion
In conclusion, DFT calculations show that incorporation of O vacancy in the zirconia lattice
lead to the formation of occupied energy states in the bandgap. Photoluminescence data
reveal no emission peak in case of 0 at.% O vacancy film while two emission peaks centered
at 388 nm and 488 nm in films containing 3 and 16 at.% O vacancies. This suggests the origin
of 388 and 488 nm peak is the formation of new states in the band gap as a result of O vacancy
147
incorporation. However, having large amount of O vacancies (16 at.%) concentration in the
zirconia lattice leads to the quenching of 388 nm peak. Transmittance data show that the
deposited films are transparent in visible-IR range. Optical data presented here is the very
preliminary data on OVSZ and may subject to more thorough investigation.
148
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149
Outlook
Tailoring the crystal structure of any material is an important parameter for controlling its
properties. In this regard, for several decades, various approaches and deposition methods
have been used to understand and control the phase formation of zirconia bulk materials and
thin film coatings; in particular, to stabilize zirconia high temperature cubic phase at room
temperature. For zirconia, phase transition as a function of temperature can lead to
mechanical failure of the material, which in turn, prevent its use in engineering applications.
In the chapters 4 and 5 of this doctoral thesis work, quantum-chemistry based calculations
and optimized reactive magnetron sputter deposition method have been combined to assess
the influence of oxygen vacancy incorporation on the phase formation of zirconia thin films.
According to the calculations, performed at the density function theory level, and their
remarkable agreement with the experimental thin film characterization data, it is concluded
that incorporating O vacancies (at least 3 at.%) in the zirconia lattice is the sole mechanism
responsible for stabilizing the zirconia high-temperature c-phase, at room temperature. The
O vacancy containing zirconia thin films were grown experimentally by dc reactive
magnetron sputtering with a tight control over the oxygen flow (the discharge current was
200 mA and total pressure 10 mTorr). The latter was enabled by implementing a feedback
control loop which links the plasma and discharge parameters to the reactive gas flow, and
to the target oxidation state. This first result highlights the importance of materials defect
150
chemistry during film growth by reactive magnetron sputtering. However, in chapter 6 it is
demonstrated, role of materials defect chemistry is only important and govern the zirconia
phase constitution only when the films are grown in moderate conditions i.e. conditions
providing global energy flux < 1000 eV/Zr atom. Above this energy, the role of materials
defect chemistry is apparently partially hindered and the monoclinic phase formations occurs
and thus, a mixed c- and m-phase is obtained. The increased global energy flux per Zr atom
leading to the build of large compressive stresses in the film is proposed to be the responsible
for such behavior. Beside controlling the materials defect chemistry and deposition
conditions, it is also shown in chapter 6 that monoclinic crystals can also appear when thicker
films are deposited. TEM analysis allowed to detect m-ZrO2 crystals 850 nm above the film-
substrate interface. This shows one also should pay attention to the film thickness when
growing thicker films even if the film deposition conditions are kept constant during the
whole deposition process. In chapter 7, the thermal stability analysis performed on pure c-
phase oxygen vacancy stabilized zirconia (OVSZ) thin films (i.e. 100 nm thin films grown at
200 mA, 10 mTorr) showed that these films are stable to rather high temperatures approx.
750 °C. The appearance of monoclinic diffraction peak after passing 750 °C is attributed to
two possible mechanisms; i) Oxygen incorporation and/or ii) thermal stresses.
Having now the understanding and the means to incorporate oxygen vacancies and to
stabilize the high temperature cubic phase of zirconia, one can further characterize this
material. In this regard, in chapter 8, lateral ionic conductivity measurements are presented.
The results reveal that the OVSZ films are ion conductors, exhibiting conductivity values
similar to those recorded for Yttria Stabilized Zirconia coatings. Furthermore, an ionic
conductivity as high as 7.4 S.cm-1 is obtained at 725 °C for 10 nm-thick OVSZ films
containing 3 at.% of vacancies. Such colossal enhancement in the ionic conductivity is
151
attributed to the lattice strain caused by film-substrate lattice mismatch resulting in higher
disorder at the film-substrate interface. Such epitaxial strain could be a key step in the design
of electrolyte exhibiting high ionic conductivity. Four-probe and Hall effect measurement
showed the OVSZ are highly resistive and exhibit no electronic conductivity at room
temperature. The high ionic conductivity and high electronic resistivity exhibited by OVSZ
thin films makes this material a promising alternate of YSZ to be used e.g. in solid oxide fuel
cells (SOFC) and in oxygen sensors. Further, photoluminescence analysis of OVSZ (chapter
9) revealed the presence of 2 emission peaks centered at 388 nm and 488 nm. This result
validates the theoretical calculations data predicting the appearance of energy states in the
band gap upon O vacancy incorporation. However, the presented photoluminescence data is
preliminary and a thorough investigation is needed to figure out the detailed mechanism and
optical transitions involved in these transitions.
As it is mentioned above, the role of materials defect chemistry is only important and govern
the zirconia phase constitution only when the films are grown in moderate conditions i.e.
conditions providing global energy flux < 1000 eV/Zr atom. Since the global energy flux per
Zr atom is the sum of energies brought to the substrate by all plasma species (Ar+, O-, e-, hv,
IR, atoms, etc.), therefore, it would be interesting to develop a strategy to unambiguously
pinpoint the influence of each energetic plasma specie on the zirconia phase constitution. One
way to achieve this could be the use of mass spectrometry in the current deposition conditions
to get the ion energy distribution function (IEDF) to find the energy of each ionized plasma
specie (Ar+, O-) and relate them with the zirconia phase formation. An ion source to see the
influence of energetic ions on the phase formation of growing film could also be used in this
regard. In similar way, an IR source could also be used separately to inquire the role of IR on
film growth and phase formation.
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Further, as the increase in film stress as a result of increased global energy flux per Zr atom
is proposed to explain the appearance of monoclinic phase, therefore, it would be interesting
to perform the stress analysis on such deposited films to calculate the critical stress value
leading to the formation of monoclinic phase in these conditions. Moreover, as it is observed
that the increase in film thickness leads to the appearance of zirconia m-phase even if the
films contain 16 at.% O vacancies. Therefore, it would be interesting to figure out the
mechanism responsible for such behavior. One way to understand this mechanism could be
the in-situ stress analysis to assess the stress evolution as the film grows.
Thermal stability analysis of the as deposited cubic zirconia thin films (100 nm thin films
grown at 200 mA, 10 mTorr) showed films are stable up-to 750 °C and after passing this
critical temperature (750 °C) m-phase start to appear. Appearance of m-phase has been
attributed to two possible mechanisms; i) Oxygen incorporation and/or ii) thermal stresses
To unambiguously separate these two mechanisms a more detailed study is needed. This can
be achieved e.g. by measuring the film chemistry before and after the annealing and also by
depositing films on various kinds of substrates characterized with different thermal expansion
coefficients. Further stress analysis should also be performed at room temperature, at critical
temperature (750 °C) and after passing critical temperature to unambiguously identify the
role of thermal stresses.
Besides, understanding of the fundamental mechanisms behind the stabilization of the high
temperature c-phase of zirconia could also pave the way to control the phase constitution of
other metal oxides such as e.g., Bi2O3 in order to promote their high temperature phase to
room temperature. Having stabilized high temperature cubic phase of Bi2O3 (δ-Bi2O3)
without any dopants might also enhance its ionic conductivity. Further, growing under-
stoichiometric CeO2 using the developed method could also help to avoid the use of reduction
153
methods currently implemented to create O vacancies in this material, to enhance its ionic
conductivity.
154
Appendix
Total energy of Zirconia polymorphs as a function of O vacancy, calculated using ab initio
DFT.
Table AI. 1. O vacancies introduced Randomly
No. of Zr
atoms
No. of O
atoms
No. of O
vacancies
Monoclinic
total energy
(eV)
Tetragonal
total energy
(eV)
Cubic total
energy (eV)
32 64 0 -30701.28469 -30701.37687 -30701.13491
32 63 1 -30258.75535 -30259.90955 -30259.28631
32 62 2 -29816.28794 -29820.36984 -29819.04874
32 61 3 -29373.56172 -29378.72436 -29379.91604
32 60 4 -28931.20262 -28936.86533 -28939.41483
32 59 5 -28488.02695 -28495.66794 -28499.19016
32 58 6 -28048.12734 -28057.42449 -28057.97726
32 57 7 -27605.83513 -27614.52185 -27616.83311
32 56 8 -27165.23928 -27174.19326 -27178.35916
32 55 9 -26723.71548 -26735.47371 -26738.11302
32 54 10 -26282.10293 -26291.39565 -26296.25087
Table AI. 2. O vacancies introduced as cluster
No. of Zr
atoms
No. of O
atoms
No. of O
vacancies
Monoclinic
total energy
(eV)
Tetragonal
total energy
(eV)
Cubic total
energy (eV)
32 64 0 -30701.28469 -30701.37687 -30701.13491
32 63 1 -30258.75535 -30259.90955 -30259.28631
32 62 2 -29815.90567 -29820.40569 -29819.03379
32 61 3 -29373.59236 -29379.98953 -29380.09725
32 60 4 -28931.62045 -28937.82173 -28934.86731
32 59 5 -28488.09909 -28498.20545 -28497.62375
32 58 6 -28043.42514 -28057.31826 -28051.31602
32 57 7 -27605.22585 -27616.22044 -27616.79238
155
32 56 8 -27160.72825 -27171.63822 -27165.15831
32 55 9 -26721.98725 -26729.82084 -26736.20141
32 54 10 -26284.63429 -26288.19487 -26294.91857
Table AI. 3. O vacancies introduced apart
No. of Zr
atoms
No. of O
atoms
No. of O
vacancies
Monoclinic
total energy
(eV)
Tetragonal
total energy
(eV)
Cubic total
energy (eV)
32 64 0 -30701.28469 -30701.37687 -30701.13491
32 63 1 -30258.75535 -30259.90955 -30259.28631
32 62 2 -29815.67681 -29819.98039 -29817.53961
32 61 3 -29376.05029 -29378.06347 -29377.39886
32 60 4 -28933.26141 -28934.9256 -28937.66177
32 59 5 -28491.0859 -28496.88414 -28499.94733
32 58 6 -28047.69346 -28055.93488 -28059.69634
32 57 7 -27605.62941 -27612.72958 -27617.79826
32 56 8 -27163.86827 -27174.49147 -27178.58252
32 55 9 -26722.18795 -26730.96115 -26734.71378
32 54 10 -26285.48753 -26292.00887 -26295.94746
156
Fig. AI. 1: Influence of O vacancies on zirconia phase constitution. The energies of each phase are
compared as a function of the concentration of oxygen vacancies. Three ways to distribute the
vacancies in the lattice are presented: (a) random, (b) apart from each others, and (c) clustered.
157
158
159
160