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Kinetics Study and Characterization of Thermally Grown Oxide on Commercial
β-(Ni,Pt)Al Bond Coats used in Thermal Barrier Coating Systems for Gas Turbine
Engine Applications
Thesis · November 2012
DOI: 10.13140/RG.2.2.11832.98563
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CENTRO DE INVESTIGACIÓN Y DE ESTUDIOS AVANZADOS
DEL INSTITUTO POLITÉCNICO NACIONAL
UNIDAD QUERETARO
Estudio Cinético y Caracterización del Oxido Crecido
Térmicamente Sobre Recubrimientos de Unión Comerciales
β-(Ni,Pt)Al Usados en Sistemas de Recubrimientos Barreras
Térmicas de Alabes de Turbinas de Gas
Tesis que presenta
Juan Manuel Alvarado Orozco
para obtener el Grado de Doctor en Ciencias
en la Especialidad de Materiales
Juan Muñoz Saldaña
Santiago de Querétaro, Qro.
Noviembre, 2012.
Kinetics Study and Characterization of Thermally Grown Oxide
on Commercial β-(Ni,Pt)Al Bond Coats used in Thermal Barrier
Coating Systems for Gas Turbine Engine Applications
by
Juan Manuel Alvarado Orozco
A dissertation submitted to the
Cinvestav, Queretaro Campus
for the degree of
Doctor in Science of Materials
Kinetics Study and Characterization of Thermally Grown Oxide
on Commercial β-(Ni,Pt)Al Bond Coats used in Thermal Barrier
Coating Systems for Gas Turbines Engine Applications
Juan Manuel Alvarado Orozco
Department of Materials Science
Cinvestav, Queretaro Campus
Abstract
The scope of the present work, dictated by the interests of GE-Aviation, was to optimize the preoxidation parameters needed to grow a homogeneous, exclusively α-Al2O3 scale (thermally grown
oxide, TGO) on the surface of a commercial β-(Ni,Pt)Al bond coat system prior to 7YSZ top coat
deposition. Based on this, the main aim of this work was to study the growth kinetics and oxidation
mechanisms, as well as to characterize the structure and microstructure of the TGO as a function
of controlled pre-oxidation treatments. The first part (Chapter 6) of this study investigated the
oxidation kinetics behavior of an as-coated commercial β-(Ni,Pt)Al bond coat system during preoxidation treatments in the range of 900 to 1200°C using an argon atmosphere flow with oxygen
partial pressure of 1 x10-5 atm. The formation of the θ-Al2O3 and α-Al2O3 phases, as well as the
transformation from θ-Al2O3 → α-Al2O3 were observed and monitored. The local parabolic fitting
methodology proposed by Monceau and Pieraggi, using a general parabolic model (t =
A+BΔm+CΔm2), was used to estimate the evolution of the parabolic rate constant, kp, during the
first 5 h of oxidation. Based on this analysis, it was concluded that all tested samples showed
deviations from the classic parabolic model established by Tamman to describe diffusioncontrolled, solid-state oxidation. A steady-state regime was established only after 4 h of treatment,
except for the sample oxidized at 1100°C, for which the steady-state was not established even after
5 h. This behavior was attributed to the Pt, which favors the growth of θ-Al2O3 after nucleation of
α-Al2O3 and consequently delays its transformation, as has been previously reported by Cadoret et
al. \
In the second part (Chapter 7), the effect of oxygen partial pressure and a grit-blasting surface
treatment on the oxidation kinetics of the commercial β-(Ni,Pt)Al bond coat system was conducted
by means of TGA analysis during isothermal treatments at temperatures from 1000 to 1150°C. This
data was also complemented by a structural analysis of the resulting oxide scale using photo
luminescence spectroscopy. The combination of techniques allowed a thorough evaluation of the
evolution of the θ-Al2O3 → α-Al2O3 transformation as a function of oxidation conditions and
surface preparation. It was found that the combined effect of low oxygen partial pressure and a
surface grit-blasting treatment stifles the growth of metastable θ-Al2O3 and accelerates the θ-Al2O3
→ α-Al2O3 phase transformation, therefore reducing the time required to reach a steady-state
regime where α-Al2O3 is the controlling phase. For the grit-blasted sample oxidized at 1000°C in
an argon atmosphere with a pO = 1 x 10-5 atm, a steady-state regime was established after ~1 h, in
2
contrast to the 4 h required for a non-grit-blasted (as-aluminized) sample. The suppression of the
growth of θ-Al2O3 was associated with the fact that at low oxygen partial pressures, gas-supply
control is established during the early stages of oxidation. Grit-blasting was associated with an
increase of α-Al2O3 nucleation sites.
i
Estudio Cinético y Caracterización del Oxido Crecido Térmicamente
Sobre Recubrimientos de Unión Comerciales β-(Ni,Pt)Al Usados en Sistemas
de Recubrimientos Barreras Térmicas de Alabes de Turbinas de Gas
Juan Manuel Alvarado Orozco
Departamento de Materiales
Cinvestav-Unidad Querétaro
Resumen
El alcance del presente trabajo fue motivado por el interés de GE-Aviation, y fue el de optimizar los
parámetros de pre-oxidación necesarios para el crecimiento exclusivo de una capa homogénea de
Al2O3-α (oxido crecido térmicamente, TGO) sobre la superficie de sistemas de recubrimientos de unión
comerciales β-(Ni,Pt)Al previo al depósito del recubrimiento aislante superior 7YSZ.
Basado en esto, el objetivo principal de este trabajo fue estudiar el crecimiento cinético y mecanismos
de oxidación, así como caracterizar las propiedades estructurales y microestructurales del TGO como
una función de los tratamientos de pre-oxidación controlados.
La primera parte (Capitulo 6) de este estudio investigo las comportamiento de las cinéticas de oxidación
de recubrimientos de unión comerciales β-(Ni,Pt)Al durante tratamiento de pre-oxidación en el rango
de 900 a 1200°C usando un flujo de argón con una presión parcial de oxigeno. La formación de las
fases Al2O3-θ y Al2O3-α, así como la transformación de θ Al2O3 → α-Al2O3 fueron observadas y
monitoreadas. La metodología propuesta por Monceau y Pieraggi basada en un ajuste parabólico local
usando un modelo parabólico general (t = A+BΔm+CΔm2) fue usada para estimar la evolución de la
constante parabólica de crecimiento, kp, durante las primeras 5 h de tratamiento. Basado en este análisis,
se pudo concluir que todas muestras oxidadas mostraron desviaciones del modelo parabólico clásico
establecido por Tamman para describir el crecimiento de una óxido controlado por la difusión en estado
sólido. Un régimen de estado estacionario fue establecido hasta después de 4 h de tratamiento, excepto
para la muestra oxidada a 1100°C. Este comportamiento fue atribuido al Pt quien favorece el
crecimiento de Al2O3-θ después de la nucleación de Al2O3-α y consecuentemente retrasa su
transformación as como fue anteriormente propuesto por Cadoret et al. En la segunda parte (Capitulo
7), el efecto de la presión parcial de oxigeno y el tratamiento superficial de la muestras por medio de
un granallado sobre la cinética de oxidación de sistemas de recubrimientos de unión comerciales β(Ni,Pt)Al fue llevado a cabo por medio de un análisis termogravimétrico durante tratamientos
isotérmicos a temperaturas de 1000 a 1150°C. Estos datos fueron al mismo tiempo complementados
por medio de un análisis estructural de las capas de óxido resultantes usando espectroscopia de
luminiscencia fotoestimulada. La combinación de técnicas permitió una evaluación completa de la
evolución de la transformación θ-Al2O3 → α-Al2O3 como función de los parámetros de pre-oxidación
y su preparación superficial del recubrimiento unión. Como resultado de este estudio, fue encontrado
que el efecto combinado de una baja presión parcial de oxigeno y un tratamiento de granallado
superficial, sofoca el crecimiento de la fase metaestable Al2O3-θ y acelera la transformación de fase θAl2O3 → α-Al2O3, por lo tanto reduce el tiempo para alcanzar un estado de crecimiento estacionario
donde Al2O3-α es la fase que controla el crecimiento del óxido. Para una muestra granallada
superficialmente y oxidada a 1000°C en una atmosfera de argón con una pO = 1 x 10-5 atm, un régimen
2
de estado estacionario fue establecido después de aproximadamente 1 h, en contraste con las 4 h
requeridas para una muestra sin tratamiento de granallado. La supresión del crecimiento de la fase de
Al2O3-θ fue asociado con el hecho de que a condiciones de baja presión parcial de oxigeno, el
crecimiento del óxido está controlado por el suministro de aire durante las primeras etapas de
crecimiento, además un incremento en la nucleación de Al2O3-α debido al proceso de granallado.
ii
PREFACE
This thesis covers the research conducted as a part of the multifunctional materials group at Centro
de Investigación y de Estudios Avanzados del Instituto Politecnico Nacional, campus Queretaro,
during the period of 2009-2012, under the supervision of Dr. Ing. Juan Muñoz Saldaña.
The following two papers are included in this thesis as chapters 6 and 7, respectively.
I.
First Stages of Oxidation of Pt-Modified Nickel Aluminide Bond Coat Systems at Low
Oxygen Partial Pressure.
J. M. Alvarado-Orozco, R. Morales-Estrella, M. S. Boldrick, J. L. Ortiz-Merino, D. G.
Konitzer, G. Trapaga-Martinez and J. Muñoz-Saldaña.
Oxidation of Metals, 2012, DOI: 10.1007/s11085-012-9305-7.
II.
A Kinetic Study of the Competitive Growth between θ-Al2O3 and α-Al2O3 during the Early
Stages of Oxidation of Pt-modified Nickel Aluminide Bond Coat Systems: Effect of Oxygen
Partial Pressure, Surface Treatment, Time and Temperature.
J. M. Alvarado-Orozco, R. Morales-Estrella, M. S. Boldrick, J. L. Ortiz-Merino, D. G.
Konitzer, G. Trápaga-Martínez, B. Gleeson, and J. Muñoz-Saldaña.
To be submitted to Oxidations of Metals.
GE Aircraft Engines, provided the samples for this study. My contributions to the appended papers
were as follow:
I.
I did all the experimental work and wrote the paper. However, it is necessary to
acknowledge the stimulating discussions of the results of this paper with professors Brian
Gleeson (at the University of Pittsburgh) and Carlos Levi (at the University of California
at Santa Barbara).
II.
I did all the experimental work and wrote the paper.
The following papers were also published during this period, but are omitted in this thesis because
they are not directly related with the topic of the dissertation.
iii
•
Inter Laboratory Comparison and Analysis on Mechanical Properties by Nanoindentation.
J. M. Alvarado-Orozco, C. Cárdenas-Jaramillo, D. Torres-Torres, R. Herrera-Basurto, A.
Hurtado-Macias, J. Muñoz-Saldaña1 and G. Trápaga-Martinez.
Mater. Res. Soc. Symp. Proc. Vol. 1243, 2010.
•
Structural evolution of B2-NiAl synthesized by high-energy ball milling.
H. Ruiz-Luna, J. M. Alvarado-Orozco, L. A. Cáceres-Díaz, López-Báez, J. MorenoPalmerín, F. J. Espinoza-Beltrán, M. S. Boldrick, G. Trápaga-Martínez and J. MuñozSaldaña.
Juan Manuel Alvarado Orozco
Queretaro, Mexico 2012.
iv
ACKNOWLEDGEMENTS
At the end of a long and enriching trip as is the development of a thesis, there are many people to
thank for their scientific contributions, friendship, and support during this time.
Firstly, I am very grateful to my supervisor, Dr. Ing. Juan Muñoz Saldaña for providing me this
opportunity, his guidance, and his confidence during these years; but more importantly, I would
like to thank him for his friendship.
Dr. Luis Gerardo Trápaga Martinez for his invaluable support and advice during the development
of this work.
Dr. Mike S. Boldrick of Peace Corps for his friendship, as well as his incredible support during the
revision of this work and advice over the last two years.
Dr. Ricardo Morales Estrella at the Metallurgical Research Institute of the Universidad Michoacana
de San Nicolás de Hidalgo, for his friendship, support, and discussions during the TGA study.
I would also like to extend my thanks to Dr. Jesús González Hernández, Dr, Francisco Javier
Espinoza Beltran, Dr. Juan Francisco Pérez Robles and Dr. José de Jesús Pérez Bueno for being
part of my thesis review committee.
Many thanks to my friends and coworkers Alma Gabriela Mora Garcia (Morita), Haidee Ruiz Luna
(Heidi), Luis Alberto Caceres Diaz (Betiko) John Edison Garcia Herrera (Chaparro), and Laura
Patricia Rivera Resendiz (Laus). I really appreciate and admire you guys. Thanks for sharing your
time and knowledge with me. We will always be a team.
I especially want to thank my friends Erika Herrera Jimenez (Ms Xangoo and many more), Adriana
Gallegos Melgar (Adna), Dr. Alicia Rodriguez Pulído (Likis), Dr. Luis Angel Gutierrez Ladron de
Guevara, Dr. Juan Jose Gervacio Arciniága (Don Gervas), and Christian Cardenas Jaramillo
(Panzon) for all laughs, advice, and quality time.
I would like to thank CONACYT and all the institutional programs that made the development of
this work possible.
Finally, but no less important, to my family - even though I am far away from you, you are part of
my heart and my motivation to keep going.
v
CONTENTS
1.
2.
Introduction ............................................................................................................................. 1
1.1
Overview ........................................................................................................................... 1
1.2
Justification ....................................................................................................................... 3
Background .............................................................................................................................. 4
2.1
Thermal Barrier Coating Systems..................................................................................... 4
2.1.1
Composition (Component Layers)......................................................................... 5
2.1.1.1
Superalloys (SA) .................................................................................................... 5
2.1.1.2
Metallic Bond Coat (BC) ....................................................................................... 7
2.1.1.2.1 Nickel Aluminides ................................................................................................. 8
2.1.1.2.1.1 Mechanisms of Diffusion Coating Formation ..................................................... 10
2.1.1.2.2 Pt-Modified Nickel Aluminides........................................................................... 13
2.1.1.2.2.1 Pt-Modified Nickel Aluminide Coating Stability ................................................ 15
2.1.1.2.3 Pt-Modified γ+γ´ Diffusion Coating .................................................................... 18
2.1.1.3
Ceramic Top Coat (TC) ....................................................................................... 19
2.1.2
Failure Mechanisms ............................................................................................. 21
2.2
Oxidation ........................................................................................................................ 22
2.2.1
Thermodynamic Principles of Oxidation ............................................................. 22
2.2.2
Mechanisms of Oxidation .................................................................................... 25
2.2.2.1
Transport Mechanisms ......................................................................................... 25
2.2.2.2
Wagner´s Theory of Oxidation ............................................................................ 28
2.2.2.3
Diffusion Controlled Oxidation [via Grain Boundary Diffusion] ....................... 34
2.2.3
High-Temperature Oxidation of Alloys ............................................................... 38
vi
2.2.4
Alumina Forming Alloys ..................................................................................... 41
2.2.4.1
Ni-Al .................................................................................................................... 41
2.2.4.2
Ni-Pt-Al ............................................................................................................... 42
2.2.5
Stability of the Alumina Phases ........................................................................... 43
2.2.6
Oxidation Kinetic Models .................................................................................... 46
3.
References ............................................................................................................................. 50
4.
Hypothesis ............................................................................................................................. 64
5.
Objectives .............................................................................................................................. 65
6.
First Stages of Oxidation of Pt-Modified Nickel Aluminide Bond Coat Systems at Low
Oxygen Partial Pressure ................................................................................................................ 66
Abstract ...................................................................................................................................... 66
6.1
Introduction ..................................................................................................................... 67
6.2
Experimental Procedure .................................................................................................. 70
6.2.1
Sample Preparation .............................................................................................. 70
6.2.2
Oxidation Treatments .......................................................................................... 71
6.2.3
Characterization ................................................................................................... 71
6.3
Results and Discussion ................................................................................................... 72
6.3.1
Thermogravimetric Analysis ............................................................................... 72
6.3.2
Microstructural and Structural Analysis .............................................................. 74
6.3.3
Kinetics Analysis ................................................................................................. 80
6.4
Conclusions ..................................................................................................................... 84
References ................................................................................................................................. 84
7.
Kinetics Study of the Competitive Growth between θ-Al2O3 and α-Al2O3 during the Early
Stages of Oxidation of Pt-Modified Nickel Aluminide Bond Coat Systems: Effects of Oxygen
Partial Pressure, Surface Treatment, and Temperature ................................................................. 88
vii
Abstract ...................................................................................................................................... 88
7.1
Introduction ..................................................................................................................... 89
7.2
Experimental Procedures ................................................................................................ 91
7.2.1
Sample Preparation .............................................................................................. 91
7.2.2
Oxidation Treatments .......................................................................................... 92
7.2.3
Characterization ................................................................................................... 93
7.3
Results and Discussion ................................................................................................... 95
7.3.1
Effect of Grit Blasting on β-(Ni,Pt)Al BC Morphology ...................................... 95
7.3.2
Effect of Oxidation Temperature and Oxygen Partial Pressure .......................... 95
7.3.3
Effect of Oxidation Temperature and Grit-Blasting Process ............................. 102
7.3.4
Mixed Effect of Low the Oxygen Partial Pressure and the Grit-Blasting Process
106
7.3.5
7.4
Analysis of Oxygen Transfer from Flowing to Sample Surface ....................... 112
Conclusions ................................................................................................................... 118
References ............................................................................................................................... 119
8.
Outlook (Future Work) ........................................................................................................ 121
viii
LIST OF FIGURES
Fig. 2.1. Cross-sectional micrograph of a thermal barrier coating system. ..................................... 5
Fig. 2.2. Progress of the temperature capability of the superalloys over the past 60 years (Y-axis
temperatures are in °C) [22]. ........................................................................................................... 6
Fig. 2.3. Examples of the phases formed in a René N5 Ni-base superalloy: (a) the high-temperature
strength phases γ and γ´, and (b) TCP precipitates. ......................................................................... 6
Fig. 2.4. Ni-Al phase diagram [27].................................................................................................. 9
Fig. 2.5. Room temperature lattice parameter and density of β-NiAl as a function of stoichiometry
[28]. ................................................................................................................................................. 9
Fig. 2.6. Microstructure and phase identities during diffusion coatings manufacture. (a) high
activity aluminide process and (b) low activity aluminide process [33]. ...................................... 11
Fig. 2.7. (a) Interdiffusion as a function of Al content and (b) DNi/DAl ratio (at 1100°C) in β-NiAl
as a function of Al content [35]. .................................................................................................... 12
Fig. 2.8. Schematic diagram of a CVD low-activity co-deposition reactor for aluminum and more
reactive elements [39].................................................................................................................... 13
Fig. 2.9. Possible microstructures of the Pt-aluminide coatings made by CVD process. (a) CVD
external single phase PtAl2 outer layer, (b) CVD external two phase platinum aluminide and (c)
MDC-150 external single phase platinum aluminide [54]. ........................................................... 14
Fig. 2.10. Schematic representation of the Pt-modified nickel aluminide bond coat degradation
[66]. ............................................................................................................................................... 16
Fig. 2.11. Ternary diagram of Ni-Al-Pt system at 1150°C (the β region at 1100°C is also included).
Adapted from [67]. ........................................................................................................................ 17
Fig. 2.12. Cross-section from the as-deposited Pt-modified γ-Ni + γ´-Ni3Al coating on Rene N5
superalloy. Adapted from [75]....................................................................................................... 18
Fig. 2.13. Advancement of the gas-turbine engine temperature capacity as a function of materials
development relative to Rene 80. Adapted from [76]. .................................................................. 19
ix
Fig. 2.14. A binary phase diagram for the ZrO2-YO1.5 system showing the region where the
metastable tetragonal prime (t´) YSZ solid solutions are stable up to 1300°C. The superposed
maroon curve represents the relative cyclic durability as a function of YO1.5 content previously
reported by Stecura [78]. This diagram was adapted from [79]. ................................................... 20
Fig. 2.15. Schematic representation of the failure mechanisms in TBC systems proposed by Evans
et al. [79] ....................................................................................................................................... 21
Fig. 2.16. Ellingham diagram showing the standard free energy change for selected oxides
formation as a function of temperature. Adapted from [88].......................................................... 24
Fig. 17. Schematic view of a Wagner’s diffusion model for a p-type oxide. ................................ 27
Fig. 18. Schematic view of a Wagner’s diffusion model for n-type oxide.................................... 28
Fig. 19. SEM micrographs of the surfaces and cross-sections of polycrystalline α-Al2O3 exposed
at 1650°C for 10 h at two different ∆pO2. (a) and (c) between 10-8 and 1 Pa respectively. (b) and
(d) between 10-8 and 105 Pa respectively. Adapted from [104]..................................................... 36
Fig. 20. Profiles of grain boundary coefficients and flux of Al3+ and O2- in polycrystalline wafer
exposed at 1650°C to pO2upper/pO2lower= 105 Pa/10-8 Pa. Adapted from [104]....................... 37
Fig. 21. Profiles of grain boundary coefficients and flux of Al3+ and O2- in polycrystalline wafer
exposed at 1200°C to pO2upper/pO2lower = 105 Pa/10-23 Pa. Adapted from [104]. ................... 37
Fig. 22. Schematic representation of (a) an internal BO particles formation beneath an AO scale
and (b) the formation of an external protective BO scale. Adapted from [108]............................ 39
Fig. 23. Schematic examples of oxidation behaviors observed in two-phase alloys for which each
phase form an oxide scale. (a) Formation of a mixed nonuniform oxide scale and (b) formation of
uniform and protective oxide scale. Adapted from [108].............................................................. 40
Fig. 24. Oxidation map for Ni-Al system under an O2 atmosphere at 0.1 atm according to Pettit
[118]. Adapted from [88]............................................................................................................... 41
Fig. 25. Oxidation map of the Ni-Pt-Al system exposed to air at 1150°C for 100 h. Adapted from
[124]. ............................................................................................................................................. 43
Fig. 26. (a) Schematic representation of the first layer α-Al2O3 formation and (b) schematic 3D
representation of α-Al2O3 structure. .............................................................................................. 43
x
Fig. 27. Schematic representation of γ-Al2O3 + 3% Pt phase transformation as a function of
temperature. Adapted from [125]. ................................................................................................. 44
Fig. 28. Arrhenius plot of the parabolic rate constants of (a) NiAl and (b) NiAl and NiAl-Cr
oxidation. The oxidation treatments were conducted in a He and O2 flow with an oxygen partial
pressure of 0.13 atm. Adapted from [21]. ..................................................................................... 45
Fig. 29. Illustration of the local-parabolic fitting analysis proposed by Monceau and Pieraggi [132].
....................................................................................................................................................... 47
Fig. 6.1. Microstructure of the as-received samples: (a) surface and (b) cross section. ................ 72
Fig. 6.2. Experimental net mass-gain curves for β-(Ni,Pt)Al bond coats during isothermal oxidation
in a purified argon stream with a pO2 = 1 x 10-5 atm.................................................................... 73
Fig. 6.3. Δm vs. t0.5 plots showing the deviations from the classic parabolic model for the net massgain curves of β-(Ni,Pt)Al bond coats during isothermal oxidation. ............................................ 74
Fig. 6.4. Surface microstructure of samples treated for 5 h at: (a) 900°C, (b) 950°C, (c) 1000°C, (d)
1050°C, (e) 1100°C and (f) 1150°C. ............................................................................................. 75
Fig. 6.5. STEM micrographs of the cross section view showing the α-Al2O3 grain morphology for
the sample treated at 1200°C for 5 h. ............................................................................................ 76
Fig. 6.6. Evolution of the α-Al2O3 phase transformation as a function of temperature after 5 h of
oxidation treatment of (Ni,Pt)Al bond coats. These results were obtained by image analysis of
PSLS mappings. A typical PSLS spectrum for a TGO composed of both α- and θ-Al2O3 phases is
also shown. .................................................................................................................................... 78
Fig. 6.7. Surface morphology of the treated samples, showing the spallation regions observed after
cooling: (a) 950°C, (b) 1000°C, (c) 1050°C, (d) 1100°C, (e) 1150°C (no spallation) and (f) 1200°C
(no spallation). ............................................................................................................................... 79
Fig. 6.8. Log kp plotted as a function of time for β-(Ni,Pt)Al bond coats during isothermal oxidation
in a purified argon stream with a pO2 = 1 x 10-5 atm: (a) transient and steady-state regimes and (b)
steady-state regime. ....................................................................................................................... 81
xi
Fig. 6.9. Arrhenius plot summarizing the kp values of β-(Ni,Pt)Al bond coats during isothermal
oxidation with a pO2 = 1 x 10-5 atm. Bold lines correspond to the γ-Al2O3, θ-Al2O3 and α-Al2O3
lines refer Brumm and Grabke [23]............................................................................................... 82
Fig. 6.10. Arrhenius plot showing a comparison between our kp values for β-(Ni,Pt)Al bond coats
during isothermal oxidation with a pO2 = 1 x 10-5 atm and literature data [56]. Bold lines
correspond to the γ-Al2O3, θ-Al2O3, and α-Al2O3 lines refer Brumm and Grabke [23]. ............... 83
Fig. 7.1. Bond coat system morphology before and after grit blasting: (a and b) surface view, (c
and d)cross-sectional view and (e and f) surface roughness. ........................................................ 94
Fig. 7.2. Δm vs. t0.5 curves showing the effect of the pO2on the oxidation behavior of as-aluminized
β-(Ni,Pt)Al BC system samples during the first 5 h of treatment with (a) pO2 = 1 x 10-5 atm and
(b) pO2 = 2.1 x 10-1 atm. ............................................................................................................... 96
Fig. 7.3. Representative luminescence (PSLS) spectra taken from the TGO surface after 5 h of
isothermal oxidation at different temperatures and oxygen partial pressure: (a) pO2 = 1 x 10-5 atm
and (b) pO2 = 2.1 x 10-1 atm. Note that θ-Al2O3 luminescence is 10-12 times weaker than α-Al2O3.
....................................................................................................................................................... 97
Fig. 7.4. Log kp plotted as a function of time for as-aluminized β-(Ni,Pt)Al BC system samples
during isothermal oxidation at different temperatures and oxygen partial pressures: (a and b)
transient and steady-state regimes and (c and d) enlargement of the steady-state regime. ........... 99
Fig. 7.5. Arrhenius plot showing the effect of low oxygen partial pressure on the kp values of asaluminized β-(Ni,Pt)Al BC system samples. Bold lines refer to Brumm and Grabke data for a NiAl
alloy [20]. .................................................................................................................................... 100
Fig. 7.6. Δm vs. t0.5 curves showing the effect of the grit-blasting process on the oxidation behavior
with pO2= 2.1 x 10-1 atm during the first 2 h of treatment: (a) grit-blasted sample and (b) asaluminized β-(Ni,Pt)Al BC samples. ........................................................................................... 102
Fig. 7.7. Representative luminescence (PSLS) spectra taken from the surface of the thermally
grown oxide on grit-blasted β-(Ni,Pt)Al BC systems samples after 2 h of isothermal oxidation at
different temperatures with pO2= 2.1 x 10-1 atm. It is important to highlight that θ-Al2O3
luminescence is 10-12 times weaker than α-Al2O3. .................................................................... 103
xii
Fig. 7.8. Log kp plotted as a function of time for grit-blasted β-(Ni,Pt)Al BC system samples during
isothermal oxidation at different temperatures and pO2= 2.1 x 10-1 atm: a) transient and steadystate regimes and c) steady-state regime. .................................................................................... 104
Fig. 7.9. Arrhenius plot showing the effect of the grit-blasting process on the kp values of β(Ni,Pt)Al BC system samples. Bold lines refer to Brumm and Grabke data for a NiAl alloy [20].
..................................................................................................................................................... 105
Fig. 7.10. Δm vs. t0.5 curves showing the combined effect of the grit-blasting process and low
oxygen pressure on the oxidation behavior of grit-blasted β-(Ni,Pt)Al BC system samples during
the first 2 h of treatment: (a) pO2= 5 x 10-5 atm and (b) pO2= 1 x 10-5 atm. .............................. 107
Fig. 7.11. Log kp plotted as a function of time for grit-blasted β-(Ni,Pt)Al BC system samples during
isothermal oxidation at different temperatures and oxygen partial pressure: (a) pO2 = 5 x 10-5 atm
and (b) pO2= 1 x 10-5 atm. .......................................................................................................... 108
Fig. 7.12. Arrhenius plot showing the combined effect of the grit-blasting process and low oxygen
pressure on the kp values of β-(Ni,Pt)Al BC systems. Bold lines refer to Brumm and Grabke data
for a NiAl alloy [20]. ................................................................................................................... 109
Fig. 7.13. Arrhenius plot summarizing the oxidation behavior for all oxidizing conditions used in
this work, once that a steady-state regime was established. Bold lines refer to Brumm and Grabke
data for a NiAl alloy [20]. ........................................................................................................... 110
Fig. 7.14. Comparison of dmO/dt vs. t curves for grit-blasted samples oxidized at 1000°C with the
available flux of oxygen during oxidation experiments for the different oxygen partial pressures
used. Red dashed curves refers to Brumm and Grabke data for a NiAl alloy [20]. .................... 117
xiii
LIST OF TABLES
Table 1.1. Polymorphs of Al2O3. Adapted from [103].44Table 6.1. Summary of rate equations for
the different parabolic models proposed by D. Monceau and B. Pieraggi (1998)+. ...................... 69
Table 6.2. Chemical composition of the single-crystal superalloy Rene N5 and the β-(Ni,Pt)Al bond
coat.70Table 7.1. Chemical composition of single-crystal superalloy Rene N5 and β-(Ni,Pt)Al
bond coat ....................................................................................................................................... 92
Table 7.2. Oxidation conditions tested in this study ..................................................................... 93
Table 7.3. Summary of rate equations derived from the general parabolic model proposed by D.
Monceau and B. Pieraggi (1998)+ ................................................................................................. 98
Table 7. 4. Summary of the kinetics and structural analysis of the thermally grown oxide on asaluminized β-(Ni,Pt)Al BC systems ............................................................................................ 101
Table 7. 5. Summary of the kinetics and structural analysis of the thermally grown oxide on gritblasted β-(Ni,Pt)Al BC systems ................................................................................................... 106
Table 7. 6. Gas Molecular Interaction Parameters [42]............................................................... 113
Table 7.7. Mass-transfer parameters for O2 in N2 (pO2 = 2.1 x 10-1 atm). ................................. 114
Table 7.8. Mass-transfer parameters for O2 in Ar. ...................................................................... 115
xiv
1. Introduction
1.1
Overview
After the invention of the first successful gas turbine with excess power by Aegidius Alling in
1903 and the co-invention of the jet propulsion engine by Hans Von Ohain (1937) and Frank
Whittle (1939), it became necessary to develop high-temperature alloys that provided strength
and environmental resistance to the parts in the hottest sections. This new group of materials
included the so-called superalloys which are a group of nickel-, cobalt- and iron-base materials
that are used at temperatures from about 540 and up to 1650°C. A noteworthy features of the
nickel-based superalloys are their excellent mechanical properties under high load bearing at
temperatures in excess of 80% of their incipient melting temperature such as creep resistance,
stress rupture strength, toughness, resistance to the thermal fatigue, as well as its metallurgical
stability [1]. In turn, the drive to improve the efficiency of the gas turbines has demanded an
increase in their operating turbine entry temperatures (TET´s) and therefore an improvement in
their constituent materials [2]. Nowadays, a large number of industrial processes, such as power
generation and propulsion applications, operate in very aggressive environments characterized
by high temperatures, large temperature gradients, oxidizing and corroding atmospheres, high
pressure, and large stresses on individual parts. The current performance of gas turbines is the
result of continuous improvements in different areas of engineering, including turbine design,
control systems, combustion systems, and materials science. Within materials science the most
significant developments have been made in the areas of alloy design, casting technology, and
coating methods [3].
1
High-temperature coatings have been widely used since the 1950s to protect the hot-sections of
aero-engines and stationary gas turbines against oxidation and corrosion phenomena by
promoting a slow-growing oxide scale, as well as against thermal fatigue of the substrate by
providing thermal insulation. These have been classified in diffusion coatings, overlay coatings,
and thermal barrier coatings [4]. Different diffusion coatings processes have been extensively
studied in the last decades, nickel aluminides being one of the most widely used in the
aeronautical industry [2,5,6]. In order to improve their oxidation and hot corrosion resistance,
diverse elements have been added during processing, including Pt, Pd, Rd, Ir, Cr and Hf [7–10].
The multilayer coatings known as thermal barrier coating (TBC) systems have been used since
the early 1980s to protect the parts exposed to the highest temperatures, increasing its lifetime.
TBC systems are usually composed of four main constituents: 1) a Ni-based superalloy (SA) as
substrate, 2) An Al2O3 forming alloys known as the bond coat (BC), 3) a thermally grown oxide
(TGO), the result of BC’s oxidation, and 4) a ceramic top coat (TC) as thermal barrier.
Notwithstanding that, TBC systems have allowed much higher TETs than were previously
possible, failures due to a mismatch between thermal expansion coefficients of TBC system
layers during operation have limited their lifetime. The failure mechanisms of TBC systems
have been classified as extrinsic and intrinsic as defined by Evans et al [11]. The extrinsic failure
mechanisms are determined by external factors (e.g., erosion and foreign object damage, etc),
while most of intrinsic failures are governed by strain misfits between the BC/TGO/TC
interfaces upon thermal cycling. These mechanisms are always manifested as TC delamination
and hence, the eventual failure of the part. Currently, much research in the high-temperature
coatings area is focused on improving the lifetime of the TBC systems by doping the BC with
reactive elements [7,12–14], using different deposition techniques [2,3] or developing oxidation
treatments prior to TC deposition in order to control the TGO properties [15–20].
2
1.2
Justification
Notwithstanding that much research has been done in this area during recent decades, the
constant demand for increased efficiency of the gas-turbine engines, and therefore, improved
performance of TBC systems at high temperatures, keeps this an active field of research. Today,
the manufacturing process of TBC systems with β-(Ni,Pt)Al diffusion coatings includes diverse
stages that can influence its oxidation behavior during service, and therefore the TBC system’s
lifetime - electrolytic Pt plating, heat treatment, and aluminizing, surface preparation, and
oxidation treatments under different oxygen partial pressures, pO , prior to TC deposition.
2
Important contributions to the understanding of the TGO kinetics for nickel aluminides have
been reported in the last two decades, some of which suggest that oxidation treatments prior to
TC deposition improve its oxidation resistance and extend the lifetime of TBC systems by a
factor of two or three [15,17,18,21]. Despite of pre-oxidation treatments having shown an
improvement in the lifetime of TBC systems, there is a lack of information concerning the effect
of oxidation conditions on the competitive growth between the metaestable phases of Al2O3 (γ
and θ) and the stable phase α-Al2O3. Furthermore, there are no reports regarding their effect on
the physical and chemical properties of the different inner layers that compose the system. The
present work considers that knowledge of these effects would help to optimize the pre-oxidation
parameters, and hence, increase the durability of TBC systems by improving TGO adhesion,
and reducing BC and SA degradation. This work is focused on the study of the effect of
temperature, grit-blasting and oxygen partial pressure on the oxidation kinetics of commercial
β-(Ni,Pt)Al BC systems manufactured by GE-Aviation, in order to optimized the region (
temperature and time) where α-Al2O3 is the phase that controls the growth of the TGO.
3
2. Background
2.1 Thermal Barrier Coating Systems
The main purpose of thermal barrier coatings (TBC) systems is to provide thermal insulation for
the parts exposed to the highest temperatures in gas turbines engines (e.g., combustor liners, rotor
blades, and stators blades), but they also provided chemical protection. TBC systems are classified
as multilayer and multifunctional high-temperature coating system formed by four different layers:
1) a nickel-based superalloys providing the mechanical resistance under load bearing at high
temperature including tensile, creep and fatigue strength during operation.
Nowadays, internal cooled superalloys are also designed to allow cool air to flow inside of the part
to restrict the metal temperatures to reasonable level [2]; 2) an Al-rich BC to allow formation of an
Al2O3 scale, which prevents substrate oxidation, and it increases hot corrosion resistance [8]; 3) a
TGO as a result of the BC oxidation, desirably α-Al2O3. TGO has two principal functions: it binds
the ceramic layer to the bond coat, and it acts as a diffusion barrier to limit further substrate
oxidation [3]; 4) a ceramic TC as a thermal insulator, typically Y2O3-doped ZrO2 (7YSZ) with a
composition leading to a non-transformable tetragonal phase [11]. A common commercial TBC
systems produced by GE-Aviation consists of about 55 μm of β-(Ni,Pt)Al BC deposited onto Rene
N5 single-crystal Ni-based superalloy, followed by an ~ 150 μm thick 7YSZ TC is shown in Fig.
2.1. The bond coat is a bilayer structure consisting of an ~ 40 μm thick β-(Ni,Pt)Al layer and an
~15 μm thick inter-diffusion zone (IDZ).
4
Fig. 2.1. Cross-sectional micrograph of a thermal barrier coating system.
2.1.1 Composition (Component Layers)
2.1.1.1 Superalloys (SA)
Superalloys are defined as a Fe-, Co- and/or Ni-base metallic alloys, which can be used at high
temperatures up to approximately 85% of their incipient melting point. Some superalloys,
particularly Ni-base alloys exhibit outstanding creep and stress-rupture properties, as well as good
resistance to static loading at temperatures up to1200°C. The used of nickel as a solvent in these
alloys can be justified on account of its stable face centered cubic (fcc) crystal structure, its
moderate cost and its low rate of thermally activated creep. A perspective about the performance
of superalloys and their manufacturing processes since they began to appear in the 1940s is shown
in Fig. 2.2. This figure, based on the most representative superalloys for each period, shows how
their maximum operating temperature (the highest temperature at which creep rupture occurs with
not less than 1000 h, at 137MPa) has increase over the years. This improvement is linked to
optimizations of their chemical compositions and improvements in their manufacturing processes
(e.g., conventional casting, directional solidification and the use of single crystals). Superalloys
derive their strength mostly from solid solution hardening in form of an austenitic fcc matrix γ
phase and precipitated phases (Fig. 2.3a). The principal strengthening precipitated phases are the
ordered fcc γ´ phase, Ni3(Al,Ti) and the ordered bct γ´´ phase, Ni3Nb. Other phases can be found
in superalloys as a result of service aging conditions or defects created during the manufacturing
process. Examples include the topologically closed packed (TCP) phases μ, σ, P, R, etc (Fig. 2.3b).
5
Fig. 2.2. Progress of the temperature capability of the superalloys over the past 60 years (Y-axis
temperatures are in °C) [22].
Fig. 2.3. Examples of the phases formed in a René N5 Ni-base superalloy: (a) the high-temperature
strength phases γ and γ´, and (b) TCP precipitates.
6
Superalloys contain a large number of elements (often more than ten) to produce the desired effects.
A first group of elements includes Ni, Co, Fe, Cr, Rh, Mo, Re, and W, which stabilizes the
formation of austenitic γ and provides solid solution strengthening. A second group of elements
includes Al, Ti, Nb, and Ta, which are precipitate γ´ and γ´´precipitate formers. The various
elements are added for one or more purposes. For instance, Mo,Ta, W, and Re provide strength, Al
and Cr provide oxidation resistance, and Ti provides hot-corrosion resistance [23].
The blades in current gas-turbine used in the aeronautic industry are mainly manufactured by
directional solidification via single-crystal technology in order to remove the grain boundaries and
hence, improve mechanical properties. The blades are grown along the [001] direction, which
exhibits the best fatigue strength. The performance of γ/γ´ Ni-base superalloys is strongly
dependent on four guidelines [2]. First the amounts of γ´- forming elements such as Al, Ti and Ta
should be sufficient, such that the γ´ fraction is about 70%, which has better performance. Second,
the chemical composition of the alloy must be optimized such that the γ/γ´ lattice misfit is as small
as possible, to minimize the interfacial energy between γ/γ´and restrict γ´ coarsening. Third the
concentration of Re, W, Ta, Mo and Ru must be sufficient to improve the creep resistance and
thermal fatigue, but no so great to promote the precipitation of topologically close-packed phases.
Finally the composition must be chosen such that surface degradation (e.g., oxidation or hot
corrosion phenomena) by the combustion gases is minimized without compromising the structural
properties. Additionally, the used of different bond coat are often needed to cover these demands
including diffusion coatings (e.g., NiPtAl, NiPdAl, etc) and overlay coatings (MCrAlYs, where M
= Ni, Co or Fe).
2.1.1.2 Metallic Bond Coat (BC)
Superalloys used in very high-temperature industrial applications are regularly designed to
optimize their structural properties while maintaining microstructural stability over a wide range
of temperatures. However, the desirable mechanical properties are achieved only at the expense of
their environmental resistance. Trying to overcome this disadvantage, different metallic bond coats
have been developed parallel to the superalloy technology to provide oxidation and hot corrosion
protection. Bond coats achieve these requirements by promoting the formation of an adherent and
7
thermodynamically stable oxide scale with a slow growth-rate (e.g., α-Al2O3, Cr2O3, SiO2), which
acts as a diffusion barrier against oxygen and corroding salts.
The coatings themselves are thermodynamically stable over a wide range of chemical composition
to avoid the formation of undesired phases and they have low diffusion rates across their interfaces
at operating temperatures. Furthermore, the exhibit a good adherence with the substrate [3,23] and
their properties match well with those of the top coat and substrate to reduce thermal stresses and
to withstand service conditions. Metallic bond coats are divided in two main categories: diffusion
coatings and overlay coatings.
Diffusion coatings are formed by the surface enrichment of an alloy with aluminum (aluminides),
chromium (chromized), or silicon (siliconized) to the formation of intermetallics compound such
via diffusion process (e.g., NiAl, CoAl, etc) [24]. Most diffusion coatings contain all the elements
present in the base superalloys, but they are by a large quantity of aluminum. Although chromizing
and siliconizing are used, aluminizing is by far the most common type of diffusion coating used in
the gas-turbine industry.
2.1.1.2.1
Nickel Aluminides
The main matrix phase in aluminide coatings used on Ni-base superalloy substrates is the β-NiAl
phase. β-NiAl is an ordered intermetallic compound that crystallizes in a primitive cubic CsCl (B2)
structure, which consists of two interpenetrating primitive cubic cells, where Al atoms occupy the
cube corners of one sublattice and Ni atoms occupy the cube corners of the second sublattice. βNiAl is strongly ordered and thermodynamically stable over a wide range of compositions, from
about 45 at.% to almost 60 at.% Ni at 400°C and as high as 70 at.% Ni at 1360°C (Fig. 2.4).
Properties such as lattice parameter and density have been shown to be highly dependent on
chemical composition as a result of structural defects (Fig. 2.5). A similar dependence has been
observed for the elastic properties of NiAl single crystals, which exhibit elastic anisotropy. At room
temperature, the elastic anisotropy factor, A=2C44/(C11-C12) is 2.56 for Ni-52% Al, while A=8.95
for Ni-36.8% Al [25,26].
8
Fig. 2.4. Ni-Al phase diagram [27].
Fig. 2.5. Room temperature lattice parameter and density of β-NiAl as a function of stoichiometry
[28].
9
Deviations from stoichiometry are accommodated with two main structural defects. At
compositions greater than 50 at.% Al, Ni vacancies (VNi) are formed, whereas for Al contents less
than 50 at.%, the structure is compensated with Ni antisites (AlNi). When deviations are large
enough to destabilize the β-NiAl, other phases can be observed. For instance, for Al-poor β-NiAl
(Al < ~37 at. %), cooling from temperatures in excess 1100°C results in a reversible martensitic
transformation from β-NiAl↔L10-NiAl. L10-NiAl phase is face-centered tetragonal (fct). The
lattice parameters for the L10 phase are a = 0.3793 and c= 0.3183 nm for the stoichiometric NiAl
coating [29]. Martensitic transformation is a diffusionless, hysteretic transformation where the βNiAl phase is the high temperature face while L10-NiAl is the lower-temperature phase[30]. Chen
et al. [31] showed that martensitic transformation results in a volume decrease of approximately
2% since the β structure is larger than L10. The linear strain generated by this phase transformation
is 0.7 %. Even although β-NiAl ↔ L10-NiAl transformation is not a requirement for the occurrence
of rumpling, repeated transformations induce cracking in the coating [32]. As a result of the Al
depletion in the bond coat, because of TGO formation (outward diffusion) and the inward diffusion
of Al into the IDZ and superalloy, the thermodynamic conditions for the formation of the γ´-Ni3Al
phase are created, and therefore the coating microstructure shifts into the two-phase β + g´ zone.
γ´-Ni3Al has an ordered fcc Cu3Au (L12) type crystal structure with a lattice parameter a = 0.3561
nm [23]. Pint et al.[32] showed that the reversible β-NiAl ↔ γ´-Ni3Al phase transformation
produces an approximately 3% volume decrease on cooling. It has been observed that sufficiently
high cooling rates are necessary to suppress the β-NiAl → γ´-Ni3Al transformation.
2.1.1.2.1.1
Mechanisms of Diffusion Coating Formation
In recent years different processing methods have been employed in the manufacture of diffusion
coatings. These methods included above-the-pack cementation process, pack cementation process
including slurry, overpack chemical vapor deposition (CVD) and gas phase CVD. The following
steps are common among all the methods: 1) generation of the vapors containing the reacting
species (e.g., Al-,Cr-,Si-,Hf-, etc); 2) transport of the reacting species to the surface substrate alloy;
3) reaction of the species with the substrate followed by their internal diffusion; 3) Heat treatments
are usually required to achieved the desired coating composition [3].
The formation mechanisms of diffusion aluminide coatings on Ni-base superalloys were first
described qualitatively by Goward and Boone [33]. They claimed that only two types of coatings
10
can be formed on Ni-based superalloys. The first type was denominated as a high-activity
aluminide process, which is characterized by inward diffusion of aluminum, with a sufficient
aluminum activity to cause the formation of the δ-Ni2Al3 phase. A subsequent stabilizing heat
treatment causes an outward diffusion of nickel to form a three-zone BC (Fig. 2.6a). An outer layer
with a β-NiAl matrix containing substrate alloying elements (e.g., α-Cr, carbides, etc), a middle
single-phase β-NiAl layer and an inner IDZ layer with a β-NiAl matrix a different dispersed with
carbides (MC and M23C6) or σ (Cr,Mo,Co). The second type was called as a low-activity aluminide
process, which is characterized by outward diffusion of nickel from the substrate alloys to react
with the aluminum to form a two-zones BC with a β-NiAl phase in the outer layer and an IDZ in
the inner layer (Fig. 2.6b). This IDZ is formed by the loss of Ni from the alloy, resulting in the
formation of NiAl, which has very low solubility form many if the alloying constituents of the SA,
which tend to precipitate in this layer.
Fig. 2.6. Microstructure and phase identities during diffusion coatings manufacture. (a) high
activity aluminide process and (b) low activity aluminide process [33].
Janssen and Rieck [34], followed later by Shankar and Siegle [35] show that the diffusivities of Ni
(DNi) and Al (DAl) are highly dependent on stoichiometry. It was shown that Ni is the predominant
11
diffusing species in low-Al NiAl with DNi/DAl ≈ 3.0-3.5. An abrupt decrease is observed in intrinsic
diffusion coefficient ratio near to the stoichiometric composition, where the Al becomes the
predominant diffusion species in high-Al NiAl with a DNi/DAl ≈ 0.1 at about 52.5 at.% Al (Fig. 2.7).
Fig. 2.7. (a) Interdiffusion as a function of Al content and (b) DNi/DAl ratio (at 1100°C) in β-NiAl
as a function of Al content [35].
Today, CVD processes are the methods of choice for applying aluminide coatings to blades because
of their capability to coat small and internal cooling passages where the traditional pack
cementation processes cannot reach. In CVD processes, an halide vapor such as AlCl3 or AlF3 is
generated by passing HCl or HF gas over an Al-containing donor alloy (e.g. pellets of 40% Cr80%Al) at a temperature around 300°C outside of the coating retort. The parts to be coated are
fixed in the CVD reactor chamber, suitably masked to protect the areas susceptible to fatigue
failure, because whereas diffusion coatings improve the environmental life, a negative effect can
be observed on fatigue life. The halide vapors are transported to the external and internal surfaces
to be coated using a carrier gas consisting in a mixture of H2 and Ar. In the presence of H2, the
AlCl3 species react with more Al to form AlCl, which in turn reacts with Ni in the surface to form
the β-NiAl phase according to the following reactions [3,7,36–38]:
6 HCl + 2Al → 2AlCl3+3H2,
2AlCl3+4Al → 6AlCl,
6AlCl+6Ni+3H2 → 6NiAl+6HCl.
Diffusion coatings manufactured by CVD methods are generally a low-activity aluminide process,
typically held at a temperature in the range of 1000-1100°C for 2-12 hour. CVD processes exhibit
12
several advantages over the traditional pack-cementation methods, the most important being the
possibility to introduce different reactive elements (Hf, Si, Zr) into the diffusion coatings (Fig. 2.8).
Fig. 2.8. Schematic diagram of a CVD low-activity co-deposition reactor for aluminum and more
reactive elements [39].
Independent of the deposition method, at temperatures greater than about 1000°C high-temperature
oxidation and hot corrosion are the most significant forms of environmental attack observed with
nickel aluminide coatings. Under these conditions, the traditional diffusion aluminide coatings
offer limited protection, and therefore modified nickel aluminide coatings have been produced [4].
The alloying elements include Hf, Zr, Y, Ce, Dy, Si, Cr, Ru, Pd, Pt, etc. For instance, Hf, Zr, Y,
and Ce enhance the oxide scale adherence and improve cyclic oxidation performance [13,40,41];
Dy has improved the oxide scale adherence, as well as lowered its growth rate of oxide scale. Si
enhances oxidation and type II hot corrosion resistance [42–44]. Cr contributes to oxidation
resistance, reduces the Al requirement for formation of Al2O3 scale and gives resistance to hot
corrosion [21,41,45,46]. Ru and Pd have been proposed as a possible solution to improve the hightemperature BC capabilities (e.g., RuAl has a higher creep strength, Pd has a high solubility in βNiAl, eliminating the precipitation of brittle intermetallic phases) and to replace Pt with less costly
material [9,47–49]. Since early 1970s Pt has been the alloying element most widely used against
oxidation and corrosion phenomena in high-temperature applications. Many of these coatings are
commercially available today; however, the most significant advances have been for the Ptmodified nickel aluminide class.
2.1.1.2.2
Pt-Modified Nickel Aluminides
13
Pt-modified nickel aluminides have proved to be more effective in improving the oxidation during
thermal cycling and hot corrosion resistance than the conventional nickel aluminides [6,50–53]. It
shows an improvement the oxidation capabilities by a factor of 4 and resistance to hot corrosion
(type I) by a factor of 2. The first commercial coating of this class was designated LDC-2 by
Bungardt et al. [52]. The Pt-deposition was made via electrolytic methods as a thin layer from 210 μm (at temperatures of 80-90°C), although physical vapor deposition or CVD methods can be
also used. In order to guarantee good Pt-diffusion into the substrate a heat treatment in vacuum or
an inert atmosphere is carried out at 850-1000°C for 1-5 h prior to the aluminizing process
including high or low-activity aluminide processes (e.g., pack cementation or CVD methods).
Depending on aluminum activity (aAl), which is defined by the processing parameters and
subsequent heat treatments, three different microstructures of platinum aluminides can be formed
(Fig. 2.9).
Fig. 2.9. Possible microstructures of the Pt-aluminide coatings made by CVD process. (a) CVD
external single phase PtAl2 outer layer, (b) CVD external two phase platinum aluminide and (c)
MDC-150 external single phase platinum aluminide [54].
Pt enters the coating as single-phase ξ-PtAl2 which is a fluorite structure (C1), as two-phase ξPtAl2 and solid solution β-(Ni,Pt)Al or as a single phase β-(Ni,Pt)Al. First-principles calculations
have determined that Pt always has preference for the Ni sublattice in Al-rich, Ni-rich and
stoichiometric β-NiAl [54]. Originally it was believed that the Pt would function as a diffusion
barrier, halting the aluminum migration from the surface where it is needed to form the alumina
scale [55]. However, Smialek and Lowell [56] showed that inward diffusion of Al from the coating
into the substrate at high-temperature conditions does indeed occur which results in coating
degradation. Although the Pt addition does not result in the formation of a diffusion barrier, it
14
nevertheless enhances the oxidation and hot-corrosion resistance by promoting the slow growth of
a pure alumina scale.
Other mechanisms have been reported trying to explain the role of Pt in the Pt-modified aluminide
coatings. Tawancy et al. [57] showed that Pt eliminates Cr-rich precipitates from the outer coating
layer, and it prevents the diffusion of refractory elements (e.g., Mo, V and W) into the outer coating.
Farrell et al. [55] and Niu et al. [58] reported that the presence of Pt significantly improves the
resistance to β-NiAl → γ´-Ni3Al transformation. It was shown that Pt accelerates aluminum
diffusion (with respect to other alloy components) and hence, the vacancy flux from the metal to
the surface is reduced, avoiding vacancy coalescence (mainly in the BC/TGO interface) and
internal void formation [59–61]. Marino and Carter [62] showed based on first principles
calculations that Pt enhances Ni and Al diffusion in β-NiAl as a result of the decrease in the
activation energy of diffusion mechanisms, which is due to the stabilizing effect that Pt has on
defects. Therefore, diffusion of Al toward the BC/TGO should be good enough to keep Al levels
sufficient to form alumina and interfacial void formation in comparison with nickel aluminides
coatings.
Fountain et al. [63] suggested that the high-temperature stability of α-Al2O3 in Pt-modified
aluminide coatings is associated with diverse stress relief mechanisms. Hou [64] reported that Pt
usually suppresses S segregation to the BC/TGO interface, however this effect is highly dependent
on the chemical composition (primarily Ni content) and/or the presence of different alloying
elements. Finally, Warnes [8] showed that the co-deposition CVD method designated MDC-150L
(Fig. 2.8) showed a four times greater cyclic oxidation resistance than the two-phase type platinum
aluminide coatings.
2.1.1.2.2.1
Pt-Modified Nickel Aluminide Coating Stability
In general, two processes control the microstructural evolution of the Pt-modified nickel aluminide
coatings during service: 1) the oxidation of the BC that lead to the growth of the Al2O3 TGO scale
layer and 2) the interdiffusion of elemental species between the BC and SA because the BC is not
in thermodynamic equilibrium with the SA. Basuki et al. [65] and Kim and Walker [66] showed
that the microstructural evolution of BCs is highly dependent on the type of thermal exposure
history and the presence of TGO. Differences in their chemical compositions coupled with the
15
high-temperature operating conditions and time, provide the driving force to activate the
interdiffusion of elements across the TGO/BC/SA interfaces, causing coating degradation (Fig.
2.10).
Fig. 2.10. Schematic representation of the Pt-modified nickel aluminide bond coat degradation
[66].
Depletion of Al due to TGO formation and the interdiffusion of Al, Ni and Pt between single phase
β-(Ni,Pt)Al BC and Ni-base SA follows a gradual transition along the Ni-Al-Pt system (Fig. 2.11)
according with the general sequence of reactions:
Al-rich β-(Ni,Pt)Al
→
Ni-rich β-(Ni,Pt)Al,
Ni-rich β-(Ni,Pt)Al
→
β-(Ni,Pt)Al + γ´
→
L10-(Ni,Pt)Al + γ´,
β-(Ni,Pt)Al + γ´
→
Pt (in alloy) + γ´,
L10-(Ni,Pt)Al + γ´
→
Pt (in alloy) + γ´,
γ´
→
Al (in alloy) + γ.
16
0.0
0.1
1.0
0.9
0.2
0.8
0.3
0.7
δ-Ni2Al3
ξ-PtAl2
0.6
0.5
Al
Ni
0.4
0.5
β-NiAl
0.6
1100°C
ε-PtAl
0.4
Martensitic Trans.
0.7
α-NiPt
Pt3Al
3
0.8
0.9
0.3
γ'-Ni Al
0.2
0.1
γ-Ni
1.0
0.0
0.0
0.1
0.2
0.3
0.4
0.5
0.6
0.7
0.8
0.9
1.0
Pt
Fig. 2.11. Ternary diagram of Ni-Al-Pt system at 1150°C (the β region at 1100°C is also included).
Adapted from [67].
Zhang and Heuer presented in [68] a study of the microstructural evolution of a β-(Ni,Pt)Al
BC+TGO+7YSZ multilayer system on a René N5 Ni-base superalloy turbine blade taken out of
service. The study was performed on samples taken from the hottest and coolest sections of the
blade, which were compared with an as-deposited system. Based on this, it was concluded that the
as-deposited BC microstructure consist of a single-phase β-(Ni,Pt)Al, a middle layer of a β(Ni,Pt)Al matrix containing a high density of μ-precipitates and an inner layer with a γ/γ´ matrix
and numerous μ- and σ-phase precipitates. During service, the microstructural changes in the hotter
sections than the cooler parts as expected. Because of the interdiffusion the inner layer extents into
γ/γ´-superalloy, as some M23C6 carbides form. The matrix in the middle layer transforms to a β/γ´
mixture, and the precipitate phases undergo coarsening. The outer layer remains single-phase β,
but since diffusion induced composition changes in the hottest sections this layer undergo
martensitic transformation on cooling. Similar results have been reported for similar systems such
as [29,31,65,69–71].
17
2.1.1.2.3
Pt-Modified γ+γγ´ Diffusion Coating
As discussed above, β-(Ni,Pt)Al BCs are designed to have relative high Al content to promote
healing of the Al2O3 TGO scale following spallation. As result of this Al-enriched composition,
these kinds of BCs are not compatible with the chemical composition of γ/γ´-superalloy substrates
and interdiffusion phenomena take place between the BC and SA trying to achieve a
thermodynamic equilibrium during service. Because of this, efforts have been conducted in the
recent years to develop a new generation of BCs that are so-called Pt-modified γ-Ni + γ´-Ni3Al
diffusion coatings, which have a higher chemical and microstructural compatibility with the
substrate than the β-(Ni,Pt)Al BCs and avoid the formation of the brittle TCP phases in the region
of the original BC/SA interface. Recently Gleeson et al. [32,72] patented a Pt + Hf-modified γ-Ni
+ γ´-Ni3Al coating that offers a viable alternative to current β-(Ni,Pt)Al BCs. Ford et al. [73]
showed that the addition of Hf reduces the oxidation rate by an order of magnitude in comparison
with β-NiAl. Deodeshmukh and Gleeson [74] reported that Pt + Hf-modified γ-Ni + γ´-Ni3Al
coatings exhibit excellent type II hot corrosion resistance when given a 6 h pre-oxidation treatment
in air at 1080°C. Fig. 2.12 shows a cross section view of an as-coated Pt-modified γ-Ni + γ´-Ni3Al
coating on Rene N5 Ni-base superalloy reported by [75], which is composed of an outer two-phase
Pt-modified γ-γ´ layer with Pt-enriched regions (lighter contrast) and an inner multi-phase IDZ.
Fig. 2.12. Cross-section from the as-deposited Pt-modified γ-Ni + γ´-Ni3Al coating on Rene N5
superalloy. Adapted from [75].
18
2.1.1.3 Ceramic Top Coat (TC)
The continuing need to increase the efficiency of gas-turbines engines has generated a race to
develop new materials including superalloys, metallic coatings, and ceramic coatings that allowed
the TET to be increased.
Fig. 2.13. Advancement of the gas-turbine engine temperature capacity as a function of materials
development relative to Rene 80. Adapted from [76].
Fig. 2.13 shows an overview of this development in the last 40 years. What is important to point
out is the impact that TBC system technology has had on the TET since the mid-1990s when
improvements in superalloys reached a plateau with the development of the single-crystals
components. Since that time efforts have been made in order to improve the high-temperature
stability of the materials used as thermal insulators, as well as the compatibility between the layers
that form the TBC systems. The materials to be used as TC should cover a number of requirements
including high melting point, low thermal conductivity, high coefficient of thermal expansion,
high-temperature phase stability, oxidation and corrosion resistance, and high strain tolerance. Not
many materials satisfy all these requirements, one exception being some zirconia solid solutions,
for instance 7YSZ, which has been the most used commercial material in gas-turbine engines.
2.1.1.3.1
Yttria-Stabilized Zirconia (YSZ)
19
Pure zirconia is monoclinic at room temperature and 1 atm of pressure. With increasing
temperature, the monoclinic (m) phase transforms to tetragonal (t) at about 1700°C and then to
cubic (c) fluorite structure at about 2370°C. Since the martensitic transformations from c → t →
m on cooling involve a volume decrease of 2.31 and 4.5% respectively, pure zirconia is unsuitable
for TC applications where an intact solid structure is required during service. With the aim of
stabilizing the t-ZrO2 (or c-ZrO2) phase different oxides with a lower valance cation (and usually
larger than Zr+4) have been used such as La2O3 and Y2O3 [77]. These lower valance cations stabilize
the high-temperature structures at room temperature (designated as t´ or c´ to denote the
nonequilibrium conditions) through the creation of oxygen vacancies, which are required for
charge compensation [76]. Nowadays, the majority of commercial TCs are 7-8 mol% YO1.5-ZrO2
coatings with a t´ structure, which have shown the longest lifetime under thermal cycling tests [78].
Considering nonequilibrium thermodynamic conditions, it is possible to establish an upper (T0(c/t))
and lower (T0(t/m)) boundary of temperature and compositions where the metaestable t´ phase can
exist (Fig. 2.14).
Fig. 2.14. A binary phase diagram for the ZrO2-YO1.5 system showing the region where the
metastable tetragonal prime (t´) YSZ solid solutions are stable up to 1300°C. The superposed
maroon curve represents the relative cyclic durability as a function of YO1.5 content previously
reported by Stecura [78]. This diagram was adapted from [79].
20
These boundaries represent the free energy at which two solid solutions with the same composition
coexist as a function of temperature. The position of the T0(t/m) curve has been confirmed through
thermal cycling tests on electron beam physical vapor deposition (EBPVD) deposited TBC systems
[80].
2.1.2 Failure Mechanisms
The durability of the TBC systems is always limited by the spallation of the insulating TC. Once
some spallation events occurs small cracks are formed and extended along the TC/TGO or
TGO/BC interface causing spontaneous delamination of large areas of TC. Evans et al. [79] based
on a very extensive compilation of failed parts (e.g., blades, shrouds, combustors) and laboratory
specimen analysis, classified the failure mechanisms in two categories: intrinsic and extrinsic
mechanisms (Fig. 2.15).
Fig. 2.15. Schematic representation of the failure mechanisms in TBC systems proposed by Evans
et al. [79]
Based on this classification, the intrinsic mechanisms are always governed by the strain misfit
observed between TBC layers upon the thermal cycling inherent during service and are always
21
related with a critic TGO thickness hcrit, which depends on the specific system and its thermal
cycling history.
Intrinsic failures include: 1) the phenomenon known as rumpling, which is manifested by the
progressively roughening of the BC surface leading to development of localized tensile stresses
perpendicular to the interface at the convex regions and subsequent crack formation along the
interface [81–83]; 2) edge-delamination in the TGO/SA or TGO/TC due to brittle failure by loss
of adhesion and/or delamination coalescence between vertical separation defects at TC columns
[84]; 3) void formation in the BC and TGO/BC interface due to stress-driven vacancy coalescence
mechanisms, related with the volume change as β-NiAl transforms to γ´-Ni3Al [69]. Extrinsic
failures are always associated with the insulating TC, and include: 1) particle impact (e.g., erosion
and foreign object damage) and 2) delamination resulting from the penetration of calciummagnesium-alumino-silicate (CMAS) deposits formed by the contact of sands and dust in the
atmosphere on the hottest sections of the gas-turbine engines and subsequent cooling and heating
cycles [85–87].
2.2 Oxidation
2.2.1 Thermodynamic Principles of Oxidation
In order to determine if a chemical reaction will occurs or not at a given constant temperature and
pressure, the thermodynamic potential known as Gibbs free energy (G) can be used. The Gibbs
free-energy change, ΔG, of a system can be defined as
∆G = ∆H - T∆S
2.1
where H is the enthalpy, S the entropy and T the absolute temperature. The second law of
thermodynamic states three cases: 1) when ΔG < 0, a spontaneous reaction is expected, 2) when
ΔG = 0 an equilibrium state is established and 3) when ΔG > 0 which represents a thermodynamic
impossible process.
For an oxidation reaction, for instance,
y M + O2 → y Mx Oy
2x
2
2.2
ΔG is expressed as
22
o
∆G = ∆G + RT ln
2/y
x y
2x/y
a M a O2
aM O
2.3
where ΔGo is the free energy change when all the species are present in their standard states; a is
the thermodynamic activity which describes the deviation from the standard state of the involve
species, for example, for a given specie i , this can be defined as
a=
pi
2.4
poi
where pi is either the vapor pressure over a condensed species or the partial pressure in the gas
species and poi is the same quantity corresponding to the standard state of i. If we assumed the metal
and oxide activities as unity, as well as the poO = 1 atm, we can used the equation 2.3 to estimate de
2
dissociation pressure of an oxide in equilibrium normalized to 1 mol O2 as
M/MO2
pO
2
= exp RT ∆Go
M/MO2
Once estimated the pO
2
4
3
2.5
for the formation of a specific oxide is estimated, for instance for
2
Al + O2 → 3 Al2 O3
2.6
it can be interpreted as follows: 1) for an environmental oxygen partial pressure larger than the
equilibrium dissociation pressure, pO > pAl/Al2 O3 , the formation of Al2O3 will take place, 2) for the
2
Al/Al2 O3
case when pO = pO
2
2
4
O2
, the system will be in a dynamic equilibrium where the metal and the
2
Al/Al2 O3
oxide co-exist, 3 Al + O2 ↔ 3 Al2 O3 , and 3) for the case when pO < pO
2
2
wherein the metal
will remain stable, without oxidizing.
This analysis s represented graphically for Al2O3 and several other oxides in the form of an
Ellingham-Richardson diagram (Figure 2.16), which ∆Go is the y-axis and temperature is the x-axis.
The usefulness of this nomograph is as follows. Firstly is possible to known the stability of several
oxides based on their ∆Go values at a given temperature. The oxide with the most negative value
will be stable. For instance, a comparison comparison between the Al2O3, SiO2 and Cr2O3 shows
that Al2O3 is more stable (has more negative ∆Go ) than SiO2 and Cr2O3 over the entire range of
M/MO2
temperatures. It is also possible to know the pO
2
at a given temperature by drawing a straight
23
line from the origin marked O through the ∆Go line at the temperature of interest and reading its
M/MO2
pO
2
at the intersection with the outer scale labeled as pO (which travels along the x- and y-axes).
2
Fig. 2.16. Ellingham diagram showing the standard free energy change for selected oxides
formation as a function of temperature. Adapted from [88].
24
M/MO2
For example, at 1000°C, the pO
2
values for the Al2O3, SiO2 and Cr2O3 are about 1 x 10-21, 1 x
10-26 and 1 x 10-34 atm respectively, which also confirm that Al2O3 formation is more stable than
SiO2 and Cr2O3.
In order to achieve the working pO diverse redox gas mixtures can be used (e.g., CO-CO2 and H22
H2O). These mixtures consist of an oxidizing and a reducing species, which equilibrate with
oxygen. Because of this, the equilibrium CO/CO2 or H2/H2O ratio with a metal at certain
temperature can also be found in the Ellingham-Richardson diagram by extending the line from
the origins marked as C or H respectively. Moreover, if the oxidation of an alloy is considered, the
activities of the metal and oxide must be taken into account to estimate the equilibrium dissociation
pressure between the alloy and its oxide.
2/y
2/y
aM O
aM O
∆Go
eq
x y M/MO2
x y
pO = 2x/y
exp RT = 2x/y
pO
2
2
aM
aM
2.7
where the aM will be a function chemical potential of the oxidized specie (μM) at the TGO/BC
interface as a result of the interdiffusion mechanisms presented through the alloy.
2.2.2 Mechanisms of Oxidation
Oxidation reactions are remarkably complex, because they are affected by several issues including
thermodynamics, alloy composition, kinetics, structure, microstructure, mechanical properties,
impurities, surface preparation, etc. Although thermodynamics is important to describe the
condition at which an oxide is stable or not, the diffusion of species through the bulk alloy (e.g.,
Al) and the oxide scale (e.g., Al3+ and O2- in an α-Al2O3) is what controls the growth kinetics of
the oxide scale during the oxidation process.
2.2.2.1 Transport Mechanisms
All crystalline materials in nature including metal and ceramics contain structural defects in their
crystallographic lattice. These defects can be classified according with their dimensionality as zero
dimensional (e.g., point defects), one dimensional (e.g., dislocations) and two dimensional (e.g.,
grain boundaries, interfaces). For their part, ceramics can also be classified as ionic crystals, and
unlike metals, point defects in ionic crystals are electrically charged. Two types of point defects
have been defined for ceramics - ionic defects and electronic defects. Ionic defects included
25
substitutional atoms, interstitials and vacancies solutes pairs. Electronic defects represent
deviations from the atomic ground state of atoms that form the crystal in which electrons are excited
into higher energy states. Such excitation results in the creation of an electron in the conduction
band or electron hole in the valence band.
Kroger [89] developed a nomenclature in order to describe the ionic defects as follows:
MM = M atom on M site
XX = X atom on X site
Mi = M atom on interstitial site
Xi = X atom on interstitial site
NM = impurity N on M site
VM = vacancy on M site
VX = vacancy on X site
Vi = vacant interstitial site
e' = electron in conduction band
h· = electron hole in valence band
In this notation, the capital letter represents the species, the subscript describes the site at which
the species is located and the superscript describes the charge on the species with respect to the
host species where (·) refers to a net positive charge and (′) refers to a net negative charge and (x)
refers to no net charge difference. Defects in ionic oxides can also be categorized as extrinsic or
intrinsic defects, based on whether stoichiometric deviations are observed in the crystal structure
(extrinsic) or not (intrinsic). The compound Al2(1-δ)O3 is an example of an ionic oxide with extrinsic
defects. It has been found that extrinsic defects are the predominant defects in a growing oxide
scale [88]. Based on this, the mechanisms by which ions and electrons are transported through
metal oxides during the growth of the oxide scale can be described using the p-type and/or n-type
semiconducting oxide models [90].
The relative charged of an ionic defects is balanced by an oppositely charged electron defect (e.g.,
electron, e´ , or electron holes, h· ). It is the predominant electronic defect that determines whether
an oxide is p-type (h· ) or n-type (e´ ). The formation of Al2O3 (Eq. 2.6) gives a good example of
these transport mechanisms.
For an oxide like θ-Al2O3 which is mainly p-type it may have predominantly Al vacancies (V'''Al )
which means that there would be a predominance of Al diffusion (Fig. 17).
26
Metal
Oxide
4
4
4
Al + 4h⋅ + VAl''' = Al AlX
3
3
3
JV '''
or
Al
pO' 2
4 X
4
Al Al + 2OOX = Al + 2Oi'' + 4h⋅
3
3
J h⋅
pO'' 2
Gas
Metal vacancies
4
O2 = 2OOX + VAl''' + 4h⋅
3
J Al 3+
or
Oxygen intertitials
O2 = 2Oi'' + 4h⋅
p- type
Fig. 17. Schematic view of a Wagner’s diffusion model for a p-type oxide.
p-type oxides exhibit a structure with negatively charged defects, for instance, metal vacancies,
where the formation of a metal-deficit semiconductor with cation vacancies and positive holes by
the incorporation of oxygen into the oxide lattice can be represented with the following reaction
O2 = 2OXO + 3 V'''Al + 4h·
4
2.8
The equilibrium constant, K2.8, for this reaction may be written as in equation 2.9 assuming that
equation 2.8 represents the only mechanism by which the defects from and they obey Henry’s law:
K2.8 =
h· 4 V ´´´
Al
pO
4/3
2.9
2
In order to keep the electrical neutrality h· =3 VAl´´´ and substituting this relation into equation
2.9 gives:
VAl´´´ = 2.8
K
81
3/16
pO 3/16
2
2.10
In contrast, for an oxide like α-Al2O3, which is mainly n-type, it may have oxygen vacancies (V··O ),
meaning there would be a predominance of oxygen diffusion (Fig. 18).
For oxides, this defect structure can be visualized as the discharge and subsequent evaporation of
an oxygen ion; the electrons enter the conduction band and a vacancy is created on the anion lattice
[88]. This process can be described by
2OXO = O2 + 2V··O + 4e'
2.11
27
and its equilibrium constant can be written as
K2.11 =VO·· 2 e' 4 pO
2.12
2
and in order to keep the electrical neutrality the following relation must be fulfilled e' = 2VO·· .
Based on this the oxygen vacancy concentration can be defined by
VO·· = 2.11
K
1/6
16
pO -1/6
2.13
2
Metal
Oxide
4
4
Al + 2VO⋅⋅ + 4e' = Al AlX
3
3
JV ⋅⋅
O
pO' 2
or
4
4
Al = Ali⋅⋅⋅ + 4e '
3
3
J e'
pO'' 2
Gas
Oxygen vacancies
2OOX = O2 + 2VO⋅⋅ + 4e'
or
J O 2−
Metal intertitials
4 X
4
Al Al + 2OOX = Al i⋅⋅⋅ + 4e ' + O2
3
3
n- type
Fig. 18. Schematic view of a Wagner’s diffusion model for n-type oxide.
Assuming that the diffusion coefficient is linearly proportional to the vacancy concentration (e.g.,
cation or anion vacancies), it is expected that growth rate constant of pure θ-Al2O3 will increase
with increasing the oxygen partial pressure to the 3/16 power, whereas the growth rate of pure αAl2O3 is expected to vary according to the negative –1/6 power.
Wagner and Grunewald [91] showed that the growth kinetic rate is only dependent on the
environmental pO , in cases where mainly negative charged defects (e.g., metal vacancies or
2
oxygen interstitials) are in oxide structure, while the kinetic rate was shown to be independent of
pO when positive charged defects (e.g., oxygen vacancies or metals interstitials) predominate as
2
will be described below.
2.2.2.2 Wagner´s Theory of Oxidation
28
In 1933, Wagner [92] presented his very well-known theory of the high-temperature oxidation of
metals, assuming a diffusion-controlled oxide scale growth. This was the first attempt to formalize
the phenomena behind the parabolic oxidation kinetics observed experimentally by Tamman in
1920 [93] for steel, and later by Pilling and Bedworth in 1922 and 1923 for heater alloys [94] by
the equation 2.14.
dx
dt
=
k´
2.14
x
where x is the oxide thickness, k´ is the parabolic rate constant and t is the time.
The Wagner oxidation theory was developed considering the following assumptions [88]:
1. The oxide layer is a compact, perfectly adherent scale.
2. Migration of ions or electrons across the scale is the rate-controlling process.
3. Thermodynamic equilibrium is established at both metal-scale and scale-gas interfaces.
4. The oxide scale shows only small deviations from stoichiometry and, hence, the ionic fluxes
are independent of position within the scale.
5. Thermodynamic equilibrium is established locally throughout the scale.
6. The scale is thick compared with the distances over which space charge effects occur.
7. Oxygen solubility in the metal is neglected.
Because a thermodynamic equilibrium is established at both metals-scale and scale-gas interfaces,
chemical activity gradients are also established across the oxide scale. The presence of these
gradients causes migration of oppositely charged species (e.g., cations, anions, vacancies, electrons
or electrons holes) in order to preserve the electroneutrality of the oxide. Because of the transport
species are charged, they react to both chemical and electrical potential gradients, which together
provided the net driving forces for ion migration.
If electrochemical potential gradients can be defined by
∇ηi = ∇μi + Zi F ∇ϕ
J mol-1cm-1
2.15
where μi is the chemical potential of species i, Zi is the charge of the transported species, F is the
Faraday´s constant and ϕ the local electrostatic potential. The equation 2.15 represents the force, fi,
acting on the mobile species i across the oxide scale, which can also be written as a function of the
transported particles using Avogadro´s number, NA, as
29
fi =
1 ∂ηi
1
NA
NA
=
∂x
∂μi
∂x
+Zi F ∂ϕ
∂x
J particle-1 cm-1
2.16
The drift velocity of a particle can then be estimated by means of equation 2.16 and the mobility
of species i as
νi = - Bi fi = -
Bi
NA
∂μi
∂x
+Zi F cm s-1
∂ϕ
∂x
2.17
where Bi is the particle mobility which is defined as the average drift velocity per unit of force,
particle cm2 J-1 s-1. The negative sing the equation arises since the drift velocity takes place in the
positive x direction for the negative electrochemical potential gradient.
Once established the drift velocity of specie i, the flux of i is obtained multiplying by the
concentration of i as
ji = Ci νi = - Ni i ∂xi +Zi F ∂x mol cm-2 s-1
CB
∂μ
∂ϕ
2.18
A
The mobility has been related to the conductivity κi, and the self-diffusion coefficient Di of the
particle for the limiting cases of zero electrical-potential gradient and zero chemical-potential
gradient to Fick´s first law and Ohm´s law, respectively, giving as a result the following
relationships
kB TBi = Di =
RTκi
2.19
Ci Z2i F2
where kB is Boltzmann´s constant. Using equation 2.19 and knowing that the gas constant R = kBNA,
the equation 2.18 can be rewritten as
ji = - 2 i 2 ∂xi +Zi F ∂x
κ
∂μ
∂ϕ
Zi F
2.20
This equation can be used to describe the flux of differently charged species through the oxide
layer. Due to their different mobilities, different species would tend to move at different rates in
order to maintain the electroneutrality throughout the scale. This condition is achieved due to the
very high mobility of the free carriers relative to those of ions, for example, Bh ≫ BVc (or
Be ≫ BVa ).
Considering the growth of a p-type oxide scale (e.g,. θ-Al2O3) sustained by metal vacancy diffusion
we can write the anion vacancy and electron holes flux using equation 2.20 as
30
jV´´´ = Al
κ ´´´
∂μ ´´´
VAl
2
9F
VAl
∂x
-3F ∂ϕ
2.21
∂x
jh· = - h2 ∂xh +F ∂x
∂μ ·
κ ·
∂ϕ
2.22
F
The local electrostatic field in equations 2.21 and 2.22 cannot be measured because any field that
develops a charge within an oxide would affect the flux of the other charged species as well. In
order to overcome this problem, the electrical neutrality condition defined by equation 2.23 can be
used
jh· -3 jV´´´ = 0
2.23
Al
which together with equations 2.21 and 2.22, allow you to eliminate the local electrostatic field via
∂ϕ
∂x
=
κV´´´
·
1
3F(κ ´´´ +κh )
VAl
∂μ ´´´
VAl
∂x
Al
-3κh·
∂μh·
∂x
2.24
Substituting equation 2.24 in equation 2.21, we can redefined the flux of anion vacancies by
equation 2.25
jV´´´ = Al
κ ´´´ κh·
2
VAl
∂xAl +3 ∂xh ·
∂μ ´´´
V
∂μ ·
9F (κ ´´´ +κh )
VAl
2.25
Considering that the local equilibrium reaction defined by equation 2.8 (a metal-deficit
semiconductor oxide) is achieved, it is possible to establish the relation given in equation 2.26
3
2
dμO = dμV´´´ + 3dμh·
Al
2.26
Using the relation of the equation 2.26, we obtain equation 2.27
1
jV´´´ = - 6 2
κ ´´´ κh·
VAl
dμO
F (κ ´´´ +κh· ) dx
Al
2.27
VAl
Assuming that the electrical conductivity of the electron holes is approximately one, compared
with aluminum vacancies conductivity, the equation 2.28 can be obtain
jV´´´ = Al
κ ´´´
VAl dμO
6F2 dx
2.28
Moreover, μO may be defined by equation 2.29
31
1
1
μO = μ°O + RTlnpO
2
2
2
2.29
2
Differentiating equation 2.29 with respect to thickness scale, and substituting this result into
equation 2.28, the equation 2.30 is obtained
jV´´´ = -
RTκ ´´´ dlnp
VAl
O2
12F2
Al
2.30
dx
Upon integration equation 2.30 from x = 0 to x = x, this yields to
jV´´´ = - ´ 2
RTκ ´´´
p´´
V
O
pO
Al
Al
12F2 x
2
2.31
dlnpO
2
If the oxygen vacancies concentration in the oxide is CV´´´ , then the flux of cation vacancies can
Al
be also written as
dx
jV´´´ = CV´´´ dt
Al
2.32
Al
and substituting the equations 2.14 in equation 2.32, we obtain
jV´´´ = CV´´´
Al
k´ ´´´
VAl
2.33
x
Al
From equations 2.31 and 2.33, the parabolic rate constant may be expressed by equation 2.34
p´´
O
k´V´´´ = - ´ 2
RTκ ´´´
VAl
pO 12F2 C ´´´
Al
2
2.34
dlnpO
2
VAl
Since electrical conductivity is expected to vary proportional to the anion vacancy concentration,
and so with the three sixteenths of the oxygen partial pressure as shown in equation 2.35
κV´´´ ∝ CV´´´ ∝ p3/16
O
Al
2.35
2
Al
the equation 2.34 can be rewritten as a function of the oxygen partial pressure by equation 2.36
p´´
O
dlnpO
k´V´´´ = ´ 2 C1 p3/16
O
Al
pO
2
2
2.36
2
which upon integration, yields equation 2.37
k´V´´´ = C1 p´´O 3/16
Al
2
- p´O 3/16
2
2.37
32
where p´´O and p´O represent the oxygen partial pressure at the working gas-oxide scale and metal2
2
oxide scale interfaces respectively.
A similar treatment can be applied now for the growth of a n-type oxide scale (e.g., α-Al2O3)
sustained by oxygen vacancies according with the local equilibrium reaction defined by equation
2.11
2OXO = O2 + 2V··O + 4e'
2.11
with the following equilibrium constant (Eq. 2.12), assuming that the defects are in a very dilute
solution to obey Henry´s law
K2.11 = C2V··O C4e´ pO
2.12
2
Excluding any impurity and intrinsic electronic defects, and in order to satisfy both stoichiometric
and electroneutrality (Ce´ = 2CV··O ), equation 2.38 can be written
κV··O ∝ CV··O ∝ pO -1/6
2.38
2
Moreover, following a procedure similar to that developed for the oxide with a p-type growth from
equations 2.21 to 2.34, the equation 2.39
p´´
O
k´V··O = ´ 2
pO
2
RTκV⋅⋅
O
2.39
dlnpO
8F2 CV⋅⋅
2
O
and using equation 2.38, the equation 2.39 can be rewritten as
k´V··O = ´ 2 C3 p-1/6
dlnpO
O
p´´
O
pO
2
2
2.40
2
which upon integration, yields equation 2.41
k´V··O = C3 p´´O -1/6
2
- p´O -1/6
2
2.41
Since generally p´´O is much greater than p´O , and so the parabolic rate constant for n-type oxides is
2
2
expected to be independent of the external oxygen partial pressure, unlike p-type oxides were
oxygen partial pressure controls its growth rate.
Notwithstanding, that Wagner´s theory does not apply to slow-growing oxides as Cr2O3 and Al2O3
(which have practical interest at high-temperature applications) since the transport properties of
33
these oxides are largely dominated by grain boundary transport, the mechanism is still diffusion,
and the basic concepts underlying Wagner’s theory can still used as basis to understand the
oxidation processes and the effect of some oxidation parameters on the oxidation behavior (e.g.,
oxygen partial pressure).
2.2.2.3 Diffusion Controlled Oxidation [via Grain Boundary Diffusion]
Many oxide scales growth too rapidly to be explained by lattice diffusion of species as assumed
oxidation Wagner´s theory, and therefore diffusion mechanisms must occur along different shortcircuit paths (e.g., chromia and alumina scales). It has been shown by TEM analysis, that only grain
boundaries could accommodate the species diffusion, since other diffusion paths, such as scale
dislocations, does not have enough density to allow the flux of ions [95–97]. Pint et al. [98] carried
out sequential oxidation experiments at 1200°C using a double oxidation experiments (using 16O
and 18O-enrichment environments) on undoped and Y- and Zr-doped β-NiAl and FeCrAl. At
1200°C, undoped α-Al2O3 showed a mixed diffusion mode involving simultaneous Al and O
transport via short-circuits pathways. Moreover, with addition of reactive elements such as Zr or
Y, Al diffusion was decreased and the α-Al2O3 scale was observed to growth mainly by inward
diffusion of O. Later, Pint et al. [99] also showed that θ-Al2O3 scales grow predominantly by
outward cation diffusion process at 950°C, since Al3+ is the primary diffusion specie, the 18O tracer
does not give reliable information about whether transport is via lattice or short-circuit diffusion.
However, based on the relative open θ-Al2O3 lattice, and its blade-like morphology, Al diffusion
through the bulk has been suggested. In 1975, Hart [100] showed that the effective diffusion
coefficient of the species i, Di,eff, in a polycrystalline materials may be expressed as weighted sum
of constants for lattice and short-circuit pathways (e.g. grain boundaries) as
Di,eff = Di,L ( 1 - f )+Di,GB f
2.42
where Di,L (T,pO ) and Di,GB (T, pO ) are the self-diffusion coefficients of the species i for the bulk
2
2
lattice and grain boundaries respectively (and are function of the temperature ), and f is the fraction
of sites in the grain boundaries, which can be defined by
f=
gδGB
λ
2.43
34
where λ is the grain size, g is a geometric factor which reflects the shape of the grain and δGB is the
grain boundary width. For grains, which are squared in cross-section and of edge length rG, the
equation 2.43 becomes
f=
2δGB
2.44
rG (x)
where, rG can be a function of scale thickness, as was showed by Naumenko et al. [101] for the
oxidation behavior of a model FeCr20Al5Y0.05 alloy, where the grain size of the formed α-Al2O3
scale exhibited a linear dependence with the distance, x, from the gas/oxide interface. Based on the
equation 2.44, the equation 2.42 can be rewritten as
Di,eff = Di,L ( 1 - f ) +
2Di ,GB δGB
rG (x)
2.45
f
Very useful reviews concerning the major diffusion processes in α-Al2O3 including lattice and grain
boundary diffusion were published by Heuer in 2008 and later complemented by Heuer et al. in
2011[102,103]. One of the most remarkable works in this field in the last years was published by
Wada et al. in 2011 [104] showing that the grain boundary diffusion mechanisms of Al3+ and O2in a non-doped polycrystalline α-Al2O3 wafers obey the following relationships
DAl,GB δ = 2.475 x 10-5 exp -604,000
DO,GB δ = 2.207 x 10-9 exp RT
-467,000
RT
pO
upper 3/16
2
plower
O
-1/6
2
2.46
2.47
Wada et al. [104] performed experiments to study the permeability of oxygen under different
oxygen partial pressure gradients from 1650 to 2000°C. Results for the samples exposed at 1650°C
for 10 h showed that under a low oxygen partial pressure gradient, e.g., between 10-8 Pa (lower
side) and 1 Pa (upper side), grain boundary grooves were observed at both surfaces of the alumina
wafer (Fig. 2.19a and c). The absence of grain boundary ridges suggests that the migration of
aluminum was scarcely in comparison with the oxygen diffusion. This surface morphology
supports the oxygen permeability results, which showed that a n-type conduction is taking place
through the alumina wafer following equation 2.11.
35
Fig. 19. SEM micrographs of the surfaces and cross-sections of polycrystalline α-Al2O3 exposed
at 1650°C for 10 h at two different ∆pO . (a) and (c) between 10-8 and 1 Pa respectively. (b) and
2
(d) between 10-8 and 105 Pa respectively. Adapted from [104].
For the sample exposed under a high oxygen partial pressure gradient, similar grain boundary
grooves were observed at the low pO (e.g., 1 Pa) surface, whereas that ridges rather than grooves
2
were observed at the high pO (e.g., 105 Pa) surface (Fig. 2.19b and d). Based on this, it is suggested
2
that oxygen permeated mainly via grain boundary diffusion of Al3+ through aluminum vacancies
from the lower pO surface to the higher pO surface, which result in the formation of grain-boundary
2
2
ridges. The Al2O3 formation would follow the defect reaction given in equation 2.8, where O2
molecules are absorbed onto the high pO surface and subsequent dissociated into O2- to react with
2
Al
3+
ions. However, the profiles of grain-boundary diffusion coefficients, and fluxes of Al3+ and
O2- showed the evidence that an interdiffusion process is really taking place along the alumina
wafer as shown in Fig. 2.20.
36
Fig. 20. Profiles of grain boundary coefficients and flux of Al3+ and O2- in polycrystalline wafer
upper
= 105 Pa/10-8 Pa. Adapted from [104].
exposed at 1650°C to pO /plower
O
2
2
An extrapolation of equations 2.46 and 2.47 to extreme oxidation conditions similar to those
upper
observed in alumina forming alloys [105], for instance at 1200°C to pO
2
/plower
= 105 Pa/10-23 Pa,
O
2
showed that the intersection of the diffusion coefficients, as well as, charged species were
upper
extremely shifted toward the pO
2
surface (Fig. 2.21).
Fig. 21. Profiles of grain boundary coefficients and flux of Al3+ and O2- in polycrystalline wafer
upper
exposed at 1200°C to pO /plower
= 105 Pa/10-23 Pa. Adapted from [104].
O
2
2
Meanwhile, Nicolas-Chaubet et al. [106] carried out a conductivity study on Al2O3 scale formed
on a β-NiAl alloy at 1100°C for 96 h, and concluded that major portion of the scale thickness is
governed by a low pO gradient, between about 10-3 and 10-5 Pa, with an abrupt pO changes
2
2
occurring at the alloy/scale and scale/gas interfaces.
Thus, based on the above evidences Heuer et al. [103] have suggested that the major portion of a
growing α-Al2O3 scale is n-type, with oxygen vacancies as the main defect structure, with p → n
37
transition occurring close to the scale-gas interface, being the responsible of grain-boundary ridges
formation.
Finally, a Wagner´s approach can be done to correlate the effective diffusion coefficient with the
parabolic rate constant assuming that Al3+ and O2- are the only species diffusing through the
alumina scale by
p´´
O
αDAl,eff
pO
f*Al
k' = ´ 2 2
+
DO,eff
f*O
p-1
dpO
O
2
2
2. 48
where α is the stoichiometric ratio of the oxide phase (for the case of Al2O3, α = 2/3), f*Al and f*O are
the correlation factors for the aluminum and oxygen ions for the self diffusion respectively [107].
2.2.3 High-Temperature Oxidation of Alloys
Most metallic materials of practical importance at high-temperature applications are alloys, which
operate at very aggressive environments that contain enough reactive species for
thermodynamically destabilize alloys and promote the growth of more stables compounds such as
oxides, nitrides, carbides, etc. Such conversion processes are generically termed oxidation.
The oxidation products and its growth rates are highly dependent on the interplay among
composition, microstructure, and surface conditions of the alloy with reactive species of the
oxidizing environment and the oxidation parameters (e.g. temperature, oxygen partial pressure,
etc.). In most cases, oxidation can lead to a loss of the structural capabilities of the alloy, because
of the thinning of the part due to the oxidation products. One of the strategies to overcome this
problematic have been the alloys design, which promote the growth of a protective oxide scale. For
long-term (>>1,000 hours) service above about 600°C, α-Cr2O3, α-Al2O3 and SiO2 are the principal
oxides used for protection because of their slow growth rate [108].
It is important to point out that maintain a protective oxide scale is often not a simple matter,
because of different problems associated with the oxide-scale adherence such as: 1) voids
formation (e.g., Kirkendall effects), 2) oxide-scale residual stresses, and 3) Al, Cr or Si depletion
to sustain the protective oxide-scale after spallation events (e.g., growth of spurious oxides).
Considering the two limit cases for an oxidation of a single-phase alloy where a binary alloy AB
for which AO, a rapid non-protective oxide, is less stable than BO, a protective scale such as Al2O3.
38
Assuming that AO and BO are mutually insoluble and oxygen partial pressure is enough to sustain
the growth of both oxides.
The two limits cases shown are the internal oxidation of the binary alloy forming internal BO
particles beneath the AO oxide (Fig. 2.22a), and the exclusive formation of a continuous protective
scale layer BO (Fig. 2.22b).
Alloy AB
BO particles
in a B-depleted
Matrix
AO
BO
Alloy AB
b)
a)
JB
Outward Flux of B
<
JO
JB
Inward Flux of O
Outward Flux of B
JO
>
Inward Flux of O
Fig. 22. Schematic representation of (a) an internal BO particles formation beneath an AO scale
and (b) the formation of an external protective BO scale. Adapted from [108].
The transition from the internal oxidation to external BO formation requires that a critical
concentration of B in the bulk alloy should be exceeded to maintain the BO stability at the AB
alloy/BO interface [108]. This critical concentration is highly dependent on thermodynamic,
kinetic, and even the surface conditions, since it have been observed than an increase on the density
of defects at the alloy surface by mechanical deformation (e.g., grit-blasting process) may promote
the preferential nucleation of the protective BO scale [18,109].
In some cases, additions of ternary (e.g., an ABC alloy) or higher order alloying elements can also
decrease the critical concentration of B to form BO, because of the so-called third-element effect
[2,110]. This concept was firstly proposed by Wagner [111], who suggested that this third element
C, acts as an oxygen getter limiting the oxygen diffusion into the alloy, and therefore the internal
oxidation of B is minimized [10,112,113].
39
Fig. 23. Schematic examples of oxidation behaviors observed in two-phase alloys for which each
phase form an oxide scale. (a) Formation of a mixed nonuniform oxide scale and (b) formation of
uniform and protective oxide scale. Adapted from [108].
On the other hand, many alloys used for high temperature applications are designed or can
developed a multiphase system during its operation. Although there are no many reports related
with the oxidation behavior of multiphase alloys, and the role of their individual phases, it can be
inferred that anything that improves the oxidation behavior for single-phase alloys, including the
additions of alloying elements to promote the third element effect, as well as increase the density
of defects at the alloy surface. Other key factors to be considered include size, shape, distribution,
volume fraction and composition of the matrix phase and minority phases [114–117].
Similarly to that observed for single-phase alloy two limit cases can be defined for ternary and
higher order two-phase alloys as is shown in Fig. 2.23. One limit case occurs when the two phases
are oxidized independently to form a non protective AO + BO scale (Fig. 2.23a), whereas that the
other limit case occurs when the solute-rich second phase acts as a reservoir for the continued
growth of the solute scale (Fig. 2.23b).
40
2.2.4 Alumina Forming Alloys
2.2.4.1 Ni-Al
In 1967, Pettit [118] presented his classic oxidation study over the Ni-Al system under a pure
oxygen atmosphere at 0.1 atm. Fig. 2.24 summarized the oxidation map obtained by Pettit as a
function of composition and temperature. Based on this, its oxidation behavior can be classified by
three different regions, corresponding to different reactions morphologies and mechanisms as
described below.
Fig. 24. Oxidation map for Ni-Al system under an O2 atmosphere at 0.1 atm according to Pettit
[118]. Adapted from [88].
Region I from about 0-13 at.% Al at 900°C, corresponds to dilute alloys which develop an external
NiO scale doped with Al and internal precipitates of Al2O3 and NiAl2O4 at temperatures between
900° to 1300°C. Region II from about 13-31 at.% Al at 900°C, consists of alloys that contain
enough Al to form initially an external Al2O3 scale, which cannot be sustained because of the Al
depletion, and is overtaken by a rapid growing of NiO + NiAl2O4 + Al2O3 from the interface gasscale to the alloy-scale interface respectively. Region III from Al contents higher than 31 at.% at
900°C, exhibit exclusive growth of an Al2O3 scale. Higher temperatures even extend this region to
lower Al contents. Similar behavior is also observed during the transition from region I to region
II which implies that a critical concentration of Al in the alloy, N*Al , be exceeded, which is a
consequence that the activation energy for Al diffusion is larger than O permeation in the alloy.
41
This behavior showed that N*Al is dependent on kinetic rather than thermodynamic factors.
[108,119]. The crosshatched regions correspond to those where a variable result can be obtained
depending upon surface conditions.
2.2.4.2 Ni-Pt-Al
It is well known that addition of Pt to nickel-aluminide alloys improves its oxidation and hotcorrosion resistance. Because of this, Pt-modified nickel aluminide alloys have been developed and
used as bond coats to protect superalloys components (e.g., blades) used at high-temperature
applications.
Copland [120–122] investigated the effect of Pt addition on the activities of aluminum (aAl ) and
nickel (aNi ) in γ-Ni, γ' -Ni3 Al, and β-NiAl alloys, showing that addition of Pt decrease aAl , and
increase the aNi . in all the cases. The decrease in the aAl have been related with the decreased in
size and number density of the interfacial voids (Kirkendall effect) observed in the traditional
nickel aluminide coatings result of the mismatch between Al and Ni alloy fluxes, whether or not
sulfur is present in the alloy [60,123]. The observed decrease in density of interfacial voids is not
result of a decrease of the amount of Al oxidation; on the contrary, Cadoret et al. [51] showed that
the oxidation rate is accelerated with the increase of Pt. Therefore, the presence of Pt depresses aAl ,
increasing the chemical potential gradient and hence the flux of Al as has been showed by Gleeson
et al. [67].
In 2009, Gleeson et al. [119] published their work entitled “Compositional factors affecting the
establishment and maintenance of Al2O3 on Ni-Pt-Al systems”. Because of this study, the oxidation
map of system Ni-Pt-Al was obtained and showed in Fig. 2.25. It is seen that Pt significantly
decreases the critical Al content for exclusive Al2O3 formation in comparison with Ni-Al system.
It is also observed that the critical Al content for the formation of external Al2O3 was found to be
about 15 at.% for alloys containing approximately 10 at.% Pt. Below this critical Al content, Al2O3
was present either as internal precipitates or as an inner layer of a duplex scale structure. Moreover,
a greater content of Pt did not showed any further beneficial effect in promoting a protective Al2O3
scale.
42
Fig. 25. Oxidation map of the Ni-Pt-Al system exposed to air at 1150°C for 100 h. Adapted from
[124].
2.2.5 Stability of the Alumina Phases
Alumina exists in a number of crystalline forms, where the trigonal α-Al2O3 is the only stable
polymorphic phase at ambient pressure (Fig. 2.26). It has a trigonal structure with space group R3c, which can be described as ABAB stacking O2- planes along normal direction with Al3+ in 2/3
of the octahedral interstitials positions.
Fig. 26. (a) Schematic representation of the first layer α-Al2O3 formation and (b) schematic 3D
representation of α-Al2O3 structure.
43
Moreover, hosts of so-called metaestable alumina phases exist, which often are formed during the
early stages of oxidation before transforming to the stable α phase as is shown in Fig. 2.27.
Fig. 27. Schematic representation of γ-Al2O3 + 3% Pt phase transformation as a function of
temperature. Adapted from [125].
All these transient aluminas exhibit a cubic or pseudo-cubic defect spinel structures as is shown in
table 1.1 [103]. Even when the phase transformations from transient aluminas to the stable α-Al2O3
have been widely studied, however, details about the transformations mechanisms such are not
completely well understood [64,126,127].
Table 1.1. Polymorphs of Al2O3. Adapted from [103].
Polymorph
Crystal system
α
δ
ε
Trigonal
Tetragonal
Hexagonal
γ
Cubic
κ´
Hexagonal
θ
Monoclinic
X
Hexagonal
Unit cell dimensions (nm)
a = 0.4758, c = 1.2991
a = 0.7943, c = 2.350
a = 0.7879 , c = 1.6183
a = 0.7949
a = 0.5544, c = 0.9024
a = 0.5620, b = 0.2906
c = 1.1790, β = 103°20´
a = 0.577, c = 0.864
Furthermore, the transformation to α-Al2O3 involves a decreases of the unit cell volume of about
8% and gives rise to substantial tensile stresses within the TGO [103]. The sequence of
transformation, as well as its kinetics of transformation can be altered significantly in the presence
of additional phases and by dissolved of doped and/or dissolved impurity species such as Pt, Cr,
Zr, Y, Hf, S, etc [13,21,41,46,60,128].
Grabke et al. [21] investigated the oxidation kinetics of NiAl and NiAl-Cr alloys. The oxidation
kinetics showed two phase transformations from γ-Al2O3 → θ-Al2O3 → α-Al2O3. The γ → θ
transformations lead to a small increase in the kp , whereas that the θ → α leads to a strong decrease
of parabolic rate constant, kp , over two orders of magnitude (Fig. 2.28a).
44
Fig. 28. Arrhenius plot of the parabolic rate constants of (a) NiAl and (b) NiAl and NiAl-Cr
oxidation. The oxidation treatments were conducted in a He and O2 flow with an oxygen partial
pressure of 0.13 atm. Adapted from [21].
At 900°C the γ → θ transformation took place after about 10 h, whereas for the θ → α
transformation, a small amount of α-Al2O3 were detected just after 60 h. Moreover, it was also
showed that the higher the chromium content and the more α-Al2O3 nuclei are present the shorter
is also the transformation time from θ → α. For instance, for a NiAl-15Cr alloy oxidized at 900 °C,
only 10 h were necessary to reach a steady state growth of α-Al2O3, compared with the about 300
h required to establish an steady state growth in binary NiAl system. An additional effect of Cr
content was reflected on the increase of kp values as is shown in Fig. 2.28b. The Cr effect on
oxidation kinetics was correlated with the Cr2O3 nuclei formation in the initial oxidation stages,
which serves as a nucleation sites for α-Al2O3.
Cadoret et al. [51] investigated the effect of Pt on the growth rate of the oxide scale formed on
single crystal of NiAl and Ni40Pt10Al50 alloy doped with different sulfur contents (e.g., < 1, 57, 90
ppmw). The results showed that the additions of Pt retards the θ-Al2O3 → α-Al2O3 transformation
at 900° and 100°C and considerably increase the total mass gain at 1100°C during the first stages
of oxidation. This effects was attenuated in the presence of high sulfur content in the alloy, which
suggest that a competitive effect of Pt and S on the segregation of Al is taking place.
45
2.2.6 Oxidation Kinetic Models
Tamman [93] and separately Pilling and Bedworth [94] developed the first systematic studies of
the kinetics of oxidation of metals and alloy during the early twentieth century. These classical
works in conjunction with subsequent experimental and theoretical investigations by other
researchers, have postulated diverse kinetics models to describe different known oxidation
behaviors. These behaviors have been found to be linear, parabolic, linear-parabolic, cubic and
logarithmic [88,125]. In most of the cases, the growth kinetics of the Al2O3 is assumed parabolic
by nature (Eq. 2.49).
∆m = kp t 0.5
2.49
A typical protocol to calculate the kp consists of plotting the net-mass gain per unit of area, Δm,
versus the square root of time, t0.5. When the later part of ∆m vs t0.5 plot becomes a straight line,
the parabolic rate constant, kp , is extracted from slope of the linear part. However, this protocol can
generated errors, particularly in those oxides that exhibit an initial faster-growing stage, which is
the case of the polymorphic transformation of Al2O3 [129]. Besides that, this protocol does not
describe the changes during the initial stage. Based on this, different efforts have been done in
order to developed a more reliable method that allow to determine the real kp for the steady state,
as well as characterized the observed changes during the first stages of oxidation [130–134].
In 1986, Pieraggi [131] established that only a very limited number of metals or binary alloy can
be described by simplest parabolic law represented by the equation 2.49. Additionally, it was
postulated that depending on the specific oxidation behavior two cases can be defined according
with the following relations
∆m2 - ∆m2i = kp (t - ti )
2.50
(∆m - ∆mi )2 = kp (t - ti )
2.51
or
For the first case, equation 2.50, the scale formed during the initial weight gain, ∆mi , after the ti,
does contribute to limiting the kinetics at the steady state for t > ti. The second case, equation 2.51,
corresponds to the case where the initial growth scale (t < ti) is less protective (e.g., NiO or FeO)
46
than the formation of continuous and adherent protective oxide scale (e.g., Cr2O3 or Al2O3) formed
at t > ti. in this case the initial scale does not contribute to the steady state rate control.
In both cases, equations 2.50 and 2.51, represents deviation from the perfect parabolic law. This
means that, if the experimental curve (∆m vs t0.5) was fitted with a power-law model described by
∆m = ktn
2.52
the instantaneous time exponent, n, will be different than 0.5. Furthermore, Pieraggi pointed out
that when a transient period of faster kinetics takes place, the plot ∆m vs t0.5 in inherently superior
to the ∆m2 vs t plot.
Monceau and Pieraggi [132] proposed a methodology to determinate the parabolic rate constants
based on a local analysis of the net-mass gain curves. This method consists on the fitting of massgain data the following parabolic function
t = A + B∆m + C∆m2
2.53
where the coefficients A, B and C are related with the kinetics parameters, for instance, kp = C-1 .
They claim that the translation of the time interval over the entire oxidation time provides an actual
instantaneous parabolic rate constant (Fig. 2.29) independently of any transient stage or
simultaneous reactions steps.
Fig. 29. Illustration of the local-parabolic fitting analysis proposed by Monceau and Pieraggi [132].
47
The completed parabolic model (Eq. 2.53) provides the advantage over the classical parabolic
model (Eq. 2.51) of considering a pure diffusion, diffusion-reaction control and transient regime.
The merit of this method is that it takes into consideration the evolution of growth rate constants
with treatment time. However, the usefulness of the method is still limited. Firstly, this method
restricts the growth kinetics from linear to parabolic behavior. Secondly, the method requires a fitting
of A, B and C coefficients from the experimental data. Due to the nature of the diffusion-controlled
process, the weight gain increase for a unit of time always becomes smaller with increasing time.
Therefore, the coefficients for higher-order terms (e.g. C ) quickly become overwhelmed by noise for
longer oxidation time [134].
Wei Zhao [134] developed a modified method to estimate the steady-state oxidation kinetics based on
a modification of the power lo model described by equation 2.52. The protocol proposed by Wei Zhao
is described below. 1) Setting the starting point of the TG curve; firstly, a zero point must established
to minimize the effect caused by the initial heating stage before to achieve the working temperature. 2)
Calculating the instantaneous time exponent by equation 2.54 (ni); this step is to have an idea of law
dictating the growth kinetics. In other words, to determine if the growth kinetics follows a constant
behavior throughout the testing time, for instance, linear (ni = 1), parabolic (ni = 0.5) or cubic (ni =
0.33).
ni =
∂( log ∆m)
2. 54
∂( log t)
3) Determination of stable growth kinetics. If the ni values are approaching one of the model described
above and these stabilizes to a constant value, the oxidation process has reached an steady state and its
growth rate constant ban be calculated by the average of the instantaneous values during the time that
ni values are approximately constants by
•
To n = 1, the linear instantaneous rate constant, kl, can be defined as kil = ∂t ∂∆m
avg
•
To n = 0.5, the parabolic instantaneous rate constant, kp, can be defined as kip = ∂t •
To n = 0.33, the cubic instantaneous rate constant, kc, can be defined as kic = ∂t ∂∆m2
avg
3
∂∆m
avg
4) Determining of the transient growth kinetics; In the case when ni values are not constants, but they
are approaching to a constant value, for instance 0.5 (assuming a parabolic kinetic), at least two
cases should be considers. Firstly, it is probably that the testing time was not long enough to
48
establish a steady-state regime, or secondly the real kinetics are more complex. In this case, it is
better to determine the additional factors affecting the kinetics behavior.
Wei Zhao claims that this protocol can be used to determine the period for steady state growth
more reliable.
Young et al.[130] proposed a near-cubic oxidation model to describe the growth kinetics and
growth mechanisms of the Al2O3 scales formed on a high purity FeCrAlY-alloy oxidized in Ar-O2
and Ar-H2-H2O atmospheres at 1200° and 1300°C. The model, which is a grain diffusion model
(Eq. 2.54), is modification of the previous model presented by Naumenko et al. [101]
a
3
X3 +r0 X2 = -
4DO,GB δ
RT
∆μO t
2.55
where X is the scale thickness, ∆μO is the oxygen potential gradient across the scale, and a and r0
are constants which defined the linear grain-size distribution as a function of scale thickness,
rG (x) = r0 + ax, encountered by Naumenko et al. [101]. The model allows to calculate the oxygen
grain-boundary diffusion parameters, DO,GB δ, based on rG(x) and the evaluated ∆μO .
49
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63
4. Hypothesis
There is a general agreement in the literature that the poor adhesion between the thermally grown
oxide and the bond coat or top coat (intrinsic problems), as well as the growth rate of the TGO, are
among the most important problems regarding the control of the thermal barrier coating systems
lifetime (intrinsic problems) under engine operating conditions. Moreover, it has been also shown
that a heat treatment of the bond coat to form an aluminum oxide scale (TGO) prior to 7YSZ top
coat deposition increase the lifetime of TBC systems. The reason why this increase occurs is
because pre-oxidation treatments allows to control and tailoring of the thermally grown oxide’s
properties, such as structure, microstructure, grain size morphology, etc.
Based on these observations, the hypothesis of this dissertation is the following:
“The controlled growth of a homogeneous α−Al2Ο3 scale formed by pre-oxidation treatments under
low oxygen partial pressure atmospheres on the surface of β-(Ni,Pt)Al bond coat systems, improve
the physical, chemical, structural and microstructural properties of the thermally barrier coating
systems, and hence, the adhesion of the refractory top coat”.
64
5. Objectives
The general aim of this work is to study the growth kinetics, as well as the structural and
microstructural properties of thermally grown oxides after controlled pre-oxidation treatments of
commercial β-(Ni,Pt)Al bond coats used in thermal barrier coating systems, in order to optimize
the oxidation parameters to grow an exclusively an α-Al2O3 scale prior to 7YSZ top coat
deposition.
Even though many important contributions have been made to the understanding of thermally
grown oxide kinetics on different intermetallic alloys, there is a lack of information about the
oxidation behavior of commercial as-coated β-(Ni,Pt)Al bond coat systems during the first stages
of oxidation (t ≤ 5 h). Based on this, the more specific objectives of the project can be formulated
as follows:
•
To study the oxidation kinetics of commercial as-coated β-(Ni,Pt)Al bond coat systems
during the first stages of oxidation (t ≤ 5 h), at temperatures from 900 to 1200°C, in an
oxidizing atmosphere with an oxygen partial pressure, pO = 1 x 10-5 atm..
2
•
To study the effect of the temperature (from 1100 to 1150°C), grit blasting (with and
•
growth between θ-Al2O3 and α-Al2O3 during the early stages of oxidation (t ≤ 2 h) of
commercial β-(Ni,Pt)Al bond coat systems.
To characterize the structural and microstructural properties of the thermally grown oxides
result of the previous pre-oxidation treatments to the better understanding of the preoxidizing parameters over the θ-Al2O3 → α-Al2O3 transformation.
without) and oxygen-partial pressure ( = 2.1 x 10-1, 5 x 10-5 atm) on the competitive
65
6. First Stages of Oxidation of PtModified Nickel Aluminide Bond Coat
Systems at Low Oxygen Partial Pressure
Abstract
The isothermal oxidation of thermally grown oxide on β-(Ni,Pt)Al bond coats during the θ→αAl2O3 phase transformation was investigated from 900 to 1200°C in an argon atmosphere stream
with the O2 partial pressure of 1 x 10-5 atm. Local parabolic fitting was used to evaluate the
evolution of kp, using a general parabolic model (t = A+BΔm+CΔm2), during the first 5 h of
oxidation. All net mass-gain curves exhibited deviations from the classic parabolic model,
Δm=kpt0.5,a steady-state regime was established only after 4 h of treatment, except for the sample
oxidized at 1100°C.
66
6.1
Introduction
Improvements in gas turbine efficiency demand an increase in operating turbine entry temperatures
(TET) [1]. The current performance of gas turbines in the aeronautic industry is the result of
continuous improvements in different areas of engineering, including turbine design, control
systems, combustion systems, and materials science. Within materials science the most significant
developments have been made in the areas of alloy design, casting technology, and coating methods
[2]. High temperature multilayer coatings known as thermal barrier coating (TBC) systems have
been used since the early 1980s to protect the parts exposed to the highest temperatures (e.g.,
combustion chamber, rotor blades and stator vanes) [3].
TBC systems are usually designed with four primary constituents: 1) a Ni-base superalloy (SA) as
the substrate, providing resistance to static load, creep, and fatigue, 2) an Al-rich bond coat (BC)
designed as an Al-reservoir to allow formation of a homogenous oxide scale, 3) a thermally grown
oxide (TGO), the result of oxidation of the BC, and 4) a ceramic top coat (TC) as a thermal barrier,
typically Y2O3-doped ZrO2 in a composition leading to a “non-transformable” tetragonal phase [4].
The TGO has two principal functions: it binds the ceramic layer to the bond coat, and it acts as
diffusion barrier to prevent substrate oxidation [4,5]. Notwithstanding that TBC systems have
allowed much higher TETs than were previously possible, failures due to a mismatch between
thermal expansion coefficients of TBC system layers during operation have limited their lifetime.
The failure mechanisms of TBC systems have been classified as extrinsic (e.g., impact damage and
molten deposits) and intrinsic (e.g., rumpling, edge-delamination and void formation), as defined
by Evans et al. [6] Most of the intrinsic TBC system failures have involved the BC/TGO/TC
interfaces, because of a group of mechanisms associated with strain misfit between layers. These
mechanisms are manifested as TC delamination. Because of this, efforts have been made to
improve the oxidation lifetimes of the TBC systems by doping the BC with reactive elements [714], using different deposition techniques [3,4,15] or developing oxidation treatments prior to TC
deposition in order to control the TGO properties [16-21].
The control parameters that determine the oxide growth properties are: 1) BC and TC deposition
methods, 2) SA, BC, and TC composition, 3) BC surface condition (e.g., grit blasting) before the
oxidation treatment, 4) oxygen partial pressure, pO , 5) temperature, and 6) heat treatment
2
procedure. Regardless of the BC used (e.g., NiPtAl, MCrAlY, etc.), the TGO must comply with
67
the following characteristics to extend the TBC systems’ lifetime: 1) promote growth of α-Al2O3
with the largest possible grain size, 2) present a uniform columnar morphology along the BC
surface, 3) have good adhesion with BC and TC, and 4) have slow growth kinetics [18-21].
In 1992 Brumm and Grabke [22] showed that oxidation of the intermetallic B2-NiAl exhibits two
polymorphic transitions (γ-Al2O3→θ-Al2O3→α-Al2O3) before reaching the stable α-Al2O3 phase.
Transformation to the stable α-Al2O3 is accompanied by a volume decrease of the lattice unit cell
[16]. For instance, the volume change from θ-Al2O3 to α-Al2O3 (∼8% reduction [22]) can lead to
defects and thermal stresses in the BC/TGO interface that may induce failure of the TBC systems.
A transient oxidation stage is generally observed during these phase transitions before reaching a
steady state. The γ- to θ-Al2O3 phase transformation leads to a small increase in the parabolic rate
constant, kp, while the θ- to α-Al2O3 phase transformation results in a large decrease of kp (almost
two orders of magnitude) [23].
It is well known that the growth rate of protective oxide scales is commonly controlled by solidstate diffusion through the oxide scale itself, and it is very well described by the parabolic rate law
∆m2 = kp t
6.1
where kp is the parabolic rate constant and Δm is the mass-gain per unit area at time t.
This behavior was first observed experimentally by Tammann [24] and, independently, by Pilling
and Bedworth [25] and subsequently was treated theoretically by Wagner [26]. Nevertheless,
important deviations from Tammann´s model have been observed in the growth rates of important
oxides like Cr2O3 and Al2O3. Such behavior may arise from different causes, including impurities,
grain-boundary diffusion, multilayer scale growth and polymorphism of the oxide (e.g., θ-Al2O3,
α-Al2O3, etc.). Different models have been proposed in the literature to consider these deviations
based on mixed kinetic behaviors [27-29].
In 1998 Monceau and Pieraggi proposed a methodology to evaluate parabolic rate constants from
a local analysis of mass-gain curves for intermetallic materials [27]. They claimed that the use of
the classic Δm2 =kpt model should be replaced by a more general model that considers mixed
(diffusion/reaction) control and a transient regime, as well as pure diffusion control
t=A+B∆m+C∆m2
6.2
68
This method is based on local parabolic fitting of (Δm,t) data over short time intervals to obtain
instantaneous values of the parabolic rate constant. This method in general has the advantage over
the usual
d∆m
k'p =2∆m dt
6.3
because it considers the non-idealities associated with oxidation tests. Table 6.1 summarizes two
cases derived from Eq. 6.2 and shows how the coefficients A, B, and C are related with kinetic
parameters. Full details of the model have been reported elsewhere [28].
Table 6.1. Summary of rate equations for the different parabolic models proposed by D. Monceau
and B. Pieraggi (1998)+.
Case
Rate Equation*
1
kp
d∆m
=
2∆m
dt
2
Kinetics law
A
B
C
∆m2
kp
0
0
1
kp
d∆m
1
∆m2 -∆m2i ∆m-∆mi
∆mi ∆m2i 1
=
1 2∆m t-ti =
+
ti dt
k + k kp kl
kp
kl
kl
l
p
1
kp
t=
*The initial conditions used to integrate the rate equations were t=ti and Δm=Δmi. +Adapted from
Ref. [27].
A power law model is also used to describe the para-linear and pseudo-cubic kinetics of Al2O3
growth, which have been respectively related with phase boundary reaction and grain boundary
diffusion (also inversely proportional to the grain size) along the scale [27,29]:
∆m=ktn
6.4
where k and n are the power law rate constant and rate exponent, respectively.
Important contributions to the understanding of the TGO kinetics for different intermetallics have
been reported in the last two decades, some of which suggest that oxidation treatments prior to TC
deposition extend the lifetime of TBC systems. However, there is a lack of information about the
kinetic behavior of (Ni,Pt)Al bond coats during the first stages of oxidation [4,5,23]. The present
study aims to contribute to the understanding of the oxidation kinetics of commercial (Ni,Pt)Al
69
bond coats, focusing on the first hours of treatment (t ≤ 5 h) as a function of temperature in an
oxidizing atmosphere with a pO = 1 x 10-5 atm.
2
6.2
Experimental Procedure
6.2.1
Sample Preparation
All samples used in this work were provided by GE Aircraft Engines (Evendale, OH) as rectangular
specimens (1.8 x 1.2 x 0.15 cm), each weighing about 2900 mg. The specimens consisted of a ∼70
μm thick bond coat deposited onto Rene N5 single-crystal Ni-based superalloy coupons. Bond
coats were produced by electroplating a thin layer of Pt onto a grit-blasted substrate. A heat
treatment was subsequently performed to diffuse Pt into the substrate, followed by a vapor phase
aluminide process (VPA) to introduce Al into the coating. After a second heat treatment the desired
β-(Ni,Pt)Al phase was achieved. The resulting bond coat is a bilayer structure consisting of a ∼50
μm thick β-(Ni,Pt)Al and a ∼20 μm thick inter-diffusion zone (IDZ). The substrate and bond coat
nominal compositions were determined using inductively coupled plasma atomic emission
spectroscopy (ICP-AES, Perkin Elmer Optima DV4300) and electron micro probe analysis
(EPMA, JEOL 8900 WD/ED microanalyzer) respectively, and are listed in Table 6.2. Samples
were ultrasonically cleaned using xylene, acetone, methanol-water (1:1), and deionized water for
15 min each to remove surface contamination prior to oxidation treatments.
Table 6.2. Chemical composition of the single-crystal superalloy Rene N5 and the β-(Ni,Pt)Al
bond coat.
Sample
Weight percent values (wt. %)
Cr
Co
Mo
Re
W
Al
Ti
Ta
Hf
Pt
Ni
Nominal
ReneN5*
7.0
8.0
2.0
3.0
5
6.2
-
7.0
0.2
-
Bal.
As-received
ReneN5**
6.19 8.27
1.39
3.23 5.03
6.37
0.01 6.93 0.15
-
Bal.
As- coated β(Ni,Pt)Al BC***
0.80 2.66 0.045
-
0.01 24.91
-
0.28
-
31.86 Bal.
* Ref. [3]; ** Measured by ICP-AES, *** Measured by EPMA.
70
6.2.2
Oxidation Treatments
Thermogravimetric analyses (TGA) were performed to follow the oxidation kinetics of β-(Ni,Pt)Al
BC systems at isothermal temperatures in the range 900–1200°C, for 5 h, under a purified argon
stream (working gas) with a pO = 1 x 10-5 atm. To reach the desired oxygen concentration, pure
2
-4
argon with 1 x 10 atm of O2 was passed through a gas purification system (Oxygon Industries
OG-120M). The thermogravimetric analyses were carried out using a thermobalance (Setaram
Setsys Evolution 16/18), which has an accuracy of ±0.03 μg. The mass change during TGA
experiments was recorded at 5 s intervals. The sample was hung from one end of the beam balance
using a 0.4 mm diameter Pt/30%Rh wire and placed in the hot zone of the vertical furnace. The
reactor furnace was made of dense alumina with an 18-mm inner diameter. The temperature of the
furnace was controlled by a Pt-Pt/13% Rh (S-type) thermocouple placed just below the sample. To
minimize any perturbation resulting from gas flow, buoyancy, drag forces or premature sample
oxidation before reaching the working temperature, the heating cycle was programmed as follows:
1) evacuation of the analysis chamber to less than 10 Pa, 2) heating to the reaction temperature at
a maximum rate of 50°C/min, 3) back filling the chamber with working gas to atmospheric pressure
using a flow rate of 200 ml/min, 4) lowering the flow of the working gas to 20 ml/min and keeping
the same at the working temperature for 5 h, and finally 5) cooling the furnace chamber to room
temperature at a maximum rate of 50°C/min with a gas flow rate of 0.3 ml/min.
6.2.3
Characterization
The microstructural evolution of the TGO was characterized by electron microscopy (field
emission gun scanning electron microscopy, FEG-SEM, Jeol JSM7401F, and field emission gun
scanning
transmission
electron
microscopy FEG-STEM, Jeol
JSM7401F). Structural
characterization by photo-stimulated luminescence spectroscopy (PSLS) was also done [30-32]. It
is desirable that the formed TGO is primarily pure α-Al2O3; however, in practice, Cr3+ ions tend to
be incorporated in the crystal structure of Al2O3. PSLS consists of irradiating Cr-doped Al2O3 with
a green (532 nm) or blue (473 nm) laser. Once illuminated, the Cr3+ ions emit a fluorescent radiation
due to radioactive decay of the excited electrons to the ground state [33].
The PSLS spectrums of α- and θ-Al2O3 consist of two major bands, with well-characterized
frequencies, the R1-R2 and T1-T2 doublets, respectively. The identification of the alumina phase
71
is done on the basis of differences in frequency of caused by different Cr3+-O2 bond distances from
the α and θ structures. PSLS measurements were performed in a micro-Raman mapping
spectrometer (Renishaw InVia) connected to a Leica microscope equipped with a 532 nm linefocus laser. The microscope stage was fitted with an automated xyz-motorized stage. Wavelength
and laser intensity were internally calibrated using a silicon standard using the peak at 520 cm-1.
PSLS spectra were collected in a CCD array using an 1800 l/mm grating. A 20 X microscope
objective was used to focus the ∼1 μm spot-sized laser beam and to collect the scattered light. The
laser power at the sample was 5 mW, and the acquisition time for each spectrum was in the 0.5-2.0
s range. PSLS spectra were obtained by mapping the sample surface over an ∼1250 μm2 area with
a pitch size of 1 μm.
6.3
Results and Discussion
6.3.1 Thermogravimetric Analysis
Figure 6.1 shows representative micrographs of the microstructure of the surface and cross section
of the as-received samples. As mentioned previously, those samples were TGA-analyzed at
isothermal temperatures from 900 to 1200°C under a purified argon stream (working gas) with a
pO = 1 x 10-5 atm.
2
Fig. 6.1. Microstructure of the as-received samples: (a) surface and (b) cross section.
Figure 6.2 shows the net mass-gain curves of β-(Ni,Pt)Al bond coats systems oxidized for 5h. The
mass-gain rates increase with the temperature from 900 to 1100°C, whereas higher temperatures
72
(1150 and 1200°C) led to decrease of the mass-gain rates. This behavior has been observed for
many Al2O3 forming alloys, and it is associated with the fact that stable α-Al2O3 is controlling the
oxidation kinetics at high temperatures [23]. It is also well known that this phase transformation is
highly dependent on time, temperature, primary chemical composition (e.g., Al, Pt, Cr), minor
additions of other elements (e.g., Y, Zr, Si, Hf) or contaminants (e.g., S), heating procedure (e.g.,
heating rate), surface preparation (e.g., polishing conditions, grit blasting process, roughness), and
pO [34-38].
2
TG Net Mass-Gain (mg/cm 2)
0.20
900°C
950°C
1000°C
1050°C
1100°C
1150°C
1200°C
0.16
0.12
0.08
0.04
0.00
0
1
2
3
4
5
t (h)
Fig. 6.2. Experimental net mass-gain curves for β-(Ni,Pt)Al bond coats during isothermal oxidation
in a purified argon stream with a pO = 1 x 10-5 atm.
2
Figure 6.3 shows the net mass-gain curves using a Δm vs. t0.5 plot, which is useful to illustrate
deviations from the straight line expected from the classic parabolic model (Eq. 6.1). Based on
these plots, one can conclude that all net mass-gain curves exhibit slight deviations from the ideal
behavior. This can be explained by inferring that initially a mixed (diffusion/reaction) control stage;
predominated followed by a transient oxidation stage that preceded the establishment of steady
state. In the present case, the transient oxidation stage is expected during the polymorphic transition
of the alumina (e.g., γ-Al2O3→θ-Al2O3→α-Al2O3). This behavior is discussed in detail in the
following sections.
73
0.20
900°C
950°C
1000°C
1050°C
1100°C
1150°C
1200°C
2
TG Net Mass-Gain ( mg/cm )
0.18
0.16
0.14
0.12
0.10
0.08
0.06
0.04
0.02
0.00
0
20
40
60
80
t 1/2 (s1/2)
100
120
140
Fig. 6.3. Δm vs. t0.5 plots showing the deviations from the classic parabolic model for the net massgain curves of β-(Ni,Pt)Al bond coats during isothermal oxidation.
6.3.2 Microstructural and Structural Analysis
The TGO morphology after oxidation treatments was studied using FEG-SEM over the surface of
the oxidized samples. Figure 6.4 shows clear changes in TGO microstructure as a result of different
mechanisms controlling the growth of alumina phases. Samples treated from 900 to 1050°C (Fig.
6.4a-d) exhibited grains in whisker and platelet morphology, which is typical of the Al2O3
metastable phases. The reason of this morphology is that γ- and θ-Al2O3 phases grow by outward
cation diffusion (p-type oxide) [39-41]. Samples treated at 1100 and 1150°C (Fig. 6.4e) showed
clear evidence of the first α-Al2O3 grains in combination with the microstructure of metastable
Al2O3 phases. Finally, the sample treated at 1200°C showed a homogeneous microstructure, which
is characteristic of the pure α-Al2O3 phase following inherent BC surface roughness (Fig. 6.1a). αAl2O3 growth is predominantly controlled by inward anion diffusion (n-type oxide) and outward
cation diffusion (p-type oxide) mainly along grain boundaries. This type of growth mechanism is
manifested by characteristic ridges in the TGO/gas and BC/TGO interfaces [42]. However, well
developed ridge structures were not observed in these samples.
74
Fig. 6.4. Surface microstructure of samples treated for 5 h at: (a) 900°C, (b) 950°C, (c) 1000°C,
(d) 1050°C, (e) 1100°C and (f) 1150°C.
The explanation of this effect is associated with the low pO used in these experiments (1 x 10-5
2
atm), as has been recently shown by Wada et al. [43]. The cracks observed in the sample treated at
75
950°C (Fig. 6.4b) are probably associated with the θ-Al2O3→α-Al2O3 transformation and will be
discussed below. The sample treated at 1000°C shows imprints resulting from the inward growth
of α-Al2O3 grains, as well as faceted voids at locally spalled areas (Fig. 6.4c). The formation of
these interfacial voids occurs during the first minutes of oxidation, as confirmed by the works of
Zimmerman et al. [44] and Smialek [45] who report the presence of faceted voids after 6 and 4.5
min, respectively, in NiAl alloys. The presence of such voids has been attributed to different causes.
For instance, the Kirkendall effects promoted by the unbalanced fluxes of aluminum (outward
diffusion) and Ni/Pt (inward diffusion) can result in interfacial void formation [46]. Vacancy
injection and evaporation of the Al from the void surface can act as void growth mechanisms [47].
The presence of third elements (e.g., Pt, Cr), reactive elements (e.g., Hf, Y, Zr) and minor amounts
of alloy impurities (e.g., S) can also affect void nucleation and growth [48]. It has been shown that
a grit-blasting process after the aluminizing process in β-(Ni,Pt)Al alloys mitigate the interfacial
void formation in service conditions [49].
Fig. 6.5. STEM micrographs of the cross section view showing the α-Al2O3 grain morphology for
the sample treated at 1200°C for 5 h.
Figure 6.5 shows a cross-sectional view of a sample oxidized at 1200°C, which was prepared using
a focused ion beam technique (FIB, Jeol JEM-9320) and observed using FEG-STEM. A uniform
scale of α-Al2O3 grains is visible along the surface with a vertical length of 0.62±0.2 µm and a
lateral length of 1.4±0.15 µm.
The α-Al2O3 content in the TGO was determined by PSLS based on the following equation
76
Cα =
(AR1 +AR2 )
(AR1 +AR2 +ϕ(AT1 +AT2 ))
6.5
where AR1, AR2, AT1 and AT2 are the integrated intensities corresponding to the characteristic R1-R2
and T1-T2 doublets for the α- and θ-Al2O3, respectively, and ϕ is a correction factor that considers
the weaker luminescence signal for θ-Al2O3. Tolpygo and Clarke [50] reported that θ-Al2O3
luminescence is 10-12 times weaker than α-Al2O3. Based on this, a constant value of ϕ = 12 was
used in this work. Fitting of the spectra was performed using Renishaw’s WiRE v3.2 software to
unambiguously determine the contributing intensity and thus the position of the R1-R2 and T1-T2
peaks.
Figure 6.6 shows the evolution of the TGO phases as obtained from PSLS analysis. A sigmoidal
behavior as a function of the treatment temperature is observed, reaching a plateau at 1150°C,
where θ-Al2O3 was no longer detected in the scanned area. Thus, one can conclude that for 900 and
950°C the TGO growth kinetics are mainly controlled by the growth of θ-Al2O3, though the growth
of γ-Al2O3 cannot be ruled out, especially during heating and very early oxidation stages. However,
γ-Al2O3 was not observed using PSLS after 5 h of treatment at any temperature. From 1000 to
1100°C the oxidation kinetics exhibit θ-Al2O3 as the initial phase, followed by a θ-Al2O3→α-Al2O3
transition stage where both phases coexist. This claim is clearly confirmed at 1100°C based on the
microstructural analysis, where both θ- and α-Al2O3 were observed (Fig. 4e).
Figure 6.6 also compares the fraction of α-Al2O3 obtained in the present work at 1100°C in a
purified argon stream with a pO = 1 x 10-5 atm with the fraction of α-Al2O3 grown on a diamond
2
polished (001) face of a β-NiAl(Zr) sample oxidized in air at 1100°C, as reported by Veal et al.
[51]. A higher transformation percent (from θ-Al2O3→α-Al2O3) is observed in the β-(Ni,Al)Pt BC
system, 68.4 %, in comparison with the 34.3% observed in the β-NiAl(Zr) after 5 h of oxidation
treatment. A comparable transformation rate was reached only after 7.68 h of treatment. These
differences are mainly associated with chemical composition because it has been reported that
larger ions such as Zr (Y, La and Hf) slow down the θ-Al2O3→α-Al2O3 transformation [52, 53],
whereas the presence of Cr (and Fe) accelerate the transformation [54, 55].
77
(Ni,Pt)Al BC after 5h
Veal et al. after 5 h
Veal et al. after 7.68 h
1.0
0.8
R1
0.6
R2
Counts (a.u)
Frac. α -Al2 O3
α - Al2 O3
T2
θ - Al2O 3
T1
0.4
0.2
14300 14350 14400 14450 14500 14550 14600 14650
-1
wavelength (cm )
0.0
900
950
1000
1050
1100
Temperature (°C)
1150
1200
Fig. 6.6. Evolution of the α-Al2O3 phase transformation as a function of temperature after 5 h of
oxidation treatment of (Ni,Pt)Al bond coats. These results were obtained by image analysis of
PSLS mappings. A typical PSLS spectrum for a TGO composed of both α- and θ-Al2O3 phases is
also shown.
The growth and transformation of an alumina scale at high temperatures on Al2O3-forming alloys
(e.g., NiAl and NiPtAl) can be divided into three regions of transformation, as has been proposed
by Veal, et al. [51]: 1) In the early stages, the alumina scale is controlled by the growth of
metastable phases (γ- and/or θ-Al2O3) via outward Al diffusion [39]. After an incubation period,
the stable α-Al2O3 nucleates at the metal-oxide (M/O) interface, and its grains grow laterally more
rapidly than they grow vertically [51]. 2) A uniform layer of α-Al2O3 is established along the M/O
interface blocking outward transport of Al and stifling the growth of θ-Al2O3. The growth of new
oxide is now mainly controlled by the slow growth kinetics of α-Al2O3. This stage is characterized
by θ-Al2O3→α-Al2O3 transformation along the scale-alloy interface. 3) Once the transformation
stage has been completed, α-Al2O3 is the only phase present in the TGO, and it controls the scale
growth kinetics. The duration of the different stages is highly dependent on the chemical
composition, surface conditions, temperature, pO , etc. [23, 50, 51].
2
78
Fig. 6.7. Surface morphology of the treated samples, showing the spallation regions observed after
cooling: (a) 950°C, (b) 1000°C, (c) 1050°C, (d) 1100°C, (e) 1150°C (no spallation) and (f) 1200°C
(no spallation).
It is worthwhile to mention that the θ-Al2O3→α-Al2O3 transformation involves an approximately
8% volume reduction [22]. Therefore, radial tensile cracks can develop in the scale during
transformation (Fig. 4b). Both the stress state to which the alumina scale is subjected during the
transformation process and the thermal stresses developed during cooling and formation of
interfacial voids can lead to cracking and spallation of the TGO scale.
79
Figure 6.7 shows the surface of the oxidized samples after cooling, where spallation regions can
be clearly identified only for the samples treated from 950 to 1100°C. The sample treated at 900°C
presented no evidence of spallation. Based on this result it can be concluded that only the samples
in the second region of transformation exhibited TGO scale spallation after cooling.
6.3.3 Kinetics Analysis
In order to study the parabolic rate constant, kp, during transient and steady-state regimes, the net
mass-gain curves where analyzed using the local fitting procedure to a general parabolic law (Eq.
2), as proposed by Monceau and Pieraggi [27].
Figure 6.8 shows the time dependence of kp for the oxidized samples. The steady-state regimes are
shown in Fig. 6.8b. According to Fig. 6.8b a steady-state regime was reached for all samples after
4 h of treatment, except for the sample treated at 1100°C, in which kp still shows a decreasing
tendency after 4.5 h.
A similar result was reported by Cadoret et al. [56] for low sulfur (S<1 ppmw) Ni40Pt10Al50 alloys
oxidized in air at 1100°C for 5 h, where the effect of Pt-additions on the growth of Al2O3 phases
was discussed. In particular, Pt favors the growth of θ-Al2O3 after nucleation of α-Al2O3 and
consequently delays its transformation. It was proposed that this effect helps to relax the stresses
during the θ-Al2O3 to α-Al2O3 transformation. This effect was attenuated in the presence of high
sulfur contents (S=90 ppmw), and it was suggested that sulfur segregation to the BC/TGO interface
limits Al segregation. However, it was reported that the transient stage never exceeded 1 or 2 h;
whereas, in the present work, this stage was extended up to 4 h. These variations can be associated
with differences in the surface preparation (as-VPA process vs. surface polished to 5 µm finish)
and chemical composition (e.g., sulfur content).
80
100
900°C
1000°C
1100°C
950°C
1050°C
1150°C
kp x 10-7 (mg2 /cm4s )
1200°C
10
1
0.1
0
1
2
3
4
5
t (h)
10
kp x 10-7 (mg 2/cm 4s)
1050°C
1100°C
950°C
1
1000°C
1200°C
900°C
1150°C
0.1
3.4
3.6
3.8
4.0
4.2
t (h)
4.4
4.6
4.8
5.0
Fig. 6.8. Log kp plotted as a function of time for β-(Ni,Pt)Al bond coats during isothermal oxidation
in a purified argon stream with a pO = 1 x 10-5 atm: (a) transient and steady-state regimes and (b)
2
steady-state regime.
81
Temperature (°C)
1200 1150 1100 1050
kp x 10-7 (mg2/cm4s)
100
950
900
θ-Al2O3
α-Al2O3
10
1000
Transient Regime in this work
1
0.1
0.01
Brumm & Grabke (1992)
Measural Kinetics
After 5 min
After 5 h
γ-Al2O3
6.6 6.8 7.0 7.2 7.4 7.6 7.8 8.0 8.2 8.4 8.6 8.8
4
Reciprocal Temperature (10 /K)
Fig. 6.9. Arrhenius plot summarizing the kp values of β-(Ni,Pt)Al bond coats during isothermal
oxidation with a pO = 1 x 10-5 atm. Bold lines correspond to the γ-Al2O3, θ-Al2O3 and α-Al2O3
2
lines refer Brumm and Grabke [23].
Figure 6.9 shows an Arrhenius plot of the kp estimated in this work along with the Brumm and
Grabke data for a NiAl alloy [23]. A good agreement is observed between the growth kinetics
during the first 5 min and the slope obtained by Brumm and Grabke for θ-Al2O3. Based on this, it
is confirmed that the metastable θ-Al2O3 controls the growth kinetics during early stages (red dots)
from 900 to 1100°C, whereas that at 1150 and 1200°C α-Al2O3 is the dominant phase. The striped
zone represents the transition regime before reaching the steady-state regime after 4 h (blue dots).
Finally, a comparison between our kp values and those reported by Cadoret et al. [56] for single
crystals (oriented along [001] direction) of Ni50Al50 and Ni40Pt10Al50 alloys (oxidized under 1 atm
of pure oxygen) doped with different concentration of sulfur is shown in Fig. 6.10.
82
Temperature (°C)
1200 1150 1100 1050
k p x 10-7 ( mg2/cm4s )
100
1000
900
θ-Al2O3
Transient state
Steady state
10
1
950
α-Al2O3
Steady state
0.1
0.01
1E-3
1E-4
γ-Al2O3
Brumm & Grabke (1992)
After 5 min (Transient State)
After 4 h (Steady State)
Ni50Al50 (S<1ppmw)
Ni50Al50 (S=57 ppmw)
Ni40Pt 10Al50(S<1ppmw)
Ni40Pt 10Al50(S=90 ppmw)
6.6 6.8 7.0 7.2 7.4 7.6 7.8 8.0 8.2 8.4 8.6 8.8
4
Reciprocal Temperature (10 /K)
Fig. 6.10. Arrhenius plot showing a comparison between our kp values for β-(Ni,Pt)Al bond coats
during isothermal oxidation with a pO = 1 x 10-5 atm and literature data [56]. Bold lines correspond
2
to the γ-Al2O3, θ-Al2O3, and α-Al2O3 lines refer Brumm and Grabke [23].
Based on this, it is observed that at 900°C the kθp is about 3 times lower that the reported by Cadoret
et al. [56] for NiPtAl alloys, while at 1100°C is only 1.5 times lower for NiPtAl alloys. This
difference is directly associated with the different pO conditions used during the oxidation
2
treatments. The decrease observed at 1100°C is also associated with a pretreatment performed on
the samples treated at 1100°C to grow a 500 nm thick oxide scale at 900°C.
83
6.4
Conclusions
All net mass-gain curves exhibited deviations from the classic parabolic model proposed by
Tammann because a transient oxidation regime (including θ-Al2O3 phase growth, α-Al2O3 grain
nucleation, θ- to α-Al2O3 phase transformation, and grain coarsening) controls the initial oxidations
kinetics at all temperatures tested (900 to 1200°C). The growth kinetics behavior of the thermally
grown oxide on the (Ni,Pt)Al bond coats can be well described by the general parabolic law t =
A+BΔm+CΔm2 using a local fitting procedure. A steady-state regime was reached after 4 h of
treatment (except for the sample treated at 1100°C), unlike the results reported in the literature [56]
for similar systems where a steady-state regime was reached in less than 2 h. This difference is
primarily associated with the surface preparation, because in the present study the samples were
oxidized after the VPA process where the characteristic ridges of the BC grains are exposed (Fig.
1), while in the reported data the samples were previously polished removing the BC ridges and
decreasing the species diffusion.
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87
7.
Kinetics Study of the Competitive Growth between
θ-Al2O3 and α-Al2O3 during the Early Stages of
Oxidation of Pt-Modified Nickel Aluminide Bond Coat
Systems: Effects of Oxygen Partial Pressure, Surface
Treatment, and Temperature
Abstract
An oxidation study of β-(Ni,Pt)Al bond coat systems was carried by means of TGA analysis during
isothermal treatments at temperatures from 1000 to 1150°C and subsequent analyses.
and TGA analysis. The effects of oxygen partial pressures and/or grit blasting on their kinetics of
oxidation were studied and complemented by photo-stimulated luminescence spectroscopy. It is
shown that the combined effect of low oxygen partial pressure and grit-blasting process stifle the
growth of metastable θ-Al2O3 and accelerate the θ-Al2O3→α-Al2O3 transformation, therefore
reducing the time required to reach a steady-state regime where α-Al2O3 is the controlling phase.
For instance, for the grit-blasted sample oxidized at 1000°C and at a pO = 1 x 10-5 atm, a steady2
state regime was established after about 1 h of oxidation whereas that for the as-aluminized (not
grit-blasted) sample oxidized in a pO = 2.1 x 10-1 atm at the same temperature, θ-Al2O3 is still
2
controlling the growth of the alumina scale after 5 h of oxidation.
88
7.1 Introduction
High-temperature coatings have been widely used since 1950s to protect the hot-sections in aeroengines and stationary gas turbines against oxidation and corrosion phenomena, such as thermal
fatigue of the substrate [1-3]. These have been classified as diffusion coatings, overlay coatings
and thermal barrier coatings [4]. Diffusion coatings have been studied extensively recent decades,
with nickel aluminides being one of the most widely used in the aeronautical industry [4-7]. In
order to improve its oxidation resistance (e.g., adherence of the alumina scale and low growth
kinetics) and hot corrosion, diverse elements have been added during its processing including Pt,
Pd, Rd, Ir, Cr, Hf, Si and Y [8-11]. The Pt-modified nickel aluminide bond coat systems, used in
the present work are diffusion coatings and are composed of three different constituents: 1) a Nibased superalloys as the substrate that provides mechanical in the form of tensile, creep and fatigue
strength during operation, 2) an alumina-forming β-(Ni,Pt)Al alloy as the bond coat (BC) that acts
as a diffusion barrier to prevent substrate oxidation, as well as to increase hot corrosion resistance
and 3) a thermally grown oxide (TGO) resulting from BC oxidation [3,12]. The TGO has two
principal functions. It acts as an additional diffusion barrier to prevent substrate oxidation, and it
binds a ceramic top coat (TC) layer, typically 7 wt% yttrium-stabilized zirconia (7YSZ) to the BC,
forming a multifunctional high-temperature coating known as a thermal barrier coating (TBC)
system [5,13]. The intrinsic failure mechanisms of TBC systems as defined by Evans et al. [12] are
always associated with the detachment along the BC/TGO/TC interfaces, as consequence of
interfacial defects, the growth stress in the TGO and the thermal expansion mismatch between the
TGO and BC, which results in large compressive stresses (3-5 GPa) in the TGO layer during
cooling [14-16]. In order to extend the lifetime of TBC systems, the TGO must: 1) be formed by
α-Al2O3 with the largest possible grain size, 2) present a uniform columnar morphology along the
BC surface, 3) have good adhesion with the BC and TC, and 4) have slow growth kinetics [17-19].
In 1992 Brumm and Grabke [20] showed that oxidation of the intermetallic β-NiAl exhibits two
polymorphic transitions (γ-Al2O3→θ-Al2O3→α-Al2O3) before reaching the stable α-Al2O3 phase.
Transformation to the stable α-Al2O3 is accompanied by a volume decrease of the lattice unit cell.
For instance, the volume change from θ-Al2O3 to α-Al2O3 (∼8% reduction [21]) can lead to defects
and thermal stresses in the BC/TGO interface that may induce failure of the TBC systems. It is
important to emphasize that the fabrication process of the TBC systems with β-(Ni,Pt)Al bond coats
89
includes diverse stages that can influence its oxidation kinetics during service, and therefore the
TBC system’s lifetime. These processes include electrolytic Pt plating [22] (e.g., sulfur impurities),
heat treatments and the aluminizing process [5,9] (e.g., high-activity and low-activity processes),
surface preparation (e.g., grit-blasting process after the aluminizing)5,9 and pre-oxidation
treatments conditions [17-18, 23-27] (e.g., oxygen partial pressure, temperature, heating rate, time)
previous to TC deposition. In the present work results of the effect of grit-blasting process and
oxygen partial pressure (pO ) conditions on the oxidation behavior of β-(Ni,Pt)Al bond coats during
2
isothermal (pre-oxidation) treatments will be discussed.
Tolpygo et al. [28] reported on the effect of grit blasting on the cyclic oxidation behavior of a β(Ni,Pt)Al coating at 1150°C, showing that the grit-blasted samples exhibited about a tenfold
increase in oxidation rate and extensive scale spallation. It was also reported that grit blasting
introduces impurities (Li, Na, K, Mg, Ca and Ti) into the bond coat surface which Tolpygo suggests
reduce the grain size of the alumina scale as compared with the as-aluminized samples. Decreasing
the grain size increases the diffusivity of oxygen and aluminum in the TGO, increasing the oxide’s
growth rate.
Haynes [29] showed that numerous elongated voids (with diameters from 0.5 to 1.9 µm and an
aspect ratio of about 3) were beneath the alumina scales that formed on as-aluminized β-(Ni,Pt)Al
BC systems produced by low-activity chemical vapor deposition after EB-PVD 7YSZ TC
deposition. The samples were grit-blasted samples prior to TC deposition had very few voids at the
oxide-metal interface. Haynes suggests that the reduction in void formation on the grit-blasted
surfaces is strongly related to with the accelerated growth of α-Al2O3 or the removal of excess
surface impurities, as void formation is highly dependent on the structure and local composition,
including Al and Pt content, S impurities and reactive element additions [30]. Meanwhile, Xie et
al. [31] reported the effect of grit blasting β-(Ni,Pt)Al BC on the durability of EB-PVD TBC
systems showing that non-grit-blasted samples exhibit a lifetime ~1.4 times higher than grit-blasted
samples. However the latter present a narrower scatter in the durability results, making TBC
systems with grit-blasting treatment more consistent in engineering practice. The idea that preoxidation treatments of β-(Ni,Pt)Al bond coats increase TBC lifetime as a result of improving TGO
properties and reducing the oxide growth rate have been proposed previously [17,18,23,27]. Preoxidation is defined as a heat treatment of the bond coat to form the aluminum oxide prior to TC
deposition [17].
90
Tolpygo and Clarke [18] demonstrated, by way of thermal cycling tests at 1150°C, that preoxidation treatments in air of β-(Ni,Pt)Al BC systems to form a thin layer of α-Al2O3 TGO prior to
TC deposition improves the lifetime of the TBC systems by a factor of two to three. At around the
same time, Spitsberg and More [17] reported on the effect of TGO microstructure on the durability
of TBC systems with β-(Ni,Pt)Al BCs. They found that BC pre-oxidation treatment, when done
under specific oxygen partial pressures conditions, can result in greater than two times
improvement in TBC system lifetime as compared with non-treated BCs. It was concluded that
pre-oxidation treatments at low pO reduce the TGO growth rate and therefore its the growth stress
2
in the TGO. This fact has a beneficial impact on the lifetime of TBC systems decreasing the rate
of formation of defects at the TGO/BC/TC interfaces. However, the optimal pre-oxidations
parameters (e.g., pO , time and temperature) were not reported in their study.
2
Notwithstanding, that the effects of these parameters have been investigated widely, there is a lack
information concerning the effect of the temperature, grit blasting, oxygen partial pressure and time
on the competitive growth between θ-Al2O3 and α-Al2O3 during the early stages of oxidation (less
that 5 h), which is important for the optimization of the TGO growth during pre-oxidation
treatments. Thus, the aim of this work is to study their effects on the oxidation kinetics of
commercial β-(Ni,Pt)Al BC systems during pre-oxidation treatments.
7.2 Experimental Procedures
7.2.1 Sample Preparation
The samples used in this work were provided by GE Aircraft Engines (Evendale, OH) as
rectangular specimens (1.8 x 1.2 x 0.15 cm), each weighing about 2.9 g. Bond coats were produced
by electroplating a thin layer of Pt onto Rene N5 single-crystal Ni-based superalloy coupons. A
heat treatment was subsequently performed to diffuse Pt into the substrate, followed by a vapor
phase aluminide process (VPA) to introduce Al into the coating. After a second heat treatment the
desired β-(Ni,Pt)Al phase was achieved. The resulting bond coat was a bilayer structure consisting
of an ∼50 μm thick β-(Ni,Pt)Al and an ∼20 μm thick inter-diffusion zone (IDZ).
In order to study the effect of the grit blasting the bond coat, some samples were subjected to a
light grit-blasting using fine corundum grit. The nominal compositions of the substrate and bond
91
coat were determined using inductively coupled plasma atomic emission spectroscopy (ICP-AES,
Perkin Elmer Optima DV4300) and electron micro probe analysis (EPMA, JEOL 8900 WD/ED
microanalyzer) respectively, and are listed in table 7.1.
Table 7.1. Chemical composition of single-crystal superalloy Rene N5 and β-(Ni,Pt)Al bond coat
Sample
Weight percent
Cr
Co
Mo
Re
W
Al
Ti
Ta
Hf
Pt
Ni
Nominal ReneN5*
7.0
8.0
2.0
3.0
5.0
6.2
-
7.0
0.2
-
Bal.
As-received ReneN5**
6.19
8.27
1.39
3.23
5.03
6.37
0.01
6.93
0.15
-
Bal.
As- coated β-(Ni,Pt)Al
BC***
0.80
2.66
0.045
-
0.01
24.91
-
0.28
-
31.86
Bal.
* Ref. [3]; ** Measured by ICP-AES, *** Measured by EPMA.
All samples were ultrasonically cleaned using xylene, acetone, methanol-water (1:1), and deionized
water for 15 min each to remove surface contamination prior to oxidation treatments. The surface
roughness of the samples prior to the oxidation treatments was evaluated using a
Mitutoyo Surftest SJ-201P surface roughness tester.
7.2.2 Oxidation Treatments
Thermogravimetric analyses (TGA) were performed on a Setaram Setsys Evolution 16/18
thermobalance, which has an accuracy of ±0.03 μg. The mass change during the TGA experiments
was recorded at 5 s intervals. The sample was hung from one end of the beam balance using a 0.4
mm diameter Pt/30%Rh wire and placed in the hot zone of the vertical furnace. To minimize any
perturbation resulting from gas flow, buoyancy, drag forces or premature sample oxidation before
reaching the working temperature, the heating cycle was programmed as follows: 1) evacuation of
the analysis chamber to less than 10 Pa, 2) heating to the reaction temperature at a maximum rate
of 50°C/min, 3) back filling the chamber with working gas to atmospheric pressure using a flow
rate of 200 ml/min, 4) fixing the flow of the working gas to 20 ml/min, and finally 5) cooling the
furnace chamber to room temperature at a maximum rate of 50°C/min with a gas flow rate of 0.3
ml/min. In order to evaluate the effect of the pO on the oxidation kinetics of β-(Ni,Pt)Al BC
2
systems, three different conditions were used: 1) pO = 2.1 x 10-1 atm, 2) pO = 5 x 10-5 atm and 3)
2
2
-5
pO = 1 x 10 atm. To reach the desired oxygen concentrations, argon with 1 x 10-4 atm of O2 was
2
92
passed through a gas purification system (Oxygon Industries OG-120M). The oxidations
parameters are summarized in table 7.2.
Table 7.2. Oxidation conditions tested in this study
As-Aluminized Samples
pO (atm) Time (h)
Temperature (°C)
2
1000
1050
1150
900
950
1000
1050
1100
1150
1200
2.1 x 10-1
2.1 x 10-1
2.1 x 10-1
1 x 10-5
1 x 10-5
1 x 10-5
1 x 10-5
1 x 10-5
1 x 10-5
1 x 10-5
5
5
5
5
5
5
5
5
5
5
Grit-Blasted Samples
Temperature (°C) pO2 (atm)
Time (h)
1000
1050
1100
1150
1000
1050
1100
1150
1000
1050
1100
1150
2.1 x 10-1
2.1 x 10-1
2.1 x 10-1
2.1 x 10-1
5 x 10-5
5 x 10-5
5 x 10-5
5 x 10-5
1 x 10-5
1 x 10-5
1 x 10-5
1 x 10-5
2
2
2
2
2
2
2
2
2
2
2
2
7.2.3 Characterization
The structural evolution of the TGO was monitored by photo-stimulated luminescence
spectroscopy (PSLS) [32-34] using a micro-Raman mapping spectrometer (Renishaw InVia)
connected to a Leica microscope equipped with a 532 nm line-focus laser. The laser power at the
sample was 5 mW, and the acquisition time for each spectrum was in the 0.5-2.0 s range. The PSLS
technique is based on photon emission from Cr3+ ions, a typical impurity incorporated in the crystal
structure of Al2O3 in TBC systems. Once illuminated, the Cr3+ ions emit fluorescent radiation due
to radioactive decay of the excited electrons to the ground state [35]. The identification of the
alumina phase is done on the basis of the well-characterized frequencies of the α- Al2O3 and θAl2O3 [33] by the different Cr3+-O2 bond distances from the α and θ structures. The microstructure
of the TGO was characterized using field emission scanning electron microscopy (Jeol JSM7401F
FEG-SEM).
93
Fig. 7.1. Bond coat system morphology before and after grit blasting: (a and b) surface view, (c
and d)cross-sectional view and (e and f) surface roughness.
94
7.3 Results and Discussion
7.3.1 Effect of Grit Blasting on β-(Ni,Pt)Al BC Morphology
Figure 7.1 shows a comparison of the microstructure of the β-(Ni,Pt)Al BC systems used in this
work before and after grit-blasting treatment including the plan view (Fig. 7.1a and 7.1b), the crosssectional view (Fig. 7.1c and 7.1d) and the average roughness (Fig. 7.1e and 7.1f). It can be seen
that alumina grits remains embedded in the BC surface, the BC grain boundaries (ridges visible in
Fig. 7.1a) were eliminated and the BC thickness was slightly reduced. The measured BC roughness
of the grit-blasted samples is similar to, but distinctly less than the as-aluminized samples, 2.03 ±
0.29 µm versus 1.76 ± 0.21.
7.3.2 Effect of Oxidation Temperature and Oxygen Partial Pressure
Figures 7.2a and 7.2b show the net mass gain (Δm) versus the square root of time (t0.5) during the
isothermal oxidation of as-aluminized β-(Ni,Pt)Al BC system samples for 5 h at different oxygen
partial pressures, pO = 1 x 10-5 atm (Fig. 7.2a) and pO = 2.1 x 10-1 atm (Fig. 7.2b). The curve for
2
2
-1
the sample treated at 1100°C with pO = 2.1 x 10 atm has not included because of atypical noise
2
during the test. Fig. 7.2a shows a mass gain rate increase from 1000 to 1100°C followed by an
abrupt decrease at 1150°C while Fig. 7.2b only show an increase from 1000 to 1150°C.
As a result of the effect of the pO increases in the net mass gain of 68, 47 and 159% were observed
2
when the samples were oxidized
in an atmosphere with a pO = 2.1 x 10-1 atm, at 1000, 1050
2
and 1150°C, respectively, and compared with samples oxidized at a pO = 1 x 10-5 atm. All the
2
curves exhibit deviations from the straight line expected for the classic parabolic model proposed
to describe the growth of a diffusion-controlled oxide, determined experimentally by Tamman [36]
and subsequently treated theoretically by Wagner [37] and described by the following equation:
∆m = kp t0.5
7.1
where kp is the parabolic rate constant. The observed deviations are associated with causes such as
BC impurities (e.g., doping elements), grain-boundary diffusion (e.g., TGO grain coarsening),
simultaneous reaction steps (e.g., diffusion/reaction control stages) and the polymorphism
95
associated with TGO and their corresponding transformations (e.g., γ-Al2O3 → θ-Al2O3 → αAl2O3) to form α-Al2O3 [20].
As-aluminized samples
-5
pO2 = 1 x 10 atm
0.22
0.18
1150°C
-1
pO = 2.1 x 10 atm
2
0.20
1100°C
1050°C
1000°C
0.18
2
TG Net Mass-Gain (mg/cm )
2
TG Net Mass-Gain (mg/cm )
0.20
As-aluminized samples
0.22
0.16
0.14
1050°C
0.12
1000°C
0.10
1150°C
0.08
0.06
0.16
0.14
0.12
0.10
0.08
0.06
0.04
0.04
0.02
0.02
(a)
0.00
0
20 40 60 80 100 120 140 160
0.5
0.5
(b)
0.00
0
20 40 60 80 100 120 140 160
0.5
0.5
(s )
t
(s )
0.5
Fig. 7.2. Δm vs. t curves showing the effect of the pO on the oxidation behavior of as-aluminized
2
β-(Ni,Pt)Al BC system samples during the first 5 h of treatment with (a) pO = 1 x 10-5 atm and (b)
2
pO = 2.1 x 10-1 atm.
2
t
In order to know if the observed deviations are related with this polymorphic transformation, a
systematic PSLS analysis was carried out on the surface (TGO) of all samples.
Figures 7.3a and 7.3b are PSLS curves that illustrate the structural evolution of the TGO from 1000
to 1150°C after 5 h of treatment in an atmosphere with a pO = 1 x 10-5 atm and pO = 2.1 x 10-1
2
2
atm respectively. It should be noted that α-Al2O3 luminescence is 10-12 stronger than for θ-Al2O3,
an important consideration during the analysis of the results [38]. Fig. 7.3a shows the presence of
both θ-Al2O3 and α-Al2O3 at 1000°C followed by a decrease in the θ-Al2O3 phase (qualitatively
observed by the change in relative peaks ratio) with increasing treatment temperature, until it
disappears at 1100°C < T ≤ 1150°C. A semiquantitative study of the evolution of θ- and α-Al2O3
phases at this pO has been reported elsewhere [39]. When using a higher pO (Fig. 7.3b), α-Al2O3
2
2
96
is not present at 1000°C, however the full transformation from θ-Al2O3 → α-Al2O3 occurs at
1050°C < T ≤ 1100°C.
As-aluminized samples
-5
pO2 = 1 x 10 atm
As-aluminized samples
(a)
α-Al2O3
α-Al2O3
-1
pO2 = 2.1 x 10 atm
(b)
1150°C
1150°C
α-Al2O3
θ-Al2O3
α-Al2O3
θ-Al2O3
1100°C
1050°C
Counts (a.u.)
Counts (a.u.)
α-Al2O3
1100°C
α-Al2O3
1050°C
θ-Al2O3
α-Al2O3
θ-Al2O3
1000°C
14350 14400 14450 14500 14550 14600 14650
θ-Al2O3
1000°C
14350 14400 14450 14500 14550 14600 14650
Wavelength (cm-1)
Wavelength (cm-1)
Fig. 7.3. Representative luminescence (PSLS) spectra taken from the TGO surface after 5 h of
isothermal oxidation at different temperatures and oxygen partial pressure: (a) pO = 1 x 10-5 atm
2
and (b) pO = 2.1 x 10-1 atm. Note that θ-Al2O3 luminescence is 10-12 times weaker than α-Al2O3.
2
In order to study the evolution of the parabolic rate constant, kp, as a function of the oxidation
conditions, a local fitting procedure proposed by Monceau and Pieraggi [40] as a general parabolic
law has been used:
t=A+B∆m+C∆m2
7.2
where the coefficients A, B, and C are related to kinetic parameters. Table 7.3 summarizes the
possible cases derived from this general parabolic law.
It is important to point out that in all cases kp is directly related with C-1. The local fitting procedure
is based on parabolic fitting of (Δm,t) data over short time intervals to obtain instantaneous values
of the parabolic rate constant.
97
Table 7.3. Summary of rate equations derived from the general parabolic model proposed by D.
Monceau and B. Pieraggi (1998)+
Case
Rate Equation*
Kinetics Law
A
B
C
1
kp
d∆m
=
dt
2∆m
t=
∆m2
kp
0
0
1
kp
2
kp
d∆m
=
dt
2∆m
t-ti =
∆m2i
kp
0
1
kp
2∆mi
kp
1
kp
∆mi ∆m2i
ti kp
kl
1
kl
1
kp
∆mi ∆m2i
+
kl
kp
1 2∆mi
+
kl
kp
1
kp
∆m2 -∆m2i
kp
ti -
3a
kp
d∆m
=
dt
2(∆m-∆mi )
(∆m-∆mi )2
t-ti =
kp
ti +
3b
d∆m
1
=
1 2∆m
dt
k + k l
p
∆m2 -∆m2i ∆m-∆mi
t-ti =
+
kp
kl
3c
d∆m
1
=
1 (2∆m-∆mi )
dt
k +
kp
l
(∆m-∆mi )2
t-ti =
kp
∆m-∆mi
+
kl
ti -
∆m2i
kp
-
* The initial conditions used to integrate the rate equations were t=ti and Δm=Δmi.
+
Adapted from Ref. [40].
Figures 7.4a and 7.4b show the time dependence of kp, including the transient and steady-state
regimes. Figures 7.4c and 7.4d are enlargements of the steady-state regimes. Based on Fig. 7.4 all
samples have reached a steady state after 4 h of treatment, with except the exception of the sample
treated at 1100°C at a pO = 1 x 10-5 atm for which kp is still decreasing after 4.5 h of treatment.
2
98
pO2 = 1 x 10-5 atm
(a)
kp x 10-7 (mg2/cm4s)
100
10
1050°C
1100°C
1
1000°C
1150°C
0.1
0
1
2
t (h)
3
4
5
As-aluminized samples pO2 = 2.1 x 10-1 atm
(b)
100
kp x 10-7 (mg2/cm4s)
As-aluminized samples
1150°C
1000°C
10
1
0.1
1050°C
0
1
2
t (h)
3
4
5
10
1050°C
1100°C
1000°C
1
1150°C
0.1
3.8
4.0
4.2
4.4
4.6
4.8
(d)
kp x 10-7 (mg2/cm4s)
kp x 10-7 (mg2/cm4s)
(c)
5.0
1000°C
10
1150°C
1050°C
1
0.1
3.8
4.0
4.2
t (h)
4.4
4.6
4.8
5.0
t (h)
Fig. 7.4. Log kp plotted as a function of time for as-aluminized β-(Ni,Pt)Al BC system samples
during isothermal oxidation at different temperatures and oxygen partial pressures: (a and b)
transient and steady-state regimes and (c and d) enlargement of the steady-state regime.
Figure 7.5 shows the effect of the pO and temperature on the oxidation kinetics of the as2
aluminized β-(Ni,Pt)Al BC system samples after 5-10 min and after 5 h of treatment. A clear effect
in the decrease of kp values (approximately one order of magnitude) was seen for the first few
minutes of oxidation for all temperatures tested at the lower oxygen partial pressure condition, pO
2
-5
-1
= 1 x 10 atm, relative to atmospheric pressure, pO = 2.1 x 10 atm.
2
Good agreement was observed between the growth kinetics (the slope) obtained by Brumm and
Grabke [20] for θ-Al2O3 from 1000 to 1100°C at approximately the same oxidizing conditions (pO
2
-1
= 1.3 x 10 atm). The decrease of kp values observed at 1150°C for both pO conditions could be
2
related with the rapid formation and high thermodynamic stability of α-Al2O3 at this temperature.
A more complex behavior should be expected for longer times (e.g., after 5 h of treatment) due to
the phenomena associated with the θ-Al2O3 → α-Al2O3 transformation as have been reported by
99
different authors [20,38,41], including strong dependence on chemical composition, surface
conditions and oxidizing parameters.
Temperature (°C)
1200 1150 1100 1050 1000
900
-1
pO = 2.1 x 10 atm
2
@ 5-10 min
100
kp x 10-7 (mg2/cm4s)
950
-5
pO = 1 x 10 atm
2
@ 5-10 min
θ-Al2O3
10
1
0.1
-1
pO = 2.1 x 10 atm
2
@5h
-5
pO = 1 x 10 atm
2
@5h
γ-Al2O3
α-Al2O3
0.01
6.6 6.8 7.0 7.2 7.4 7.6 7.8 8.0 8.2 8.4 8.6 8.8
Reciprocal Temperature (104/K)
Fig. 7.5. Arrhenius plot showing the effect of low oxygen partial pressure on the kp values of asaluminized β-(Ni,Pt)Al BC system samples. Bold lines refer to Brumm and Grabke data for a NiAl
alloy [20].
Veal, et al. [41] proposed that oxide growth and transformation can be divided into three regions.
The first stage when TGO growth is dominated by the growth of metastable phases (e.g., θ-Al2O3).
After an incubation period, stable α-Al2O3 nucleates at the metal-oxide (M/O) interface. In the
second stage, a uniform layer of α-Al2O3 is established along the M/O interface (because the rate
lateral growth rate of the grains is faster than the vertical), blocking outward transport of Al and
stifling the growth of θ-Al2O3 (which is defined as a p-type oxide). Once a uniform α-Al2O3 scale
is established at M/O interface, new oxide formation is now mainly controlled by the slow growth
100
kinetics of α-Al2O3. This stage is characterized by θ-Al2O3→α-Al2O3 transformation along the
M/O interface. Finally during the third stage, after transformation is completed, α-Al2O3 is the only
phase present in the TGO, and it controls the scale growth kinetics.
Table 7. 4. Summary of the kinetics and structural analysis of the thermally grown oxide on asaluminized β-(Ni,Pt)Al BC systems
Tested parameters
TGA
Kinetics Analysis
Net mass gain
T(°C)
pO2
time
(atm)
(h)
(mg/cm2)
@2h
@5h
PSLS
kp x 10-7 (mg2/cm4s)
@ 5-10 min
θ-Al2O3
@2h
@5h
α-Al2O3
@5h
As-Aluminized Samples
1000
2.1 x 10-1
5
0.114
0.178
37.012
14.953
6.337
Yes
Yes
1050
2.1 x 10-1
5
0.158
0.201
132.103
8.352
1.752
Yes
Yes
1100
2.1 x 10-1
5
-
-
-
-
-
No
Yes
1150
2.1 x 10-1
5
0.148
0.212
84.669
26.612
3.354
No
Yes
1000
1 x 10-5
5
0.07
0.106
11.604
4.045
1.242
Yes
Yes
1050
1 x 10-5
5
0.092
0.137
24.149
7.548
5.107
Yes
Yes
1100
1 x 10-5
5
0.145
0.192
76.947
9.271
1.366
Yes
Yes
1150
1 x 10-5
5
0.065
0.082
8.871
1.093
0.198
No
Yes
In order to have a better understanding of the effect of the pO on the kp values after 5 h of oxidation
2
treatment these they were correlated with the PSLS measurements (Table 7.4). For the samples
treated at 1000°C, where θ-Al2O3 was the dominant phase, a decrease in kp is observed at lower
oxygen concentration. At 1050°C an opposite behavior in kp was registered. At first glance, this
result might be considered contradictory; however, this behavior can be explained if it is assumed
that different regions of growth (based on the regions proposed by Veal et al.[41]) are controlling
the formation of the alumina scale. For instance, for the sample oxidized at a pO = 1 x 10-5 atm,
2
the TGO is controlled by the growth of θ-Al2O3 (region 1), whereas for the sample oxidized at a
pO = 2.1 x 10-1 atm it is mainly controlled by the growth of α-Al2O3 (region 2). For the samples
2
treated at 1150°C where α-Al2O3 controls the oxide growth, a decrease in the kp value is observed
at the lower oxygen concentration, as expected.
101
7.3.3 Effect of Oxidation Temperature and Grit-Blasting Process
In order to study the effect of the grit-blasting process on the oxidation behavior of the β-(Ni,Pt)Al
BC systems, a group of samples were subjected to a grit-blasting step after the aluminizing process
in order to remove the characteristic BC grain boundaries (Fig. 7.1). After a cleaning step the
samples were oxidized with pO = 2.1 x 10-1 atm for 2 h at different temperatures and compared
2
with as-aluminized samples oxidized for 5 h under the same oxidizing atmosphere.
Grit-blasted samples
0.16
0.16
-1
pO = 2.1 x 10 atm
2
0.14
1000°C
0.10
0.08
1150°C
0.06
1100°C
0.04
2
1050°C
0.12
TG Net Mass-Gain (mg/cm )
2
TG Net Mass-Gain (mg/cm )
0.14
0.12
As-aluminized samples
1050°C
-1
pO = 2.1 x 10 atm
2
1150°C
1000°C
0.10
0.08
0.06
0.04
0.02
0.02
(a)
0.00
0
12 24 36 48 60 72 84 96
0.5 0.5
t
(s )
(b)
0.00
0
12 24 36 48 60 72 84 96
0.5 0.5
t
(s )
Fig. 7.6. Δm vs. t0.5 curves showing the effect of the grit-blasting process on the oxidation behavior
with pO = 2.1 x 10-1 atm during the first 2 h of treatment: (a) grit-blasted sample and (b) as2
aluminized β-(Ni,Pt)Al BC samples.
A comparison of the Δm vs t0.5 curves with and without grit blasting after 2 h of treatment is
presented in Fig. 7.6. Fig 7.6a (grit-blasted samples) shows a mass gain rate increase from 1000 to
1050°C. An abrupt decrease was observed at 1100°C followed by an increase at 1150°C. A similar
behavior was observed for the as-aluminized samples (Fig. 7.6b). However, the grit-blasting
process results in a decrease of the net mass gain after 2 h of treatment of 11.40, 25.32 and 52.02%
at 1000, 1050 and 1150°C respectively. The correlation between the TG results (Fig. 7.6a) and
102
PSLS measurements (Fig. 7.7) for the grit-blasted samples show that the abrupt decrease in the net
mass gain at 1100 and 1150°C is because α-Al2O3 is controlling the growth of the TGO.
Grit-blasted samples
-1
pO2 = 2.1 x 10 atm
α-Al2O3
1150°C
Counts (a.u.)
α-Al2O3
1100°C
α-Al2O3
1050°C
θ-Al2O3
α-Al2O3
1000°C
θ-Al2O3
14350 14400 14450 14500 14550 14600 14650
Wavelength (cm-1)
Fig. 7.7. Representative luminescence (PSLS) spectra taken from the surface of the thermally
grown oxide on grit-blasted β-(Ni,Pt)Al BC systems samples after 2 h of isothermal oxidation at
different temperatures with pO = 2.1 x 10-1 atm. It is important to highlight that θ-Al2O3
2
luminescence is 10-12 times weaker than α-Al2O3.
Furthermore , grit-blasting process accelerates the θ-Al2O3→α-Al2O3 transformation as shown by
comparing Fig. 7.3b and Fig. 7.7. In the as-aluminized sample treated at 1000°C for 5 h, only θAl2O3 is observed, whereas in the growth of nearly pure α-Al2O3 is observed after only 2 h for the
grit-blasted sample. Once again, it is important to remember that α-Al2O3 luminescence is 10-12
times stronger than θ-Al2O3, and thus θ-Al2O3 is still controlling the TGO growth of the grit-blasted
sample treated at 1000°C.
103
Tolpygo and Clarke [38] reported that the θ-Al2O3→α-Al2O3 transformation is highly dependent
on surface roughness showing that a smooth surface delays the transformation relative to a rough
surface. They postulated that scratches and other microscopic defects on the BC surface provide
sites for α-Al2O3 nucleation. This argument has been used by different authors to explain the effect
of grit blasting on the pre-oxidation behavior of β-(Ni,Pt)Al BC systems prior to TC deposition
[17,23]. In our case the surface roughness is very similar for the as-aluminized and grit-blasted
samples (Fig. 7.1), but the density of surface defects is clearly higher on the grit-blasted surfaces,
thus supporting the postulation of Tolpygo and Clarke [28].
1000°C
Grit-blasted samples
-1
pO2 = 2.1 x 10 atm
10
1150°C
1100°C
1050°C
1
0.0
1000°C
kp x 10-7 (mg2/cm4s)
kp x 10-7 (mg2/cm4s)
100
Grit-blasted samples
-1
pO2 = 2.1 x 10 atm
10
1150°C
1100°C
1050°C
0.5
1.0
t (h)
1.5
2.0
0.4
0.6
0.8
1.0
1.2 1.4
t (h)
1.6
1.8
2.0
Fig. 7.8. Log kp plotted as a function of time for grit-blasted β-(Ni,Pt)Al BC system samples during
isothermal oxidation at different temperatures and pO = 2.1 x 10-1 atm: a) transient and steady-state
2
regimes and c) steady-state regime.
Figure 7.8 shows the time dependence of kp for the grit-blasted samples oxidized with pO = 2.1 x
2
10-1 atm. Based on these curves, it can be concluded that the transition from the θ-Al2O3-controlled
to the α-Al2O3-controlled has changed from approximately 4 h for the as-aluminized β-(Ni,Pt)Al
BC system samples (Fig. 7.4b and Fig. 7.4d) to less than 2 h for the grit-blasted samples oxidized
at T > 1000°C. For example, the grit-blasted sample treated at 1100°C reached a steady state after
0.5 h of oxidation, while the as-aluminized sample was still in the transient regime after 4.5 h of
treatment. Similarly the grit-blasted sample oxidized at 1150°C reached a steady state after 1.2 h
while the as-aluminized sample required 4.5 h to reach the same regime.
Figure 7.9 shows an Arrhenius plot comparing the kp values after 5-10 min and after 2 h of treatment
for the grit-blasted samples with the kp values after 5-10 min, 2 h and 5 h of treatment for the as104
aluminized samples. All these samples had been treated with pO = 2.1 x 10-1 atm. The graph also
2
includes Brumm and Grabke´s data for a NiAl alloy [20].
Temperature (°C)
1200 1150 1100 1050 1000
900
As-aluminized S.
@ 5-10 min
100
kp x 10-7 (mg2/cm4s)
950
Grit-blasted S.
@ 5-10 min
As-aluminized S.
@2h
10
θ-Al2O3
Grit-blasted S.
@2h
As-aluminized S.
@5h
1
0.1
α-Al2O3
γ-Al2O3
-1
pO2 = 2.1 x 10 atm
6.6 6.8 7.0 7.2 7.4 7.6 7.8 8.0 8.2 8.4 8.6 8.8
Reciprocal Temperature (104/K)
Fig. 7.9. Arrhenius plot showing the effect of the grit-blasting process on the kp values of β(Ni,Pt)Al BC system samples. Bold lines refer to Brumm and Grabke data for a NiAl alloy [20].
After 5-10 min of oxidation the as-aluminized β-(Ni,Pt)Al BC system samples have higher kp values
than those that were grit-blasted. This result can be explained by a decrease in the outward diffusion
of Al caused by removal the BC ridges during the grit-blasting process. After 2 h of oxidation, the
as-aluminized samples still have not reached a steady state with the exception of the sample treated
at 1000°C, which exhibits an intermediate steady state (corresponding to the θ-Al2O3 phase) from
1.5 to 3.5 h of (Fig. 7.4b), so the estimated kp values are useful only for comparative purposes. In
order to understand the oxidation behavior once a steady-state regime was established, the kp values
of the samples with and without grit-blasted process were compared after 2 and 5 h of isothermal
105
oxidation, respectively. The grit-blasted samples have higher kp values than the as-aluminized
samples for all temperatures tested except for the sample treated at 1000°C where the as-aluminized
sample has a slightly higher value. This result is in good agreement with the findings of Tolpygo
et al. [18, 28], who showed that the oxide formed on the grit-blasted β-(Ni,Pt)Al BC surface
exhibits a higher growth rate. The trend in kp as a function of temperature is similar for both tested
conditions, decreasing from 1000 to 1050°C, followed by an increase in the parabolic rate constant.
A summary of the kinetics behavior results for the grit-blasted samples is presented in table 7.5.
Table 7. 5. Summary of the kinetics and structural analysis of the thermally grown oxide on gritblasted β-(Ni,Pt)Al BC systems
Tested parameters
TGA
Kinetics Analysis
Net mass gain
T (°C)
pO2
time
(atm)
(h)
2
(mg/cm )
@2h
1000
1050
1100
1150
1000
1050
1100
1150
1000
1050
1100
1150
2.1 x 10-1
2.1 x 10-1
2.1 x 10-1
2.1 x 10-1
5 x 10-5
5 x 10-5
5 x 10-5
5 x 10-5
1 x 10-5
1 x 10-5
1 x 10-5
1 x 10-5
2
2
2
2
2
2
2
2
2
2
2
2
kp x 10-7 (mg2/cm4s)
@ 5-10 min
Grit-Blasted Samples
0.101
36.231
0.118
52.971
0.051
10.671
0.071
46.677
0.028
4.614
0.030
28.323
0.040
31.924
0.066
195.971
0.025
2.461
0.027
18.784
0.034
7.195
0.055
44.647
PSLS
θAl2O3
@2h
@2h
5.471
3.402
5.182
6.548
0.312
1.572
3.485
9.987
0.395
0.886
1.019
5.216
α-Al2O3
Yes
Yes
No
No
No
No
No
No
No
No
No
No
Yes
Yes
Yes
Yes
Yes
Yes
Yes
Yes
Yes
Yes
Yes
Yes
7.3.4 Mixed Effect of Low the Oxygen Partial Pressure and the Grit-Blasting Process
This final group of experiments was carried out with the purpose of studying the combined effect
of low pO and grit-blasting process on the oxidation behavior of β-(Ni,Pt)Al BC systems as a
2
function of isothermal treatment at different temperatures.
106
0.07
0.07
0.06
Grit-blasted samples
-5
pO = 1 x 10 atm
2
1150°C
Grit-blasted samples
-5
pO = 5 x 10 atm
2
0.06
0.05
1100°C
0.04
1050°C
0.03
1000°C
0.02
2
TG Net Mass-Gain (mg/cm )
2
TG Net Mass-Gain (mg/cm )
1150°C
0.01
0.05
0.04
1100°C
0.03
1050°C
1000°C
0.02
0.01
(a)
0.00
0
20
40
60
0.5 0.5
t
(s )
80
(b)
0.00
100
0
20
40
60
0.5 0.5
t
(s )
80
100
Fig. 7.10. Δm vs. t0.5 curves showing the combined effect of the grit-blasting process and low
oxygen pressure on the oxidation behavior of grit-blasted β-(Ni,Pt)Al BC system samples during
the first 2 h of treatment: (a) pO = 5 x 10-5 atm and (b) pO = 1 x 10-5 atm.
2
2
Figure 7.10 shows the TGA curves of the 2 h pre-oxidation treatments performed from 1000 to
1150°C at two different oxidizing atmospheres, pO = 5 x 10-5 atm (Fig. 7.10a) and pO = 1 x 10-5
2
2
atm (Fig. 7.10b). Clear deviations from the expected linear behavior for the classic parabolic model
(Eq. 1) were observed for all samples during the early stages of isothermal oxidation (~20 min).
For the sample treated at 1000°C, at least two distinct transitions (infection points) were observed
(Fig. 7.10a). In the same way, Fig. 7.10b shows two transitions for the samples treated at 1000 and
1100°C. Reducing the amount of O2 in the working gas led to a decrease of 11, 10, 15 and 17% in
the net mass gain at 1000, 1050, 1100 and 1150°C respectively.
There was an increase in the net mass gain as a function of the treatment temperature for both
oxidizing conditions. This behavior suggests that a same phase (presumably α-Al2O3) is dominating
the oxide growth, even during very early stages of oxidation. This assumption was confirmed by
PSLS measurements and a kinetics analysis using the local parabolic fitting previously described
107
[40]. Based on, the PSLS measurements, it can be concluded that α-Al2O3 is the only phase present
in the TGO after 2 h of oxidation treatment for all the evaluated conditions (Table 7.5).
The kinetics analysis presented in the Fig. 7.11 shows the time dependence of kp as a result of the
pre-oxidation treatments performed from 1000 to 1150°C with pO = 5 x 10-5 atm and pO = 1 x 102
2
atm.
Grit-blasted samples
-5
pO2 = 5 x 10 atm
(a)
kp x 10-7 (mg2/cm4s)
100
1150°C
10
1100°C
1
0.1
0.0
1050°C
1.0
t (h)
1.5
10
1150°C
1100°C
1
1050°C
1000°C
1000°C
0.5
Grit-blasted samples
-5
pO2 = 1 x 10 atm
(b)
100
kp x 10-7 (mg2/cm4s)
5
2.0
0.1
0.0
0.5
1.0
t (h)
1.5
2.0
Fig. 7.11. Log kp plotted as a function of time for grit-blasted β-(Ni,Pt)Al BC system samples
during isothermal oxidation at different temperatures and oxygen partial pressure: (a) pO = 5 x 102
5
atm and (b) pO = 1 x 10-5 atm.
2
Based on this analysis, for the sample oxidized at 1000°C and a pO = 5 x 10-5 atm, the transient
2
regime is composed of three transition zones, with a stepwise decrease of kp before reaching a
steady-state regime after about ~1.25 h where α-Al2O3 controls the oxidation kinetics. A similar
behavior was observed in Fig. 7.8 for the grit-blasted samples oxidized in a pO = 2.1 x 10-1 atm at
2
1000 and 1050°C, however a longer time was required to reach steady state (~ 1.7 h). The samples
treated at 1000 and 1100°C with a pO = 1 x 10-5 atm (Fig. 7.11b) also have a transient regime, but
2
theirs are composed of only two clear transition zones before reaching steady state after about 0.5
and 0.75 h respectively. Steady state was reached after about 1.25 h, without obvious transient
regimes, for the samples treated at 1050°C and 1150°C. For the samples treated at 1050, 1100 and
1150°C with a pO =5 x 10-5 atm (Fig. 7.11a), a well developed steady-state regime was observed
2
after about 1, 0.5 and 0.75 respectively. The combined effects of grit blasting and low pO on the
2
108
oxidations kinetics, after steady state was well developed are shown in an Arrhenius plot (Fig.
7.12).
Temperature (°C)
1200 1150 1100 1050 1000
kp x 10-7 (mg2/cm4s)
α-Al2O3
10
950
900
θ-Al2O3
-5
pO = 5 x 10 atm
2
@2h
1
-5
pO = 1 x 10 atm
2
@2h
γ-Al2O3
0.1
Grit-blasted samples
6.6 6.8 7.0 7.2 7.4 7.6 7.8 8.0 8.2 8.4 8.6 8.8
Reciprocal Temperature (104/K)
Fig. 7.12. Arrhenius plot showing the combined effect of the grit-blasting process and low oxygen
pressure on the kp values of β-(Ni,Pt)Al BC systems. Bold lines refer to Brumm and Grabke data
for a NiAl alloy [20].
An increase in the kp values as a function of temperature was observed for both oxidizing
conditions. Increasing the pO results in little change in kp at 1000°C, but a significant increase was
2
observed at 1050 (~1.7X), 1100(~ 3.42X) and 1150°C (~1.9X). Good agreement with the slope
obtained by Brumm and Grabke [20] for the growth α-Al2O3 was observed for both oxygen partial
pressures.
109
Figure 7.13 summarizes the oxidation behavior for all tested conditions presented in table 7.2. At
this point, it is possible to compare the combined and separate effects of grit-blasting process and
low-oxygen partial pressure on the competitive growth between θ-Al2O3 and α-Al2O3 as function
time and temperature.
Temperature (°C)
1200 1150 1100 1050 1000
950
900
100
kp x 10-7 (mg2/cm4s)
α-Al2O3
10
θ-Al2O3
1
0.1
0.01
Brumm & Grabke
As-aluminized samples @ 5 h
-5
1 x 10 atm
-1
2.1 x10 atm
Grit-blasted samples @ 2h
-5
1 x 10 atm
-5
5 x 10 atm
-1
2.1 x10 atm
γ-Al2O3
6.6 6.8 7.0 7.2 7.4 7.6 7.8 8.0 8.2 8.4 8.6 8.8
Reciprocal Temperature (104/K)
Fig. 7.13. Arrhenius plot summarizing the oxidation behavior for all oxidizing conditions used in
this work, once that a steady-state regime was established. Bold lines refer to Brumm and Grabke
data for a NiAl alloy [20].
For the samples treated at 1000°C the following observations can be made:
1) At a pO = 2 x 10-1 atm θ-Al2O3 phase is controlling the growth of the TGO, consistent with
2
the PSLS results (Fig. 7.3b and Fig. 7.7).
110
2) Although no significant differences were observed in the kp values, grit blasting reduced the
time required to reach steady state from about 4 to 1.6 h at a pO = 2 x 10-1 atm.
2
-5
3) At a pO = 1 x 10 atm and 5 h of oxidation, an abrupt decrease in kp was registered for the
2
as-aluminized sample as a result of the oxygen diinution in the oxidizing atmosphere
decreasing the growth of θ-Al2O3 ( an outward cation diffusion oxide, or p-type oxide).
4) For the low pO conditions combined with the grit-blasting process, a synergistic effect to
2
stabilize the growth of α-Al2O3 was observed, suppressing the growth of metaestables
phases.
For the samples treated at 1050°C the following observations can be made:
1) At a pO = 2 x 10-1 atm, a clear transition was observed as a result of the θ-Al2O3 → α-Al2O3
2
transformation through the thickness of the TGO as was confirmed by PSLS measurements
(Fig. 7.3b and Fig. 7.7).
2) At a pO = 1x10-5 atm and 5 h of oxidation, kp increases with respect to the sample treated at
2
1000°C, because θ-Al2O3 is still the controlling the growth of the oxide scale.
3) As with the sample treated at 1000°C, suppression of the growth of metastable phases was
observed for the combined effect of low pO conditions and grit-blasting. The transient state
2
regime decreased from about 4 h for the as-aluminized sample treated in a pO = 2.1 x 10-1
2
-5
atm to1 h for the grit-blasted sample oxidized at a pO = 5 x 10 atm.
2
4) A large decrease of kp as a function of diminution in oxygen concentration was observed
for the grit-blasted samples (e.g., approximately 400% decrease from a pO = 2.1x10-1 atm
2
to pO = 1x10-5 atm). This behavior is also related to the stifling of the growth of θ-Al2O3.
2
For the samples treated at 1100°C the following observations can be made:
1) α-Al2O3 is the controlling phase for all the grit-blasted conditions.
2) The decrease in kp values as a function of temperature could be associated with the α-Al2O3
microstructure for each pO (e.g., density of nucleation site, grain coarsening, etc.).
2
Additional characterization is required to corroborate this assumption.
3) For the as-aluminized sample an steady-state regime had not been reached after 4.5 h of
treatment, which is in good agreement with Cadoret et al [26], who reported that at this
111
temperature Pt favors the growth of θ-Al2O3 and therefore the time to reach the full θAl2O3→α-Al2O3 transformation is increased.
Finally, for the samples treated at 1150°C the following observations can be made:
1) Variations in kp of more than an order of magnitude are observed with respect to the asaluminized sample oxidized at a pO = 1 x 10-5 atm. For instance, an increase of about
2
2600% in kp value is registered as a result of the combined effect of the grit-blasting process
and low oxygen partial pressure ( pO = 1 x 10-5 atm).
2
2) Meanwhile, a maximum variation of only about 300% increase was observed for all the
grit-blasted conditions with pO = 5 x 10-5 atm with respect to the as-aluminized sample
2
oxidized with pO = 2.1 x 10-1 atm. These results confirm that the grit blasting, even at low
2
pO , increases the oxidation rate, as has been previously reported by Tolpygo et al. [28].
2
3) Additionally, a low dependence of the kp on the pO was observed at 1150°C for the grit2
blasted conditions in comparison with the as-aluminized samples.
7.3.5 Analysis of Oxygen Transfer from Flowing to Sample Surface
In order to better understand the combined effect of low oxygen partial pressure and grit blasting
on shortening the transient-stage regime (e.g. stifling the growth of θ-Al2O3) during the early stages
of oxidation, a first approach to calculated the flux of oxygen to a flat surface from a gas flowing
parallel to it was carried out based on the theory of mass transfer in dilute gases [42]. The present
analysis assumed that the working gas contains only one reactive species to promote the sample
oxidation. In our case, once the treatment temperature is achieved in the furnace chamber, the
vacuum is broken, the working gas flows past the β-(Ni,Pt)Al BC samples and oxygen molecules
(the reactive species) are transferred from the bulk gas through the boundary layer to the samplegas interface by diffusion and convection phenomena [43]. Once the oxygen molecules have
reached the sample-gas interface, the reaction between the Al and O takes place to form an alumina
scale according to the reaction:
2Al (s) + 1.5O2 (g) Al2O3 (s)
7.3
If the oxygen flux, J*O , is given by
112
k
(o)
(i)
RT
2
2
J *O = m (pO -pO ) (mg O /cm2 s)
7.4
where km is mass-transfer coefficient, T absolute temperature, and R the universal gas constant. The
superscripts (o) and (i) represent the values at the working gas and solid surface interfaces
(o)
respectively. Three working gas mixtures, pO , were used in the present study, as described above;
2
(i)
pO was assumed to be zero; and km can be evaluated using:
2
km = 0.664 1
D4AB 6 vl 1/2
ν
l
(cm/s)
7.5
where DAB is the diffusion coefficient in an A-B binary gas, v the kinematic viscosity, vl the linear
velocity of the gas (fixed at 0.131 cm/s during the TGA analysis) and l the length of the surface
parallel to gas flow (1.8 cm for all samples). DAB and v values were estimated using the kinetic
theory of gases, using the Chapman-Enskog formulation and the gas molecular interaction
parameters shown in table 7.6 [42].
Table 7. 6. Gas Molecular Interaction Parameters [42].
Species
σ (Å)
ε/k (K)
M (g/mol)
O2
3.433
113
32.00
N2
3.667
99.8
28.01
Ar
3.432
122.4
39.948
Air
3.617
97
28.97
The DAB is defined as
DAB =
1.858 x 10-3 T3 1
1
+ MA MB
Pσ2AB ΩD,AB
σ +σB
σAB = A
2
1.06036
T
k
= √ A B (K-1)
T* =
ε ε
k
kT
εAB
7.6
(Å)
ΩD,AB = *0.1561 +
εAB
(cm2/s)
7.7
0.193
exp(0.47635T* )
+
1.03587
exp(1.52996T* )
+
1.76474
exp(3.89411T* )
7.8
7.9
7.10
113
where σ is the collision diameter, ΩD the diffusion collision integral, ε the characteristic energy of
interaction (used in the Lennard-Jones potential), M the molecular weight, k Boltzman´s constant
and P the system pressure (set to 1 atm during TGA analysis). The subscripts A and B represent the
gaseous species considered during the interactions.
Meanwhile v is given by
η
ν=
ρ
(cm2/s)
7.11
where η and ρ are the average values of viscosity and density of the species in the working gas,
respectively. η is given by
!M T
ηi = 2.6693 x 10-5 σ" i (poise)
7.12
#
xα ηα (poise)
α=1∑ x ϕ
β β αβ
7.13
ηmix = ∑N
1.16145
Ωη = *0.14874 +
T
ϕαβ =
1
√8
1+
0.52487
*
exp(0.7732T )
Mα
Mβ
-1/2
+
2.16178
7.14
exp(2.43787T* )
1/2
1+ ηα
ηβ
1/4 2
Mα
7.15
Mβ
where ηi and is the viscosity of the pure species i, ηmix the viscosity of mixed species and Ωη the
viscosity collision integral. The mass-transfer parameters for the different oxidizing conditions
used in this work were calculated from equations 7.4-7.15 and are tabulated in tables 7.7 and 7.8.
Table 7.7. Mass-transfer parameters for O2 in N2 (pO = 2.1 x 10-1 atm).
2
(vl/l)1/2
km
J *O
(cm/s1/2)
(s1/2)
(cm/s)
(mg O/cm2s)
1.18
1.16
0.27
3.13 x 10-1
3.02 x 10-2
2.57
1.27
1.20
0.27
3.23 x 10-1
2.99 x 10-2
1000
2.73
1.37
1.23
0.27
3.32 x 10-1
2.97 x 10-2
1150
2.90
1.47
1.27
0.27
3.41 x 10-1
2.95 x 10-2
T
DAB
(°C)
(cm2/s) (cm2/s)
1000
2.41
1050
Ν
0.664 (D4AB /ν)
1/6
114
Table 7.8. Mass-transfer parameters for O2 in Ar.
1/6
(vl/l)1/2
km
J *O (mg O/cm2s)
(cm/s)
pO = 1 x 10-5 atm pO = 5 x 10-5 atm
T
DAB
Ν
0.664 (D4AB /ν)
(°C)
(cm2/s)
(cm2/s)
(cm/s1/2)
1000
2.33
1.28
1.12
0.27
0.30
1.39 x 10-6
6.94 x 10-6
1050
2.48
1.37
1.15
0.27
0.31
1.38 x 10-6
6.89 x 10-6
1100
2.64
1.47
1.19
0.27
0.32
1.37 x 10-6
6.83 x 10-6
1150
2.80
1.58
1.22
0.27
0.33
1.36 x 10-6
6.78 x 10-6
(s1/2)
2
2
It is accepted that during the very early stages of oxidation of NiAl and NiPtAl alloys, the alumina
scale is mainly controlled by the growth of metastable phases (e.g., undoped γ- and/or θ-Al2O3
phases) via outward Al diffusion [44-46]. Accordingly, three main mass-transport processes are
taking place during the growth of metastable alumina phases: 1) Al diffusion within the alloy, 2)
Al transport across the TGO and 3) the flux of O2 from the working gas to the oxide surface (Eq.
7.4).
Figure 7.14 shows a comparison between the dmO/dt vs. t curves for the grit-blasted samples
oxidized at 1000°C along with the same curves for pure θ-Al2O3 and α-Al2O3 based on the
oxidation data reported by Brumm and Grabke [20] for NiAl alloys oxidized in a He- and O2 flow
with a pO = 1.28 x 10-1 atm. The kp values reported in the literature were used to calculated their
2
corresponding dmO/dt using:
JO = 2 =
1
kp
d(∆mO )
t
dt
7.16
The fluxes of oxygen available in the working gas used in this work are also indicated in the figure.
It is this flux that sustains the growth of the thermally grown oxide for the oxygen partial pressures.
∗
An intersection between a dmO/dt vs. t curve and its corresponding JO
represents the transition from
gas-supply control (for times before the intersection) to metal supply control (for times after the
intersection).
The grit-blasted sample oxidized with pO = 2.1 x 10-1 atm is in good agreement with Brumm and
2
∗
Grabke’s data for θ-Al2O3 [20]. In both cases the JO
is high enough that metal supply controls
oxidation from the start (Fig. 7.14a). The grit-blasted samples oxidized at low oxygen partial
115
pressures exhibit clearly different oxidation behavior with respect to the sample oxidized with pO =
2
-1
2.1 x 10 atm.
As a result of the diminution of oxygen fluxes (for the different pre-oxidation conditions), dmO/dt
decreases as a function of oxidation time, by approximately one order of magnitude after 2 h of
oxidation. In order to explain this behavior some assumptions must be made: 1) all the samples had
the same surface conditions (roughness, area, impurities, etc) and 2) the chemical composition, and
therefore chemical potential of the species that make up the BC, are the same after grit blasting.
With these assumptions, it can be established that as a result of the heat treatment at 1000°C, all
samples had the same Al outward diffusion from the BC to TGO/BC interface. When the oxygen
partial pressures was dramatically decreased from 2.1 x 10-1 atm to 5 x 10-5 atm and 1 x 10-5 atm,
gas-supply control was established at t ≤ 0.35 h and t ≤ 1.75 h respectively (Fig. 7.14b). This effect
is associated with the suppression of the growth of θ-Al2O3 observed after 2 h of oxidation and
confirmed by PSLS measurements and the kinetics analysis presented above.
116
0.1
dmO / dt (mg O/cm2s)
0.01
-1
*
JO @ 2.1 x10 atm
-1
*
JO @ 1.58 x10 atm
Grit-blasted samples
-5
1 x 10 atm
-5
5 x 10 atm
-1
2.1 x 10 atm
1E-3
θ-Al O (Brumm & Grabke)
2 3
1E-4
-5
*
JO @ 5 x10 atm
1E-5
1E-6
-5
*
JO @ 1 x10 atm
α-Al O (Brumm & Grabke)
2 3
1000°C
1E-7
0.0
0.5
1.0
1.5
2.0
t (h)
1E-4
Grit-blasted samples
-5
1 x 10 atm
-5
5 x 10 atm
-1
2.1 x 10 atm
θ-Al O (Brumm & Grabke)
2 3
dmO / dt (mg O/cm2s)
(a)
-5
*
JO @ 1 x10 atm
1E-5
1E-6
-5
*
JO @ 1 x10 atm
α-Al O (Brumm & Grabke)
1000°C
0.0
2 3
0.5
1.0
1.5
(b)
2.0
t (h)
Fig. 7.14. Comparison of dmO/dt vs. t curves for grit-blasted samples oxidized at 1000°C with the
available flux of oxygen during oxidation experiments for the different oxygen partial pressures
used. Red dashed curves refers to Brumm and Grabke data for a NiAl alloy [20].
117
7.4 Conclusions
Pre-oxidation treatments performed on as-aluminized samples at a pO2 = 1 x 10-5 atm did not
accelerate the θ-Al2O3→α-Al2O3 transformation, even after 5 h of oxidation relative to samples
oxidized in a pO2 = 2.1 x 10-1 atm, with the exception of the sample treated at 1000°C, where
apparently α-Al2O3 nucleation started earlier for the sample treated at a low pO2. Furthermore, the
samples treated at higher pressure, pO2 = 2.1 x 10-1 atm, reached full transformation from θAl2O3→α-Al2O3 at lower temperature than those treated at low pressure, pO2 = 1 x 10-5 atm
(1050°C versus 1100°C).
Grit blasting resulted in a decrease in the incubation time for the α-Al2O3 nucleation at all tested
temperatures (from 1000 to 1150°C). The time required to transition from θ-Al2O3-controlled to αAl2O3-controlled kinetics was reduced from about 4 h for the as-aluminized samples to less than 2
h for the grit-blasted samples when oxidized with pO2 = 2.1 x 10-1 atm at T >1000°C. This behavior
is associated with an increase in density of surface defects, which provide a greater number of αAl2O3 nucleation sites as suggested previously by Tolpygo and Clarke [38]. Moreover, an increase
in kp was observed as a result of grit blasting for the samples treated from 1050 to 1150°C. This
result was also in good agreement with Tolpygo et al. [28], who reported a tenfold increase in kp
for grit-blasted β-(Ni,Pt)Al coatings oxidized at 1150°C.
As a result of the synergistic effect of low pO2 (1 x 10-5 and 5 x 10-5 atm) and grit blasting, a
suppression of the transient kinetics of oxidation prior to reaching a steady state, where α-Al2O3 is
controlling the growth of alumina scale was establish during the early stages. For instance, for the
grit-blasted samples oxidized at 1000°C with pO2= 1 x 10-5 atm a steady-state regime was
established after about 1 h of oxidation, whereas that, for the grit-blasted sample oxidized with a
pO2= 2.1 x 10-1 atm, α-Al2O3 phase was not observed after 2 h of oxidation. Similar results were
observed for as-aluminized samples oxidized at 1000°C in a pO2= 2.1 x 10-1 atm and at a pO2= 1
x 10-5 atm, for which θ-Al2O3 is still the controlling phase in the thermally grown oxide after 5 h
of oxidation. The growth suppression of θ-Al2O3 has been associated with the fact that at low pO2
conditions gas-supply control is established during the early stages of oxidation. This condition is
enhance by the increase in α-Al2O3 nucleation sites as a result of the grit-blasting process.
118
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120
8. Outlook (Future Work)
Despite our efforts, the effect of pre-oxidation parameters and surface conditions (grit-blasting
process) on the competitive growth between θ-Al2O3 and α-Al2O3 is still not fully elucidated.
Additional measurements and analyses that would complement this work are suggested below.
Kinetics Analysis
In this work, we only studied the structural and microstructural properties of the TGO after 2 and
5 h of isothermal oxidation. A set of experiments to study the θ-Al2O3 → α-Al2O3 transformation
as a function of time (during the first hour of oxidation) is strongly recommended to evaluate the
transformation rate for given pre-oxidation parameters (e.g., oxygen partial pressure and
temperature). This information will allow the implementation of more robust kinetics models that
consider the simultaneous growth of θ- and α-Al2O3.
System Degradation Analysis
In order to complement the kinetics analysis in choosing the best pre-oxidation parameters, a study
to understand the effect of the pre-oxidation treatments on BC and SA degradation (including
chemical, structural, and mechanical properties) will be necessary.
121
EL JURADO DESIGNADO POR LA UNIDAD QUERÉTARO DEL CENTRO DE INVESTIGACIÓN
Y DE ESTUDIOS AVANZADOS DEL INSTITUTO POLITÉCNICO NACIONAL, APROBÓ LA
TESIS DOCTORAL DEL (LA) C. (NOMBRE DEL ESTUDIANTE) TITULADA: (“TÍTULO DE LA
TESIS”), FIRMAN AL CALCE DE COMÚN ACUERDO LOS INTEGRANTES DE DICHO
JURADO, EN LA CIUDAD DE QUERÉTARO, QRO., A LOS (FECHA CON LETRA).
Dr. Juan Muñoz Saldaña
Cinvestav-Querétaro
Dr. Luis Gerardo Trápaga Martínez
Cinvestav-Querétaro
Dr. Francisco Javier Espinoza Beltrán
Cinvestav-Querétaro
Dr. Juan Francisco Pérez Robles
Cinvestav-Querétaro
Dr. José de Jesús Perez Bueno
CIDETEQ
Dr. Jesús González Hernández
CIMAV-Unidad Chihuahua
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