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TRANSCRIPT
My professional path within science started with a high school assignment about Marie Curie.
Other important women for my scientific journey have been my high school chemistry teacher
Kerstin Göras and my master thesis supervisor Dr. Karin Persson who have been a great
influence on me. Therefore, I would like to dedicate this work to these women.
“You cannot hope to build a better world without improving the individuals. To that end, each
of us must work for his own improvement and, at the same time, share a general responsibility
for all humanity, our particular duty being to aid those to whom we think we can be most
useful.”
-Marie Curie
“I tried out various experiments described in treatises on physics and chemistry, and the results
were sometimes unexpected. At times, I would be encouraged by a little unhoped-for success;
at others, I would be in the deepest despair because of accidents and failures resulting from my
inexperience.”
-Marie Curie
1
Abstract Super duplex stainless steel consists of two phases: austenite and ferrite, and is a highly
corrosion resistant material, with a wide range of applications. The corrosion resistance of super
duplex stainless steel is determined by the performance of the passive oxide film formed
spontaneously on the surface. This is a 1-3 nm thick and Cr oxide rich film, which protects the
material from further oxidation and the steel is said to be in the passive state. The passive state
can be investigated by electrochemical polarization methods and surface analysis techniques,
which make it possible to detect the formation and degradation of the passive film. This thesis
work contains two parts: in-situ/operando synchrotron measurements to study the degradation
mechanism during anodic polarization; and ex-situ measurements to map the lateral thickness
and microscopic elemental distribution in the passive film.
The in-situ/operando synchrotron experiments combined several experimental techniques,
including X-ray reflectivity (XRR), X-ray diffraction (XRD), X-ray fluorescence (XRF) and
electrochemical impedance spectroscopy (EIS), to simultaneously characterize the structure
and chemical/electrochemical properties, as well as their changes in the surface region of the
samples. During anodic polarization, the passive film thickness, density, roughness, crystalline
structure, and the electrochemical properties of the film were measured at each stepwise
increased polarization potential. Furthermore, the dissolved metal elements were probed by
XRF from the electrolyte above the sample surface. It was found that the oxide film became
more defective with increasing potential, leading to a decreasing density of the passive film.
On the other hand, the Ni rich alloy surface layer below the oxide film showed an increasing
density, indicating an increased concentration of heavy elements (Mo and Ni). Fe was the first
and main element detected, and the significantly enhanced metal dissolution above 1000 mV vs
Ag/AgCl indicates that the material entered the so called transpassive state. The XRD data showed
evidence of nanocrystalline Cr and Fe oxidic components in the passive film, whereas the
amorphousness of the passive film increased with increasing potential. Moreover, the surface
strain induced by mechanical grinding was found to affect the crystalline nature, making the
film more amorphous. In short, the passivity breakdown is a continuous degradation process of
the passive film over a potential range, involving structural and compositional changes of the
passive film and the underlying alloy surface layer associated with enhanced Fe dissolution
before rapid Cr dissolution (at ≥1300 mV vs Ag/AgCl).
The ex-situ investigations employed hard X-ray photoemission electron microscopy
(HAXPEEM), providing X-ray photoelectron spectroscopy (XPS) data from individual grains
with crystallographic orientations of (111), (101) and (001), parallel to the sample plane. The
experimental approach enabled the analysis of the same sample area before and after
polarization. The XPS data was used to evaluate the thickness and Cr content of the native
passive film on individual grains, with particular grain orientations, of the ferrite and austenite
phases, and of all the analyzed grains, respectively. The results reveal lateral variations in the
native passive film between the two phases and among the three grain orientations. The Cr
content was higher on the ferrite than the austenite, whereas the thickness was rather uniform.
The grain orientation has a small but detectable influence on the thickness and Cr content of
the native passive film. For example, Ferrite (111) grains had a lower Cr content in the outer
layer of the passive film than the other ferrite grains.
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Populärvetenskaplig sammanfattning En av de största anledningarna till att rostfria stål är så pass korrosions resistenta är på grund
av en spontan oxidskiktsbildning på deras yta. Denna metalloxid består i huvudsak av Cr och
Fe oxid och hydroxid men även utav Mo föreningar. Oxidskiktet är mycket tunt, mellan 1-3 nm
tjockt och det bildas och ändras ständigt. Dess dynamiska karaktär i samband med att det är så
tunt gör det svårt att definiera dess tjocklek, sammansättning och nedbrytningsprocess. Den
passiva filmen består av två lager, ett oxidskikt och ett hydroxidskikt. På ett mycket
rostbeständigt stål så som det höglegerade och tvåfasiga super duplex rostfritt stål 2507, visas
inga tecken på lokal nedbrytning under normala förhållanden. Dock så visar dess
polarisationskurva att den går från passiv till upplösning i det så kallade transpassiva området,
där den passiva filmen samt materialet upplöses. Hur nedbrytningen sker på en atomär nivå och
hur duplex stålets uppbyggnad samt sammansättning påverkar den passiva filmens nedbrytning
eller avstående från nedbrytning är inte helt förstått.
I denna avhandling har den passiva filmens nedbrytning följts utifrån flera aspekter så som den
passiva filmens tjocklek, densitet och struktur samt genom att följa den kemiska avfällingen
från ytan och därmed kunnat bättre beskriva nedbrytningsprocessen. Detta möjliggjordes
genom att förskjuta potentialen eller den elektromotoriska kraften i riktning mot oxidering
genom polarisation av materialet.
Den passiva filmen kunde också beskrivas på lokal mikroskopisk nivå genom fotonemission
elektronmikroskopi (HAXPEEM) analys, som gör det möjligt att korrelera kemiska spektra från
mikroskopiska bilder. De mikroskopiska bilderna är uppbyggda från elektronspektra som kan
härledas till specifika bindningsenergier för undersökta element så som Cr, Fe, Ni och O.
Tekniken gör det möjligt att välja ett specifikt område så som en fas nivå eller enbart en
kornorientering och undersöka den passiva filmen på lokal nivå.
Resultaten av dessa mätningar visade att ett oxidskikt finns på ytan upp till höga potentialer på
1300 mVvs Ag/AgCl. Efter denna potential är nedbrytningen snabb och inget skyddande oxidskikt
finns på ytan. Med ökande potential blir den passiva filmen mer och mer defekt vilket minskar
dess barriäregenskaper och dess passivitet minskar. Under oxidskiktet finns ett skikt med högre
densitet med ungefär samma tjocklek. Densiteten av detta lager ökar vilket indikerar en
berikning av Ni och Mo. Denna undersökning visade att nedbrytningen av den passiva filmen
är en process som sker under ett potentialförlopp.
De lokala resultaten från kornen visade att det förmodligen finns en liten skillnad mellan den
passiva filmen på faserna ferrit och austenit och dess tillhörande korn. Sammansättningen skilde
sig något då ferrit som hade ett högre innehåll av Cr i dess passiva film än austenit. Detta är
rimligt då ferritens kemiska sammansättning består av mer Cr än vad austenit gör.
Tjockleksmässigt är skillnaden på oxidfilmen mellan de två fasernas försumbart liten. Dock så
verkar det som att en viss skillnad kan finnas mellan oxidskiktets två lager på korn-nivå. Ferrit
(111) hade en lägre Cr koncentration än övriga ferritiska kornorienteringar.
3
List of publications included in the thesis Paper I: In-situ synchrotron GIXRD study of passive film evolution on duplex stainless
steel in corrosive environment
Cem Örnek, Marie Långberg, Jonas Evertsson, Gary Harlow, Weronica Linpé, Lisa Rullik,
Francesco Carlà, Roberto Felici, Eleonora Bettini, Ulf Kivisäkk, Edvin Lundgren, and
Jinshan Pan
Corrosion Science (2018), 141, p. 18-21
Paper II: Redefining passivity breakdown of super duplex stainless steel by
electrochemical operando synchrotron near surface X-rays
Marie Långberg, Cem Örnek, Jonas Evertsson, Gary S. Harlow, Weronica Linpé, Lisa Rullik,
Francesco Carlà, Roberto Felici, Eleonora Bettini, Ulf Kivisäkk, Edvin Lundgren, and Jinshan
Pan
NPJ Material degradation (2019), 3
Paper III: Influence of Surface Strain on Passive Film Formation of Duplex Stainless Steel
and Its Degradation in Corrosive Environment
Cem Örnek, Marie Långberg, Jonas Evertsson, Gary Harlow, Weronica Linpé,
Lisa Rullik, Francesco Carlà, Roberto Felici, Ulf Kivisäkk, Edvin Lundgren,
and Jinshan Pan
Journal of the Electrochemical Society (2019), 166, p. 3071-3080
Paper IV: Characterization of Native Oxide and Passive Film on Austenite/Ferrite Phases
of Duplex Stainless Steel Using Synchrotron HAXPEEM
Marie Långberg, Cem Örnek, Fan Zhang, Jie Cheng, Min Liu, Elin Grånäs, Carsten Wiemann,
Andreii Gloskovskii, Yury Matveyev, Satishkumar Kulkarni, Heshmat Noei, Thomas. F.
Keller, David Lindell, Ulf Kivisäkk, Edvin Lundgren, Andreas Stierle, and Jinshan Pan
Journal of the Electrochemical Society (2019), 166, p. C3336-C3340
Paper V: Lateral Variation of Native Passive Film on Super Duplex Stainless Steel Resolved by
Synchrotron Hard X-Ray Photoelectron Emission Microscopy
Marie Långberg, Cem Örnek, Fan Zhang, Jie Cheng, Min Liu, Elin Grånäs, Carsten Wiemann,
Andreii Gloskovskii, Yury Matveyev, Satishkumar Kulkarni, Heshmat Noei, Thomas. F.
Keller, Christoph Schlueter, David Lindell, Ulf Kivisäkk, Edvin Lundgren, Andreas Stierle,
and Jinshan Pan
Corrosion Science (2020), 174, 108841
4
Contribution Influence of Surface Strain on Passive Film Formation of Duplex Stainless Steel and Its
Degradation in Corrosive Environment
Contributed to sample preparation, experiment execution, data analysis, scientific discussion,
and manuscript preparation.
Redefining passivity breakdown of super duplex stainless steel by electrochemical
operando synchrotron near surface X-rays
Contributed significantly to sample preparation, experiment execution, data analysis including
a major part of XRF, XRR, EIS data, and major contribution to manuscript preparation.
In-situ synchrotron GIXRD study of passive film evolution on duplex stainless steel in
corrosive environment
Contributed to sample preparation, experiment execution, data analysis and evaluation of the
data especially of current transient and XRF data, scientific discussion, and manuscript
preparation.
Characterization of Native Oxide and Passive Film on Austenite/Ferrite Phases of Duplex
Stainless Steel Using Synchrotron HAXPEEM
Actively participated in sample preparation and experiment execution. Main contributor to data
analysis and manuscript preparation.
Lateral Variation of Native Passive Film on Super Duplex Stainless Steel Resolved by Synchrotron
Hard X-Ray Photoelectron Emission Microscopy
Led sample preparation and actively participated in the experiment execution, main responsible
for data analysis and manuscript preparation.
Work not included in the thesis:
Integration of electrochemical and synchrotron-based X-ray techniques for in-situ
investigation of aluminum anodization
Fan Zhang, Jonas Evertsson, Florian Bertram, Lisa Rullik, Francesco Carlà, Marie Långberg,
Edvin Lundgren, Jinshan Pan
Electrochemica Acta (2017), 241, p. 299-308
Operando time- and space-resolved high energy X-ray diffraction measurement to
understand hydrogen-microstructure interactions in duplex stainless steel
Cem Örnek, Timo Müller, Ulf Kivisäkk, Fan Zhang, Marie Långberg, Ulrich Lienert, Ki-Hwan
Hwang, Edvin Lundgren, Jinshan Pan
Corrosion Science (2020), 175, 108899
5
List of abbreviations
AC – Alternating current
AES – Auger electron spectroscopy
BCC – Body centered cubic
DC – Direct current
CCD – Charge coupled device
EBSD – Electron backscatter diffraction
EIS – Electrochemical impedance spectroscopy
FCC – Face centered cubic
FIB – Focused ion beam
FOV – Field of view
HAXPEEM – Hard X-ray photoemission electron microscopy
OCP – Open circuit potential
PEEK – Polyether ether ketone
PEEM – Photoemission electron microscopy
PREN – Pitting resistant equivalent number
PTFE – Polytetrafluoroethylene
ROI – Region of interest
SEM – Scanning electron microscopy
TEM – Transmission electron spectroscopy
UHV – Ultra high vacuum
UV – Ultraviolet
XRD – X-ray diffraction
XRR – X-ray reflectivity
XRF – X-ray fluorescence
XPS – X-ray photoelectron spectroscopy
6
Contents
Abstract ...................................................................................................................................... 1
Populärvetenskaplig sammanfattning ........................................................................................ 2
List of publications included in the thesis .................................................................................. 3
Contribution ............................................................................................................................... 4
List of abbreviations ................................................................................................................... 5
1. Introduction ........................................................................................................................ 8
2. Material ............................................................................................................................ 11
2.1 Duplex Stainless steel ................................................................................................ 11
2.2 Role of alloying elements .......................................................................................... 13
3. Passivity and breakdown of stainless steel ....................................................................... 15
3.1 The passive film and passivity ................................................................................... 15
3.2 Passivity breakdown and passive film degradation ................................................... 16
4. Experimental techniques .................................................................................................. 18
4.1 Electrochemical techniques ....................................................................................... 18
4.1.1 Direct current techniques ................................................................................... 19
4.1.2 Electrochemical Impedance Spectroscopy (EIS) ............................................... 20
4.2 Scanning Electron Microscopy / Electron Backscatter Diffraction ........................... 23
4.3 Synchrotron based techniques ................................................................................... 25
4.3.1 X-Ray Diffraction (XRD) .................................................................................. 26
4.3.1.1 XRD Principle ...................................................................................................... 26
4.3.1.2 Surface sensitivity of XRD .................................................................................. 27
4.3.1.3 Measurement of strain by XRD ........................................................................... 28
4.3.2 X-Ray Reflectivity (XRR) ................................................................................. 29
4.3.3 X-Ray Photoelectron Spectroscopy (XPS) ........................................................ 30
4.3.3.1 XPS Principle ....................................................................................................... 30
4.3.3.2 Surface sensitivity of XPS.................................................................................... 31
4.3.3.3 Chemical shifts and spectra fitting ....................................................................... 31
4.3.3.4 XPS data analysis ................................................................................................. 32
4.3.4 Photoemission Electron Microscopy (PEEM) ................................................... 34
4.3.5 X-ray Fluorescence (XRF) ................................................................................. 35
5. Results & Discussion ....................................................................................................... 36
5.1 Passive film formation, stability, and degradation .................................................... 36
5.1.1 In-situ/operando experimental setup ....................................................................... 36
7
5.1.2 Summary of results from in-situ/operando measurements ..................................... 37
5.1.2.1 Thickness and density of the surface layers ......................................................... 38
5.1.2.2 Structural changes during anodic polarization ..................................................... 39
5.1.2.3 Dissolution and dealloying ................................................................................... 41
5.1.3 Summary of in-situ/operando experiment results ................................................... 44
5.2 Local chemical composition and thickness of passive film ...................................... 45
5.2.1 HAXPEEM experimental setup .............................................................................. 45
5.2.2 Thickness and Cr content of passive film and lateral variations ........................ 47
5.2.2.1 Composition ratios of single grains ...................................................................... 47
5.2.2.2 Lateral variation in thickness and Cr content of native passive film ................... 49
6. Conclusions ...................................................................................................................... 54
7. Outlook and future work .................................................................................................. 55
Acknowledgements .................................................................................................................. 56
References ................................................................................................................................ 57
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1. Introduction Protective outer surfaces are a natural occurring phenomenon that inhibit a system’s rapid
interaction with surrounding elements. Trees have bark, humans have skin and metallic systems
have a protective surface oxyhydroxide called the passive film. Many metals form an oxide film
spontaneously on their surface through oxidation of surface atoms. However, only films with
barrier properties that significantly decrease the oxidation/dissolution rate of the metal are
called a passive film. Passive films are dynamic surface oxide layers that are forming and
dissolving simultaneously. Their formation and dissolution rate are dependent on several
factors related to the material and its environment. The stability and performance of a material
is often intricately connected to its tendency to effectively form a passive film. The obtained
performance, structure, and thickness of the film depend on the properties of the underlying
material and its environment. Regulating factors are environmental immersion time, pH,
temperature, aggressive ion concentration, potential and the material composition [1-3].
The phenomenon of passivity was first recognized around 200 years ago, when the dissolution
behavior of Fe immersed in a concentrated and a diluted of nitric acid solution was described.
The two solutions had different oxidizing strengths on the Fe metal. While Fe was immersed in
the concentrated solution, the metal showed signs of drastic dissolution at first indicated by
bubble formation from the metal surface. However, the dissolution was reduced after some
time, which much later was found to be due to a formation of an iron oxide covering the surface.
In the dilute solution, the oxidation was weaker and the metal continued to dissolve because no
effective protecting film was able to form [4-6]. However, even though the surface oxide layers
are not protective, does not exclude a formation of oxide layers.
Oxide films appear within milliseconds when metal resides in air under ambient conditions.
However, the oxidation rate of metals under ambient conditions is relatively slow compared to
immersed metal surfaces, leading to an oxide growth over a few days. Native oxide films grow
because of an electric field that establishes through the oxide film. The growth is thereby
possible by an electron tunneling effect and ion migration facilitated by the electric field. The
strength of the electric field decreases with the film thickness until a critical thickness is
established and the oxide growth ceases to occur [1].
Resistive properties of a protective surface film on a metal are easy to determine when the
metal is immersed in an electrolyte solution. This can be measured by electrochemical
techniques, which detect ion and electron transport through the metal/electrolyte interphase.
Passive metals have a specific pattern during anodic polarization, where the current deviates,
from high, in the active state, to low, in the passive state, with an increasing applied potential.
A decrease in current means that the electron flow between working electrode, the sample, and
the counter electrode becomes limited by a stable oxide film on the sample surface [1, 7, 8].
There are three states during anodic polarization, which are illustrated schematically in Figure
1. In the potential range a), the metal is in the active state, where the metal dissolves actively in
the form of metallic ions, releasing electrons, and the current is increasing with the applied
potential. The passive state refers to the potential range b) in Figure 1, where formation of a
stable oxide takes place, and the current density decreases with increasing potential initially,
and becomes relatively stable for a considerable range of potentials. The third state, in the
9
potential range c), is the so called transpassive state, when the current density increases
drastically with potential, pushing the metal oxides towards higher oxidation states. The metal
oxides become unstable in the transpassive region because their solubility increases. Moreover,
the drastic current increase is furthermore induced, at least partially, by oxygen evolution that
occurs in this region [5, 6, 9].
Figure 1: Schematic illustration of the anodic polarization curve of a passive metal showing
three different states: a) an active state, with a high current and rapid dissolution; b) a passive
state, with a lower current due to a formation of a stable oxide film; and c) a transpassive state,
where further oxidation of the oxidic metal species occurs, leading to a more soluble and
unstable oxide film, as well as oxygen evolution. This figure is adapted from ref [10].
There have been many studies reported regarding passive films on a wide variety of metals and
environments. To achieve specific desirable properties, it has been shown that several metallic
elements can be added in alloys, such as Fe-based, Ni-based and Co-based alloys. Corrosion
resistance is determined by the spontaneous formation of the passive film, which remains stable
in a large variety of environments. One material with high corrosion resistance performance is
stainless steel, which is a group of Fe-based alloys with the common factor of at least 13 wt%
of Cr of its bulk composition. Cr is the key element for the passivity of stainless steels, where
the lower limit of 13 wt% is necessary to create the onset of passivation by forming a stable
surface film enriched by Cr3+ oxide and hydroxide. Through effective alloying, it has also been
possible to further improve the corrosion resistance by the development of highly alloyed
stainless steels with a Cr content up to 20-30 wt%. As an example, duplex stainless steels, with
both high corrosion resistance and high mechanical strength, have become some of the most
used industrial steels [11].
The improvement of the corrosion resistance of stainless steel has become possible through
a deep understanding of the role of alloying elements and the influence of the microstructure
on the passive properties. Due to their tremendous industrial importance, passive films,
10
especially the ones formed on stainless steels, have been extensively investigated. More
recently, the passivity of duplex stainless steel has also received considerable attention due to
its high corrosion resistance and multiphase system. However, the complex and diminutive
nature of its passive films, sustaining formation and dissolution reactions through a thickness
of 1-3 nm makes it challenging to study. To this end, several studies have been conducted using
surface sensitive techniques, including X-ray photoelectron spectroscopy XPS [12-14] and
Auger electron spectroscopy AES [15-17], to determine the chemical composition and
thickness of the passive films on duplex stainless steels. During the recent few years high
resolution transmission electron microscope TEM [18, 19] has also been used for studying
passive films in detail. Such ex-situ techniques have been very useful for determining the
thickness and composition of the passive films; however, they often demand ultra-high vacuum
for the measurement and using ideal and ultra clean samples. Many of these techniques have
been combined with electrochemical techniques, such as impedance and polarization [20, 21],
in the study of passivity and passive films. More advanced techniques, such as the ones
available at synchrotron radiation facilities, provides the possibility to investigate samples in
situ (i.e. in solution environments), and with a high energy beam to unveil physical and
chemical properties of passive films as well as mechanisms of passivity and its breakdown.
With such capabilities, several important questions can be addressed, for instance how the alloy
microstructure, alloying elements and which critical conditions effect the passive properties of
duplex stainless steel [22]. The overall purpose for studying the passive film of duplex stainless
steel is to increase the knowledge of the protective surface layers of this type of alloy, and how
the environment influences the passive film on an atomic level, which can facilitate the
development of high-performance corrosion resistant alloys. A more resistant alloy will lead to
less consumption of materials and less repairs and thereby contribute to a more sustainable
society following the UN global goals.
The specific aims of this thesis were to:
i) define the critical conditions that lead to passivity breakdown.
ii) gain detailed knowledge of the composition and thickness of the passive film of
the austenite and ferrite phases. As well as, if and how the passive film can vary
between the two phases.
iii) derive the atomic changes that leads to breakdown of the passive films.
iv) determine how the microstructure influences the passivity breakdown.
This work has been a multi-collaboration project founded mainly by Vetenskapsrådet and
partly by Swerim. The duplex stainless steel used in the studies was SAF2507, supplied by AB
Sandvik Materials Technology. Material preparation and post analyses were mostly done at
Swerim. The primary techniques used were synchrotron based XRR, XRD, XRF and EIS and
the in-situ/operando synchrotron measurements were performed at beamline I03 at ESRF in
Grenoble in collaboration with Lund University. The HAXPEEM measurement was performed
at beamline P22 Petra III DESY in Hamburg, in collaboration with Forschungszentrum Jülich
and DESY Nanolab. This thesis includes five published papers (see the list of the publications),
based on the results from the synchrotron measurements.
11
2. Material
2.1 Duplex Stainless steel In this thesis, the focus is on super duplex stainless steel. Super duplex stainless steel has a high
resistance towards local corrosion compared to other duplex stainless steel grades. This type of
steel is commonly used within nuclear, petroleum and medical applications, which involve
corrosive environments and high safety requirements. The use of this type of steel is also
favored for economic reasons. There are several materials, such as Ni-based steels, that have
the same or even better corrosion performance, but they are much more expensive. The dual
phase microstructure of duplex stainless steel, with a relatively low Ni content, has made it
possible to decrease the Ni content in the bulk, but still achieve high mechanical strength and
corrosion resistance. The microstructure consists of a body-centered cubic (BCC) ferrite phase
(δ) and the face center cubic (FCC) austenite phase (γ), with equal volume fraction [23]. The
typical microstructure and individual unit cell for each phase are illustrated in Figure 2.
Figure 2: a) A SEM micrograph of the microstructure of duplex stainless steel with austenite
phase, γ, and ferrite phase, δ; b) the BCC unit cell of ferrite phase; and c) the FCC unit cell of
austenite phase. The unit cells were made in illustration program Avogardo [24]. The SEM
image was a part of an image taken by Cem Örnek.
Duplex stainless steel often exceeds the corrosion resistance of its ferrite and austenite
counterparts, especially in stress corrosion cracking, indicating an intrinsic synergistic effect
between the two phases in the microstructures [25]. Within this steel family there are three
groups of steel grades: low alloyed lean duplex stainless steel, standard duplex stainless steel,
and the high alloyed super duplex stainless steel. The different grades have deviating corrosion
resistance performance. As an example, Figure 3 shows the potentiodynamic polarization
curves of lean (green), standard (red) and super (blue) duplex stainless steel immersed in 0.1 M
NaCl aqueous solution with pH 4, at ambient temperature.
12
Figure 3: Potentiodynamic curves for super (2507, blue), standard (2205, red) and lean (2304,
green) duplex stainless steel in 0.1 M NaCl. The curve of the lean duplex stainless steel show
high current spikes and sadden current increase at a relatively low potential, indicating pitting
corrosion, while the curves of the standard and super duplex stainless steel only show smooth
current increase only at high potentials indicating absence of pitting events, and transpassive
breakdown occurring at high potentials.
In Figure 3, the lean duplex stainless steel shows a clear sign of metastable pitting by current
density increase spikes. Metastable pitting refers to pitting initiation events that becomes
prohibited by repassivation. After exceeding the breakdown potential, EB, the current density
drastically increases with the increasing potential during potentiodynamic measurements in
ambient temperatures. Above the breakdown potential, the potentiodynamic polarization curve
shows discontinuities due to metastable pitting formation and repassivation events, which
influence the measured current density. The standard grade, shown as the red curve, have some
irregularities in the curve, which possibly indicate metastable pitting events, but no pitting
corrosion occurred. On the other hand, super duplex stainless steel, shows a low and stable
current density up to the transpassive region, where the current density increases drastically.
For super duplex stainless steel, no sign of pitting events is visible in the potentiodynamic curve.
For resistive materials, such as super duplex stainless steel, the high breakdown potential
becomes hidden by the occurrence of oxygen evolution, which induces a drastic increase of
current density. This complication makes it difficult to understand the passivity breakdown for
the super duplex stainless steel. To observe the pitting potential by electrochemical techniques
requires an increase of temperature, above a critical pitting temperature. The measurement of
critical pitting potentials is a common industrial practice, but this will not be covered in this
thesis. Instead other techniques (synchrotron-based analysis) are applied to investigate passivity
breakdown of super duplex stainless steel.
13
2.2 Role of alloying elements The elemental composition of super duplex stainless steel alloy enables it to perform well in a
large variety of industrial applications and sustain highly corrosive environments. The name
specification of super is due to the superior resistance towards local corrosion, such as pitting
corrosion, which has been accomplished by effective alloying with N, Mo, and Cr. Empirically,
the pitting corrosion resistance in Cl- environments can be estimated (here from the elemental
composition in wt%) by a pitting resistance equivalent number, PREN [11]:
𝑃𝑅𝐸𝑁 = 𝐶𝑟% + 3.3 ∙ 𝑀𝑜% + 16 ∙ 𝑁% (1)
The PREN parameter for stainless steels is in a range between ca. 17-48, where the super grade
is above 38. Super duplex stainless steel 2507 has a PREN of ca. 42.5 [26], so it is entitled the
grade of super [11]. The ferrite and austenite phases of super duplex stainless steel have a
designed composition and annealing temperature to give both phases a PREN above 40 [27].
The role of alloying elements is of critical importance for production of corrosion resistant
alloys. The PREN is calculated from the bulk composition and does not consider the surface
film or the calculated value during heat treatment, as can be seen in Figure 4.
Figure 4: The PRE number variation for the austenite and ferrite vs. annealing temperature.
Image provided by Ulrika Borggren at Sandvik AB.
Table 1: Average elemental composition (wt%) of super duplex stainless steel 2507 at 1075˚C,
calculated by Ulrika Borggren at Sandvik materials Technology, using Thermo-calc software.
Element Cr Ni Mo Mn Si N Cu C Fe PREN
Global 24.9 6.90 3.90 0.80 0.30 0.30 0.30 0.03 62.6 42.6
Ferrite 27.2 4.79 5.17 0.71 0.27 0.06 0.19 0.01 61.6 45.2
Austenite 23.6 8.08 3.19 0.85 0.32 0.44 0.36 0.04 63.1 41.1
14
Table 1 presents the calculated composition and the PREN number for the average (global) and
for each phase of super duplex stainless steel 2507. The Cr content of duplex stainless steel is
between 20 wt% and 30 wt%, which far exceeds the stainless requirement of at least 13 wt%.
It has been shown that the Cr content has a large influence on the corrosion resistance, that is,
the higher proportion the better resistance properties. The high corrosion resistance is due to
the enrichment of stable Cr3+ oxide in the passive film, a bulk with a higher proportion of Cr
leads to a more protective oxide covering the surface [28, 29]. Moreover, besides the role in
passive film formation, the high amount of Cr in the bulk also facilitates the solubility of
nitrogen. Nitrogen is a strong austenite former and increases the austenite phase resistance
towards local corrosion, such as pitting. The mechanism for increasing pitting resistance by
nitrogen is not completely understood, but existing hypotheses include nitrogen-induced
ammonium ions reacting with aggressive ions and nitride formation with molybdenum on the
bulk/oxide interface [30, 31]. Molybdenum can also contribute to an enhanced corrosion
resistance during exposure at ambient and elevated temperature. Similar effects have also been
found by the addition of W and Cu [31, 32]. Mo exists in the passive film (to a small extent) in
the oxidation states Mo4+ and Mo6+. The amount of metallic Cr and Mo in the bulk decreases
the amount of Ni needed in the alloy. Ni is an important austenite former which increases the
corrosion resistance significantly. Both Ni and Mo are mostly enriched in the bulk layers
beneath the oxide due to preferential dissolution of Cr and Fe, which are the main constituents
of the passive film. The bulk layer beneath the passive film is called the “alloy surface layer”
which, during polarization, eventually inhibits the dissolution of Cr and Fe [1, 30, 32, 33].
15
3. Passivity and breakdown of stainless steel
3.1 The passive film and passivity Passivity is a metal state reached by the spontaneous formation of a passive oxide film on the
metal surface, and thus the current density decreases drastically due to the oxide film acting as
a barrier, increasing the corrosion resistance of the metal. Compared to the active state, the
current density drops approximately 100 -1000 orders of magnitudes across to the passive film.
The current density remains at that low level over a range of potentials, often a few hundred
millivolts. The current density drop is caused by the decreased metal dissolution rate, meaning
a reduced charge transfer at the metal/electrolyte interphase and ion transport across the passive
film. The dissolution rate in the passive state is independent on the applied potential, which is
the electrochemical driving force for the charge transfer reaction (oxidation and corrosion) and
ionic migration. The formation of a passive film on the metal surface establishes a potential
gradient at the material/electrolyte interphase. An increased anodic potential lead to a larger
potential divergence between the material and electrolyte. The potential gradient results in the
growth of the passive film, as described by the high-field theory [1, 34]. In the passive state,
there is no change of the oxidation state of the metal oxides which remain stable [7].
Stainless steel is a Fe-based alloy, and Fe dissolves quickly from the outer most surface when
immersed in an aggressive environment. Cr, which is the second most abundant element in the
bulk, also oxidizes quickly by the oxidant present in the environment. While the Fe is
preferentially dissolved, Cr remains at the surface forming an oxide layer. However, not all the
Fe ions leave the surface, some instead become part of the passive film, where Fe stays in the
Fe2+and Fe3+ oxidation states. The protective barrier of stainless steels is enriched by Cr3+
which, in the passive film, far exceeds the bulk concentration. For duplex stainless steels
containing a high amount of Cr, the oxide film has a Cr composition between 50%-70% in
acidic solutions. As mentioned above, Fe2+ and Fe3+ as well as Mo4+ and Mo6+ species are also
incorporated into the oxide film. During anodic polarization, the chemical composition changes
so that the content of the Cr3+ species decreases due to increased oxidic Fe species [1, 7, 32].
There is a general agreement that the passive film of stainless steel has a bilayer structure, with
the oxide layer providing the protective barrier properties, and a top layer of oxyhydroxide. The
oxide layer is enriched of Cr3+, which is the main component of the passive film on high alloyed
stainless steels. The composition of the passive film depends on the chemical composition of
the alloy, and is also influenced by the electrolyte, pH and potential [7]. There have also been
observations that while Cr is enriched in the oxide layer, Fe is enriched in the top layer during
anodic polarization [18]. The passive films formed on stainless steels are initially amorphous,
but crystallize with time, and become nanocrystalline [2, 35].
Figure 5 schematically shows the structure of the passive film consisting of two oxide layers.
While the top layer is hydroxide rich, the oxide layer at the interface to the metal contains
almost exclusively metal oxides. The Ni enriched alloy surface layer below the oxide layer is
also included. However, the exact details of the thickness and composition of the multiphase
duplex stainless steel is still debated in the literature [16, 36].
16
Figure 5: Model of the near-surface region of duplex stainless steel. The bulk contains of the
austenite (γ) and the ferrite (δ). Between the bulk and the oxide layer, there is a Ni rich alloy
surface layer. The oxide layer is composed of Cr3+ and Fe2+oxides, while the top layer contains
of Cr3+hydroxide and a mixture of oxide and hydroxide Fe3+ species.
3.2 Passivity breakdown and passive film degradation The demise of the passive state of metal/alloy, and thereby disruption of the passive film, will
eventually occur with time under certain environmental conditions. The transition from passive
state to breakdown is frequently discussed in term of the concept of local corrosion, such as
pitting or crevice corrosion. In the literature the critical factors of passivity breakdown have
been discussed extensively. For instance, whether the pit formation and breakdown are caused
by the failing film accompanied with its decreasing capability of protecting the bulk, or if it is
the pit formation that is the reason for the breakdown of the film. An alternative argument is
that both are true and are intertwined [37, 38].
Pitting initiation and local raptures by pitting corrosion are not fully understood. The presence
of aggressive ions, such as chlorides, increase the probability for passivity breakdown and
material failure to occur. Local disruptions, such as pitting, often become autocatalytic due to
local pH changes (i.e. acidification) within the pit [39-41]. However, the stability of the passive
film and its failure mechanism depends on several factors, including the electrolyte pH,
temperature, concentration of aggressive ions and the metal or alloy immersing time, potential
and composition [30, 42].
Passive films are dynamic systems where dissolution and formation occur simultaneously. In
the passive state, the dissolution rate decreases so that the formation rate exaggerates the
dissolution. The passive film is stable when both rates are equal, i.e., the system is in a steady
state. This is true until the metal reaches a critical potential, beyond this point the dissolution
rate will increase and surpass the formation rate, leading to breakdown [43]. The theory of how
passivity breaks down by initiation of pitting is categorized into three mechanisms which
include: i) mechanical (both mechanical ruptures and mechanical film thinning); ii) aggressive
ion adsorption, and iii) aggressive ion penetration [8, 41, 44]. Breakdown caused by mechanical
ruptures occurs only in the presence of aggressive ions, such as Cl-, which prevent repassivation
processes by their adsorption on the film. Further, aggressive ions are adsorbed on the film
17
surface leading to formation of dissolvable complexes which enhance local dissolution and
therefore local thinning of the film.
Aggressive ion adsorption can occur via vacancies created by oxide ion displacements [7, 41].
The surface adsorption is the first step of the penetration mechanism, where the aggressive ions
are transported and accumulated at the film/bulk interface. This process has been reported for
stainless steel by investigators showing TEM images of accumulated Cl- ions after anodization
[18]. However, the transportation process is not fully understood. One possible model for the
mechanism is the point defect model (PDM) where the ion transport is assumed to occur
through defects, such as oxygen vacancies [41]. PDM is based on a flow of vacancies and ions
and their flowrates depend on the electric field through the film [5, 45].
As shown in Figure 3, super duplex stainless steel does not experience any pitting corrosion
under ambient conditions. With increasing potential, passivity breakdown will occur, which is
commonly believed to be caused by further oxidation of the stable Cr3+ oxide to the more
solvable Cr6+ oxide. The properties of the electrolyte, such as the temperature and pH influence
the composition and stability of the passive film. For example, in more acidic electrolytes is Cr
more predominant in the passive film of stainless steel, while in more alkaline electrolytes is
the content of Fe higher [46, 47]. The temperature of the electrolyte also influences the
properties of passive films. Passivity breakdown occurs at lower potentials that may be caused
by an increase of oxide defects, such as vacancies, due to enhanced temperatures [20, 48].
In this thesis, the focus is to monitor the changes of the passive film, aiming to achieving an
improved understanding of the mechanism leading to passivity breakdown during anodic
polarization at ambient conditions in NaCl solutions. To describe different aspects of the
passive film, several different synchrotron methods have been used in the study.
18
4. Experimental techniques
4.1 Electrochemical techniques Electrochemical techniques that have been used in this thesis are direct current (DC)
potentiodynamic and current measurements, and alternating current (AC) electrochemical
impedance spectroscopy (EIS). These techniques are based on the transfer of charged particles,
that can be regulated by electrical devices [49, 50].
For a metal immersed in an electrolyte containing anions and cations, an electrochemical double
layer is formed at the metal/electrolyte interphase. The charge difference across the double layer
leads to establishment of an equilibrium potential of the metal in the electrolyte. Metal corrosion
in an electrolyte involves both anodic (oxidation) reaction and cathodic (reduction) reaction,
and each of these electrochemical reactions has its own equilibrium determined by the
thermodynamics of the reaction. Without any external circuit, i.e. under open-circuit condition,
the anodic and cathodic reactions are coupled, meaning that the electrons released from the
anodic reaction must be consumed in the cathodic reaction. This constrain leads to a change (by
i.e., polarization) of the electrochemical potential from the equilibrium potential of the anodic
and cathodic reactions, respectively, towards a so-called mixed potential (also called corrosion
potential). Where the anodic reaction rate and cathodic reaction rate reach an equal point is the
mixed potential. Through an external electrical circuit, electrons can be injected to or withdrawn
from the metal surface, leading to a perturbation of the equilibrium and change of the potential,
and thus a current flow in the electrical circuit. The kinetic theory of the electrochemical
reactions, i.e., the relationship between the electrode potential (and polarization potential) and
the electrochemical current (a measure of reaction rate) form the basis for electrochemical
techniques used for the study of electrochemical phenomenon. In DC electrochemical
techniques, the electrochemical current and the resistance follow Ohms’ law as in electrical
engineering.
Ohms’ law states the relationship between resistance, R, voltage, E, and current, I [49, 50]:
𝑅 ≡𝐸
𝐼 (2)
To investigate electrochemical processes, associated with passivity breakdown occurring at
anodic polarization, one must push the system from the equilibrium (strictly speaking, open -
circuit potential) towards higher potentials. Higher potentials increase the kinetic rates of anodic
reactions, which make it possible to study the oxidation processes as a function of applied
potential like in the potentiodynamic curves, which is schematically shown in Figure 1 and
measured for different stainless steel grades in Figure 3. The anodic reaction taking place during
increasing potentials depends on the material and electrolyte as discussed in previous chapters.
In general, by using an electrochemical cell and applying a polarization potential to the sample
using an electrochemical instrument, the electrochemical reactions can be pushed towards
positive (over potential) or negative (under potential) directions, which promotes either anodic
oxidation or cathodic reduction reactions [49, 50].
Anodic reactions in a corroding system involve ionization of metal and emission of electrons:
𝑀 → 𝑀𝑛+ + 𝑛𝑒− (3)
19
The corresponding cathodic reactions occurring at the same time at the metal surface
(metal/electrolyte interphase) [50]:
𝐻𝑦𝑑𝑟𝑜𝑔𝑒𝑛/𝑝𝑟𝑜𝑡𝑜𝑛 𝑟𝑒𝑑𝑢𝑐𝑡𝑖𝑜𝑛: 𝐻+ + 𝑒− →1
2𝐻2 (4)
𝑂𝑥𝑦𝑔𝑒𝑛 𝑟𝑒𝑑𝑢𝑐𝑡𝑖𝑜𝑛: 1
2𝑂2 + 2𝐻+ + 2𝑒− → 𝐻2𝑂 (5)
The rate of electrochemical anodic and cathodic reactions can be measured by either DC or AC
techniques. DC measurements determine the net current flow resulted from the electrochemical
reactions during the applied potential [49, 50].
4.1.1 Direct current techniques
When applying a constant potential, to the sample, the current transient (current vs. time) shows
the electronic and ionic responses of the sample surface, as schematically shown in Figure 6a.
In potentiodynamic polarization measurements, the potential is changed, and the resulting
current is recorded continuously in a certain potential range, as illustrated in Figure 6b.
Figure 6: Schematic curves of a) current transient response showing the ionic and electronic
responses to an applied potential. Adapted from ref [51]. b) A potentiodynamic curve, where
the current is measured during potential sweep in the anodic or cathodic domain. At the open
circuit potential, OCP, no potential is applied, and the net current is zero. At this point the rate
is the same between the anodic and cathodic reactions, which is also the corrosion rate.
It is possible to gain information of the metal/electrolyte interphase qualitatively (e.g. oxide
growth or instability) and quantitatively (e.g. oxidation rate) from the current transient. The
decrease of current density with time, could be a result of a formation or growth of a resistive
surface layer increasing the resistance. The increase of the current density with time indicates
a decreasing resistance, which could be due to a decrease of the barrier properties of the surface
film. Another reason for a current increase at high potentials can also be related to oxygen
evolution reactions, which increases the current density [51].
20
Potentiodynamic curves show the current density versus applied potential, as exemplified in
Figure 6b. The polarization curves provide information regarding the different metal states,
ranging from the active state, where the oxidation rate is high and metal dissolves quickly into
ions releasing electrons, to the passive state, where a stable oxide covers the surface decreasing
the oxidation rate. At high potentials, the metal reaches transpassive state, where the barrier
oxide is further oxidized and dissolved, and the current transient is also increased drastically by
the oxygen evolution [7].
4.1.2 Electrochemical Impedance Spectroscopy (EIS)
EIS is a technique based on alternating current (AC). The measurements are made by applying
a small sinusoidal AC perturbation potential of varying frequency, often from high to low, and
measure the current response from the metal/electrolyte interphase. At each frequency, a
potential signal is applied to the system and the current density output signal is recorded. By
analyzing (transformation) the output signal compared to the input signal, the difference in
phase (offset) and the amplitude are obtained, as shown in Figure 7 [49, 50]. The measurements
yield electrochemical impedance spectra, i.e., impedance as a function of frequency, which
provide information of different kinds of electrochemical processes, with different relaxation
time, of the metal/electrolyte interphase.
Figure 7: The applied sinusoidal potential input signal and measured current output signal.
The amplitude and the phase shift of the signal makes it possible to probe electrochemical
reactions. Adapted from ref [50].
The frequency dependence of the impedance makes it possible to detect and analyze fast
processes, such as reaction kinetics, at high frequencies, and slow processes, such as mass
transport (diffusion), at low frequencies. In the frequency domain Ohms laws is still valid, and
the resistance becomes impedance instead. The EIS spectra can be plotted in Nyquist plot,
where the imaginary of the impedance is plotted against the real part of the impedance.
Moreover, the real and imaginary variables of the impedance can be transformed into the
amplitude and phase angle variables by Euler’s relationship [52]:
𝑍(𝜔) =𝐸(𝜔)
𝐼(𝜔)= 𝑍0 (cos(𝜙) + 𝑖 ∗ sin(𝜙)) (6)
21
where Z0 is the amplitude, ϕ the phase shift and the wave vector; angular frequency ω=2∙πf,
where f is the frequency. The amplitude and phase shift between the input and output signal are
commonly plotted as a function of frequency in Bode plots. The data analysis is commonly
made by creating a descriptive equivalent electrical circuit of the investigated system. The
electrical circuit is modeled by elements such as resistors, capacitors, and inductors,
representing the resistance, capacitance and inductance of the system, respectively [53, 54].
The electrochemical resistance is a measure of the resistance towards the electrical current flow,
which is directly proportional to the real impedance Zw [53]. The capacitance is a measure of
the accumulation of electrical charge at the interphase, and the impedance of a capacitor is
inversely proportional to the frequency [53, 55]:
𝑍𝑐 =1
𝑗𝜔𝐶 (7)
In this thesis, the capacitance of the system is represented by a constant phase element (CPE)
which is used instead of pure capacitance [53]. This element takes into account non-ideal
response of the electrochemical capacitor caused by distributed features (e.g., roughness and
other heterogeneities) of the interphase [54]:
𝑍𝑄 =1
𝑌0(𝑗𝜔)𝑛 (8)
in which Y0 is the admittance of an ideal capacitor and the n a fitting parameter, which can vary
between 0 to 1, where 0 represents a pure resistor and 1 a pure capacitor [53].
The inductance is a measure of the preventions towards system changes, which creates a
magnetic field that is proportional to frequency. It is often associated with adsorption processes
of charged species at the metal surface [53-55]:
𝑍𝐿 = 𝑗𝜔𝐿 (9)
22
Figure 8 shows an example of a simple equivalent circuit and corresponding EIS spectra.
Figure 8: a) Equivalent circuit describing a film on a substrate in an electrolyte, b) the Nyquist
plot, c) Bode modules, and d) Bode phase.
Figure 8a shows an equivalent circuit describing a film (e.g., passive film) on a substrate in an
electrolyte, b) Nyquist plot, c) Bode modulus and d) Bode phase plots of the impedance
spectrum, from which the resistance and capacitance can be obtained. The solution resistance
RS is the resistance of the electrolyte between the working and reference electrode. The charge
transfer resistance (R) is the measure of the resistance against the electrochemical reaction (e.g.,
metal oxidation). When R is measured at the open-circuit potential, it is the polarization
resistance of the corroding system, a measure of the corrosion resistance. In quantity, R is the
difference between the solution resistance and the total resistance [56]. In the Nyquist plot, each
frequency is one data point, which is measured from high to low frequencies [57]. In the Bode
plots, the modulus shows the amplitude of the impedance, while the phase shows the offset of
the phase between the input and output signal [57].
23
4.2 Scanning Electron Microscopy / Electron Backscatter Diffraction Scanning electron microscopy (SEM) is a versatile surface imaging technique. In this thesis,
SEM was used for identifying surface areas for further investigation by synchrotron techniques
and post analysis of surfaces.
In this technique, the incidence electrons are emitted from a heated filament, such as a
tungsten wire. The electron beam is focused on a specific area of the surface by condenser and
objective lenses. The electrons are used to probe the surface by interacting with its electrons
within a depth of several µm. The type of information gained is dependent on the origin of
detected signal and the detector used. A common way to describe the electron detection of SEM
is by the illustration of the pear-shaped interaction seen in Figure 9 [58]:
Figure 9: The” SEM-pear” illustrating how information is gained from different sample depths.
Different kinds of information can be obtained from different depths [58]. Electron backscatter
diffraction (EBSD) is a SEM technique, from which information about a surface’s phases and
grain orientation can be gathered [59]. The distribution of elements within the depths of 1-2 μm
can be obtained by energy dispersive X-ray analysis [60]. It is possible to differentiate between
the phases of stainless steel from the backscattered electrons. In backscatter mode of SEM
images, the austenite has a brighter contrast because of the elemental portioning of the heavier
element, Ni.
In EBSD, the electrons are inelastically scattered from several tens of nanometer beneath the
surface and from all directions. The diffracted signals reaching the detector that satisfy Braggs
law are recorded. Because the sample has a three-dimensional structure, it is possible to
determine several diffracted signals with the same angle, due to symmetry in the arrangement.
24
The detector registers lines called “Kikushi bands”, where each line is related to a direction of
a crystallographic plane, such as (001) [61].
In this thesis, EBSD was measured to detect the grain orientations over the surface of super
duplex stainless steel samples. All mentioned grain orientation (hkl) are hereafter assigned as
parallel to the sample surface. The measured grains were divided into three general directions
of (001), (101) and (111), where the atomic arrangement is shown in Figure 10a. How the
crystals are located at the stainless steel surface is shown in Figure 10b.
Figure 10: a) The atomic arrangement for each orientation and crystal structure of the
austenite (FCC) and ferrite (BCC) phases of stainless steel and b) how the crystallographic
orientations are placed at the surface. Adapted from ref [62].
25
4.3 Synchrotron based techniques X-rays were first detected in 1895 by Wilhelm Conrad Röntgen, who was awarded the Nobel
prize in 1901 for his discovery. X-rays can be produced when high energy electrons bombard
a metal target (anode). For in-house sources, the targets can be Mg and Al, providing X-rays of
1253.6 eV (Mg Kα) and 1486.6 eV (Al Kα) respectively. Today, it is possible to produce X-
rays of significantly higher energies at synchrotron facilities, which provides numerous
benefits, such as (1) tuneable photon energy, (2) high intensity (brilliance), (3) small spot size,
and (4) polarization control. These facilities contain a storage ring, where electrons are kept
orbiting at a velocity close to the speed of light in vacuum by several bending magnets. The
bending magnets deflects the electrons, leading to the emission of X-rays. The generated
spectrum produces photons with energies spanning between single eV and 100 of keV
simultaneously. Higher intensities for selected photon energies can be achieved by using
different insertion devices, such as wigglers and undulators, in the straight sections. An
undulator is a device that consists of a periodic array of alternating magnets. When electron
bunches pass through the undulator in an oscillating movement, due to the alternating magnetic
field, photons are emitted. The photons are emitted at different locations and with different
wavelengths along the undulator and interfere constructively to a resulting polychromatic beam.
The energy spectrum of an undulator has a series of sharp peaks (harmonics) and the photon
energy position for these harmonics can be tuned by changing the gap between the magnetic
arrays [63].
An overview of a synchrotron facility can be seen in Figure 11. At the beamlines, the X-rays
are further tuned to the energy required for a specific experiment by monochromators which
can be made of for example Si crystals. The beam is also focused by several mirrors [63].
Figure 11: Illustration of a synchrotron facility with a storage ring containing bending
magnets, insertion devices, such as undulators and wiggler. The photons emitted from the
electrons are used at beamlines where the X-rays are monochromatized and focused by mirrors
on the sample.
26
4.3.1 X-Ray Diffraction (XRD)
XRD is one of the most commonly used methods for the determination of the atomic structure
of materials in many fields of science [64-66]. This chapter will briefly describe how XRD has
been used to study passive films and transpassive layers in corroding environments. The
technique can detect crystalline structures, as Bragg peaks, in the diffractogram. It is also
possible to study amorphous structures, such as amorphous surface oxides, which appears as a
diffuse background. This background can be probed at low incidence angles. In this thesis,
XRD was used to gain structural information about the passive films and the bulk during the
in-situ/operando measurement. It was also used to determine strain in the sample.
4.3.1.1 XRD Principle
X-rays can interact with electrons in an atom leading to scattering, resulting in new
electromagnetic waves of other frequencies and directions. However, for atomic nuclei,
scattering can be considered negligible since the nuclei is many orders of magnitudes heavier
than electrons, and therefor are less accelerated compared to the electrons. When atoms,
arranged periodically in a crystal, are irradiated by a polarized X-rays, of a certain frequency,
the electrons surrounding the atoms will generate beams of scattered X-rays of the same
frequency. In certain directions, X-rays beams scattered from different atoms will be in phase
and interfere constructively with each other. As the atoms are arranged in a periodic array, the
scattered beams will create an interference pattern, and form Bragg peaks due to the
constructive interference. From this information, it is possible to measure the lattice plane
distances, d, which is described by Braggs law [67]:
2𝑑𝑆𝑖𝑛(𝜃) = 𝑛𝜆 (10)
where θ is the angle of the reflected beam, λ is the wavelength and n is an integer. Figure 12
shows the geometrical relationship of Braggs law when periodically arranged atoms are
irradiated by X-rays [63].
Figure 12: Bragg diffraction geometry. When two beams of the same wavelength are incident
on a crystalline solid they will be scattered by the atoms. The lower beam will travel 2d sin(θ)
more than the one above and constructive interference will occur when this extra length is equal
to an integer multiple of wavelength.
27
4.3.1.2 Surface sensitivity of XRD
X-rays are well-known for their penetrating capabilities and can penetrate through a liquid
electrolyte to probe a sample during in-situ/operando measurements. However, under certain
conditions, it is also possible to probe the very surface of a sample, for instance the surface of
a corroding steel sample. The reasons for this ability are explained below.
The refractive index determines how fast light travels through different mediums, the higher
the index, the lower the velocity. When light transfers from one medium to another with a
different refractive index, the light will change the trajectory direction. X-rays can penetrate
deep into matter due to its low interaction, which gives large signals from the bulk excelling
the contribution from the surface. To enable data collection from the surface, gracing incidence
angles are used for the measurement which increase the contribution from the surface. At a
certain low angle, called the critical angle, it is possible to obtain total reflection at the
interphase of the two medias, such as an alloy surrounded by air. If the light reaches the critical
angle the X-rays will create a standing wave with a penetration depth corresponding to the
wavelength. For angles below the critical angle, the refractive index will reach a value below
1, unity, which means that the X-rays will all be reflected and not refracted, then what is known
as the total external reflection condition is reached. Materials that are conducting, like most
metals, are interacting with the light through their electrons that absorb and reflect the light.
The general formula of refractive index, n, is:
𝑛 = 1 − 𝛿 + 𝛽𝑖 (11)
where δ is the dispersion coefficient and β is the absorption. However, because X-rays interfere
lightly with matter, these coefficients are small [63].
Snells’ law describes the angle relationship between the incident αi and reflected αr radiation:
cos(α𝑟) = n ∗ cos(α𝑖) (12)
When the incident angle is very small the critical angle can be described as:
cos(α𝑐) = n (13)
By expanding Snells’ law the critical angle can be estimated to [68]:
α𝑐 ≈ √2𝛿 (14)
Therefor critical angles can be obtained from reflectivity measurements. In this thesis, the
critical angle for stainless steel, measured using 20.5 keV, was 0.25˚.
28
4.3.1.3 Measurement of strain by XRD
Strain can be detected from changes in the form of crystalline Bragg peaks if the strain causes
changes in the crystalline structure. A regular crystalline peak from a non-strained material is
shown in Figure 13a. If the crystalline structure is under a strain that, for example expands the
inter-layer distance in the crystal, the Bragg peak is shifted towards smaller values, as illustrated
in Figure 13b. For a crystal contraction the peak instead shifts towards higher values. Non-
uniformed strain, caused by local crystal imperfections, such as impurities or vacancies in the
lattice, will instead cause a broadening of the Bragg peaks, as in Figure 13c. The broadening
increase, β, can be calculated quantitatively by the following equation:
𝛽 =𝜆
𝜏∗cos (𝜃) (15)
where λ is the wavelength, θ is the Bragg position and τ is the average grain size [69].
Figure 13: Illustration of the influence of strain on the Bragg peak position and shape. a) No
strain results in a well-defined Bragg peak at certain peak position. b) under uniform strain,
such as expansion d1<d2, the Bragg peak position shifts to lower Bragg positions. c) In the case
of non-uniform strain, e.g., one plane in the crystal is more strained than another, the peak
becomes broader. Adapted from ref [70].
29
4.3.2 X-Ray Reflectivity (XRR)
XRR measurements provide information about surface film thickness, electron densities and
roughness. The experiments are performed in angles smaller than 5o, starting close to the critical
angle, αc, where the X-rays are totally reflected from the sample surface. The angle of the
detected X-rays in XRR is equivalent to the angle of incidence, as shown in Figure 14a. With
increasing angle, the intensity of the reflected beam drops and creates an interference pattern
called Kiessig fringes, as shown in Figure 14b. From the fringes it is possible to estimate the
film thickness since the period of the fringes is inversely proportional to the film thickness. The
amplitude of the fringes is a way to estimate the electron density of the film. The decay of the
slope shows the roughness of the film [51, 56, 71].
Figure 14: An overview of X-ray reflectivity showing a) the reflectance during a measurement
of a thin film, b) an example of the resulting spectra and the information that can be gained.
Adapted from ref [71].
30
4.3.3 X-Ray Photoelectron Spectroscopy (XPS)
4.3.3.1 XPS Principle
XPS is commonly used to investigate the chemical composition of surfaces layers and the
thicknesses of thin films. The photoelectric effect was discovered by Hertz in 1887 [72] and
explained by Einstein in 1905 [73]. The photoelectric effect is the process of emitting electrons
from atoms by irradiation of photons creating an electric current. XPS was developed in the
1950s by K. Siegbahn and co-workers. K. Siegbahn found that each element has a characteristic
electron binding energy that gives rise to a characteristic set of peaks in the photoelectron
spectrum [74].
A schematic representation of the principle of XPS is shown in Figure 15. When photons are
impinging on a material, they are absorbed by the electrons of the elements in the material. If
the energy of the photons is higher than the binding energy of an electron, the atom becomes
ionized and the electron is ejected from the atom. Electrons that escape the energy of the
vacuum barrier, EV, become free electrons that can be detected. The mathematical relationship
between the binding and kinetic energy, EB and Ek, are:
𝐸𝐵 = ℎ𝑣 − 𝐸𝑘 − 𝜙𝑠𝑝𝑒𝑐𝑡𝑟𝑜𝑚𝑒𝑡𝑒𝑟 (16)
where hv is the photon energy and 𝜙𝑠𝑝𝑒𝑐𝑡𝑟𝑜𝑚𝑒𝑡𝑒𝑟 is the work function of the spectrometer [6].
In XPS, core level electrons are detected, i.e., the detected electrons are originating from the
inner filled shells of the atom [63]. In this thesis, the peak position (binding energy) and peak
intensities of various core-levels were measured by XPS.
Figure 15: Illustration of the X-ray photoemission process. The atoms are irradiated by a
photon beam with the energy hv. If the electron absorbs an energy hv ≥ EB (binding energy of
a core electron), the atom becomes excited and an electron is ejected with a unique kinetic
energy, referred to the specific element. Image inspired by ref. [63].
31
4.3.3.2 Surface sensitivity of XPS
XPS uses low energy electrons to probe surfaces. When electrons are emitted from a metal at
low kinetic energies, they originate mostly from the surface atoms, due to their short mean free
path. The mean free path is the distance that an electron travels in a solid before it experiences
inelastic scattering and it depends on the electron’s kinetic energy. Figure 16 shows the inelastic
mean free path of different materials as a function of the kinetic energy of the electron [75]. As
it depends, to a less extent, on the specific material, the curve shown in Figure 16 is also called
the universal curve for the electron mean free path. The curve demonstrates that the mean free
path of electrons in the energy range of 10 to 1000 eV is maximum 5 nm, providing the surface
sensitivity of XPS at these energies. In later years the mean free path for energies higher than
1000 eV has been calculated with mathematically algorithms with the premises from the
universal curve [76, 77].
Figure 16: The universal curve of inelastic mean free path. Adapted from ref. [75].
4.3.3.3 Chemical shifts and spectra fitting
If an electron is bound to an atom in an element that is surrounded by another chemical
environment other than the pure element, a so-called “chemical shift” can be detected in the
XPS spectra [63]. This shift is a fingerprint for different species present on the surface. The
number of components in a core-level spectrum corresponds to the number of chemically
different atoms, the peaks in the spectrum can also include that from spin doublets. Therefore,
the name ESCA (Electron Spectroscopy for Chemical Analysis) is also used for this technique
[78]. The binding energy shifts are often correlated to the charge of an atom by electrostatic
effects since the removal of a valence electron leads to a stronger binding energy to the nucleus
of the remaining electrons. For instance, the higher the positive charge of an ion in a molecule
is, the higher the binding energy will shift. Another example is a metals’ core level in an oxide,
which is shifted towards higher binding energies, compared to the pure metal core level binding
energy, due to the transfer of electrons from the metal to the oxygen atoms in the oxide. Figure
17 shows the core level of 2p3/2 for Cr and Fe. For this core level, the spectra can be
deconvoluted into three peaks, one metal peak and two oxidic peaks for both elements. The
deconvolution of Fe spectra is complex and difficult because it contains many different possible
oxidic peaks e.g. FeO, Fe2O3, Fe3O4, as well as satellite peaks. In this thesis work the spectra
fitting has been simplified to Fe2+ and Fe3+ components.
32
Figure 17: The XPS spectra of core levels of 2p3/2 for Cr and Fe, respectively, obtained at 4
keV. The Cr spectrum can be deconvoluted into Cr metal (Crmet), Cr oxide (Crox) and Cr
hydroxide (Crhy). The Fe spectrum, for simplification, can be deconvoluted into Fe metal
(Femet), and oxidic components in Fe2+ and Fe3+ states.
4.3.3.4 XPS data analysis
A core level spectrum is often a convolution of several binding energy peaks due to the presence
of atoms in different chemical surroundings. To quantitatively analyze a core-level spectrum, a
deconvolution is often required, as important information is contained in the shape of a
spectrum. There are various factors that contribute to the shape of a core-level photoemission
peak. For solids, a Gaussian distribution accounts for the broadening of the line due to the
experimental resolution (which is dependent on the monochromator and analyzer), excitations
of quantized vibrations in the solid lattice (phonons) and disorder [67, 79, 80], while the lifetime
broadening has a Lorentzian shape [75]. Each core-level has an intrinsic width due to the finite
lifetime of the hole state according to Heisenberg’s uncertainty principle:
∆𝐸 =ħ
∆𝑡 (17)
where E is the energy, t time and ħ is the Planck’s’ constant divided by 2π.
The spectra deconvolution in this work was performed using both Doniach-Sunjic and
asymmetric Lorentzian line shapes to describe core-level line shapes from the metallic elements
[81], and Gaussian/Lorentzian, (30%/70%) line shapes for the oxidized elements [82, 83]. The
asymmetry for photoelectrons signal from metals was shown already in the 1970’s, which is a
result of the large emission of electrons and electron interactions [84]. However, there are many
ways to describe this asymmetry. Doniach-Sunic peak shapes, DS(α, n) contain one asymmetry
parameter α and one convolution width parameter n [85]. For asymmetric Lorentzian peak
shapes are in the form of LA(α, β, m) where α and β are the Lorentzian tail shape of either side
of the highest point of the peak and m is the Gaussian width [86]. The later has been commonly
used in the investigations about passive films of stainless steel [36, 83].
33
XPS can also be used to determine composition and thickness of surface films, such as oxidized
layers of metals. For native and passive films of stainless steel the thickness calculations can
be done as follows [36, 87]:
𝑑𝑖𝑛 = 𝜆𝑗𝑖𝑛𝑠𝑖𝑛(𝜃) ln (1 + (
𝐼𝑗𝑖𝑛𝐷𝑗
𝑚𝑒𝑡𝜆𝑗𝑚𝑒𝑡
𝐼𝑚𝑒𝑡𝑀 𝐷𝑜𝑥
𝑀 𝜆𝑜𝑥𝑀 )) (18)
𝑑𝑜𝑢𝑡 = 𝜆𝑗𝑜𝑢𝑡𝑠𝑖𝑛(𝜃) ln (1 + (
𝐼𝑗𝑜𝑢𝑡𝐷𝑗
𝑖𝑛𝜆𝑗𝑖𝑛
𝐼𝑗𝑖𝑛𝐷𝑗
𝑜𝑢𝑡𝜆𝑗𝑜𝑢𝑡)) [1 − 𝐸𝑋𝑃 (
−𝑑𝑖𝑛
𝜆𝑗𝑖𝑛𝑠𝑖𝑛(𝜃)
)] (19)
where d is the thickness, λ the inelastic mean free path, I the intensity, and Dj=ρj/Mj∙xj the
elemental concentration, in which ρ is the density, M the molar mass and x the atomic% of the
component j. The subscript met stands for metal, in for the inner oxide layer and out for the
outer oxide/hydroxide layer components that are expected in the passive film. These equations
are derived from the intensity decay of the metal signal when electrons from the metal atoms
travel through surface layers, one advantage using the metal/oxidic ratios is that several factors
such as cross section and transmission factors are not needed in these equations [87-89].
From spectra containing both metallic and oxidic peaks it is possible to calculate the thickness
of the oxide film. The spectra, such as those in Figure 17, for the Cr and Fe species are obtained
from the same area. With two elements, the thickness of oxide can be calculated using the
intensity of the second peak assuming that Crox and Fe2+ are only present in the oxide layer
(model in Figure 5). The dinCr are calculated by inserting one dox equation with all the content
related to Cr/ Crox and a second dinFe with its content related to Fe/Fe2+. The variable here is the
elemental concentration, x. Because this is calculated only from Cr and Fe data, the function
will contain two equations with the two unknowns, dox and x, (e.g. xCr=1-xFe). At the elemental
concentration where dinCr=din
Fe is the calculated thickness and content of the layer. The same
procedure is applied for the dout equation. The assumptions for these equations are that one
knows the layer partitioning within the film and in this thesis are only two elements (Cr and Fe)
are considered.
34
4.3.4 Photoemission Electron Microscopy (PEEM)
PEEM, is a microscopy technique where spectroscopy and microscopy are combined. Here, the
sample is irradiated by an X-ray beam over a large area, but a small area can be probed, which
is defined by the field of view (FOV) of the electron lenses in the microscope. The
photoelectrons emitted from the surface are transported towards several cathodic lenses onto a
CCD detector that creates an image from the FOV. The same electrons can then be further
analyzed based on their kinetic energy so that specific elements can be detected, following the
fundamentals of XPS. The image and spectroscopic data are then recorded as a stack of images,
in which spatial and spectroscopic data are combined. Each image in the stack represents one
measured binding energy. This makes it possible to obtain microscopic spectroscopic data, as
shown in Figure 18a, providing spatially resolved chemical information on a microscopic scale
[90, 91]. Using hard X-ray PEEM measurement (HAXPEEM), it is possible to probe larger
depths and avoid contamination of surface carbon that are present on surfaces and becomes a
problem in soft X-ray PEEM measurements.
Figure 18: Example of HAXPEEM measurement of stainless steel a) scan image stack for Fe
and local XPS spectra for the marked areas; and b) Work function PEEM image from the FOV.
Another way of creating a surface image is to use the work function, i.e. the minimum energy
needed to eject an electron from the surface. The work function can be used as a contrast
mechanism to image different phases or other variations on the surface. This approach can be
used to image the austenite and ferrite phases in duplex stainless steel, since the two phases
have different work functions as shown in Figure 18b.
A number of different light sources can be used for PEEM, ranging from UV light (5-10 eV) to
hard X-rays (up to 10 keV [92]). In this thesis, a UV light source was used to find platinum
markers (and thus the region of interest) that shine bright when irradiated by UV light due to
lower work function of Pt compared to the steel. The marker was placed on the steel surface to
facilitate repeatedly finding the same area on the surface. The experiment was performed using
energies between 3-6 keV to gain depth-dependent information, however, only results from 4
keV are included in this thesis.
35
4.3.5 X-ray Fluorescence (XRF)
XRF is a technique that can be used to investigate the chemical composition of a sample. The
technique uses X-rays with energies higher than the binding energy of a core level electron to
ionize an atom. The created electron vacancy will generate a core electron-hole, which will be
filled by an electron from a higher electron shell, de-exciting to the electron-hole in the lower
shell. As this happens, a fluorescence X-ray will be emitted, with an energy corresponding to
the energy difference between the higher and lower electron shell, as illustrated in Figure 19.
The de-excitation can also happen by the ejection of an Auger electron. The most prominent
process is determined by the fluorescence cross section of the element [67]. The data obtained
from an XRF measurement is plotted as a 2D diagram (intensity vs. energy) showing peaks for
which energy is unique for each element.
Figure 19: Overview of the XRF principle.
In this thesis, only the fluorescence X-rays from the de-excitation process were detected. XRF
was used to detect the dissolution of different elements in duplex stainless steels at stepwise
increased potential, by probing the electrolyte close to the sample surface with the X-ray beam.
In this way, the onset of the dissolution of various metals in the sample could be detected. For
instance, in a 1 M NaCl solution, Fe was detected in the electrolyte when the sample was
polarized at 900 mV vs Ag/AgCl.
When performing such measurements, a number of experimental issues needs to be taken into
account in order to quantify the amount of dissolution such as cross sections, attenuation, etc.
[93] and often the data needs to be normalized. Since the electromagnetic radiation is emitted
due to the electron relaxation the XRF measurements can be performed in ambient conditions
and even in liquids [94], as performed in this thesis work.
36
5. Results & Discussion Most of the results in this thesis were obtained from two synchrotron beamtime measurements
at two different beamlines, ID03 in ESRF and P22 at PETRA III, DESY. At ID03 in ESRF,
four techniques (XRD, XRR, XRF and EIS) were combined for the in-situ/operando
measurement, giving structural, chemical, and electrochemical information during anodic
polarization. At P22 in PETRA III, an ex-situ HAXPEEM measurement was preformed to
obtain microscopic chemical information of the steel surface. These two sets of different
synchrotron measurements provided information of different aspects of the passive film and its
breakdown of the super duplex stainless steel. This chapter summarizes the two experimental
setups, results, and discussions. Details can be found in the five published papers enclosed in
the thesis.
5.1 Passive film formation, stability, and degradation Information concerning the formation, stability and degradation of the passive film was
obtained in-situ/operando XRD, XRR, XRF and EIS experiments. The experiment made it
possible to determine changes in the structure, chemical/electrochemical properties, and metal
dissolution during a stepwise increase in applied potential. The comprehensive results from the
in-situ/operando experiments have generated new knowledge and improved our understanding
of the passive film, passivity breakdown, and metal dissolution of the super duplex stainless
steel in near neutral NaCl solutions.
5.1.1 In-situ/operando experimental setup The detailed experimental setup has been described in Paper (I-III). In short, the measurement
utilized a three-electrode electrochemical cell where the steel sample surface was exposed to
the electrolyte. The electrochemical cell, illustrated in Figure 20a, was fabricated from
polyether ether ketone (PEEK). The PEEK cell is translucent to X-rays, allowing investigation
of an electrochemical controlled steel surface using X-rays. The electrochemical cell consisted
of the steel sample as the working electrode placed at the bottom of the cell, a saturated
Ag/AgCl reference electrode, and a glassy carbon counter electrode, placed above the sample.
Inlet and outlet tubes of polytetrafluoroethylene (PTFE) were connected to the cell sides to be
able to exchange gas, pure water, and the corrosive electrolyte. The electrodes were connected
to a potentiostat to be able to control the electrochemical condition of the sample. The samples
were “hat shaped” form to ease the mounting inside the cell. The sample surface was polished
to be ultra-flat to facilitate the beam alignment and to be able to draw accurate conclusions from
the measurement. Hereafter will all mentioned potentials be in mV in respect to Ag/AgCl.
The experiment was arranged by recording electrochemical and X-ray signals during a
stepwise increase of the applied anodic potential, accelerating the surface oxidation/dissolution
reactions. Electrochemical current transient and impedance data were recorded at each applied
potential to investigate the electrochemical response and the resistive and capacitive properties
of the sample surface. The emitted reflection, diffraction, and fluorescence signals from the
sample during X-ray irradiation were recorded by 2D detectors, yielding XRR, XRD and XRF
data from the sample surface at each potential, as illustrated in Figure 20b-c. For the XRF
measurement at each potential step, the beam was raised to 3 mm above the sample surface to
37
record XRF signals from the electrolyte to detect metal species dissolved from the surface, as
illustrated in Figure 20d. The applied potential was increased from the open circuit potential up
to 1400 mV (transpassive regime).
In Papers I-II, the XRD was measured for the same sample first exposed in air and then
immersed in electrolyte (1 M NaCl solution). In Paper III, the electrolyte was 0.1 M NaCl, and
XRD was also measured in water in addition to the other conditions. The sample preparation
differed between the two experiments. In Paper I-II, the sample was prepared by grinding and
polishing first, followed by a fine polishing procedure to remove the surface strain. In Paper
III the sample was prepared only by grinding and polishing.
Figure 20: Schematic illustration of the experimental setup used for the in-situ/operando
measurements a) The electrochemical cell used to combine the X-ray scattering and
electrochemical measurements; and schematics of the b) XRR, c) XRD and d) XRF
measurement.
5.1.2 Summary of results from in-situ/operando measurements The traditional way to study the active, passive and transpassive states of metals is to perform
potentiodynamic polarization measurements [43]. During this measurement, the potential of the
sample is increased, and the resulting current is measured continuously in a large potential range
covering the different states. Even though this type of experiment can determine the potential
regions for different states, it does not give any structural or chemical information of the
changes of the surface layer of the metal. In Figure 21a, the potentiodynamic polarization curve
for duplex stainless steel in 1 M NaCl during ambient temperature shows the current response
to applied anodic potential.
The current vs. time and EIS spectra were measured at each applied potential (Paper II). The
current transient at the applied potentials, accompanied with the EIS spectra, revealed
decreasing resistive properties of the surface with increasing potential. Figure 21b shows that,
the current density decreases with time for potentials up to 900 mV. At 1000 mV the current
38
started to increase after an initial decrease, indicating an instability of the passive film.
Increasing current with time was detected from 1100 mV up to 1400 mV, indicating breakdown
of the passive film. However, in potentials from 1200 mV and above the current density is very
high, suggesting that oxygen evolution reaction also contributes to the measured current. The
resistive properties towards charge transfer, measured by impedance, decreased at the increased
applied potential above 900 mV, as shown in Figure 21c. The large decrease in the resistance
indicates that the barrier properties is lost, due to degradation of the passive film (became
defective or dissolved).
Figure 21: a) A potentiodynamic curve, b) current transients and c) EIS spectra of the super
duplex stainless steel polarized in 1 M NaCl.
5.1.2.1 Thickness and density of the surface layers
The XRR curves displayed clear, intensity oscillations, i.e., Kiessig fringes, up to a potential of
1300 mV, as shown in Figure 22, indicating the presence of a passive/oxide film on the surface.
However, at 1300 mV and 1400 mV, the fringes became indistinct and vanished at 1400 mV.
The decrease and eventual disappearance of the fringes from the XRR curve indicates a thinning
and finally vanishing of the passive film on the sample surface.
39
Figure 22: XRR results showing the fringes for the sample polarized at a) 1100 mV, b) 1300
mV and c)1400 mV.
Through fitting the XRR curve suing a model with a passive film on the top and an alloy surface
layer between the passive film and the bulk metal, quantitative information was obtained
regarding the thickness, density (strictly speaking electron density) of the passive film, and the
roughness of the surface. The passive film has a lower density while the alloy surface layer has
a higher density then the bulk. The passive film density had a decreasing trend with increasing
potential, indicating dissolution or an increasing amount of vacancies in the oxide lattice. The
thickness of the passive film reached a maximum of 4 nm for 1200 mV, and then decreased
with further increase of the potential. The passive film density is similar to that of approximately
that of oxides of Cr and Fe.
The XRR results also showed a higher density of the alloy surface layer beneath the passive
film, in accordance with the enrichment of Ni (higher density). The density of the alloy surface
layer also increased with the applied potential. This increase in the density of this layer with
increasing potential indicates an increase of elements with higher densities such as Ni and Mo,
even though the XRR technique do not give direct chemical information. The presence of the
alloy surface layer, enriched in Ni, has been reported in literature [1, 18]. The increase of Ni
and Mo with the applied potential can be explained by preferential oxidation of Fe and Cr on
the surface, and subsequent dissolution of these elements.
5.1.2.2 Structural changes during anodic polarization
The structure of the passive film depends on the surface state of the steel and the exposure
conditions. The GIXRD is a surface sensitive technique, and the results show diffraction peaks
with low intensity in addition to the high intensity peaks originated from the two bulk phases,
as shown in Figure 23. These low intensity diffraction peaks were consistently observed when
measured in air, in electrolyte during an applied potential of 900 mV and after transpassive
dissolution in electrolyte. In Paper I the peaks matched the positions of several Cr-, Fe-, Cr-
Mo, Fe-Mo- and Fe-Mn- oxides, where Cr oxide was the major component giving the largest
diffraction peak. The number of peaks indicates a multiphase and mixed composition of a
nanocrystalline structure within an amorphous matrix.
40
Figure 23: GIXRD results showing lower intensity diffraction peaks from the surface oxides.
The measurements were done in air (black), in electrolyte with an applied potential of 900 mV
(red) and after termination of the anodic polarization up to 1400 mV (blue).
The crystallinity of the passive film decreased when the sample was exposed to the electrolyte
and under anodic polarization, proved by an increased background, as seen in Figure 24a. The
decreased crystallinity indicates an increase of amorphous compounds in the passive film.
Moreover, surface strain induced by grinding and polishing has influenced the structure of the
passive film, which was observed by comparing the results from the sample with Paper III and
without surface strain Paper I. The background signal at lower Bragg angles (2-Theta 10˚-20˚),
are significantly higher for the strained sample in Figure 24b than for the sample with “normal”
strain in Figure 24a, indicating a larger amorphous portion on the sample.
Figure 24: a) GIXRD results from a fine polished sample (no surface strain) showing the
overall diffraction pattern, and b) GIXRD results from the surface strained sample showing
larger background, indicating a higher amorphous part of the passive film.
41
Furthermore, the surface strain also affects the dissolution properties of the passive film. In
Paper (III), a few peaks could be indexed to Cr and Fe oxides and hydroxides, while the
background signal indicated the presence of an amorphous portion of the passive film. Upon
immersion in 0.1 M NaCl and anodic polarization, the Cr hydroxide peak disappeared, as shown
in Figure 25a, while the Fe hydroxide peak broadened, shown in Figure 25b.
Figure 25: XRD signals of a) Cr-hydroxide, which disappears upon immersion in the of 0.1 M
NaCl electrolyte, due to either dissolution or becoming amorphous, and b) Fe hydroxides, with
broadened peak after polarization up to 900 mV, indicating change of the passive film.
It has been shown that the solubility of amorphous Cr hydroxide is higher than for Fe
hydroxides [95]. However, no dissolved Cr species was detected during immersion in the
electrolyte and the anodic polarization, except in the transpassive state. Therefore, the vanishing
of Cr hydroxide peaks could also be caused by a change toward an amorphous structure. The
broadening of FeOOH peak indicated a strained crystal structure.
5.1.2.3 Dissolution and dealloying
The surface preparation resulted in different amorphous compounds in the passive film, which
influenced the dissolution properties. Such change in the dissolution properties was also
observed in another study on an aluminum alloy, where a mechanically altered surface showed
a different corrosion behavior compared to a polished sample [96]. For the unstrained sample,
the anodic metal dissolution from the bulk showed peak shifts of the bulk toward lower Bragg
angles, as shown in Figure 26a, indicating lattice expansion. The diffraction peaks of the bulk
phases became distorted and broader after polarization to higher anodic potentials. The
distortion was more significant for austenitic grains, indicating more induced strain in this
phase. The super duplex stainless steel accommodates a strain gradient between the two phases,
42
seen in Figure 26b, which is caused by the difference of mechanical strength, with the softer
austenite phase being more strained after mechanical treatments [97].
Figure 26: GIXRD results showing the shifts of the peaks toward smaller Bragg angles for a)
the normal strained sample and b) the surface-strained sample. The peaks show distortion,
which is more enhanced for the strained sample.
A peak shift was also observed in the signals originating from the passive film, showing an
opposite trend than those form the bulk, i.e., the oxide peaks shifted towards higher Bragg
angles, shown in Figure 27a-b.
Figure 27: GIXRD results show that the oxide peak shifts towards higher Bragg angles,
indicating a compressed oxide lattice that could be caused by an increased imperfection of the
crystals, possibly be due to an increased amount of vacancies.
43
The results from XRF measurement of metal dissolution in 1 M NaCl are summarized in Figure
28, which reveled preferential dissolution of Fe over the entire range of applied potentials.
Dissolution of Ni was detected at 1200 mV, while an enhanced dissolution of Cr and Mo was
detected at and above 1300 mV. During transpassive dissolution, Cr dissolution, was greatly
enhanced, which occurred at 1300/1400 mV.
Figure 28: The metal dissolution detected by XRF: a) elemental %; and b) the elemental
dissolution ratio, at different applied potentials.
44
5.1.3 Summary of in-situ/operando experiment results In summary, the breakdown of the protective oxide is associated with several processes. The
structural changes include loss of crystallinity and compression of the oxide lattice upon
increased anodic polarization of the sample. These structural changes can be caused by an
increased amount of vacancies. The underlying bulk alloy has an opposing trend, i.e., lattice
expansion, which can be caused by dissolution of small elements, such as Fe. This is supported
by the detected preferential dissolution of Fe.
At 1300 mV, Cr dissolution is enhanced, which is the definition of the transpassive dissolution.
At 1400 mV, the passive oxide layer vanished, and active dissolution of the bulk occurs, with
the ferrite being preferably dissolved. The surface strain leads to less crystallinity of the passive
film. The strain is preferably accumulated on austenite, and the dissolution of ferrite is retarded.
Moreover, after termination of anodic polarization at 1400 mV, a new passive film forms on
the surface, which is nanocrystalline but less crystalline than the air-formed oxide film.
The findings are schematically summarized in Figure 29, in which the enhanced dissolution
with increased potential is indicated by the broadening arrows. The figure also shows the
preferred dissolution of the ferrite (δ) phase. The oxide layer has an increasing concentration of
defects, with increasing potential. The alloy surface layer, with an enrichment of Ni and Mo,
has a similar thickness as the oxide film.
Figure 29: Illustration of the oxide layer, alloy surface layer and the bulk of the super duplex
stainless steel during increasing potential, showing different processes occurring in passive,
transpassive and active states. The thicker arrows indicate enhanced dissolution.
45
5.2 Local chemical composition and thickness of passive film
5.2.1 HAXPEEM experimental setup The HAXPEEM setup and measurement protocol were described in Paper IV, which also
showed an example for how to extract chemical information from individual grains to evaluate
the passive film thickness and composition. The detailed data analysis was reported in Paper V
focusing on lateral variation in the thickness and Cr content of the passive film, over the ferrite
and austenite phases, and the influence of grain orientation.
The six-step measurement protocol, shown in Figure 30, starts by marking the sample with
three fiducial Pt markers with a FIB (focused ion beam) in a SEM apparatus in DESY Nanolab
[98]. The markers were made with L shapes, of different sizes. The first was 60 μm × 40 μm
and the second was 20 μm × 10 μm. The third smaller marker of 2 μm × 1 μm was made close
to the region of interest (ROI), as shown in Figure 30a. The Pt markers enabled easy
identification of the ROI in the HAXPEEM microscope, and the ultra violet (UV) image, as
shown in Figure 30 b-c. A work function image was also made over the ROI to identify the
phases of the measured area.
The HAXPEEM measurement was performed using X-rays of an energy of 4 keV and the
lateral resolution was ca. 1 × 1 μm. The Fe 2p, Cr 2p, Ni 2p, O 1s and Mo 3p signals were
measured by XPS with an energy resolution of 0.2 eV. Mo 3p gave a very weak signal, which
was difficult for quantitative analysis. Surface adsorbed carbon was not detected due to low
intensity; therefore, the measured XPS peak of Pt was used for internal calibration of the energy
of the above core-levels. The sample was anodically polarized in steps, ex-situ at 600, 900 and
1000 mV in 1 M NaCl. The same ROI was measured before and after the electrochemical
polarization, and electron backscattering diffraction (EBSD) measurement was done to gain
information of crystallographic orientation of the grains of the two phases.
Using the work function image, it was possible to hand pick individual grains and extract
XPS spectra of the investigated elements. The fitting results of the XPS data allow the
calculation of the thickness and elemental concentration on chosen local and general scales.
46
Figure 30: A schematic of the measurement procedure showing a) the feducial Pt markers made
by FIB-SEM; b) a picture of one allocated marker in HAXPEEM by UV light; c) the PEEM
image over the ROI, showing different phases, reflected by their differenting workfunction; d)
measuring the surface for the binding energies of Cr, Fe, Ni and O; e) post measurement over
the ROI with EBSD, making it possible to obtain the phase and grain orientation information;
and f) XPS spectra that can be extracted on a local or global level.
In Paper V, based on the XPS data extracted from a selection of totally 58 grains, as shown in
Figure 31, the comparison was made for the two phases, and for three crystallographic
orientations. The aim was to investigate how the phase and crystallographic orientation
influence the thickness and composition of the native passive film.
47
Figure 31: Selected grains from the measured surface. a) 36 ferrite grains and 22 austentite
grains were selected; and b) corresponding grain orientations parallell to the sample surface.
5.2.2 Thickness and Cr content of passive film and lateral variations
The thickness and chemical composition of passive films has previously been investigated for
multi-phase systems. In these investigations, the method used required interference with the
material. For duplex stainless steel, the experimental approaches have included preferential
etching of one phase [99], local sputtering and depth profiling using AES [100] or by production
of corresponding single phases of the stainless steel [36]. In our HAXPEEM experiment, the
protocol minimized the manipulation of the material, and the Pt marker allowed measurement
of the same area by different techniques, e.g., HAXPEEM and EBSD to establish one-to-one
correlation between the analyzed site with the microscopic features (Papers IV-V).
5.2.2.1 Composition ratios of single grains
Paper (IV) demonstrated that quantitative analysis of XPS spectra and comparable estimations
could be done on two individual grains before and after anodic polarization to 1 V in 1 M NaCl,
as shown in Figure 32. The Cr spectra for the native oxide was fitted with three components
assigned to Cr metal, Cr3+ oxide and Cr3+ hydroxide, respectively. After the polarization, the
hydroxide peak disappeared while an increase of the Cr oxide peak was observed. The Fe
spectra for the native oxide could be fitted using two components, Fe metal and Fe2+ oxide.
After the polarization, a third peak of state Fe3+ component could be ascertained. Deconvolution
of Fe spectra is complicated because oxidized Fe can exist in many forms, and the spectra can
also contain satellite peaks, which were neglected in this study.
XPS data enables the estimation of elemental concentrations of the analyzed material. The
calculated the (Cr oxide)/(Cr metal) ratio for the individual grains suggest that the native passive
film of ferrite (001) grain contains a higher amount of Cr oxide than comparable to the austenite
(001) grain. After the polarization, the Cr oxide content increased for both phases and the
difference was less pronounced. The Cr hydroxide signal was not detectable after polarization.
Fe3+ appeared after polarization and was evenly distributed on both grains.
48
Figure 32:The XPS spectra of Cr 2p3/2 from a (001) grain from austenite (blue) and ferrite (red)
phases, respectively, before and after polarization.
49
It should be noted that, the energy resolution of HAXPEEM is limited and XPS signals of
individual grains (micrometer in size) are relatively weak, resulting in a low signal to noise
ratio, so in some cases it is difficult to obtain precise values from the spectra fitting.
Nevertheless, the study showed that it is possible to investigate the lateral variation of the
passive film non-intrusively, without modifying the surface by for example etching.
5.2.2.2 Lateral variation in thickness and Cr content of native passive film
Paper (V) reports the lateral variations in thickness and Cr content of the native passive film on
the super duplex stainless steel, regarding the difference between the ferrite and austenite
phases, and the influence of grain orientation grouped as (001), (101) and (111). Based on the
fitting results of individual grains, and the summed spectra of the specific groups of grains, the
thickness and Cr content of the passive film were calculated for all individual grains, and also
for each phase and each grain direction, utilizing equations (18) and (19). A two-layer model
of the passive film, shown in Figure 33, was used for the quantitative analysis. In the model,
the inner layer contains Cr3+ and Fe2+ oxides, and the outer layer contains Cr3+ hydroxide and
Fe3+ oxyhydroxide.
Figure 33: The model used to calculate the thickness and Cr content of the native passive film.
The spectra fitting was made from sum spectra from all included grains, grains from each phase
and from each grain direction. Individual grain spectra were also fitted, which had a lower
signal to noise ratio than the sum spectra shown in Figure 34, to gain an error estimation of the
fitting.
50
Figure 34: XPS spectra of Cr 2p3/2 for a) sum of all grains; b) sum of austenite phase; c) sum
of all austenite grains of (101) direction; and d) one single austenite (101) grain.
The results are presented in two ways: 1) data from the summed spectra of all analyzed grains,
the summed spectra of each phase, and the summed spectra of each orientation; 2) data from
the average of individual grains of specific groups, giving standard deviations for each phase
and for each grain orientation.
The summed spectra of all analyzed grains gave a thickness of 2.1 nm and a Cr content for the
total film of ca. 80 at%. The summed spectra of each phase gave a lower thickness of 1.8 nm
for and 1.6 nm, and Cr content of 75 at% and 79 at%, for the austenite and the ferrite,
respectively. From the average values of individual for the two phases, taking into account the
scattering of the data, it could be seen that the thickness is similar for the two phases, but it is
likely that the Cr content of the whole film is higher on the ferrite than on the austenite, as
shown in the box plot in Figure 35, which is consistent with the observations in Paper IV. The
ferrite phase contains more Cr than the austenite phase due to elemental partitioning [101],
which can be seen in Table 1 for the material used in this study. This could explain the higher
Cr content in the passive film on the ferrite. No significant difference in the passive film
thickness was found between the two phases, which in agreement with Rahimi, et al. [19], but
contradicted to Gardin, et al. [36]. The later used samples of two single phase materials.
51
Figure 35: Box plot of Cr content data for the inner layer, outer layer and total film, following
the model in Figure 33, calculated for the austenite (blue) and ferrite (red) phases.
At the grain level, the results from the summed spectra show that the grain orientation has a
small but detectable influence of the passive film, in particular, the Cr content of the outer layer
is lower for ferrite (111) orientation compared to other ferrite grains Paper V. The data from
individual grains show relatively large standard deviations, probably due to weak signals in
some cases. Nevertheless, the same observation can be seen in the box plot in Figure 36, i.e.,
consistent between two ways of analysis. This grain orientation has also shown to be more
corrosion prone compared to the other orientations in ferritic steel [102]. From Figure 36, it is
also likely that austenite (111) grains has a higher Cr content in the outer layer compared to
other austenitic grains.
52
Figure 36: Cr content of the inner layer, outer layer, and average native passive film for each
grain orientation (001), (101) and (111), for the austenite (A) and ferrite (F) phases.
The results from summed spectra showed that the outer layer of the ferrite grains had a thinner
outer film, as shown in Table 2. This may explain the preferential dissolution observed for the
ferrite phase during corrosion of duplex stianless steels [27, 103].
Table 2: Thickness of the inner layer, outer layer and total native passive film for each grain
orientation of the austenite and ferrite phases.
Summed spectra Mean of individual grains
Inner
layer
Outer
layer Total
Inner
layer
Outer
layer Total
Austenite
(001) 1.3 1.1 2.4 1.2 (±0.4) 1.2 (±0.1) 2.4 (±0.4)
(101) 1.2 1.0 2.2 1.2 (±0.4) 0.8 (±0.3) 1.9 (±0.4)
(111) 1.2 1.0 2.2 1.4 (±0.3) 1.2 (±0.3) 2.6 (±0.4)
Ferrite
(001) 1.3 0.8 2.1 1.2 (±0.4) 0.9 (±0.1) 2.1 (±0.5)
(101) 1.2 0.7 1.9 1.3 (±0.3) 0.9 (±0.2) 2.2 (±0.3)
(111) 1.1 0.8 1.9 1.2 (±0.2) 0.8 (±0.2) 2.0 (±0.3)
53
The observed influence of the grain orientation cannot be explained by diffusion anisotropy
since cubic crystals (austenite and ferrite phases) do not have any diffusion anisotropy.
However, the influence of orientation on oxidation and oxide growth has been observed for Fe
crystals [104]. In this work, it is possible that the grain orientation affects the surface reactivity
and thereby leads the observed differences. To explain the observed differences in the native
passive film between the two phases and the influence of grain orientation needs a deeper
understanding of the film growth mechanism of the system. This oxide film formation at
ambient temperature has been described by the high-field theory [1, 34], and passive film
formation and breakdown in aqueous electrolytes has been described by the point defect model
[105], for pure metals and simple alloy systems. The duplex stainless steel contains many
different elements and two phases and thereby is a very complex system. Many different aspects
must be considered, and further research efforts are needed to reach a fundamental
understanding of the passive film formation and breakdown. The experimental works in this
thesis have shown possibilities to gain detailed knowledge of such complex systems.
54
6. Conclusions The work in this thesis has led to the following conclusions:
• Critical conditions causing degradation of the passive film on super duplex stainless
steel in near-neutral NaCl solutions:
The passive film degradation under anodic polarization in 1 M NaCl stretches over a
potential range of more than 200 mV. Enhanced dissolution of Cr occurs at and above
1300 mV, leading to passive film breakdown.
The surface strain induced by mechanical grinding and polishing results in a higher
degree of amorphousness and affects the degradation of the passive film. The strain
leads to enhanced metal dissolution of the austenite phase.
• Atomic changes within the passive film leading to breakdown:
The passive film breakdown is a result of several chemical compositional and structural
changes of the film and the underlying alloy surface layer caused by the increasing
anodic potential. At applied potentials between 1000-1200 mV, a preferential
dissolution of Fe occurs, more pronounced on ferrite than on austenite. Dealloying,
generates defects, such as vacancies, leading to the strain in the surface region. At these
potentials, the passive film thickens but the density decreases due to the dissolution of
Fe. The density of the underlying alloy surface layer also increases as a result of
enrichment of heavier elements (Ni and Mo) in this layer.
Air-formed and aged passive film has a nanocrystalline structure. Under anodic
polarization the passive film became more amorphous. The passivated film formed in
the electrolyte is less crystalline than the film aged in air.
Repassivation of a strained sample surface results in a higher crystalline content of
the passive film than a non-strained sample, because the strain leads to enhanced
dissolution which in turn cause more dissolution-induced stress relaxation.
• The influence of bulk microstructure on passive film thickness and Cr content:
The native passive film has a bilayer structure, with an inner oxide layer and outer
oxyhydroxide layer. The oxide layer on the ferrite phase has a higher Cr content than
that on the austenite. The thickness of the whole film for the two phases does not
differentiate significantly, whereas the ferrite grains likely has a thinner outer layer
compared to the austenite grains. The grain orientation has a small but detectable
influence of the passive film. Most notably, ferrite (111) grains have a lower Cr content
in the outer layer of the passive film than other ferrite grains.
• The influence of bulk microstructure on passivity breakdown:
The ferrite phase is more susceptible to dissolution than the austenite phase, most likely
due to the differences (thickness and Cr content) in the passive film between the two
phases. A strained sample resulted in an enhanced degradation of the austenite phase,
due to that it is more susceptible to compressive strain, compared to a non-strained
sample.
55
7. Outlook and future work This thesis has provided more knowledge about the passivity degradation processes and related
structural changes of stainless steels. The study has also revealed lateral differences in the
passive film related to the microstructure. A deeper understanding of passivity and passive film
degradation can bring guidelines of how to further increase the corrosion resistance of stainless
steels. Development and utilization of the analytical techniques can disclose the complex
process of passivity breakdown of advanced alloys. Although some new insights have been
presented in this thesis, there are still several issues that are not fully understood.
• The compositional changes during passivity breakdown are important for the stability
of the passive film and thus the corrosion resistance. In this work, it was possible to
detect mainly the structural changes in real time during passivity breakdown, and the
metal dissolution in the electrolyte. However, information concerning ongoing chemical
changes in the passive film during anodic polarization is still to a large extent missing.
Electrochemical XPS measurements at synchrotron facilities could facilitate in-situ
studies of these changes, to gain real time information. In particular, the measurement
of detailed chemical changes of Cr, Fe and Mo oxides should be performed during
anodic oxidation and transpassive dissolution.
• The effect of the alloy surface layer underneath the passive film, if any, on passivity
breakdown is not understood fully yet. Ni enrichment in this layer has been reported,
and this study also suggests Mo enrichment for the super duplex stainless steel, but there
is no information about the structure and properties of this alloy surface layer. More
detailed analysis by surface sensitive XPS studies are needed to gain knowledge of
compositional changes of the alloy surface layer during passivity breakdown. The role
of this layer in the passivity breakdown needs to be further clarified.
• It has been debated whether Ni oxides are present in the passive film. To clarify this
question requires high-sensitivity XPS measurements and deep understanding of the
XPS spectra of Ni for the duplex stainless steel. This knowledge is crucial for the
understanding of synergism of the alloying elements in the passivity.
56
Acknowledgements A lot of gratitude goes to Professor Jinshan Pan who has supervised me through this PhD
journey. I would also like to thank my other supervisors: Professor Edvin Lundgren at Lund
University, Dr. Ulf Kivisäkk at Sandvik, Dr. David Lindell at Swerim, and Dr. Mari Sparr and
my closest coworkers in Professor Pan’s group; Dr. Cem Örnek, Dr. Fan Zhang, Dr. Min Liu
and Dr. Jie Cheng.
I’m also truly grateful for all the funding supports received from the Swedish Research Council
(Vetenskapsrådet) with project no. 2015-04490, the Röntgen-Ångström Cluster “In-situ High
Energy X-ray Diffraction from Electrochemical Interfaces” (HEXCHEM) with project no. 2015-
06092 and Swerim.
During my four years of PhD studies, I have worked alongside many highly intelligent people
in Professor Lundgren’s group who helped with the synchrotron experiments: Dr. Jonas
Evertsson, Dr. Gary Harlow, Dr. Lisa Rullik and Weronica Linpé in the division of Synchrotron
Radiation Research at Lund University
For their help during the beam times, a great thanks to the beam scientists: Dr. Francesco Carlá,
Dr. Roberto Felici, Dr. Carsten Wiemann, Dr. Andrei Gloskovskii, Dr. Yuri Matveyev and Dr.
Christof Schlueter. Also, I would like to thank Professor Andreas Stierle, Dr. Thomas Keller,
Dr. Elin Grånäs, Satishkumar Kulkarni, Dr. Heshmat Noei, in the Desy Nano Lab.
To all my coworkers at Swerim and Swerea KIMAB: Nuria Fuertes, Sara Munktell, Clara
Linder, Emil Stålnacke, Hanna Nilsson Åhman, Shirin Nouhi, Leyla Wickström, Konstantin
Simonov, Jesper Flyg, Miroslava Sedlakova and Sarahe Göthelid.
And of course, my office mates: Gen Lee, Weije Zhao, Seiya Watanabe, and my division and
PhD friends: Tingru Chang, Georgia Pilkington, Peter Dömstedt, Erik Bergendal, Adrian
Stoher, Sanghamitra Segupta, Anna Oleshkevych, Patricia Pedraz and Natalia Wojas. I also
extend my gratitude to all the other people I crossed paths with during my time as a PhD study:
Laetitia, Maria, Sara, Zahra, Akanskha, Sulena, Dima, Krishnan, Illia, Amanda, Xiujing,
Nanyan and Chung.
And to the department seniors: Mark Rutland, Inger Odervall Wallinder, Magnus Johansson,
Robert Corkery, Per Claesson, Mattew Fielden, Yolanda Hedberg, Jonas Hedberg, Eva
Blomberg, Gunilla Hurtig, Bruce Lyne, Eric Tyrode, Christoffer Leygraf, Rachel Pettersson,
Peter Szakalos and Mats Lundberg.
Hopefully, I have spelled your name correctly. If not, I hope you forgive me!
An extra thanks to my friends Georgia Pilkington and David Elhammer who spent time to
correct my linguistic errors in this thesis. Figure 4 and Table 1 in this thesis were provided by
M. Sc. Ulrika Borggren at Sandvik AB, which I’m of course also are grateful for.
Last but not least, I’d like to thank my family and friends who have been there for me during
these tough years and passing through the personal states of Year 1: “I don’t know anything!”,
Year 2: “I maybe know something” Year 3: “I know something” and this final year: “What do
I really know?”.
57
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