Design of a pre-oxidation treatment for a NiCoCrAlY bond coating in view of enhancing the lifetime of high temperature coatings systems

Carlos Serrano Vergel

Master Project Report

Supervisors: Thijs Nijdam

Wim Sloof

Surface and Interfaces

Materials Science and Engineering

Delft University of Technology

Netherlands 2000

CONTENTS

ABSTRACT...........................................................................................................................i

1. INTRODUCTION .......................................................................................................ii

2. BACKGROUND........................................................................................................ 1

2.1. HIGH TEMPERATURE OXIDATION ....................................................................... 1

2.1.1. THERMODYNAMICS ............................................................................................... 1

2.1.2. KINETICS ................................................................................................................. 7

2.2. HIGH TEMPERATURE COATINGS ...................................................................... 11

2.2.1. HIGH TEMPERATURE ENVIRONMENT .............................................................. 11

2.2.2. THE HTC SYSTEM ................................................................................................ 12

2.2.3. SUPERALLOY........................................................................................................ 12

2.2.5. THERMALLY GROWN OXIDE .............................................................................. 15

2.2.6. FAILLURE OF THE SYSTEM ................................................................................ 17

3. EXPERIMENTAL.................................................................................................... 24

3.1. MATERIALS............................................................................................................ 24

3.2. FURNACE............................................................................................................... 24

3.3. REACTION CHAMBER .......................................................................................... 24

3.4. XPS......................................................................................................................... 24

3.5. XRD......................................................................................................................... 24

3.6. SEM ........................................................................................................................ 25

3.7. EPMA...................................................................................................................... 25

3.8. EXPERIMENT FLOW CHART ............................................................................... 25

4. RESULTS ............................................................................................................... 28

4.1. INITIAL CONDITIONS............................................................................................ 28

4.2. 10 MINUTES ISOTHERMAL HEATING IN Ar-H2 .................................................. 30

4.3. ISOTHERMAL HEATING UNDER VACUUM ........................................................ 31

4.4. 10 MINUTES ISOTHERMAL OXIDATION............................................................. 32

4.5. 60 MINUTES ISOTHERMAL OXIDATION............................................................. 39

4.6. ISOTHERMAL OXIDATION AT 800°C .................................................................. 41

4.7. ISOTHERMAL OXIDATION AT 1100°C ................................................................ 44

4.8. SUMMARY OF RESULTS...................................................................................... 53

5. DISCUSSION.......................................................................................................... 55

5.1. STRUCTURE OF THE LPPS NiCoCrAlY BONDCOATING.................................. 55

5.2. EXTERNAL OXIDATION OF ALUMINUM ............................................................. 55

5.2.1. THEMODYNAMICS................................................................................................ 55

5.2.2. KINETICS ............................................................................................................... 57

5.3. NICKEL, COBALT AND CHROMIUM.................................................................... 60

5.3.1. THERMODYNAMICS ............................................................................................. 60

5.3.2. KINETICS ............................................................................................................... 62

5.4. YTTRIUM................................................................................................................ 63

5.4.1. THERMODYNAMICS ............................................................................................. 63

5.4.2. INTERNAL YTTRIUM OXIDATION........................................................................ 64

5.5. SUMMARY.............................................................................................................. 69

5.5.1. THERMODYNAMICS ............................................................................................. 69

5.5.2. KINETICS ............................................................................................................... 70

5.6. TGO CRITERIA FOR AN IMPROVED LIFETIME ................................................. 72

5.7. TGO FORMATION ................................................................................................. 73

6. CONCLUSION: Proposed pre-oxidation treatment ............................................... 79

7. RECOMMENTDATIONS........................................................................................ 80

REFERENCES ................................................................................................................. 81

i

ABSTRACT

High temperature coating (HTC) systems frequently fail when the thermally grown

oxide (TGO) reach a critical thickness. Pre-oxidation treatments can modify the

mature TGO structure and therefore its growth rate. Structural modifications induced

by isothermal oxidation of a NiCoCrAlY bond coating were studied, in order to define

the pre-oxidation conditions that can lead to an extended HTC lifetime.

The initial structure of a NiCoCrAlY bond coating deposited by low pressure plasma

spray (LPPS) is fined grained and meta-stable. Y is evenly distributed among in the

three initial metallic phases present in the BC (e.g. !-(Ni,Co)/!’-(Ni,Co)3(Al,Cr), "-

(Ni,Co)(Al,Cr) and #-(Co)) . As a consequence of the meta-stable nature of the bond

coating structure and the high density of grain/phase boundaries, closed alumina

layers of uniform thickness can be formed and the initial state of Y (i.e. distribution

and bonding) can be swiftly modified upon short periods of isothermal oxidation. The

modification of the state of Y on the oxidation behaviour is two-fold. On the one hand,

the oxide scale growth kinetics is reduced when Y is bonded to oxygen inside the

coating, because a smaller amount of Y-rich oxides (pegs) are included in the scale.

On the other hand, the oxide scale adhesion may be reduced If Y is internally

oxidized, because then Y is not available for the scavenging of impurities in the

coating (e.g. S, C, Ca). Therefore the state of yttrium in the bond coating has a major

effect in the lifetime.

Based on the oxide layer structures formed by isothermal oxidation upon variation of

the oxidation parameters, such as temperature, time and oxygen partial pressure, a

short pre-oxidation treatment is proposed in view of enhancing the lifetime of the high

temperature coating system.

1. INTRODUCTION

ii

1. INTRODUCTION

In gas turbines cycles the exergy of the fuel is converted into shaft work (land based)

or trust power (aero). The open gas turbine cycle has three main steps compression,

combustion and expansion. The three main turbine components are shown in Figure

1. The losses associated with combustion can be reduced by burning the fuel closer

to its adiabatic flame temperature, this implies higher inlet temperature for the turbine

engine. Furthermore, if the ratio between the inlet and outlet temperature of the

expansion cycle is increased, the expansion losses are reduced. Reduction of these

losses means that fuel can be saved [1, 3, 4].

Figure 1: Aeroderivative landbased turbine designed by General Electric

The hot flue gas contains oxygen tied to hydrogen, carbon, sulfur, and nitrogen, plus

solid silicates containing Ca, Mg and Al oxides (CMAS) in the case of aero engines.

This highly aggressive work environment demands: high strength, creep, wear and

oxidation resistance of turbine components [1, 23].

For example, the rotor/stator blades are nowadays made of Ni/Co based superalloys.

After manufacturing these components, the superalloy is heat treated to create the

required strength. This treatment comprises: (i) solution treatment at 1150°C for 2

hours and (ii) precipitation hardening at 850°C for 24 hours. The microstructure

developed after the heat treatment is consists of !-Ni phase (matrix) and !´-Ni3Al

phase (precipitates).

1. INTRODUCTION

iii

To maintain the structural integrity of the Ni/Co based superalloy under the described

environmental conditions, turbine blades are isolated by a ceramic layer and cooled

by air or steam. Figure 2 show the air cooling channels of a turbine blade. The

ceramic layer (TBC) is commonly made up of zirconium oxide stabilized with 7-8 wt%

of yttrium oxide. This material has low thermal conductivity, high strain compliance

and toughness, besides an excellent erosion resistance. However, this layer allows

the penetration corrosive species (e.g. O and S). Hence the superalloy needs to be

protected against corrosion. This is accomplished by forming an oxide layer between

the ceramic coating and the metallic blade.

Figure 2: Turbine blade drawing showing internal cooling channels

Alumina (Al2O3), chromia (Cr2O3) and silica (SiO2) are used as protective oxides.

Chromia scales have the highest growth rate, and tend to from volatile oxy-

hydroxides at temperatures above 800°C [3]. Silica layers have the lowest growth

rate in dry oxygen. However, when water vapor is present in the atmosphere, Si

forms volatile hydroxides [34]. Therefore, under the conditions that occur in gas

turbines, alumina is the best option to from a protective oxide layer eventhough it

offers lower protection against hot corrosion (i.e. attack by sulfur [4]) compared to

chromia or silica [1, 4].

Based on their composition, superalloys have the capability to form a protective

alumina layer, but their ability to maintain it will be the limiting factor, compromising

the mechanical strength of the superalloy in a reduced time span. Hence, to expand

the lifetime of the expensive superalloys used to make the turbine rotor/stator blades,

1. INTRODUCTION

iv

bond coatings (BC) are used. A BC is basically an alloy designed to form a protective

alumina scale by the selective oxidation of Al. Two different types of bond coatings

are employed: diffusion and overlay. Overlay bond coatings are usually composed of

MCrAlRE, where the M stands for Fe, Ni, Co or a combination of Ni and Co and RE

is a reactive element such as: Y, Hf, Ce or Zr. The advantages of overlay over

diffusion coatings are: (i) MCrAlRE bond coating can easily be replace while keeping

the coated component undamaged; (ii) MCrAlRE bond coating composition can be

tailored to specific environmental requirements [3, 15].

The protective Al2O3 layer forms during superalloy precipitation hardening and TBC

deposition and grows during operation. This protective oxide layer is commonly

referred to as the thermally grown oxide (TGO).

The growth of the TGO eventually cause the failure of the high temperature coating

system (HTC=BC+TGO+TBC). When the strain energy (function of the scale

thickness) stored in the oxide layer upon cooling (due to the different thermal

expansion coefficients of the BC and the oxide layer) is larger than the fracture

thoughness of the TGO/BC interface, the oxide layer detaches from the BC.

Therefore, the failure of a high temperature coating system (HTC) is influenced by

the growth rate of the oxide layer and the fracture toughness of the TGO/BC interface

[1, 4].

The reactive elements (RE´s) such as Y and Hf interact with impurities such as S and

C in the BC. Parameters such as: temperature, time and partial oxygen pressure can

be used to modify RE’s and impurities state (bonding and distribution) in the BC. This

modifies the TGO structure and can influence the fracture strength of the TGO/BC

interface.

The aim of the present project is to design a pre-oxidation treatment for a LPPS

NiCoCrAlY BC in view of enhancing the life time of complete high temperature

coating system. Different pre-oxidation treatments were performed on a free standing

as-deposited BC in order to understand the influence of oxidation parameters (e.g.

oxidation time, oxidation temperature and oxygen partial pressure) on the initial TGO

structure. Afterwards, isothermal oxidation experiments were performed for a long

oxidation time in order to study the consequences of the pre-oxidation treatments on

the mature TGO structure. The obtained results were used to design the proposed

pre-oxidation treatment.

1. INTRODUCTION

v

Pre-oxidation treatments performed on an as-deposited BC prior to the super-alloy

precipitation hardening that effectively increased the high temperature coating

system lifetime give a dual advantage: (i) the increase in lifetime reduces overhaul

expenses for the turbine user and (ii) short pre-oxidation times and the possibility to

use a higher pressure set up can reduce the manufacturing cost for the coating

producer.

This document is structured as follows: First, in Section 2, the relevant theoretical

background is presented. Then, in Section 3, a description is given of the

experiments performed and the characterization techniques used. Next, in Section 4,

the obtained results are presented and in Section 5 these results are discussed.

Finally, in Section 6, a pre-oxidation treatment is proposed as a conclusion.

2. BACKGROUND

1

2. BACKGROUND

2.1. HIGH TEMPERATURE OXIDATION

In order to understand how the thermally grown oxide (TGO) is formed during the

turbine operation, it is necessary to mention some fundamental concepts related to

the high temperature oxidation of metals. The fundamental aspects that will be

discussed are: the thermodynamical basis of oxide formation, the kinetics of external

and internal oxidation, and the principle of selective oxidation.

In order to form a metal-oxide molecule two reactions have to take place,

O2 reduction (consumption of electrons):

$$ %+ 22 2

2bObeO

b (1a)

M oxidation (generation of electrons):

$$ +%+ beOMbOaM ba 22 (1b)

The total reaction is shown in equation 1c.

2.1.1. THERMODYNAMICS

Thermodynamical calculations predict whether the reaction can occur in the given

direction. Eq. 1c is an oxidation reaction describing the formation of a pure oxide. a

and b are the number of metal and oxygen atoms involved in the formation of one

metal-oxide (MaOb) molecule, respectively, i.e.:

baOMOb

aM %+ 22

(1c)

The change in standard free energy of formation ( OG& ) is a function of temperature,

if the composition of the metallic phase is constant, the oxide is stoichiometric and

the process occurs at a constant pressure [1]. The standard free energy for the

elements and compounds of interest as function of temperature used in Eq. 2 can be

found in references 36 and 37. Taking into account the previous assumptions, Eq. 2

gives the standard free energy of formation for a metallic oxide [1 and 4].

2. BACKGROUND

2

A plot of the change in standard free energy as a function of temperature aids in the

analysis of pure and mixed oxide formation. Reactions with a large negative standard

free energy of formation, compared to those with a small negative standard free

energy of formation, are energetically favorable since the magnitude of their driving

force is larger (Graphs 1-4).

( ) ( ) ( ) ( )

+$=& TGb

TaGTGTG OO

OM

OOM

O

ba 22 (2)

Graph 1 shows the standard free energy of formation of the relevant oxides, it can be

seen that Y and Al are the more reactive components of the NiCoCrAlY coating,

while Ni and Co are the more noble ones. In Graph 2, it can be seen that formation of

NiAl2O4 (spinel) by solid state reaction has a much less negative standard energy of

formation than formation by solute incorporation to a pure oxide or by the reaction

between pure elements. Graphs 3 and 4 show the same comparison for YAlO3 and

Y3Al5O12, respectively.

-1500

-1000

-500

0

500 1000 1500

Temperature (oC)

&G

O (

kJm

ol-1

) 2Cr+1.5O2=Cr2O3

2Al+1.5O2=Al2O3

Ni+0.5O2=NiO

Co+0.5O2=CoO

2Y+1.5O2=Y2O3

Graph 1: Plot of the standard free energy of formation as a function of temperature for pure metals. Data

taken from references 36 and 37.

2. BACKGROUND

3

-1500

-1000

-500

0

500 1000 1500

Temperature (°C)

&G

O (

kJm

ol-1

)

-50

-40

-30

-20

-10

0

&G

O (

kJm

ol-1

)

Ni+2Al+2O2=NiAl2O4 NiO+2Al+1.5O2=NiAl2O4 NiO+Al2O3=NiAl2O4

Graph 2: Plot of the standard free energy of formation as a function of temperature for NiAl2O4. The right

side axis applies to the solid state reaction only. Data taken from references 36 and 37.

-1500

-1000

-500

500 700 900 1100 1300 1500

Temperature (°C)

&G

O (

kJm

ol-1

)

-50

-45

-40

-35

-30

-25

&G

O (

kJm

ol-1

)

Y+Al+1.5O2=YAlO3 0.5Y2O3+Al+0.75O2=YAlO3 0.5Al2O3+0.5Y2O3=YAlO3

Graph 3: Plot of the standard free energy of formation as a function of temperature for YAlO3. The right

side axis applies to the solid state reaction only. Data taken from references 36 and 37.

2. BACKGROUND

4

-1500

-1000

-500

0

500 1000 1500

Temperature (°C)

&G

O (

kJm

ol-1

)

-250

-200

-150

-100

-50

0

&G

O (

kJm

ol-1

)

3Y+5Al+6O2=Y3Al5O12 1.5Y2O3+5Al+3.75O2=Y3Al5O12

2.5Al2O3+1.5Y2O3=Y3Al5O12 Al2O3+3YAlO3=Y3Al5O12

Graph 4: Plot of the standard free energy of formation as a function of temperature for Y3Al5O12. The

right side axis applies to the solid state reaction(s) only. Data taken from references 36 and 37.

The thermodynamical analysis for a mixed-oxide is analogous to the one explained

for pure oxides and follows the same assumptions. For mixed-oxides Eq. 1 is then

modified by simply adding a second reactant, which can either be a metal or a metal-

oxide.

Three different formation routes are plotted for the formation of mixed oxides (cf.

Graphs 2, 3 and 4): (i) solid state reaction between pure oxides, (ii) reaction between

pure metals and oxygen and (iii) solute incorporation into a pure oxide.

The change in the standard free energy can also be written as follows [4]:

( ) ( )KRTTGO ln$=& (3)

where R is the ideal gas constant kJmol-1K-1, T is the temperature in Kelvin and K is

the reaction constant, which in this case equals:

2222

1b

O

b

OaM

OM

ppa

aK ba == (4)

2. BACKGROUND

5

An oxidation reaction can occur if the standard free energy of formation is negative

and the oxygen pressure is above the equilibrium oxygen partial pressure [4]. If the

thermodynamic activities are assumed to be unity, the reaction constant K is

inversely proportional to the equilibrium oxygen partial pressure (2Op ). Then the

oxygen partial pressure equals:

( )

&=

bRT

TGp

O

O

2exp

2 (5)

If the equilibrium oxygen partial pressure is plotted versus 1/T the slope of the line is

proportional to 2&GO/bR. For the relevant oxidation reactions, these relations are

plotted in Graph 5. In this graph it can be seen that the oxidation of metals with less

negative standard free energy of formation have a higher equilibrium pressure.

Furthermore, the equilibrium oxygen partial pressure increases as temperature

increases. Thus, in general, a metal-oxide is less stable at higher temperature

(Graphs 5-8).

1.E-150

1.E-100

1.E-50

1.E+00

0.5 1 1.5 2

1x103/Temperature (K)

log (

pO

2)

2Cr+1.5O2=Cr2O3

2Al+1.5O2=Al2O3

Ni+0.5O2=NiO

Co+0.5O2=CoO

2Y+1.5O2=Y2O3

Graph 5: Plot of the oxygen partial pressure of equilibrium as a function of temperature for pure oxides.

Data taken from references 36 and 37.

For the specific case of NiCoCrAlY alloy oxidation, Graph 5 indicates that if a close

aluminum oxide layer is formed, underneath this layer only Y can be oxidized.

Graphs 6, 7 and 8 show that less oxygen is required to from a mixed oxide by solute

2. BACKGROUND

6

incorporation into a pure oxide compared with the reaction of pure metals with

oxygen.

1E-100

1E-50

1

0.5 1 1.5 2

1x103/Temperature (K)

log(p

O2)

Ni+2Al+2O2=NiAl2O4 NiO+2Al+1.5O2=NiAl2O4

Graph 6: Plot of the oxygen partial pressure of equilibrium as a function of temperature for NiAl2O4. Data

taken from references 36 and 37.

1E-100

1E-50

1

0.5 1 1.5 2

1X103/Temperature (K)

log(p

O2)

Y+Al+1.5O2=YAlO3 0.5Y2O3+Al+0.75O2=YAlO3

Graph 7: Plot of the oxygen partial pressure of equilibrium as a function of temperature for YAlO3. Data

taken from references 36 and 37.

2. BACKGROUND

7

1E-100

1E-50

1

0.5 1 1.5 2

1x103/Temperature (K)

log(p

O2)

3Y+5Al+6O2=Y3Al5O12 1.5Y2O3+5Al+3.75O2=Y3Al5O12

Graph 8: Plot of the oxygen partial pressure of equilibrium as a function of temperature for Y3Al5O12.

Data taken from references 36 and 37. 2.1.2. KINETICS

EXTERNAL OXIDATION

Figure 3: Stages of external alloy oxidation.

2. BACKGROUND

8

The kinetics of oxide scale growth by external oxidation of a metallic component is

usually displayed as the oxide scale thickness versus oxidation time. Figure 3 shows

the three stages identify in the external oxidation of alloys.

Initially the oxide layer is thin and contains a high density of grain boundaries [4].

Then, the diffusion through the oxide layer is fast and the kinetics is determined by

the absorption and dissociation of oxygen. The growth of the oxide layer at this stage

is usually linear with oxidation time t [4] :

tkx LTMO = (6)

where TMOx is the oxide layer thickness (transient stage) and kL is the linear growth

rate.

The following analysis applies (steady state) to an external oxide layer growing by

inward diffusion of atomic oxygen through grain boundaries (cf. Scheme 1).

Scheme1: Growth of an external oxide layer

The flux of oxygen atoms through the MO/M interface can be written in two

analogous ways. (i) Einstein’s general approach to describe forced diffusion. In this

case the flux of oxygen atoms is the product between the average migration velocity

of the oxygen atoms OMO

OU m/s through the MO/M interface and the oxygen

concentration at the same interface MMO

OC in at% and (ii) Fick’s first law in terms of

the chemical potential gradient [4, 59]:

MMO

OM

MO

OO

MOO

MMO

OMMO

O UCdx

d

RT

DCJ =$=

µ (7)

In Equation 7, MOOD is the diffusion coefficient of oxygen ions through the oxide layer

in m2/s. If the existing oxide layer has a thickness equal to xMO and local

thermodynamic equilibrium is assumed at the MO/O2 and the MO/M interfaces, then:

2. BACKGROUND

9

MMO

OM

MO

O

MO

MMO

OO

MO

OMOO

MMO

OMMO

O UCxRT

DCJ =

$

$=

µµ 2

(8)

The average migration velocity of the oxygen atoms through the MO/M interface

OMO

OU equals the growth rate of the oxide layer dxMO/dt. Because, once an oxygen

atom crosses this interface it reacts to form MO becoming part of the oxide layer,

then:

MO

MMO

OO

MO

OMOOMOM

MO

OxRT

D

dt

dxU

µµ $==

2

(9)

If the initial conditions state that at t=0 the oxide layer thickness is zero, then:

tktRT

DxP

MMO

OO

MO

O

MOOMO =

$= µµ 2

2

2

(10)

The coefficient for oxygen diffusion through the scale can be related to the coefficient

for oxygen diffusion along the oxide layer grain boundaries by considering the grain

boundary width ' and the grain size r :

tktRT

D

r

xP

MMO

OO

MO

O

GBOMO =

$= µµ'

22

2

2

(11)

Equation11 states that for thermodynamic equilibrium at the interfaces and if the

oxide growth is controlled by inward oxygen diffusion, the oxide layer thickness is

proportional to the square root of time.

The parabolic constant kP can be determined from measurements of the oxide layer

thickness for a given oxidation time using Eq. 11 and has units of m2s-1. Alternatively,

the mass increment &m can be measured upon oxidation and a parabolic growth

constant is deduced from Equation 12, according to the mass change:

tkA

m mp=

&

(12)

where A is the area over which the reaction occurs. Then, the parabolic rate constant

calculated from mass gain has units of g2cm-4s-1. The relation between the two

constants in given by Equation 13:

2. BACKGROUND

10

( )mP

O

OMO

p kw

nvk

2

2

21

= (13)

where vMO is the molar volume of the oxide, nO is the oxygen valence and wO is the

atomic mass of oxygen.

After the element forming the protective oxide layer is consumed, the oxide layer is

no longer self-healing, this leads to a rapid degradation of the bond coating which is

known as breakaway oxidation [1, 4 and 55].

INTERNAL OXIDATION

The internal formation of oxides in an alloy occur, when the oxygen partial pressure

in the alloy is high enough to oxidize one or more alloy components that are present

in concentrations below the limit for the transition to external oxidation [4]. Internal

oxidation form oxide precipitates (see Scheme 2), in contrast with the continuous

layer formation by external oxidation.

Scheme 2: Internal oxidation of metal Z bellow an external MO layer

The following analysis of internal oxidation can be found in appendix B of Ref. 4. If

the oxidation front is said to be controlled by atomic oxygen diffusion, then the

penetration depth of the internal oxidation front (i.e. distance between the MO/M

interface and the I.O.F.) xIOF can be written as:

tDx MOIOF !2= (14)

2. BACKGROUND

11

where MOD is the effective oxygen diffusion coefficient in the metal and ! is an

undetermined factor that gives the proportionality between xIOF and the characteristic

diffusion length tDMO . Fick’s second law applies if

( )0($=

dt

dC

dx

xdJ (15)

hence, Fick’s second law is the first derivate of the first law with respect to position x,

assuming that the diffusion coefficient is constant with respect to x. For oxygen

diffusion:

2

2

x

CD

t

C OO

O

)

)=

)

) (16)

where CO is the concentration of oxygen.

Initial conditions at=0:

x=0 MMO

OO CC = concentration of oxygen at the MO/M interface

x>0 0=OC No oxygen is dissolved in the metal

Boundary conditions at t=t:

x=0 MMO

OO CC =

x=xIOF 0=OC

The solution of Fick’s second law under the mention conditions is:

( )( )

$=!erf

tD

xerf

CxCMOM

MO

OO

21 (17)

Next, Fick’s second law is applied to the solute Z that is being internally oxidized.

Initial conditions at t=0:

2. BACKGROUND

12

x<0 0=ZC no metallic Z above the MO/M interface

x>0 *= ZZ CC concentration of Z away from the I.O.F.

Boundary conditions at t=t:

x=0 0=ZC no metallic Z at the MO/M interface

x= * *= ZZ CC

Then the solution of Fick’s second law is

( )

+

$

$=

*

*

!21

21

erfc

tD

xerfC

CxC

MZ

M

ZZ (18)

Where + = MOD / M

ZD . The internal oxidation kinetics (i.e., !) are obtaining from writing

the flux balance at the internal oxidation front (i.e. x=xIOF). At this interface the flux of

oxygen equals , (i.e. , is the ratio between oxygen and Z atoms in one molecule of

the formed oxide) times the flux of Z atoms, since it is assumed that at the I.O.F. the

concentration of metallic Z and oxygen atoms drop to zero. The flux balance yields

to:

( ) ( )( )2

2

21

21 exp

exp

!

!!,

! +

++

=*

erfc

erf

C

C

M

MMO

O (19)

An analytical solution of Eq. 19 can be obtain for significant counter diffusion of metal

Z using the following approximations 1<<! and 121

<<+! . Since

( ) !-

!2

=erf (20)

121

.

+!erfc (21)

( ) ( ) 1expexp 22 .+. !! (22)

2. BACKGROUND

13

Substituting Equations 20, 21 and 22 into 19 leads to Eq. 23:

*

=Z

MMO

OMZ

MO

C

CD

D

,

-

!2

21

(23)

( )tkt

CD

CDx

p

ZZ

MMO

OOIOF int

22

22

2

22=

=*,

- (24)

Eq. 24 is obtain by substituting Eq. 23 into Eq. 14 and states that the penetration

depth of the internal oxidation front xIOF is proportional to the square root of time.

SELECTIVE OXIDATION

As was shown in Section 2.1.1, the driving force for oxidation of metallic elements

depends on the standard free energy of formation. Thus, the selective oxidation

depends on the oxygen partial pressure, the temperature and the thermodynamic

activity of the metal in the alloy. Furthermore, the diffusitivity of the elements in the

alloy, which in turn depend on aspects of the alloy microstructure and composition,

have a strong influence on the kinetics.

All these parameters allow to control the selective oxidation of an alloy component in

order to form a protective oxide layer. However, this process is complicated

(impaired) by, formation of transient and mixed oxides and internal oxidation; cf.

Sections 2.1.1 and 2.2.5.

Selective oxidation of elements that form a slowly growing stable oxide layer is the

basis for corrosion/oxidation protection of alloys used at high temperature [4]. Three

elements are considered to form protective scales, Si, Al and Cr. From these three,

the best protection is provided by Al when the temperature exceeds 800°C and water

vapor is present [1].

2.2. HIGH TEMPERATURE COATINGS

The conditions under which high temperature coating (HTC) operate, the

components comprising the HTC system and aspects related to the failure of the

HTC system, will be discussed in this section.

2. BACKGROUND

14

2.2.1. HIGH TEMPERATURE ENVIRONMENT

The operation environment of turbine engines is characterized by: high temperature

and pressure, plus the presence of oxidizing species and solid particles (aero) in the

atmosphere [1, 2].

Due to these operation conditions, materials used in turbine components need to

withstand: thermally induced fatigue, creep damage, oxidative/corrosive degradation

and erosive wear. These degradation processes limit the operation temperature, and

hence have an impact on the power generation efficiency of the turbine [3].

The enhancement of high temperature strength of super-alloys has lead to a

reduction in their oxidation/corrosion resistance [1]. In order to produce components

that can withstand more severe conditions for longer times, coating systems that

provide thermal insulation and enhanced oxidation/corrosion resistance are required.

2.2.2. THE HTC SYSTEM

The function of the HTC system is to provide protection to the super-alloy against the

operation environment. There have bean three fundamental developments

concerning the materials used in high temperature environments [23,4]:

(I) Optimization of superalloy high temperature strength and the coating

corrosion/oxidation resistance.

(II) Improve casting technologies that have lead to the production higher creep

resistance components.

(III) Development of technologies that allowed coating of complex geometry

components.

The high temperature coating is composed by a series of layers of different materials

with different functions see Figure 4.

The top layer is the thermal barrier coating (TBC). This is a ceramic layer made of

Zr2O3 stabilized with Y2O3. The function of the TBC is to thermally insulate the

underlying layers and the substrate from the environment.

The second layer is also a ceramic and is composed mainly of Al oxide. This layer is

not deposited, but it is formed and grows during the deposition of the TBC and during

2. BACKGROUND

15

operation. The function of the thermally grown oxide (TGO) is to prevent the

oxidation of the underlying substrate.

Figure 4: Schematic illustration of the High temperature coating (HTC).

The third layer shown in Figure 4 is the bond coating (BC). This metallic alloy, is

corrosion/oxidation resistant due to its ability to (re)form the TGO under operation

conditions. The two main functions of the BC are to maintain together the TBC and

the superalloy and to serve as a reservoir of Al.

2.2.3. SUPERALLOY

Figure 5: Nickel aluminum phase diagram. Taken from Ref. 1

2. BACKGROUND

16

Superalloys have evolved over the years to optimize high temperature strength,

creep and fatigue resistance at the expense of corrosion/oxidation resistance. This

has led to the development of nickel based superalloys that are able to operate at

80% of their melting point [1].

Superalloys are essentially a ! -Ni matrix hardened by !’-Ni3Al precipitates. As

indicated in Figure 5, both phases have face centered cubic structure. The slow

growing cubic (coherent) !’-Ni3Al precipitates are formed by thermal treatment of the

superalloy, commonly after bond coating deposition. This heat treatment comprises:

(i) solution at 1150°C for 2 hours and (ii) precipitation at 850°C for 24 hours [1].

2.2.4. THE BOND COATING

The bond coating (BC) connects the ceramic TBC to the superalloy and serves as an

aluminum reservoir. In order for an alloy to be used as a bond coating, it must be

able to: (i) form a thermodynamically stable, protective and slow growing oxide scale,

(ii) avoid phase transformations in the BC during operation, (iii) accommodate the

strains experience by the TBC and the superalloy upon thermal cycling, (iv) endure

creep damage and thermal fatigue.

In order to comply with the first requirement, sufficient Al has to be supplied to the

TGO. Oxide maps such as the one shown in Figure 6, have been constructed for

different temperatures and alloy systems, in order to aid the design of alloys that form

protective TGO’s.

There are two types of bond coatings: diffusion and overlay. In both cases the idea is

to increase the aluminum concentration at the component surface.

Diffusion coatings can either be formed by inward aluminum diffusion into the super-

alloy or by outward nickel diffusion. Ni-Al compounds have poor solubility of other

elements. Hence, there is a strong interdependence between the bond coating and

the super-alloy composition. This limits both oxidation/corrosion resistance and the

mechanical properties of diffusion bond coatings [1,3].

Overlay coatings give the possibility of tailoring the properties of the BC to the

environmental conditions with much less restrictions as compared to diffusion

coatings. Overlay coatings are comprised of a !-Ni and/or "-NiAl phase(s), and their

chemical components are: MCrAl-RE, where M can be iron, nickel, cobalt or a

combination of Ni and Co and RE is a reactive element such as Y, Hf, Ce or Zr [1].

2. BACKGROUND

17

Figure 6: External and internal oxides formed after isothermal oxidation of NiCrAl alloys. Taken from

Ref. 1.

There are two widely used techniques for the deposition of overlay BC’s: spraying

and evaporation. Plasma spray can be carried out with high or low pressure. When

the process takes place at low pressure (LPPS) some technical advantages can be

obtained, such as:

(I) The amount of Al loss by oxidation is reduced.

(II) Since the residence time is longer, powders attain higher temperatures. In

addition there is no convective cooling, hence the surface temperature of the

substrate is higher. Both these conditions aid to reduce BC porosity.

(III) Bond strength is enhanced due to higher particle velocity.

ROLE OF DIFFERENT ALLOYING CONSTITUENTS IN THE BC

Every element of an NiCoCrAlY alloy has an specific function. Ni provides the

strength. Co provides strength plus thermal stability, i.e. a NiCrAl alloy has a "-NiAl/!-

Ni structure at the operating temperature, but upon cooling "$NiAl transforms into

/$NiCr and !$Ni transforms into !’-Ni3Al. Adding Co to the alloy prevents these

transformations [1].

2. BACKGROUND

18

Cr improves hot corrosion resistance and promotes the selective oxidation of Al by

reducing the oxygen diffusivity in the BC. A drawback is that Cr has a negative effect

on the alloy creep resistance [1, 4].

Low additions of Y, improve the adhesion between the alloy and the oxide scale.

These elements prevent segregation of impurities such as S and C to the BC/TGO

interface. The mechanism is not well understood. However, it has been observed,

that Y-oxides dissolve at high temperature and metallic Y segregates to the surface,

implying that Y-sulfides should dissolve also [2, 14, 29]. Another relevant observation

is that Y-pegs can mechanically enhance BC/TGO fracture toughness [1, 30].

Finally, the function of Al is to improve oxidation resistance, by forming a protective

and slow growing alumina layer. In addition, it increases the coating toughness and

refines the coating microstructure. Microstructure refinement increases the number of

Al diffusion paths to the surface, enhancing the formation of the alumina scale. A

drawback is that, with increasing aluminum content the melting point of the alloy

decreases [1, 6, 31].

2.2.5. THERMALLY GROWN OXIDE

Generally, the thermally grown oxide forms while other steps on the component

manufacturing process are in progress [1]. Since the growth rate of the TGO controls

the HTC durability [3, 20-22, 30], more attention is now paid to the initial structure of

the TGO and pre-treatments are design to alter the mature TGO structure.

TGO FORMATION

The structure of the TGO depends on the oxidation temperature, oxidation time and

the oxygen partial pressure. As the temperature increases process kinetics are

enhanced and the stability of oxides is reduced (increased oxygen equilibrium

pressure), allowing selective oxidation of the more reactive elements to take place in

a controlled manner. Diffusion of metal cations and oxygen anions takes time, hence

longer oxidation time leads to thicker oxide layer. At lower partial oxygen pressure,

oxide formation is hindered allowing Cr evaporation and Y migration to the surface

[21].

2. BACKGROUND

19

Figure 7: Oxide phase present in the TGO.

The presence of a native oxide lowers the oxygen partial pressure underneath this

layer. Hence, its presence can reduce Cr evaporation and Y migration towards the

surface [21]. Figure 7 shows a schematic representation of the TGO structure.

TRANSITION ALUMINAS

Transition aluminas (! and 0) form epitaxial films on "-NiAl. The orientation relation

makes their nucleation energetically favorable over the randomly oriented alpha

alumina at temperatures below 1100°C [9]. The precise temperature depends on the

alloy composition.

It has been observed that the orientation relation between the transition aluminas and

the substrate is lost, after the transformation to /-alumina [7]. The 0$alumina to

/$alumina transformation is enhanced as temperature increases [17]. Also, as

oxidation time increases, 0-alumina transforms into /-alumina. Then, the presence of

transition aluminas (!, ' and 0) is only reported at the early oxidation stages for

temperatures bellow 1100°C [7-9].

The transformation sequence !-'-0-/ occurs for pure alumina and is studied by

differential scanning calorimetry [35]. However, this transformation sequence does

not occur upon isothermal oxidation of NiAl as shown by Pint et al. [27].

There are two basic structural differences between the /-alumina and the transition

aluminas: (i) /-alumina has an hcp oxygen sublattice, while transition phases exhibit

a slightly disordered fcc sublattice, (ii) In /-alumina, the Al cations occupy only

2. BACKGROUND

20

octahedral sites, while in transition aluminas, they are distributed between octahedral

and tetrahedral sites.

The ratio between octahedral to tetrahedral occupied sites (i.e. a spinel oxide has a

ratio of 2) in transition aluminas is as follows: for ! and ' aluminas this ratio is

between 1 and 2, and in 0-alumina is exactly one. The occurrence of tetrahedrally

coordinated Al cations is the reason for the stabilization of the fcc oxygen sublattice

[10, 12].

Grabke et al. concluded that transition aluminas have a large component of outward

growth due to the high diffusivity of Al cations, while the /-alumina grow by inward

diffusion of oxygen in NiAl and NiFeAl alloys [9]. Due to the outward growth of

transition aluminas the oxide exhibits a needle like surface morphology. This does

not mean that the surface morphology of the transitional phase can be used as an

indicator of its presence, but it can be considered as evidence of past existence. It

also has been seen that the alumina transformation from 0 to / is fast when it grows

on top of the !’-Ni3Al phase and slow when it grows on top the "-NiAl phase. This is

said to be due to the larger amount of Al, which gives the " phase the ability to

sustain outward growth for longer times [20].

De Wit et al. [32] observed that the addition of Y to "-NiAl delays the transformation

from 0$alumina to /$alumina, since it retards oxygen diffusion and at the same time

it enhances Al diffusion through the oxide layer, although this theory is not proven as

discussed by Pint [33]. Grabke showed that the presence of Cr enhances the growth

rate of /-alumina, but has little effect on the growth of 0. He also observed that as the

content of Cr increases, the transition time from 0 to / decreases [8].

Rhüle et al. [27] formed a !-alumina layer by isothermal oxidation at 800°C and

observed the transformation to /-alumina without formation of 0-alumina, while

Doychak formed oxide layers exclusively composed of 0-alumina on top of the same

substrate after ten hours of isothermal oxidation at 800°C. The difference between

the experiments is attributed to what can be called a ‘nucleation step’; Rhüle oxidized

his samples for ten minutes at 870°C, while Doychak oxidized his samples at 1100°C

[28].

The presence of transitional aluminas (!, ', 0) can promote the formation larger

alumina grains as observed by Nijdam et al. [30]. This can reduce the mature TGO

2. BACKGROUND

21

growth rate, since diffusion of atomic oxygen through grain boundaries is considered

the main transport mechanism.

SPINEL OXIDES

(Ni,Co)(Al,Cr)2O4 spinel forms on top of the /-Al2O3 layer, due to the lack of

aluminum to sustain the exclusive formation of alumina. Initially, pure oxides are

formed and is commonly argued that, the following formation of mixed oxides occur

by solid state reactions between NiO or CoO and Al2O3 or Cr2O3 [40].

Spinel formation has been studied by Pettit [43] at 1200-1500°C in diffusion couples

of polycrystalline samples of pure NiO and Al2O3. The kinetics of the transformation is

described by a parabolic relation.

Hence, spinel formation by solid state reaction is a diffusion controlled process. The

parabolic growth constant is largely dependent on temperature at which the reaction

takes place, but independent from oxygen partial pressure. The growth constant

increases by three orders of magnitude from 1200 to 1400°C and is controlled by

aluminum diffusion in this range [43].

YTTRIUM ALUMINATES

Y2O3 is the first oxide to form since Y is the most reactive element. After Y, Al is

oxidized and a closed alumina layer forms. As the process continues both of

elements (Al and Y) are driven towards the surface by the change in free energy

associated to oxide formation. The formation of Y-aluminates at the BC/TGO

interface is commonly explained by the solid state reaction between Y and Al oxides.

When the driving force decreases or the kinetics are reduced (i.e. by reducing the

temperature), less segregation of Y to the surface takes place and consequently the

density of Y-aluminates at the BC/TGO interface decreases. When the formation of

the protective alumina layer is hindered at low oxygen pressure, high Y migration to

the surface is observed.

Impurities such as S and C may segregate to the BC/TGO interface, internal pores

surfaces and grain boundaries. The idea is that, by segregation to these locations,

the impurities have the opportunity to reduce their surface energy. Y prevents the

detrimental effects of the impurities by preventing their migration to the BC/TGO

interface. If S segregates to the BC/TGO interface, it decreases the TGO adhesion,

2. BACKGROUND

22

because the interfacial covalent-ionic Ni-O bonds are replaced by weaker ionic-

covalent S-Al bonds [21, 22, 30 and 49].

2.2.6. FAILLURE OF THE SYSTEM

Lifetime is determined by intrinsic and extrinsic processes. Extrinsic degradation

processes are related to particle impact and high temperature corrosion. Intrinsic

degradation processes arise from misfit strains between the different layers.

The misfit strains develop due to differences in the thermal expansion coefficient as

well as from phase transformations and interdiffusion.

The TBC can easily accommodate misfit strains because of its porosity. This strategy

can not be used for the BC or the TGO because to serve their function they need to

be dense [49]. No large difference exists between the thermal expansion coefficients

of the BC and the substrate. Only between the TGO and the BC an important

difference between thermal expansion coefficients exists. Therefore, failure is most

probable to occur along the TGO/BC interface.

The oxide layer grows at high temperature and at the oxidation temperature only

growth stresses are present. The growth stresses in an oxide layer are related to: (i)

volume differences between the oxide layer and the alloy from which it forms, (ii)

epitaxial relations between oxide layer and alloy, (iii) compositional changes in the

alloy or oxide layer, (iv) recrystallization and (v) mixed-oxide formation within the

oxide layer.

If the volume difference is assumed to be the main source of stress, then the sign of

the stress is given by the ratio between the molar volume of the oxide and the molar

volume of the metal vMO/vM. If this ratio, the so called Pilling-Bedworth ratio is larger

than 1, then compressive stresses develop in the oxide scale. This applies when the

oxide layer grows by inward oxygen diffusion. When the oxide layer grows by

outward diffusion no growth stresses result from this mechanism [4].

Upon cooling the thermal contraction of the metallic BC is larger than the thermal

contraction of the oxide layer. Hence, the thin oxide layer is forced to match the BC

reduction. Forced contraction of the oxide generates inplane compressive stresses.

The Young´s modulus EMO and the Poison´s ratio ,MO of alumina are 340 GPa and

0.28, respectively and the difference between the thermal expansion coefficient of

2. BACKGROUND

23

the BC (/M) and the TGO (/MO) equals -8x10-6. The stress in the oxide layer follows

from:

( )( )TE

MMO

MO

MOMO &$

$= //

,#

1 (25)

This inplane stress is around -4 GPa for the relevant temperature change

( .&T 1000°C). Consequently, although the oxide layer may be relatively thin at the

end of the life cycle ( critMOx = 5-7 µm), the energy store in it is large.

As the oxide layer grows, the elastic strain energy per unit of interface area G stored

in this layer increases, since it is a function of the oxide layer thickness [49]:

( )MOMO

MO

MO xE

G 21#

,$= (26)

In Eq. 26 MOx is the scale thickness. Delamination of the TGO occurs when the

strain energy per unit of interface area (Jm-2) exceeds the TGO/BC interface

toughness. At a thickness of 6 µm the stored energy is in the order of 200 Jm-2 [49

and 55].

Intrinsic failure mechanisms have a characteristic critical TGO thickness critMOx , which

depends on BC composition and microstructure as well as on thermal cycling history.

3. EXPERIMENTAL

24

3. EXPERIMENTAL

A description of the experiments performed and the characterization techniques used

to study the changes that had taken place upon oxidation of the bond coating are

described in the following sections.

3.1. MATERIALS

A Ni-23Co-17Cr-12.5Al-0.3Y (wt %) bond coating was plasma sprayed at low

pressure (LPPS) on top of stainless steel substrate. The bond coating with a

thickness of 0.7 mm was removed from the substrate and cut into disks with a

diameter of 10 mm using spark erosion.

Samples were grinded using every SiC paper until P4000 roughness, cleaned

ultrasonically in ethanol and dried with compressed nitrogen prior to the oxidation

treatments.

3.2. FURNACE

The samples were isothermally oxidized in a horizontal alumina tube furnace (Lenton

PTF 16/75/610). The oxidation treatment was carried out at a constant oxygen partial

pressure of 101 kPa, obtained by blowing pure oxygen (99.998 vol% O2) with a flow

rate of 100 ml/min through an alumina tube with an inner diameter of 100 mm. A heat

treatment was also performed under Ar with 5 vol.% H2. The gas mixture flow rate

was again 100 ml/min.

3.3. REACTION CHAMBER

Low oxygen partial pressure experiments were performed in a Ultra High Vacuum

chamber with a base pressure below 1x10-7 Pa, coupled to an instrument for XPS

analysis (base pressure below 1x10-8 Pa), cf. Section 3.4.

The samples were prepared as described for the high pressure experiments (cf.

Section 3.1.). A thermocouple was spot welded to the bottom of the sample in order

to monitor the specimen temperature. The segregation experiments were carried out

at the reaction chamber base pressure.

3. EXPERIMENTAL

25

3.4. XPS

The elements present at the oxide surface were identified by X-Ray Photoelectron

Spectroscopy (XPS), using non-monochromated Mg K/ radiation (1253.6 eV).

Spectra were recorded with a PHI 5400 ESCA instrument. The pass energy of the

spherical capacitor analyzer was set as 35.75 eV for, Al, Cr, Ni, Co, C and O and to

71.55 eV for, Y and S. the energy was incremented in 0.2 eV steps and the detection

angle was set to 45° with respect to the sample surface.

From the recorded spectra, the positions of the peaks were determined after satellite

subtraction and charge correction by repositioning the carbon peak at 284.80 eV [57].

Photoelectron peak positions were determined by curve fitting to a Gauss-Lorentz

distribution and using an iterated Shirley background subtraction [58]. Auger

electrons peak positions were determined from the second derivate of the recorded

spectra [58].

3.5. XRD

X-Ray Diffraction (XRD) measurements were performed using a BRUKER AXS D8,

diffractometer equipped with a Co anode. These measurements were used to identify

the crystalline phases present.

The divergent beam emitted by the X-ray source was converted into a parallel beam

(laminar) by using a polycap. The angle of incidence 0 was fixed to 5° while the gas

field detector 20 range was rotated from 10 to 110° at 0.034° 20 steps.

The recorded diffractogram was plotted as total counts versus the inter-planar

distance d and analyzed using the ICDD database [PDF2].

3.6. SEM

The morphology, structure and thickness of the oxide layers were studied using a

JEOL JSM 6500F Scanning Electron Microscope (SEM), equipped with an Autrara

backscatter electron detector and a Noran-Pionner 30 mm2 Si(Li) detector for Energy

Dispersive X-Ray Spectroscopy (EDS).

The cross-sections of the samples were prepared by coating the oxide layer surface

with a sputtered thin film of gold. Then, the sample was embedded in a epoxy resin

and cut perpendicular to the scale surface with a diamond saw. The cross-section

surface was grinded with every SiC paper until P4000 roughness. After grinding, the

3. EXPERIMENTAL

26

samples were cleaned ultrasonically in ethanol and dried with compressed nitrogen.

To obtain high resolution images, a second thin film of gold was sputter onto the

cross-sections, next the cross sections were polished with a JEOL precision ion

polishing system for 16 hours, using a 5 keV Ar+ beam. The error in the microscope

magnification was less than 3% as determined with a calibration sample.

3.7. EPMA

Electron Probe X-Ray Micro-Analysis (EPMA) was carried out in regions of 100x50

µm containing the oxide layer and the under laying substrate using a 2x2 µm grid.

The measurements were performed with a JEOL JXA 8900R microprobe using an

electron beam with energy of 15 keV and a beam current of 50 nA employing

Wavelength Dispersive Spectrometry (WDS) for Ni, Co, Al, Cr, Y, Fe and Si.

Carbon and oxygen were measured separately by recording the full spectra of the C-

K/ and O-K/ peaks using multi-layered analysing crystals (LDEC or LDE1). The

electron beam had an energy of 15 keV and a beam current of 500 nA.

The composition at each analysis point of the selected area was determined using

the X-Ray intensities of the constituent elements after background correction relative

to the corresponding intensities of reference materials. The obtained intensity ratios

were processed with a matrix correction program CITZAF [39]. For C and O the last

procedure was performed by determining the area of the spectra recorded. As

standard materials Fe3C and YAlO3 were used for calibration of the C-K/ and O-K/

intensities, respectively.

The data from the quantitative measurement and the spectroscopy were recalculated

with the PROZA96 algorithm of Bastin and Heijligers [39] taking the experimental

intensity ratios as input.

3.8. EXPERIMENT FLOW CHART

A summary of the oxidation experiments performed with the free standing as-spayed

NiCoCrAlY bondcoating (cf. Section 3.1.) is presented in Flow chart 1.

The experiments were performed at low pressure inside a high vacuum chamber and

at atmospheric pressure in the tube furnace. Low pressure experiments were

performed for the oxidation times listed in the flow chart. After reaching the oxidation

time, the surface was analysed using XPS.

3. EXPERIMENTAL

27

Atmospheric pressure experiments were performed in oxidizing and in reductive

atmospheres, cf. Section 3.3. The treatment in the reductive atmosphere was

intended to reveal the changes induced by heating without sample oxidation.

Experiments at high oxygen partial pressure were performed at four different

temperatures for ten minutes. All sample were then oxidized for 16 hours at 800°C

and then for 96 hours at 1100°C. By performing the described oxidation treatments,

the effects that short term pre-oxidation treatments have on the mature oxide scales

can be studied.

Flow chart 1: Oxidation experiments performed with a free standing as sprayed NiCoCrAlY bond coating

alloy.

4. RESULTS

28

4. RESULTS

The experimental results will be presented as follows: first, the initial conditions of the

NiCoCrAlY bond coating. Next, the changes of the BC microstructure induced by

heating in a non-oxidizing atmosphere. Finally, the oxide formation and

microstructure development upon isothermal oxidation in order of increasing

oxidation time.

4.1. INITIAL CONDITIONS

The composition of the NiCoCrAlY alloy is given in mass fractions in Table 1. The

tabulated composition was determined from the EPMA results, cf. Section 3.7. The

distribution of elements in the BC is displayed in Graphs 9a and 9b. The average of

the metallic elements include every measured spot (i.e. 1250 spots) while in the case

of oxygen spots above 0.5 wt% where left out (pores). Table 1: Mass fractions of the NiCoCrAlY alloy components (wt%) as determined with EPMA.

Al Y Cr Co Ni Average 12.6 0.4 17.4 24.9 44.6

3# 0.8 0.1 0.8 2.6 2.3

0

10

20

30

40

50

0 10 20 30 40 50

Depth (µm)

Mass F

ract

ion (

wt%

)

Al Cr Co Ni

Graph 9a: Cr, Al, Ni and Co depth profile determined with EPMA.

4. RESULTS

29

0

0.2

0.4

0.6

0.8

1

0 10 20 30 40 50

Depth (µm)

Mass F

ractio

n (

wt%

)

0

1

2

3

4

5

Num

ber

of

Y-o

xid

e s

po

ts

Y O #

Graph 9b: Yttrium and oxygen depth profile determined with EPMA. The x signs denote the number of

Y-oxide rich spots (right side axis).

The Y distribution in the bond coating as a function of depth is constant as evidenced

in the profile shown in Graph 9b. Y-oxide rich zones where not observed in EPMA

measurements. An Y-oxide rich zone was defined as a zone that contain Y in excess

of 1 wt% and enough oxygen after subtracting oxygen adsorbed at the sample

surface (0.24 wt%) to oxidize all the Y present.

The amount of oxygen adsorbed at the sample surface was determined from the

measurements of the as-sprayed BC, since in this sample the least internal oxidation

is expected. The average oxygen concentration in the as-sprayed BC equals 0.24

wt%, if the points containing more than 0.5 wt% (pores) are not taken into account

which is about20% of the measured spots.

The structure of the as deposited BC can be observed in Figure 8. Al-poor

precipitates are present as needles inside Al-rich zones as well as small grains.

According to the XPS analysis the samples at the beginning of every treatment are

covered with an oxide layer of less than 10 nm, which consists of: Al, Cr, Co and Ni

oxides. The composition of this native oxide is given in Table 2. Y is below the

detection limit and C was identified as the main surface contaminant. C is not taken

into account for composition analysis, but its photoelectron peak is used for charge

correction (cf. Section 3.4).

4. RESULTS

30

Figure 8: Backscatter image of the of the sample surface prior to heat treatment.

Table 2: Composition of the sample surface prior to heat treatment determined with XPS and binding

energy (BE) of the observed photoelectrons.

Y3dAl Al-O C O-M Cr-O Co-O Ni-O Y-O

BE (eV) 72.0 73.9 284.8 530.3 576.7 780.8 855.9 -C (at.%) 10.4 19.7 - 57.8 4.2 2.6 3.1 -

Co2p Ni2p

AS

Al2p C1s O1s Cr2pSTATEPEAK

4.2. 10 MINUTES ISOTHERMAL HEATING IN Ar-H2

Upon heating the BC in Ar with 5 vol.% H2 at 1100°C for ten minutes, the !-(Ni,Co)

phase initially present dissolves (Fig. 9). This observation is an indication of the

meta-stable nature of the as-deposited structure. The occurrence of an initial meta-

stable structure is caused by the high cooling rate associated with deposition method

(i.e. Low Pressure Plasma Spray). During sample introduction ingress of air into the

furnace caused the oxidation of Y and Al.

From backscatter electron images it can be seen that upon heating the initial needle

like precipitates dissolved (cf. Figs 8 and 10). Two main phases made up the

structure: the dark gray phase is "-(Ni,Co)(Al,Cr) and the light gray is !’-

(Ni,Co)3(Al,Cr)/!-(Ni,Co).

4. RESULTS

31

Figure 9: X-ray diffractogram of the NiCoCrAlY alloy prior to heat treatment (bottom) and after 10

minutes of isothermal heating at 1100°C under an Ar-H2 atmosphere (top).

Figure 10: Backscatter image of the of the sample surface (grinded) after heating for 10 minutes at

1100°C in an Ar-H2 atmosphere, showing the resulting structure. The light gray is !’-(Al/Cr)(Ni/Co)3 and

dark gray is "-(Al/Cr)(Ni/Co).

4. RESULTS

32

4.3. ISOTHERMAL HEATING UNDER VACUUM

After heating the samples to 1000°C at low oxygen partial pressure, XPS analysis

(see table 3) indicate that the TGO growth is hindered and Y is observed in the

surface after 10 minutes. As a function of time the amount Y and S at the surface

remain constant. Furthermore, from the sulfur 2s binding energy it can be said that it

is present as a Metal-S compound. Al in the surface is partially oxidized while the

other less reactive elements remain in the metallic state. The summary of these

observations can be found in Table 3.

Table 3: Composition of the sample surface determined with XPS and binding energy (BE) of the

observed photoelectrons. Sample isothermally heated at an oxygen partial pressure below 1x10-7 Pa.

S2s Y3dAl Al-O S-M C O-M Cr Co Ni Y-O

BE (eV) 71.9 73.8 225.3 284.8 530.2 573.6 777.6 852.2 157.9C (at.%) 10.7 5.8 0.5 - 35.9 2.4 3.0 11.7 30.1BE (eV) 71.7 - 225.3 0* 530.2 573.5 777.6 852.3 157.7C (at.%) 19.2 - 0.6 - 24.0 1.7 3.9 17.2 33.4BE (eV) 71.9 - 225.4 0* 530.3 573.4 777.6 852.3 157.7C (at.%) 18.2 - 0.6 - 25.0 1.4 4.2 19.7 31.0BE (eV) 71.6 73.0 225.4 284.8 530.0 573.2 777.3 852.0 157.4C (at.%) 15.0 3.1 0.5 - 32.5 1.1 3.6 15.9 28.4

*Substraction of 0.5 [eV] for charge correction

240 min

10 min

30 min

60 min

Co2p Ni2pAl2p C1s O1s Cr2pSTATEPEAK

4.4. 10 MINUTES ISOTHERMAL OXIDATION

The BC was isothermally oxidized for ten minutes at 900, 1000, 1100 and 1200°C.

The surface changes were studied using XPS, bulk changes were study with XRD

and the TGO structure was analyzed using electron backscatter images in

combination with EDS. Table 4 shows the surface composition at the four different

temperatures. The highest amount of Y at the surface is observed at the highest

temperature (i.e. 1200°C); the same holds for the other less reactive non-protective

elements (i.e. Cr, Ni, Co). The least amount of non-protective elements is observed

in the surface of the TGO grown at 1000°C.

Table 4: Composition of the samples surface isothermally oxidized for 10 minutes at an oxygen partial

pressure of 1x105 Pa determine from XPS.

S2s Y3dAl-O S C O-M Cr-O Co-O Ni-O Y-O

C (at%) 40.1 0.0 - 59.4 0.0 0.4 0.1 0.0C (at%) 46.3 0.1 - 53.4 0.0 0.3 0.0 0.0C (at%) 41.2 0.0 - 57.9 0.2 0.5 0.2 0.1C (at%) 38.9 0.0 - 57.2 0.4 0.4 0.3 2.9

Co2p Ni2p

900°C

Al2p C1s O1s Cr2pSTATEPEAK

1000°C1100°C1200°C

4. RESULTS

33

On the basis of the diffractograms recorded of the oxidized BC (Figure 11), The TGO

grown at 900 and 1000°C consist of metastable (!,0)-Al2O3. No inclusions were

observed in the backscatter electron images of cross-sections, Figs. 12 and 13.

However Y2O3 is formed internally in the BC.

Figure 11: Diffractogram of the NiCoCrAlY alloy prior to heat treatment (bottom), after 10 minutes of

isothermal oxidation at 900°C (middle), after 10 min of isothermal oxidation at 1000°C (top). The TGO grown at 900°C has an average thickness of 180 nm (Fig. 12a) while the

average thickness for the TGO growth at 1000°C is 330 nm (Fig. 13). The reported

average thickness was calculated from backscatter electron images using image

analysis software (photoshop). Micron 1 for every EPMA depth profile is the first measured line with an overall mass

fraction above 80 wt%. In Graph 10a it can be seen that only the first micron of BC

the composition is altered, compared to the as-deposited state. This observation can

be related to NiO nodules at the sample surface. In Graph 10b it can be seen. Y-

oxide can be found at a depth of 20 µm.

After ten minutes of isothermal oxidation at 1100°C a dense TGO consisting of /-

Alumina (Figs 14 and 15) is formed. Internal oxidation of Y is also observed. The

average thickness of the scale determine from Figure 8 is 330 nm.

4. RESULTS

34

Figure 12: Backscatter electron image of the cross section of the bond coating isothermally oxidized for

10 min at 900°C. (a) oxide layer composed of alumina and (b) actual oxide layer formed. Further

characterization of the composition of the top layer is required.

4. RESULTS

35

Figure 13: Backscatter electron image of the cross section of the bond coating isothermally oxidized for

10 min at 1000°C.

0

10

20

30

40

50

0 10 20 30 40 50

Depth (µm)

Mass

Fra

ction

(w

t%)

Al Cr Co Ni

Graph 10a: Ni, Co, Cr and Al depth profiles of the cross section of the NiCoCrAlY alloy isothermally

oxidized for 10 minutes at 1000°C.

4. RESULTS

36

0

0.2

0.4

0.6

0.8

1

0 10 20 30 40 50

Depth (µm)

Ma

ss F

ractio

n (

wt%

)

0

2

4

6

8

10

Nu

mbe

r Y

-oxid

e s

pots

O Y #

Graph 10b: Y and oxygen depth profiles of the cross section of the NiCoCrAlY alloy isothermally

oxidized for 10 min at 1000°C. x signs indicate the number of Y-oxide rich spots at a given depth.

Figure 14: X-ray diffractogram of the NiCoCrAlY alloy prior to heat treatment (bottom), after 10 minutes

of isothermal oxidation at 1100°C under an O2 atmosphere (middle), after 10 min of isothermal oxidation

at 1200°C (top).

4. RESULTS

37

Figure 15: Backscatter electron image of the cross section of the bond coating isothermally oxidized for

10 min at 1100°C.

Figure 16: Backscatter electron image of the cross section of the bond coating isothermally oxidized for

10 min at 1200°C.

4. RESULTS

38

0

10

20

30

40

50

0 10 20 30 40 50

Depth (µm)

Mass

Fra

ction

(w

t%)

Al Cr Co Ni

Graph 11a: Ni, Co, Cr and Al depth profiles of the cross section of the NiCoCrAlY alloy isothermally

oxidized for 10 minutes at 1200°C.

0

1

2

3

4

5

0 10 20 30 40 50

Depth (µm)

Mass F

raction

(w

t%)

0

5

10

15

20

25

Num

be

r Y

-oxi

de s

po

ts

Y O #

Graph 11b: Y and oxygen depth profiles of the cross section of the NiCoCrAlY alloy isothermally

oxidized for 10 min at 1200°C. x signs indicate the number of Y-oxide rich spots at a given depth.

4. RESULTS

39

For the highest temperature (i.e. 1200°C) the TGO consists mainly of stable /-Al2O3,

but is porus and non-protective spinel (Ni,Co)(Al,Cr)2O4 oxides are present at the top.

YAP (YAlO3) and YAG (Y3Al5O12) oxides identified by XRD (Figure 14) are observed

at the BC/TGO interface along with large Y2O3 particles bellow the scale (Fig. 16).

The average thickness of this TGO is 620 nm.

In Graph 11a it can be seen that the depletion region extends around 3 µm into the

BC for every element. The depletion is related to the formation of non-protective

(Ni,Co)(Al,Cr)2O4 spinel oxides at the sample surface. At 1200°C Y-oxides are found

until 3 µm underneath the surface. The penetration depth of oxygen is defined as the

depth at which the mass fraction of oxygen drops below 1 wt%, in this case this is

observed to occur at 3 µm also (cf. Figure 11b). This criterion is adopted to determine

the penetration depth of oxygen. 4.5. 60 MINUTES ISOTHERMAL OXIDATION

The surface composition of samples oxidized for one hour at 1000 and 1100°C is

shown in Table 5. Y is observed only at 1100°C. Taking into account the surface

composition after 10 minutes at 1100°C, it can be said that Y can diffuse through this

TGO and reach the surface (compare Y concentration at 1100°C for 10 and 60 min).

The TGO after one hour of oxidation at 1000°C has an average thickness of 660 nm

and is composed of / and (!,0) transition aluminas, Y2O3 is formed internally (Figs 17

and 18). After 60 min oxidation at 1100°C the TGO consists of /-Al2O3 and

(Ni,Co)(Al,Cr)2O4 spinel oxides in the top part. YAP oxides are found at the TGO/BC

interface and large Y2O3 particles can be seen below the TGO (Figs 19 and 20). The

average thickness of the TGO growth at 1100°C is 1570 nm after one hour (Fig. 20). Table 5: Composition of the sample surface determined with XPS and binding energy of the observed

photoelectrons. Sample isothermally oxidized for 60 minutes at an oxygen partial pressure of 101 kPa.

S2s Y3dAl-O S C O-M Cr-O Co-O Ni-O Y-O

BE (eV) 74.4 - 284.8 531.3 577.2 780.6 855.7 -C (at.%) 42.7 - - 56.7 0.1 0.3 0.1 -BE (eV) 74.0 - 284.8 530.8 576.6 780.2 856.0 157.9C (at.%) 41.5 - - 56.3 0.5 0.3 0.2 1.3

Co2p Ni2pAl2p C1s O1s Cr2pSTATEPEAK

1000°C

1100°C

4. RESULTS

40

Figure 17: X-ray diffractogram of the NiCoCrAlY alloy prior to heat treatment (bottom), after 10 minutes

of isothermal oxidation at 1000°C under an O2 atmosphere (middle), after 60 min of isothermal oxidation

at 1000°C (top).

Figure 18: Backscatter electron image of the cross-section of the bond coating isothermally oxidized for

60 min at 1000°C (top).

4. RESULTS

41

Figure 19: X-ray diffractogram of the NiCoCrAlY alloy prior to heat treatment (bottom), after 10 minutes

of isothermal oxidation at 1100°C under an O2 atmosphere (middle), after 60 min of isothermal oxidation

at 1100°C (top).

Figure 20: Backscatter electron image of the cross section of the bond coating isothermally oxidized for

60 min at 1100°C.

4. RESULTS

42

4.6. ISOTHERMAL OXIDATION AT 800°C

After the 10 minutes of isothermal oxidation at the four selected temperatures (i.e.

900, 1000, 1100, 1200°C) the samples were subjected to a second oxidation period

of 16 hours at 800°C. The follow-up of the changes have been investigated only for

the sample initially oxidized at 1000°C.

After the 16 hours oxidation at 800°C, no further surface enrichment of either Y, Cr,

Co or Ni is observed as can be seen from the XPS results in Table 6.

Table 6: Composition of the sample surface determined with XPS and the binding energy of the

observed photoelectrons. Sample isothermally oxidized at 800°C for 16 hours after 10 minutes of

isothermal oxidation at 1000°C.

S2s Y3dAl-O S C O-M Cr-O Co-O Ni-O Y-O

BE (eV) 74.7 - 284.8 531.8 576.9 781.0 0.0 0.0C (at.%) 40.9 - - 58.5 0.3 0.2 0.0 0.0

800°C

Co2p Ni2pAl2p C1s O1s Cr2pSTATEPEAK

Figure 21: X-ray diffractogram of the NiCoCrAlY alloy. Sample isothermally oxidized for 10 minutes at

1000°C (bottom). Sample oxidized for 16 hours at 800°C after 10 minutes of isothermal oxidation at

1000°C (top).

4. RESULTS

43

Figure22: Backscatter electron image of the cross section of the NiCoCrAlY alloy isothermally oxidized

for 16 hours at 800°C after a 10 minutes oxidation at 1000°C.

0

10

20

30

40

50

0 10 20 30 40 50

Depth (µm)

Ma

ss

Fra

ction

(w

t%)

Al Cr Co Ni

Graph 12a: Ni, Co, Cr and Al depth profiles of the cross section of the NiCoCrAlY alloy isothermally

oxidized for 10 minutes at 1000°C initially and then oxidized for 16 hours at 800°C.

4. RESULTS

44

0

2

4

6

8

10

0 10 20 30 40 50

Depth (µm)

Ma

ss F

ractio

n (

wt%

)

0

1

2

3

4

5

Nu

mb

er

Y-o

xid

e s

pots

Y O #

Graph 12b: Y and oxygen depth profile of the cross section of the NiCoCrAlY alloy isothermally oxidized

for 10 min at 1000°C initially and then oxidized for 16 hours at 800°C. x signs indicate the number of Y-

oxide rich spots.

The TGO average thickness calculated from Fig. 22 is 345 nm, which is slightly

thicker than after the initial ten minutes of isothermal oxidation at 1000°C. From the

XRD diffractogram and the backscatter electron image of cross-section (Figs 21 and

22), it can be seen that the TGO still consists of metastable (!,0)-Al2O3. Y2O3 is

present internally in the BC (Fig. 21).

Graph 12a shows no nickel depletion, although (NiO,CoO) nodules are present at the

surface. Y-oxides are observed until 11 µm bellow the surface as can be seen in

Graph 12b. Oxygen concentration above 1 wt% was measured only in the first

micron.

Hence, it can be said that no changes are observed in the TGO structure after 16

hours of isothermal oxidation at 800°C. However, slow growth of the oxide layer and

some increase of the internal Y oxidation occur.

4.7. ISOTHERMAL OXIDATION AT 1100°C

In order to study the effect of the initial 10 minutes oxidation on the structure of the

mature TGO, the following oxidation treatments were executed. After the two

oxidation treatments, i.e. (i) Isothermal oxidation at the four selected temperatures for

4. RESULTS

45

10 minutes, (ii) Isothermal oxidation for 16 hours at 800°C, the NiCoCrAlY alloy was

oxidized for 96 hours at 1100°C (cf. Section 3.8).

Figure 23: Backscatter electron images of the cross section of the NiCoCrAlY alloy isothermally oxidized

for 96 hours at 1100°C with an initial oxidation of 10 minutes at 900°C.

4. RESULTS

46

0

10

20

30

40

50

0 10 20 30 40 50

Depth (µm)

Mass F

ractio

n (

wt%

)

Al Y O Cr Co Ni

Graph 13a: Elemental depth profile of the cross section of the NiCoCrAlY alloy isothermally oxidized for

96 hours at 1100°C with an initial 10 minutes oxidation at 900°C. Right side axis applies only for Y. The

dotted lined indicate the approximate position of the TGO/BC interface.

0

1

2

3

4

5

0 10 20 30 40 50

Depth (mm)

Mass F

ract

ion (

wt%

)

0

10

20

30

40

50

Num

be

r Y

-oxid

e s

pots

O Y #

Graph 13b: Yttrium and oxygen depth profiles of the cross section of the NiCoCrAlY alloy isothermally

oxidized for 96 hours at 1100°C with an initial 10 minutes oxidation at 900°C. x signs indicate the

number of Y-oxide spots.

4. RESULTS

47

Figure 24: Backscatter electron images of the surface (a) and cross section (b and c) of the NiCoCrAlY

alloy isothermally oxidized for 96 hours at 1100°C with an initial oxidation of 10 minutes at 1000°C. In

24a, dark gray is alumina and light gray are mixed oxides.

4. RESULTS

48

0

10

20

30

40

50

0 10 20 30 40 50

Depth (µm)

Mass F

ractio

n (

wt%

)

Al Y O Cr Co Ni

Graph 14a: Elemental depth profiles in the NiCoCrAlY alloy isothermally oxidized for 96 hours at 1100°C

with an initial oxidation of 10 minutes at 1000°C. The dotted lined indicate the approximate position of

the TGO/BC interface.

0

1

2

3

4

5

0 10 20 30 40 50

Depth (µm)

Mass

Fra

ctio

n (

wt%

)

0

10

20

30

40

50

Nu

mbe

r Y

-oxid

e s

pots

O Y #

Graph 14b: Yttrium and Oxygen depth profile in the NiCoCrAlY alloy isothermally oxidized for 96 hours

at 1100°C with an initial oxidation of 10 minutes at 1000°C. x signs indicate the number of Y-oxide spots

at a given depth.

4. RESULTS

49

The TGO formed on the BC that was initially oxidized at 900°C, achieved an average

thickness of 5.35 µm with a fraction of inclusions ((Ni,Co)(Al,Cr)2O4, pores and Y-

aluminates) of 22%. Both the average thickness and the percentage of inclusions

were determined from Figure 23 using image analysis software.

In Graph 13a the TGO/BC interface can be located at 5 µm. At which point the

oxygen concentration is around 30 wt% (~40 at.%). Oxygen concentration drops

bellow 1 wt% at a depth of 21 µm. This is the oxygen penetration depth (as defined

previously). Y-oxides still can be found until 35 µm bellow the surface; see Graph

13b.

The TGO formed on the BC that was initially oxidized at 1000°C, achieved an

average thickness of 5.33 µm with a fraction of inclusion of 27%. A large amount of

non-protective oxides are observed at the sample surface (cf. Figure 14).

In Graph 14a the TGO/BC interface can be located at 3 µm. At which point the

oxygen concentration is around 25 wt% (~30 at.%). The oxygen penetration depth is

19 µm. Although Y-oxides still can be found until 35 µm bellow the surface; see

Graph 14b.

4. RESULTS

50

Figure 25: Backscatter electron images of the cross section of the NiCoCrAlY alloy isothermally oxidized

for 96 hours at 1100°C with an initial oxidation of 10 minutes at 1100°C.

0

10

20

30

40

50

0 10 20 30 40 50

Depth (µm)

Ma

ss

Fra

ction

(w

t%)

Al Y O Cr Co Ni

Graph 15a: Elemental depth profiles in the NiCoCrAlY alloy isothermally oxidized for 96 hours at 1100°C

with an initial oxidation of 10 minutes at 1100°C. The dotted lined indicates the position of the TGO/BC

interface.

4. RESULTS

51

0

1

2

3

4

5

0 10 20 30 40 50 60

Depth (µm)

Ma

ss F

ractio

n (

wt%

)

0

10

20

30

40

50

Nu

mbe

r Y

-oxid

e s

pots

Y O #

Graph 15b: Yttrium and Oxygen depth profile in the NiCoCrAlY alloy isothermally oxidized for 96 hours

at 1100°C with an initial oxidation of 10 minutes at 1100°C. x signs indicate the number of Y-oxide spots

at a given depth.

The TGO formed in the sample that was initially oxidized at 1100°C achieved an

average thickness of 5.30 µm with a fraction of inclusion of 30% (Figure 25).

In Graph 15a the TGO/BC interface can be located at 5 µm. At this point the oxygen

concentration is around 20 wt% (~25 at.%). The oxygen penetration depth is 19 µm.

Although yttrium oxides still can be found until 25 µm bellow the surface; see Graph

15b.

The TGO formed in the sample that was initially oxidized at 1200°C achieved an

average thickness of 7.17 µm with a calculated fraction of inclusion of 40%. A large

amount of non-protective oxides are observed at the sample surface (Figure 26).

In Graph 16a the TGO/BC interface can be located at 7 µm. At which point the

oxygen concentration is around 25 wt% (~30 at.%). The oxygen penetration depth is

19 µm. Although Y-oxides still can be found until 25 µm bellow the surface; see

Graph 16b.

4. RESULTS

52

Figure 26: Backscatter electron images of the sample surface (a) and cross section (b and c) of the

NiCoCrAlY alloy isothermally oxidized for 96 hours at 1100°C with an initial oxidation of 10 minutes at

1200°C. In 26a, dark gray is alumina, light gray are mixed oxides and lighter regions are Y-oxides.

4. RESULTS

53

0

10

20

30

40

50

0 10 20 30 40 50

Depth (mm)

Ma

ss F

raction (

wt%

)

Al Y O Cr Co Ni

Graph 16a: Composition depth profiles in the MCrAlY alloy isothermally oxidized for 96 hours at 1100°C

with an initial oxidation of 10 minutes at 1200°C. The dotted lined indicates the approximate position of

the TGO/BC interface.

0

1

2

3

4

5

0 10 20 30 40 50

Depth (µm)

Ma

ss F

ract

ion (

wt%

)

0

10

20

30

40

50

60

Nu

mb

er

Y-o

xide

sp

ots

O Y #

Graph 16b: Yttrium and Oxygen depth profiles in the NiCoCrAlY alloy isothermally oxidized for 96 hours

at 1100°C with an initial oxidation of 10 minutes at 1200°C. x signs indicate the number of Y-oxide spots

at a given depth.

4. RESULTS

54

4.8. SUMMARY OF RESULTS

The results obtain from the performed experiments are summarized in Tables 7-10.

The tabulated observations are: (i) atomic fraction of Y at the samples surface

determined from XPS measurements, (ii) oxide layer thickness and fraction of

inclusions determined from low magnification backscatter electron images, (iii)

phases identified by XRD and EDS and (iv) oxygen penetration determined from the

EPMA data analysis.

Table 7: Summary of the results for the 10 minutes isothermal oxidation

T (°C) Oxygen10 min Y-Surface (at.%) Tickness (nm) Phases Penetration (µm)

900 - 185 (!$0)-Al2O3

1000 - 330 (!$0)-Al2O3 <11100 0.1 365 /-Al2O3

/-Al2O3

(Ni, Co)(Al, Cr)2O4

YAlO3

Y3Al5O12

3

Ext. Oxidation

6402.91200

Table 8: Summary of the results for the 60 minutes isothermal oxidation

T (°C) Oxygen60 min Y-Surface (at.%) Tickness (nm) Phases Penetration (µm)1000 - 670 (!$0)-Al2O3

/-Al2O3

(Ni, Co)(Al, Cr)2O4

YAlO3

1100 1.3 1520

Ext. Oxidation

Table 9: Summary of the results for the 16 hours isothermal oxidation

Oxygen10 min 16 h Y-Surface (at.%) Tickness (nm) Phases Penetration (µm)1000 800 - 345 (!$0)-Al2O3 3

T (°C) Ext. Oxidation

Table 10: Summary of the results for the 96 hours isothermal oxidation

Ext. Oxidation Oxygen10 min 16 h 96 h Tickness (nm) Phases* Inclusions (vol.%) Penetration (µm)

900 5350 /-Al2O3 22 211000 5330 (Ni, Co)(Al, Cr)2O4 27 191100 5300 YAlO3 30 191200 7170 Y3Al5O12 40 19

* The forth phases are present in the four different temperatures

800 1100

T (°C)

5. DISCUSSION

55

5. DISCUSSION

Pre-oxidation treatments have proved to be and effective way to improve the lifetime

of high temperature coating systems [50]. Commonly, such a thermal treatment is

performed after a vacuum annealing which besides strengthening the substrate,

modifies the as deposited structure of the BC [16 and 17]. The aim of this discussion

will be to propose a pre-oxidation treatment before vacuum annealing of the BC, in

view of enhancing the lifetime of high temperature coating systems.

The discussion of the obtained results will be organized as follows: (i) structure of the

LPPS MCrAlY bond coating, (ii) external oxidation of Al, (iii) formation of

(Ni,Co)(Al,Cr)2O4 spinel oxides, (iv) Internal oxidation of Y, (v) summary of the

NiCoCrAlY alloy oxidation, (vi) TGO criteria in view of achieving an improved lifetime

of the high temperature coating system (cf. Section 2.2.6) and (vii) TGO formation.

5.1. STRUCTURE OF THE LPPS NiCoCrAlY BONDCOATING

Three phases where identified in the as sprayed bond coating, !-Ni/!’-Ni3Al which

have a face centered cubic structure, simple cubic "-NiAl and a hexagonal #-Co

phase which is meta-stable and dissolves readily upon heating for ten minutes at

1100°C in an Ar-H2 atmosphere.

Different authors report the presence of #-(Co, Cr) tetragonal phase and intermetallic

yttrium compounds (M5Y) located at phase boundaries, besides the !/!’ and " phases

[16-18]. The mentioned structures belong to samples that have been deposited by

LPPS but heat treated. The heat treatment is performed to simulate precipitation

hardening of the Ni-based super-alloy. During the thermal treatment the sample is

heated at 1150°C for 2 hours and later at 850°C for 24 hours, under vacuum.

5.2. EXTERNAL OXIDATION OF ALUMINUM

5.2.1.THEMODYNAMICS

At the beginning of the oxidation treatment a 10 nm thick native oxide (NO) layer is

present. The amorphous native oxide transforms to crystalline alumina upon thermal

oxidation [48].

A closed (!,0) transition alumina layer is formed after ten minutes of isothermal

oxidation, at 900 and 1000°C [cf. Figures 11-13]. Differentiation between /-Al2O3 and

5. DISCUSSION

56

meta-stable aluminas at the surface was attempted by determination of the Auger

parameter from XPS data [24-26]. Differences were observed only in the Auger

parameter of oxygen for the lowest kinetic energy line (i.e.

( )1111 1 LKLsO

OLKL BEhBEa $+= , . However, the observed differences are too small for

reliable identification.

Low angle XRD measurements were performed to reduce the depth of analysis. This

technique aids in the differentiation of the transition phases (reduction of the BC

signal), but the results did not correlate with the backscatter electron images or

XRD measurements performed at a higher incidence angle.

-900

-850

-800

-750

800 1000 1200 1400

Temperature (°C)

&G

(kJm

ol-1

)

2Al+1.5O2=a-Al2O3 2Al+1.5O2=g-Al2O3

Graph 17: Standard free energy plot as a function of temperature for /-Al2O3 and !-Al2O3.

The change in the standard free energy of formation is more negative for /-Al2O3 at

every temperature compared to !-Al2O3, as shown in Graph 17.

It has been shown that (!,0)-Al2O3 has excellent lattice matching with "-NiAl at given

relative orientations, this follows from the studies carried out by Doychack on 0-Al2O3

[12] and by Rhüle on !-Al2O3 [40]. Good lattice matching decreases the strain energy

at the TGO/BC interface and results in a reduction of the total standard free energy of

formation of transition aluminas to levels bellow that of /-alumina. This makes the

5. DISCUSSION

57

nucleation of transition aluminas thermodynamically possible. The reason for

nucleation of transition aluminas on NiCoCrAlY alloys is assumed to be analogous to

nucleation on "-NiAl.

The transition alumina(s) grows mainly by outward cation diffusion [9]. Alumina

growth by outward aluminum diffusion generates a needle like morphology at the

sample surface and a flat TGO/BC interface (cf. Figures 12 and 13 where the

alumina layer is composed of transition alumina(s)).

At the two higher temperatures (1100 and 1200°C) an /-Al2O3 layer was formed after

ten minutes of isothermal oxidation. This indicates that lattice matching no longer

provides the required reduction of the total standard free energy of formation. It is

also mentioned in the literature that the presence of Y-oxides aids in the nucleation of

/-Al2O3 [22]. The XPS measurements performed show the presence of Y-oxides at

the sample surface if the oxide layer consists of /-Al2O3.

It has been argued by several authors [2-4] that the growth of /-Al2O3 occurs

principally by inward oxygen diffusion. However, some dissident work is also found,

in the article written by K. Messaudi [45] important Al diffusion is suggested. It has

also been shown that as the alumina layer grows, the grain size increases. Hence,

the density of grain boundaries as well as the growth rate decreases [41].

Roughening of the TGO/BC interface can also be attributed to inward growth of the

oxide layer [cf. Figures 15 and 16].

5.2.2. KINETICS

Closed alumina layers of homogeneous thickness at the surface of an as-sprayed BC

can be formed after 10 minutes of oxidation (see Figs 12, 13 and 15). At 900 and

1000°C the oxide layer is made up of transition aluminas (cf. Section 4.4.), the

alumina structure is maintained after 16 hours of oxidation at 800°C (i.e. alumina

remains meta-stable), the oxide layer grows from 330 to 345 nm (cf. Section 4.6). If

the oxidation time is increased to 60 minutes at 1000°C, then the scale is composed

of stable /-Al2O3 and (!,0)-Al2O3 (cf. Figs 17 and 18) At higher temperatures only /-

Al2O3 is identified after ten minutes of oxidation (cf. Section 4.4.). Hence, during the

96 hours oxidation at 1100°C only /-Al2O3 should be formed.

Assuming that /-Al2O3 grows exclusively by inward oxygen diffusion, the oxide

growth constant can be related to the oxygen diffusion coefficient through alumina

5. DISCUSSION

58

grain boundaries. Several authors have studied the diffusion of oxygen through alpha

alumina. Most of these experiments where performed with bulk aluminas. These

results, when extrapolated to low temperatures, give diffusion coefficients that are

several orders of magnitude lower than those deduced from isothermal oxidation

experiments [44-46].

From backscatter electron images (see Figs 23-26), the oxide thickness can be

determined and the oxide growth constant can be estimated. Using the diffusion

model presented in section 2.1.2 for external oxidation, the oxygen diffusion

coefficient through the grain boundaries of the alumina layer can be determined, if

the change in oxygen potential is known.

The fraction of inclusions f after 96 hours of oxidation at 1100°C is presented in Table

9, for samples initially oxidized for 10 minutes at 900-1200°C. The thickness of a

clean /-alumina layer 32OAlx can be calculated using:

( )fxx MOOAl $= 132

(27)

where MOx is the determined oxide layer thickness.

0

2.5E-17

5E-17

7.5E-17

900 1000 1100 1200

Temperature (°C)

kP (

m2s-1

)

Oxide Al2O3

Graph 18: Determined parabolic growth constants for imaginary pure alumina layers and for the

determined oxide layer thickness for the different pre-oxidation temperatures (initial 10 minutes).

5. DISCUSSION

59

A common thickness increase of 3.5 µm is observed for samples initially oxidized at

1000 and 1100°C. For samples initially oxidized for ten minutes at 900 and 1200°C

the thickness increase was 4.0 µm, upon oxidation at 1100°C for 96 hours.

Using the determined oxide layer thickness increase upon 96 hours of oxidation at

1100°C, a parabolic growth constant can be calculated (Eq. 10). This constant is in

the order of 2.5x10-17 m2s-1 (4x10-14 gr2cm-4s-1), see Graph 18. This value is 1 order of

magnitude less than the growth constant of /-alumina reported by Grabke [9] at

1100°C on a "-NiAl substrate, e.g. 2.5x10-16 m2s-1 (5x10-13 gr2cm-4s-1). This reduction

can be associated with Y doping of the alumina grain boundaries. As can be seen in

Graph 19, Y reduces the diffusion coefficient of oxygen through alumina grain

boundaries.

1E-26

1E-23

1E-20

1E-17

1E-14

1000 1100 1200 1300 1400 1500

Temperature (°C)

D0

2 (m

2s

-1)

Nakagaw a Nakagaw a (Y) Prot (Y) Prot

Graph 19: Diffusion coefficients of oxygen along /-Al2O3 grain boundaries. Data taken from reference

43.

Using Eq. 11 in section 2.1.2 and the change in oxygen chemical potential gradient

given in Refs. 53 and 54, i.e. 75 and 750 kJmol-1, the effective diffusion coefficient of

oxygen through /-Al2O3 can be determined.

5. DISCUSSION

60

The oxygen chemical potential gradient can be approximate as follows based on

local thermodynamic equilibrium at the G/MO interface and the MO/M interface:

=&

MMO

O

MOGas

O

a

aRT lnµ (28)

MMO

O

MOGas

O

MMO

O

MOGas

O

C

C

a

a= (29)

The change in oxygen chemical potential ( reduction from 750 to 75 kJmol-1) can be

directly related to the increase in oxidation time, i.e. reduced Al activity and increased

oxygen penetration, both lead to a higher MMO

OC . The reduction of the driving force

for oxygen diffusion and the decrease in the density of grain boundaries in the oxide

layer (cf. images in Ref. 38), reduce the oxide layer growth rate as oxidation time

increases.

Using the growth constant determined for a pure alumina layer (e.g. 2.5x10-17 m2s-1),

an average grain size of 5x10-7 m and a grain boundary width equal to 1x10-9 nm [56]

GBOD' at 1100°C would be between 7x10-26and 7x10-25 m3s-1, the difference is due to

the different chemical potential gradient of oxygen used. The determined values

enclosed the GBOD' value calculated by Clemens in reference 46, i.e. 'DO=1x10-25

m3s-1. Then /-Al2O3 growth at 1100°C can be assumed be controlled by oxygen

inward diffusion through alumina grain boundaries, and Y doping of the alumina grain

boundaries reduce the oxygen inward diffusion (cf. graph 19).

This calculation procedure gives results of the right order of magnitude. Hence,

GBOD' can be calculated for the oxide layer formed on the BC.

After an initial oxidation of 10 minutes at 1000°C the oxide thickness is 0.33 µm (cf.

Table 7), this thickness is not significantly modified by the 800°C oxidation for 16

hours ( cf. Table 9). The scale growth after the complete treatment (see Flow chart in

Section 3.8) is 5.33 µm. Hence, the growth during 96 hours of oxidation at 1100°C is

5 µm (the same growth is observed for the samples initially oxidized at 900 and

1100°C). Applying the same analysis to the sample initially oxidized at 1200°C the

estimated growth at 1100°C after 96 hours of oxidation equals 6.5 µm.

5. DISCUSSION

61

This difference in thickness is due to the density of inclusions in the scale, and

doubles the growth constant. The growth constants are: 3.5x10-17 m2s-1 (7x10-14

gr2cm-4s-1) and 6x10-17 m2s-1 (1x10-13 gr2cm-4s-1) for samples initially oxidized for 10

minutes at 900-1100°C and 1200°C respectively, see Graph 18. The growth

constants were calculated using Eqs. 13 and 15, the values used in Eq. 15 are those

of /-Al2O3, i.e. =32OAlMOv 25.6 cm3mol-1 and =32OAl

MOw 16 grmol-1. The determined

values agree with the growth constant calculated from gravimetric experiments

performed by D. Toma [17] on VPS NiCoCrAlY coatings (e.g. 7.8x10-14 gr2cm-4s-1

after 260 hours at 1050°C).

The GBOD' values for the samples initially oxidized for ten minutes at 900-1100°C are:

1.4x10-25 m3s-1 (&µ=750 kJmol-1) and1.4x10-24 m3s-1 (&µ=75 kJmol-1).

The GBOD' values for the sample initially oxidized for 10 minutes at 1200°C are:

2.3x10-25 m3s-1 (&µ=750 kJmol-1) and 2.3x10-24 m3s-1 (&µ=75 kJmol-1). The presence

of oxide inclusions in the scale can be related to an increase of the oxygen diffusion

coefficient through grain boundaries by one order of magnitude.

5.3. NICKEL, COBALT AND CHROMIUM

5.3.1. THERMODYNAMICS

Ni (Co) oxide nodules are observed at the sample surface after ten minutes of

isothermal oxidation at every temperature, except at 1200°C (Ni,Co)(Al,Cr)2O4 spinel

oxides are found after ten minutes of oxidation (Figure 27).

Growth of these nodules is a common transient phenomenon of alloy oxidation, if

large differences exist in diffusion and thermodynamic properties of the oxides. For

example, nodules of less stable less protective oxides, have been observed in films

of more stable oxides formed on: "-NiAl, Ni-Cr, Co-Cr, Fe-Cr, Fe-Al, and Fe-Si alloys

[41]. As evidenced by Hindam and Smeltzer [41], NiO nodules grew in the early

exposure periods, these nodules grew preferentially at alloy grain boundaries.

However, as oxidation time was increased these nodules transformed into NiAl2O4

spinel oxides. They proposed that the transformation takes place by solid state

reaction between NiO and Al2O3 [41].

5. DISCUSSION

62

Figure 27: Backscatter electrons image of MCrAlY BC oxidized for ten minutes at 1100 nickel oxide

nodule (27a) and 1200°C spinel oxide formation (27b)

5. DISCUSSION

63

-1000

-750

-500

-250

0

800 1000 1200 1400

Temperature (°C)

&G

(kJ

/mol O

2)

-50

-40

-30

-20

-10

0

&G

(kJ

/mol O

2)

Ni+2Al+2O2=NiAl2O4 NiO+2Al+1.5O2=NiAl2O4

2Al+1.5O2=a-Al2O3 NiO+Al2O3=NiAl2O4

Graph 20: Plot of the standard free energy of formation Versus temperature for spinel oxides.

The following thermodynamic considerations should be taken into account to explain

the formation of the observed oxide structures:

(i) the change in the standard energy of formation for a solid state reaction is

negative but closed to zero (cf. Graph 20) (ii) after short oxidation periods NiO

nodules are present at the surface for oxidation temperatures as high as 1100°C (cf.

Figure 27), (iii) the change in standard free energy is slightly more negative for

spinel formation by doping of NiO than for formation of Al2O3 per mole of O2

consumed in the reaction (cf. Graph 20), (iv) /-Al2O3 formation is preferred over

spinel if it takes place at the TGO/BC interface, because the oxygen partial pressure

is too low to form NiO, (v) (!,0)-Al2O3 formation is preferred at the gas phase/TGO

interface over spinel formation due to the reduction in the standard free energy of

formation related to good lattice matching.

These considerations and the formation of voids in the upper part of the oxide layer,

lead to the conclusion that (Ni,Co)(Cr,Al)2O4 spinel oxides are formed by solid state

reactions.

5. DISCUSSION

64

5.3.2. KINETICS

The surface concentration of non-protective oxides (Ni, Co and Cr) is below 1 at.%

for every temperature after 10 minutes of isothermal oxidation. By increasing the

oxidation time to 60 minutes no increase of the surface concentration of Ni, Co or Cr

(cf. Tables 4 and 5) was detected. The same can be said for a long oxidation at

800°C. After heating at low oxygen partial pressure for ten minutes, the concentration

of non-protective elements at the surface rises over 15 at.% (cf. Tables 3 and 4).

From the EPMA measurements (cf. Graph 10) it can be seen that after ten minutes of

oxidation at 1000°C Ni is depleted from the first micron bellow the surface , this

depletion is believed to be related with the presence of NiO nodules at the sample

surface ( analogous to Figure 27). No Al depletion is observed since oxide layer in

included in the measured region. The same is observed for the sample oxidized at

1200°C, in this case depletion of Ni, Co and Cr is observed (cf. Graph 11b), which is

congruent with the presence of (Ni,Co)(Al,Cr)2O4 spinel oxides in the top part of the

oxides layer.

The increase in the amount of non-protective oxides (specially NiAl2O4) at the top of

the oxide layer, correlates with the depletion of Ni observed after 96 hours of

oxidation at 1100°C in the region bellow the scale (cf. Graphs 13a-16a)..

As determined by Pettit [43], the NiAl2O4 spinel oxide growth constant by solid state

reaction is around 7x10-17 m2s-1 at 1200°C. Then, after 96 hours of isothermal

oxidation the penetration of Al into NiO would be close to 7 µm. The size of the

observed (Ni,Co)(Al,Cr)2O4 spinel oxide inclusions is well below this limit (cf. Figures

23-26). It is also worth nothing that pores in the scale are located in the top part of

the TGO, next to (Ni,Co)(Al,Cr)2O4 spinel oxides (cf. Figs 23-26). These pores may

be due to condensation of vacancies plus volume changes. This is congruent with

spinel formation by a solid state reaction, taking place mainly by Al diffusion from

Al2O3 into NiO [43].

Although, after the pre-oxidation treatment the concentration of non-protective oxides

is very low (cf. Table 4), their concentration increases after 96 hours of isothermal

oxidation at 1100°C (cf. figs 23-26). Then, the presence of a protective oxide layer

reduce the diffusion Ni, Co and Cr towards the surface, but do not prevent it form

happening after prolonged oxidation periods.

5. DISCUSSION

65

5.4. YTTRIUM

5.4.1. THERMODYNAMICS

Y is found as Y-aluminates and as pure Y2O3. The location of the different types of

oxide depends on the oxygen partial pressure below the TGO.

At the TGO/BC interface, initially the oxygen partial pressure can be assumed to be

equilibrium pressure of alumina at the oxidation temperature, but as oxidation time

increases the oxygen partial pressure increases. Oxide formation below the close

alumina layer occurs, if the oxide being formed has a lower equilibrium pressure than

that at the TGO/BC interface or if it is formed by a solid state reaction. Since solid

state reactions are independent of oxygen partial pressure.

Y aluminates are observed at the TGO/BC interface and grow deep into the BC

(Figure 16). The standard free energy of formation of YAlO3 (YAP) and Y3Al5O12

(YAG) by solid state reactions is close to zero (Graphs 3 and 4). Also it is relevant to

point out that the solid state reaction between Al2O3 and Y2O3 powders start at

temperatures above 1200°C [42].

1E-75

1E-50

1E-25

0.7 0.8 0.9 1 1.1 1.2

1x10-3/Temperature (K)

log(p

O2)

0.5Y2O3+Al+0.75O2=YAlO3 1.5Y2O3+5Al+3.75O2=Y3Al5O12

2Al+1.5O2=Al2O3 2Y+1.5O2=Y2O3

Graph 21: Plot of the equilibrium oxygen pressure of Y-Al oxides versus inverse of temperature (700-

1400°C).

5. DISCUSSION

66

Al incorporation into Y2O3 posses a standard free energy of formation in between the

pure oxides (i.e. Al2O3 and Y2O3). Hence, the equilibrium oxygen pressure for the

reaction is below that of alumina, and the reaction can take place underneath the

oxide layer after short oxidation periods (cf. Graph 21).

YAP was identified as the only Y-aluminate after 1 hour of oxidation at 1100°C (cf.

Figure 19), while YAP and YAG are both present after ten minutes of oxidation at

1200°C (cf. Figure 14). This is due to the fact that the formation YAG requires higher

Y concentrations and Al incorporation. The Y concentration at the surface increases

at higher temperature (cf. Table 4 and 5) as the diffusivity of Al. Further away from the TGO/BC interface the amount of available oxygen drops, and

only Y2O3 is formed (see Figures 14 and 19).

5.4.2. INTERNAL YTTRIUM OXIDATION

Y segregation to grain/phase boundaries relaxes the lattice of the BC [21]. Once Y is

at the grain/phase boundaries, it diffuses towards the surface regardless of the

oxygen partial pressure magnitude at the surface. Size Y-oxide precipitates and

penetration of the internal oxidation front depends mainly on the diffusion of Y along

grain boundaries towards the surface.

Temperature has a large influence in the movement of Y towards the surface. After

ten minute of oxidation at 1100°C the surface concentration is only 0.1 at.%, but

when the temperature is raised to 1200°C the concentration increases to 3 at.% (cf.

Table 4). The same effect is observed if the oxidation time is increased from 10 to 60

minutes at 1100°C. In this case the concentration changes from 0.1 to 1.3 at.%.

However, after 16 hours at 800°C no Y is detected at the surface.

The enhancement of the kinetics due to oxidation temperature can also be realized

by comparing the backscatter electron images of the sample oxidized for ten

minutes at 1200°C (Fig. 16) with the ones oxidized at lower temperatures 1100-

900°C (Figs 12, 13 and 15). At 1200°C the surface highly enriched with Y containing

oxides, while in the samples oxidized at lower temperatures no Y-aluminates or large

Y2O3 particles can be observed, or by comparing the backscatter electron image of

the sample oxidized for ten minutes at 1200°C and the sample oxidized for 1 hour at

1100°C (Figure 20), i.e. at that point the structures observed are equivalent, i.e.

presence of (Ni,Co)(Al,Cr)2O4 spinel oxides at the top part of the scale, an /-Al2O3

5. DISCUSSION

67

oxide layer, Y-aluminates at the TGO/BC interface and pure Y2O3 deeper into the

BC.

At low oxygen pressure conditions the amount of Y segregation is striking. Low

oxygen pressure hinders the formation of an external alumina layer and the internal

oxidation of Y. This allows Y enrichment of the BC surface. After ten minutes at

1000°C the atomic fraction measured was 30 at.%, while at high oxygen pressure (1

bar) no Y was detected at the surface for the same temperature (Tables 3 and 4).

Hence, it can be said that oxidation temperature and oxygen partial pressure have a

much stronger effect than oxidation time in the Y distribution and that internal Y

oxidation hinders Y segregation.

0

1

2

3

4

5

0 10 20 30

Depth (µm)

Mass F

ract

ion

12

00

°C (

wt%

)

0

0.2

0.4

0.6

0.8

1

Mass F

ract

ion

10

00

°C (

wt%

)

1200-10M 1000-10M

Graph 22: Y distribution after ten minutes of isothermal oxidation at 1200°C (left side axis) compared

with the Y distribution of the sample oxidized for 10 minutes at 1000°C (right side axis).

From the EPMA measurements of the BC oxidized at 1000 and 1200°C the

characteristic diffusion length of Y can be determined. :

tDl YY 2= (30)

The characteristic Y diffusion length lY after 600 seconds is around 7.5 µm at 1000°C

and 15 µm at 1200°C (Graph 21).This leads to effective diffusion coefficients in the

order of 4.72.3x10-144 m2s-1 and 1.9x10-13 m2s-1 for 1000 and 1200°C, respectively.

5. DISCUSSION

68

Since D=Doexp(-E/RT) the diffusion coefficient at 1100°C can be estimated from:

$

=

21

11

ln2

1

TT

DDR

ETY

TY

= 108 kJmol-1 (31)

then, OD = 1.5x10-9 m2s-1 and ( )CTDEffY °= 1100 = 9.8x10-14 m2s-1.

The diffusion coefficient of yttrium through grain boundaries equals '2/YGBY rDD = .

The average grain size can be taken as 1x10-7 m (as-deposited BC). The grain

boundary width is usually set as 1x10-9 m [56]. This leads to effective diffusion

coefficients in the order of 5x10-12 m2s-1. The Y diffusion coefficient in NiCoAlCrY

alloys has not been thoroughly investigated and the only value found in literature was

reported by T. Nijdam [47] 5x10-12 m2s-1 for a EB-PVD NiCoCrAlY bond coating

where Y was present as an intermetallic compound at grain/phase boundaries.

Table 11: Penetration of the internal oxidation front xIOF for the different pre-oxidation temperatures after

the complete oxidation treatment.

Tpre-o (°C) PO-pre (µm) PMO/M (µm) PO (µm) xIOF (µm)

900 3 5 21 131000 3 3 19 131100 3 5 19 111200 3 7 19 9

The penetration of the internal oxidation front xIOF after 96 hours of oxidation can be

determined by subtracting the position of the MO/M interface PMO/M , and the oxygen

penetration after the 2 previous pre-oxidation treatments (cf. Section 3.8) PO-pre , from

the oxygen penetration depth after the complete oxidation treatment PO (cf. Table

11). In order to analyze the internal oxidation of Y the model presented in Section

2.1.2 can be used (solute Z=Y). In this model inward diffusing oxygen atoms react

with the outward diffusion Y atoms at the I.O.F., at this plane the concentration of

both atomic oxygen and metallic Y drops to zero. The concentration of oxygen does

not drop to zero but it is assumed that this plane is located at PO (oxygen

concentration is bellow 1%). The oxygen diffusion coefficient in the BC can be

estimated from Eq. 32 using the experimentally determined data. :

5. DISCUSSION

69

( )tkt

CD

CDx

p

YY

MMO

OOIOF int

22

22

2

22=

=*,

- (32)

where, DO and DY are the oxygen and Y diffusion coefficients in m2s-1 respectively

and , is the ratio of oxygen to Y atoms and equals 1.5 for Y2O3. The relevant

concentrations are determined with EPMA: MMO

OC =30-40 at.% (atomic fraction of

oxygen at the TGO/BC interface) and =*YC 0.25 at.% (atomic fraction of yttrium in as

deposited BC). Kofstad et. al. has assumed an oxygen concentration in the metal at

the MO/M interface of 30 at.% for the oxygen dissolution analysis in Ti, Zr and Hf

[59]. The calculated effective diffusion coefficient for oxygen in the BC is presented in

Table 12.

Table 12: Estimated Y and O effective diffusion coefficients in the BC. Minimum value of DO is in red. , CO (at.%) Initial T (°C) x (µm) DO (m2s-1) DY/DO

1000 13 6.8E-17 14341200 9 4.7E-17 20711000 13 5.1E-17 19121200 9 3.5E-17 2761

1.530

40

The DY/DO ratio is expected to be larger than 1. Eq. 30 indicates that the

characteristic diffusion length l and the diffusion coefficient D can be related as

follows:

2lD 1 (33)

for a given temperature and oxidation time. Thus the ratio between the diffusion

coefficients is given by:

2

=

O

Y

BCO

BCY

l

l

D

D (34)

The characteristic diffusion length of Y is observed to be 15 µm and the oxygen

penetration into the BC is observed to be around 3 µm for 10 minutes of oxidation at

1200°C, this values give a ratio equal to 25. No reference was found of this ratio. The

calculated oxygen diffusion coefficients can be used to determined an oxygen

distribution as a function of depth x and compare this with the oxygen penetration

curves as determined from EPMA measurements (underneath the oxide layer).

5. DISCUSSION

70

The oxygen penetration profiles are assumed to be of the following form:

( )

+=

tD

xBerfAxC

2 (35)

The boundary conditions are:

C(x=0)= MMOOC / =25 wt%

C(x=xIOF)=*OC <1 wt%

For these conditions the oxygen diffusion profile is:

( )

$$= *

BCO

IOF

BCO

OM

MO

OM

MO

OO

tD

xerf

tD

xerf

CCCxC

2

2 (36) (39)

0

5

10

15

20

25

0 10 20 30 40 50

Depth (µm)

Mass

Fra

ctio

n (

wt%

)

EPMA O-Profile Calculated O-Profile

Graph 23: Oxygen penetration profile of the sample initially oxidized for ten minutes at 1200°C. The

dotted red line indicates the approximate position of the TGO/BC interface.

5. DISCUSSION

71

0

5

10

15

20

25

30

0 10 20 30 40 50

Depth (µm)

Mass

Fra

ctio

n (

wt%

)

EPMA O-Profile Calc O-Profile

Graph 24: Oxygen penetration profile of the sample initially oxidized for ten minutes at 1000°C. The

dotted red line indicates the approximate position of the TGO/BC interface. The diffusion coefficients OD used were the ones presented in Table 12 (Written is

red). The calculated oxygen concentration profile resembles those measured by

EPMA (Graphs 21-24). This means that diffusion of Y through the BC is faster than

the oxygen penetration by three orders of magnitude.

An analytical solution can also be found for the internal oxidation model presented in

Section 2.1.2 neglecting the counter diffusion of Y. In such a case the relevant

solution would be Eq. 37 [appendix B Ref. 4]. The effective oxygen diffusion

coefficient through the BC in such a case would be in the order of 5x10-13 m2s-1. This

magnitude of the oxygen diffusion coefficient is closer to that observed in non-

protective Ni-Al alloys [56]

*=

Y

OM

MO

OIOF

C

DCx

,

2

2

2

(37)

There are three main reasons for the difference: (i) Fast counter diffusion of Y tied up

the penetrating oxygen closer to the surface, (ii) NiCoAlCrY alloys formed a

protective external Al2O3 layer and (iii) oxide particles formed in the NiCoCrAlY alloy

are larger in size and present in lower density, compared to the oxide particles

5. DISCUSSION

72

formed in non protective Ni-Al alloys. The reduced density of fast diffusion paths for

oxygen (i.e. MO/M interfaces) in the NiCoCrAlY alloy reduces the effective oxygen

diffusion coefficient [56].

The difference in the oxygen penetration observed between the sample oxidized at

1000°C and the one oxidized at 1200°C, can be attributed to the difference in Y

distribution after the initial 10 minutes oxidation. The large Y surface enrichment at

the higher temperature prevents oxygen penetration due to oxide formation. This

effect has also been observed in internal oxidation experiments [56], where the

penetration of the internal oxidation front decreases as the solute concentration

increases.

5.5. SUMMARY

5.5.1. THERMODYNAMICS

The alumina layer consists of meta-stable or /-alumina depending on oxidation

temperature. A high oxidation temperature reduces the strain energy due to lattice

mismatch. Hence, it reduces the effect that good lattice matching have on the

standard free energy of formation for transition aluminas and allows direct nucleation

of /-Al2O3. Nucleation of /-Al2O3 correlates with the presence of yttrium at the

sample surface (cf. Tables 4 and 5).

The closed alumina layer reduces the oxygen partial pressure below the TGO. The

conditions to from Y-aluminates by solute incorporation are: (i) outward diffusion of

aluminum (depletion), (ii) inward oxygen diffusion (/-Al2O3 is present when Y-

aluminates are present), (iii) previous existence of pure Y2O3 and (iv) high enough

oxygen partial pressure. All the mentioned conditions are met, if Y-aluminates are

identified. It is proposed that their formation takes place by Al and O incorporation

into pure yttrium oxide. Further away from the TGO/BC interface only pure Y2O3 is

present.

The thermodynamic considerations to explain the observed oxide layer structure

have been discussed (cf. Sections 5.2.1, 5.3.1 and 5.4.1). Spinel oxides are formed

at the upper part of oxide layer (cf. Figures 23-26). These oxides are formed by solid

state reaction between NiO/CoO and Al2O3/Cr2O3. This is concluded from the

presence of voids in the oxide layer when (Ni,Co)(Al,Cr)2O4 spinel oxides are

present (cf. Figures 16, 20, 23-26). If the formation occurs by solute incorporation no

5. DISCUSSION

73

voids should form in the oxide layer, since the doping elements (Al, Cr) would diffuse

from the BC.

5.5.2. KINETICS

As the alumina layer thickens and Y-aluminates are formed Al is depleted from the

region underneath the oxide layer. Al is incorporated into the protective alumina layer

and into the Y2O3 particles to form Y-aluminates. The region where the "-NiAl phase is dissolved below the oxide layer formed after 10

minutes of oxidation at 1200°C (cf. Figure 29) is considerably larger than bellow the

oxide layer formed after 10 minutes of oxidation at 1000°C (cf. Figure 28). From

Graph 18 it can be seen that the amount of Al2O3 formed is the same in every

sample. However, by plotting the Al depth profile underneath the TGO/BC interface it

can be observed that depletion region is larger at the higher temperature. Upon

oxidation at 1200°C the segregation of Y to the TGO/BC interface is higher as

compared with oxidation at 1000°C, this caused formation of a larger amount of Y-

aluminates (cf. Figures 13 and 16). This attributes to a higher growth rate of the oxide

layer and also to the larger Al depletion of the BC oxidized at 1200°C.

Figure 28: Backscatter electron image of the sample oxidized for 10 minutes at 1000°C. The bright

areas correspond to !-(Ni)/!´-(Ni3Al) and the dark areas correspond to "-(NiAl).

5. DISCUSSION

74

Figure 29: Backscatter electron image of the sample oxidized for 10 minutes at 1000°C. The bright

areas correspond to !-(Ni)/!´-(Ni3Al) and the dark areas correspond to "-(NiAl). In the "-phase depleted

region Y-oxides are observed.

0

10

20

30

40

0 10 20 30 40 50

Depth (mm)

Ma

ss F

rac

cti

on

(w

t%)

Al-1200°C Al-1000°C ASD

Graph 26: Al depletion underneath the TGO/BC interface of the samples initially oxidized for 10 minutes

at 1000 and 1200°C.

5. DISCUSSION

75

The oxide layer growth can be approximate by assuming that it grows exclusively by

inward oxygen diffusion. In the case of internal Y oxidation, both inward oxygen and

outward Y diffusion are relevant. Y diffuses three orders of magnitude faster than

oxygen in the BC (cf. Table 11). The oxidation temperature has a strong influence on

the Y diffusion, but not on the oxygen penetration (cf. Table 7).

The oxygen penetration depth is approximately the same after 96 hours of oxidation

regardless the pre-oxidation treatment (cf. Graphs 23 and 24). Since the segregation

of Y increases with the oxidation temperature (see Graph 22), the amount of Y-

oxides is larger at higher temperature.

5.6. TGO CRITERIA FOR AN IMPROVED LIFETIME

In the performed experiments only intrinsic failure mechanisms are operative and

since all intrinsic mechanisms have a characteristic TGO thickness ( )critMOx , the life

time analysis will be based on this feature, see Section 2.2.6. Failure along the

TGO/BC interface can be considered as an upper limit for the lifetime of the high

temperature coating system. If failure occurs within the TGO, then the lifetime of the

high temperature coating system is shorter and the TGO thickness after failure is

smaller than critMOx [50].

Figure 30: Backscatter electron image of Y-Aluminates formed at phase/grain boundaries and inside

grains.

5. DISCUSSION

76

Failure within the TGO is associated with progressive damage build up in the vicinity

of defects such as voids and micro-cracks which are commonly observed in the

upper part of the oxide layer where (Ni.Co)(Al,Cr)2O4 spinel oxides are formed (cf

Figures 23-26). Thus, the formation of these oxides should be minimized.

As discuss in Section 5.2.2 and 5.5.2, the growth of the oxide is faster and the "-NiAl

phase depletion is enhanced, if the amount of Y-rich oxide inclusions in the oxide

layer increases. Thus, to reduce the growth of the scale, the segregation of Y to the

surface (top 5 µm of the BC) must be minimized.

As mentioned in section 4.4., Y2O3 forms after short oxidation times when the BC

microstructure is fine. During extended periods of high temperature (1100°C)

oxidation, recrystalization and grain growth in the BC occurs [52]. Then, a fraction of

the initially formed Y2O3 particles become part of the growing " or !/!’ grains. The

oxidation lifetime of ODS alloys is usually 2-3 time lower than that of FeCrAlY alloys

of similar composition [51]. This difference is linked to the benefit of stress relaxation

in the weaker wrought material. If internal oxidation of Y occurs, and the formed

oxides become part of the new grains, the elastic modulus of the BC in the internal

oxidation zone (IOZ) may increase. Consequently, the BC deformation is hindered

and hence relive of the stresses generated in the scale upon cooling is impaired. The

reduction of stress release capability causes an increase of the strain energy in the

oxide layer. This can be directly related to a reduction in the oxidation lifetime of the

alloy (i.e. smaller critical thickness).

Also internal Y oxidation hinders the formation of Y-sulfides since more stable Y-

oxides are being formed (reactive element effect, cf. Section 2.2.5.). If S segregates

to the TGO/BC interface its fracture toughness will be reduced. Thus, to increase the

lifetime, oxygen penetration should be minimized.

Failure along the TGO/BC interface is promoted upon increasing oxidation time

since: (i) as the TGO thickens the strain energy stored increases, (ii) the fracture

toughness of the TGO/BC interface may decrease due to segregation impurities from

the BC, (iii) the density of Y-rich oxide inclusions in the oxide layer increases, and (iv)

the amount of (Ni,Co)(Al,Cr)2O4 spinel oxides at the top of the oxide layer increases.

So, to extend the lifetime of the high temperature coating the following four points

should be addressed: (i) the TGO should have the lowest growth rate at the

temperature of interest (900-1200°C), (ii) the BC/TGO interface must be as strong as

5. DISCUSSION

77

possible in order to increase the critical TGO thickness, (iii) large Y enrichment of the

surface and extended oxygen penetration should be avoided and (iv) the density of

defects within the TGO should be small in order to suppress the number of crack

initiation sites for premature failure.

5.7. TGO FORMATION

To from an oxide layer on the TGO that grows as slowly as possible during

isothermal oxidation, the presence of other oxides besides Al2O3 in the initial

oxidation stage should be minimized.

Closed oxide layers that exclusively consist of alumina can be formed after short

oxidation times, cf. Table 7 and Figures 12, 13 and15. Then, it is an advantage to

have a fine BC structure with a high density fast diffusion paths for Al and oxygen.

This is the case directly after BC deposition (see Figure 8). However, Y segregation

towards the surface is also facilitated.

Extending the oxidation time to 1 hour and still form a pure alumina layer is possible

at 1000°C (cf. Figure 18). The formation of closed oxide layers of transition aluminas

may be of advantage (formed at CT 1000= ), because of two reasons: (i) the

transition alumina will transform to coarse grained [30] stable /-Al2O3 upon solution

treatment (first part of the precipitation hardening) of the substrate precipitation and

(ii) metastable aluminas grow mainly by outward Al diffusion, which reduces the

mass fraction of oxygen underneath the TGO after the pre-oxidation treatment.

The formation of (Co, Ni) oxide nodules can not be avoided by shortening the

oxidation time or by decreasing the oxidation temperature (Figure 27). However, their

density is minimal after the 10 minutes oxidation treatments (cf. Table 4). Although,

the alumina layer is closed and (Co, Ni) oxide concentration at the surface is low, a

significant amount of (Ni,Co)(Al,Cr)2O4 spinel oxides are formed at the sample

surface after 96h hours of isothermal oxidation (cf. Figures 23-26). This means that

metallic elements (e.g. Ni and Co) diffuse through the closed alumina layer towards

the surface. A possible solution can be to increase the thickness of the outer alumina

layer.

Thickening of the external alumina layer is considerably slow at low oxidation

temperature (800°C). Long oxidation periods at low temperature will form pure

alumina scales. Since Y diffuses slower at lower temperatures, the density of Y-rich

inclusions in the scale, the size of the Y-oxides (cf. Figures 16 and 20) and the

5. DISCUSSION

78

growth rate of the scale decreases. But, since oxidation time is longer the oxygen

diffusion into the BC will increase (cf. Table 10) as well as "-NiAl depletion.

This approach was initially considered (cf. Figures 13 and 22). Instead, thickening of

the external scale is better achieved by increasing oxidation time at temperatures

below 1100°C to minimize Y segregation towards the surface (cf. Table 5), but above

900°C in order to enhance the oxide layer growth kinetics.

Pre-oxidation treatments above 1100°C are not recommended, massive Y

segregation and rapid formation fast growing oxides takes place even after a very

short time plus large depletion of the "-NiAl phase, which is the main Al source for

the formation of the protective alumina layer (c.f. Figure 16). After just 10 minutes,

heating at 1000°C in the vacuum chamber large amounts of Y can be found at the

sample surface. The Y concentration at the surface remains constant with time

during the low pressure heating. At high oxygen partial pressure and the same

temperature no Y at the sample surface is observed. Then, pre-oxidation treatments

(with native oxide on BC with fined grained structure) at very low partial oxygen

pressures should be avoid, since they cause rapid Y diffusion towards the surface.

After long time oxidation periods, the oxide structure show the same features for any

pre-oxidation treatment: i.e. a thin layer with a high density of (Ni,Co)(Al,Cr)2O4 spinel

oxides on top of thick layer of alumina with Y-rich oxide inclusions (Figures 23-26).

The initial treatment has proved to be able to change the density of inclusions

present in the mature scale and thereby its growth rate.

6. CONCLUSION

79

6. CONCLUSION: Proposed pre-oxidation treatment

As the main conclusion of the research executed, a pre-oxidation treatment is

proposed:

Perform the pre-oxidation before precipitation hardening of the super alloy substrate.

In this way the fined grained structure of the as deposited BC helps in forming a

closed alumina layer of uniform thickness. Furthermore, the effects of the substrate

strengthening on the oxide layer formed after pre-oxidation may be of advantage for

the lifetime of the high temperature coating system. Because upon solution heat

treatment of the super alloy substrate, the metastable oxides formed after pre-

oxidation transform into stable /-Al2O3 with relatively large grains.

After very short oxidation times the Y distribution is effectively modified. Since Y

movement towards the surface should be minimized, the oxygen partial pressure

should be high and the temperature should not exceed 1100°C during the pre-

oxidation treatment. To reduce oxygen penetration and "-NiAl phase depletion, the

pre-oxidation time should short.

The oxidation treatment should be carried out between 1000 and 1100°C at an

oxygen pressure of about 1 bar. A lower temperature will reduce Y diffusion towards

the surface. The oxidation time can be varied between 15 and 25 minutes depending

on the temperature. In this way the formed oxide layer will be composed of pure

alumina and thick enough to minimize the formation of non-protective oxides at the

top part of the oxide layer.

Thermal cycle oxidation experiments are necessary to determine the effect of the

proposed pre-oxidation treatment on the lifetime of the high temperature coating.

7. RECOMENDATIONS

80

7. RECOMMENTDATIONS

After the pre-oxidation treatment, the heat treatment of the super-alloy should be

simulated. Since the initial two hours at 1150°C of heating in argon/vacuum may

modify the Y distribution and the structure of alumina. The second part of the super-

alloy strengthening treatment, will not affect the Y distribution but, can modify the

nickel distribution as was evidenced after the 16 hours of isothermal oxidation at

800°C (cf. Graphs 10a and 12a), the Ni depletion region observed after ten minutes

is not present after the second 16 hours oxidation period.

The next step should be to compare the lifetime in a cyclic oxidation experiment for

samples pre-oxidized before strengthening of the substrate at 1000°C (low density Y-

rich oxides in the oxide layer) and 1200°C (high density of Y-rich oxides in the oxide

layer). Also the ability of the BC, after precipitation hardening of the substrate, to form

closed alumina scales after short oxidation time should be investigated, and a

comparative analysis of the TGO performance when the pre-oxidation is done before

and after precipitation hardening of the substrate.

81

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