High Temperature Oxidation and Electrochemical
Studies on Novel Co-Base Superalloys
(Hochtemperaturoxidation und elektrochemische
Untersuchungen an neuartigen Co-BasisSuperlegierungen)
Der Technischen Fakultät der
Universität Erlangen-Nürnberg
zur Erlangung des Grades
DOKTOR-INGENIEUR
vorgelegt von
Dipl.-Ing. Leonhard Klein
Erlangen, 2013
Als Dissertation genehmigt von
der Technischen Fakultät der
Universität Erlangen-Nürnberg
Tag der Einreichung:
05.12.2012
Tag der Promotion:
27.02.2013
Dekanin:
Prof. Dr.-Ing. M. Merklein
Berichterstatter:
Prof. Dr. sc. techn. S. Virtanen
Prof. Dr.-Ing. U. Glatzel
Parts of this work were already published or submitted to Corrosion Science and
Electrochimica Acta:
L. Klein, A. Bauer, S. Neumeier, M. Göken, S. Virtanen, High temperature oxidation of
γ/γ'-strengthened Co-base superalloys, Corrosion Science 53 (2011), 2027-2034.
L. Klein, Y. Shen, M.S. Killian, S. Virtanen, Effect of B and Cr on the high temperature
oxidation behaviour of novel γ/γ'-strengthened Co-base superalloys, Corrosion Science 53
(2011), 2713-2720.
L. Klein, S. Virtanen, Electrochemical characterisation of novel γ/γ'-strengthened Cobase superalloys, Electrochimica Acta 76 (2012), 275-281.
L. Klein, S. Virtanen, Corrosion properties of novel γ'–strengthened Co-base superalloys,
Corrosion Science 66 (2013), 233-241.
L. Klein, M.S. Killian, S. Virtanen, The effect of nickel and silicon addition on some
oxidation properties of Co-based high temperature alloys, Corrosion Science (2013),
accepted manuscript.
Table of contents
Abstract …………………………………………………………………………………..1
Zusammenfassung ………………………………………………….…………..………..3
1. Introduction …………………………..…………………………………………….....5
2. Theoretical background ………………...………………..……………..…….…..…..9
2.1 High temperature materials ……………………………………………………...…9
2.2 Superalloys ……………………………………………………………………......11
2.2.1 Historical development …………..………………………………………....11
2.2.2 Microstructure of γ'-strengthened Co- and Ni-base superalloys …...….........13
2.3 High temperature oxidation in air ..………………...……………………...…...…15
2.3.1 General aspects …………………………..……………………...………….15
2.3.2 Oxidation of pure metals .…....…………..……………………...………….16
2.3.3 Influence of different alloying elements …..……………..……..………….21
2.4 Aqueous corrosion ………………………………………...………………….......23
2.4.1 Significance of electrochemical investigations for superalloys …………....23
2.4.2 Thermodynamic fundamentals ……...…...….……...……....……………...24
2.4.3 Reaction kinetics ………………………...………………………....……....28
2.4.4 Passivity and localised corrosion ...……….….……..………………...……33
2.5 Scientific aims of the present work ………………………………………………40
3. Experimental ……………………………...……………………..………………...…42
3.1 Investigated materials ……………………………………………………………..42
3.2 Heat treatment ………………...………………………………………...……..….43
3.3 Sample preparation ....…………………………………………………..……...….43
3.4 High temperature oxidation ……………………………………………..…….......44
3.5 Characterisation methods ………………………………………………………....46
3.5.1 Electrochemical investigations …….………………….………………....…46
3.5.2 Scanning electron microscopy ……….……………….…………….…...….47
3.5.3 Electron channelling contrast imaging ……………………………..………48
3.5.4 Energy dispersive X-ray spectroscopy ……………….……………....…….48
3.5.5 Optical microscopy ……………………………………………….……..…48
3.5.6 Time-of-flight secondary ion mass spectrometry …….……………...…….48
3.5.7 X-ray diffraction ……….……………………………………………...…...49
4. Results …..…………………………………………………………………………....50
4.1 High temperature oxidation of a quaternary Co–Al–W–B superalloy …..….....…50
4.1.1 Long-term oxidation behaviour …………………………………...…..…....50
4.1.2 Microstructure and composition of the formed oxide layers ……….....…...51
4.1.3 Influence of oxidation time and temperature on oxide layer formation …....54
4.2 Influence of different alloying elements on high temperature properties …….….57
4.2.1 Influence of the boron content ……....…………………....…………....…..57
4.2.2 Effect of chromium addition ……….....………………………………...…60
4.2.3 Effect of nickel addition ……………..…...…………………………....…..64
4.2.4 Effect of silicon addition ……………..…...……………………...…….….66
4.2.5 Effect of yttrium addition …………………...…………………...…….…..69
4.2.6 Effect of titanium addition …………………..……………………….…....71
4.3 Electrochemical studies on the quaternary Co–Al–W–B superalloy …….......…..74
4.3.1 Corrosion of the unoxidised alloy in aqueous solutions of different pH ….74
4.3.2 Corrosion of the oxidised alloy in aqueous solutions of different pH .........78
4.3.3 Corrosion of unoxidised alloys in chloride containing solution ……....…..82
4.3.4 Corrosion of oxidised alloys in chloride containing solution ….…….........84
5. Discussion ………………...…………………………...………………………….......88
5.1 Oxidation properties of Co-base superalloys ………………………...…...……....88
5.2 Influence of alloying elements ……………………………………...…………....93
5.2.1 Effect of boron addition …………………………………………………….93
5.2.2 Effect of chromium addition ………………………………………………..95
5.2.3 Effect of yttrium addition …………………………………………..……..100
5.2.4 Effect of silicon and nickel addition ………………………………..……..100
5.2.5 Effect of titanium addition ………...……………………………………....102
5.2.6 Summary of the alloying effects ………………………………………..…103
5.3 Corrosion protection of high temperature oxides in aqueous solutions …….…..105
5.3.1 Corrosion properties in aqueous solutions of different pH …………….…105
5.3.2 Corrosion properties in chloride containing solution ……………………..110
5.4 Comparison of high temperature oxidation and aqueous corrosion …………….112
6. Conclusion and outlook ….....……………………………...…………………..…..114
References ………..………………..…………………………………………………..119
Acknowledgement ………..…………………………………………………………...127
Abstract
1
Abstract
Isothermal oxidation in air was carried out on novel γ'-strengthened Cobalt-base
superalloys of the system Co–Al–W–B. After fast initial oxide formation, a multi-layered
structure establishes, consisting of an outer cobalt oxide layer, a middle spinel-containing
layer, and an inner Al2O3-rich region. Ion diffusion in outward direction is hindered by
the development of Al2O3, that can be either present as a continuous and protective layer
or as a discontinuous Al2O3-rich area without comparable protective effect. Furthermore,
high temperature oxidation leads to phase transformation (from γ/γ' into γ/Co3W) at the
alloy/oxide layer interface due to aluminium depletion.
Pure cobalt and ternary Co–Al–W alloys exhibit parabolic oxide growth due to the
lack or insufficient amounts of protective oxides, whereas quaternary Co–Al–W–B alloys
possess sub-parabolic oxidation behaviour (at 900 °C). At lower temperatures (800 °C),
even a blockage of further oxidation can be observed. High amounts of B (0.12 at%)
significantly improve oxidation resistance mainly due to its beneficial effect on inner
Al2O3-formation at the alloy/oxide interface. Furthermore, B prevents decohesion of high
temperature scales due to the formation of B-rich phases (presumably tungsten borides) in
the middle oxide layer. Appropriate amounts of chromium (8 at%) as additional alloying
element to Co–Al–W–B alloys lead to the formation of an inner duplex layer composed
of protective Cr2O3 and Al2O3 phases. In this respect, chromium also benefits selective
oxidation of aluminium, which results in higher Al2O3-contents compared to chromiumfree alloys. Major drawbacks of chromium additions are, on the one hand, the formation
of volatile chromium-containing species at temperatures exceeding 1000 °C and on the
other hand, the instability of the γ/γ'-microstructure. Titanium and silicon additions lead
to improved oxidation resistance due to their beneficial effect on Al2O3 formation
(especially at 900 °C and higher) and due to additional generation of titanium- and
silicon-rich phases, respectively, without altering the γ/γ'-microstructure. Moreover, the
titanium-containing alloy is reported to exhibit excellent creep properties at 850 °C and
hence this material is expected to be the most promising alloy system for further
optimisation. In contrast, additions of silicon lead to silicon-containing phases at the
oxide/alloy interface, within precipitates, and at the grain boundaries, which are expected
2
Abstract
to impair the mechanical properties. Additions of nickel most probably enhance solubility
of boron within the alloy matrix and therefore the previously described positive boroneffect gets eliminated. Based on knowledge of other alloy systems, small amounts of the
rare earth element yttrium are reported to improve the oxidation resistance. However,
0.005 atomic percent of yttrium in Co–Al–W–B alloys do not lead to the expected effect,
presumably due to insufficient amounts of the minor element.
Electrochemical measurements on the unoxidised Co–Al–W–B superalloy and
pure cobalt in aqueous solutions of different pH reveal significantly improved corrosion
resistance with increasing pH value due to the formation of a duplex layer, i.e. Co3O4 or
CoOOH species on top of a Co(OH)2 film. Upon polarisation, both materials show
primary and secondary passivation in alkaline 0.1 M NaOH solution, whereas limited
passivation can be observed in neutral 1 M Na2SO4, and active dissolution in acidic
0.5 M H2SO4 solution. Further investigations in neutral 0.5 M NaCl solution reveal
limited initial passivation followed by severe pitting corrosion at higher potentials.
High temperature oxide scales on the alloy surface are highly efficient barriers
against corrosive attack over the entire polarisation range from -1 V to +2 V (vs.
Ag/AgCl), indicating remarkable retardation of cathodic and anodic reactions. The best
protective properties of high temperature oxides can be achieved in acidic solution. Even
for thin oxide layers formed during short-term oxidation of only 1 hour at 900 °C, no
breakdown events can be observed in chloride containing solution. The addition of silicon
to the Co–Al–W–B alloy considerably increases corrosion resistance of the oxidised
specimen in 0.5 M NaCl aqueous solutions due to the formation of protective Al2O3 and
silicon-rich species. This demonstrates that not only the thickness of high temperature
scales, obtained by longer oxidation times and temperatures, but also their detailed nature,
such as chemical composition or oxide stabilities in solutions of different pH, dictates the
corrosion performance.
Zusammenfassung
3
Zusammenfassung
In der vorliegenden Arbeit wurden neuartige γ'-gehärtete Kobalt-Basis Superlegierungen
des Systems Co–Al–W–B in Luftatmosphäre isotherm oxidiert. Die Anfangsphase ist
geprägt durch schnelle Oxidation der Legierungsbestandteile, die im weiteren Verlauf ein
Mehrschichtsystem ausbilden, bestehend aus einer äußeren Kobaltoxidschicht, einem
mittleren spinellhaltigen Bereich und einer inneren Al2O3-reichen Zone. Das Al2O3 kann
entweder als durchgängige innere Schicht vorliegen, die weitere Oxidation verhindert,
oder diskontinuierlicher Natur sein und somit keinen vergleichbaren Schutz gegen
Ionendiffusion
in
auswärtige
Hochtemperaturoxidation
Richtung
aufgrund
bieten.
der
Des
Weiteren
Aluminiumverarmung
führt
zu
die
einer
Phasentransformation an der Grenzfläche zwischen Oxid und Legierung. Hierbei wandelt
sich das γ/γ' Gefüge zu einem zweiphasigen γ/Co3W Gebiet um.
Reines Kobalt und ternäre Co–Al–W Legierungen zeigen parabolisches
Oxidschichtwachstum, wohingegen bei quaternären Co–Al–W–B Legierungen bei 900 °C
subparabolisches Verhalten zu beobachten ist und nach ca. 100 Stunden bei 800 °C keine
weitere Oxidation stattfindet. Ein hoher Borgehalt (0.12 Atomprozent) wirkt sich hierbei,
hauptsächlich
aufgrund
begünstigter
innerer
Al2O3-Bildung,
positiv
auf
das
Oxidationsverhalten aus. Des Weiteren führt das Vorhandensein borhaltiger Phasen
(vermutlich Wolframboride) in der mittleren Oxidschicht zu einer beträchtlichen
Verbesserung der Schichthaftung. Durch das Zulegieren von Chrom (8 Atomprozent)
kann die Bildung einer inneren Doppelschicht erreicht werden, bestehend aus
schützendem Cr2O3 and Al2O3. Hierbei wirkt sich Chrom vorteilhaft auf die selektive
Oxidation von Aluminium aus, was im Vergleich zu den chromfreien Legierungen höhere
Gehalte an innerem Al2O3 zur Folge hat. Als Hauptnachteile der Chromzugabe können
zum einen die Bildung flüchtiger chromhaltiger Oxide bei Temperaturen oberhalb von
1000 °C, zum anderen eine resultierende Instabilität des γ/γ' Gefüges angesehen werden.
Durch das Zulegieren von Titan und Silizium wird das Oxidationsverhalten ebenfalls
aufgrund begünstigter innerer Al2O3-Bildung (insbesondere bei Temperaturen oberhalb
von 900 °C) und der Ausscheidung titan- und siliziumreicher Phasen verbessert, jedoch
bleibt das γ/γ' Gefüge stabil. Zudem wurden für titanhaltige Legierungen bereits
4
Zusammenfassung
ausgezeichnete Kriecheigenschaften bei 850 °C festgestellt, weshalb diese Materialien
durch weitere Optimierungen vielversprechende Ergebnisse erwarten lassen. Im
Gegensatz dazu werden für siliziumhaltige Legierungen verschlechterte mechanische
Eigenschaften berichtet. Des Weiteren wird angenommen, dass die Zugabe von Ni zu
Co–Al–W–B Legierungen zu einer erhöhten Löslichkeit des Bors im Korninneren führt
und somit der beschriebene Bor-Effekt wieder aufgehoben wird. Geringe Mengen
(0.005 Atomprozent) an Metallen der seltenen Erden, wie zum Beispiel Yttrium,
verbessern die Oxidationsbeständigkeit der untersuchten Legierungen nicht, was aber
vermutlich auf den zu geringen Gehalt an Yttrium zurückzuführen ist.
Elektrochemische
Messungen
an
unoxidierten
quaternären
Co–Al–W–B
Legierungen und reinem Kobalt lassen eine verbesserte Korrosionsbeständigkeit mit
zunehmendem pH-Wert der wässrigen Lösung erkennen. Dies ist durch die Bildung einer
Doppelschicht zu erklären, bestehend aus einer Co(OH)2 Passivschicht, die im weiteren
Verlauf der Polarisation durch Abscheidung von Co3O4 bzw. CoOOH erweitert wird.
Sowohl die Legierung als auch reines Kobalt zeigen in alkalischen Lösungen diese
primäre und sekundäre Passivierung, wohingegen in neutralem Medium lediglich ein
limitierter Passivbereich und in sauren Lösungen aktive Materialauflösung beobachtet
werden kann. Weitere Untersuchungen in neutraler chloridhaltiger Lösung zeigen eine
anfängliche Passivierung der Materialien, gefolgt von großflächiger Lochkorrosion bei
Überschreitung eines kritischen Potenzials.
Oxidierte Legierungen besitzen aufgrund der schützenden Oxidschicht vor allem
in
sauren
Lösungen
eine
ausgezeichnete
Korrosionsbeständigkeit,
wobei
die
Barriereschichten kathodische und anodische Oberflächenreaktionen in einem weiten
Polarisationsbereich (-1 V bis +2 V vs. Ag/AgCl) außerordentlich stark hemmen. Selbst
dünne Oxidschichten (nach nur einstündiger Oxidation bei 900 °C) zeigen kein Versagen
in chloridhaltigen Lösungen. Durch Zugabe von Silizium ist es möglich die
Korrosionsbeständigkeit der oxidierten Legierungen in Natriumchloridlösung sogar
weiter zu steigern. Dies zeigt, dass nicht nur die Dicke, sondern auch die chemische
Zusammensetzung der Schicht und die entsprechenden Oxidstabilitäten in Lösungen
unterschiedlichen pH-Wertes und Chloridgehaltes die Korrosionseigenschaften diktieren.
Introduction
5
1. Introduction
In 2012, the world population exceeded the threshold of seven billion people for the first
time. A continuously growing population will result in a higher demand for primary
energy, at which fossil fuels will remain the most important energy source. The ethics
commission for a safe energy supply highlighted in May 2011 that more efficient supply
technologies demand innovative materials, for instance for wind turbines or high
temperature power plants. Therefore, the strengthened expansion of research was
recommended, since it can provide urgently needed high-performance materials for
energy systems [1].
In the field of high temperature materials, Ni-base superalloys are commonly
applied, for instance for turbine vanes and discs, due to their excellent mechanical
properties (creep strength, ductility) and good oxidation resistance at high
temperatures [2]. A proper choice of alloying elements enables the required oxidation
resistance whereas high temperature strength results from a stable ternary Ni3Al
compound with the L12 structure, which precipitates as a cuboidal intermetallic γ'-phase
in the fcc γ-Ni solid solution [3, 4]. In 2006, a novel γ'-strengthened Co–Al–W superalloy
was developed showing a high temperature strength comparable to Ni-base
superalloys [5]. Strengthening was achieved by a Co3(Al,W) compound with the L12
structure, comparable to the Ni3Al-phase in Ni-base superalloys. The addition of boron to
polycrystalline alloys is reported to improve the mechanical properties such as creep
strength [6], rupture strength, and ductility [3, 7]. Recent publications [8, 9] confirm that
enhanced creep properties can be achieved for novel B-containing Co-base superalloys
due to grain boundary strengthening, as well. Therefore, this new class of high
temperature material has the potential to compete with the conventional Ni-base
superalloys as load bearing material for high temperature applications. In addition, Co has
a higher melting point than Ni (1495 °C vs. 1455 °C) which may further increase the
process temperature of the turbines. However, the high temperature oxidation behaviour
of recently developed Co-base superalloys is still relatively poor and therefore has to be
improved considerably by suitable alloying strategies, since lifetime of high temperature
materials is significantly decreased by oxidation, as it impairs the mechanical properties.
6
Introduction
Typically, a thermal barrier coating (TBC) is deposited on the superalloy before service.
However, it is important to possess an oxidation resistant substrate material, as well.
Compared to Ni-base superalloys [10-14], only very little research has been carried out
concerning oxidation behaviour of new γ'-strengthened Co-base superalloys.
Aluminium is reported to be the most important alloying element for high
temperature alloys since it may form a protective Al2O3 layer on the alloy surface, which
is characterised by a slow growing rate, high stability, and a compact structure with very
few defects [10]. Therefore, transport of species through the oxide layer is highly
inhibited resulting in the superior protective properties of Al2O3. Small amounts of rare
earth elements, such as yttrium, enhance adhesion of the oxide layer, especially under
cyclic conditions, and increase the oxidation resistance of the base alloy due to selective
Al oxidation [15-21]. Other important elements are chromium and silicon, which form
protective Cr2O3 and SiO2, respectively. In addition, Cr forms stable and non-liquid Crcontaining sulphides, which improve hot corrosion properties [22]. However, protection
of Cr2O3 is reduced at temperatures higher than 1000 °C and in oxidising atmospheres
with high flow rates due to the formation of volatile CrO3 species [23-25]. In contrast,
silicon enhances the oxidation resistance at elevated temperatures (T ≥ 900 °C), not only
because of SiO2 formation [26, 27] but also due to the beneficial effect of silicon on the
formation of protective Al2O3 [28].
In addition to their high temperature oxidation properties, the general corrosion
behaviour of novel γ'-strengthened Co-base superalloys is of interest for their application
in turbines. Electrochemistry is a typical tool for investigating the corrosion behaviour of
materials at ambient temperature. The reaction process takes place in an electrolyte which
is usually an aqueous medium. In contrast to Ni-base superalloys [29-33], conventional
Co alloys [34-39], and pure Co [40-48], very little electrochemical studies have been
carried out for new γ'-strengthened Co-base superalloys. Those alloys, in addition to their
special microstructure, show different chemical compositions than the conventional highcorrosion resistant Co-base superalloys and the binary Co-containing alloys. Therefore, a
direct prediction of the corrosion properties of these novel alloys, based on the literature,
is not possible. The electrochemical examination at ambient temperature is an important
issue which has to be addressed in addition to typical hot corrosion testing, since turbine
Introduction
7
vanes, for example in aircraft engines, may suffer from corrosion in aggressive media
(e.g. sulphate or chloride containing media of different pH values) after operation, as
well. In this case, high temperature scales are present on the oxidised alloy surface, which
leads to an altered corrosion behaviour of the material. Sulphates are typical reaction
products of the combustion process and the saline marine air leads to the presence of
aggressive chlorides, which may attack the surface of the turbine vanes.
The goal of the present PhD study was to systematically investigate the oxidation
properties of novel γ'-strengthened Co-base superalloys at high temperatures, since very
little research has been carried out in this field so far [49]. For this, different alloying
elements (B, Cr, Ni, Si, Y, and Ti) were added to the Co–Al–W ternary alloy in varying
amounts and isothermal oxidation in air was carried out at 800, 900, and 1000 °C with
exposure times up to 500 h. Additions of Ni, B, and Ti are usually known for adjusting
the microstructure and improving the mechanical properties of superalloys [3, 6-9, 50].
However, their influence on the high temperature properties of the novel Co-base
superalloys was of particular interest for the present study. Several analytical techniques
were applied, such as thermogravimetry, scanning electron microscopy (SEM), electron
channelling contrast imaging (ECCI), optical microscopy, energy dispersive X-ray
spectroscopy (EDX), time-of-flight secondary ion mass spectrometry (ToF-SIMS), and
X-ray diffraction (XRD), in order to characterise the oxidation process, the microstructure
of the alloys and oxide layers, as well as the elemental distribution within the oxide layers
and at the grain boundaries. Furthermore, electrochemical experiments were carried out at
ambient temperature on bare and oxidised alloys in several aqueous solutions of different
pH and chloride containing solutions. For this, acidic 0.5 M H2SO4, neutral 1 M Na2SO4,
neutral 0.5 M NaCl, and alkaline 0.1 M NaOH solutions were chosen as electrolytes.
Obtained results were compared to the behaviour of pure Co. Measurements were
performed in order to investigate the protective effect of the high temperature oxide
scales on the metal surface, in comparison with the passive layers on non-oxidised
samples.
Based on the potential of novel Co-base superalloys and on current research
progress, including the findings of the present work, the development of sophisticated γ'strengthened alloy systems in the foreseeable future may result in novel high temperature
8
Introduction
materials, which possess superior mechanical and corrosive properties compared to
common γ'-strengthened Ni-base superalloys.
Theoretical background
9
2. Theoretical background
2.1 High temperature materials
Bürgel et al. presented a detailed overview of high temperature materials, their properties
and the resulting application fields, which is analogously summarised in the
following [22]. High temperature materials, such as ceramics, certain metals, and
intermetallic compounds, are usually used above approximately 500 °C and have to
exhibit sufficient corrosion resistance and mechanical properties in order to be suitable
for high temperature components. Materials can be used in the field of power
engineering, drive engineering, metallurgy, or in the chemical industry. Examples for
such applications may be steam and gas turbines, nuclear power plant reactors, jet
turbines, components for raw materials production and processing, and components used
for high temperature pyrolysis of chemical compounds.
High temperature materials have to withstand a combination of corrosive attack,
thermal stress, and mechanical loading. As a result, such materials have to meet a variety
of demands, such as long-term structural stability at high temperatures, sufficient
mechanical properties (creep and fatigue strength, ductility), machinability, and high
temperature corrosion resistance. In this respect, the formation of well adhesive and
protective surface layers on non-ceramic high temperature materials is aspired. As for
ceramic materials, application fields are relatively limited due to their brittle character
and low defect tolerance. For instance, ceramics are unsuitable for most structural
applications [51].
Base elements for high temperature alloys have to exhibit melting points above
approximately 1400 °C in order to ensure sufficient long-term creep strength. Several
metals, such as Ni, Co, Fe, Ti, or Pt meet this requirement and therefore are suitable as
base elements. Nb, Ta, Mo, and W possess very high melting points, as well. However,
poor oxidation resistance is one major drawback for the use of refractory pure metals and
alloys as high temperature materials. Ti-alloys reveal low creep strength, embrittlement,
and decreased oxidation resistance at higher temperatures, which limits their application
range below 600 °C. Recently developed intermetallic compounds based on γ-TiAl and
10
Theoretical background
Ti3Al may be used at higher temperatures but still have to be optimised for industrial use.
So far, a multitude of intermetallic phases is known (e.g. Ni3Al, NiAl, MoSi2, or Co3Ti),
characterised by strong atomic bonds which lead to a high melting point and Young’s
modulus at high temperatures. However, creep rupture strength, ductility, and oxidation
resistance have to be improved by adequate alloying and hardening strategies. It is
expected that, with respect to their high temperature properties, intermetallic construction
materials may close the gap between classic high temperature alloys and ceramics.
Precipitation hardening in Ni-base superalloys is accomplished by means of high
volume fractions of the intermetallic Ni3Al phase, which enables excellent mechanical
properties at high temperatures, such as high creep strength. Even though the melting
point of Ni is the lowest compared to Co and Fe, Ni-base superalloys represent the most
widely applied type of superalloys, since the best combination of mechanical properties,
machinability, corrosion resistance, and thermal properties can be achieved. However, the
composition is frequently optimised with respect to the mechanical properties but at the
expense of corrosion resistance or the other way around. Alloy properties can also be a
reasonable compromise, which takes costs and manufacturability into account [51]. High
temperature alloys based on precious metals (such as Pt, Ir, Ru, or Rh) show good high
temperature properties but are limited to special applications due to superior material
costs. Recently, novel γ'-strengthened Co-base superalloys were developed by Sato et al.,
which exhibit creep properties comparable to Ni-base superalloys. The addition of W to
the binary Co–Al system resulted in stabilisation of a strengthening γ' Co3(Al,W)
compound similar to the strengthening in Ni-base superalloys [5]. Therefore, this
promising novel class of high temperature material is currently investigated intensely. In
the following, γ'-strengthened Co- and Ni-base superalloys will be addressed in detail due
to their outstanding position as high temperature materials and the huge development
potential.
Theoretical background
11
2.2 Superalloys
2.2.1 Historical development
The driving force for superalloy development in the 20th century was the increasing
demand for novel materials which can withstand high process temperatures and
mechanical loads in gas turbine and jet engines [51]. In the 1920s, a binary Ni–20Cr alloy
(“Nichrome”; value in mass percent) was patented and provided a pathway for later
superalloys, such as Nimonics and Inconels [3, 52]. In the following, small amounts of Al
and Ti were added to Ni–20Cr alloys for further strength improvements by means of
relatively small amounts of γ' (Ni3Al) precipitation (Nimonic 80) [51]. At the beginning
of the last century, Co–Cr and Co–Cr–W alloys were developed for applications in
aggressive environments and as wear-resistant materials [3]. Furthermore, the Co-base
superalloy Stellite 21 was the first turbine vane material utilised in jet engines [51].
Between the 1930s and 1950s, the development of superalloys was based on
microstructural improvements, such as carbide strengthening or solid solution hardening
by means of refractory elements [3, 51]. Novel vacuum melting techniques developed in
the 1950s enabled an optimisation of the casting process and compositional control, and
therefore the material properties [53]. Especially the possibility of adding higher amounts
of Al and Ti, which do not oxidise in the vacuum atmosphere, lead to the development of
γ'-strengthened Ni-base superalloys [51]. Due to a decrease in Cr-content from 20 to
approximately 10 weight percent, coatings had to be applied in order to assure
environmental resistance of the alloys [51]. As a result, process temperatures up to
1200 °C could be achieved in turbine engines [54]. It has been reported in the 1970s, that
Co-base superalloys may form a γ' Co3Ti phase, as well. However, even though the effect
of various alloying elements on the morphology and stability of Co3Ti was studied, the
application of such alloys was restricted to temperatures up to approximately
750 °C [55, 56]. Nevertheless, higher melting temperatures, superior hot corrosion
resistance due to high Cr-contents, superior thermal fatigue resistance, better weldability,
and improved castability due to limited alloying amounts of Al and Ti (and therefore
lower vacuum demands) are some advantages of Co-base superalloys compared to Nibase superalloys. Those properties justify the use of Co-base superalloys in mechanically
12
Theoretical background
lower loaded parts of turbines where environmental resistance is required [3]. In the late
1960s, Pratt and Whitney introduced a new casting technique which enabled directional
solidification (DS) of alloys [57], characterised by a columnar crystal structure. In the late
1970s, the first single crystalline (SC) superalloy PWA-1480 was cast, exhibiting superior
creep properties and improved operation temperatures [57] due to an elimination of grain
boundaries. In Figure 2.1, cast turbine blades are illustrated in the (a) polycrystalline, (b)
directionally solidified, and (c) single-crystalline form according to [58].
Figure 2.1: Turbine blades in the (a) polycrystalline, (b) directionally solidified, and (c)
single-crystalline form [58].
In Figure 2.2, the previously described development of Ni-base superalloys is
illustrated since 1940 [4]. The application of novel casting technologies and a continuous
alloy development lead to a significant increase in process temperature (here temperature
for 1000 h creep life at 137 MPa). Current research is based on optimising the casting
process and material properties of SC superalloys, as well as the development of novel
Co-base superalloys, which can be strengthened by a stable γ'-phase similar to
conventional Ni-base superalloys.
Theoretical background
13
Figure 2.2: Development of Ni-base superalloys since 1940 [4].
2.2.2 Microstructure of γ'-strengthened Co-and Ni-base superalloys
Novel γ'-strengthened Co-base superalloys (see Figure 2.3) possess very similar γ/γ'microstructures compared to conventional Ni-base superalloys. The cuboidal γ'-phase
distributes homogeneously within the γ-matrix and the precipitates align along the [001]
directions [5].
Figure 2.3: γ/γ'-microstructure of a novel Co-base superalloy of the system Co–Al–W.
14
Theoretical background
Within the ternary Co–Al–W system, the stable intermetallic Co3(Al,W)
compound exhibits the L12 crystal structure (see Figure 2.4) similar to the Ni3Al phase in
Ni-base superalloys. Here, Al and W atoms are situated at the corners whereas Co atoms
are situated at the planes of the unit cell.
Figure 2.4: Unit cell of the fcc γ-Co (A1) matrix on the left hand side and the
γ' Co3(Al,W) phase (L12) on the right hand side.
Since the γ'-phase (L12) is embedded coherently in a face-centred cubic (fcc) γ-Co (A1)
matrix [5], the lattice misfit is relatively low. However, coherency strains are present at
the γ/γ'-interface, which are minimised by a cubic shape of the γ'-precipitates (see
Figure 2.3). Further reduction of the lattice misfit, e.g. by Ni-addition to the Co–Al–W
system, would lead to a spherical geometry of the γ'-precipitates in order to reduce the
interfacial energy [50]. Furthermore, Shinagawa et al. reported that at 900 °C, the twophase γ/γ'-area in the isothermal section diagram can be broadened with increasing Nicontent. As a result, Ni-addition to the Co–Al–W system may avoid an instability of the
γ'-phase when adding high amounts of further alloying elements.
Even though γ'-precipitates are embedded coherently in the γ-matrix, significant
hardening can be achieved because penetration of the ductile and ordered γ'-phase by
dislocations from the γ-matrix is hindered. For this reason, deformation mainly occurs in
the matrix channels [4]. Strengthening due to the γ'-phase is affected by its volume
fraction and size [59], but also by the presence of secondary fine dispersed γ'-precipitates
within the γ-matrix [3].
Theoretical background
15
B-containing polycrystalline Ni-base superalloys exhibit further improvements in
mechanical properties due to B-rich grain boundary precipitates which prevent grain
boundary decohesion [3, 6, 7]. As already mentioned in chapter one, this strengthening
effect can also be observed for novel Co-base superalloys. However, up to now only very
little research has been carried out in this field [8, 9]. In addition, no publications can be
found addressing the influence of B on high temperature oxidation behaviour of Co-base
alloys. Therefore, one main focus of this work is to investigate the effect of different
alloying elements, such as B, Cr, Ni, or Si on the oxidation properties of novel Co-base
superalloys in air.
2.3 High temperature oxidation in air
2.3.1 General aspects
High temperature alloys, e.g. used for turbine blades, have to withstand metal loss due to
oxide layer formation, since reduced thickness of the components, and therefore the loadbearing cross-sections, may lead to failure. For this reason, extensive efforts are made in
order to develop not only superalloys with outstanding mechanical but also excellent
oxidation properties for service temperatures exceeding 1100 °C [22].
With respect to metals, high temperature oxidation refers to the reaction of the
sample with oxygen from the surrounding air in order to form an oxide scale on the
surface. Bare metal surfaces react rapidly and establish an oxide layer, which may
decelerate further oxidation depending on the scale’s porosity, thickness, chemical
composition and adherence. If the oxide layer is cracking, it reforms spontaneously in air.
Higher oxidation temperatures result in thicker oxide layers due to higher reaction
kinetics. In general, high temperature oxidation of a metal is controlled by
thermodynamic and kinetic factors. A detailed examination of the oxidation process will
be presented in the following chapter.
16
Theoretical background
2.3.2 Oxidation of pure metals
In order to investigate the oxidation behaviour of complex alloy systems, such as Co- or
Ni-base superalloys, it is important to understand the oxidation mechanisms of rather
simple systems, i.e. pure metals. In the following, high temperature oxidation in air of
pure Co will be presented in detail. Oxidation mechanisms are similar to other pure
metals, such as Ni, Fe, or Cr.
Oxidation properties of pure Co have been investigated extensively in the past due
to the importance of the cobalt-oxygen system as an oxidation model [60-71]. Isothermal
oxidation at high temperatures and under constant pressure can only occur spontaneously
if Gibbs free energy (or “free enthalpy”) ∆G is negative. If ∆G = 0, a thermodynamic
equilibrium is present and if ∆G > 0, the oxidation reaction is thermodynamically not
possible, i.e. not spontaneous. At a constant temperature and pressure, ∆G can be
expressed by the following expression [72]:
∆G = ∆H - T·∆S
(2.1)
∆H is the change in enthalpy [J], T the absolute temperature [K] and ∆S the entropy
change [J/K]. Gibbs free energies ∆G0 for the formation of several important oxides are
given in the Ellingham-Richardson diagram in Figure 2.5 [73]. However, the diagram
does not take reaction kinetics into account. Metals plotted at the top of the diagram are
easier to reduce (noble metals) than those plotted lower on the diagram, which naturally
tend to exist in very stable oxide forms. The oxide layer forms only if the partial pressure
of the surrounding oxygen is higher than the dissociation pressure of the oxide in
equilibrium with the metal [74]. Hence, a metal (in this case Co) can be oxidised to CoO
under the following condition:
⎡ 2 ⋅ ∆G 0 (CoO ) ⎤
p o2 > exp ⎢−
⎥
RT
⎣
⎦
(2.2)
Theoretical background
17
R is the gas constant [J/(mol·K)] and T the temperature [K]. The dissociation pressure of
the oxide and the standard Gibbs free energy ∆G0 of oxide formation at a given
temperature can be read off the Ellingham-Richardson diagram in Figure 2.5.
Figure 2.5: Ellingham-Richardson diagram of some important oxides [73].
18
Theoretical background
Depending on the oxidation temperature, cobalt oxide can be present either in the
form of mainly Co3O4 below 900 °C or CoO above 900 °C [75]. It has been reported that
oxidation of pure Co shows a parabolic time dependency and the cobalt oxide single-layer
grows on the metal surface as a result of outward diffusion of cobalt ions through the
growing compact oxide scale [64-71]. Parabolic kinetics indicate that diffusion of the
reactant is rate determining [63]. This will be discussed in more detail in the present
chapter. Mass gain or oxide layer growth decelerates with increasing oxidation time
because outward diffusion of cobalt ions is hindered by a thick, compact, and adherent
oxide layer on the metal surface. The layer growth is proportional to the square root of
time, therefore the parabolic rate law can be written as follows:
x2 = 2k'·t
(2.3)
The oxidation time is represented by t. The thickness of the oxide layer at a specific
oxidation time t is expressed by x and the parabolic rate constant by k'. Since the
measured mass gain per unit area (∆m/A) is proportional to the oxide layer thickness, x
can be substituted by ∆m/A in equation (2.3). A graphic account of the equation shows
that the parabolic rate constant k' represents the slope of the straight line. In Figure 2.6,
rate constants of different important oxides on their corresponding metals are given as a
function of temperature [3, 76]. For instance, Al2O3 exhibits a lower rate constant
compared to Cr2O3 at a given temperature and therefore thinner oxide layers form on the
metal surface. However, CoO on pure Co shows a relatively high rate constant compared
to other oxides.
Theoretical background
19
Figure 2.6: Rate constants of some important oxides as a function of temperature [3, 76].
Furthermore, cobalt in the hexagonal form (cold-worked specimen) oxidises faster than in
the cubic form (annealed specimen) [77]. In case of cold-worked cobalt, grains are fine
and the boundaries act as easy diffusion paths for the transport of cobalt. The rate
constant of CoO in Figure 2.6 obeys an Arrhenius dependency (in the same way as the
other oxides) and can be expressed by the following equation [78]:
⎛ Q ⎞
k = k0 ⋅ exp⎜ −
⎟
⎝ RT ⎠
(2.4)
Q is the activation energy [J/mol], R the gas constant [J/(mol·K)] and T the
temperature [K]. In Figure 2.7, the oxidation mechanism of pure Co at 900 °C in air is
schematically illustrated:
20
Theoretical background
Figure 2.7: Oxidation mechanism of pure Co at 900 °C in air.
Initially, oxygen from the surrounding air gets adsorbed on the metal surface and reacts
with cobalt ions from the bulk metal. Adsorption depends on the surface conditions (such
as roughness, defects, impurities), and the metal and gas impurities. As a result, a thin
film of cobalt oxide forms on the sample surface which separates the metal from the hot
air. Therefore, further oxidation can only occur via Co2+ diffusion on lattice vacancies of
the oxide layer (in addition to electron diffusion on electron holes). At the inner
metal/oxide interface, Co gets ionised from the Co matrix (Figure 2.7) and migrates into
the oxide lattice. At the outer oxide/gas interface, O2- ions react with diffused Co2+ ions
and form new cobalt oxide in the outer direction. A higher oxygen partial pressure would
lead to the formation of thicker oxide layers. Co2+ lattice vacancies get produced at the
outer interface and eliminated at the inner interface resulting in a concentration gradient
of the Co2+ vacancies. This gradient represents the driving force of the oxidation reaction
and remains present as long as interfacial reactions are faster than diffusion. The rate
determining diffusion of the Co2+ cations on the lattice vacancies can be described by
Fick’s first law [79]:
jCo 2+ = − DCo 2+ ⋅
dcCo 2+
dx
(2.5)
Theoretical background
21
Here, jCo2+ is the diffusion flux [mol/(cm2·s)] and DCo 2+ the temperature dependent
diffusion coefficient [cm2/s], which exhibits the same Arrhenius dependency as the rate
constant k (see equation 2.4). The concentration gradient of the Co2+ cations in the oxide
layer is expressed by the term
dcCo 2+
dx
[mol/cm4].
Metal oxides are generally semiconductors at high temperatures, divided in n-type
and p-type semiconductors [80]. According to the oxidation mechanism, i.e. the diffusion
of Co2+ cations on lattice vacancies within the formed oxide layer, cobalt oxide is
considered as a metal deficient p-type semiconductor. In contrast, examples for oxygen
deficient n-type semiconductors may be TiO2 or SiO2. For those oxides, ionic transport is
accomplished by oxygen vacancies [22].
2.3.3 Influence of different alloying elements
Many factors which affect high temperature oxidation of pure Co can be adopted in order
to interpret the oxidation properties of Co-based alloys. However, oxidation is very
complex due to the variety of alloying elements which have different affinities to oxygen
as already pointed out by different standard Gibbs free energies of oxide formation in
Figure 2.5 [73]. In addition, ternary oxide formation, different diffusivities of ions in the
alloy and the oxide, and selective inner oxidation of elements due to diffusion of oxygen
into the alloy contribute to the complexity of alloy oxidation. Therefore, a discussion of
the oxidation behaviour of homogeneous Co–Cr, Co–Cr–X (X = Ternary Element), and
Ni–Cr–Al alloys is carried out in the following, before approaching complex Co-base
superalloy systems.
In most binary Co–Cr alloys, a maximum of about 20 to 25 weight percent of Cr
has to be present, since parabolic rate constants of such alloys posses a minimum value
within this compositional range [81, 82]. As a result, predominantly protective Cr2O3
forms on the alloy surface in addition to small amounts of CoO and CoCr2O4. At lower
Cr-contents, high rate constants and mainly CoO formation can be observed. Very high
Cr-contents lead to an increase in corrosion rate due to dissolution of N2 within the Cr2O3
scale. Additions of 0.5 weight percent of B, V, Be, or Nb to a Co–32Cr alloy are reported
to result in decreased oxidation resistance since low-melting oxides form in addition to
22
Theoretical background
Cr2O3 [82]. Furthermore, protective properties of Cr2O3 are decreased due to oxide
decomposition. However, results of the present work show that lower amounts of B
(0.02 weight percent) in Co-base alloys significantly improve oxidation properties and
therefore seem to be contrary to the findings of Preece and Lucas [82]. According to
literature, additions of 0.5 weight percent of Ti or Zr favoured additional XCr2O4 spinel
formation (X = Alloying element) which however barely affects the oxidation resistance
of the alloy [82]. Al or Si as alloying element lead to the formation of protective oxides
(Al2O3 and SiO2) in addition to Cr2O3, which further increases the oxidation properties.
Later studies [83, 84] also reported the formation of a double layer of protective Cr2O3
and Al2O3 beneath an external layer of CoO in Co–Cr–Al alloys. However, protective
properties of Cr2O3 are reduced at temperatures exceeding 1000 °C and in an oxidising
atmosphere with high flow rates because of Cr2O3 oxidation to volatile CrO3 [23-25]. For
better understanding of the formation of differently composed oxides, a thermodynamic
assessment was carried out recently by Wu et al. [85] in order to study the oxidation
mechanism of a Ni–26Cr–4Al (weight %) ternary alloy. It is demonstrated that the
layering of different oxides can be determined. The most stable oxide (Al2O3) exists
directly above the metal/oxide interface and the least stable oxide (NiO) represents the
outermost layer. A Cr2O3 and a spinel containing layer may form in between the
innermost and outermost layer. A comparison between calculated results and
experimental observations shows excellent agreement. Stability of the important oxides,
such as Al2O3, Cr2O3, or NiO at different temperatures and oxygen pressures can be read
off the Ellingham-Richardson diagram in Figure 2.5 [73] and the corresponding rate
constants which may help determining the thickness of the layers are summarised in
Figure 2.6 [3, 76].
In order to improve oxide layer adhesion and oxidation resistance, sufficient
amounts of Y (at least 0.05 weight percent) in Co–30Cr alloys are reported to be
beneficial [86]. As a result, very little amounts of CoO and CoCr2O4 can be detected
besides Cr2O3. Co/Y-rich phases in the matrix lead to a blockage of Cr3+-diffusion at the
metal/oxide interface and therefore decelerated reaction kinetics. In Al-containing alloys,
small amounts of Y (or other rare earth elements) are not only reported to enhance the
Theoretical background
23
adhesion of the oxide scales, especially under cyclic conditions, but also to increase the
oxidation resistance of the base alloy due to selective Al oxidation [15-21].
Compared to the previously discussed effects of alloying elements and oxide scale
formation on simple binary and ternary alloys, complex superalloys have been optimised
concerning elemental composition in order to obtain improved mechanical and oxidation
properties. Therefore, mechanisms of oxide scale formation, especially the study of
Wu et al. [85], are applicable to complex superalloys, as well. Furthermore, in addition to
the findings of Preece and Lucas on the Si-effect in ternary Co–Cr–Si alloys [82], silicon
is reported to improve the oxidation resistance of Al-containing superalloys at elevated
temperatures (T > 900 °C) not only due to the formation of SiO2 [26, 27, 82] but also due
to the beneficial effect of silicon on the formation of protective Al2O3 on the alloy
surface [28]. Therefore, oxidation resistance of Co-and Ni-base superalloys results mainly
from the formation of protective slow growing Cr2O3, Al2O3, or SiO2 species. Those
oxides (in addition to Ni- and Co-oxides) further decrease oxidation rates of the alloy
significantly, due to their compact structure with very little and barely mobile defects [10]
leading to suppressed ion diffusion through the scales.
2.4 Aqueous corrosion
2.4.1 Significance of electrochemical investigations for superalloys
In addition to high temperature oxidation in air it is also important to investigate the
corrosion behaviour of unoxidised and oxidised superalloys at ambient temperature in
aggressive aqueous environments, such as sulphate or chloride containing solutions or
solutions of different pH values. Sulphates are typical products of the combustion process
and the saline marine air leads to the presence of chlorides on the alloy surface. Since
superalloys may also suffer from corrosion prior or after operation at high temperatures,
additional examination of the corrosion properties under the mentioned conditions
(ambient temperature, aqueous solution) is indispensable. No studies can be found in
literature dealing with the protective properties of high temperature oxide layers on
superalloys in corrosive aqueous environments. Depending on the chemical composition,
24
Theoretical background
thickness, porosity, and adhesion, oxide layers may have very different protective
properties.
Electrochemical measurements are typical tools for investigating the corrosion
behaviour of materials. Examinations are carried out in an aqueous corrosive medium at
ambient temperature in an accelerated manner. For better understanding of the matter,
electrochemical fundamentals according to [87-93] are presented in the following
sections.
2.4.2 Thermodynamic fundamentals
Noble metals are thermodynamically stable in a variety of corrosive media since their
oxidation potential is lower than the reduction potential of the surrounding species. In
contrast, less noble metals with higher oxidation potentials may corrode due to the
difference in redox potential of the two phases in contact. This is the case for most of the
materials used in industry. The driving force for oxidation can lead to either active
corrosion (dissolution) of the metal or formation of a stable oxide layer (passivation) on
the metal surface. A possible anodic oxidation reaction (active corrosion) of e.g. Co can
be written as follows:
Co( s ) → Co 2 + (aq ) + 2e −
Ea0 = -0.28 V
(2.6)
Comparing the standard electrode potential Ea0 (vs. NHE) of the anodic metal oxidation
with the potential of the reduced species shows the thermodynamic stability of the metal
and the probability of corrosion. Of course, alloying elements have to be taken into
account for Co-base superalloys (such as Al, W or Cr) and may have significant influence
on overall corrosion behaviour. Depending on the pH value of the aqueous solution,
following possible reduction reactions may take place and drive above mentioned metal
oxidation:
Acidic solution:
Theoretical background
2 H + (aq ) + 2e − → H 2 ( g )
25
Ec0 = 0 V (at aH + = 1 )
(2.7)
Ec0 = 0.4 V (at aOH − = 1 )
(2.8)
Alkaline and neutral solution:
2 H 2O(l ) + O2 ( g ) + 4e − → 4OH − (aq)
Ec0 = 1.23 V (at aH + = 1 )
In (2.8), Ec0 is the standard cathodic potential. However, the effective electrode potentials
follow the Nernst equation (2.9), since the values above are only valid for standard
conditions. Therefore, the anodic oxidation potential Ea and the cathodic reduction
potential Ec can be both expressed by the general Nernst equation:
k
E = E0 +
RT
ln
zF
∏a
vi
i , Ox
i =1
k
∏a
(2.9)
vi
i , Re d
i =1
R is the gas constant [J/(mol·K)], T the temperature [K], z the number of electrons, F the
Faraday constant [96485.34 J/(V·mol)], E 0 the standard electrode potential ( Ec0 or Ea0 ),
ai,Ox and ai,Red the activity of the involved oxidised and reduced species, and vi the
stoichiometric coefficient of species i (see reaction equation). For the coupled redox cell,
the electromotive force Etotal can be calculated as:
Etotal = Ec – Ea
(2.10)
and the Gibbs free energy ∆G as:
∆G = – zFEtotal
(2.11)
26
Theoretical background
Therefore, it is possible to predict under which conditions (concentration of species
involved, temperature) the Co oxidation reaction (2.6) proceeds thermodynamically. The
oxidised species (Co2+) can either be dissolved and form Co( H 2O)62 + complexes or react
with species from the electrolyte in order to form a cobalt hydroxide (e.g. Co(OH)2) or
cobalt oxide (e.g. Co3O4) layer on the metal surface. Such an electrochemically formed
passive layer can be an effective barrier against further corrosion reactions depending on
its chemical composition, adhesion, porosity, and thickness. In order to predict the
thermodynamics of metal corrosion (here Co) at specific pH and potential values (in a
given electrolyte), so-called Pourbaix diagrams are reliable tools. In Figure 2.8, the
simplified Pourbaix diagram for pure Co in water (T = 25 °C) is shown [94, 95]. The total
concentration of dissolved Co species in water accounts for 10-6 mol/kg.
Figure 2.8: Pourbaix diagram for pure Co in water (T = 25 °C) [94, 95].
Theoretical background
27
Figure 2.9: Potential-pH diagram for pure Co in water (T = 25 °C) according to [94].
The Pourbaix diagram of Co shows different regions of existence, which means that for a
given pH value and potential (Etotal) it is possible to predict if Co is stable, actively
dissolving or passive. The dotted lines (a and b) correspond to the reactions (2.7) and
(2.8) and indicate the region of water stability. The solid lines represent the stability
ranges of Co and its corrosion products. Stability of Co in acidic environment is limited to
potentials below approximately – 0.45 V. At higher pH values (9 to 16), stability is
continuously shifted to lower potentials with a minimum at approximately – 0.95 V at
pH 16. Active dissolution of Co can be observed at potentials higher than – 0.45 V and
pH values lower than 9. However, active dissolution is limited at high potentials due to
the formation of Co(OH)3 (see Figure 2.8), which is unstable only at very high potentials
and low pH values at the same time. For instance, Co would dissolve in an acidic solution
at pH 1. As Ec of reaction (2.7) accounts for approximately – 0.1 V, Co dissolves and
forms Co2+ species. Furthermore, Co2+ can be oxidised to Co3+ if O2 is present in the
solution due to reaction (2.8). Passivation of the Co metal due to cobalt hydroxide and
28
Theoretical background
cobalt oxide formation is present at pH values between approximately 9 and 13 and
potentials higher than – 0.45 V (at pH 9) and – 0.7 V (at pH 13), respectively.
In Figure 2.9, the potential-pH diagram for pure Co in water (T = 25 °C) is shown
according to [94]. In contrast to Figure 2.8, which represents the potential-pH diagram at
25 °C according to Pourbaix, Figure 2.9 takes the presence of stable CoOOH into account
by using new thermodynamic data. As a result, the stability region of Co3O4 increases
considerably whereas Co(OH)2 decreases and Co(OH)3 gets replaced by CoOOH. At this
temperature, CoOOH has a higher thermodynamic stability than Co(OH)3. In addition, the
predominance region of Co2+ and Co3+ species are reduced and therefore the passivity
area of Co is increased compared to the passivity domain given in Figure 2.8.
Furthermore, no stability region of CoO species can be observed. These results [94] show
the importance of the acquisition of new and reliable thermodynamic data in order to
predict the corrosion behaviour of Co or other metals.
However, potential-pH diagrams only predict corrosion properties with respect to
thermodynamics but not kinetics. Therefore, no information on the rates of the possible
corrosion reactions are given and thus further kinetic considerations are necessary which
will be presented in the following chapter.
2.4.3 Reaction kinetics
A key consideration in aqueous corrosion is how fast the reaction proceeds. Mostly, the
corrosion reaction is activation energy controlled, i.e. the rate determining step is the ion
transfer through the Helmholtz double layer on the metal surface. The transition state of
the reaction is referred to as “activated surface complex”. However, the mass transport in
the electrolyte can become rate determining, as well. In this case, the corrosion reaction is
diffusion controlled.
As already mentioned, the oxidised Co2+ species can be dissolved and form
Co( H 2O)62 + complexes. The transition species Co(OH) and CoOH+ represent the
“activated surface complex”. The rate constant k for generation and degeneration of those
species can be described as follows:
k = Be− ∆G*/ RT
(2.12)
Theoretical background
29
B is a constant and ∆G* is the activation energy, which can be reduced or increased by
applying a potential ∆Φ. Equation (2.12) can also be written in terms of the so-called
“Butler-Volmer equation”, i.e. the reaction rate is described by a current flow j.
Therefore, the anodic oxidation and cathodic reduction reaction of Co can be written as:
ja = j0eb∆Φ
(2.13)
jc = − j0e−b' ∆Φ
(2.14)
ja is the anodic current density and jc the cathodic current density. Furthermore, j0, b, and
b' are constants and ∆Φ represents the applied potential. In absence of an external
potential, j0 is the reaction current density. In this case, ja = │jc│ = j0 applies and the
adjusted equilibrium potential of Co immersed in an aqueous solution is Φ0. This
potential would be equal to the Nernst potential of the Co2+/Co redox couple. However,
the prevailing reduction reaction is according to either equation (2.7) or (2.8), since Co is
immersed in an electrolyte which is involved in the overall reaction. For this reaction
equilibrium (ja* = │jc*│ = jcorr), ECo is the equilibrium potential of the oxidised Co and EH
the equilibrium potential of the reduced H2. Figure 2.10 (in accordance with Schmuki and
Graham [87]) shows the schematic polarisation curves of this mixed electrode, i.e. Co
immersed in an aqueous electrolyte in the presence of hydrogen ions:
30
Theoretical background
Figure 2.10: Schematic polarisation curves of a mixed electrode.
The Co2+/Co redox system is now coupled with the H+/H2 redox system. The resulting
equilibrium potential of the mixed electrode is called Ecorr and the corresponding current
density is jcorr. In Figure 2.10, js* stands for the sum of the anodic and cathodic current
density (dotted line). Experimentally, the corrosion current density jcorr and the corrosion
potential Ecorr can be determined by measuring js* and plotting it as a function of an
externally applied potential. This is the so-called “Tafel plot” (see Figure 2.11, in
accordance with Schmuki and Graham [87]). Ecorr can be read off at the minimum of the
curve and jcorr can be obtained by extrapolating the linear parts of the ja and jc branches to
the corrosion potential. The intersection point represents the value of jcorr. If jcorr is
smaller than approximately 10 µA/cm2, the metal is assumed to be passive. In order to
acquire the polarisation curve in Figure 2.11, a three electrode set-up is applied. The
system refers to a reference electrode (e.g. Ag/AgCl) and the potential between the
sample and the reference electrode is varied by means of a counter electrode (mostly
platinum) while measuring the current. More detailed information about the
electrochemical measurements and the set-up used in this study can be obtained in
chapter 3.5.1.
Theoretical background
31
Figure 2.11: Tafel plot of the measured js*.
The corrosion current can be converted to metal loss ∆m by applying Faraday’s law as
follows:
⎛M ⎞
∆m = ⎜ ⎟ j ⋅ t
⎝ zF ⎠
(2.15)
M is the molar mass of the metal, z the number of electrons, F the Faraday constant
[96485.34 J/(V·mol)], j the measured current density, and t the time. However, this
assumes 100 per cent current efficiency with respect to metal dissolution and therefore no
other reactions take place. Furthermore, it should be mentioned that an externally applied
potential leads to cathodic reduction reactions (hydrogen evolution) at the Pt counter
electrode whereas metal oxidation occurs at the surface of the working electrode. This is
in contrast to the (unbiased) open circuit equilibrium case.
There is also a variety of other techniques for determination of the corrosion rate,
such as the electrochemical impedance spectroscopy (EIS), which is usually performed at
the corrosion potential. In the present study, this method is utilised for the
electrochemical characterisation of pure Co and Co-base superalloys, as well. From the
32
Theoretical background
impedance spectra (see Figure 2.12, in accordance with [96]), the polarisation resistance
(i.e. total impedance RE + Rct) of the system can be obtained, which is indirectly
proportional to the corrosion current density jcorr. Rct is the charge transfer resistance, RE
the resistance of the electrolyte, and Cdl the double layer capacity on the metal surface. As
shown in Figure 2.12, impedance [Ω cm2] is plotted as a function of frequency [Hz], both
on logarithmic axes. This is the so-called “Bode plot”. The amplitude of the corrosion
potential is kept small in order to avoid removal from equilibrium conditions. A reduced
double layer capacity would lead to enhanced polarisation resistance and therefore
reduced jcorr values due to thicker passive layers. By applying the corrosion potential, the
phase shift φ [°] of the current with respect to the potential can be measured as a function
of frequency, as well (see Figure 2.12). The extent of the phase shift (maximum at the
frequency f ') and its region of occurrence (frequency range) may give information about
the corrosion resistance of the material.
Figure 2.12: Schematic diagram of a Bode plot.
So far, the discussed electrochemical processes are activation controlled.
However, as already mentioned, the mass transport in the electrolyte can become rate
determining if the formation of the activated complex is fast compared to the diffusion of
anions to the metal surface or dissolved cations away from the surface. In this case, the
corrosion reaction is diffusion controlled and the reaction rate depends only on the supply
Theoretical background
33
of O2 (g) to the metal surface. The diffusion flux is again determined by Fick’s first law
(see also equation (2.5)). Here, the concentration gradient of oxygen is within the “Nernst
diffusion layer”. In the limiting case, the surface concentration of the reacting species
equals zero, since all arriving ions react instantly with the metal ions. As a consequence,
the limiting current density jL becomes potential independent and can be written as
follows:
⎛ zFD ⎞
jL = ⎜
⎟⋅c
⎝ δ ⎠
(2.16)
Equation (2.16) shows that jL depends on the diffusion coefficient D, the thickness of the
Nernst diffusion layer δ, and the anion concentration c in the bulk solution. Therefore, jL
depends only on diffusion. In the Tafel plot (see Figure 2.11), the linear slope in the
cathodic and anodic regime indicates activation control, whereas at high cathodic and
anodic potentials current density would become potential independent which indicates
diffusion control. This would result in hardly increasing current density values (jL) in the
Tafel plot (not shown in Figure 2.11) with increasing cathodic or anodic potentials.
2.4.4 Passivity and localised corrosion
In contrast to active corrosion (dissolution) of a metal, the formation of a second phase
film (usually an insoluble three dimensional surface oxide layer) is favoured. This passive
film can form at the equilibrium, i.e. at Ecorr, but also as a result of an external anodic
potential. In the corresponding potentiodynamic polarisation curve, anodic passivation
can be observed as an abrupt decrease of corrosion current density (several orders of
magnitude possible) at a specific onset potential Ep. Here, the metal changes from an
active to a passive state (see Figure 2.13). In terms of an electrochemical treatment,
passivation of a metal surface represents a significant deviation from ideal electrode
behaviour [93] (compare Figure 2.11). Generally, the reaction scheme for passivation can
be divided into the active, transition/prepassive, and passive range [97, 98], followed by
transpassivity at higher potentials.
34
Theoretical background
Figure 2.13: Schematic polarisation curve of a passivating metal.
In the transition/prepassive range, metal hydroxide adsorbates increasingly cover the
metal surface and therefore progressively block active dissolution. The passivation
potential Ep is reached when the surface is entirely covered with metal hydroxides.
Deprotonation may lead to the formation of a passive layer mainly consisting of metal
oxide [93]. According to the revised Pourbaix diagram in Figure 2.9 [94], Co forms stable
Co(OH)2, Co3O4, and CoOOH phases during anodic passivation (depending on the
applied potential), in which Co3O4 is expected to possess the highest protective
properties. Furthermore, the current density in the passive range (jp) is a measure for the
protective properties of the passive layer. The composition and thickness of such an oxide
layer vary with applied potential, polarisation time, temperature, and environmental
conditions (species in solution, pH value, etc.). The situation of passive layer growth on a
metal immersed in an electrolyte is given schematically in Figure 2.14 (in accordance
with Schmuki [93]):
Theoretical background
35
Figure 2.14: Growth of a passive layer on the metal surface in the presence of an
externally applied anodic potential.
The metal (such as Co) gets oxidised at the inner interface according to equation (2.6). If
no external potential is applied, i.e. under open circuit conditions, growth mechanism is
similar to high temperature oxidation (see Chapter 2.3.2 and Figure 2.7). In this case,
metal oxidation at the inner interface is coupled with oxygen reduction at the outer
interface by ion and electron diffusion through the layer in order to enable oxide growth.
If an outer anodic potential is applied, layer growth can additionally occur without
electron migration through the oxide since reduction of oxygen takes place at the counter
electrode (see Figure 2.14). Therefore, layer growth can take place with OH¯ species
present in the solution, involving a (field-aided) deprotonation of the layer [93]. The
driving force is the applied potential Eappl. Since ions and electrons have to migrate
through the oxide layer, its conductance is essential for the process. This is comparable to
the conductance of ions and electrons through the oxide layer during high temperature
oxidation. Therefore, oxides such as Al2O3, SiO2, or Cr2O3 are expected to show
comparable protective effects at ambient temperature in aqueous environments and at
high temperature in air. These oxides possess very little amounts of barely mobile
defects [10], which lead to suppressed ion and electron diffusion, respectively (compare
36
Theoretical background
chapter 2.3.3). In comparison to thin passive films on metals, which may form due to
anodic polarisation, oxides formed at high temperatures are considerably thicker.
Independent of their composition, those oxides may also show good barrier properties
against electrochemical dissolution if being immersed in an aqueous solution, simply due
to their compact structure and the high oxide thickness.
Quantitative characterisation of passive layer formation kinetics dates back to the
studies of Cabrera and Mott [99] and Vetter [100], stating that at low temperatures and
high field strengths F (> 106 V/cm), the extent of exclusive ion diffusion is much lower
than field-aided transport. Growth models consider different processes to be ratedetermining [99, 101]. However, all approaches can be expressed by the following
equation:
j = A ⋅ exp (β F )
(2.17)
j is the current density, F is the field strength over the oxide, and A and β are constants.
According to (2.17), layer growth is self-limiting since F gets decreased with increasing
layer thickness (valid for a constant applied potential), resulting in the “inverse
logarithmic” growth law:
1
= A − B ⋅ log(t )
x
(2.18)
x is the layer thickness, t is the time, and A and B are constants. However, there are
several further approaches for describing the mechanisms of passive layer growth on
metals. Therefore, it can be concluded that up to date, this issue is not completely
understood and further mechanistic studies have to be carried out.
The protective properties of a passive layer are determined by ion transfer through
the film, as well as the stability of the layer against dissolution [93]. Several factors can
influence ion diffusion within the layer, such as its chemical composition, crystal
structure, grain boundaries, and the amount of defects and pores. For passivated pure
metals, such as Co, information mainly on the layer’s structure and chemical composition
Theoretical background
37
can lead to a reliable explanation of the protective properties and the stability of the
passive film. However, for alloy systems the situation is more complicated. Explanations
are based on the already mentioned ion mobility in the oxide layer [102, 103], percolation
effects [104, 105], structural arguments [106], and charge distribution [107, 108].
Percolation effects are based on the formation of an insoluble oxide network (e.g. Cr2O3)
within the soluble and less protective main species (e.g. Co oxide) of the passive
layer [104, 105]. The formation of the protective species depends on the amount of Cr
content within the alloy and a continuous and protective network is formed once a
specific content is reached. However, the critical value depends on the electrolyte
chemistry and differs in case of different dissolution reactions [109]. Therefore, in
addition to the geometrical considerations, a model has to consider the dissolution
chemistry, as well. Structural arguments [106] are based on an increased tendency of the
oxide to become more disordered with increasing Cr content, suggesting that amorphous
layers are more protective than crystalline layers due to bond and structural flexibility.
Explanations based on charge distribution [107, 108] assume the presence of a bipolar
passive layer composed of an anion-selective region (on the metal side) and a cationselective region (on the solution side). The bipolar layer retards the anodic ion flux. An
external anodic potential would lead to dehydration of the layer and H+ movement
through the cation-selective region to the electrolyte. Simultaneously, metal ion diffusion
is retarded in the anion-selective region on the metal side. All presented mechanisms may
explain the protective properties of passive layers on some alloys immersed in certain
electrolytes. However, still no overall model exists due to the fact that the composition
and structure of the passive layer depends on the passivation parameters. Therefore,
comparing different results is very complicated.
Under certain circumstances, passivity of a metal is susceptible to breakdown, for
instance due to localised dissolution of the passive layer [110-117]. Distinct anodic
regions (“pits”) on the metal surface form where dissolution reactions according to
equation (2.6) are dominant. Those active pits are surrounded by the intact cathodic
passive layer where the reduction reaction occurs according to equation (2.8). Generally,
pit initiation and pit growth can be distinguished. Initiation of pits can be ascribed to
inhomogeneities of the bulk metal (such as precipitates, grain boundaries, or dislocations)
38
Theoretical background
or to passive layer properties (such as local composition or structure variations) [93].
Pitting corrosion of a passivated metal immersed in NaCl aqueous solution is
schematically illustrated in Figure 2.15. Pitting occurs for a variety of metals, such as Fe,
Al, or Co, in halide containing solutions. However, almost no literature [118] can be
found concerning pitting corrosion of conventional and novel Co-base superalloys.
Figure 2.15: Schematic cross section through an actively growing pit at the surface of a
metal immersed in NaCl solution.
The pitting corrosion process is autocatalytic, i.e. initial dissolution conditions stimulate
further dissolution of the metal inside of the pit. Metal ions form an aqueous complex
M ( H 2O)62 + which reacts with the halide ions (here Cl¯) in order to form the chlorocomplexes M ( H 2O ) 4 (OH )Cl and M ( H 2 O) 5 Cl + , as shown in Figure 2.15. For this, Cl¯
ions have to migrate into the pits. The following equations summarise the chemical
reactions in the pits [93]:
Theoretical background
39
M 2 + + 6 H 2O → M ( H 2O)62 +
(2.19)
M ( H 2 O) 62+ + Cl − → M ( H 2 O) 5 Cl + + H 2 O
(2.20)
M ( H 2 O) 5 Cl + ↔ M ( H 2 O) 4 (OH )Cl + H +
(2.21)
As a result of the reactions above, Cl¯ and H+ concentrations within the pits increase,
which further accelerate metal dissolution. In addition, the ratio of the cathodic to the
anodic area is crucial for the localised dissolution rate. High cathodic currents result from
a large cathodic area and therefore equally high anodic currents have to be provided by
the small anodic pit area in order to ensure electroneutrality. As a result, the local
corrosion rate at the pits increases with a larger cathode/anode ratio. In the polarisation
curve of an initially passive metal (Figure 2.16, in accordance with Schmuki [87]), the
onset of localised corrosion can be observed as a strong increase in current density at a
specific anodic potential, Epit, far below the occurrence of oxygen evolution or
transpassive dissolution. If the applied anodic potential is higher than Epit, stable pit
growth occurs (curve 1b). Current transients due to metastable pits are represented by
curve 1a. Passivity breakdown due to passive layer dissolution would lead to the
polarisation curves 2b (if secondary passivation takes place) and 2a (in the absence of
secondary passivation). If a reductive or oxidative reaction of the passive layer is
favoured under the present environmental conditions (see Pourbaix diagram) and the
resulting species are soluble, depassivation will take place (curve 2a). Furthermore, a
strong increase of current density can also be observed due to oxygen evolution (at EO2 )
or at very high applied potentials (Ebd) due to dielectric breakdown.
40
Theoretical background
Figure 2.16: Different polarisation curves of a passive metal showing typical passivity
breakdown phenomena.
2.5 Scientific aims of the present work
Based on well-investigated pure cobalt and simple alloy systems, such as binary and
ternary Co-alloys, fundamental assumptions regarding electrochemical properties and
oxidation behaviour of more complex Co-base alloys can be drawn. However, the
chemical composition and microstructure of novel γ'-strengthened Co–Al–W superalloys
strongly differ from conventional Co-base alloys. Therefore, a reliable prediction of the
resulting properties, based on the literature, is not possible.
The aim of the present work was to systematically investigate the isothermal high
temperature oxidation behaviour of novel γ'-strengthened Co-base superalloys. In this
regard, the effect of different alloying elements (B, Cr, Ni, Si, Y, and Ti), oxidation
temperatures, and exposure times on oxide scale formation and grain boundary
precipitation were of particular interest. Furthermore, occurrence of microstructural
changes, as a result of exposure to high temperatures and alloying element additions, had
to be investigated. Electrochemical characterisation was carried out at ambient
Theoretical background
41
temperature on bare and oxidised alloys in several aqueous solutions of different pH, in
order to explore the protective properties of the formed passive layers and high
temperature oxide scales, respectively. Furthermore, susceptibility to pitting corrosion in
chloride containing solution was of particular interest.
So far, almost no research on high temperature oxidation and aqueous corrosion of
novel γ'-strengthened Co-base superalloys has been conducted. Therefore, the findings of
the present work are expected to contribute to the development of new Co-based alloys,
which have the potential to exceed the high temperature properties of commonly used Nibase superalloys.
42
Experimental
3. Experimental
3.1 Investigated materials
The investigated polycrystalline Co-base superalloys were vacuum arc melted as ingots
from commercially pure metals by “Hauner Metallische Werkstoffe” (HMW, Röttenbach,
Germany). The extremely high melting point of W (3422 °C) and the evaporation of Al
due to its relatively low melting point (660 °C) lead to slight differences between the
nominal and measured compositions of the experimental alloys after the melting process.
The corresponding values measured by energy dispersive X-ray spectroscopy (EDX, see
also section 3.5.4) and the used alloy abbreviations are summarised in Table 3.1.
Table 3.1: Nominal and measured compositions of the experimental superalloys and
utilised abbreviations.
Utilised
abbreviations
Nominal composition
in at%
9W
0.04B
0.08B
0.12B
0.04B4Cr
0.04B8Cr
0.12B9Ni
0.12B18Ni
0.12B2Si
0.12B0.005Y
0.12B2Ti
Co-9Al-9W
Co-9Al-9W-0.04B
Co-9Al-9W-0.08B
Co-9Al-9W-0.12B
Co-9Al-9W-0.04B-4Cr
Co-9Al-9W-0.04B-8Cr
Co-9Al-9W-0.12B-9Ni
Co-9Al-9W-0.12B-18Ni
Co-9Al-9W-0.12B-2Si
Co-9Al-9W-0.12B-0.005Y
Co-9Al-9W-0.12B-2Ti
Measured composition in at%
Co Al W Cr
Ni
Si
83.4 8.1 8.4
81.9 8.4 9.7
81.5 8.9 9.5
81.8 8.8 9.3
77.8 8.6 9.5 4.1
73.8 8.7 9.3 8.2
71.4 9.6 9.6
9.3
63.2 9.2 9.6
18.0
80.2 8.1 9.4
2.3
82.9 7.9 9.2
79.8 8.4 9.6
-
Ti
2.2
B-contents were confirmed by supplementary glow discharge optical emission
spectroscopy (GDOES, RF GD Profiler, HORIBA Jobin Yvon GmbH) since EDX is not
suitable for detecting B and other elements with an atomic number lower than six.
GDOES measurements were carried out by Natalie Kömpel at “Neue Materialien Fürth
GmbH” (NMF GmbH, Fürth, Germany). Qualitative EDX investigations (see chapter 4)
Experimental
43
confirmed the presence of Y in the experimental 0.12B0.005Y alloy due to Yenrichments in the oxide layer and at the oxide/alloy interface. However, the exact
content in the alloy could not be determined due to the low amount of Y (0.005 at%).
In addition, pure cobalt (hexagonal α-cobalt) and a cast Co-base superalloy
(MAR-M-509®) were obtained from the Institute of General Material Properties (WW1,
University of Erlangen-Nürnberg, Germany) and investigated in this study. MAR-M-509
is a conventional corrosion resistant γ'-free superalloy with the following composition
(in at%): 57.85Co–24.86Cr–10.25Ni–2.29W–1.16Ta–3.00C–0.33Zr–0.25Ti.
3.2 Heat treatment
Heat treatment of the as-cast experimental alloys (see Table 3.1) was performed in
cooperation with Alexander Bauer from WW1. The alloys had to undergo a solution heat
treatment for 12 h at 1300 °C in argon atmosphere and a subsequent ageing for 200 h at
900 °C in air. Solution heat treatment was performed for homogenising the material and
ageing was carried out in order to enable the growth of the γ'-phase and to obtain the
desired γ/γ'-microstructure. Differential scanning calorimetry (DSC, Netzsch STA 409
CD) measurements on the as-cast alloys with a heating/cooling rate of 5 K/min were
carried out by Peter Randelzhofer from the “Institute of Metals Science and Technology”
(WTM, Friedrich-Alexander-University of Erlangen-Nürnberg, Germany). For this, cubic
samples with an edge length of about 3 mm were utilised. DSC enabled the determination
of the characteristic temperatures for both heat treatment steps.
3.3 Sample preparation
For sample preparation, discs with a thickness of 2-3 mm were cut (Secotom-10,
Struers GmbH) from the heat treated alloy ingots and the pure Co sample, respectively.
For high temperature oxidation experiments, it was necessary to drill a pin hole in each
sample. The small plates were wet ground using SiC paper up to grit 1200. Subsequent
44
Experimental
polishing of the samples using lubricant and diamond suspensions of different grain sizes
(DP-Suspension-P 6 µm, 3 µm, and 1 µm, Struers GmbH) removed residual scratches
without causing excessive deformation or smearing of the surface microstructure.
Furthermore, the samples were cleaned with ethanol between the different polishing steps
in order to remove residues from the sample surface.
Oxidised alloys were embedded into a fast-curing resin (Technovit 4071, Heraeus
Kulzer GmbH) and cross-sections were prepared by grinding and polishing in order to
characterise the oxide layers by SEM and EDX (see chapters 3.5.2 and 3.5.4). Prior to
that, samples were masked with an aluminium tape with conductive adhesive
(Plano GmbH) and carbon was vapour-deposited on the surface (MED 010,
Balzers Union) in order to provide electrical conductance.
For time-of-flight secondary ion mass spectrometry investigations (ToF-SIMS, see
chapter 3.5.6) the resin was removed after the grinding and polishing step. In addition, no
vapour-deposition of carbon or masking was carried out. Since ToF-SIMS techniques are
very surface sensitive, it was of particular importance to provide clean samples without
any organic contamination on the surface.
For X-ray diffraction studies (XRD, see section 3.5.7) no cross-sections of the
oxide layers had to be prepared. Instead, a thin alloy sheet of about 250 µm thickness was
isothermally oxidised for 500 h at 900 °C in air (see also chapter 3.4). The metallic
material was completely transformed into an oxide and by means of milling (using a
mortar) an appropriate powder for subsequent XRD investigations could be obtained.
3.4 High temperature oxidation
High temperature oxidation of the alloys and pure cobalt was carried out isothermally in
air at 800, 900, and 1000 °C for different oxidation times (between 1 and 500 h) in a
vertical tube furnace. The specimens were attached to a microbalance by means of a heat
resistant freely suspended platinum or NiCr wire and introduced into the furnace after the
set temperature was reached. A relatively constant temperature, which corresponds to the
set temperature of the furnace control, was solely present in a narrow area within the
Experimental
45
vertical tube. Therefore, specimens were always suspended at this specific height of the
tube in order to assure the desired oxidation temperatures. Mass change (in mg/cm2), i.e.
oxide layer growth on the sample surface, was determined simultaneously with increasing
oxidation time by thermogravimetry (TGA). Reproducibility of the curves was checked
by performing at least three measurements on one alloy at a specific oxidation
temperature. In Figure 3.1, the set-up for the isothermal oxidation experiments in air is
schematically illustrated. The surface area of the samples was measured prior to oxidation
in order to provide comparability of the results. After finishing the oxidation experiments,
samples were immediately removed from the vertical furnace and cooled off at room
temperature.
Figure 3.1: Experimental set-up for the isothermal oxidation experiments in air.
46
Experimental
3.5 Characterisation methods
3.5.1 Electrochemical investigations
Electrochemical characterisation of the unoxidised (polished) and oxidised materials were
carried out at ambient temperature in an electrochemical O-ring cell with a conventional
three-electrode set-up (see Figure 3.2), open-to-air. For that, 0.159 cm2 of the sample
surface was in contact with the aqueous solution. The test sample was used as the
working electrode, platinum as the counter electrode, and Ag/AgCl (immersed in
3 M KCl; Electrode potential is + 208 mV vs. SHE) as the reference electrode. As
electrolytes 0.5 M H2SO4 (pH 0.6), 1 M Na2SO4 (pH 5.9), 0.5 M NaCl (pH 5.8), and
0.1 M NaOH (pH 12.8) aqueous solutions were used. In order to perform the
electrochemical studies, an electrochemical workstation (“IM6eX”), a potentiostat
(“XPot”),
and
the
corresponding
software
(“Thales”)
from
Zahner-Elektrik
GmbH & Co. KG were used. Each experiment was carried out three times in order to
ensure reproducibility. Electrochemical characterisation of the samples consisted of three
different measurements. In the first experiment, the potential of the system was recorded
while exposing the sample to the electrolyte. After 3 h, a nearly constant open circuit
potential (OCP vs. Ag/AgCl) was reached. Subsequent electrochemical impedance
spectroscopy (EIS) measurements were performed at the constant OCP value, with a
frequency range between 5 mHz and 1 MHz (starting at 1 MHz) and an amplitude of
± 10 mV, in order to record the total impedance (in Ω cm2) of the passive film or high
temperature oxide layer on the specimen surface. The phase shift of the current with
respect to the applied potential is not presented in this study, since oxide layers on Cobase superalloys are very complex and therefore a reasonable evaluation and discussion
of the resulting complicated phase angle curves was not possible. After the OCP and EIS
experiments, potentiodynamic j/E (current density vs. potential) measurements were
conducted applying a constant scan rate of 10 mV/s in a potential range between – 1 and
+ 2 V vs. Ag/AgCl. The potentiostat was used in order to control the potential between
the working electrode (sample) and the reference electrode while measuring the current
between the Pt counter electrode and the sample.
Experimental
47
Figure 3.2: Electrochemical three-electrode set-up.
3.5.2 Scanning electron microscopy
The morphology of the high temperature oxide layers and the alloy microstructure were
characterised by means of SEM of prepared cross-sections using Jeol JSM 6400, Zeiss
Cross-Beam 1540 EsB FIB (both at WW1), FE-SEM Hitachi S-4800 (at the Institute for
Surface Science and Corrosion, University of Erlangen-Nürnberg, Germany), and FEI
Quanta 3D FEG (at the Monash Centre for Electron Microscopy, Melbourne/Clayton,
Australia) microscopes. Imaging was accomplished in backscattered electron mode (BSE)
in order to achieve elemental contrast. Furthermore, BSE micrographs and the image
processing software “ImageJ” were utilised for determining the mean oxide layer
thicknesses. For this, at least 10 separate measurements of the oxide layer thickness were
performed.
48
Experimental
3.5.3 Electron channelling contrast imaging
Electron channelling contrast images (ECCI) of oxide layers in cross-section were taken
using a FEI Quanta 3D FEG electron microscope. Due to the channelling of electrons
along the crystal planes of the investigated sample, changes in crystallographic
orientation, e.g. distortion of the crystal planes, produce changes in grey scale in the
ECCI. For this imaging set-up, samples had to be tilted and low-angle backscattered
electrons, i.e. channeled electrons, were detected. By this method, it was possible to
investigate the crystallinity of the different obtained oxide layers on the alloy surface.
3.5.4 Energy dispersive X-ray spectroscopy
Qualitative and quantitative EDX measurements (Oxford Instruments and EDAX/TSL
Genesis 4000) across the prepared cross-sections were carried out in order to study the
composition of the alloys and the oxide scales. EDX mappings also allowed the
determination of elemental distribution within the oxide layer and the alloy, respectively.
3.5.5 Optical microscopy
After the electrochemical studies of unoxidised (polished) and oxidised materials in
0.5 M NaCl aqueous solution, optical microscopy (Nikon Instruments Inc., Eclipse
LV 150) was applied in addition to SEM for characterising the corrosive attack of
chloride ions on the sample surface (top-view).
3.5.6 Time-of-flight secondary ion mass spectrometry
Positive and negative static ToF-SIMS (TOF.SIMS 5 spectrometer, ION-TOF GmbH)
allowed for qualitative determination of the spatial B- and Si-distribution by means of
ToF-SIMS mapping, since it was not possible to detect B or hard to distinguish Si peaks
from W peaks by means of EDX. The oxidised alloys were irradiated with a pulsed
25 keV Bi+ ion beam. Signals were identified according to their accurate mass and
isotopic pattern in the high mass resolution mode and chemical maps were recorded in
high lateral (and low mass) resolution mode. Therefore, a lateral resolution of about 200
to 300 nm and a sampling depth of about 10 to 20 Å was achieved with this setting.
ToF-SIMS images were recorded in positive and negative polarity and care was taken not
Experimental
49
to exceed the static limit of 1013 ions/cm2. The brightest pixel corresponds to the largest
number of counts for the respective fragment.
3.5.7 X-ray diffraction
The crystalline phases within the oxide layer on top of the alloy surface were identified
by using XRD (Philips X’Pert). The incidence angle of the X-rays on the milled oxide
powder was 5°. The step size for the measurement amounts to 0.01° and the time per step
was set to 10 seconds.
50
Results
4. Results
4.1 High temperature oxidation of a quaternary Co–Al–W–B superalloy
4.1.1 Long-term oxidation behaviour
Figure 4.1 shows the isothermal oxidation behaviour of the quaternary 0.12B alloy at 800
and 900 °C in air. Long-term thermogravimetric measurements are carried out for 500 h
and representative oxidation curves are illustrated. For comparison, oxidation behaviour
of pure Co at 900 °C is given in Figure 4.1, as well.
Figure 4.1: Isothermal oxidation of the 0.12B alloy at 800 and 900 °C in air.
A higher oxidation temperature leads to enhanced mass gain of the investigated 0.12B
alloy. As a result, much thicker oxide layers of 126 ± 18 µm form after 500 h of exposure
at 900 °C, compared to 23 ± 4 µm after 500 h at 800 °C (see Table 4.1). At first, a
transient oxidation period can be observed at both temperatures characterised by strong
mass gain. After about 100 h of exposure, continuous oxide layer growth takes place at
Results
51
900 °C, whereas further oxidation stops at 800 °C. In contrast, pure Co forms only an
outer CoO single-layer after oxidation at 900 °C [119] and exhibits significantly stronger
mass gain (see Figure 4.1). Therefore, a thick layer (231 ± 27 µm) can be measured at
900 °C after already 65 h of exposure.
Table 4.1: Mean oxide layer thickness (and standard deviation) of the investigated alloys
after oxidation for 500 h at 800, 900, and 1000 °C in air.
Alloy
9W
0.04B
0.08B
0.12B
0.04B4Cr
0.04B8Cr
0.12B9Ni
0.12B18Ni
0.12B2Si
0.12B0.005Y
0.12B2Ti
Mean oxide layer thickness in µm after
500 h at 800 °C
500 h at 900 °C
500 h at 1000 °C
143 ± 21 µm
259 ± 11 µm
91 ± 14 µm
187 ± 14 µm
29 ± 6 µm
139 ± 11 µm
23 ± 4 µm
126 ± 18 µm
713 ± 39 µm
32 ± 6 µm
96 ± 13 µm
14 ± 3 µm
150 ± 10 µm
242 ± 9 µm
230 ± 7 µm
72 ± 10 µm
58 ± 7 µm
275 ± 95 µm
24 ± 5 µm
134 ± 10 µm
29 ± 5 µm
55 ± 9 µm
4.1.2 Microstructure and composition of the formed oxide layers
Figure 4.2 shows a representative cross-sectional SEM micrograph (BSE mode) of the
long-term oxidised 0.12B alloy after 500 h at 800 °C, which reveals three different oxide
layers on the alloy surface. The outermost layer (1) is mostly composed of Co3O4, which
can be identified based on the EDX measurements in Figure 4.3 and supplementary XRD
data (not shown). The middle mixed oxide layer (2) consists of several oxides of all
alloying constituents (e.g. CoWO4 and CoAl2O4) and the innermost layer (3) of Al2O3. As
shown in Figure 4.2, a phase transformation occurs between the base alloy and the Al2O3
layer. In this area no γ' but a Co-rich zone and elongated precipitates can be observed.
In Figure 4.4, the relative crystal orientation within the different oxide layers is
illustrated by means of ECCI. The outer Co3O4 layer possesses a granular structure
whereas the mixed oxide layer may be either amorphous or nanocrystalline.
52
Results
Figure 4.2: a) Oxide layers on the surface of a 0.12B alloy after 500 h at 800 °C in air.
b) γ'-depleted area between the γ/γ'-matrix and the innermost Al2O3 layer.
Figure 4.3: EDX point scan of the oxidised 0.12B alloy surface (500 h at 800 °C in air).
Figure 4.4: ECCI of the oxidised 0.12B alloy surface (500 h at 800 °C in air).
Results
53
Long-term oxidation for 500 h at 900 °C leads to comparable oxide layers on the 0.12B
alloy surface, with respect to morphology and composition. In Figure 4.5, a representative
SEM micrograph of such an oxidised alloy is presented. The outermost oxide layer (1)
can be identified as a mixture of mainly CoO and residues of Co3O4, based on XRD data
and EDX point scans in Figures 4.6 and 4.7, respectively. The middle mixed oxide layer
(2) and the inner Al2O3–rich region (3) are supposed to be composed of the same oxide
species as the oxide layers formed at 800 °C. However, compared to the continuous
Al2O3 layer formed at 800 °C (Figure 4.2), a discontinuous Al2O3–rich region can be
observed after long-term oxidation at 900 °C. In addition, a discontinuous γ'-free area
next to the formed Al2O3 can be detected in Figure 4.5.
Figure 4.5: Oxide layers on the surface of a 0.12B alloy after 500 h at 900 °C in air.
54
Results
Figure 4.6: XRD of the 0.12B alloy after oxidation for 500 h at 900 °C in air.
Figure 4.7: EDX point scan of the oxide layers on the 0.12B alloy surface after 500 h at
900 °C in air.
4.1.3 Influence of oxidation time and temperature on oxide layer formation
In addition to long-term experiments for 500 h, oxidation is also carried out for 5, 24, and
100 h at 800 °C and 900 °C in air, in order to determine the influence of oxidation time
and temperature on oxide layer formation. Figures 4.8 and 4.9 show a comparison of
Results
55
SEM micrographs and EDX mappings of the differently oxidised 0.12B alloy. Longer
exposure times in air result in thicker oxide layers. Furthermore, measured oxide layer
thicknesses in Figure 4.10 correlate very well with the thermogravimetric curves in
Figure 4.1. After 5 h of exposure, cobalt oxide formation occurs at the initial alloy surface
due to reaction of Co with the surrounding oxygen. This results in a Co-depleted zone
beneath the outer cobalt oxide scale. Due to diffusion, Al and W get enriched in the Codepleted region and form mixed oxides (spinels). Oxidation for at least 24 h leads to the
formation of inner Al2O3 and transformation of the γ/γ' microstructure adjacent to the
Al2O3. At 800 °C, a continuous Al2O3 layer establishes after about 100 h of oxidation. In
contrast, it can be observed that the layer remains discontinuous at 900 °C.
Figure 4.8: SEM micrographs and EDX mappings of the 0.12B alloy surface after 5, 24,
100, and 500 h at 800 °C in air.
56
Results
Figure 4.9: SEM micrographs and EDX mappings of the 0.12B alloy surface after 5, 24,
100, and 500 h at 900 °C in air.
Figure 4.10: Measured mean oxide layer thicknesses and standard deviations after
different oxidation times at 800 and 900 °C in air.
Results
57
4.2 Influence of different alloying elements on high temperature properties
4.2.1 Influence of the boron content
Representative oxidation curves of isothermally oxidised B-free and B-containing Cobase superalloys (9W, 0.04B, 0.08B, and 0.12B) are given in Figures 4.11 and 4.12.
Long-term oxidation of the novel alloys is performed at 800 °C and 900 °C, respectively.
For comparison, a conventional MAR-M 509 superalloy is oxidised at 800 °C for 100 h
and obtained results are presented in Figure 4.11, as well. It can be observed that the Crrich MAR-M 509 alloy exhibits superior oxidation behaviour compared to the novel Bcontaining superalloys. However, the 0.12B alloy with the highest B-content (see also
chapter 4.1) shows the lowest measured mass gain at both oxidation temperatures
compared to the 9W, 0.04B, and 0.08B alloys. Therefore, the thinnest oxide layers of this
alloy series can be measured for the 0.12B alloy after long-term oxidation for 500 h (see
Table 4.1). All alloys continuously oxidise at 800 °C and 900 °C, except for the 0.08B
and 0.12B alloys at 800 °C. In this case, no further mass gain can be observed after
exposure times longer than 100 to 150 h. Generally, a decreased B-content results in
reduced oxidation resistance at both temperatures and therefore a parabolic time
dependency of the oxide layer growth can be measured for B-free 9W alloys, whereas
subparabolic oxide layer growth or even a stop in oxide layer growth after a certain
oxidation time can be obtained by adding sufficient amounts of B (at least 0.08 at%). In
Figure 4.13, SEM micrographs of the 9W, 0.04B, 0.08B, and 0.12B alloys (after longterm oxidation for 500 h at 800 °C) reveal that a higher B-content results in more
pronounced inner Al2O3 formation and thus lower oxide scale thicknesses. In addition,
improved oxide layer adhesion can be obtained by adding B. These observations also
apply for the different B-containing alloys oxidised at 900 °C for 500 h.
58
Results
Figure 4.11: Isothermal oxidation of the different B-containing alloys and the
conventional MAR-M 509 alloy at 800 °C in air.
Figure 4.12: Isothermal oxidation of the different B-containing alloys at 900 °C in air.
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59
Figure 4.13: SEM micrographs of the different B-containing alloys after 500 h at 800 °C
in air.
Determination of the spatial B-distribution can be achieved by ToF-SIMS measurements
on the different B-containing alloys. In Figure 4.14, mappings of the B-distribution in the
0.04B (a) and 0.12B alloy (b) are presented after oxidation for 500 h at 800 °C in air.
Occurrence of B-containing grain boundary precipitates is independent of the B-amount,
i.e. even the 0.04B alloy shows sufficient precipitation. In addition, B-rich precipitates
within the grains can be detected. Due to high temperature treatment, W and B
accumulate in the inner oxide layer and within precipitates (Figures 4.14 and 4.15),
whereas almost none of those elements can be detected in the outer oxide layer.
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Results
Figure 4.14: ToF-SIMS mappings of the B-distribution in the 0.04B alloy (a) and 0.12B
alloy (b) after oxidation for 500 h at 800 °C in air.
Figure 4.15: ToF-SIMS mappings of the W (a) and B (b) distribution in the 0.12B alloy
after oxidation for 500 h at 800 °C in air.
4.2.2 Effect of chromium addition
In Figure 4.16, thermogravimetric results of the oxidation experiments on the 0.04B,
0.04B4Cr, and 0.04B8Cr alloys are illustrated (representative curves). The alloys are
isothermally oxidised for 500 h at 800 and 900 °C in air. Results of the Cr-free 0.04B
Results
61
alloy are also presented in Figures 4.11 and 4.12. Generally, longer exposure times lead to
increased mass gain and therefore increased oxide layer thickness. Oxidation of the high
Cr-containing 0.04B8Cr alloy results in the lowest mass gain after 500 h at 800 and
900 °C, and therefore the lowest mean oxide layer thicknesses (see Table 4.1), compared
to the other investigated 0.04B and 0.04B4Cr alloys. This indicates that an appropriate
amount of Cr as additional alloying element substantially increases the oxidation
resistance. With longer oxidation times and therefore thicker oxide layers on the sample
surface, further mass gain gradually decreases in case of all three alloys. Oxidation
mechanisms seem to be similar to the previously investigated Cr-free alloys. In this
respect, no further mass gain is measurable for the 0.04B8Cr alloy at 800 °C after
exposure times longer than approximately 300 h. After oxidation for 500 h at 900 °C, no
oxide layer thickness can be measured for this alloy (see Table 4.1) due to extensive scale
spalling. Figure 4.16 also shows that oxidation at 900 °C leads to significantly higher
mass gain and therefore thicker oxide layers compared to oxidation at 800 °C.
Figure 4.16: Isothermal oxidation of the different Cr-containing 0.04B, 0.04B4Cr, and
0.04B8Cr alloys at 800 and 900 °C in air.
62
Results
Figure 4.17 shows a comparison of SEM micrographs and EDX mappings of the
0.04B, 0.04B4Cr, and 0.04B8Cr alloys after oxidation in air for 500 h at 800 °C. The
0.04B8Cr alloy exhibits a continuous inner Al2O3 layer, whereas oxidation of the 0.04B
and 0.04B4Cr alloys results in discontinuous Al2O3. In addition, the formation of a
protective Cr-rich layer above the Al2O3 layer can be observed. The Cr-rich phase, which
slowly establishes with increasing oxidation time, seems to have a positive effect on the
formation of a continuous and thicker inner Al2O3 layer, as well. Furthermore, scale
delamination can be observed in the SEM micrographs of the investigated 0.04B and
0.04B4Cr alloys. Additional SEM/EDX experiments on the 0.04B, 0.04B4Cr, and
0.04B8Cr alloys after long-term oxidation at 900 °C reveal the same Cr-effect on
oxidation resistance and scale composition. In contrast to the observed improvements in
oxidation properties, Cr addition also leads to an altered γ/γ'-microstructure, i.e. the γ'phase gets instable and loses its cubic shape (Figure 4.18). In this regard, γ'-phases seem
to grow together in order to form larger agglomerates.
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63
Figure 4.17: Comparison of SEM micrographs and EDX mappings of the (a) 0.04B, (b)
0.04B4Cr, and (c) 0.04B8Cr alloys after oxidation for 500 h at 800 °C in air.
Figure 4.18: Comparison of the γ/γ'-microstructure of the (a) 0.04B, (b) 0.04B4Cr, and
(c) 0.04B8Cr alloys.
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Results
4.2.3 Effect of nickel addition
In Figure 4.19, representative oxidation curves of the 9W, 0.12B, 0.12B9Ni, and
0.12B18Ni alloys are summarised. Isothermal oxidation in air is carried out for 500 h at
800 and 900 °C. As already observed for the previously investigated B-free 9W and Bcontaining 0.12B alloys, Ni-containing 0.12B9Ni and 0.12B18Ni alloys exhibit an initial
transient period at both temperatures after short exposure times, as well. This period is
characterised by high mass gain, followed by nearly parabolic oxide layer growth.
Considerably lower oxidation resistances and therefore thicker oxide layers can be
observed by adding 9 at% of Ni to the 0.12B alloy (see Table 4.1). Since the presence of
Ni results in continuous oxidation of the alloy, further mass gain is not inhibited at
800 °C, as it is the case for the 0.12B alloy. Further increase of the Ni-content (18 at%)
leads to comparable oxidation behaviour of both Ni-containing alloys at 800 °C and a
slight decrease in oxidation resistance of the 0.12B18Ni alloy at 900 °C. However, layer
thickness cannot be determined after long-term oxidation of the 0.12B18Ni alloy at
800 °C due to severe oxide spalling. Compared to the 9W alloy, comparable oxidation
behaviour at 800 °C and improved oxidation resistance at 900 °C can be observed for the
Ni-containing alloys.
Figure 4.20 shows an SEM micrograph and EDX mappings of the 0.12B9Ni alloy
after long-term oxidation for 500 h at 900 °C in air. Since the Ni-containing samples
exhibit comparable oxidation behaviour, layer morphology, and layer composition, solely
results of the 0.12B9Ni specimen are presented in Figure 4.20. The inner oxide layer
consists of oxides of all alloy constituents, whereas predominantly Ni-rich oxides can be
detected at the alloy/oxide interface instead of the protective Al2O3 layer (see also
chapter 4.1). ToF-SIMS measurements on the same alloy reveal that, in contrast to the Nifree 0.12B sample, typical B-containing precipitates do not form within the inner oxide
layer or at the grain boundaries (Figure 4.21a). In addition, Figure 4.21b confirms the
presence of Ni-rich oxides at the oxide/alloy interface.
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65
Figure 4.19: Isothermal oxidation of the 9W, 0.12B, 0.12B9Ni, and 0.12B18Ni alloys at
800 and 900 °C in air.
Figure 4.20: SEM micrograph and EDX mappings of the 0.12B9Ni alloy after oxidation
in air for 500 h at 900 °C.
66
Results
Figure 4.21: ToF-SIMS mappings of the B (a) and Ni (b) distribution in the oxide layer
and at the grain boundaries of the oxidised 0.12B9Ni alloy (500 h at 900 °C in air).
4.2.4 Effect of silicon addition
Isothermal oxidation of the 0.12B and 0.12B2Si alloys was performed for 500 h at 800,
900, and 1000 °C in air. In Figure 4.22, thermogravimetric results are summarised. The
Si-containing alloy exhibits inferior oxidation resistance at 800 °C, characterised by
continuous oxidation and a higher mass gain, whereas oxidation at 900 and 1000 °C leads
to a significantly improved oxidation resistance and thus thinner oxide scales in
comparison with the Si-free 0.12B alloy (see also Table 4.1). Different EDX mappings of
the 0.12B2Si alloy (oxidised in air for 500 h at 800 °C (a + b), 900 °C (c + d), and
1000 °C (e + f)) in Figure 4.23 reveal that Si accumulates in the inner oxide layer and at
the grain boundaries due to high temperature oxidation. In addition, at elevated
temperatures (T ≥ 900 °C) higher amounts of Al can be detected at the oxide/alloy
interface compared to the same alloy oxidised at 800 °C. At 1000 °C, there are still
several Al-rich zones visible near the interface, which improve the oxidation resistance
compared to the Si-free material. However, scale spalling becomes an issue after
oxidation at 1000 °C (see crack in Figures 4.23e and f) due to high strains during sample
cooling in air.
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67
Figure 4.22: Isothermal oxidation of the 0.12B and 0.12B2Si alloys for 500 h at 800,
900, and 1000 °C in air.
Figure 4.23: EDX mappings (Al and Si) of the 0.12B2Si alloy after oxidation for 500 h at
800 °C (a + b), 900 °C (c + d), and 1000 °C (e + f) in air.
68
Results
In Figure 4.24, grain boundary precipitates of the oxidised 0.12B2Si alloy (500 h
at 900 °C in air) are characterised by means of SEM and EDX mappings. Adjacent to the
grain boundaries, high Co- and low W-contents indicate the presence of γ'-free γ-Co,
whereas Co- and Al-containing phases accumulate at the grain boundaries in addition to
W- and Si-rich precipitates.
Supplementary ToF-SIMS experiments were carried out for the same alloy (see
Figure 4.25). Mappings of the Si-distribution confirm the presence of Si-rich phases at the
grain boundaries, within precipitates, and at the oxide/alloy interface.
Figure 4.24: SEM micrograph and EDX mappings of grain boundary precipitates in the
0.12B2Si alloy after oxidation for 500 h at 900 °C in air.
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69
Figure 4.25: ToF-SIMS mappings of the Si distribution (near the sample surface (a) and
at the grain boundaries (b)) in the 0.12B2Si alloy after oxidation for 500 h at 900 °C.
4.2.5 Effect of yttrium addition
Long-term oxidation of the 0.12B0.005Y alloy was carried out isothermally for 500 h at
800 and 900 °C in air (Figure 4.26) in order to explore the influence of Y-addition on the
oxidation properties. Measured mean oxide layer thicknesses and standard deviations are
given in Table 4.1. Compared to the 0.12B alloy, 0.12B0.005Y shows similar oxidation
resistance at 800 °C, whereas inferior oxidation properties can be determined for the Ycontaining alloy at 900 °C. However, oxidation mechanisms of both materials seem to be
identical. Cross-sectional SEM micrographs and EDX mappings of the long-term
oxidised 0.12B0.005Y alloy in Figures 4.27 and 4.28 confirm the similarity in oxide layer
morphology and composition when comparing them to the results of the 0.12B sample.
Additional Y in the 0.12B0.005Y alloy diffuses mainly into the inner oxide layer, very
likely forming Y2O3. However, results indicate that Y-addition does not affect oxidation
resistance at 800 °C, but at 900 °C a negative effect can be observed. These findings are
contrary to frequently reported positive effects of Y on oxidation resistance of other
superalloys [17-21].
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Results
Figure 4.26: Isothermal oxidation of the 0.12B0.005Y and 0.12B alloys at 800 and
900 °C in air.
Figure 4.27: SEM micrograph and EDX mappings of the 0.12B0.005Y alloy after 500 h
at 800 °C in air.
Figure 4.28: SEM micrograph and EDX mappings of the 0.12B0.005Y alloy after 500 h
at 900 °C in air.
Results
71
4.2.6 Effect of titanium addition
Due to Ti addition to the 0.12B alloy, improvements in oxidation resistance can be
achieved especially at 900 °C. However, oxidation behaviour seems to be instable and
therefore no specific time dependency can be determined. In Figure 4.29 and Table 4.1,
results of the isothermal oxidation in air for 500 h at 800 and 900 °C are summarised.
Mean oxide layer thicknesses of 29 ± 5 µm (500 h at 800 °C) and 55 ± 9 µm (500 h at
900 °C) can be determined for the 0.12B2Ti alloy in contrast to 23 ± 4 µm (500 h at
800 °C) and 126 ± 18 µm (500 h at 900 °C) for the Ti-free 0.12B alloy (see also
chapter 4.1.1).
Figure 4.29: Isothermal oxidation of the 0.12B2Ti and 0.12B alloys at 800 and 900 °C
in air.
SEM micrographs and EDX mappings of the 0.12B2Ti alloy after long-term oxidation for
500 h at 800 and 900 °C are presented in Figures 4.30 and 4.31, respectively. Protective
72
Results
Al- and Ti-rich phases form at both temperatures. However, it seems that a completely
continuous Al2O3 layer does not develop at 800 °C, whereas at 900 °C a continuous but
relatively thin layer can be detected. In addition, EDX mappings show that Ti
accumulates at the oxide/alloy interface forming most likely protective Ti-rich oxides.
Higher oxidation temperatures seem to benefit the formation of a more continuous Ticontaining oxide layer.
Figure 4.30: SEM micrograph and EDX mappings of the 0.12B2Ti alloy after 500 h at
800 °C in air.
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73
Figure 4.31: SEM micrograph and EDX mappings of the 0.12B2Ti alloy after 500 h at
900 °C in air.
A summary of the presented alloying effects is given in Table 4.2 and will be
discussed in detail in chapter 5.
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Results
Table 4.2: Influence of alloying elements on oxidation resistance and microstructure.
Alloying
element
Influence on oxidation
Influence on microstructure
+ Improved inner Al2O3 formation
+ Formation of B-rich phases (mainly at
grain boundaries and in the inner oxide
layer)
+ Improves oxide layer adhesion
+ Formation of Cr-oxides
+ Improved Al2O3 formation
– Instability of the γ'-phase
– Impaired Al2O3 formation
– No borides at grain boundaries and in the
inner oxide layer
+ Formation of Si-rich phases (mainly at
grain boundaries and in the inner oxide
layer)
+ Improved Al2O3 formation
B
+ Strongly improved
oxidation resistance
Cr
+ Strongly improved
oxidation resistance
Ni
– Strongly impaired
oxidation resistance
Si
+ Improved oxidation
resistance at T ≥ 900 °C
Y
○ No positive effect at
800 °C
– Slightly impaired
oxidation resistance at
900 °C
○ Formation of Y-rich phases in the inner
oxide layer
Ti
+ Improved oxidation
resistance at 900 °C
+ Formation of Ti-oxides in the inner
oxide layer
+ Improved Al2O3 formation
4.3 Electrochemical studies on the quaternary Co–Al–W–B superalloy
4.3.1 Corrosion of the unoxidised alloy in aqueous solutions of different pH
Results of the OCP, EIS, and potentiodynamic j/E measurements of the unoxidised
Co–Al–W–B superalloy (0.12B) in aqueous solutions of different pH are presented in
Figures 4.32, 4.33, and 4.34. Characterisation is carried out in 0.5 M H2SO4 (pH 0.6),
1 M Na2SO4 (pH 5.9), and 0.1 M NaOH (pH 12.8) solutions. In each figure,
representative curves are presented and compared to the corrosion behaviour of pure Co.
At the beginning of exposure in acidic 0.5 M H2SO4 solution, measured potentials
of both specimens start at approximately –250 to –300 mV and remain almost constant
during the measurement (see Figure 4.32). After 3 h, the OCP value of the alloy is
Results
75
slightly higher compared to the OCP of pure Co. In neutral 1 M Na2SO4 solution, the
potentials start at –520 and –570 mV for pure Co and the superalloy, respectively (see
Figure 4.33). For pure Co, a temporary potential shift in anodic direction can be detected
with a maximum value after about 1.5 h of measurement. After 3 h in neutral solution, the
measured OCP of pure Co is slightly higher than the value of the alloy and compared to
the results obtained in acidic solution, OCP values are shifted in cathodic direction.
Figure 4.32: Electrochemical characterisation of the unoxidised alloy (a) and pure Co (b)
in 0.5 M H2SO4 aqueous solution (pH 0.6).
76
Results
At the beginning of exposure in 0.1 M NaOH (see Figure 4.34), the measured potentials
of pure Co and the alloy start at –600 and –400 mV, respectively, and then shift in anodic
direction until an approximately constant OCP is reached after 3 h of immersion. The
steady-state OCP value of pure Co is slightly higher than the OCP of the alloy, similar to
the obtained ranking of OCP values in neutral solution. However, potentials in alkaline
solution are more anodic than in neutral solution.
Figure 4.33: Electrochemical characterisation of the unoxidised alloy (a) and pure Co (b)
in 1 M Na2SO4 aqueous solution (pH 5.9).
Results
77
EIS measurements at the corrosion potential reveal that the total impedance of the
alloy is considerably enhanced in 0.1 M NaOH (approximately one order of magnitude)
compared to pure Co. However, in 1 M Na2SO4 impedance values of both materials are
comparable and in 0.5 M H2SO4 total impedance of the alloy is even one order of
magnitude lower compared to pure Co. Therefore, in case of the alloy a decreased
impedance can be observed with decreasing pH, whereas impedance values of pure Co
seem to be pH independent.
Figure 4.34: Electrochemical characterisation of the unoxidised alloy (a) and pure Co (b)
in 0.1 M NaOH aqueous solution (pH 12.8).
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Results
After carrying out OCP and EIS investigations, j/E measurements were performed
in all three aqueous solutions. In alkaline solution, the alloy shows better corrosion
resistance, i.e. lower corrosion current densities jcorr and higher corrosion potentials Ecorr
compared to pure Co. However, both materials possess primary and secondary
passivation regions. As for pure Co, oxygen evolution occurs after secondary passivation,
whereas for the alloy, current density abruptly increases at lower potentials compared to
pure Co. In neutral solution, pure Co and the alloy exhibit comparable corrosion
behaviour characterised by very limited passivity, in contrast to primary and secondary
passivity in alkaline solution. However slightly higher jcorr and higher Ecorr values can be
measured for pure Co. This is in good agreement with the obtained impedance data in
both solutions. Furthermore, small drops in current density are visible at higher potential
values. In acidic solution, both materials show active dissolution, i.e. a lack of
passivation, with comparable jcorr and Ecorr values and small drops in current density at
higher potentials. However, dissolution of the alloy is slightly decelerated due to a less
steep anodic branch in the measured polarisation curve (see Figure 4.32). As a conclusion
from the electrochemical investigations, the corrosion resistance increases with increasing
pH value of the aqueous solution, whereas the 0.12B alloy in alkaline solutions exhibits
the best corrosion behaviour.
4.3.2 Corrosion of the oxidised alloy in aqueous solutions of different pH
Figures 4.35, 4.36, and 4.37 compare the electrochemical behaviour of the oxidised
quaternary superalloy in three aqueous solutions of different pH, depending on prior
oxidation time and temperature. In each figure, representative curves are presented and
compared to the corrosion behaviour of the unoxidised alloy. A considerable increase in
OCP and total impedance of the oxidised compared to the unoxidised alloys can be
observed. After immersion in acidic solution, potential values of the long-term oxidised
alloys continuously decrease during the OCP measurement (Figure 4.35). After 3 h of
immersion, values of 200 mV and –130 mV can be measured for the 500 h/800 °C and
500 h/900 °C alloys, respectively. In contrast, after an initial exposure period, alloys
oxidised for 24 h at 800 °C and 24 h at 900 °C establish relatively constant OCP values
which amount to 175 mV and 140 mV, respectively.
Results
79
Figure 4.35: Electrochemical characterisation of the oxidised alloy in 0.5 M H2SO4
aqueous solution (pH 0.6). Samples are oxidised in air for 500 h at 800 °C (a), 500 h at
900 °C (b), 24 h at 800 °C (c), 24 h at 900 °C (d), and unoxidised (e).
In neutral solution, potential values of the oxidised alloys are in the same range as
in acidic solution and decrease with increasing exposure time, as well (Figure 4.36). In
alkaline solution, differences in OCP between differently oxidised alloys are rather low
and values are in the range of 0 to –100 mV (Figure 4.37). Furthermore, for the shortterm oxidised alloy (24 h) at 800 °C potential fluctuations can be observed during the
OCP measurement. In neutral and alkaline solutions, longer oxidation times and higher
temperatures lead to an increase of total impedance values and a decrease of corrosion
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Results
density. However, in acidic solutions only oxidation time seems to have a positive
influence on the measured impedance since oxidising the alloys for a specific oxidation
time at different temperatures results in more or less comparable impedance values and
potentiodynamic curves. Furthermore, the differently oxidised alloys reveal significantly
decreased capacities compared to the unoxidised alloy in all three utilised solutions of
different pH.
Figure 4.36: Electrochemical characterisation of the oxidised alloy in 1 M Na2SO4
aqueous solution (pH 5.9). Samples are oxidised in air for 500 h at 800 °C (a), 500 h at
900 °C (b), 24 h at 800 °C (c), 24 h at 900 °C (d), and unoxidised (e).
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81
Regarding the potentiodynamic curves, the highest decrease of corrosion rate can be
observed between the unoxidised alloy and the alloy oxidised for 24 h. Further oxidation
of the alloy prior to the electrochemical studies still increases the corrosion resistance but
differences between the short-term and long-term oxidised alloys (or between specimens
oxidised at different temperatures) are less significant than differences between the
unoxidised 0.12B alloy and any of the oxidised specimens.
Figure 4.37: Electrochemical characterisation of the oxidised alloy in 0.1 M NaOH
aqueous solution (pH 12.8). Samples are oxidised in air for 500 h at 800 °C (a), 500 h at
900 °C (b), 24 h at 800 °C (c), 24 h at 900 °C (d), and unoxidised (e).
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Results
A noteworthy feature of the j/E curves of oxidised 0.12B alloys are the remarkably flat
cathodic and anodic branches measured in all three solutions. In addition, an onset of
oxygen evolution cannot be detected in the anodic region. No considerable differences in
corrosion properties can be observed for the oxidised alloys in neutral and alkaline
aqueous solution. However, long-term oxidised alloys in acidic solution reveal improved
corrosion resistance, i.e. higher total impedance and lower jcorr values, compared to the
alloys immersed in neutral and alkaline solution. As a result, the alloy oxidised for 500 h
at 800 °C shows the best corrosion behaviour of all investigated alloys in the different
solutions.
4.3.3 Corrosion of unoxidised alloys in chloride containing solution
In addition to electrochemical characterisation in solutions of different pH, two superalloy
variants and pure Co are investigated in chloride containing 0.5 M NaCl aqueous solution
(pH 5.8) in order to study the susceptibility to pitting corrosion. In Figure 4.38,
representative curves of the OCP, EIS, and potentiodynamic measurements of unoxidised
0.12B0.005Y (a), 0.12B (b), and pure Co (c) are presented. At the beginning of
immersion, the measured potentials of all samples are in the range of –550 to –600 mV,
and then rapidly shift in anodic direction. However, during the measurement, the potential
of pure Co drops significantly after about 5 minutes of immersion. After 3 h,
approximately constant OCP values are reached and the steady-state potentials of both
alloys (0.12B and 0.12B0.005Y) are more anodic (–370 mV and –355 mV) than the
potential of pure Co (–435 mV). EIS measurements show that at the corrosion potential,
total impedance values are comparable and amount to approximately 4 to 5 kΩ·cm2 in
case of all three samples. In comparison, pure Co and the 0.12B alloy exhibit comparable
impedance spectra in neutral (chloride-free) sodium sulphate solution, as well.
Subsequent potentiodynamic j/E studies (Figure 4.38) reveal comparable corrosion
properties of pure Co, 0.12B, and 0.12B0.005Y alloys, characterised by limited
passivation at potentials between –600 mV and –300 mV, which is followed by strong
corrosion at more anodic potentials. In addition, a temporary current drop at high
potentials can be measured, most pronounced for pure Co. This seems to be similar to the
measured drops in acidic and neutral solution.
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83
Figure 4.38: Electrochemical characterisation of the unoxidised 0.12B0.005Y (a), 0.12B
(b), and pure Co (c) in 0.5 M NaCl aqueous solution (pH 5.8).
According to SEM and Optical Microscopy images (see Figure 4.39 a-d), which
are taken after performing electrochemical investigations, the measured increase in
current density at anodic potentials of at least –300 mV, i.e. after the limited initial
passivation, can be ascribed to severe pitting of all materials studied in chloride
containing solution. The observed pits are homogeneously distributed over the metal
surface
without
any
preferred
initiation sites.
Supplementary
electrochemical
measurements with short polarisation in the pitting region and subsequent optical
microscopy (not shown in Figures 4.38 and 4.39) also indicate no preferred initiation sites
of the pits. Figure 4.39b reveals that the formed pits possess a wide and shallow
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Results
morphology and do not penetrate deep into the investigated metals. Based on the
observed pit morphology, which is comparable for all the investigated metals in NaCl
solution, pit propagation seems preferential in the lateral direction.
Figure 4.39: Optical Microscopy images of 0.12B0.005Y (a), 0.12B (c), and pure Co (d)
after performing electrochemical investigations. An additional SEM micrograph of the
0.12B0.005Y alloy is given in (b).
4.3.4 Corrosion of oxidised alloys in chloride containing solution
Since Co-base superalloys are promising materials for high temperature applications,
oxidation of two novel alloys (0.12B0.005Y and 0.12B) is carried out for 65 h at 800 and
900 °C in air, prior to electrochemical characterisation in 0.5 M NaCl aqueous solution.
For comparison, pure Co is pre-oxidised at 900 °C for 65 h and characterised in NaCl
solution, as well. In Figure 4.40, representative curves of the OCP, EIS, and j/E
measurements are summarised. Since representative oxide layers already form during
Results
85
short-term oxidation, exposure of the materials was limited to 65 h instead of performing
long-term experiments for 500 h in air.
Figure 4.40: Electrochemical measurements of (a) oxidised 0.12B0.005Y (65 h at
900 °C), (b) 0.12B (65 h at 900 °C), (c) 0.12B0.005Y (65 h at 800 °C), (d) 0.12B (65 h at
800 °C), and (e) pure Co (65 h at 900 °C) in 0.5 M NaCl aqueous solution.
After an initial potential drop, constant OCP values can be reached for all
investigated metals after 3 h of immersion in chloride containing solution. Due to a higher
oxidation temperature, considerable increase of the measured OCP and the total
impedance can be observed. In addition, a decrease of current density of approximately
one order of magnitude can be measured. Oxidised pure Co seems to be less corrosion
86
Results
resistant in chloride containing solution compared to the alloys, whereas both alloys show
comparable corrosion properties. Similar to previously presented electrochemical results
of oxidised alloys, polarisation curves again show flat cathodic and anodic branches over
the whole potential range from –1 to +2 V (vs. Ag/AgCl), especially in case of the
oxidised 0.12B0.005Y and 0.12B alloys.
Further electrochemical experiments were carried out for an oxidised Sicontaining alloy (0.12B2Si) and for a short-term oxidised 0.12B0.005Y alloy (1 h at
900 °C) in order to investigated the influence of an additional alloying element and other
oxidation conditions on the previously observed extraordinary protection against pitting
corrosion. In Figure 4.41, results (OCP, EIS, and j/E) of the (a) oxidised 0.12B0.005Y
(65 h at 900 °C), (b) 0.12B2Si (65 h at 900 °C), and (c) 0.12B0.005Y (1 h at 900 °C)
alloys in chloride containing solution are presented. Si-addition leads to further
remarkable improvements of the corrosion protection, i.e. higher impedance and lower
current density values, compared to the Si-free alloy. Another interesting finding is the
effect of short-term oxidation (1 h at 900 °C) on the corrosion resistance of the
0.12B0.005Y alloy in chloride containing solution. Results show that even pre-oxidation
of 1 h is sufficient for enabling good protective properties of the formed oxides without
passivity breakdown. In fact, EIS and j/E results of both alloys oxidised for 1 h and 65 h
at 900 °C do not show significant differences.
Results
87
Figure 4.41: Electrochemical measurements of the (a) oxidised 0.12B0.005Y (65 h at
900 °C), (b) 0.12B2Si (65 h at 900 °C), and (c) 0.12B0.005Y (1 h at 900 °C) alloys in
0.5 M NaCl aqueous solution.
88
Discussion
5. Discussion
5.1 Oxidation properties of Co-base superalloys
High temperature oxidation of the quaternary Co–Al–W–B superalloy at 800 °C and
900 °C in air indicates the presence of a transient period at short exposure times followed
by a “steady-state” oxidation period. Generally, oxidation at a higher temperature leads to
the formation of thicker oxide scales due to higher reaction kinetics, since the diffusion
coefficient D of metal ions (for instance Co2+) exhibits an Arrhenius dependency
according to equation (2.4) and therefore increases with higher temperatures. However,
the formation of different protective oxides, such as Al2O3, Cr2O3, or SiO2, may have a
tremendous influence on the total scale thickness and mass gain of the alloy system,
depending on the content of the corresponding alloying element. Those effects will be
discussed in detail in the course of this chapter.
At both oxidation temperatures, the reaction kinetics of the transient oxidation
period are more rapid compared to the steady-state oxidation. Since outward diffusion of
metal ions is hindered by the formation of a thick, compact, and adherent oxide layer on
the alloy surface, the measured mass gain (or oxide layer growth) decelerates with
increasing oxidation time. The transient period is characterised by oxidation of the alloy
constituents, i.e. Co, Al, and W. CoO formation at 900 °C is due to exceeding the
decomposition temperature of Co3O4 [75]. At steady-state conditions, the different
formed oxides coexist and show the typical multi-layered structure. Ion diffusion in
outward direction is hindered by the development of the innermost Al2O3 layer, leading to
further inhibition of oxide growth. Initially, Al2O3 forms at isolated spots (see Figures 4.8
and 4.9), which broaden with increasing oxidation time in order to form the protective
layer. Depending on the moment of Al2O3 formation, varying scale thicknesses above the
Al2O3 can be observed, which may be the reason for the curved alloy/oxide layer
interface. The initiation sites of Al2O3 formation are very likely governed by local
inhomogeneities of the Al concentration within the alloy, for instance caused by the
formation of dendrites during the cooling process after casting [28] and by an insufficient
ensuing solution heat treatment. In case of Ni-base superalloys, interdendritic regions are
Discussion
89
reported to exhibit higher Al concentrations compared to the innerdendritic core
regions [28]. This may be similar for Co-base superalloys and therefore explain the
preferred initiation sites of Al2O3 formation. By oxidising the Co–Al–W–B alloy at
800 °C, a continuous and protective inner Al2O3 layer forms after approximately 100 h of
exposure. The formed Al2O3 layer possesses a slow growth rate, high stability, and a
compact structure with very few defects [10]. In contrast, oxidising the alloy at 900 °C in
air causes the formation of a discontinuous Al2O3 layer without comparable protective
effect, which leads to a continuous oxidation of the superalloy. The high diffusion rate of
mainly Co ions in outward direction may be a reason for the disturbed formation of a
continuous inner Al2O3 layer at 900 °C. Due to the alloying elements, at which Al plays
the most important role, oxidation of the quaternary Co–Al–W–B alloy at 900 °C exhibits
a sub-parabolic time dependency. The presence of B within the alloy is crucial for the
Al2O3 layer formation, which will be discussed in detail in chapter 5.2.1. In contrast to the
Co-base alloy, pure Co shows highly decreased oxidation resistance and parabolic oxide
layer growth [64-66, 68-71] due to the lack of alloying elements, which may form oxides
(such as Al2O3) acting as barriers against further oxidation.
In addition to the innermost Al2O3 and the outer Co-oxide layer, an amorphous or
nanocrystalline middle oxide layer can be observed by ECCI (see Figure 4.4). Therefore,
it is difficult to determine if the crystalline ternary oxides detected by XRD (Figure 4.6)
are present in the middle oxide layer (if this layer is nanocrystalline) or if small fractions
may also exist within the outer Co-oxide layer (if the middle layer is amorphous).
Nevertheless, an amorphous structure of the middle oxide layer would be favourable with
respect to oxidation resistance, since outward ion diffusion is expected to be decelerated
compared to diffusion in a crystalline layer.
For a better understanding of the formation of multi-layered oxide structures,
thermodynamic studies were recently performed by Wu et al. [85] on a Ni–26Cr–4Al
(weight %) ternary alloy. It was demonstrated that the layering of different oxides within
the total oxide scale can be predicted. The most stable oxide (here Al2O3) exists directly
above the metal/oxide interface and the least stable oxide (NiO) represents the outermost
layer. A comparison between calculated data and experimental results shows very good
agreement. The principles of this study can be applied to oxide structures on novel Co-
90
Discussion
base superalloys, as well. In the present work, a Co–Al–W base alloy is under
investigation and results show that, in accordance to Wu et al. [85], the most stable oxide
(Al2O3) forms at the innermost metal/oxide interface whereas the least stable oxide (CoO
and Co3O4, respectively) represents the outermost layer. In between, several spinels can
be found which constitute the middle mixed oxide layer. By this method, it can also be
predicted how additions of Cr, Ni, Si, or Ti would affect the layering structure of the
oxide scale. However, a rough estimation can also be drawn by regarding the stabilities of
the important oxides at different temperatures and oxygen pressures, which can be read
off the Ellingham-Richardson diagram in Figure 2.5 [73]. The corresponding rate
constants of oxide formation are summarised in Figure 2.6 [3, 76] and can help
determining the resulting layer thicknesses.
In addition to the formation of the oxide layer itself, oxidation of the quaternary
Co–Al–W–B superalloy causes a pronounced phase transformation at the alloy/oxide
layer interface. Due to a continuously decreasing Al content of the alloy matrix during
oxidation, the Al depleted area transforms from a two-phase γ/γ' to a two-phase γ/Co3W
microstructure adjacent to the formed Al2O3. In the Al depleted region, concentration of
Al is less than 7 at% and therefore the γ' phase becomes instable. As a result, the observed
elongated Co3W phases form. In Figure 5.1 [5, 120], the partial phase diagram of the
ternary Co–Al–W system at 900 °C illustrates the origin of this phase transformation.
Previous studies of Kobayashi et al. [120] and Tsukamoto et al. [121] on phase equilibria
in ternary Co–Al–W alloy systems reveal similar precipitate microstructures compared to
the investigated alloys of the present work. Figure 5.2 [120] shows a representative BSE
image of the precipitates after heat treatment (2000 h at 900 °C). In accordance with
Figure 4.2b, transformation from γ/γ' to γ/Co3W can be observed. However, oxidation of
novel Co-base superalloys does not lead to the formation of an additional CoAl phase
adjacent to γ/Co3W, as observed in Figure 5.2. This is because of outward diffusion of Al
in order to form Al2O3, instead of CoAl.
Discussion
91
Figure 5.1: Partial phase diagram of the ternary Co–Al–W system at 900 °C [5, 120]. Al
depletion leads to the formation of a γ/Co3W microstructure.
Figure 5.2: BSE image of the precipitates observed in a Co–Al–W system heat treated
for 2000 h at 900 °C [120].
92
Discussion
Beyond that, the formation of a so called secondary reaction zone (SRZ) beneath
aluminide coatings in several single crystal Ni-base superalloys has been reported by
Das et al. [122]. Furthermore, an interdiffusion zone (IDZ) develops between the SRZ
and the coating during heat treatment (see Figure 5.3 [122]).
Figure 5.3: Cross-sectional microstructure of a coated single crystal “F-26” Ni-base
superalloy [122].
Discussion
93
The observed elongated phases within the SRZ possess very similar morphologies
compared to precipitates in Figure 4.2b. However, phases in Figure 5.3 (so called P TCP
phases) terminate on the γ-phase of the γ/γ'-substrate and the area in between is reported
to be the γ'-phase. This differs from the precipitation process observed in the present
work. Furthermore, the SRZ is believed to form as a result of inward Al diffusion from
the aluminide coating during high temperature exposure [123]. Das et al. also suggest that
the P TCP phase precipitates from the supersaturated γ-matrix in order to form a three
phase γ-γ'-P structure [122]. In case of precipitates which form in novel Co-base
superalloys, outward diffusion of Al is the reason for the observed transformation.
Therefore, resulting morphologies of the two precipitation processes are very similar but
transformation mechanisms and resulting phases differ considerably. In the SRZ observed
by Das et al. [122], γ'-phases are detected in addition to TCP phases. In the present study,
no γ'-phases can be detected within the transformation zone and therefore resulting
mechanical properties of the oxidised superalloy may be impaired by the lack of the
strengthening γ'-phase.
5.2 Influence of alloying elements
5.2.1 Effect of boron addition
The motivation of B-addition to ternary Co–Al–W superalloys is based on its positive
effect on the mechanical properties of Ni-base superalloys at high temperatures, such as
enhanced creep and rupture strength, as well as improved ductility due to Baccumulations at grain boundaries [3, 6, 124-126]. This strengthening effect has been
reported for Ni3Al intermetallic compounds, as well [127]. However, the present study
shows that also enhanced oxidation resistance and scale adhesion can be achieved by
appropriate amounts of B within Co-base superalloys. Alloys containing 0.12 at% of B
exhibit the best results compared to 0.08B, 0.04B, and 9W alloys. B-accumulation within
the middle oxide layer can be considered as the main reason for improved scale adhesion
due to the formation of B-containing intermetallic phases (such as tungsten borides)
which enhance cohesion [128]. This is also the reason for grain boundary strengthening in
94
Discussion
high temperature alloys due to B addition. Furthermore, B seems to benefit the formation
of protective inner Al2O3 at the alloy/oxide layer interface due to selective oxidation of
Al. As a result, a transition from rather parabolic oxidation behaviour of 9W to
subparabolic behaviour of B-containing alloys can be observed. However, Al
concentration in the 9W alloy is decreased by approximately 0.7 to 0.8 at% compared to
the other B-containing alloys (see Table 3.1). This may have a negative effect on Al2O3
formation and therefore be an additional reason (besides the lack of B) for impaired
oxidation resistance of 9W alloys. B-contents higher than 0.12 at% may lead to further
improvements of the oxidation resistance compared to the investigated 0.04B, 0.08B, and
0.12B alloys due to enhanced selective oxidation of Al and therefore stronger inner Al2O3
formation. As a result, a continuous Al2O3 layer may also form at 900 °C and shorter
oxidation times may be sufficient in order to establish a continuous Al2O3 layer at 800 °C.
However, it is recommended that the content of B should be lower than 1 at% in order to
avoid embrittlement and the formation of low-melting oxides [82]. In addition to the
observed B-rich phases within the middle oxide layer, ToF-SIMS investigations show that
precipitates in the alloy matrix exhibit high amounts of B and therefore may improve
oxidation resistance and high temperature strength, as well. In this respect, B-rich
precipitates (most likely tungsten borides) may also act as obstacles for ion diffusion and
dislocation movement. It can be assumed that B-depletion underneath the middle oxide
layer does not have a negative effect on the mechanical properties. However, the already
mentioned elongated Co3W precipitates may decrease high temperature strength due to a
lack of the γ'-phase. Ojo et al. [126] and Zhang et al. [129] recently reported the
formation of fine M5B3-type borides along the grain boundaries of the directionally
solidified (DS) superalloy Rene 80. Those boride types have been observed in Hf-doped
IN 792 [130] and Ni–Al–Cr–B alloys [131], as well. In addition, well known M3B2-type
borides, which form during ingot solidification [132-134], have been detected within the
grains of the Rene 80 alloy. As a conclusion, it can be assumed that fine borides in the
middle oxide layer and at the grain boundaries of the investigated Co-base superalloys
mostly consist of W5B3 and larger borides in the grains of W3B2.
Discussion
95
5.2.2 Effect of chromium addition
High temperature oxidation of Cr-containing Co–Al–W–B–Cr alloys at 800 °C and
900 °C reveals that an appropriate amount of Cr as additional alloying element leads to
improved oxidation properties due to Cr-rich oxide scales and the positive effect of Cr on
the formation of protective inner Al2O3. Compared to conventional corrosion and
oxidation resistant Co-base superalloys, such as MAR-M-509, which exhibit Cr-contents
of approximately 25 at%, less than 10 at% of Cr in novel Co-base superalloys is sufficient
in order to achieve satisfying oxidation protection due to the formation of Al2O3 and Crrich oxides. It is assumed that Cr-rich oxides on 0.04B4Cr and 0.04B8Cr alloys are
mainly composed of CoCr2O4 and protective Cr2O3, according to previous research on a
γ'-free Co-base superalloy [135]. As a result, scale growth on the 0.04B8Cr alloy is highly
decelerated in comparison with the 0.04B4Cr and 0.04B alloy. Previous studies on
Co(Ni)-Cr-Al alloys [83, 84] reported an inner double layer of protective Cr2O3 and
Al2O3 beneath an outer Co(Ni)O layer, similar to the oxide layering on novel Co-base
superalloys. New findings of Wu et al. [85] confirm this behaviour and demonstrate that
based on thermodynamic calculations, the layering of different oxides on complicated
alloy systems can be predicted in excellent agreement with experimental observations.
SEM micrographs reveal scale delamination on the surface of 0.04B4Cr and
0.04B alloys after removing the specimens from the oxidation furnace. However, this
finding is due to the low B-content of the alloys and not due to a decrease in Cr-content.
Improved scale adhesion in case of 0.04B8Cr alloys may be ascribed to the relatively low
scale thickness compared to 0.04B4Cr and 0.04B alloys, which leads to reduced strain
during the cooling process. However, during oxidation increased scale spalling occurs for
the 0.04B8Cr alloy, which can be recorded by thermogravimetry as an interim decrease in
mass gain (see Figure 4.16).
Even though Cr-addition leads to considerable improvements concerning
oxidation resistance of Co-base superalloys, several drawbacks have to be highlighted, as
well. Protection of Cr2O3 is reported to be decreased at temperatures higher than 1000 °C
and in oxidising atmospheres with high flow rates due to the formation of volatile CrO3
species [23-25]. Furthermore, SEM studies on Cr-containing alloys (Figure 4.18) reveal
an altered γ/γ'-microstructure compared to the Cr-free alloy, which is expected to reduce
96
Discussion
high temperature creep properties considerably. Therefore, γ'-stabilisation is necessary by
means of additional alloying elements such as Ni, Mo, or Ta [16], in order to maintain
high temperature creep strengths which are already comparable to Ni-base
superalloys [8]. With respect to oxidation resistance, a lack of γ'-precipitates due to prior
dissolution seems to impair diffusion processes and therefore improve oxidation
properties. Supplementary experiments on two different 0.08B alloys for 500 h at 800 °C
(see Figure 5.4) confirm this assumption. One sample was solution heat treated and aged
in the same way as all other samples of this study, whereas the comparison specimen was
solely solution heat treated in order to avoid γ'-formation. Results show that the γ'-free
0.08B alloy exhibits considerably improved oxidation resistance due to lower ionic
diffusion rates. For this reason, a completely continuous Al2O3 layer is expected to be
formed after exposure times longer than 500 h and therefore mass gain of the γ'-free
0.08B alloy is still slightly increasing in Figure 5.4.
Figure 5.4: Comparison of a γ'-containing 0.08B alloy and a γ'-free 0.08B alloy, obtained
by different heat treatments.
Discussion
97
The ternary Co–17Re–23Cr alloy is an experimental alloy recently investigated by
Mukherji et al. [136] with poor ductility in the polycrystalline state due to grain boundary
failure. The addition of B changes the fracture mode from intergranular to transgranular
and therefore dramatically improves the strength and ductility. Even small amounts of B
(0.02 at%) leads to the formation of fine Cr2B-type borides at the grain boundaries. As a
result, a strengthening effect can be observed similar to the B-containing Co-base alloys
investigated in this study and by Bauer et al. [8, 9]. However, it is reported that borides
are not present at all grain boundaries of the experimental Co–Re–Cr–B alloys [136]. It
can be assumed that in addition to the previously discussed occurrence of W5B3 and W3B2
borides, fine Cr2B-type borides may also be present at the grain boundaries of novel Cobase superalloys of the system Co–Al–W–B–Cr.
A comparison of the investigated B- and Cr-containing alloys with the
conventional γ'-free Co-base superalloys MAR-M-509 reveals several notable differences
in oxidation behaviour. Figures 5.5 and 5.6 show cross-sectional SEM micrographs and
EDX mappings of the MAR-M-509 alloy after oxidation for 100 h at 800 °C. In addition
to considerable amounts of Cr-rich species within the inner oxide layer, which provide
excellent oxidation resistance at high temperatures [135], most likely NiO and Co3O4
species are present within the mixed outer layer. Similar to the investigated γ'strengthened Co-base alloys, complex oxides may form especially in the inner layer, as
well. Precipitates containing Ta, Ti, and Zr (see Figure 5.6), which have previously been
reported by Beltran et al. [137], are necessary for precipitation strengthening of the
MAR-M-509 alloy [135]. Larger precipitates have grown at the expense of smaller ones
and dominate the microstructure of the investigated alloy (see Figure 5.5). Because of
non-uniform oxide layer formation and scale spalling no reasonable mean oxide layer
thickness can be calculated. Compared to the oxidation behaviour of the different Bcontaining alloys (0.04B, 0.08B, and 0.12B), MAR-M-509 exhibits superior oxidation
resistance at 800 °C in air. However, compared to the Cr-containing alloys, 0.04B8Cr
shows comparable mass gain due to high-Cr contents whereas 0.04B4Cr possesses
inferior oxidation resistance. An increase in B-content of both Cr-containing alloys may
even further enhance oxidation resistance and scale adhesion. As a result, experimental
98
Discussion
0.12B8Cr alloys are expected to outrange conventional Co-base superalloys, such as
MAR-M-509, with regard to oxidation resistance at 800 °C in air.
Figure 5.5: SEM micrographs of MAR-M 509 after oxidation for 100 h at 800 °C in air.
Discussion
99
Figure 5.6: SEM micrograph and EDX mappings of MAR-M 509 after oxidation for
100 h at 800 °C in air.
100
Discussion
5.2.3 Effect of yttrium addition
It is reported that small amounts of rare earth elements, such as Y, enhance oxide layer
adhesion and increase oxidation resistance of the alloy due to selective Al oxidation [1521]. However, present results reveal no measurable improvements due to the addition of
0.005 at% of Y to the 0.12B alloy. Previously, Yu et al. investigated the oxidation
behaviour of the Ni-base superalloy K38 and showed that alloys containing 0.1 wt% of Y
exhibit excellent oxidation behaviour compared to Y-free and 0.05 wt% containing
alloys [17]. However, excessive amounts of Y (0.5 wt%) decrease oxidation resistance of
the K38 alloys due to the formation of Y-rich phases. Therefore, a further increase in Ycontent of the current 0.12B0.005Y alloy may help improving the oxidation resistance
compared to the 0.12B alloy. Furthermore, EDX measurements revealed that the Y-free
alloy possesses a higher Al-content, which is assumed to result in slightly improved
oxidation properties.
5.2.4 Effect of silicon and nickel addition
In contrast to undesired modifications of the γ/γ'-structure due to Cr addition, Si and Ni
additions to the 0.12B alloy lead to comparable high volume fractions and similar sizes of
the cubic γ'-phase, as shown in the supplementary SEM micrographs (Figure 5.7).
However, as a result of Ni- and Si-addition considerable differences in oxidation
behaviour and grain boundary precipitation can be observed. In case of 0.12B alloys,
borides at the grain boundaries and in the inner oxide layer lead to improved creep
properties [8, 9], oxidation resistance, and scale adhesion. Additions of 9 at% and
18 at% of Ni to the 0.12B alloy, however, seem to enhance solubility of B within the
alloy matrix and therefore an absence of B at grain boundaries and the inner oxide layer is
detected by ToF-SIMS. The resulting decrease in layer cohesion leads to the observed
oxide spalling on the high Ni-containing 0.12B18Ni alloy after oxidation for 500 h at
800 °C. SEM/EDX investigations show that the B-effect on Al2O3 formation is
suppressed by Ni-addition and therefore no continuous inner Al2O3 layer can be detected.
However, Ni-rich species (Ni-containing spinels or NiO) are present at the alloy/oxide
layer interface. Since no protective inner Al2O3 scale establishes, nearly parabolic oxide
growth can be observed by thermogravimetry, comparable to the oxidation behaviour of
Discussion
101
the ternary 9W alloy (Co–Al–W) at 800 °C. However, at 900 °C, both Ni-containing
alloys seem to be more oxidation resistant than the 9W alloy. A lower Al-content of 9W
(Table 4.1) may be a possible explanation for this finding. Furthermore, lower mass gain
and a simultaneously thicker oxide scale is measured on the 0.12B9Ni alloy in
comparison with the 0.12B18Ni alloy (after 500 h at 900 °C) probably due to higher
oxide porosity and different oxide composition. According to the measured values in
Table 4.1, a higher Al-content and lower Ni-content of 0.12B9Ni may lead to increased
Al2O3 formation but less Ni oxides, compared to oxidised 0.12B18Ni. As a conclusion of
the present findings, Ni-additions of 9 at% and 18 at% to Co–Al–W–B alloys is not
suitable for improving oxidation resistance and high temperature strength due to the
observed negative effect of Ni on B-distribution and Al2O3 formation. Even though Ni is
reported to be more oxidation resistant than Co (the rate constant of NiO growth is two
orders of magnitude lower than for CoO according to Figure 2.6 [3, 76]), no evidence of a
positive influence can be found based on the present findings.
Figure 5.7: SEM micrographs of the γ/γ'-microstructure of a) 0.12B9Ni, b) 0.12B18Ni, c)
0.12B2Si, and d) 0.12B alloys.
102
Discussion
However, further small additions of B or additional strengthening elements, such as Zr or
C [16], may be a way to compensate the presented negative effects of Ni, since Niaddition is important for broadening the two-phase γ/γ'-area of the novel Co-base
superalloys [50]. In order to improve oxidation resistance of Ni-containing Co–Al–W–B
alloys, additions of Si is recommended due to the formation of protective and stable
oxides [26-28] and due to the fact that, in contrast to Cr, it does not alter the γ/γ'microstructure of the alloy (see Figure 5.7).
At high temperatures, i.e. 900 and 1000 °C, the investigated Si-containing
0.12B2Si alloy exhibits improved oxidation resistance compared to the 0.12B alloy due to
the formation of Si-containing phases (according to literature mostly SiO2) [26, 27] and
the beneficial effect of Si on inner Al2O3 formation [28], similar to the mechanism of Cr.
The 0.12B2Si alloy exhibits a lower Al-content compared to the Si-free 0.12B alloy
(Table 4.1). Nevertheless, improvements in oxidation resistance are notable. Si-rich
phases which can be detected at the alloy/oxide layer interface, within the grains in terms
of precipitates, and at the grain boundaries, are assumed to act as diffusion barriers and
therefore increase oxidation resistance. However, a decrease in creep strength at 850 °C
was reported for novel Co–Al–W–Si superalloys due to Si addition [9]. Recent creep
experiments of A. Bauer reveal that also 0.12B2Si alloys possess worse high temperature
properties compared to Si-free 0.12B alloys because of lower ductility. At 1000 °C,
further decrease in high temperature strength is expected since long-term oxidation of the
0.12B2Si alloy leads to γ'-dissolution due to exceedance of the γ'-solvus temperature. Taaddition is reported to increase the γ'-solvus temperature of Co–Al–W superalloys [138]
and therefore may help stabilising the γ'-phase of the investigated alloys at 1000 °C.
Furthermore, higher B-contents and additions of rare earth elements, such as Y, are
expected to improve scale adhesion on 0.12B2Si alloys at such temperatures.
5.2.5 Effect of titanium addition
High temperature oxidation of the 0.12B2Ti alloy at 800 and 900 °C reveals that
significant improvements at 900 °C compared to the Ti-free alloy can be ascribed to the
formation of a thin and continuous inner Al2O3 layer in addition to Ti-rich phases at the
alloy/oxide layer interface. Therefore, Ti seems to influence oxidation properties of 0.12B
Discussion
103
alloys similar to Si and Cr, i.e. it forms protective Ti-rich phases (most likely Ticontaining ternary oxides and TiO2) and exhibits a positive effect on Al2O3 formation at
oxidation temperatures of at least 900 °C. In contrast to the 0.12B2Si alloy, no negative
influences on high temperature strength are reported due to Ti-addition [9]. According to
Bauer et al. [9], Ti-addition to 0.12B does not alter the γ/γ'-microstructure and 0.12B2Ti
exhibits even the best creep properties of all investigated alloys at 850 °C with values
comparable to the Ni-base superalloy IN 100. Therefore, optimisation of the 0.12B2Ti
alloy may be the most promising way in order to obtain superalloys with sufficient
oxidation resistance and mechanical properties.
5.2.6 Summary of the alloying effects
In Table 4.1, the influence of alloying elements on oxidation resistance of novel Co-base
superalloys is summarised. For this, the mean oxide layer thicknesses (and the standard
deviations) are determined and compared after oxidation experiments for 500 h at 800,
900, and 1000 °C in air. Based on the knowledge of conventional Co- and Ni-base
superalloys, not all alloying elements behave as expected. For instance, B is well known
for improving the mechanical properties of Ni-base superalloys at high temperatures due
to B-segregation at the grain boundaries [3, 6, 123-125]. However, results of the present
study also reveal significantly enhanced oxidation properties of B-containing Co-base
superalloys due to a positive effect of B on inner Al2O3 formation. Furthermore, Baccumulations within the inner oxide layer (most likely tungsten borides) are assumed to
be the reason for improved scale adhesion. Additional B-rich precipitates at the grain
boundaries and within the grains may not only act as obstacles for dislocation movement
but also hinder ion diffusion and hence further improve oxidation resistance.
Results of the investigated Cr-containing samples confirm the expected oxidation
properties due to the formation of protective Cr-rich oxides and the positive effect of Cr
on inner Al2O3 formation. However, in contrast to conventional Co-base superalloys, the
present results show that it is sufficient to add less than 10 at% of Cr in order to achieve
an appropriate oxidation resistance. Since Cr is known for stabilising the fcc matrix [5],
the observed instability of the γ'-phase due to Cr-addition was expected.
104
Discussion
According to literature [15-21], small amounts of Y improve oxide layer adhesion,
especially under cyclic conditions, and enhance oxidation resistance due to selective Al
oxidation. In the present study, oxidation properties of the investigated 0.12B0.005Y
alloy are comparable to the Y-free 0.12B alloy and therefore differ from information
presented in literature. However, those findings are most likely attributed to the low
amount of Y in the 0.12B0.005Y alloy and hence a higher Y-content is recommended for
future investigations.
Additions of Ni to 0.12B alloys are expected to be beneficial since Ni is reported
to be more oxidation resistant than Co [3, 76]. However, results of this study reveal that
9 and 18 at% of Ni strongly impair oxidation resistance of novel B-containing alloys. It is
assumed that Ni has a negative effect on B-accumulation within the inner oxide layer and
at the grain boundaries. As a consequence of higher B-solubility in the matrix,
significantly less Al2O3 can be detected at the alloy/oxide interface. This finding is new
and could not have been predicted from previous knowledge.
Si as additional alloying element in 0.12B alloys not only benefits the formation
of Si-containing phases (at the alloy/oxide interface, within precipitates, and at the grain
boundaries) which may act as diffusion barriers, but also improves the generation of
Al2O3 at oxidation temperatures of at least 900 °C. In literature, this behaviour was
already reported for Ni-base superalloys [26-28]. However, for novel γ'-strengthened Cobase superalloys, the influence of Si on the oxidation behaviour was not investigated, so
far.
The investigation of Ti-containing alloys reveals significantly improved oxidation
resistance at 900 °C compared to Ti-free 0.12B alloys. Ti-addition is assumed to benefit
inner Al2O3 formation, similar to the effect of Si-addition. Furthermore, protective Ti-rich
phases (TiO2 and ternary oxides) are formed in the inner oxide layer. However, in
contrast to the other investigated alloy variants, Ti-containing Co-base superalloys are the
most promising alloy systems for further investigation and optimisation due to good
mechanical and oxidation properties. Hence, the results of the performed oxidation
experiments on novel 0.12B2Ti alloys provide important information for the development
of the mentioned new alloy system.
Discussion
105
5.3 Corrosion protection of high temperature oxides in aqueous solutions
5.3.1 Corrosion properties in aqueous solutions of different pH
In addition to high temperature oxidation properties, corrosion behaviour of Co-base
superalloys in aggressive aqueous solutions at ambient temperature is of particular
interest. Comparing results of oxidised and unoxidised alloys reveals that high
temperature oxide scales are extraordinary protective in aqueous solutions. At first,
results of unoxidised alloys (and pure Co) in solutions of different pH values are
discussed before comparing them to oxidised alloys. In the following, the influence of
chloride ions on the corrosion behaviour of several Co-base superalloys will be discussed,
as well.
At the beginning of immersion in alkaline solution, the initial and irreversible
anodic shift of the measured potentials is assumed to be a result of native passive layer
formation or modification (CoO or hydrated CoO · H2O [139]) on the alloy surface.
However, this initial behaviour can only be observed in alkaline solution, indicating a
stabilisation of the native passive film with increasing pH of the solution [40].
Nevertheless, in the course of OCP measurement in neutral solution, a temporary shift in
anodic direction with a maximum value after 1.5 h can be observed for pure Co, as well.
However, this shift is most probably due to instable Co-containing species on the
specimen surface. Passivation of pure Co and the Co-base alloy shows the typically
observed regions of primary and secondary passivation of pure Co in an alkaline solution
due to CoO (or CoO · H2O) and Co3O4 formation [41, 42]. Furthermore, it is proposed
that anodic passivation of pure Co in alkaline solution is based on Co(OH)2 as first
oxidation product followed by subsequent formation of Co3O4 and CoOOH at more
anodic potentials, which is represented by anodic peaks in the polarisation
curves [140, 141]. Those phases are also reported by Chivot et al. [94] in the form of a
revised Pourbaix-diagram, which takes new thermodynamic data into account. Therefore,
the formation of either Co3O4 or CoOOH is assumed to be present as a second layer on
top of the native passive film of pure Co and the Co-base alloy. The presence of such a
duplex film was also proposed by Badawy et al. [40]. However, it strongly depends on the
pH of the solution. In acidic solution, continuous dissolution of Co or its native passive
106
Discussion
layer is reported in accordance with the results of the present study [40, 47]. However, it
seems that according to the OCP curves, an initially stable passive layer establishes which
dissolves in the course of the following potentiodynamic experiment. Passive layers on
pure Co immersed in acidic solution exhibit lower capacities compared to the alloy
(according to EIS data), indicating a thicker film on pure Co. Nevertheless, anodic
dissolution of the Co-base alloy is slightly less pronounced, probably due to the presence
of passivating alloying elements, such as W. According to Pourbaix [142], W forms
stable WO3 at approximately pH < 4 and therefore may have a positive influence on the
corrosion behaviour of the alloy. Furthermore, small current drops at high potential values
in acidic and neutral solutions are assumed to be related to CoOOH (acidic solution) and
Co3O4 or CoOOH formation (neutral solution), respectively. However, solubility of the
phases is too high in order to enable sufficient passivation. In addition, current drops may
also result from re-precipitation of corrosion products on the sample surface. Another
noteworthy finding of the alloy investigated in alkaline solution is that, after reaching the
secondary passivation range, the measured current abruptly increases at lower potentials
compared to pure Co. It is assumed that the low onset potential for transpassivity might
be due to W-dissolution. Studies of Kellner et al. on WC-Co hard metals report similar
results [143]. Here, limited passivation in alkaline solution can be explained by the onset
of WC-dissolution in the secondary passivation range of Co. According to
Pourbaix [142], Al and W are not passive in alkaline solution and therefore improved
corrosion resistance due to the addition of such alloying elements to pure Co seems
surprising. For instance, binary Co–W alloys show passive behaviour in alkaline (and
acidic) solutions even if W-contents are higher than 20 % [38, 39]. Passive films in
alkaline solutions are composed of Co-hydroxides and -oxides, similar to the layers
observed for pure Co. The exact mechanism of passivity enhancement in alkaline solution
due to W addition remains to be investigated. However, it may be similar to the beneficial
effect of W on passivity of duplex stainless steels due to the adsorption of W into the
passive layer without modifying its oxide state [144]. In this respect, it is reported that W
additions reduce the passive current density of stainless steels and increase the passive
range in anodic direction [145]. In Figure 5.8, a schematic summary of the effects of
Discussion
107
different alloying elements on the anodic polarisation curve is illustrated (after
Sedriks [145]).
Figure 5.8: Influence of alloying elements on anodic polarisation (according to
Sedriks [145]).
In addition to W, Cr and Si are effective elements for improving passivation of stainless
steels and most likely also of Co-base alloys, due to the similar chemical nature of Co and
Fe. Appropriate amounts of Cr and Si lead to the formation of a passive Cr-rich
oxyhydroxide film [146] and a SiO2 passive layer, respectively. Furthermore, Ni addition
to stainless steel is reported to improve corrosion resistance due to its effects on the
crystallographic structure of the steel. In this respect, the main role of Ni is to control
element partitioning and phase balance [147]. However, a considerably improved
corrosion resistance of the currently investigated γ/γ' 0.12B alloy is not expected, as a
consequence of Ni addition.
High temperature oxidation of Co-base superalloys leads to the formation of thick
and highly stable oxide scales, composed of Co-oxides, Al2O3, and ternary oxides
108
Discussion
(spinels), which provide excellent corrosion protection in aqueous solution and act as
effective dissolution barriers in a large potential range. The best protective properties of
high temperature (HT) oxides can be achieved in acidic solution. This may be due to the
presence of tungsten oxides within the HT-scale, which are particularly stable in acidic
solution [142]. In contrast, thin passive layers on unoxidised Co-base superalloys exhibit
several orders of magnitude lower protective properties in all tested aqueous solutions. As
a result of pre-oxidation, OCP ennoblement, increased total impedance values, lower
capacities, and decreased current densities can be measured by means of electrochemical
techniques. Generally, longer oxidation times and higher oxidation temperatures result in
improved corrosion properties of superalloys in neutral and alkaline solution. In acidic
solution, oxidation at 800 and 900 °C leads to comparable corrosion properties, whereas
merely longer oxidation times seem to be favourable for higher protection. Furthermore,
in alkaline solution the 500 h/800 °C alloy exhibits almost identical corrosion behaviour
as the 24 h/900 °C alloy in the anodic region, even though the HT-oxide scale formed in
the latter case is about twice as thick as the scale formed during long-term oxidation at
800 °C. This demonstrates that not only the thickness of HT-scales but also their detailed
nature (Al2O3-content, morphology, defect structure, porosity, and oxide stabilities in
solutions of different pH) dictate the corrosion performance. Therefore, it is very
complicated to predict the extent of the protective effect of complex oxide scales
depending on oxidation conditions and utilised aqueous solution. However, the highest
increase in corrosion resistance can always be measured between the unoxidised 0.12B
alloy and the alloy oxidised for 24 h, which is due to the initial formation of a HT-oxide
scale (depending on the temperature either Co3O4 or CoO). Potential fluctuations during
the OCP measurement of the 24 h/800 °C alloy can be detected in alkaline solution,
probably due to a relatively unstable scale compared to thicker HT-oxide scales obtained
by longer oxidation times and higher temperatures. Further oxidation decreases the
corrosion rate due to an increased total oxide thickness and the additional formation of
protective inner Al2O3, which is more or less continuous depending on oxidation time and
temperature. Moreover, the effect of further alloying elements, such as W and B, which
are of importance for the HT-oxidation behaviour, should be considered for the corrosion
properties in aqueous solution. As already mentioned, W is passive in acidic solution and
Discussion
109
therefore may have a beneficial effect on the corrosion behaviour of Co–Al–W–B alloys
in 0.5 M H2SO4 aqueous solution [142]. B is reported to be a strong corrosion resistant
material [95], and hence small amounts of B-containing species within the HT-oxide
scale or the passive layer are expected to contribute to the corrosion properties of the
superalloy, as well.
In comparison to the Co–Al–W–B superalloys of the present study, ObigodiNdjeng [29] previously investigated the electrochemical behaviour of HT-oxides on
polycrystalline Ni-base superalloys (Ni–13.6Cr, Ni–13.6Cr–8.8Co, Ni–13Cr–7.4Al,
in at%) and a commercial single crystalline Ni-base superalloy “PWA 1483”
(Ni–13.7Cr–8.9Co–7.8Al–5Ti–1.6Ta–1.2W–1.2Mo–0.3C, in at%) in neutral borate buffer
solutions (pH 8.4) using the same electrochemical set-up as presented in Figure 3.2. The
model alloys possess approximately the same Cr, Co, and Al content as the commercial
PWA 1483 alloy. Oxide scales were obtained by oxidising the samples for 100 h at
800 °C in air. The measured impedance values were extraordinary high (> 100 MΩ cm2),
whereas maximum impedance values were not reached within the frequency range of the
EIS measurements (5 mHz – 100 kHz). In the present study, total impedance values of
novel Co–Al–W–B superalloys oxidised for 100 h at 800 °C are high (about 2 MΩ cm2 in
neutral solution) but two orders of magnitude lower than the corresponding values
reported by Obigodi-Ndjeng. Furthermore, polarisation curves of oxidised Ni-base alloys
(from -1 V to +2 V vs. Ag/AgCl) exhibit comparable corrosion potentials but corrosion
current densities of approximately 1·10-8 A/cm2, i.e. one order of magnitude lower than
the values reported for Co–Al–W–B superalloys. It is noteworthy that oxidised Ni-base
alloys, in contrast to Co-base alloys of the present study, exhibit a steady increase in
current density upon anodic polarisation up to 10 mA/cm2 at +2 V (vs. Ag/AgCl), which
indicates a higher electrochemical reactivity of the HT-oxides, for instance oxidation
reactions of some scale constituents. In this respect, oxide scales on novel Co–Al–W–B
superalloys seem to be better dissolution barriers over the entire polarisation range from
-1 V to +2 V (vs. Ag/AgCl), presumably due to their considerably higher thickness
(approximately one order of magnitude) and the presence of other oxide species (Co3O4,
CoO, CoWO4, CoAl2O4, etc.) compared to HT-scales on Ni-base alloys. The use of a
borate buffer as neutral electrolyte in the studies of Obigodi-Ndjeng is expected to have a
110
Discussion
minor effect on the statements above. As a conclusion, oxidised Ni-base alloys reveal
superior corrosion resistance in neutral aqueous solution but barrier properties of HToxides on Co-base alloys seem to be improved over the entire potential range. Since Nibase alloys possess high Cr- and Ti-contents, additions of those alloying elements may
further enhance corrosion properties of oxidised Co–Al–W–B superalloys. However,
appropriate amounts of Ti and especially Cr have to be chosen in order to maintain
mechanical properties. Further additions of γ'-stabilising elements may be necessary as
previously discussed in chapter 5.2.2.
5.3.2 Corrosion properties in chloride containing aqueous solution
Pitting corrosion of metals in aqueous solutions is an important issue which has to be
addressed in addition to pH depending corrosion. In comparison to the neutral
1 M Na2SO4 electrolyte, the presence of chloride ions in neutral 0.5 M NaCl solution
leads to decreased corrosion resistance of the investigated unoxidised specimens,
characterised by lower impedance values and a limited passivation range, which is
followed by rapid passive layer destruction due to pitting corrosion. Due to the limited
passivation by Co-oxides and -hydroxides, the entire metal surface is homogeneously
attacked by chloride ions and the resulting pits can easily grow in lateral direction. This
leads to the wide and shallow pit morphology observed by optical microscopy. In
contrast, additions of Cr as further alloying element would lead to sufficient passivation
of the surface and hence rather deep and narrow pits would be expected in the course of
chloride attack. Furthermore, the high amount of pits on the investigated sample surfaces
is a result of long polarisation up to potentials of +2 V (vs. Ag/AgCl). Typical pitting
corrosion products, such as the hydrated chloro-complex CoOHCl · 4H2O [118], are
formed on the sample surface but cannot be analysed by means of surface analytical tools
since they do not attach to the surface or re-precipitate at higher potentials. As a result of
severe pitting corrosion, the investigated Co-base superalloys have to be protected in
chloride containing environments, for example by pre-oxidation of the samples in order to
form protective oxide layers. Electrochemical data shows that thick HT-oxide scales are
effective dissolution barriers that provide excellent protection against chloride attack. Flat
cathodic and anodic branches of the polarisation curves indicate suppression of
Discussion
111
electrochemical processes at the surface, such as reduction or oxidation reactions of the
scale components. It is noteworthy that due to a thick HT-scale, oxidised pure Co exhibits
corrosions properties almost comparable to oxidised 0.12B and 0.12B0.005Y alloys.
However, similar to previous investigations in solutions of different pH, corrosion
properties not only depend on the HT-scale thickness but also on the incorporated oxide
species. In this respect, the addition of Si to 0.12B alloys has the potential for
considerable improvements in corrosion resistance due to the formation of protective
SiO2 [26, 27] and the beneficial effect of Si on the formation of Al2O3 on the alloy
surface [28]. As a result, superior impedance values and decreased current densities can
be measured over the entire polarisation range of the pre-oxidised 0.12B2Si alloy (65 h at
900 °C). Due to the formation of protective Ti-rich phases, 0.12B2Ti alloys are assumed
to possess excellent corrosion properties in chloride containing solutions, as well. Even
short-term oxidation of the 0.12B0.005Y alloy for 1 h at 900 °C leads to the formation of
a sufficiently thick and compact CoO scale (in addition to further ternary oxides). Higher
Y-contents in the 0.12B0.005Y alloy are assumed to enhance selective oxidation of Al, in
addition to enhanced layer adhesion [15-21], and therefore improved Al2O3 formation on
the alloy surface is expected to result in even higher pitting corrosion resistance.
Representative top-view optical microscopy images of the homogeneous outer CoO scale
on pure Co, 0.12B, 0.12B0.005Y, and 0.12B2Si samples (see Figure 5.9) visualise high
protection against pitting corrosion since no changes of the morphology or corrosion
initiation sites can be observed after performing electrochemical experiments. However,
since passive layer breakdown is frequently time dependent, this matter may need further
assessment, such as constant anodic potential tests for long periods of time. If the thick
HT-scale breaks down locally, it might self heal or act as a crevice leading to higher
localised corrosion susceptibility. Therefore, the role of the HT-oxide layer after a
breakdown event needs further assessment in future work.
112
Discussion
Figure 5.9: Top-view optical microscopy images of the oxidised sample surface
(a) before and (b) after electrochemical characterisation.
5.4 Comparison of high temperature oxidation and aqueous corrosion
In the present work, pure Co and several Co-alloys are exposed to a variety of aqueous
solutions at ambient temperature and oxidised at high temperatures in air. It turns out that
the influence of alloying elements on the material properties is significantly stronger in
the latter case. Compared to pure Co, Co-base superalloys (0.12B and 0.12B0.005Y)
exhibit comparable corrosion properties in neutral aqueous solution and slightly improved
passivation in acidic and alkaline solution. In contrast, high temperature oxidation of the
alloys results in remarkably improved oxidation resistance compared to pure Co. This
leads to the assumption that alloying elements, such as B and Al, have a strong impact on
high temperature oxidation by forming protective Al2O3 in the inner oxide layer and Bcontaining grain boundary phases, which both act as effective diffusion barriers during
the oxidation process. In comparison, pure Co merely forms an outer non-protective Cooxide layer. Based on the results of the present work, an outer Co-oxide layer, which also
establishes during aqueous corrosion, seems to be the dominant factor for determining the
corrosion properties of the investigated samples at ambient temperature. For this reason,
alloying elements are supposed to be less effective for improving passivation in aqueous
solutions compared to their effect on high temperature oxidation resistance.
Discussion
113
High temperature oxidation of 0.12B alloys at 800 °C leads to the formation of
considerably thinner oxide layers than oxidation at 900 °C. Subsequent electrochemical
investigations of those pre-oxidised 0.12B alloys reveal that thicker oxide layers often
provide better corrosion resistance in aqueous solution. However, differences in corrosion
resistance are less distinctive than differences in oxidation resistance. Oxidation at 800 °C
already leads to sufficiently thick and compact oxide scales which provide good barrier
properties against electrochemical dissolution. For this reason, further improvements due
to higher oxidation temperatures, and therefore thicker oxide layers, are limited. Results
of oxidised 0.12B2Si alloys reveal that the presence of SiO2 within the surface layer leads
to noticeable improvements of the corrosion resistance in aqueous solution. Furthermore,
the difference in corrosion resistance of oxidised Si-free and Si-containing alloys in
aqueous solution is comparable to their difference in oxidation resistance at high
temperature. As a result, some alloying elements may considerably improve oxidation
resistance at high temperatures in air and corrosion resistance of oxidised alloys in
aqueous solution, but effects of alloying elements on corrosion properties of unoxidised
alloys in aqueous solution are rather low.
114
Conclusion and outlook
6. Conclusion and outlook
High temperature oxidation in air of novel γ'-strengthened Co–Al–W based superalloys
indicates the presence of a transient oxidation period at short exposure times followed by
a steady-state oxidation range. After fast initial oxide formation, a multi-layered oxide
structure develops, composed of an outer Co-oxide layer (Co3O4 and CoO), a middle
mixed oxide layer consisting of presumably amorphous or nanocrystalline spinels (e.g.
CoWO4 and CoAl2O4), and an innermost Al2O3-rich area. Due to the formation of Al2O3,
outward ion diffusion is impaired resulting in improved oxidation resistance. In case of a
continuous inner Al2O3 layer, no further oxidation can be detected. As a result of outward
Al-diffusion during high temperature oxidation, a phase transformation at the alloy/oxide
interface from γ/γ' into γ/Co3W can be observed. However, it is assumed that the Aldepleted γ'-free region beneath the Al2O3 layer may have negative effects on mechanical
properties at high temperatures, such as creep strength.
Investigation of pure Co and ternary Co–Al–W alloys reveals a parabolic timedependency of oxide layer growth. However, the addition of appropriate alloying
elements has a strong influence on oxidation behaviour. For instance, quaternary
Co–Al–W–B alloys show significantly improved oxidation resistance resulting in subparabolic oxidation at 900 °C and, after a certain exposure time, complete blockage of
further oxidation at 800 °C. Appropriate amounts of B (0.12 at%) benefit the formation of
Al2O3 at the alloy/oxide interface and prevent decohesion of the scale due to Baccumulation (most likely tungsten borides) in the middle oxide layer. Further addition of
Cr to the Co–Al–W–B system leads to the development of an inner double layer of
protective Cr2O3 and Al2O3. Since Cr benefits selective oxidation of Al, a higher amount
of inner Al2O3 can be detected compared to Cr-free alloys. In contrast to conventional
corrosion-resistant Co-base superalloys (such as MAR-M-509), relatively low amounts of
Cr (8 at%) are sufficient in order to achieve good oxidation resistance. However, Cr leads
to the formation of volatile Cr-rich phases at temperatures higher than 1000 °C and to an
altered γ/γ'-microstructure. As a result, further alloying elements are required, such as Ni,
Mo, or Ta in order to stabilise the γ'-phase and maintain the desired mechanical
properties. In contrast to Cr-addition, Ti and Si do not alter the γ/γ'-microstructure but
Conclusion and outlook
115
lead to enhanced oxidation resistance, as well. This behaviour can be ascribed to Ti- and
Si-rich phases and the beneficial effect of Ti and Si on Al2O3 formation at oxidation
temperatures of at least 900 °C. In this regard, both elements seem to operate similar to
Cr. However, it is reported that the Ti-containing Co-base alloy possesses outstanding
creep properties at 850 °C, compared to other novel Co-based alloy systems, and hence
may be the most promising alloy for further optimisation (concerning mechanical and
corrosion properties). In contrast, Si-rich phases at the alloy/oxide interface, within
precipitates, and at the grain boundaries are expected to improve the oxidation resistance
but impair the mechanical properties. Since Ni plays a fundamental role in the field of
superalloys, its influence on high temperature oxidation of Co–Al–W–B alloys was
investigated, as well. Results reveal that enhanced solubility of B in the alloy matrix can
most likely be ascribed to the addition of Ni. This leads to an absence of B at the grain
boundaries and within the inner oxide layer and therefore to impaired oxidation
resistance. Further additions of B or other strengthening elements, such as Zr or C, may
be a way to compensate the negative effects of Ni. According to experimental results of
the present work, small amounts of the rare earth element Y (0.005 at%) in Co–Al–W–B
alloys do not improve the oxidation resistance as expected, most likely due to insufficient
amounts of the alloying element. Therefore, the Y-content has to be adjusted for the novel
alloy system in order to achieve the desired properties.
In addition to high temperature oxidation, pure Co and Co–Al–W–B alloys are
further characterised by electrochemical measurements in aqueous solution of different
pH. Results of the unoxidised samples reveal that corrosion resistance improves with
increasing pH value due to the formation of a double layer consisting of Co3O4 or
CoOOH species on top of a Co(OH)2 film. In alkaline 0.1 M NaOH solution, both
materials show primary and secondary passivation while performing polarisation
experiments from -1 to +2 V (vs. Ag/AgCl). In contrast, limited passivation is observed in
neutral 1 M Na2SO4 solution and active dissolution in acidic 0.5 M H2SO4 solution.
Experiments in neutral chloride containing solution (0.5 M NaCl) lead to severe pitting
corrosion after limited initial passivation of the metal surface. Chloride ions
homogeneously attack the entire sample surface due to limited protection of Co-oxides
116
Conclusion and outlook
and Co-hydroxides. As a result, preferential pit growth is in lateral direction and hence
wide and shallow pits can be observed on the sample surface.
High temperature oxidation of the investigated samples leads to the formation of
highly effective barriers against aqueous corrosion without any breakdown events in the
course of the performed measurements. The significant retardation of cathodic and anodic
reactions can be ascribed to the formation of thick oxide layers on the metal surface,
which possess best barrier properties in acidic solution. Tungsten oxides within the scale
are particularly stable in acidic solution and hence may explain this finding. Experiments
in 0.5 M NaCl aqueous solution reveal that Si-addition to quaternary Co–Al–W–B alloys
leads to remarkably improved corrosion properties of the oxidised samples due to the
formation of Al2O3 and Si-rich phases. Therefore, not only oxide layer thickness but also
the chemical composition and the stability of the scale dictate corrosion properties in
different aqueous solutions. However, further long term assessment at constant anodic
potentials is inevitable since oxide layer breakdown is frequently time dependent.
Moreover, detailed electrochemical investigations addressing the effect of further alloying
elements, such as Ti, Cr or Si in Co–Al–W–B alloys, may give important insights for the
development of novel corrosion resistant Co-base superalloys.
With regard to high temperature oxidation properties, it would be of interest to
investigate new Co-based alloys under cyclic oxidation conditions and to test different
corrosive atmospheres, for instance sulphur-containing hot gases. Furthermore, the
influence of different microstructures, such as grain size, single crystallinity or γ'morphology on oxidation properties, as well as initial oxidation stages after very short
exposure times should be examined in future work.
In the field of polycrystalline alloys, it would be important to discover the exact
mechanisms associated with the positive effect of B on high temperature strength and
oxidation resistance. Furthermore, even though single-crystalline superalloys exhibit
superior high temperature properties due to a lack of grain boundaries, it is recommended
to add small amounts of B for further improvements in oxidation resistance.
Qualitative phase analysis of grain boundaries, precipitates, and surface layers
may be achieved by using transmission electron microscopy (TEM) after sample
preparation by focussed ion beam (FIB) techniques. However, oxide detection within the
Conclusion and outlook
117
middle surface layer is expected to be difficult due to overlapping of diffraction pattern of
several co-existing oxide species.
Thermodynamic and kinetic simulation of the oxidation behaviour and phase
stability within novel Co-base superalloys is an important tool which has to be applied in
future, similar to the modelling of Ni-base superalloys. A project in the new
SFB/Transregio 103 addresses this issue by developing an appropriate database in
addition to commonly used ThermoCalc and DICTRA, in order to assure reasonable
simulations.
Based on the potential of Co–Al–W based superalloys and on current research
progress (including the findings of the present work), the development of sophisticated
single-crystalline Co-base alloys in the foreseeable future may result in novel high
temperature materials, which possess superior mechanical and corrosive properties
compared to common Ni-base superalloys.
118
Conclusion and outlook
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Acknowledgement
127
Acknowledgement
First of all, I would like to thank my advisor Prof. Dr. S. Virtanen for giving me the
opportunity to work on this interesting topic and for her guidance and support during the
last three years. I am also very grateful for the valuable discussions and for her
confidence in me and my work.
I would like to thank Prof. Dr. U. Glatzel for his commitment within the graduate school
(Graduiertenkolleg 1229/2) and for being my co-advisor and second examiner of this PhD
thesis.
In addition, I would like to acknowledge the German Research Foundation (Deutsche
Forschungsgemeinschaft, DFG) for the financial support considering the graduate school.
A special thank goes to my colleague Alexander Bauer for proofreading the present PhD
study, but also to all the members of the graduate school for their help, advice, the
valuable discussions, and the pleasant atmosphere at our meetings and summer schools.
Furthermore, I would like to thank my colleagues and the technicians at the Institute for
Surface Science and Corrosion (LKO), especially the members of the corrosion group and
my student assistants/bachelor students Christian Roy, Sebastian Martin, Florian Fischer,
Yilei Shen, Dagmar Rückle, and Manuel Wagner, who all contributed to the success of
this work.
I am also very thankful to Dr. N. Birbilis for accepting to be my advisor at Monash
University in Clayton, VIC, Australia, and for his support before and during my stay
abroad. Furthermore, I would like to thank the staff of the Department of Materials
Engineering and the members of the Monash Centre for Electron Microscopy (MCEM)
for assisting me with sample preparation and electron microscopy.
128
Acknowledgement
Furthermore, I would also like to thank the Institute of General Material Properties
(WW1) for support, valuable discussions, and especially for giving us the opportunity to
use some of their equipment (SEM, EDX, carbon evaporator, high temperature oven, etc.)
for our research.
I am extremely thankful to my whole family, especially my beloved wife Tina for her
encouragement, help, and for always being there for me.
Last but not least, I would also like to thank my friends and everybody who accompanied
me during the last three years of my PhD study.