Enhancement of Hot Corrosion Resistance of Thermal Barrier ...
148
i Enhancement of Hot Corrosion Resistance of Thermal Barrier Coatings through Modified Configuration of Bondcoat By Imran Nazir Qureshi School of Chemical and Materials Engineering (SCME) National University of Sciences and Technology (NUST) 2017
Transcript of Enhancement of Hot Corrosion Resistance of Thermal Barrier ...
i
Enhancement of Hot Corrosion Resistance of
Thermal Barrier Coatings through Modified
Configuration of Bondcoat
By
Imran Nazir Qureshi
School of Chemical and Materials Engineering (SCME)
National University of Sciences and Technology (NUST)
2017
ii
Enhancement of Hot Corrosion Resistance of
Thermal Barrier Coatings through Modified
Configuration of Bondcoat
Imran Nazir Qureshi
Reg.No. 2012-NUST-DirPhD-MS-E-09
This work is submitted as PhD thesis in Partial Fulfillment of The
requirements for the degree of
PhD in Materials and Surface Engineering
Supervisor: Dr. Muhammad Shahid
School of Chemical and Materials Engineering (SCME)
National University of Sciences and Technology (NUST), H-12
Islamabad, Pakistan
2017
iii
iv
v
vi
vii
viii
Dedication
I dedicated this work to my parents and family, whose affectionate support was a
continuous supply of motivation
ix
Acknowledgement
All thanks to Almighty Allah who gave me the ability to complete the project successfully.
It is a great honor and privilege to present a deep sense of gratitude to my supervisor Dr.
Muhammad Shahid for his kind support and positive criticism. He always encouraged me in
developing critical thinking not only for this project but also in other scientific areas. I feel
obliged for the time he devoted in successful completion of this dissertation.
My gratitude also goes to my Co-Supervisor Dr. Aamer Nusair for his guidance and technical
help in processing of coatings by plasma spraying at the institute of industrial control systems
(IICS), Rawalpindi. I have learnt a lot from him in the field of coatings.
I also express most cordially thanks GEC members: Dr. Muhammad Mujahid, Dr. Adeel umer
and Dr. Iftikhar us Salam for their excellent advice and selfless succor.
I pay my salutations and recognitions to Dr. Shaheed, Dr. Shabbar Abas Rizvi, Dr. Ahnaf Usman
and Mr. Khalid Mahmood for their complete co-operation during my research work. I cannot
forget the commendable assistance of KRL staff especially, Mr. Khawar, Mr. M. Irfan, Mr.
Muhammad Imran Baig, Mr. M. Ghafoor, Mr. Munir, Mr. Shahid Mahmood, Mr. Shahab, Mr.
Shams ud Din, Mr. Shaqib, Mr. Ijaz Hussain and Mr. Mukhtar Ahamad.
I believe it won’t be justified not to concede the great endurance of my compassionate wife and
adoring daughters for their inestimable support.
x
Abstract The Trend of applying coatings is widely increasing in industrial and aero turbines. Thermal
barrier coatings (TBCs) are widely used to resist high temperatures and to protect the base metal
from exposure to high temperatures. TBCs are becoming increasingly essential to protect the hot
section components of gas turbine engines against oxidation and hot corrosion. Invariably, it is
the corrosion of the coating that determines the need for refurbishment and the life of a
component, and not the loss of high-temperature mechanical strength.
In present work, a novel technique has been introduced to enhance the life of TBCs by applying
thin layer of TiN (by physical vapor deposition method) on bondcoat for improving the oxidation
and hot corrosion resistance of bondcoat. Standard TBC samples (Yttria stabilized zirconia
thermal barrier coating (TBC) along with CoNiCrAlY bondcoat) were compared with TiN
modified bondcoat TBC samples. Both TBC systems were exposed to high temperature under
the presence of corrosive salts i.e. a mixture of V2O5 and Na2SO4 (50:50) for 50 hours. The
characterizations of the coatings included X-ray diffraction analysis, scanning electron
microscopy and optical microscope. It was observed that TiN modified samples showed better
results in terms of oxidation resistance and delamination. The formation of Cr2Tin-2O2n-1 phases
at the interface of topcoat-bondcoat, in TiN modified samples, were found responsible to
enhance the thermal and oxidation properties of the TBC. The durability of coatings is evaluated
by thermal cycling. After 225 cycles, the standard TBC samples spalled 30% of the topcoat,
whereas, the TiN-modified samples spalled only 5%. Based on above results it could be
concluded that TiN-modified coating interface was better thermal shock resistant and enhanced
the life of coatings against hot corrosion.
The effect of bondcoat thickness on the hot corrosion resistance was studied. Results indicated
that TBCs with thick bondcoat exhibited superior hot corrosion resistance to the TBCs with
conventional bondcoat. The reaction products were microscopically investigated and a detailed
elemental diffusion of different alloying elements was investigated.
xi
Contents
Contents xi
List of Figures xiv
List of Tables xxi
1 Introduction 1
1.1 Objectives of Work ....................................................................................................................... 3
1.2 Outline of Thesis ........................................................................................................................... 4
2 Literature Survey 6
2.1 Diffusion coatings ......................................................................................................................... 8
2.2 Overlay coatings ........................................................................................................................... 8
2.3 Thermal barrier coatings ............................................................................................................... 8
2.3.1 Ceramic materials for TBCs. .............................................................................................. 11
2.3.2 Bondcoat Materials ............................................................................................................. 16
2.3.3 TBCs Methods of Deposition ............................................................................................... 17
2.3.4 Oxidation and hot corrosion of TBCs ................................................................................. 27
2.3.5 Selection of TiN for modification of bondcoat ................................................................... 32
3 Experimental Procedures 33
3.1 Process Flow Chart ..................................................................................................................... 33
3.2 Experimental ............................................................................................................................... 34
3.2.1 Material of the Substrate ..................................................................................................... 34
3.2.2 Cutting and Cleaning .......................................................................................................... 34
3.2.3 Cylindrical Holder............................................................................................................... 35
3.2.4 Grit Blasting ........................................................................................................................ 35
3.2.5 Coating Materials ................................................................................................................ 37
3.2.6 Deposition Process .............................................................................................................. 38
3.2.7 Cathodic Arc Physical Vapor Deposition of Titanium Nitride ............................................. 41
3.2.8 Hot Corrosion ...................................................................................................................... 42
3.2.9 Thermal Cycling Treatment ................................................................................................ 43
xii
3.2.10 Adhesion Test ..................................................................................................................... 43
3.2.11 Delamination of Topcoat after Hot Corrosion .................................................................... 44
3.3 Characterization .......................................................................................................................... 44
3.3.1 Stereo Microscopy .............................................................................................................. 44
3.3.2 Sample Preparation (Metallography) .................................................................................. 44
3.3.3 Optical Microscopy ............................................................................................................. 45
3.3.4 Scanning Electron Microscopy (SEM).................................................................................. 45
3.3.5 X-Ray Diffraction (XRD) Study ............................................................................................. 46
4 Results and Discussion-I 48
4.1 Hot corrosion of yttria-stabilized zirconia coating, in a mixture of sodium sulfate and vanadium
oxide at 950oC. ........................................................................................................................................ 48
4.1.1 Microscopy ......................................................................................................................... 48
4.1.2 Chemical Composition Profile ............................................................................................ 56
4.1.3 X-Ray Diffraction Analysis ................................................................................................ 62
4.1.4 Delaminated topcoat sample ............................................................................................... 67
4.2 Conclusions ................................................................................................................................. 68
5 Results and Discussion-II 69
5.1 Evaluation of titanium nitride modified bondcoat system used in thermal barrier coating in
corrosive salts environment at high temperature .................................................................................... 69
5.1.1 Results and Discussion........................................................................................................ 70
5.1.2 Delaminated Topcoat Obtained after Hot Corrosion .......................................................... 80
5.1.3 X-Ray Diffraction Analysis ................................................................................................ 82
5.1.4 Conclusion .......................................................................................................................... 84
5.2 Thermal cycling behavior of air plasma sprayed thermal barrier coatings on Inconel X750 alloy
with and without TiN modification of ‘bondcoat’. ................................................................................. 85
5.2.1 Result and Discussion ......................................................................................................... 85
5.2.2 Conclusion .......................................................................................................................... 95
6 Results and Discussion-III 96
6.1 Effect of bondcoat thickness on hot corrosion of ZrO2-8Y2O3 thermal barrier coating ............. 96
6.1.1 Result and discussion .......................................................................................................... 96
6.1.2 Conclusion ........................................................................................................................ 104
xiii
6.2 Behavior of air plasma sprayed thermal barrier coatings with different bondcoat thicknesses,
subject to intense thermal cycling ......................................................................................................... 105
6.2.1 Result and Discussion ....................................................................................................... 105
6.2.2 Conclusion ........................................................................................................................ 112
7 Summary 113
8 Publications and presentations 116
9 References 117
xiv
List of Figures
Fig.2.1: Increase of operational temperature of turbine components made possible
by alloy development and TBCs
6
Fig. 2.2: Relative service temperature enhancement as a result of improvement in
coating technologies
7
Fig. 2.3: Temperature drop substrate’s surface by YSZ coating
9
Fig.2.4: Thermal conductivities of various polycrystalline oxides in as a function
of temperature
11
Fig.2.5: Thermal conductivity vs. thermal expansion coefficient of various ceramic
materials
12
Fig.2.6: Yttria - Zirconia phase diagram. Note that the shaded region indicates the
region where the formation of the metastable t’ phase occurs upon
cooling.
13
Fig. 2.7: Schematic illustration of plasma spraying process
20
Fig. 2.8: Energy contents of various plasma forming gases as a function of
temperature
23
Fig.2.9: Thermal spray coating (a) plasma sprayed as-sprayed coating [117] (b)
coating build up
25
Fig. 2.10: The EB-PVD process (a) schematic of the electron beam process (b)
TBCs deposited by EB-PVD
28
Fig. 2.11: Schematic illustration of temperature effect on rate of damage to
superalloys based on type I and II hot corrosion superimposed on
contribution due to oxidation
29
Fig. 3.1: The process flow chart
33
Fig. 3.2: Aluminum cylindrical holder for holding substrate samples during grit
blasting and coating
35
Fig. 3.3: Grit blasting set up used for samples cleaning
36
Fig. 3.4: Morphology of coating powders used for plasma coating (a) YSZ ceramic
powder (b) Co32Ni22Cr8Al0.5Y metallic powder.
37
xv
Fig. 3.5: Schematic diagram showing the setup of APS
38
Fig. 3.6: A schematic diagram showing experimental setup for APS coating used in
this study.
39
Fig. 3.7: A thin layer of TiN (golden color) was deposited by physical vapor
deposition method after the deposition of bondcoat for TiN modified
samples
40
Fig. 3.8: TBC coating systems after the deposition of topcoat
40
Fig. 3.9: DREVA RC 400 Coating plant used for deposition of Titanium Nitride
films on the samples.
41
Fig.3.10: Standard TBC samples, placed in a stainless steel tray, with salt mixture
on the top surface of the samples for hot corrosion test at 950˚C
42
Fig. 3.11: A schematic setup used for evaluation of bonding strength: (a) showing
the parts used for testing; (b) the assembly used for tensile adhesion
testing.
43
Fig. 3.12: BX51 Olympus optical microscope with digital camera and image
analyzer
45
Fig. 3.13: Scanning Electron Microscope (SEM) JEOL JSM 5910 LV with EDS
46
Fig. 3.14: X-Ray Diffractometer used for phase analysis of coatings
47
Fig.4.1: The as-sprayed topcoat surface exhibiting a rough surface with few semi-
molten particles (SMP), cracks and porosity like features
49
Fig.4.2: Topcoat surface after 50 hours exposure to hot corrosion environment,
containing YVO4 rods and agglomerated crystals of ZrO2 (b) high
magnification, rod like and agglomerated crystals
50
Fig.4.3: Topcoat surfaces after, exposing to 950°C in a hot corrosion environment
for various time intervals, showing increased concentration of YVO4 rods
with higher exposure times.
50
Fig.4.4: (a) Cross-section of as sprayed sample (b) high magnification of bondcoat
showing grayish flake-like features and porosity
51
Fig.4.5: Cross section of samples exposed to hot corrosion environment at 950°C,
showing effect of exposure time on thermally grown oxide (TGO).
52
xvi
Fig.4.6: SEM micrograph showing cross-section of bondcoat in a sample exposed
to hot corrosion testing for 50 hours: Site-1 and 2 are locations of EDS
analyses performed; data is given in Table-4.1
53
Fig.4.7: SEM micrograph showing discrete sites in the bondcoat in a sample
exposed for 50 hours showing knife-like features (Site-1 and 5). Site-1 to
5 represent locations for EDS analyses reported in Table No.4.2.
54
Fig.4.8: Cross-section of the coating showing increase in deterioration (arrows) of
the top surface with time.
55
Fig.4.9: Schematic diagram (a-d) showing various oxidation reactions during hot
corrosion testing.
55
Fig.4.10: Elemental distribution after 10 hours exposure in hot corrosion
atmosphere at 950ºC
58
Fig.4.11: Profile of elements after 30 hours exposure in hot corrosion atmosphere at
950ºC
59
Fig.4.12: Profile of elements after 40 hours exposure in hot corrosion atmosphere at
950ºC
60
Fig.4.13: Profile of elements after 50 hours exposure in hot corrosion atmosphere at
950ºC
61
Fig.4.14: Titanium-rich precipitates (arrows) close to the substrate-bondcoat
interface.
62
Fig.4.15: XRD graph showing patterns of as-sprayed coating and samples exposed
to hot corrosion for various durations.
63
Fig.4.16: Effect of exposure time on the percentage of YVO4
64
Fig.4.17: Effect of exposure time on the percentage of m-ZrO2 and t-ZrO2
64
Fig.4.18: XRD patterns showing shift of (200) plane with exposure time at high
temperature.
65
Fig.4.19: Effect of exposure time on the lattice parameters of m-ZrO2 and YVO4
66
Fig.4.20: XRD pattern showing multiple phases formed in the delaminated topcoat of the sample exposed at 950ºC for 50 hours
67
Fig.5.1: Standard TBC samples (row-1) and TiN modified sample (row-2), placed
in a stainless steel plate, with salt mixture on the top surface of the
samples
69
xvii
Fig.5.2: S2 (Standard TBC) and TiN modified samples showing condition of top
surfaces after different time intervals, treated at 950°C in a hot corrosion
environment
71
Fig.5.3 Top surface of standard TBC (a) low magnification (b) high magnification
showing rod like features (YVO4) after 50 hours exposure
72
Fig.5.4 Top surface of TiN-modified sample (a) low magnification (b) high
magnification showing rod like features (YVO4) after 50 hours exposure
72
Fig.5.5 (a) SEM micrograph showing typical structure of as sprayed TBC
coating. (b) high magnification of bondcoat showing lamellar structure
73
Fig.5.6 Optical micrograph showing layer of TiN (arrows) and interface. Vertical
cracks (box) are also present at some locations.
73
Fig.5.7 Optical micrograph showing layer of TiN which was not deposited
properly at some locations, (arrows)
74
Fig.5.8 Cross-section of both standard TBC (a) and TiN modified (b) samples,
after 50 hours exposure in hot corrosion environment
75
Fig.5.9 Cross-section of TiN modified sample after 50 hours exposure in hot
corrosion environment demonstrating dense and uniform oxide layer at
interface of bondcoat-topcoat
75
Fig.5.10 Non-uniform oxide layer in Standard TBC sample after 50 hours exposure
(a-low mag. and b-high mag.)
76
Fig.5.11 EDS analysis at the boundaries of the splats in bondcoat showed no
“vanadium” is present in TiN-modified samples (a), standard TBC
samples demonstrated the presence of vanadium (b) near topcoat-
bondcoat interface
76
Fig.5.12: Different sites are marked from where EDS analysis at bondcoat-topcoat
interface were taken, in TiN modified sample after 50 hours exposure.
EDS analysis are represented in Table 5.1.
78
Fig.5.13: Standard TBC sample (after 50 hours exposure) demonstrating topcoat-
bondcoat interface, site 1 to 5 are shown from where the EDS analysis
was taken and reported in Table 5.2.
79
Fig.5.14: Delaminated topcoat of TiN modified sample (after 50 hours exposure)
demonstrating the regions which were broken away from the bondcoat.
80
xviii
Two regions rich in alumina and chromium oxides
Fig.5.15: (a) TiN modified sample (after 50 hours exposure) showing, patches (box)
of chromium-titanium phases having crystals like structure
(b)Schematic representation of delaminated topcoat showing layers of
alumina and chromium oxide and layer with chromium-titanium phase
81
Fig.5.16: XRD patterns comparing the scans of as-sprayed coating with the sample
exposed at 950ºC for 50 hours.
82
Fig.5,17: XRD pattern showing different phases formed in the delaminated coating
of TiN modified sample exposed at 950ºC for 50 hours.
83
Fig.5.18: TiN-modified samples and (b) Standard TBC samples, showing
photographs of top surfaces after intense thermal cycling
86
Fig.5.19: Cross-section of (a) TiN modified samples and (b) Standard TBC samples
showing delamination of topcoat
87
Fig.5.20: Cross-sections of (a) TiN modified samples and (b) Standard TBC
samples showing intensity of spalling near edges during cycling
87
Fig.5.21: Cross-sections of (a) TiN modified samples and (b) Standard TBC
samples showing cracks at the interface of topcoat and bondcoat
88
Fig.5.22: Cross-sections of (a) TiN-modified sample and (b) Standard TBC sample:
shows bondcoat cracking and the oxidation penetration into the substrate;
TiN-modified sample exhibits better resistance
88
Fig.5.23: (a) Standard TBC samples and (b) TiN modified samples, showing
samples’ warping at edges; the curvature increased with increase in
numbers of cycles
91
Fig.5.24: (a) Standard TBC samples and (b) substrate sample without coating,
showing that the sample without coating deformed in multiple directions
92
Fig.5.25: Schematic illustration to calculate the stress in the coatings by using
curvature
92
Fig.5.26 Residual stresses as a result of thermal cycling
93
Fig.5.27: Weight loss of the samples during thermal cycling
94
Fig.5.28 XRD patterns comparing the scans of Standard TBC coating with TiN-
modified coating samples after thermal cycling
94
xix
Fig.6.1: S1, S2 and S3 samples showing appearance of top surfaces after varying
exposure times, treated at 950°C in a hot corrosion environment
97
Fig.6.2: SEM micrographs showing surface morphology of samples (a) as sprayed
and (b) after 50 hours exposure to hot corrosion environment
98
Fig.6.3: SEM micrographs showing cross section of as-sprayed samples (a) S1 (c)
S2 (e) S3 with different bondcoat thicknesses (b,d,f) high magnifications
of bondcoats showing typical lamellar structure
99
Fig.6.4: SEM micrographs after 50 hours exposure to hot corrosion environment
showing, cross section of samples (a) S1 with thick TGO and cracks
(arrows) and (b) S3 with thin TGO
100
Fig.6.5: SEM micrograph showing discrete sites at the interface in a sample
exposed for 50 hours showing different features (Site-1 and 5). Site-1 to 5
represent locations for EDS analyses reported in Table 6.1
101
Fig.6.6: XRD pattern showing various phases formed in the delaminated coating
of S1 sample exposed at 950ºC for 50 hours
102
Fig.6.7: Samples exposed at 950ºC for 50 hours, indicating more oxidation in (a)
S1 with thin bondcoat as compared to (b) S2 with thick bondcoat
103
Fig.6.8: Samples exposed at 950ºC for 50 hours, EDS analysis at the boundaries of
the splats showed that “vanadium” is present in (a) sample S1, whereas,(b)
sample S2 revealed no “vanadium” in analysis
103
Fig.6.9: (a) S1 samples and (b) S3 samples, showing condition of top surfaces
after 225 intense thermal cycles
106
Fig.6.10: Cross-section of (a) S1 samples and (b) S3 samples showing delamination
of topcoat
106
Fig.6.11: Cross-sections of (a) S1 samples and (b) S3 samples showing intensity of
spalling near edges during cycling
107
Fig.6.12: Cross-sections of (a) S1 samples and (b) S3 samples, showing that after
spallation of the topcoat the bondcoat also started to spall near cracked
edges
107
Fig.6.13: S1 samples started to curve from edges during intense thermal cycling, it
was noted that curvature increased with increase in numbers of cycles
108
Fig.6.14: S3 samples started to curve from edges During intense thermal cycling, it
was noted that curvature increased with increase in numbers of cycles
109
xx
Fig.6.15: Substrate sample without coating, showing deformed in multiple
directions
109
Fig.6.16: Stresses calculated using the curvature of samples as a result of thermal
cycling for the two systems (S1 and S3 samples)
111
Fig.6.17 Weight loss of the samples (S1 and S3) during thermal cycling
111
xxi
List of Tables
Table 2.1: Thermal and mechanical properties of new potential TBC materials
15
Table 2.2: Process parameters for plasma spraying
21
Table 2.3: Process parameters for plasma spraying
22
Table 2.4: Comparison of properties (melting point & thermal expansion coefficient) 32
Table 3.1: Chemical composition of Inconal-X750 and substrate 34
Table 3.2: Grit blasting of substrate (Inconal-X750).
36
Table 3.3: Chemical compositions (wt. %), particle size range and morphology of the
spraying powders
37
Table 3.4: Air plasma spraying parameters for TBCs system
39
Table-4.1: Chemical composition of various phases marked in Fig.4.6 (exposed for 50
hours at 950⁰C)
53
Table-4.2: Chemical compositions of various phases marked in Fig.4.7 (exposed for
50 hours at 950⁰C)
54
Table 5.1: Average chemical compositions (% atomic) of various phases marked in
Fig.5.12 (TiN modified sample exposed for 50 hours at 950⁰C)
78
Table 5.2: Average chemical compositions (% atomic) of various phases marked in
Fig.5.13 (standard TBC sample exposed for 50 hours at 950⁰C)
79
Table-6.1: Chemical composition of various phases marked in Fig.6.5 (exposed for 50
hours at 950⁰C)
101
1
Chapter-1
Introduction
There has been a tremendous progress in the field of gas turbines over the last 70 years. This
owes itself to academic and industrial research efforts that were mainly focused on increasing the
efficiency of the engines through an increase in their working temperature. The operational
temperature limit of aero turbines and industrial gas turbines (IGTs) has been significantly
increased through advancements in such as single-crystal turbine blades and thermal barrier
coatings (TBCs). High temperature oxidation is becoming the prime reason for degradation of
turbines, due to ever increasing operational temperature demands. However, the research is
motivated by ever increasing market of IGTs [1-7].
Thermal barrier coatings (TBC) provide better thermal insulation to hot sections of gas turbines
which result in increased operating temperature and consequently a higher efficiency [1, 4-26]. A
typical TBC has a duplex structure, i.e., metallic bondcoat and a ceramic topcoat. The metallic
bondcoat, which generally is 70-150 µm thick, is sandwiched between the substrate and the
ceramic topcoat and protects the substrate from high temperature oxidation and corrosion [27,
28]. Moreover, it also improves the adhesion of the ceramic topcoat due to its rough surface that
provides a mechanical bonding and also by reducing the coefficient of thermal expansion (CTE)
mismatch between the substrate and topcoat [4, 29-31]. The classic bondcoats consist of Ni-Cr,
Ni-Al, Ni-Cr-Al, Ni-Al(Pt), MCrAIY (M = Ni, Co or Ni + Co) [1, 2, 6, 29, 30, 32]. Recently,
MCrAlY having (wt.%) Ni-22Co-12Al-18Cr-0.5Y has been shown to be a highly successful
material for bondcoat [33-43].
High melting point, high coefficient of thermal expansion, excellent thermal stability, good
erosion resistance and low thermal conductivity are the characteristic properties of the
insulating topcoats [6]. One of the best candidates that can be employed for the topcoat is based
on stabilized Zirconia. Zirconia can be stabilized by CaO, MgO, Y2O3, CeO2 etc. and among
them yatria stabilized zirconia (YSZ or ZrO2-Y2O3) has been recognized as an outstanding
material in the last few decades [2, 20, 29, 32, 44-46]. The main reason for YSZ as top
2
insulating coat material is its low thermal conductivity and high coefficient of thermal expansion
(10.7x10-6
/K) [46], which closely matches that of the Ni based substrate (12.6x10-6
/K) [47]. The
small mismatch in coefficient of thermal expansion (CTE) can be further decreased by
incorporation of the bondcoat (MCrAlY) [27, 28].
Another layer, that is present between the topcoat and the bondcoat in a TBC system, is the
thermally grown oxide (TGO). It forms during TBC deposition and its thickness increases during
normal operation by diffusion of oxygen through the topcoat [11, 12]. It acts as diffusion barrier
and reduces the speed of reaction between oxygen and the elements of the bondcoat, such as Cr
and Al. The α-Al2O3 is the most preferred phase of this layer, as it protects the bondcoat against
oxidation above 900°C [14]. The increase in the thickness of TGO is accompanied by stress at
the interface of bondcoat and topcoat. Sometimes this stress is more than the tolerance of the
TBC system, resulting in delamination of the coating at the interface.
There are two methods that are mostly used for deposition of both metallic and ceramic coatings:
(1) Air plasma spraying (APS); (2) Electron beam physical vapor deposition (EB-PVD) [1, 2, 30,
48]. Coatings produced through APS have a lamellar structure containing intra-lamellar and
inter-lamellar cracks and pores; the individual lamella consists of columnar grains. On the other
hand, EB-PVD grown coatings have columnar microstructures and possess better durability [2,
16, 32, 44]. However, APS technique is more versatile and cost-effective as compared to EB-
PVD process and, in addition to this, the coatings produced by this technique have low thermal
conductivity compared with those prepared by EB-PVD [49, 50]. Air plasma sprayed TBCs are
used in numerous engineering systems such as gas turbine engines (e.g. blades, vanes etc.),
diesel engines (e.g. piston head, valves), petrochemicals, heat exchangers, etc. [51, 52].
Industrial turbines and diesel engines employing low quality fuel are prone to damage from hot
corrosion at 700–900ºC. Low-quality fuels usually contain impurities such as sodium and
vanadium which can result in deposition of oxides/salts such as Na2SO4 and V2O5 on the surface
of the turbine components. These salts, in their fused state, can react with YSZ and may
transform tetragonal zirconia to monoclinic phase during cooling. This transformation is
accompanied with 3-5% increase in volume, leading to cracking and spalling of TBC [53, 54].
The life of the bondcoat is limited due to oxidation/corrosion and also by inter-diffusion between
3
the substrate and bondcoats. Hot corrosion then becomes predominant and drastically reduces the
life of the TBC.
The oxidation behavior of the bondcoat strongly determines the life of the TBCs. The
oxidation/hot corrosion resistance of the bondcoat is affected by several factors, including: (i)
bondcoat and substrate materials; (ii) bondcoat and substrate inter-diffusion; (iii) characteristics
and morphologies of the thermally grown oxides (TGO); (iv) phase transformation and cracking
of the coating [55-60]. MCrAlY coatings have been studied extensively for the last two decades
[39-42]. Significant amount of work was carried out on development of MCrAlY bondcoats
resistant to hot corrosion [33-37, 39-43]. Literature reveals that efforts to improve hot corrosion
resistance showed inconsistent results, limiting the understanding of hot corrosion behavior of
TBCs [54, 61]. Additionally, the role of elemental diffusion from substrate to bondcoat and vice
versa have not been discussed in detail, which is considered to play an important role in
degradation of various coating systems [54, 62].
In the current thesis, Yttria Stabilized Zirconia TBCs along with CoNiCrAlY bondcoat was
deposited using air plasma spray on Inconel-X750 superalloy and subsequently tested for hot
corrosion behavior. The reaction products were microscopically investigated and a detailed
elemental diffusion of different alloying elements was investigated. An attempt has been made to
enhance the life of TBCs by applying thin layer of TiN (by physical vapor deposition) on
bondcoat for improving the oxidation and hot corrosion resistance of bondcoat. Effect of
bondcoat thickness on the hot corrosion of TBCs was also studied. The durability of coatings
was evaluated by thermal cycling.
1.1 Objectives of Work
The objectives of this research work are:
1. To produce Yttria stabilized zirconia thermal barrier coating (TBC) along with CoNiCrAlY
bondcoat on Inconel-X750 superalloy by employing the conventional, flexible and low cost air
plasma spraying.
2. To study the hot corrosion behavior of locally deposited TBCs exposed at high temperatures
and investigation of the reaction products. Further, to examine the detailed elemental chemical
4
composition profile of all samples to analyze the diffusion behavior of various alloying elements
of bondcoat and substrate.
3. Introduction of modified bondcoat in TBCs by applying thin layer of TiN on bondcoat to
improve the oxidation and hot corrosion resistance and make comparison with standard TBCs by
studying their microstructures and evolution of phases.
4. To evaluate the hot corrosion and oxidation behavior of TBCs in term bondcoat thickness.
5. Thermal cycling tests of both the TiN-modified bondcoat-thermal barrier coating system and
TBCs with thick bondcoat, for the assessment of their durability.
1.2 Outline of Thesis
“Chapter 2 (Literature survey) Summary of previous research work done about the thermal
barrier coatings (TBCs) systems for the references which comprise brief description of high
temperature coatings, characteristic properties of TBCs, hostile environments, methods and
materials used for depositions of TBCs, oxidation and hot corrosion of TBCs and their
estimation.”
Chapter 3 (Experimental) Validates the experimental work carried out for the deposition of
TBCs systems, hot corrosion testing, heat treatment used for rapid thermal cycling and finally
techniques used for coatings characterization and tools used to evaluate the TBCs systems.
Chapter 4 (Result & discussion) Yttria stabilized zirconia thermal barrier coating (TBC) along
with CoNiCrAlY bondcoat was deposited using air plasma spray on Inconel-X750 superalloy.
The coated samples were exposed to a mixture of Na2SO4 and V2O5 at 950⁰C. The exposed
specimens were investigated using XRD and SEM. The formation of spinel and perovskite
structures was revealed at the interface of topcoat and the bondcoat. Further, the chemical
composition profile of all samples helped to analyze the diffusion behavior of different
constituent elements of bondcoat and substrate. XRD analyses of the samples confirmed phase
transformation of the tetragonal zirconia into monoclinic zirconia and yttrium vanadate. The shift
of high angle peaks indicated lattice distortion, which was directly related to the stresses in the
coating.
5
“Chapter 5 (Result & discussion) In this chapter, standard TBC samples were compared with
TiN modified bondcoat TBC samples. Titanium nitride was deposited on bondcoat by utilizing
physical vapor deposition technique. Both TBC systems were exposed to high temperature under
the presence of corrosive salts i.e. a mixture of V2O5 and Na2SO4 (50:50) for 50 hours. It was
observed that TiN modified samples showed better results in terms of oxidation resistance and
delamination. The formation of Cr2Tin-2O2n-1 phases at the interface of topcoat-bondcoat, in TiN
modified samples, were found responsible to enhance the thermal and oxidation properties of the
TBC.”
Chapter 6 (Result & discussion) The effect of bondcoat thickness on the hot corrosion resistance
was studied. Hot corrosion test were carried out in 50 wt:% Na2So4+50 wt:% V2O5 molten salt at
950ºC for 50 hours. The characterizations of the coatings included X-ray diffraction analysis,
scanning electron microscopy and optical microscope. Results indicated that TBCs with thick
bondcoat exhibited superior hot corrosion resistance to the TBCs with conventional bondcoat.
Chapter 7 (Summary) Results are summarized in terms of major conclusions and findings.
Chapter 8 (References) All references are given in this chapter
6
Chapter-2
2 Literature Survey
In industry the components may face highly aggressive environment (which is also oxidizing and
corrosive in nature) i.e fluctuating stresses due to mechanical loads, elevated temperature, high
pressure, etc. Turbine engine used in power generation, to produce electricity, marine propulsion
and aircraft are common examples of these components used in these processes [3, 15, 63].
Engineers made it possible that metallic parts exposed to such harsh environment, should have
the proficiency to withstand without failure. Stainless steels and superalloys are the contestant
materials for gas turbine engines and diesel engines. Since 1965 significant efforts have been
made to improve the mechanical properties at high temperature by altering the alloys,
compositions [4, 5]. Fig. 2.1 presents a concise summary of development of these alloys over the
years from wrought to cast, then to directionally solidified alloys and single crystal materials [7,
64]. Although superalloys with excellent hot corrosion resistance and outstanding high-
temperature strength, have been in use for high temperature applications, however, sometimes
fail under extreme conditions where temperature reached close to their melting points [3, 6].
Fig.2.1: Increase of operational temperature of turbine components made possible by alloy
development and TBCs [63].
7
Fig.2.1 demonstrates that mere strengthening the alloys cannot perfectly contribute to cope with
extreme conditions and avoid components’ destruction. The potential solution to enhance and
prolong the life of materials is to bring innovations in ‘coating technology’.
The provision of coatings and methods to ensure that applied coatings remain intact during
service conditions has become a critical issue in the field of gas turbine. The components like in
the combustor and turbine sections would degrade quickly if proper coating has not been applied
[65]. Fig.2.2 summaries the various coating types [66, 67] and subsequent relative improvement
in component’s life by specific coatings.
Fig. 2.2: Relative service temperature enhancement as a result of improvement in coating
technologies [66].
8
2.1 Diffusion coatings
The most common form of surface protection is diffusion coatings. For example, aluminum is
commonly deposited by this technique onto the surface of the superalloys by pack cementation
process. In this process, the component is packed in aluminum powder with some activators and
subsequently heat treated to promote interdiffusion by forming a typical β–NiAl phase [68] on
component’s surface, subsequently improving oxidation resistance of the alloy. It has been
demonstrated that prior deposition of 5-10 μm thick platinum layer before aluminization
improves high-temperature oxidation and hot corrosion resistance of the alloy [69]. Superalloys
are known to be protected by aluminide or platinum-aluminide coatings. In fact platinum
aluminide coatings have become standard industrial treatment for many high temperature
materials.
2.2 Overlay coatings
It had been demonstrated that overlay coatings provide a better oxidation resistance then
diffusion coatings by virtue of unique chemical composition; these coating are generally
deposited by air or vacuum plasma spraying (APS/VPS) [70], or also by electron beam physical
vapour deposition (EB-PVD) [71]. The MCrAlX-type coatings in various compositional
variations have become standard overlay coatings for gas turbine applications [72, 73]; ‘M’
refers to Ni or Co or combinations of these, and ‘X’ is a ‘reactive’ element (or mixture of them)
at concentrations up to 0.5 wt.% (yttrium, hafnium and silicon are commonly used). The overlays
have greater flexibility of composition; chemical and mechanical behavior is largely independent
of the substrate as compare to diffusion coatings which inherently strongly depend on substrate’s
composition.
2.3 Thermal barrier coatings
The thermal barrier coatings (TBCs) are designed to withstand the harsh environments and thus
find a widespread usage in turbine engines [3]. These coatings system protects the metallic
hardware from heat, wear, oxidation, and corrosion [6, 18-20, 74, 75]. The TBC makes it
possible to improve the life of component by reducing the exposed surface temperature up to
200ºC, depending upon the thickness of coating (Fig.2.3); these coating systems also help to
9
improve the engine performance by allowing higher gas temperature (~100–300ºC) resulting in
improved engine thrust and efficiency up to 5% and 1%, respectively [4-7, 15, 18-26, 64].
Fig. 2.3: Temperature drop substrate’s surface by YSZ coating [74].
A typical thermal barrier consists of a ceramic layer which when deposited on a superalloy,
provides thermal insulation and lowers the substrate’s temperature. The first thermal barrier
coating system developed in the 1960s; it was produced by plasma spraying of calcia- or
magnesia-stabilized zirconia [76]. These coating performed well, below ∼1000ºC, whereas at
10
higher temperatures they became unstable by formation of MgO and CaO as a result of diffusion
of Mg2+
and Ca2+
ions. The formation of the monoclinic form of zirconia was promoted causing
a four-fold increase in the thermal conductivity [76].
Modern TBC’s are required to design in a way that they not only minimize the heat transfer but
also protect the underlying hardware from oxidation and hot corrosion. No single coating can
meet these multifunctional requirements. The last two decades’ research has led to preferred
coating systems consisting of three distinct layers [18] to achieve required characteristics i.e long
term effectiveness at high temperatures and, oxidative and corrosive environment for which they
are intended to be used (Fig.2.3).
Bondcoat–It serves as an oxidation and hot corrosion resistant layer which is meant to protect the
metallic components from aggressive environment. This layer is required to remain relatively
stress-free, chemically stable and remain adherent to the substrate during long thermal exposure
to avoid premature failure of the TBC system. This layer must also provide an adherent surface
for the ceramic topcoat. Normally, a thin aluminum rich oxide (< 1 µm) thermally grown upon
the bondcoat is utilized for the protective purpose [20]. Since modern nickel based superalloys
are not typically high enough in the aluminum content to form a fully protective alumina scale
therefore the bondcoats are designed in a way that thermally grown alumina oxide may form
[75]. A 50-70 µm thick layer of platinum aluminide [21] or MCrAlY (where M is Ni or Co) [22]
are developed for this purpose. However the applied coating systems are expected to be thin and
low in density in order to limit centrifugal load exerted on rotating engine components and
therefore have good thermal and mechanical compatibility.
Topcoat-It is a thermally stable TBC layer with low thermal conductivity in order to provide
maximum thermal drop across the thickness of coating. Since topcoat is a ceramic layer which
may have a thermal expansion coefficient different from the substrate to which it is applied.
Therefore, the layer should have a high in-plane compliance to accommodate the thermal
expansion mismatch between the TBC and the underlying the superalloy. Moreover, the top
layer must have a property to retain the above characteristics, to have low thermal conductivity
and to demonstrate thermal stability during prolonged high temperature exposure [7].
11
Modern TBCs that are commonly used to protect the high temperature components like turbine
blades, nozzle guide vanes and other combustor sections components are based upon zirconia
containing about 7 wt.% Yttria, known as yttria-stabilized zirconia (YSZ). This composition was
originally identified in 1970s by NASA after giving the best thermal cycle results conducted on a
burner rig [77, 78]. Enhancing TBC’s life and the design of procedures to estimate it are
currently main challenges faced by materials scientists and engineers working in the gas turbine
field [8].
2.3.1 Ceramic materials for TBCs
It is well known that ceramics have low thermal conductivity but to select an appropriate system
TBC remains a question. Fig.2.4, presents thermal conductivities of few ceramic compounds [3,
79].
Fig.2.4: Thermal conductivities of various polycrystalline oxides in as a function of
temperature [3, 79]
12
The thermal conductivities of various ceramics range from 1 to 30 W/(m K), whereas these
values decrease with increasing temperature; zirconia (ZrO2) demonstrates the lowest
conductivity of 2 W/(m K). Moreover, no strong temperature dependence associated with it,
therefore it cannot suffer from the drawback of a sharply increasing conductivity at lower
temperatures. These findings explain why ZrO2-based ceramics have become a choice for TBCs
applications.
However, pure ZrO2 cannot be used directly for TBCs due to its phase transformation from
tetragonal to monoclinic (‘t’ to ‘m’) phase; this change is associated with 4% volumetric change
and may exhibits poor thermal cycling resistance. Modern TBCs are thus stabilized with yttria
(Y2O3) 6 to 11% by weight [80]. YSZ has become a standard material due to its low thermal
conductivity and relatively high thermal expansion coefficient as shown in Fig. 2.5 [24]. This
significantly reduces the mismatch of thermal expansion.
Fig.2.5: Thermal conductivity vs. thermal expansion coefficient of various ceramic materials
[24].
13
YSZ also presents good erosion resistance which is important for engine components since high
velocity particles may cause impingement [81]. The chemical composition Zirconia stabilized
with 7% Yttria turned out to be optimum for forming metastable (non-transformable) tetragonal
(t/) phase when the ceramic is quenched from the cubic phase field present at high temperature
(Figure 2.6). The t/ phase is an equilibrium tetragonal polymorph having an entirely changed
microstructure containing of anti-phase domain boundaries and numerous twins [82]. When
Y2O3 is combined with ZrO2 it gives t/ phase, the resulting solution is known as ‘partially
stabilized zirconia’; the use of the adjective ‘partially’ distinguishes it from ‘fully stabilized
zirconia’ which formed at higher Y2O3 concentration. For this reason, partially stabilized zirconia
is used as the standard for TBC material. This can be deposited by air Plasma Spraying (APS)
and by Electron Beam Physical Vapor Depositing (EB-VPD) methods. However, Yttria
stabilized zirconia is thermally stable up to 1200ºC, beyond this temperature it exhibits two
problems.
Fig.2.6: Yttria - Zirconia phase diagram. Note that the shaded region indicates the region where
the formation of the metastable t’ phase occurs upon cooling. [24].
14
(1). partial reversion from non-equilibrium (t/ phase) to an equilibrium ‘t’ phase; the tetragonal
(t) phase tends to transform to monoclinic upon subsequent thermal cycling [83].
(2) prolonged exposure to high temperatures may cause sintering of the YSZ, subsequently
increasing the elastic modulus [84]. The combination of both effects deteriorates the thermal
cycling properties.
Much research has been carried out to explore the new chemical composition distinct from the
7% yttria stabilized zirconia. This is because the demand of high temperature stability and low
thermal conductivity are increasing. For example, Ceria (CeO2) is recommended for Yttria
(Y2O3) to stabilize the zirconia, since no monoclinic phase was observed up to 1400°C and
during limiting exposure to 1600°C, but unfortunately, the ceria stabilized zirconia did not show
good erosion resistance [85]. Similarly, some researcher demonstrated that 90% substitution of
Scandia (Sc2O3) with Yttria, so called scandria-yttria stabilized zirconia has been to have
significantly better tetragonal (t/) phase stability at 1400°C [86, 87]. Alternatively, ZrO2 based
TBCs can be replaced entirely with HfO2. HfO2-Y2O3 TBCs have been demonstrated comparable
thermal cycling resistance to 7% YSZ but to higher Y2O3 (i.e upto 27% wt.) when the crystal
structure is fully cubic [88]. Further, lanthanum zirconate (La2Zr2O7) demonstrated the
pyrochlore crystal structure which have an excellent thermal conductivity (1.6 W/mk), has been
proposed as a TBC material [89, 90], Similarly other rare earth zirconates based on Gd and Sm
can be a good candidates for TBCs [91]. Finally lanthanum hexa-aluminate, which has
“magnetoplumbite” structure, is considered as a good competitor to partially stabilized zirconia
for operations above 1300°C, due to its good resistance against sintering [92]. The science of
these ceramics for thermal barrier applications has been reviewed recently in [93]. Pratt and
Whitney claim to be using Gd2Zr2O3 based TBCs for niche applications in military turbines. The
major problem in new ceramic materials have inadequate resistance to erosion which can occur
due to foreign object damage (FOD) accumulation and sand ingestion. Table -2.1 , shows the
thermal and mechanical properties of some widely used TBCs materials reported by Robert et
al.[21].
15
Table 2.1: Mechanical and Thermal properties of potential TBC materials [21].
Materials Coefficient of
thermal
expansion at
30°C – 1000°C,
10−6
/K
Thermal
conductivity
at 1000 °C,
W/m K
Melting
temperat
ure °C
Fracture
toughness,
MPa m1/2
Young's
modulus,
GPa
YSZ 11.5 2.12 2680 1–2 210±10
Perovskites
Zirconates
BaZrO3
7.9 3.42 2690 181±11
SrZrO3 10.9 2.3 2800 1.5±0.1 170±4
CaZrO3 8.4-8.9 2.0 2550
Complex form
Ba(Mg1/3Ta2/3)O3 10.9 2.71 3100 ~0.7 186±2
La(Al1/4
Mg1/2Ta1/4)O3
9.7 1.82 ~0.8 174±2
BaLa2Ti3O10 10-13 0.7
Hexaaluminates
LaMgAl11O19 9.6 2.6 130±11
GdMgAl11O19 9.5 2.7
Gd0.7Yb0.3MgAl11O1
9
9.6 1.9
LaLiAl11O18.5 10 3.9
Pyrochlores
La2Zr2O7 9.1
1.56 2300 175
Gd2Zr2O7
10.4 1.6
TBCs Cluster 11.6-13.5 1.7-2.1
16
2.3.2 Bondcoat Materials
In thermal barrier coating system, bondcoat is one of the most important components. Its choice
is very conclusive in determining the resilience of TBCs. It basically defines the spalling
characteristics of TBC system [3]. It has been noted that the mismatch of thermal expansion co-
efficient between the metallic substrate and the depositing ceramic can develop high interfacial
shear stress, during spraying process. It is also observed that the ceramic coat delaminated during
the operational thermal cycling [44, 94]. Therefore, an intermediate layer between the metallic
substrate and the ceramic coat may decrease the miss-match of co-efficient of thermal expansion
and extends the life of TBCs system [6, 15, 29, 30]. Further ceramic topcoat in TBC is usually
porous in nature and allows oxygen to pass through at elevated temperatures. Therefore,
bondcoats must have a good oxidation and corrosion resistance for the protection of under laying
substrate. Further, bondcoat strengthens the bonding between ceramic topcoat and the substrate
[6, 15, 19, 49, 95-98]. In thermal barrier coating system, a thermally grown oxide (TGO) layer is
produced due to the oxidation of oxidizable elements present in bondcoat. The nature of TGO
depends upon the chemical composition of bondcoat. This TGO grows at the interface of
bondcoat and topcoat. It has been reported that the morphology, thickness and microstructure of
TGO greatly influence the durability of the TBC system, which linked with chemistry and
microstructure of the bondcoat. [1, 6, 8, 15, 29, 94, 96, 98]. The life of TBC system depends
upon the stress/strain developed during the coating process and the bonding of TGO to bondcoat
[99]. Wortman and Miller et al [96] demonstrated that strong bondcoat with good creep
resistance can improve the durability of the TBC during the thermal cycling. Other characteristic
features of good bondcoat are that it prevent the formation of brittle phases and inhibit elemental
interdiffusion between the metallic substrate and the bondcoat. In this regard element like Mo,
Ta and S diffuse towards from the substrate at high temperature and can create voids close to the
interface either by themselves or by their oxides which can damage the bonding of the coatings
[96, 99]. The most commonly used bondcoat materials, used in TBCs are MCrAlY (where M=
Ni or Co or both Ni and Co), NiAl, Ni(Pt)Al etc. [1, 2, 29, 30, 95, 100].
2.3.2.1 MCrAlY System
Various TBC system employ MCrAlY (where M = Co, Ni or Ni+Co) overlay coating as a
bondcoat. MCrAlY systems used for dual purpose i.e it protect the underlying hardware from
17
high temperature corrosion and improve the adhesion of ceramic coat [6, 15, 96, 98-100]. In
MCrAlY, Cr (20-30 wt. %) and Al (5-10 wt.%) forms oxides, when exposed at high temperature
(˃ 900°C), and inhibit the further oxidation of the metallic materials, whereas Y stabilizes these
protective oxides [2, 96, 99-101]. It is described that the presence of elements like, Y, Hf and Zr
in MCrAlY play a vital role in improving the oxides bonding e.g Hf and Y help to form fine
grains nucleation sites [99]. . Microstructure of MCrAlY consists of gamma phase (matrix) and
β-phase [63]. Gamma phase is consisting of Ni or Co while β-phase is of Ni/CoAl. Some other
phases are also reported which include γ’- Ni3Al, α-Cr, Ni-γ and σ-CoCr etc. [102-104], the
formation of these phases depends upon the exposure temperature. Formation of various oxides
at high temperature are also reported i.e α-Al2O3, θ/γ- Al2O3, Cr2O3, NiO, (Ni, Co) (Cr, Al)3O4
etc. [105, 106] and spinel like Ni(Al, Cr)2O4 also formed [107]. Spinel are compounds having
AB2O3 or AO.B2O3 (A2+
and B3+
) stoichiometry formula [108-110].
2.3.3 TBCs Methods of Deposition
Thermal barrier coating are frequently deposited by following methods [1, 2, 32, 94, 111-118].
1. Air plasma spraying (APS) method
2. Electron beam-physical vapour deposition (EB-PVD) method
APS and EB-PVD methods are used to deposit ceramic topcoat in the TBC system while
metallic bondcoat can be deposited by APS, HVOF (High velocity oxy-fuel), EB-PVD, LPPS
(low pressure plasma spraying), VPS (vacuum plasma spraying) and CVD (chemical vapour
deposition) etc. [94, 95].
APS and EB-PVD method s are used to deposit ceramic topcoat in the TBC system [125] while
metallic bondcoat can be deposited by APS, HVOF (High velocity oxy-fuel), EB-PVD, LPPS
(low pressure plasma spraying), VPS (vacuum plasma spraying) and CVD (chemical vapor
deposition) [111].
2.3.3.1 Plasma Spraying
Plasma spraying used for depositing TBCs is the most common method and has been used since
1950s for enhance the efficiency and protection of hot section components e.g. flare head,
18
blades, vanes and primary zone sections of combustors within turbines, diesel engines (pistons
and valves) [119-124]. Plasma spraying is the most flexible and adoptable technique to deposit
high performance metallic and non-metallic materials. This method of coating deposition has
been utilized for many decades in different industries all the world [126]. Plasma spraying
method has several aspects to understand. This includes formation of plasma, injection of
powders, heating and propulsion of powder particles, impact of semi-molten/molten particles on
the substrate and powder morphology i.e. shape, size and distribution [127].
Advantages of plasma spraying [96, 123, 128]
Following are the advantages associated with this technique.
1. Variety of materials can be deposited by this technique. This includes ceramics, pure
meals, alloys, cermet and plastics.
2. Excellent feature of this technique is that the chemical composition of deposited coating
remains unchanged as of original powder.
3. This technique produces coating having appreciable bond strength with the underlying
substrate.
4. Plasma spraying is frequently utilized to spray high melting point materials.
5. Very high deposition rate can be achieved with plasma spraying technique e.g ˃ 4Kg/h.
6. One of the most significant advantages of this technique is that it can be utilized for any
shape and size of component. It can coat the internal and external surfaces.
7. Plasma spraying can be carried out in different environments such as air, inert gas and
under vacuum.
2.3.3.1.1 Plasma Spraying Process
Various steps are involved to deposit the spraying material on the substrate.
(i) Surface preparation
(ii) Plasma formation
19
2.3.3.1.1.1 Surface Activation
The surface preparation is an important step for plasma coatings. The surface of substrate
was activated so that it can adhere the coming semi-melted particles. In this regard, the
substrate surface thoroughly cleaned and then roughened. The most common means
employed for this purpose are grid blasting, water jet treatment, laser ablation and chemical
process. Grit blasting, however, is more popular for its ease and efficiency. In this technique
sharp and angular abrasive particles are bombard, at high speed, on the substrates which upon
impact create roughness. Grit blasting can serve both the purpose of cleaning and
roughening. The coating and the substrate adhesion during plasma spraying reply primarily
upon the mechanical bonding. In other words, rough surface increases the surface area
(particle to substrate) and crates the surface irregularities into which semi-molten particles
can anchor to substrate.
2.3.3.1.2 Plasma Formation
Plasma is the fourth form of matter and can be formed by ionization of gases. In plasma
spraying, the plasma is formed mainly by argon gas, which is ionized by using electric arc. This
arc is produced when D.C current is applied between thoriated tungsten cathode and copper
anode[63, 114]. An inert carrier gas usually, Nitrogen, is used to insert the coating material in the
form of powder in the producing plasma. The producing plasma arc, produced between the
cathode and anode, having high enthalpy and kinetic energy, act like a heat source and help to
melt down the injecting powder very quickly. In case of ceramic these particles are in semi-
molten condition. Since the kinetic energy of the producing plasma is very high it propelled the
molten/semi-molten particles supersonically towards the substrate [129]. These particles when
struck with relatively cold substrates are quickly solidified to build up coating which is primarily
mechanically bonded to the rough surface of substrate [114, 125]. Fig. 2.7 demonstrates the
schematic diagram of plasma sprayed process.
2.3.3.1.2.1 Plasma Spraying Parameters
The properties of the final sprayed coating greatly depend upon the number of spraying
parameters and the materials being sprayed. The plasma spraying parameter are described in
various literatures [124, 130, 131]. According to Nusair.A.Khan te al [130], the plasma spraying
20
parameters can be divided into three classes which are further sub-classified in various groups
(Table 2.2). Similarly, Lugscheider et al [132] reported that the various spraying parameters
(Table 2.3) which affect the quality of the coating.
Fig. 2.7: Schematic illustration of plasma spraying process [96].
Anode Electrical insulation
Powder injection
Electric arc
Plasma
forming
gas
Cooling water P
lasm
a
flam
e
Cooli
ng
wat
er
Cathode
21
Table 2.2: Parameters used for plasma spraying Technique [130].
Parameters of process Remarks
Plasma formation
Plasma formation comprises
Plasma gun design
Currents and voltages,
Flow rates of plasma producing gases and their
Composition
spraying or standoff distance (Distance of the
plasma gun from the substrate)
Flow rate of carrier gas and its nature.
Powder and powder
transportation
It includes
The sizes and shapes of powder used
Powder chemical composition and residence time
which is called dwell time
Temperature of beam used in process.
Enveloping atmosphere
It defines
The nature of oxidation of powder.
The composition and pressure of gases used during
coating process
The temperature and length of plasma jet used
22
Table 2.3: Parameters for plasma spraying technique [132].
Parameters Remarks
Nozzle and burner chamber
The proficiency of nozzles and burner chamber depends
on different parameters i.e nozzle geometry, plasma gas,
mass flow rate of plasma gas, cooling , Mass flow rate
cooling fluid and power supply
Feed of powder
Powder feed is also important which depends upon
thermal properties of powder material, powder fraction
and shape, injection geometry, mass flow rate carrier gas
and carrier gas
Plasma jet
The characteristics of plasma jet depends upon the jet
velocity and temperature, Particle velocity and
temperature and Particle trajectory
Particle impact
The particle impact determine the quality of coating and
effected by impact distribution, velocity at impact, particle
impact angle, molten state of particle at impact, substrate
type and substrate temperature
2.3.3.1.2.2 Morphology of Spraying Powders
Morphology of the spraying powder is one of important parameter which defines the
characteristics of the final coating. Morphology of the powder includes the shape, size and size
distribution. The size of powder particles are usually in the range of 40±10 µm [63], where the
powder size deviation from -45+5 µm to -25+5 µm and -100+45 µm to -100+10 µm is also
reported [128]. Fine particles are sprayed to produce dense coatings while coarse powder is used
23
to produce porous coatings, as are required in case of TBCs [96, 128]. Oversized particles are
usually remain un-melted or in semi molten condition during the stay time in plasma stream
[133] Fine particles i.e less than 10 µm (if are not dense) typically fail to infiltrate into the hottest
zone of the plasma beam or may quickly evaporate before reaching the targeted substrate
surface, and therefore, have less contribution in the deposited coating [123]. Particle shape
control the flow properties of the powder and thus also an important parameter in controlling the
finial coating characteristics [128]. Additionally, the powder particles have wide variation in the
shapes due to the different methods used for their production.
2.3.3.1.2.3 Plasma Gases
There are two types of plasma gasses:
1. Monoatomic gases (Argon, Helium etc.)
2. Diatomic gases (Nitrogen, Hydrogen etc.)
Fig. 2.8: Energy contents of various plasma forming gases as a function of temperature
[123].
Gas Temperature (oC)
E
NE
RG
Y C
ON
TE
NT
OF
GA
S (
kcal /m
ole
)
24
The selection of plasma gasses depends upon the required temperature and velocity in the plasma
plume [124]. The energy contents, of some plasma making gases, as a function of temperature
are shown in Fig. 2.8 [123, 124]. The relation between the energy contents and the temperature
of the gasses is not ideally linear because of the ionization and dissociation behavior of the gases
during plasma making. Diatomic gases are first dissociated and then ionized. Whereas,
monoatomic gasses are directly ionized. This demonstrates that diatomic gasses need more
energy to enter into the plasma state which intron increases the enthalpy of the plasma [123].
Similarly, when the metastable state of ion, with in the plasma, recombine into atomic and
molecules of gasses it generate heat energy. The latter energy is utilized to meet down the solid
powder particles for plasma spraying [124].
2.3.3.1.2.4 Microstructure of Plasma Coatings
Complex microstructure bas been developed in plasma spray coatings. The coatings consist of
splats, micro cracks, porosity and un-melted particles. Splats are formed after striking the semi-
molten powder particles on the substrate and described as lamellar structure. The trapped air
during spraying caused voids, Fig. 2.9a [117]. Oxides and inclusion are also present in these
coatings which are formed during interaction of powder particles with oxygen at high
temperature. While un–melted particles are generated due to presence of larger or oversized
particles or may be owing to injection of powder particles into the colder region of plasma
plume, Fig. 2.9b. This type of microstructure can make these coatings less stiff but high strain
tolerant [2, 3, 29, 32, 95, 112, 114]. Moreover, this type of microstructure compromise low
thermal conductivity (0.8-1.7 W/mK) as a result of porosity and voids between the lamella [95].
For the duration of the plasma spraying millions of semi-molten and molten particles, in various
sizes, are propelled towards the substrate per second. This built up coating as shown in Fig. 2.9b
[70]. As soon as molten particles strikes the surface of substrate it spreads out and deformed to
procedure a splat [125]. Attachment of deposited particles during solidification is happening by
setting onto the asperities on the substrate surface and then theses splats inter lock with one
another. It is proven that there is a association between the microstructure and the properties [70,
96, 123, 124]. these properties depends upon various factors and summarized as follow [123]:
1. Motion of plasma gun and substrate.
25
2. Rate of cooling of the coated material and the substrate
3. Velocity of particle and its temperature during the coating process.
The coating deposited in plasma spraying are fabricated particle by particle [134]. The time for
the solidification of molten particles, after striking the substrate, is less in comparison to the time
of arrival of molten particles from the gun, consequently, this reduces the encounter time
between the liquid pools. Furthermore, the solidification time of molten particles is about two
times longer and thus supports the diffusion and stress relief processes to some degree [96, 123,
134]. The porosity of the coating also depends upon wetting and flowing properties of the molten
particles. Further it is also affect the adhesion and morphology of coating-substrate interface
[123].
Fig.2.9: Thermal spray coating (a) plasma sprayed as-sprayed coating [117] (b) coating build up
[70].
(a) (b)
26
2.3.3.1.2.5 Adhesion in Plasma Spray Coatings
The most important condition for an extraordinary quality plasma spray coating is good bonding
to the substrate. The adherence ability of the coating in plasma spray may be defined as, ‘it is
the force required to tear off a unit area of the coating from the substrate’. The bond formation
between the coating and substrate is somehow multifarious in nature during plasma coating;
following types of bonds are thought to establish adherence [63, 70, 96, 123, 124].
1. When molten/semi molten particles are deposited on the substrate surface their
mechanical interlocking and anchorage into the substrate asperity takes place.
2. The physical interactive forces also put share in the coating adherence.
3. The micro-welding of particles which involves the formation of covalent or metallic
chemical bonding.
The other factors which affect the bonding of the coating are:
1 Existence of residual stresses inside the coating.
2 Existence of localized alloying due to melting at the contact surfaces between the
particles and substrate.
3 Elemental species Diffusion across the boundaries of splat.
Depending on the nature of the coating the acute role of the specific bond type is changed,
substrate material and spraying parameters [96, 124]. It is described in literature that
presumably for the metallic coatings deposited on the metallic substrate metallic bonds exist
whereas in case of the ceramic coatings deposited on the metallic substrate mechanical bonds
are preferably made [124].
27
2.3.3.2 Electron Beam Physical Vapors Deposition
“The distinguishing feature of the Electron beam physical vapor deposition (EB-PVD) process
is the use of an electron beam which is used to vaporize an ingot of the coating material, that is
typically held in a water-cooled copper crucible; this causes a vapour cloud to be made above
the ingot in which the component for coating is handled. The schematic of the electron beam
process is shown in Fig. 2.10a [114]. As the coating material, start to evaporate, the feedstock is
fed into the chamber to sustain a persistent feedstock height it is important to generate a coated
surface completely made from molten vapors, free of splashing so that the coating process is
stable. Therefore, in thus process no chemical reactions are involved. Both the ceramic and
metallic coatings can be deposited by EB-PVD process e.g. MCrAlY, Zirconia, Zirconia
stabilized, etc. [96].
This is comparatively expensive and complex method of TBCs deposition and is usually used
for coating on highly stressed gas turbine components due to its exceptional strain tolerant
columnar structure and good corrosion resistance [45, 117, 118, 135]. On the other hand,
thermal conductivity of EB-PVD coatings is high (twice) as compared to APS coatings [136].
By design, the resulting morphology then consists of a series of columnar colonies (Fig. 2.10b)
which grow competitively in a direction perpendicular to the surface of the substrate [2, 32,
112]. The bonding strength of EB-PVD coated YSZ with the substrate is 10 times higher than
APS [114, 137].”
2.3.4 Oxidation and hot corrosion of TBCs
“In hot corrosion, the metallic elements of a material are converted into their oxides when the
material is exposed to high temperature in oxygen or oxygen-containing environment. Corrosion
and erosion of a materials may occur when oxygen and some other environmental constituents
such as CO2 and SO2, fused or molten salts and sand simultaneously act on it [138]. This
phenomenon is in focus as serious problem boilers, power generation equipment, gas turbines,
internal combustion engines, fluidized bed combustion, industrial waste incinerators and the
paper and pulp industries, since 1940s. However, it was not until 1960s, that the problem got
more attention when gas turbine engines in military aircraft suffered severe corrosion while
operating over the sea during the Vietnam conflict [139-141].”
28
.
Fig. 2.10: The EB-PVD process (a) schematic of the electron beam process (b) TBCs deposited
by EB-PVD [63].
(b)
(a)
29
“In hot corrosion, molten salts attack the material in oxidizing environment at high temperature
[142-145]. It occurs when a metal or alloy is covered with salt films at a temperature typically
between 700 and 900°C. It is reported in literature by Hancock [146] and Eliaz et al [147] that,
hot corrosion is an accelerated form of oxidation that occurs when metals are heated at a
temperature of 700–900°C in the presence of sulphate deposits. The sulphates themselves are
produced through a reaction between sodium chloride and sulphur compounds present in the gas
phase that surround the metal. Corrosion which occurs above the melting point of the salt is
called ‘type-I hot corrosion’, while that which takes place at the lower end of the temperature
range is called ‘type-II hot corrosion [138]. The temperature range of the said types of corrosion,
as measured in terms of metal loss under the influence of temperature, is shown schematically in
Fig.2.11 [138]. In both types of corrosion, the corroding salts render the protective oxide scale,
which forms on superalloys and coatings, as useless owing to their fluxing action. The hot
corrosion takes place in two stages. In first stage; the initiation involves the breakdown of
protective oxide scale, while the second stage of corrosion ensues involving attack in the form of
oxidation and sulphidation at exceedingly high rates.”
Fig. 2.11: Schematic illustration of temperature effect on rate of damage to superalloys based on
type I and II hot corrosion superimposed on contribution due to oxidation [138].
30
“The degradation of a TBC system may take place by a number of mechanisms. The deposits
may affect zirconia topcoat as well as the bondcoat, chemically or mechanically [61, 148-151].
Stresses can build up in the TBC due to transformation of zirconia from metastable tetragonal
form to its monoclinic phase and/or due to its accelerated sintering as the stabilizing phase
present in it reacts with deposits to form oxides. [10, 54, 152-154]. The transformation of
tetragonal zirconia to monoclinic zirconia as well as the enhanced sintering contributes to the
increased stress build-up within the zirconia layer. The phase transformation in zirconia is
accompanied by approximately 4% volumetric expansion and thus severely undermines the
structural integrity of the ceramic coating.”
“An increase in the Young’s modulus of zirconia coating as a result of sintering, adversely
affects its mechanical properties [155]. Molten salt can penetrate into the YSZ coatings along the
cracks thus produced as well as along the pores which may be already present in the coating, and
attack the metallic bondcoat [156, 157]. As the thermal expansion coefficient of the deposits is
different from zirconia, their infiltration through the pores of YSZ creates additional stresses.
This leads to the suggestion that non-infiltrating deposits may be considered as benign. However,
such deposits may still take part in local chemical reaction of yttria from YSZ leading to a phase
transformation from tetragonal to monoclinic phase. Borom et al [126] have proposed a damage
map comprising of four in-service failure regimes for APS YSZ TBCs. They characterized the
regimes as “
(i) infant mortality
(ii) particle erosion
(iii) infiltration of molten particles
(iv) Thermochemical phenomena, such as, sintering, phase changes and bondcoat
oxidation, etc.
“Notably, for APS TBCs, hot corrosion attack of the bondcoat seemed to play a significant
role only after a very long operating period. A model that requires a negative solubility
gradient has been proposed by Rapp and Goto [158] to sustain hot corrosion i.e.”
{𝑑(𝑜𝑥𝑖𝑑𝑒 𝑠𝑜𝑙𝑢𝑏𝑖𝑙𝑖𝑡𝑦)
𝑑𝑥 }
𝑎𝑡 𝑡ℎ𝑒 𝑜𝑥𝑖𝑑𝑒−𝑠𝑎𝑙𝑡 𝑖𝑛𝑡𝑒𝑟𝑓𝑎𝑐𝑒
< 0
31
“where ‘x’ denotes the distance from the oxide–salt interface. This mechanism does not consume
the molten salt and therefore will not be sustained until either the melt becomes more basic and
oxide ions are not produced or the oxide scale is completely removed and the metal becomes
accessible to the salt. The formation of continuous and protective Al2O3 is prevented by self-
sustaining corrosion process.”
“Typically, alkali and alkaline earth sulphates are the salts responsible for hot corrosion. The
exact composition of the salts is determined by various factors including industrial process,
fuel/air/coolant compositions and impurities. According to Khanna and Jha [159], combustion of
sulphur containing coal and fuel oils produce SO2, which in turn partially changes to SO3. NaCl,
either present as impurity in the fuel or already present in the air, reacts with SO3 and water
vapor at the combustion temperature to yield Na2SO4. The product deposits on the surface of
metal/alloy and, with its melting point being 884 °C, may liquefy upon exposure to high
temperature thereby initiating accelerated attack.
An attack is significant only when a sodium sulphate deposit is fused rather than solid as
demonstrated by DeCrescente and Bornstein [116]. An attack was significant only when a
sodium sulphate deposit was fused rather than solid. V2O5 would be formed by small amounts of
vanadium present in the fuel during combustion, which can react with Na2SO4 to form low
melting sodium vanadates that are highly corrosive.”
“Other salts, viz. vanadates or sulphates–vanadate mixtures or in the presence of solid or gaseous
salts, such as chlorides can also cause hot corrosion [160]. Vanadium and sodium are present in
low grade petroleum fuels as impurities. Molten sulphate–vanadate deposits resulting from the
condensation of combustion products of such fuels are extremely corrosive to high-temperature
materials in combustion systems [161]. It was demonstrated by Goebel et al [151] that when the
experimentation was carried out in air, basic fluxing in combination with sulphidation could
cause the hot corrosion of Ni base alloys. It was noted that the failure occurred within the topcoat
and close to the YSZ/TGO interface during the cyclic oxidation of TBC in atmosphere of NaCl
vapors. Due to the above explained mechanisms, the formation of voluminous and non-
protective oxide scales and the increased TGO thickness occurred which may be responsible for
the accelerated failure of TBC [162].”
32
2.3.5 Selection of TiN for modification of bondcoat
In the current thesis, Yttria Stabilized Zirconia TBCs along with CoNiCrAlY bondcoat was
deposited using air plasma spray on Inconel-X750 superalloy. An attempt has been made to
enhance the life of TBCs by applying thin layer of TiN (by physical vapor deposition) on
bondcoat for improving the oxidation and hot corrosion resistance of bondcoat.
The TiN was selected due to its high melting point and high coefficient of thermal expansion
which is comparable to YSZ, see Table 2.4 [6, 15, 96, 98-100].
Table 2.4: Comparison of properties (melting point & thermal expansion coefficient) [6, 15, 96,
98-100]
Materials Thermal Expansion
Coefficient
Melting Point
YSZ 9.2x10-6/K 2780⁰C
TiN 9.35x10-6/K 2930⁰C
Al2O3 8.1x10-6/K 2015⁰C
33
Chapter-3
3 Experimental Procedures
3.1 Process Flow Chart
The processing flow chat followed for the deposition of TBCs systems and consequent heat
treatments, characterization and evaluation of coatings is shown in Fig.3.1.
Substrate samples cutting & edge
rounding
Grit blasting of substrate samples
by Al2O3
Characterization of selected
Powders
Selection of Powders
Bondcoat (MCrAlY), Topcoat (YSZ)
Air Plasma Spray (APS) coatings of
bondcoat on Substrates
Samples with standard range of
bondcoat thickness
Samples with thickness greater
than standard range of bondcoat thi
APS coating of
topcoat
APS coating of
topcoat
TiN coating on
bondcoat
APS coating of
topcoat
Hot corrosion testing Thermal cycling
Stereo Microscopy, Optical Microscopy, SEM, XRD
Fig. 3.1: The process flow chart
34
3.2 Experimental
3.2.1 Material of the Substrate
The material of substrates, used in the experiments, was Inconel-X750 in the form of sheet
having thickness 2 mm. This superalloy is used in applications that require relatively high
oxidation and corrosion resistance and substantially high strength at elevated temperature. It is
commonly used for gas turbines, turbine blades, seals, and rotors. It is also used for aircraft
structures and rocket engines. The chemical composition of the material is given in the table 3.1.
3.2.2 Cutting and Cleaning
Samples of size 25X25X2 mm were cut from a sheet (Inconal-X750) with the help of rapid cutter
machine and the edges were subsequently tapered in order to avoid delamination of coating
during the spraying process. The samples were thoroughly cleaned with acetone in order to
remove any contamination.
Table 3.1: Chemical composition of Inconal-X750 and substrate
Elements Weight percentage
Substrate Inconel-X750*
Ni Bal Balance
Cr 15.41 15.5
Fe 6.99 7.00
Ti 2.46 2,5
Al 0.57 0.7
C 0.04 0.04
* Nearest standard
35
3.2.3 Cylindrical Holder
An aluminum cylindrical holder for holding substrate samples, having 200 mm
diameter and thickness of 8 mm was used to get uniform coatings and good control of
substrate temperature. This cylindrical holder had two parts; the lower part and the
upper ring which were screwed together, Fig.3.2. The specimens were placed between
the lower part and the ring and screws were tightened. It was designed to hold16
uniformly distributed substrate samples. Substrate samples were fixed in the cylindrical
holder which rotated during the spraying operation, making possible uniform coat on
all the samples during coating process.
3.2.4 Grit Blasting
“In order to get rough surface for mechanical bonding between coating and substrate, the samples
were grit blast/sand blast by holding them in the fixture as shown in the Fig. 3.3. The sand
blasting machine furnished with a jig to retain the distance and to regulate the angle between the
Fig. 3.2: Aluminum cylindrical holder for holding substrate samples during grit
blasting and coating
36
nozzle and the samples was used for sand blasting of the samples. The surface of the samples
(substrates) was sand blasted with Al2O3 particles for 5 minutes to achieve roughness (Ra) of 3.0
to 3.5 µm, which is fairly adequate for subsequent coating. Portable surface roughness gauge
(RugoSurf, model 100 S) was used to measure the Ra value. In order to clean the samples, dry
compressed air was used. To avoid surface oxidation and contamination, the samples were
immediately coated with bondcoat after cleaning. Without dismantling the samples the same
fixture was used for plasma coating. The parameters of grit blasting process are given in Table.
3.2.”
Table 3.2: Sand blasting of substrate (Inconal-X750).
Abrasive
Material Air pressure
(psi)
Grit blasting
time (min)
Angle between
substrate &
nozzle
Substrate surface
roughness, Ra
(µm)
Al2O3 81 5 90 3-3.5
Fig. 3.3: Grit blasting set up used for samples cleaning.
Fixture axis
Nozzle axis
Angle of 90
0 between substrates’
surface & nozzle
37
3.2.5 Coating Materials
For topcoat (YSZ) and bondcoat (Co32Ni22Cr8Al0.5Y in wt.%) Metco-204B and AMDRY-
995C powders were used, respectively. Details of the powders are given in the Table 3.3 and
morphology is show in Fig. 3.4.
Powder Chemistry Particle size range Particle
morphology
Metco-204B
(Topcoat)
Zirconia stabilized
with 8% yttria
45 to 75 µm Spherical
AMDRY-955C
(Bondcoat)
Co32-Ni21-Cr8-
Al0.5-Y
45 to 75 µm Spherical
Table 3.3 Chemical compositions (wt. %), particle size range and morphology of the
spraying powders
Fig. 3.4: Morphology of coating powders used for plasma coating (a) YSZ ceramic powder (b)
Co32Ni22Cr8Al0.5Y metallic powder.
38
3.2.6 Deposition Process
For the deposition of both the metallic bondcoat and the ceramic topcoat, air plasma spraying
(APS) technique was used. The Fig. 3.5 shows the experimental set up of plasma coating.
The freshly blasted surfaces with fixture were immediately held in a chuck which was later
rotated at an optimized speed of 120 rpm. For plasma spraying 9MB Sulzer Metco gun was
selected to deposit both topcoat and bondcoat. The spraying gun was adjusted at 90º to the
substrates. In order to get uniform thickness, the gun was moved to and fro, relative to the
rotating samples, Fig. 3.6. Substrate’s temperature was maintained at ~170oC using a constant
flow of compressed air during the coating process. In this regard an IR-camera (IRtech-
P1000+Mk2) was also installed at a distance of 4 meter to monitor the temperature of the
coatings. The distance between the substrate holder and the compressed air nozzle was
maintained at 110 mm, to ensure the reproducibility of the coatings. All the important spraying
parameters both for topcoat and bondcoat are mentioned in the Table 3.4. All the samples were
Fig. 3.5: Schematic diagram showing the setup of APS
coatings.
39
preheated by utilizing the same plasma gun as was used for coating, so that no moisture was left
on the surface before depositing the spraying powder. For topcoat and bondcoat Metco-204B and
AMDRY-995C powders were used, respectively. Argon (Ar) was used as primary gas, for the
generation of plasma whereas hydrogen (H2) was used as secondary gas in plasma flame to
increase its enthalpy. The carrier gas Argon was used for the inoculation of coating powders.
After the deposition of bondcoats with different thicknesses, one set of samples was removed
from the fixture and a thin layer of TiN was deposited by physical vapor deposition method,
Fig.3.7. Finally topcoat was applied by air plasma spraying on all samples, Fig.3.8.
Parameters Spraying Powder
CoNiCrAlY YSZ
Current (A) 600 600
Voltage (V) 66 66
Primary gas, Ar (SLPM) 55 32
Secondary gas, H2 (SLPM) 8 10
Powder feed rate (g/min) 150 150
Spray distance (mm) 110 110
Table 3.4: Air plasma spraying parameters for TBCs system.
Fig. 3.6: A schematic diagram showing experimental setup for APS coating
used in this study.
Pla
sm
a g
un
mo
tio
n
40
Fig. 3.7: A thin layer of TiN (golden color) was deposited by physical vapor deposition
method after the deposition of bondcoat for TiN modified samples
Fig. 3.8: TBC coating systems after the deposition of topcoat
41
3.2.7 Cathodic Arc Physical Vapor Deposition of Titanium Nitride
“Titanium Nitride film was deposited on the surface of samples by using DREVA RC 400
coating plant. The prepared substrates with bondcoat deposition were fixed on a rotating sample
holder at an angle of 30 with the normal to the cathode. The cathodic arc physical vapor
deposition coating machine is shown in Fig. 3.9, equipped with a disc shaped cathode (titanium
target) attached on water cooled copper stage. A disc shape anode was mounted at a
perpendicular distance of 300 mm on top of the cathode. The whole assembly was enclosed in a
double walled stainless steel jacket. After evacuation, the chamber was cleaned with the help of
hollow cathode argon plasma discharge. The arc was triggered by using grounded copper wire.
For the ingress of nitrogen into the chamber with controlled partial pressure, an automatic
microprocessor controlled feeding system was used. The parameters used for deposition of
coating were optimized in order to deposit a thin coating with low density of permeable defects.”
Fig. 3.9: Cathodic Arc Physical Vapor Deposition Coating plant (DREVA RC 400) used
for deposition of Titanium Nitride films.
42
Fig.3.10: Standard TBC samples, placed in a stainless steel tray, with salt mixture on
the top surface of the samples for hot corrosion test at 950˚C
Salt mixture
TBC coated
sample
3.2.8 Hot Corrosion
For the hot corrosion testing, V2O5 and Na2SO4 were mixed in a ball mill in 1:1 ratio (by
weight). The mixture was spread over the as sprayed samples with a concentration of 30 mg/cm2,
leaving 3 mm surface/space free of salt from the edges to avoid edge effect as per procedure
mentioned by Chen et al [163]. The melting point of Na2SO4 and V2O5 is 884⁰C and 690⁰C,
respectively [53]. The samples were placed in a stainless steel tray before loading into the
furnace (Fig.3.10). The samples were heated up to 950ºC, at the rate of 20ºC/min. The hot
corrosion tests were run in cycles of ten hours durations. After each heating cycle the furnace
was shut down to let the samples cool to room temperature. They were then visually inspected
and each set of samples was characterized using XRD, optical metallography and SEM, while
the remaining samples were re-exposed to same environment, without replenishing the salts. A
total of 5 such cycles (50 hours) were given to samples before concluding the results.
43
3.2.9 Thermal Cycling Treatment
The samples were placed in an electric box type furnace at 950°C for about 5 min. After this they
were quenched in water (to ambient temperature). This thermal cycling process for life
estimation and evaluation was repeated until 30% spallation of the coating topcoat was observed.
3.2.10 Adhesion Test
Bonding strength between TBC coatings layer and substrate was evaluated as per standard
ASTM C 633. The schematic of bonding strength measurement set up is shown in the Figure
3.11.
Fig. 3.11: A schematic setup used for evaluation of bonding strength: (a) showing the
parts used for testing; (b) the assembly used for tensile adhesion testing.
Epoxy wafer
(a)
(a)
(b)
(a)
44
The elastomeric adhesive FM 1000 (1 inch in diameter and 0.02 inch thick) was applied on
coated coupons (5 mm thick) at the bottom and top to the pulling rods. At 180 °C for 3 hours, the
adhesive was cured. The tests were performed on a tensile testing machine at a cross-head speed
of 2 mm/min. and average of three measurements is reported. Prior to testing of the coatings,
bonding strength of the cured adhesive (FM 1000) with stainless steel pulling rods was
measured.
3.2.11 Delamination of Topcoat after Hot Corrosion
In order to confirm the formation of different phases in samples after exposure to hot corrosion,
the topcoat was delaminated by chemical etching process. For this purpose 50% diluted HCl was
used, whereas, the process was done at room temperature. The chemical attacked the interface of
both bondcoat-topcoat and substrate-bondcoat. Due to relatively porous nature of the bondcoat,
as compared to the substrate, the chemical dissolved the bondcoat preferentially. As a result, a
delaminated topcoat was obtained with attached phases that were formed during the hot
corrosion. The delaminated topcoat was washed with water and was preserved for further study.
3.3 Characterization
3.3.1 Stereo Microscopy
Top surface of as coated samples were observer under a stereo microscope. During hot corrosion
testing, the samples were taken out after every 10 hours exposure and top surface was observed
to detect spalling and initiation of micro cracking. The samples after the intense thermal cycle
testing were also observed after fixed interval of cycling, to detect spalling and beginning of
cracking.
3.3.2 Sample Preparation (Metallography)
“Slow Speed Diamond Cutter was used for the cutting of samples to reduce the damage by wear
and heat (produced due to frictional). After cutting, samples were mounted in mounting machine.
For grinding, wet grinding technique was used to grind the samples and diamond paste was used
for polishing of sample. The polished samples were cleaned with the help of ultrasonic
machine.”
45
3.3.3 Optical Microscopy
In order to study microstructures of air plasma sprayed coatings, metallographic examination
was carried. Olympus BX51 microscope (Fig. 3.12) was used to observe the cross section of
mounted and polished samples. An image analyzer which was attached to the optical microscope
was utilized to calculate the percentage of overall porosity and defects present in the coating. The
interfaces of substrate and coating, bondcoat and topcoat were also observed. Moreover, the
thicknesses of the bondcoat, topcoat and the overall coating were also determined by the image
analyzer.”
3.3.4 Scanning Electron Microscopy (SEM)
“After optical microscopy, the structure of air plasma sprayed coatings was observed in Jeol
Scanning Electron Microscope (Fig. 3.13). Both the top surface and the cross-section of the
samples after hot corrosion, thermal cycling and as coated samples were examined. In order to
analyze the ceramic coatings in SEM, gold sputtering was done on the polished surfaces of the
samples to make them conductive. Chemical composition profiles in treated and as-sprayed
samples were determined by the Energy Dispersive Spectroscopy (EDS). EDS analysis was
Fig. 3.12: BX51 Olympus optical microscope with digital camera and image analyzer
46
taken at different points on the interfaces of the bondcoat and topcoat. EDS analysis, in atomic
percentage was carried out in order to have an idea of the stoichiometry of the compound present
at interfaces.”
3.3.5 X-Ray Diffraction (XRD) Study
“JEOL JDX-8030 machine was used for X-ray diffraction analysis, Fig.3.14. For the scanning of
spraying powder and each coating system, Cu-K radiations with Ni filter were used. The scan
step used in the study was 0.05º whereas the scanning range was from 20º to 100º. The
‘integrated intensity ratio’ method was used to calculate the volume of each relevant phase by
using XRD pattern.”
Fig. 3.13: JEOL JSM 5910 LV - Scanning Electron Microscope (SEM) with EDS
47
X-ray diffraction analyses were performed on the ‘as coated’ sample as well as on the exposed
samples to study different changes in the phases. XRD of the delaminated topcoat portion that
faced the bondcoat was done to study the phases that were produced by diffusion and oxidation
processes occurring through the interface, during the hot corrosion at high.
Moreover, the values of lattice parameters were measured and were used to calculate the residual
stresses.
Fig. 3.14: X-Ray Diffractometer used for phase analysis of coatings
48
Chapter-4
4 Results and Discussion-I
4.1 Hot corrosion of yttria-stabilized zirconia coating, in a mixture of sodium
sulfate and vanadium oxide at 950oC.
Chemical compositions (wt.%), particles size range and morphology of the spraying powders
utilized to deposit standard TBC (MCrAlY-bondcoat + YSZ-topcoat) are given in Table 3.3. All
the important spraying parameters both for topcoat and bondcoat are mentioned in the Table 3.4.
The hot corrosion test (detail of test and heat treatment cycle is given in experimental 3.2.8) of
the samples was performed to determine the effects on their microstructures, phases and
delamination behavior. Furthermore, the chemical composition profile of all samples was done
to analyze the diffusion behavior of different constituent elements of bondcoat and the substrate.
Results and Discussion
4.1.1 Microscopy
4.1.1.1 Surface morphology
The top surface of YSZ coating of the as-sprayed sample demonstrated rough surface with few
semi-molten particles (SMP) as shown in Fig. 4.1. Additionally, cracks and porosity like features
were also observed on the surface. These features are typically observed in similar ceramic
coatings [164].
The surface of the topcoat after hot corrosion testing showed numerous rod-like and
agglomerated crystals (Fig.4.2). EDS analysis of these agglomerates and rod-like structure
revealed ZrO2 and YVO4, respectively. It was found that degradation of the topcoat started by the
formation of YVO4 crystals as some of the crystal rods were found on the surface while others
were firmly adhered to the topcoat; evident from Fig.4.2. During exposure the salt mixture (50%
Na2SO4 and 50% V2O5) demonstrated a chemical reaction and formed a eutectic compound
NaVO3, as per following equation:
49
Na2SO4 + V2O5 →2 NaVO3 + SO2 +1/2 O2 (Eq. 1)
The NaVO3 compound acted as a corrosion catalyst and served as an oxygen carrier. The
compound which was believed to enter into the pores present in the plasma-sprayed coating,
reacted vigorously with Y2O3 forming YVO4 as per following reaction [165]:
ZrO2 (Y2O3) +2NaVO3 → ZrO2 + 2YVO4 + Na2O (Eq. 2)
It was noticed that formation of YVO4 increased with an increase in exposure time, as evident
from Fig.4.3. In first 10 hours the concentration of the YVO4 needles/rods was very low whereas
the size and density of these rods increased rapidly with increase in exposure temperature.
Additionally, no sodium peaks were observed during EDS analyses, as expected from Equation
(2); probably by removal of Na2O, due to sublimation process at high temperature [166].
Fig.4.1: The as-sprayed topcoat surface exhibiting a rough surface with few semi-
molten particles (SMP), cracks and porosity like features
Crack
Porosity
SMP
50
Fig.4.3: Topcoat surfaces after, exposing to 950°C in a hot corrosion environment for various
time intervals, showing increased concentration of YVO4 rods with higher exposure times.
Fig.4.2: Topcoat surface after 50 hours exposure to hot corrosion environment, containing YVO4
rods and agglomerated crystals of ZrO2 (b) high magnification, rod like and agglomerated crystals
51
4.1.1.2 Cross sectional observations
Cross section of an as-sprayed sample (Fig.4.4) showing general features of topcoat and
bondcoat. In general, topcoat exhibited porosity with lamellar structure and in bondcoat some
grayish flake-like features were observed with porosity. Figure 4.5 reveals morphology of the
exposed samples showing presence and growth of thermally grown oxide (TGO) with a
prolonged exposure time. The oxide penetration in the bondcoat was relatively higher for longer
holding times. It was observed that a crack started to appear in samples after 30 hours exposure
(Fig.4.5) which became more pronounced in the samples treated for 40 and 50 hours, Fig.4.5.
In order to understand the cracking phenomena at the interface between the bondcoat and
topcoat, the cross section of the longest exposed (50 hours) sample was further characterized.
Various distinct features have been marked on Fig.4.6 and 4.7, whereas, the corresponding
chemical compositions and the possible compounds are given in Table-4.1 & 4.2, respectively.
Formation of Al2O3 was noticed away from the bondcoat-topcoat interface (Fig.4.6, site-2),
demonstrating that initially alumina was formed. Whereas, perovskite-like structure AlCrO3
(Fig.4.6, site-1) was observed closed to the interface. It was believed that Al, Cr after diffusing
Fig.4.4: (a) Cross-section of as sprayed sample (b) high magnification of bondcoat showing
grayish flake-like features and porosity
52
into topcoat initially converted into Al2O3 and Cr2O3, followed by formation of perovskite of Al
and Cr [165]. The possible reactions are given below:
2Al+ 3/2O2 → Al2O3 (Eq. 3)
2Cr+3/2O2 → Cr2O3 (Eq. 4)
Al2O3+Cr2O3→2AlCrO3 (Eq. 5)
The interface between the diffused layers (Fig.4.7, light grey) surrounded by topcoat and the
bondcoat showed presence of retained Al2O3 at discrete locations which did not crack or dissolve
Fig.4.5: Cross section of samples exposed to hot corrosion environment at 950°C, showing
effect of exposure time on thermally grown oxide (TGO).
53
during the chemical reactions. Figure 4.7 also demonstrates knife-like features, penetrated into
the bondcoat forming perovskite of CoNiO3.
Toshio et al [167] reported a reaction between Cr2O and NaVO3 which was believed to dissolve
Cr2O3 forming Na2CrO4 along with V2O5 but no evidence of Na2CrO4 formation was observed.
EDS analysis (atomic %) revealed formation of both, spinels and perovskite structures during the
hot corrosion at bondcoat-topcoat interface and within the bondcoat; these include AlCrO3,
NiCr2O4, NiCrO3, NiCrO4, CoNiO3 CoCr2O4. The formation of these structures generated cracks
at the interface as a result of volumetric changes [168]. Furthermore, the cross sectional samples
%age, atomic Possible Phases
Location Al Cr Co Ni
1 37 37 10 6 AlCrO3
2 29 17 22 27 Al2O3
Table-4.1: Chemical composition of various phases marked in
Fig.4.6 (exposed for 50 hours at 950⁰C)
Fig.4.6: SEM micrograph showing cross-section of bondcoat in a sample exposed to hot
corrosion testing for 50 hours: Site-1 and 2 are locations of EDS analyses performed; data
is given in Table-4.1
54
revealed that during hot corrosion at high temperature, the surface layer constantly eroded due to
the formation of various phases. It was noticed that deterioration of the top surface increased
with increasing time (Fig.4.8). The thickness of the topcoat decreased from ~230 µm to ~210
µm. The sequence of hot corrosion process is schematically presented in Fig.4.9.
%age, atomic Possible Phase
Location Al Cr Co Ni
1 9 20 35 36 CoNiO3
2 43 35 11 10 AlCrO3
3 38 46 9 7
4 49 14 19 17 Al2O3
5 6 23 37 34 CoNiO3
Table-4.2: Chemical compositions of various phases marked in
Fig.4.7 (exposed for 50 hours at 950⁰C)
Fig.4.7: SEM micrograph showing discrete sites in the bondcoat in a sample exposed for
50 hours showing knife-like features (Site-1 and 5). Site-1 to 5 represent locations for EDS
analyses reported in Table No.4.2.
55
Fig.4.9: Schematic diagram (a-d) showing various oxidation reactions during hot
corrosion testing.
O2 O2
Cr+2
Cr+2
O2
O2 O2
Al2O3 Al & Cr depleted region
Cr2O3
Bondcoat
Al2O3
Co+1
Ni+1
Bondcoat
O2
O2 O2
AlCrO3
O2
O2 O2
Al2O3
Cr2O3
Bondcoat
CoO NiO
AlCrO3 NiCr2O4
Topcoat Topcoat
Topcoat Topcoat
NiCrO3 and NiCrO4
CoNiO3, CoCr2O4 Al, Ni, Co & Cr depleted region
Bondcoat
(a) (b)
(d) (c)
Fig.4.8: Cross-section of the coating showing increase in deterioration (arrows) of
the top surface with time.
56
4.1.2 Chemical Composition Profile
Significant diffusion phenomena were observed at the bondcoat-substrate interface. Chemical
composition profile revealed that the alloying elements diffused both from substrate to the
bondcoat and from bondcoat to the substrate. The alloying elements diffused from higher
concentration to lower concentration sites. Thus iron, nickel and titanium diffused towards
bondcoat, whereas, cobalt and aluminum diffused towards substrate.
Iron was present up to ~7% in the substrate and its concentration in bondcoat after 10 hours was
observed to be substantially decreased, as shown in Fig.4.10. It was observed that iron reached
much deeper in bondcoat with increased exposure time, however, no evidence was available to
show that it approached to topcoat. It can also be seen from the slope of the curve in Fig.4.10-
4.13, that iron concentration curve decreased with an increase in exposure time i.e. the
percentage of iron close to substrate-bondcoat interface increased with time. The continuous
increase in concentration with time indicated that iron was not exposed to oxygen and
presumably did not form an oxide when reached to the bondcoat. The available oxygen, at that
time, close to the substrate within the bondcoat preferentially was utilized to oxidize aluminum
and titanium (diffused from the substrate).
Titanium, which was present in the substrate (~2.5 wt. %) also diffused to the bondcoat.
Although the atomic size of Ti atom is comparatively larger (i.e 2 Aº), however, it diffused at
high temperature. This was due to the availability of vacancies created by the diffusion of cobalt
to the substrate [169]. It was observed that the concentration of titanium increased in first 30
hours and then no titanium activity was noticed within the substrate at higher exposure time. This
may be due to either the presence of very small amount of titanium concentration left within the
substrate or due to the oxidation of titanium, which diffused into the bondcoat, forming titanium
oxide. This oxidation prohibited further transport of titanium atoms at higher temperatures.
Additionally, titanium-sulphur rich precipitates (titanium sulphide) were observed close to the
interface of substrate-bondcoat, within the substrate, as shown in Fig.4.14. It was noticed that
formation of these precipitates increased with an increase in exposure time (Fig.4.14 (a-d)). The
availability of sulphur was probably occurred due to the transportation of Na2SO4, present in hot
corrosion salts.
57
Nickel was another alloying element which diffused out from the substrate to bondcoat. The
concentration of nickel was more than 70% in substrate whereas, 32% nickel was present in the
bondcoat. The diffusivity of nickel in the first 10 hours can be seen in Fig.4.10. It can be seen
that some peaks are present in the chemical composition profile at lower exposure times. The
peaks may refer to the fact that nickel oxides were formed at some instance but this oxide
formation did not hinder the nickel diffusion for the next 30 hours (Fig.10). It can also be noted
in Fig.4.12 that up to 20% nickel depleted in region close to the bondcoat in the substrate after 40
hours exposure at high temperature. Similarly, the regions close to the bondcoat having high
nickel concentration zone.
In the bondcoat cobalt and aluminum were in high concentration as compared to the substrate
(i.e. 38% Co and 8% Al in bondcoat); thus both elements were expected to diffuse towards the
substrate. Chemical composition profiles for aluminum at low exposure time (i.e. 10 and 30
hours) indicate peaks of aluminum, indicating formation of aluminum oxide. At longer exposure,
however, these peaks disappeared (Fig.4.12-4.13) while the overall weight percentage of
aluminum in bondcoat decreased, no matter the region was either closer to the substrate or near
the topcoat. In case of cobalt, the chemical composition profile exhibited (Fig.4.10-4.13) that it
was depleted in the regions closer to the bondcoat and contrarily enriched in regions of bondcoat,
near the topcoat (Fig.4.10-4.13).
Interface between the bondcoat and topcoat also demonstrated a diffusion phenomenon at high
temperature (950⁰C). It was revealed that chromium, cobalt, nickel and aluminum diffused
towards the topcoat. The diffusion of these elements became possible upon the cracking or
dissolution of the alumina layer called TGO (thermally grown oxide) [170]. The oxidation of
these elements then led to the formation of spinels and perovskite structure. It was observed that
the diffusion of above elements continuously increased in the topcoat and thus the thickness of
spinels and perovskite structure constantly increased with an increase in exposure time.
58
Fig.4.10: Elemental distribution after 10 hours exposure in hot corrosion atmosphere at 950ºC
5 10 15 20 25 30 35
40 45 50
(µm)
5 10 15 20 25 30 35
40 45 50
(µm)
59
Fig.4.11: Profile of elements after 30 hours exposure in hot corrosion atmosphere at 950ºC
5 10 15 20 25 30 35
40 45 50
(µm)
5 10 15 20 25 30 35
40 45 50
(µm)
60
Fig.4.12: Profile of elements after 40 hours exposure in hot corrosion atmosphere at 950ºC
5 10 15 20 25 30 35
40 45 50
(µm)
5 10 15 20 25 30 35
40 45 50
(µm)
61
Fig.4.13: Profile of elements after 50 hours exposure in hot corrosion atmosphere at 950ºC
5 10 15 20 25 30 35
40 45 50
(µm)
5 10 15 20 25 30 35
40 45 50
(µm)
62
4.1.3 X-Ray Diffraction Analysis
4.1.3.1 Phase analysis
X-ray diffraction (XRD) analyses were performed on the exposed samples as well on as-sprayed
samples. It was noticed that three main phases formed on the surface during the long exposure
(Fig.4.15), in hot corrosion environment. In as-sprayed surfaces no phase other than tetragonal
ZrO2 (t-ZrO2) was present [171]. The exposure to hot corrosion environment caused
transformation of t-ZrO2 to monoclinic-ZrO2 (m-ZrO2) and yttrium vanadate (YVO4).
Fig.4.14: Titanium-rich precipitates (arrows) close to the substrate-bondcoat interface.
63
It was noticed that concentration of YVO4 was the highest after 10 hours which was maintained
on further exposure. , It can be observed in Fig-4.16 that the concentration of vanadate
continuously increased till 30 hours and then decreased up to 40 hours. The decrease in
percentage was believed to be due to spalling of topcoat (Fig.4.8), associated with high stress
imposed by the formation of YVO4 on the surrounding coating. After 50 hours a fresh vanadate
formation was noted.
Comparison of t-ZrO2 phase and m-ZrO2 phase concentrations (Fig.4.17) reveals that formation
of m-ZrO2 was directly related to the t-ZrO2 i.e. t-ZrO2 rendered with time at high temperature
and transformed into m-ZrO2 phase. It was also observed that the formation of m-ZrO2
consistently increased with increasing exposure time, however, a decrease in m-ZrO2 phase was
observed after 40 hours exposure; this was associated with the spalling of topcoat, due to
excessive stresses exerted by the volumetric changes associated with m-ZrO2 phase.
Fig.4.15: XRD graph showing patterns of as-sprayed coating and samples exposed to hot
corrosion for various durations.( PCPDF#●82196 ◊ 830939 ○811548 )
2Ɵ →
I/I0
● ●YVO4(tetragonal)
◊ZrO2(monoclinic)
○ ZrO2(tetragonal)
●
○
● ●
●
◊
○
● ◊
◊ ●
◊
○
64
4.1.3.2 Lattice distortion
Lattice distortion was directly related to the development of stresses. As the distortion of the
system increases, the stress in the system would also increase. Therefore, the shift of high angle
peaks, in XRD scan, for YVO4 and m-ZrO2 were taken into account while calculating lattice
parameters for the respective systems, Fig.4.18. In this regard (200) and (002) crystallographic
planes were considered for YVO4 and m-ZrO2, respectively. In case of YVO4, (200) plane is the
strongest but (002) is not the strongest for m-ZrO2; however, the reason for selection of (002) for
Fig.4.16 Effect of exposure time on the percentage of YVO4
Fig.4.17 Effect of exposure time on the percentage of m-ZrO2 and t-ZrO2
65
m-ZrO2 was to simplify the calculations. YVO4 has a tetragonal structure with a=b=7.118 nm,
c=6.289 nm and α=β=γ=90⁰, in stress free condition [172]. The following equation was utilized
to calculate the value of “a” [173]:
1
𝑑2 = 1
𝑎2(ℎ2 + 𝑘2) +
1
𝑐2 𝑙2 (Eq. 6)
Figure 4.19, represents change in the value of lattice parameter “a” with exposure time. It was
observed that the distortion in the lattice increased with increasing the exposure time. The
distortion in lattice represented the stress of the system which consistently increased with time,
as shown in Fig.4.19. It was also noticed that the curve progressed with exposure time, was of
the order of 3rd
polynomial i.e. in the first 10 hours the distortion of the lattice was higher
compared with the later stages; this was believed to be due to the integrity of the coating present
around the YVO4 rods. Thus, the crystals of yttrium vanadate suffered from a constraint around
them which was relaxed at longer exposure time due to partial spalling of the coating in the
vicinity. Similarly, after 40 hours the distortion in the vanadate crystals increased relatively
Fig.4.18: XRD patterns showing shift of (200) plane with exposure time at high
temperature.
● ●YVO4(tetragonal)
◊ZrO2(monoclinic)
○ ZrO2(tetragonal)
◊ ●
○
66
rapidly which could be related to the boundary conditions provided by the newly transformed
monoclinic phase.
Monoclinic zirconia (m-ZrO2) has the following lattice parameters: a=5.149 nm, b=5.207 nm,
c=5.316 nm and α=90⁰=γ, β=99.225⁰ in stress free condition [167]. The hot corrosion cycles
forced to transform the topcoat into other constituents and thus lattice parameters were affected.
Following equation was used to calculate the value of “c” [173]:
1
𝑑2 =1
𝑎2
ℎ2
𝑆𝑖𝑛2𝛽+
1
𝑏2 𝑘2 +1
𝑐2
𝑙2
𝑆𝑖𝑛2𝛽−
2ℎ𝑙𝐶𝑜𝑠𝛽
𝑎𝑐𝑆𝑖𝑛2𝛽 (Eq. 7)
It is evident from Fig.4.19 that the lattice parameter “c” for m-ZrO2 gradually increased for first
30 hours. This increase in lattice parameter was attributed to the constant formation of m-ZrO2.
The m-ZrO2 evolved from t-ZrO2 on loosening the Y2O3, a stabilizing agent present in solid
solution. On plotting a trend line, it was revealed that the lattice parameter “c” of monoclinic
phase followed a 2nd
order polynomial i.e. in the beginning it increased very rapidly and then
decreased after 30 hours exposure time. A decrease in “c” was presumably associated with
delamination of the coating due to the development of high stresses within the zirconia coating.
Fig.4.19: Effect of exposure time on the lattice parameters of m-ZrO2 and YVO4
(te
trag
on
al)
a, n
m
(mo
no
clin
ic)
c, n
m
M-ZrO2
YVO4
67
4.1.4 Delaminated topcoat sample
The portion of topcoat facing the bondcoat was considered important because of diffusion and
oxidation processes occurred through this interface during the hot corrosion at high temperature.
It was also important to investigate the phases formed during hot corrosion.
Furthermore, it was necessary to confirm the EDS analysis (% atomic) results showing the
formation of spinel and perovskite structures.
In this regard, the portion next to bondcoat of the delaminated topcoat, obtained after chemical
etching, was exposed to X-rays. The results demonstrated that after 50 hours at high temperature
in hot corrosion environment, multiple phases were formed at the bondcoat-topcoat interface. In
these phases, alumina, oxides (NiCrO4, CoNiO2, CoNiO4) and the formation of Spinel NiCo2O4,
CoCr2O4 and perovskite NiCrO3 was confirmed, Fig.4.20.
2Ɵ →
I/I0
Fig.4.20: XRD pattern showing multiple phases formed in the delaminated topcoat of the sample exposed at 950ºC for 50 hours. (JCPDF# •760144 ●210596 ♦220748 ○750198
♥801668 ♠100188 ♣731704)
68
4.2 Conclusions
Yttria stabilized zirconia along with CoNiCrAlY bondcoat were deposited by air plasma spaying
on Inconel-X750 coupons. The coatings were exposed to hot corrosion using corrosive salts of
Na2SO4 and V2O5. The results demonstrated that the Y2O3 present in solid solution of ZrO2
reacted with the salt mixture and formed rods of yttrium vanadate (YVO4). Reaction of oxygen
with various metallic elements in the bondcoat resulted in formation of spinels consisting of
NiCr2O4 and CoCr2O4 along with perovskite structure of AlCrO3, NiCrO3 and CoNiO3 and
oxides (NiCrO4, CoNiO2, CoNiO4) at the interface of bondcoat and the topcoat. Development of
these structures forced to crack the interface.
Chemical composition profile revealed that various alloying elements diffused from bondcoat to
the substrate and from substrate to the bondcoat which might had altered mechanical properties
of the interface. Similarly diffusion also took place at interface of bondcoat and the topcoat. The
diffusion of elements (chromium, cobalt, nickel and aluminum) from bondcoat to topcoat
became possible upon the cracking or dissolution of TGO. The oxidation of these elements then
leads to the formation of spinels and perovskite structure.
XRD analyses determined that m-ZrO2 formed along with YVO4 with increasing exposure time
at high temperature. Moreover, a shift in high angle peaks indicated high level of stresses present
in the coating due to the formation of YVO4 and m-ZrO2.
69
Chapter-5
5 Results and Discussion-II
5.1 Evaluation of titanium nitride modified bondcoat system used in thermal
barrier coating in corrosive salts environment at high temperature
“Thermal barrier coating (TBC) systems were produced by air plasma spraying on nickel base
superalloy. These coatings were composed of Y2O3 stabilized ZrO2 topcoat and CoNiCrAlY
bondcoat and were given the name as Standard TBC. Chemical compositions (wt.%), particles
size range and morphology of the spraying powders utilized to deposit Standard TBC (MCrAlY-
bondcoat + YSZ-topcoat) coatings system are given in Table 3.3. All the important spraying
parameters both for topcoat and bondcoat are mentioned in the Table 3.4. In this Chapter,
standard TBC samples were compared with TiN modified bondcoat TBC samples. Titanium
nitride was deposited by utilizing physical vapor deposition technique on the bondcoat (see,
3.2.7). Both TBC systems were exposed to high temperature under the presence of corrosive
salts i.e. a mixture of V2O5 and Na2SO4 (50:50) for 50 hours (Fig.5.1). It was observed that TiN
modified samples showed better results in terms of oxidation resistance and delamination. The
formation of Cr2Tin-2O2n-1 phases at the interface of topcoat-bondcoat, in TiN modified samples,
were found to enhance the thermal and oxidation properties of the TBC.”
Fig.5.1 Standard TBC samples (row-1) and TiN modified sample (row-2), placed in a
stainless steel plate, with salt mixture on the top surface of the samples
Mixture of salts
70
5.1.1 Results and Discussion
5.1.1.1 Surface of topcoat after hot corrosion
A little spalling was observed at the edges of the topcoat in both types of samples after 10 hours
exposure. These edges were probably spalled under the thermal stresses due to initial direct
exposure of samples at high temperature. After 30 hours of exposure, it was revealed that the
standard samples spalled to a greater extent than the TiN-modified samples, Fig.5.2. After the
exposure of 50 hours, the standard samples showed about 10-12% spalling, whereas the TiN
modified samples spalled far lesser at the edges, Fig.5.2.
High magnification images of the topcoat of the two systems revealed the rods of YVO4 (yttrium
vanadate). In both cases, the rods were randomly dispersed on the surface of the topcoat, Fig.5.3
and 5.4. It appeared that the salt mixture reacted and formed a eutectic compound NaVO3:
Na2SO4 + V2O5 →2 NaVO3+ SO2 +1/2 O2 (Eq. 1)
The NaVO3 compound acted as an oxygen carrier and entered into the pores of the plasma-
sprayed topcoat. It reacted with Y2O3 (present in the solid solution of Y2O3 stabilized ZrO2)
forming YVO4 as per following reaction [165].
ZrO2 (Y2O3) +2NaVO3 → ZrO2 + 2YVO4 + Na2O (Eq. 2)
The loose powder, which was found as debris on the topcoat surface, was analyzed as ZrO2 that
was left after the formation of YVO4 rods.
Microstructural Analysis
5.1.1.2 Cross-section of as-Sprayed Coatings
The cross-section of the two systems demonstrated typical air plasma sprayed coatings features
i.e. micro-cracks, lamella of semi-molten particles and shrinkage cavities. It was estimated that
about 8-12% pores were present in the topcoat. Typical lamellar structure was observed
predominantly in the bondcoat after the air plasma spraying process, Fig.5.5.
71
Fig.5.2: S2 (Standard TBC) and TiN modified samples showing condition of top surfaces after
different time intervals, treated at 950°C in a hot corrosion environment
72
It was noticed that the thickness of TiN deposited on bondcoat varied from 6 to 10 µm, Fig.5.6.
It was found that the sputtered layer was not deposited properly at some locations, Fig.5.7.
Furthermore, few vertical cracks were also observed within the TiN thin layer, Fig.5.6. This
cracking might be due to stress relaxation of the coating during pre-heating before the deposition
of topcoat.
Fig.5.3 Top surface of standard TBC (a) low magnification (b) high magnification showing
rod like features (YVO4) after 50 hours exposure
Fig.5.4 Top surface of TiN-modified sample (a) low magnification (b) high magnification showing
rod like features (YVO4) after 50 hours exposure.
73
Fig.5.5 (a) SEM micrograph showing typical structure of as sprayed TBC coating.
(b) high magnification of bondcoat showing lamellar structure
Fig.5.6 Optical micrograph showing layer of TiN (arrows) and interface. Vertical cracks
(box) are also present at some locations.
74
5.1.1.3 Cross-section after Hot Corrosion
TiN-modified samples, after 50 hours of hot corrosion testing, demonstrated that overall
oxidation condition of bondcoat were relatively less severe as compared to the standard samples,
Fig.5.8. This confirmed that TiN acted as a good barrier against the oxygen, at high temperature.
It was observed that TiN modified system had formed denser and uniformly thick oxides layer at
the interface of topcoat-bondcoat, Fig.5.9. However, in the case of standard system, the oxide
layer was irregular and scattered within the bondcoat and topcoat, Fig.5.10. The inherent defects
of plasma sprayed coatings such as porosity and splat boundaries acted as a diffusion channels
for corrosive liquids. EDS analyses of bondcoat showed that no “vanadium” was present, in TiN-
modified samples, whereas, bondcoat of standard samples revealed “vanadium” in analyses,
Fig.5.11. This concluded that vanadium oxide crossed the diffusion barrier of alumina (TGO) in
case of standard samples, whereas, TiN layer offered resistance against its penetration.
Fig.5.7 Optical micrograph showing layer of TiN which was not deposited properly at some
locations, (arrows)
75
Fig.5.8 Cross-section of both standard TBC (a) and TiN modified (b) samples, after 50 hours
exposure in hot corrosion environment
Fig.5.9 Cross-section of TiN modified sample after 50 hours exposure in hot corrosion environment
demonstrating dense and uniform oxide layer at interface of bondcoat-topcoat
76
Fig.5.10 Non-uniform oxide layer in Standard TBC sample after 50 hours exposure (a-low
mag. and b-high mag.)
Fig.5.11 EDS analysis at the boundaries of the splats in bondcoat showed no “vanadium” is present in
TiN-modified samples (a), standard TBC samples demonstrated the presence of vanadium (b) near
topcoat-bondcoat interface
77
At high magnification, the TiN-modified samples revealed different features at the interface of
bondcoat-topcoat. It was observed that at high temperature, TiN destabilized and formed other
compounds in the presence of abundant oxygen. TiN is not stable above 600⁰C [174]. It seemed
that above this temperature, TiN oxidized to Magneli phase TinO2n-1 (4 ≤ n ≤ 9) [175] which had
a complicated defect structure [176]. The point defects present in these system are dominated by
oxygen vacancies and titanium interstitials [175, 177]. The so-called Magneli phases have long
order defect structures [175]. These phases further react with chromium oxides and may form a
series of homologous structures i.e. Cr2Tin-2O2n-1 (6 ≤ n ≤ 9). Cr2Tin-2O2n-1 phase is known from
its stability against thermal stresses and oxidation [178].
EDS analysis, in atomic percentage, may give idea of atomic ratios in the compound. Thus the
most abundantly formed phase, present at the boundaries of topcoat-bondcoat, revealed that the
atomic ratio of Ti:Cr is 3. Figure 5.12 shows three different points from where EDS analysis was
taken. The %atomic composition obtained from these points is shown in the Table 5.1. This
indicated that Cr2Ti5O13 phase is predominately formed at the interface of topcoat-bondcoat, after
50 hours of exposure. Cr2Ti5O13 is known as stable up to 1485⁰C [179].
The formation of Cr2Ti5O13 at the interface of topcoat-bondcoat can be explained by two
mechanisms. Firstly, at high temperature, the destabilized TiN transformed to titanium oxide that
further reacted with the underlying chromium of bondcoat thus forming: Cr2Tin-2O2n-1. The
formation of these compounds by the same mechanism is supported by the work of C.Winde
[180] where the Cr-films were deposited on the surface of TiO2-crystals and then the influence
of temperature and surface stoichiometry was studied.
TinO2n-1 + Cr → Cr2Tin-2O2n-1 (Eq. 3)
Another mechanism of formation of these types of compounds can be explained by the fact that
some places were left uncoated during TiN deposition while the other sites demonstrated cracks
within the TiN coating. These sites provided paths for oxygen and as a result formation of
alumina and chromium oxides took place at the interface of the topcoat-bondcoat. Thus the other
mechanism of the formation of Cr2Tin-2O2n-1 could be that the chromium oxides directly reacted
with titanium oxides at high temperature.
78
TinO2n-1 + Cr2O3 → Cr2Tin-2O2n-1 (Eq. 4)
These types of reactions are also explained by other researchers like Somyia et al. [179] and
Harju et al. [181].
Spectrum Al Ti Cr Fe Co Ni Y Zr
Spectrum 1 3.34 59.17 20.07 0.12 6.44 3.19 0.51 4.54
Spectrum 2 5.62 44.87 13.98 0.37 7.26 4.14 6.77 14.12
Spectrum 3 3.69 58.45 17.02 0.24 5.55 2.78 1.49 8.47
Table 5.1: Average chemical compositions (% atomic) of various phases marked in Fig.5.12
(TiN modified sample exposed for 50 hours at 950⁰C)
Fig.5.12: Different sites are marked from where EDS analysis at bondcoat-topcoat
interface were taken, in TiN modified sample after 50 hours exposure. EDS analysis are
represented in Table 5.1.
79
In case of standard bondcoat system, the oxides formation was rather different. It was observed
that after 50 hours exposure at high temperature, spinels and perovskite type structures were
formed. In this regard, aluminum and chromium oxides were formed firstly and then reacted with
the oxides of Ni and Co forming: CoCr2O4, NiCr2O4, NiCrO3, CoNiO3 and NiCrO4 phases [165].
These phases formed at the interface of topcoat-bondcoat. The formation of some of the above
mentioned spinels and perovskite structures in standard bondcoat system are demonstrated in
Fig.5.13, whereas, Table 5.2 shows the atomic percentage of these compounds.
%age, atomic Possible Phase
Location Al Cr Co Ni
1 8 19 36 37 CoNiO3
2 43 36 10 10 AlCrO3
3 39 45 9 7
4 49 15 18 17 Al2O3
5 6 21 38 35 CoNiO3
Table 5.2: Average chemical compositions (% atomic) of various phases marked in Fig.5.13
(standard TBC sample exposed for 50 hours at 950⁰C)
Fig.5.13: Standard TBC sample (after 50 hours exposure) demonstrating topcoat-bondcoat
interface, site 1 to 5 are shown from where the EDS analysis was taken and reported in Table 5.2.
80
5.1.2 Delaminated Topcoat Obtained after Hot Corrosion
The topcoat of the TiN modified sample that was subjected to hot corrosion test for 50 hours,
was delaminated by chemical etching and observed in SEM. The light and dark grey regions
were observed as attached to the topcoat. EDS observations revealed that these grey regions were
rich in alumina and chromium oxides, Fig.5.14. It seemed that the layers of alumina and
chromium oxide were present on that side of topcoat which faced the bondcoat. Further,
investigation revealed that patches of chromium-titanium phases having distinct physical
characteristics were also present. The chromium-titanium rich phase with crystals like structure
was present closer to the topcoat surface, Fig.5.15. Chemical composition, in %atomic, analysis
showed the similar composition as was determined in the cross sections of the same samples,
Fig.5.12. This confirms the presence of chromium-titanium phase with Ti:Cr ratio in the range of
2.5 to 4, referring to the fact that Cr2Tin-2O2n-1 (Cr2Ti6O15, Cr2Ti5O13 and Cr2Ti7O17) phases were
formed.
Fig.5.14: Delaminated topcoat of TiN modified sample (after 50 hours exposure) demonstrating the
regions which were broken away from the bondcoat. Two regions rich in alumina and chromium
oxides
81
(b)
Fig.5.15: (a) TiN modified sample (after 50 hours exposure) showing, patches (box) of chromium-
titanium phases having crystals like structure
(b)Schematic representation of delaminated topcoat showing layers of alumina and chromium oxide
and layer with chromium-titanium phase
Topcoat
Chromium-titanium rich
layer
Alumina and
chromium oxide
Substrate
Bondcoat
Delaminated
topcoat
(a)
82
5.1.3 X-Ray Diffraction Analysis
5.1.3.1 As-sprayed topcoat
X-Ray diffraction analysis of as-sprayed TiN modified samples along with standard samples was
made. It was observed that 100% tetragonal-ZrO2 structure was formed after spraying. However,
after 50 hours, at high temperature in hot corrosion environment, presence of monoclinic-ZrO2
phase along with yttrium vanadate (YVO4) was observed. Vanadate salts attacked Y2O3, present
in the solid solution of ZrO2, and thus destabilized tetragonal-ZrO2, Fig.5.16. The formation of
monoclinic phase is associated with volumetric changes [165] and thus leading to delamination
of the top surface.
Fig.5.16: XRD patterns comparing the scans of as-sprayed coating with the sample exposed
at 950ºC for 50 hours.( PCPDF#●821968 ◊ 830939 ○811548 )
83
5.1.3.2 Delaminated Coating of TiN Modified Sample
The portion of delaminated topcoat that faces the bondcoat, (Fig. 5.15b), is important because
diffusion and oxidation processes occurred through this interface during the hot corrosion. It is
also important to investigate the reasons why TiN-modified bondcoats demonstrated relatively
better results than standard bond-coated samples. Furthermore, it was necessary to confirm the
EDS analysis (% atomic) results showing the formation of Cr2Tin-2O2n-1 phase, as discussed in
previous section.
In this regard the grey surface (portion next to bondcoat) of the delaminated topcoat, obtained
after chemical etching, was analyzed with XRD. The results demonstrated that after 50 hours at
high temperature in hot corrosion environment, multiple phases were formed at the bondcoat-
topcoat interface. In these phases, alumina, chromium oxide (Cr3O4, Cr2O3) and the formation of
Cr2Ti5O13 was confirmed, Fig.5.17.
Fig.5,17: XRD pattern showing different phases formed in the delaminated coating of TiN
modified sample exposed at 950ºC for 50 hours. .( PCPDF#♦ 290460 ●290063 ○120559
•060504 )
84
5.1.4 Conclusion
It was demonstrated in this study that TiN-modified bondcoat could enhance the oxidation
properties of TBC systems, in hot corrosion environment. This improvement was attributed to
the formation of Cr2Tin-2O2n-1 phase exhibiting reasonably good stability at high temperature
against oxidation and delamination. In comparison to the above, standard TBC systems i.e
without bondcoat modification delaminated faster due to the formation of spinels and perovskite
structures.
85
5.2 Thermal cycling behavior of air plasma sprayed thermal barrier coatings
on Inconel X750 alloy with and without TiN modification of ‘bondcoat’.
Thermal shock behavior is very important among the properties of coatings related to the high
temperature application of gas turbines. Many researches have indicated that the thermal shock
resistance of partially stabilized zirconia coatings mainly depends on the heating conditions and
microstructure. In this study, thermal shock behavior of two types of thermal barrier coating
(TBC) systems developed by air plasma spraying system on nickel base superalloy, were studied.
These coatings were composed of a Y2O3-stabilized ZrO2 topcoat and a CoNiCrAlY bondcoat
known as standard TBC and TiN-modified bondcoat TBC in which bondcoat was modified by
deposition of thin layer of titanium nitride on bondcoat, employing a physical vapor deposition
technique. Both TBC systems were exposed to intense thermal cycling consisting of direct
exposure of samples to 950˚C, holding up to 5 min and then water quenching. Residual stresses
at the interface of topcoat and bondcoat were also determined by specimen curvature method in
both coating systems.
5.2.1 Result and Discussion
5.2.1.1 Cross-section of as-sprayed coatings
The cross-section of the coated samples revealed typical APS coating features i.e. micro-cracks,
lamella of semi-molten particles and shrinkage cavities. It was estimated that about 8-12% pores
were present in the topcoat. Typical lamellar structure was observed predominantly in the
bondcoat after the APS process, Fig.5.5. It was noticed that the thickness of TiN deposited on
bondcoat varied from 6 to 10 µm, Fig.5.6; moreover, the sputtered layer was observed to be
deposited non-uniformly at few locations, Fig.5.7.
5.2.1.2 Thermal Cycling
Both standard TBC and TiN modified samples were simultaneously exposed to thermal cycling.
It was observed that both systems sustained up to 116 thermal cycles without any noticeable
damage, Fig.5.18. However, further cycling caused initiation of edge cracking/spalling in both
samples. This degradation further increased as the number of thermal cycles increased, Fig.5.18.
After 225 cycles, the standard TBC samples spalled 30% of the topcoat, whereas, the TiN-
modified samples spalled only 5%, Fig. 5.18.
86
Fig.5.18: TiN-modified samples and (b) Standard TBC samples, showing photographs of top surfaces
after intense thermal cycling
87
The cross-sections of both the systems were also observed, after thermal cycling. It was revealed
that the topcoat began to delaminate from the top surface (Fig.5.19) in the form of fine lamella,
due to severe thermal shocks. The edges of the cross-section demonstrated severe spalling in
both the systems. It was observed that in standard TBC system, not only the topcoat but also
bondcoat spalled during cycling. However, in case of TiN-modified samples, this spalling was
not significant, Fig.5.20. In both systems, cracks at the interface of topcoat and bondcoat were
revealed as shown in Fig.5.21. It was further noticed that in case of standard TBCs, the bondcoat
also cracked and the oxidation penetrated into the substrate, after spallation of topcoat. Similar,
findings were also observed in the TiN-modified samples, however, intensity of the damage was
relatively mild compared with standard TBC samples, Fig.5.22.
.
Fig.5.20: Cross-sections of (a) TiN modified samples and (b) Standard TBC samples
showing intensity of spalling near edges during cycling
a b
Fig.5.19: Cross-section of (a) TiN modified samples and (b) Standard TBC samples
showing delamination of topcoat
88
It was observed that that both types of samples started to curve from edges after 25 cycles. This
warping further increased due to intense thermal cycling which was aggravated by quenching
from high temperature, Fig.5.23; both standard TBC samples and TiN-modified samples
demonstrated similar type of warping.
In order to observe the coating effect on bending, an uncoated sample was also exposed to
thermal cycling along with both types of coated systems. It was observed that the uncoated
samples deformed non-uniformly showing multi-dimensional warping as shown in Fig.5.24.
Fig.5.21: Cross-sections of (a) TiN modified samples and (b) Standard TBC samples
showing cracks at the interface of topcoat and bondcoat
Fig.5.22: Cross-sections of (a) TiN-modified sample and (b) Standard TBC sample: shows bondcoat
cracking and the oxidation penetration into the substrate; TiN-modified sample exhibits better
resistance
89
However, the coated samples demonstrated only uni-directional bending effect (Fig.5.24). It was
concluded that the coating restricted the bending in the material to some extent.
5.2.1.3 Residual Stress Measurement
Residual stress can be measured by different methods like XRD, hole-drill method, fringes
method and from the knowledge of specimen curvature [182-185]. Specimen curvature method
can directly help to determine stress level at the interface of coating and the substrate. In other
methods like XRD and hole-drill method, it is difficult to measure the stresses at the interface of
substrate and coating. In specimen curvature method, the largest stress in the coating-substrate
specimen exists at the interface which is given by
𝜎𝑐(𝑖𝑛𝑡𝑒𝑟𝑓𝑎𝑐𝑒) = −𝑃
𝑏ℎ− 𝐸′𝑐𝛿
where 𝛿 is the vector and can be measured from the interface to the free surface and can be
represented as
𝛿 =ℎ2𝐸′𝑐 − 𝐻2𝐸′𝑠
2(ℎ𝐸′𝑐 + 𝐻𝐸′
𝑠)
In this equation h and H are the thicknesses of the coating and substrate respectively, 𝐸′𝑐 and 𝐸′𝑠
are the modified young’s modulus for coating and substrate, respectively, and can be expressed
as
𝐸′𝑐 =𝐸(𝑐𝑜𝑎𝑡𝑖𝑛𝑔 )
1−𝑟(𝑐𝑜𝑎𝑡𝑖𝑛𝑔) and 𝐸′𝑠 =
𝐸(𝑠𝑢𝑏𝑠𝑡𝑟𝑎𝑡𝑒)
1−𝑟(𝑠𝑢𝑏𝑠𝑡𝑟𝑠𝑡𝑒)
where 𝐸(𝑐𝑜𝑎𝑡𝑖𝑛𝑔) and 𝐸(𝑠𝑢𝑏𝑠𝑡𝑟𝑎𝑡𝑒) are the young modulus for coating and substrate, 𝑟(𝑐𝑜𝑎𝑡𝑖𝑛𝑔) and
𝑟(𝑐𝑜𝑎𝑡𝑖𝑛𝑔) are the Poisson’s ratio of the coating and substrate
The curvature (𝐾) produced due to thermal cycling can be calculated by using the equation
𝐾 =8𝑙
𝐿2
where ′𝑙′ is the height of the curvature and ′𝐿′ is the traversing length of the sample, Fig.5.25.
The curvature of the sample can be used to calculate the stress within the coating [186].
90
The misfit strain i.e. ∆∈ between the sprayed coating and the substrate produced due to thermal
contraction of the molten splats can be expressed as
where, h is the thickness of the sprayed coating and H is the thickness of substrate. The misfit
strain (∆∈) thus can be calculated as
∆∈= ∆∝ ∆𝑇
where ∆𝑇 is increase in the temperature
From the above strain, the force (P) per unit width of the specimen (b) is given by
In the above two systems, the interfacial stress between the topcoat and bondcoat was calculated,
for which the thickness of the topcoat (h) was taken as 200 µm each for standard as well as TiN-
modified coatings. Bondcoat is considered as substrate for topcoat, since the interfacial stress is
determined only for topcoat-bondcoat interface. This interface is considered generally, since
spalling is usually believed to initiate from this region/plane [165] Bondcoat thickness for
standard and TiN-modified system was considered as 100 µm and 350 µm, respectively. The
stiffness (𝐸𝑐 and 𝐸𝑠) of the topcoat and substrate (which is bondcoat in this case) were previously
found to be 40000 N/mm2
and 86000 N/mm2, respectively [187]. Stresses calculated from the
above equations for the two systems are shown in Fig.5.26, as resulted from number of cycles. It
can be deduced that both coating systems demonstrated almost similar type of stresses. TiN-
modified coating interface demonstrated better thermal shock resistance for equal magnitude of
interfacial stress. This can be attributed to formation of thermal shock resistant phase during
oxidation of TiN, as explained below.
91
Fig.5.23: (a) Standard TBC samples and (b) TiN modified samples, showing samples’ warping at edges; the
curvature increased with increase in numbers of cycles
92
Fig.5.24: (a) Standard TBC samples and (b) substrate sample without coating, showing that the sample
without coating deformed in multiple directions
𝐿 (25mm)
Substrate 𝑙
H
h
1
𝐾
-P
P -P
P
Fig.5.25: Schematic illustration to calculate the stress in the coatings by using curvature
93
It has been discussed in our earlier study [188] that in TiN-modified TBCs TiN destabilizes and
forms other compounds in the presence of abundant oxygen when exposed to high temperature.
TiN oxidizes to Magneli phase (TinO2n−1 ; 4 ≤ n ≤ 9) having a complicated defect structure [174-
176]. The point defects present in these system are dominated by oxygen vacancies and titanium
interstitials [175, 177]. These so-called Magneli phase have long order defect structures [175].
These phases also react with chromium oxides and may form a series of homologous structures
i.e. Cr2Tin-2O2n-1 (6 ≤ n ≤ 9). Cr2Tin-2O2n-1 phase is known for its stability against thermal stresses
and oxidation [178].
5.2.1.4 Effect of Thermal cycling on weight loss
During thermal cycling, weight loss of the samples was determined and has been shown in
Fig.5.27. It can be observed that both types of coating systems behaved similarly up to 100
cycles, however, at prolonged exposure, the weight loss of standard TBC system increased
rapidly compared with TiN-modified coating system. This indicated that the interface between
topcoat and bondcoat was better in case of TiN-modified system.
Fig.5.26 Residual stresses as a result of thermal cycling
94
5.2.1.5 XRD observation
XRD analyses of as-sprayed TiN-modified samples were performed along with standard
samples. It was observed that 100% tetragonal-ZrO2 structure was formed after spraying; the
structure remained unchanged after 225 thermal cycles (Fig. 5.28) because samples were
exposed to high temperature for shorter duration.
Fig.5.27: Weight loss of the samples during thermal cycling
Fig.5.28 XRD patterns comparing the scans of Standard TBC coating with TiN-
modified coating samples after thermal cycling
95
5.2.2 Conclusion
The experimental results showed that both systems exhibited good resistance against thermal
shocks for up to 116 cycles. After 225 cycles, the standard TBC samples spalled 30% of the
topcoat, whereas, the TiN-modified samples spalled only 5%. Residual stresses determined by
Specimen curvature method at the interface of topcoat and bondcoat, were found to be similar in
both coating systems. Based on above results it could be concluded that TiN-modified coating
interface was better thermal shock resistant. This is believed to be due to formation of a better
thermal shock resistant phase during oxidation of TiN; Cr2Tin-2O2n-1 (6 ≤ n ≤ 9). Cr2Tin-2O2n-1
phase is known for its stability against thermal stresses and oxidation.
96
Chapter-6
6 Results and Discussion-III
6.1 Effect of bondcoat thickness on hot corrosion of ZrO2-8Y2O3 thermal
barrier coating
Hot corrosion over extended exposures reduces durability of the system. Therefore, there is a
requirement, to develop new design approaches for TBCs in order to operate in hostile
environment at high temperatures. In this Chapter, the effect of bondcoat thickness on the hot
corrosion resistance has been discussed. Material used for substrate is Udmet-720 (wt.%:
Co=14.7, Cr=17.9, Mo=3.0, W=1.3, Ti=5.0, Al=2.5, C=0.03 Ni=balance). Chemical
Compositions (wt.%), particle size range and morphology of the spraying powders utilized to
deposit TBCs (MCrAlY-bondcoat + YSZ-topcoat) are given in Table 3.3. All important
spraying parameters for both topcoat and bondcoat are mentioned in the Table 3.4. Three set of
samples with varying bondcoat and thicknesses and same topcoats thickness were produced. Hot
corrosion test of these samples were carried out in 50wt.% Na2SO4 + 50wt.% V2O5 molten salt at
950ºC for 50 hours. The characterizations of the coatings included XRD, SEM and optical
microscopy. Results indicated that TBCs with thick bondcoat delayed the hot corrosion process
as compared to the TBCs with conventional bondcoat.
6.1.1 Result and discussion
After the exposure of first 10 hours the edge of the topcoat of all the samples demonstrated
insignificant spalling, Fig.6.1. These edges were presumably spalled due to the thermal stresses
generated in samples by directly charging them at high temperature. In further 20 hours it was
revealed that the samples S1 & S2 spalled relatively greater than the samples S3. After 50 hours
the S1 & S2 samples showed about 12-15% spalling, whereas S3 samples demonstrated only 8-
10% spalling from the edges, Fig.6.1.
97
Fig.6.1: S1, S2 and S3 samples showing appearance of top surfaces after varying exposure times,
treated at 950°C in a hot corrosion environment
98
The surface of the topcoat of as-sprayed samples revealed splats of YSZ, porosities and some
semi-melted particles during SEM observation (Fig.6.2) whereas, after the exposure of 50 hours,
in the hot corrosion environment rod-like features and agglomerated crystals were observed on
the surface, Fig.6.2. EDS analyses of these agglomerates and the rod, revealed as ZrO2 and
YVO4, respectively. It seemed that degradation of the topcoat began by formation of YVO4 rods.
Some of these crystal rods found to be present on the surface while others were observed as
firmly adhered to the topcoat. During exposure of salt mixture i.e. 50wt.% Na2SO4 and 50wt.%
V2O5 reacted with each other and formed a eutectic compound NaVO3.
Na2SO4 + V2O5 → 2NaVO3 + SO2 + 1/2O2 (1)
Fig.6.2: SEM micrographs showing surface morphology of samples (a) as sprayed and
(b) after 50 hours exposure to hot corrosion environment
99
The NaVO3 formed during the above reaction, acted as a corrosion catalyst and served as an
oxygen carrier. NaVO3 entered into the crevices and pores, present within the plasma sprayed
coatings, reacted vigorously with Y2O3, added in the pure zirconia solid solution for stabilizing
the tetragonal crystal structure. The resultant reactant yielded YVO4 [165].
ZrO2 (Y2O3) +2NaVO3 → ZrO2 + 2YVO4 + Na2O (2)
C
CC
Substrate
290±10µm
Substrate
Topcoat
Fig.6.3: SEM micrographs showing cross section of as-sprayed samples (a) S1 (c) S2 (e) S3 with
different bondcoat thicknesses (b,d,f) high magnifications of bondcoats showing typical lamellar
structure
c
b a
f e
d
Topcoat
290±10µm Bondcoat
Substrate
100
6.1.1.1 Cross sectional observation
The cross-section of as-sprayed samples revealed that the thickness of the bondcoat in S1 was
100±10 µm, S2 was 290±10 µm and S3 was 380±10 µm (Fig.6.3). Further, the cross-section of
all the systems exhibited typical air plasma sprayed coatings features i.e. the topcoat having
porosity ranging from 7-12%. The bondcoats have typical lamellar structure surrounded with
thin oxide layer, formed during the air plasma spraying process, Fig.6.3. It was noticed that the
interfaces between bondcoat-topcoat and bondcoat-substrate were generally reasonably good.
After 50 hours exposure at 950°C in hot corrosion environment, cross sections of the two
extreme case samples (S1 & S3) revealed that samples (S1) with thin bondcoat had relatively
thick thermally grown oxide (TGO) at discrete places as compared to sample (S3) with thinner
bondcoat, Fig.6.4. Further, it was noticed that the oxidation of the bondcoat of sample S3 is
comparatively less severe compared with sample S1, also cracks started to appear in sample S1,
Fig.6.4a.
In order to understand the oxidation phenomenon at the interface of bondcoat and topcoat, the
cross sections were analyzed in detail. The cross section of the sample revealed multiple features,
Fig.6.5. EDS analyses showed that each feature had its own distinct chemical composition. The
chemical composition (in %atomic) of these features is presented in Table 6.1. In order to
confirm the formation of various phases, the topcoat was delaminated by chemical etching
Fig.6.4: SEM micrographs after 50 hours exposure to hot corrosion environment
showing, cross section of samples (a) S1 with thick TGO and cracks (arrows) and (b) S3
with thin TGO
101
process. As a result, a delaminated topcoat was obtained with attached phases, formed during the
hot corrosion. After washing with water, the XRD of the delaminated topcoat was carried out.
From XRD analysis it revealed that the formation of spinels consisting of NiCr2O4 and CoCr2O4
along with perovskite structure of NiCrO3 at the interface of bondcoat and the topcoat, Fig.6.6.
Volumetric changes occurred due to development of these structures and forced to crack the
interface. Samples, after 50 hours hot corrosion testing, demonstrated that overall oxidation
condition of bondcoat was more severe in sample S1 as compared to the sample S3, Fig.6.7.
%age, atomic Possible
Phase Location Al Cr Co Ni
1 9 20 35 36 CoNiO3
2 43 35 11 10 AlCrO3
3 38 46 9 7
4 49 14 19 17 Al2O3
5 6 23 37 34 CoNiO3
Table-6.1: Chemical composition of various phases marked in
Fig.6.5 (exposed for 50 hours at 950⁰C)
Fig.6.5: SEM micrograph showing discrete sites at the interface in a sample exposed
for 50 hours showing different features (Site-1 and 5). Site-1 to 5 represent locations
for EDS analyses reported in Table 6.1
102
EDS analysis at the boundaries of the splats showed that “vanadium” was present near the
interface of both samples. However the EDS analysis in the depth of depth of bondcoat revealed
that no “vanadium” was present in sample S3, whereas, sample S1 revealed “vanadium” in
analysis, Fig.6.8, this demonstrates the delay of vanadium oxide in crossing the diffusion barrier
of alumina in case of sample S3. This may be due to initially formation of NaVO3 takes place
and it start dissolving alumina and chromia, the bondcoat continues to promote their formation,
thereby prevent the diffusion of corrosive species. If a point is reached where chromium and
aluminum levels in the bondcoat falls below the level at which protective alumina and chromia
scales cannot be formed preferentially, faster inward diffusion of corrosive species and outward
diffusion of alloying elements of super alloy takes place, interact with each other to form
corrosion products causing Volumatic changes which promote delamination of TBC [189]. As
indicated in Fig.4.13 that the Al diffused from bondcoat to both topcoat and substrate during hot
corrosion. In case of thick bondcoat more reservoir of Al is present and it prevent the diffusion of
corrosive species in bondcoat by maintaining the formation of alumina and chromia for longer
time as compared to thin bondcoat. More ever, inter-diffusion between the bondcoat and the
Fig.6.6: XRD pattern showing various phases formed in the delaminated coating of S1
sample exposed at 950ºC for 50 hours. (JCPDF# •760144 ●210596 ♦220748 ○750198
♥801668 ♠100188 ♣731704)
2Ѳ
I/I0
103
underlying superalloy because minor alloy additions (S, Ta and W) increased the growth rate of
the TGO layer and promoted the formation of non-protective oxide scale [190] because of this
sample with thin bondcoat was more affected.
Fig.6.7: Samples exposed at 950ºC for 50 hours, indicating more oxidation in (a) S1 with thin
bondcoat as compared to (b) S2 with thick bondcoat
Fig.6.8: Samples exposed at 950ºC for 50 hours, EDS analysis at the boundaries of the splats
showed that “vanadium” is present in (a) sample S1, whereas,(b) sample S2 revealed no
“vanadium” in analysis
104
6.1.2 Conclusion
The coatings were exposed to hot corrosion using corrosive salts of Na2SO4 and V2O5. The
results demonstrated the delay of corrosive species in crossing the diffusion barrier of alumina in
case of sample S3. This may be due that more reservoir of Al is present in case of thick bondcoat
and it prevent the diffusion of corrosive species in bondcoat by maintaining the formation of
alumina and chromia for longer time as compared to thin bondcoat.
105
6.2 Behavior of air plasma sprayed thermal barrier coatings with different
bondcoat thicknesses, subject to intense thermal cycling
Two sets of samples (S1 & S3) having same topcoat thickness but with different bondcoats
thicknesses were produced by air plasma technique. The cross section of as-sprayed samples
revealed that the thickness of bondcoat of sample S1 was 100±10 µm and S3 was 380±10 µm
(Fig.6.3).
6.2.1 Result and Discussion
6.2.1.1 Thermal cycling
“All TBC systems were exposed to intense thermal cycling which consisted of directly exposing
samples to 950ºC and holding up to 5 min followed by water quenching. It was noticed that both
system S1 and S3 sustained up to 100 thermal cycles without damage. However, after 100 cycles
first cracking appeared at the extreme edges of the samples; the cracking then propagated around
the sample on further cycling. After 225 cycles it was observed that 30-35% of the topcoat
spalled in case of S1 samples while the topcoat of S3 sample spalled about 10-15%, Fig. 6.9. In
general, the spallation process in thermal barrier coatings was progressive and analogous to the
fatigue in metals. Localized damage might be initiated at a relatively low fraction (5–20%) of
total life. Damage progression was often by micro-cracking extension followed by mergers into
larger cracks. After initiation of a single crack, the crack encountered stress fields associated
with the geometrical arrays of cavities at most of the grain boundaries; rapid crack link up and
macro-crack propagation occurred due to the extreme heating and cooling conditions
encountered at the edges. The cross section of samples after 225 cycles demonstrated that the
surface of the topcoat (ZrO2–7%Y2O3) was continuously delaminated during the course of
thermal cycling Fig.6.10.”
106
The cross-section of the both systems demonstrated severe spalling at the edges of samples,
Fig.6.11. Bondcoat oxidation was also associated with high temperatures. Bondcoat oxidation
has been clearly linked to spallation of the ceramic topcoat [191-193]. Bondcoats have enough
aluminum which act as a reservoir and provide a protective oxide layer against oxidation. These
layers were believed to be composed of alumina, as indicated by few researchers [194, 195].
These thermally grown oxides (TGOs) were thought to grow due to the oxidation of the
aluminum in the plasma flame and during splat formation. During thermal exposure of a TBC at
high temperature for a long time the TGO grew further. This TGO growth was believed to play a
crucial role to the life of the coating [196-199] and induced the strain energy for the crack
propagation during the spallation. After the completion of 225 cycles of water quenching the
Fig.6.9: (a) S1 samples and (b) S3 samples, showing condition of top surfaces after 225
intense thermal cycles
Fig.6.10.: Cross-section of (a) S1 samples and (b) S3 samples showing delamination of
topcoat
107
total exposure for the TBC at high temperature (950°C) was approximately 1125 min. The TGO
seems to be thicker close to the cracked edges of the TBC than at the center of the sample, after
225 cycles, it was also observed that in both systems near the cracked edges, after spallation of
the topcoat the bondcoat also started to spall out from the surface of the substrate, shown in
Fig.6.12. It was observed that both types of samples began to curve from edges during intense
thermal cycling just after 25 cycles. This curvature further increased with increase in number of
cycles upon quenching from high temperature as evident from Fig.6.13 and Fig.6.14.
Fig.6.11: Cross-sections of (a) S1 samples and (b) S3 samples showing intensity of
spalling near edges during cycling
Fig.6.12: Cross-sections of (a) S1 samples and (b) S3 samples, showing that after spallation of
the topcoat the bondcoat also started to spall near cracked edges
108
In order to observe the coatings’ effect on warping, a substrate sample without coating was also
thermally cycled along with both types of coating systems. It was observed that the uncoated
sample without coating deformed in almost all directions, Fig.6.15. Whereas, coated samples
demonstrated only one side folding/bending effect, Fig.6.13 and Fig.6.14. This indicated that the
coating restricted the movement of the bending of the substrate material. The results revealed
that the relatively increased bending /curvature was observed in samples with thinner bondcoat,
Fig.6.13 and Fig.6.14, indicating that thick bondcoat restricted bending to a larger extent in
samples compared with thin bondcoat.
Fig.6:.13 S1 samples started to curve from edges during intense thermal cycling, it was
noted that curvature increased with increase in numbers of cycles
109
Fig.6.14: S3 samples started to curve from edges during intense thermal cycling, it was
noted that curvature increased with increase in numbers of cycles
Fig.6.15: Substrate sample without coating, showing deformed in multiple directions
110
6.2.1.2 Residual Stress Measurement
Residual stress can be measured from the knowledge of specimen curvature [185]. Specimen
curvature method can directly indicate the stress level at the interface of coatings and the
substrate. In specimen curvature method, the largest stress in the coating-substrate specimen
existed at the interface as calculated by method described in chapter no.5, section 5.2.1.3.
In the above two systems, for the calculation of the interfacial stress between the topcoat and
bondcoat, the thickness of the topcoat (h) was taken as 200 µm each for S1 and S2 coatings
(Fig.5.25). Bondcoat was considered as substrate for topcoat, since the interfacial stress was
calculated only for topcoat-bondcoat interface. This interface was considered because the
spalling during thermal cycling or isothermal treatment is generally initiated from this
region/plane [200]. Bondcoat thickness for S1 and S3 coatings was taken as 100 µm and 400 µm,
respectively. The stiffness (𝐸𝑐 and 𝐸𝑠) of topcoat and the substrate (which is bondcoat in this
case) are 40000 N/mm2 and 86000 N/mm
2,respectively [187]. Stresses calculated for the two
systems against number of cycles were plotted in Fig.6.16. It can be seen that S1 coating system
had more residual stress as compared to S3 coating system. This might be due to lesser bending
in S3 coatings because S3 system has thick coating (topcoat 200 µm +bondcoat 400 µm) as
compared to S1 system (topcoat 200 µm +bondcoat 100 µm); thicker coating restricted the
movement during warping to a greater extent as described earlier.
The interface in S3 coating system was more resistant to thermal shock because of low value of
residual stresses at the interface of topcat and bondcoat due to thick bondcoat; moreover, the
thick bondcoat prevented the diffusion of corrosive species into it by maintaining the continuous
formation of alumina and chromia for longer durations as compared to thinner bondcoat.
6.2.1.3 Effect of Thermal cycling on weight loss
During thermal cycling, weight loss of the samples was measured after each 10 cycles and is
plotted in Fig.6.17. It can be seen that both type of coating systems almost behaved similar for up
to 100 cycles but afterwards the weight loss of S1.system increased relatively more rapidly as
compared to S3 coating system. This indicated that the spalling in S1 samples was greater as
compared to S3 samples. The weight loss in samples during thermal cycling was from ablation
and spalling of the ceramic coating.
111
Fig.6.16: Stresses calculated using the curvature of samples as a result of thermal
cycling for the two systems (S1 and S3 samples)
S3 S1
Fig.6.17 Weight loss of the samples (S1 and S3) during thermal cycling
S1 S3
112
6.2.2 Conclusion
Two sets of samples (S1 & S3) having same topcoat thickness but with different bondcoats
thicknesses were produced by air plasma technique. These were exposed to intense thermal
cycling consisting of directly exposure of samples to 950ºC, holding up to 5 min and then water
quenching. Residual stresses were measured from the specimen curvatures. The results indicated
that the interface in S3 coating system was relatively more thermal shock resistant because of
low residual stress at the interface of topcoat and bondcoat due to thick bondcoat because
warping in the material was greater restricted in thick bondcoat as compared to thin bondcoat.
Moreover, thick bondcoat delayed the diffusion of corrosive species in it by maintaining the
formation of alumina and chromia for longer time as compared to thin bondcoat.
113
Chapter-7
Summary
The research work undertaken could be divided into three segments and the conclusions
regarding each portion are as follows.
Hot corrosion of yttria-stabilized zirconia coating, in a mixture of sodium sulfate and
vanadium oxide at 950oC
Yttria stabilized zirconia along with CoNiCrAlY bondcoat were deposited by air plasma
spaying on Inconel-X750 coupons. The coatings were exposed to hot corrosion using
corrosive salts of Na2SO4 and V2O5. The results demonstrated that the Y2O3 present in
solid solution of ZrO2 reacted with the salt mixture and formed rods of yttrium vanadate
(YVO4). Reaction of oxygen with various metallic elements in the bondcoat resulted in
formation of spinels consisting of NiCr2O4 and CoCr2O4 along with perovskite structure
of AlCrO3, NiCrO3 and CoNiO3 and oxides (NiCrO4, CoNiO2, CoNiO4) at the
interface of bondcoat and the topcoat. Development of these structures forced to crack the
interface.
Chemical composition profile revealed that various alloying elements diffused from
bondcoat to the substrate and from substrate to the bondcoat which might had altered
mechanical properties of the interface. Similarly diffusion also took place at interface of
bondcoat and the topcoat. The diffusion of elements (chromium, cobalt, nickel and
aluminum) from bondcoat to topcoat became possible upon the cracking or dissolution of
TGO. The oxidation of these elements then leads to the formation of spinels and
perovskite structure.
XRD analyses determined that m-ZrO2 formed along with YVO4 with increasing
exposure time at high temperature. Moreover, a shift in high angle peaks indicated high
level of stresses present in the coating due to the formation of YVO4 and m-ZrO2.
114
Evaluation of titanium nitride modified bondcoat system used in thermal barrier coating in
corrosive salts environment at high temperature
It was demonstrated in this study that TiN-modified bondcoat could enhance the
oxidation properties of TBC systems, in hot corrosion environment. This improvement
was attributed to the formation of Cr2Tin-2O2n-1 phase exhibiting reasonably good stability
at high temperature against oxidation and delamination. In comparison to the above,
standard TBC systems i.e without bondcoat modification delaminated faster due to the
formation of spinels and perovskite structures.
The results of intense thermal cycling (consisting of directly exposure of samples to
950ºC, holding up to 5 min and then water quenching) showed that both systems
exhibited good resistance against thermal shocks for up to 116 cycles. After 225 cycles,
the standard TBC samples spalled 30% of the topcoat, whereas, the TiN-modified
samples spalled only 5%. Residual stresses determined by Specimen curvature method at
the interface of topcoat and bondcoat, were found to be similar in both coating systems.
Based on above results it could be concluded that TiN-modified coating interface was
better thermal shock resistant. This is believed to be due to formation of a better thermal
shock resistant phase during oxidation of TiN; Cr2Tin-2O2n-1 (6 ≤ n ≤ 9). Cr2Tin-2O2n-1
phase is known for its stability against thermal stresses and oxidation.
Effect of bondcoat thickness on hot corrosion of ZrO2-8Y2O3 thermal barrier coating
Two sets of samples (S1 & S3) having same topcoat thickness but with different
bondcoats thicknesses were produced by air plasma technique. The coatings were
exposed to hot corrosion using corrosive salts of Na2SO4 and V2O5. The results
demonstrated the delay of corrosive species in crossing the diffusion barrier of alumina in
case of sample S3. This may be due that more reservoir of Al is present in case of thick
bondcoat and it prevent the diffusion of corrosive species in bondcoat by maintaining the
formation of alumina and chromoia for longer time as compared to thin bondcoat.
The samples were exposed to intense thermal cycling and residual stresses were
measured from the specimen curvatures. The results indicated that the interface in S3
115
coating system was relatively more thermal shock resistant because of low residual stress
at the interface of topcoat and bondcoat due to thick bondcoat because warping in the
material was greater restricted in thick bondcoat as compared to thin bondcoat. Moreover,
thick bondcoat delayed the diffusion of corrosive species in it by maintaining the
formation of alumina and chromia for longer time as compared to thin bondcoat.
116
Publications and presentations
Journal Publications
1. Imran Nazir Qureshi, Muhammad Shahid, A. Nusair Khan, Evaluation of titanium nitride
modified bondcoat system used in thermal barrier coating in corrosive salts environment
at high temperature, journal of thermal spray technology, volume24, Number 7,
October 2015 (published online- DOI 10.1007/s11666-015-0344-x)
2. Imran Nazir Qureshi, Muhammad Shahid, A. Nusair Khan, Hot corrosion of yttria-
stabilized zirconia coating, in a mixture of sodium sulfate and vanadium oxide at 950oC,
journal of thermal spray technology, volume24, Number 8, December 2015 (published
online- DOI 10.1007/s11666-015-0374-4)
3. Imran Nazir Qureshi, Muhammad Shahid, A. Nusair Khan, Effect of Bondcoat Thickness
on High Temperature Hot Corrosion of ZrO2-8Y2O3 thermal barrier coatings, Actca
Physica Polonica-A, 2015 (DOI:10.12693/APhysPolA.128.B-314)
Conference Publication
1. Imran Nazir Qureshi, Muhammad Shahid, A. Nusair Khan, Effect of Na2SO4-
V2O5 Mixture on Y2O3 Stabilized ZrO2 Thermal Barrier Coatings Exposed at High
Temperature, Advanced Materials Research, Vol. 1101, pp. 423-427, Apr. 2015
Conferences presentations
1. International Conference on Computational and Experimental Science and Engineering
(ICCSEN-2014), Anatalya, Turkey
2. 2015 3rd Conference on Nano and Materials Science (ICNM-III 2015), Zhuhai, China
(Excellent oral presentation award)
117
References
1. P. C. Tsai, C.F. Tseng, C. W. Yang, I. C. Kuo, Y. L. Chou and J. W. Lee, Surface &
Coatings Technology, 2013. 228: p. S11–S14.
2. L. Wang, Y. Wang, X.G. Sun, J.Q. He, Z.Y. Pan and C.H. Wang, Vacuum 2012. 86: p.
1174-1185.
3. Natin P. Padture, Maurica Cell, Eric H. Jordan, Thermal Barrier Coatings for Gas-
Turbine Engine Applications. Science, 2002. 296: p. 280-284.
4. G. Guidoni, A.Dudek, S. Patsias, and M. Anglada, Fracture Behaviour of Thermal Barrier
Coatings After High Temperature Exposure in Air. Mater. Sci. Eng. A, 2005. A397: p.
209-214.
5. M.J. Pomeroy, Coatings for Gas Turbine Materials and Long Term Stability Issues.
Mater. Des., 2005. 26: p. 223-231.
6. Y. Q. Wang, G.Sayre, Surface & Coatings Technology 2009. 203 p. 2186-2192.
7. U. Schulz, C. Leynes, K. Fritscher, M. Peters, B.S. Brings, O. Lavigne, J.M. Dorvaux, M.
Poulain, R. Mevrel, and M. Caliez, Some Recent Trends in Research and Technology of
Advanced Thermal Barrier Coatings. Aerosp. Sci. Technol., 2003 7: p. 73-80.
8. A. G. Evans, D.R.Mumm, J.W. Hutchinson, G. H. Meier and F. S. Pettit, Mechanisms
controlling the durability of thermal barrier coatings. Progress in Materials Science, 2001.
46 p. 505–553.
9. J.R. Brandon, R.Taylor, Microstructure, composition and property relationships of
plasma-sprayed thermal barrier coatings, Surf. Coat. Techol., 1992. 5(2): p. 141.
10. I.Gurrappa, Thermal barrier coatings for hot corrosion resistance of CM 247 LC
superalloy. J. Mater. Sci. Lett., 1998. 17: p. 1267-1269.
11. R.L.Jones, Some aspects of the hot corrosion of thermal barrier coatings. J.Therm.Spray
Technol, 1997. 6(1): p. p 77.
12. B.A.Nagaraj, D.J.Wortman, Development of Corrosion Resistant Coatings for Marine
Gas Turbine Applications Trans. ASME, 1990. 112 p. 536.
13. R. Srinivasan, J.M.Merrilea, The hot corrosion resistance of 20 mol% YTaO4 stabilized
tetragonal zirconia and 14 mol% Ta2O5 stabilized orthorhombic zirconia for thermal
barrier coating applications. Surf. Coat. Technol., 2002, . 160: p. 187.
14. R.L.Jones, India as a Hot Corrosion-Resistant Stabilizer for Zirconia. J. Am. Ceram.
Soc., 1992. 75(7): p. 1818.
15. A. Rico, J.G.Garcia, C.J. Munez, P. Poza, and V. Utrilla, Mechanical Properties of
Thermal Barrier Coatings After Isothermal Oxidation. Depth Sensing Indentation
Analysis. Surf. Coat. Technol., 2009. 203: p. 2307-2314.
16. E. Garcia, C.Cano, T. W. Coyle, M.I. Osendi and P. Miranzo, Journal of Thermal Spray
Technology 2008 (17(5-6) ): p. 865-871.
118
17. E. Withey, C.Petorak, R. Trice, G. Dickinson and T. Taylor, Journal of European
Ceramic Society 2007. 27: p. 4675-4683.
18. Y. Bai, Z.H.Han, H. Q. Li, C. Xu, Y. I. Xu, C. H. Ding and J. F. Yang Surface &
Coatings Technology, (2011). 205 (13-14) p. 3833-3839.
19. M. Saremi, A.Afrasiabi, and A. Kobayashi, Microstructural Analysis of YSZ and
YSZ/Al2O3 Plasma Sprayed Thermal Barrier Coatings After High Temperature
Oxidation. Surf. Coat. Technol., 2008. 202: p. 3233-3238.
20. A. Portinha, V.Teixeira, J. Carneiro, J. Martins, M.F. Costa, R. Vassen, and D. Stoever,
Characterization of Thermal Barrier Coatings with a Gradient in Porosity. Surf. Coat.
Technol, 2005. 195: p. 245-251.
21. R. Vaben, M.O. Jarligo, T. Steinke, D.E. Mack, and D. Stover, Overview on Advanced
Thermal Barrier Coatings. Surf. Coat. Technol., 2010. 205: p. 938-942.
22. P. Mohan, T. Patterson, V.H. Desai, and Y.H. Sohn, Degradation of Free-Standing Air
Plasma Sprayed CoNiCrAlY Coatings by Vanadium and Phosphorus Pentoxides. Surf.
Coat. Technol., 2008. 203: p. 427-431.
23. H. Liu, S. Li, Q. Li, Y. Li, and W. Zhou, Microstructure, Phase Stability and Thermal
Conductivity of Plasma Sprayed Yb2O3, Y2O3 Co-Stabilized ZrO2 Coatings. Solid State
Sci., 2011. 13(3): p. 513-519.
24. T. Patterson, A. Leon, B. Jayaraj, J. Liu, and Y.H. Sohn, Thermal Cyclic Lifetime and
Oxidation Behavior of Air Plasma Sprayed CoNiCrAlY Bondcoats for Thermal Barrier
Coatings. Surf. Coat. Technol., 2008. 203: p. 437-441.
25. K.W. Schlichting, N.P.Padtrue, E.H. Jordan, and M. Gell, Failure Modes in Plasma-
Sprayed Thermal Barrier Coatings. Mater. Sci. Eng., 2003. A342: p. 120-130.
26. A. J. Burns, R.Subramanrian, B. W. Kempshall and Y. H. Sohn, , Surface and Coatings
Technology, (2004). 177-178 p. 89-96.
27. A.Nusair Khan, J.Lu, H. Liao, Effect of residual stresses on air plasma sprayed thermal
barrier coatings. Surf. Coat. Technol, 2003. 168: p. 291-299.
28. H.Edris, D.G.McCartney, A.J.Sturgeon, Microstructural characterization of high velocity
oxy-fuel sprayed coatings of Inconel 625,. J. Mat. Sci., 1997. 32: p. 863 - 868.
29. Y. Bai, L. Zaho, J.J.Tang, S.Q.Ma, C.H.Ding, J.F.Yang, L.Yu and Z.H.Han, Ceramics
International, 2013. 39: p. 4437–4448.
30. H. Jamalin, R.Mozafarinia, R. S. Razavi and R. Ahmadi-Pidani, Ceramics International,
2012. 38 p. 6705–6712.
31. A. N. Khan and J. Lu, Surface & Coatings Technology (2007) 201 p. 4653-4658.
32. J. H. Qiao, R. Bolot, H. Liao, P. Bertrand and C. Coddet, International Journal of
Thermal Sciences 2013. 65: p. 120-126.
33. I.Gurrappa, Identification of hot corrosion resistant MCrAlY based bondcoatings for gas
turbine engine applications. Surf. Coat. Technol., 2001. 139: p. 272.
119
34. I.Gurrappa, Influence of alloying elements on hot corrosion of superalloys and coatings:
necessity of smart coatings for gas turbine engines. Mater. Sci. Technol., 2003. 9: p. 178.
35. I.Gurrappa, Hot corrosion behavior of Nimonic-75. J. High Temp. Mater. Sci., 1997. 38:
p. 137.
36. I.Gurrappa, Effect of aluminum on hot corrosion resistance of MCrAIY-based
bondcoatings, . J. Mater. Sci. Lett., 2001. 20: p. p 2225.
37. I, Gurrappa, Overlay coatngs degradation—an electrochemical approach. J. Mater. Sci.
Lett., 1999. 18: p. 1713.
38. K.L. Luthra, D.A. Shore, Mechanism of Na2SO4 Induced Corrosion at 600°-900°C. J.
Electro Chem., 1980. 127: p. 2202.
39. C. Leyens, K.Fritcher, M. Peters, Oxide scale formation on a MCrAlY coating in various
H2-H2O atmospheres, . Surf. Coat. Technol., 1996. 139(1-2): p. 133.
40. W. Beele, N.Czech, W.J. Quadakkers, W. Stamn, Long-term oxidation tests on a re-
containing MCrAlY coating, . Surf. Coat. Technol., 1997. 94-95( ): p. 41.
41. A. Strawbridge, H.E.Evans, C.B. Ponton, Spallation of oxide scales from NiCrAlY
overlay coatings. Mat. Sci. Forum, 1997. 251- 254: p. 365.
42. C. Leyens, K.Fritscher, M. Peters, W.A. Kayser, Transformation and oxidation of a
sputtered low-expansion Ni-Cr-Al-Ti-Si bondcoating for thermal barrier systems, . Surf.
Coat. Technol., , 1997. 94-95(155).
43. S.R.J.Saunders, J.R.Nicholl, Hot salt corrosion test procedures and coating evaluation,.
Thin Solid Films, 1984. 119(1): p. 247.
44. X. Zhang, M.Watanabe, S.Kuroda, Acta Materialia, 2013. 61: p. 1037–1047.
45. W. Beele, G.Marijnissen., A. van Lieshout, Surface &Coating Technology, 1999. 120-
121: p. 61-67.
46. X. Q. Cao, R.Vassen, D. Stoever, Journal of European Ceramic Society, (2004) 24: p. 1-
10.
47. Inconal-X750 data sheet, the special metals corporations, U.S.AF
48. G. Levi, Current Opinion in Solid State and Materials Science, 2004. 8: p. 77-91.
49. H. Zhou, F.Li, B. He, J. Wang and B-De Su, Surface & Coating Technology, 2007. 201
p. 7360-7367.
50. A.J. Burns, R.Subramanian, B.W. Kempshall, and Y.H. Sohn, Microstructure of As-
Coated Thermal Barrier Coatings with Varying Lifetimes. Surf. Coat. Technol., 2004.
177-178: p. 89-96.
51. P. Scardi, Surface and Coatings Technology, , 1994 68/69 p. 106-112.
52. M. Tanaka, M.Hasegawa, A.F. Dericioglu and Y. Kagawa, Materials Science and
Engineering A, 2006. 419: p. 262-268.
53. A. Afrasiabi, M.Saremi, A.Kobayashi, Material Science and Engineering-A, 2008. 478:
p. 264-269.
120
54. S.Y. Park, J.H.Kim, M.C. Kim, H.S. Song, C.G. Park, Surface and Coatings Technology,
2005. 190 (2-3): p. 357-365.
55. B.A. Pint, I.G.Wright, W.Y. Lee, Y. Zhang, K. Prüßner, K.B. Alexander, Substrate and
bondcoat compositions: factors affecting alumina scale adhesion. Mater. Sci. Eng. A,
1998. 245: p. 201–211.
56. B. Lowrie, D.H.Boone, Composite coatings of CoCrAlY plus platinum. Thin Solid Flims,
1977. 45 (3) p. 491–498.
57. H.M. Tawancy, N.Sridhar, B.S. Tawabini, N.M. Abbas, T.N. Rhys-HJones, Thermal
stability of a platinum aluminide coating on nickel-based superalloys. J. Mater. Sci. Lett.,
1992. 27 p. 6463–6474.
58. H.M. Tawancy, N.M. Abbad, T.N. Rhys-Jones, Role of platinum in aluminide coatings, .
Surf. Coat. Technol. , 1991. 49 (1–3): p. 1–7.
59. H.J. Kim, M.E.Walter, Characterization of the degraded microstructures of a platinum
aluminide coating. Mater. Sci. Eng. A 2003. 360 (1–2) p. 7–17.
60. G. Naves, L.Viloria, Laboratory and field corrosion behavior of coatings for turbine
blades. Surf. Coat. Technol., 1997. 94–95: p. 161–167.
61. C. Leyens, I.G.Wright., B.A. Pint, Hot corrosion of an EB–PVD thermal-barrier coating
system at 950 °C. Oxid. Met., 2000. 54 (5–6) p. 401–424.
62. J.H. Kim, C.G.Park, M.C. Kim, Metals Mater., 2001. 7: p. 557.
63. S. Bose, High Temperature Coatings, 2007: Elsevier Science & Technology Books
64. P. Morell, D.S.Rickerby, AGARD Report 823. Neuilly-sur-Seine, France, 1998. 20: p. 1-
9.
65. Y. Tamarin, Protective Coatings for Turbine Blades, Materials Park, OH: ASM
International, 2002.
66. J.H. Wood, E.Goldman, Protective coatings, in C. T. Sims, N. S. Stoloff and W. C.
Hagel, eds, Superalloys II, 1987: New York: John Wiley and Sons.
67. J.R. Nicholls, Advances in coating design for high performance gas turbines. MRS
Bulletin, 2003. 28: p. 659–670.
68. G.W.Goward and D.H. Boone, Mechanisms of formation of diffusion aluminide coatings
on nickel-base superalloys. Oxidation of Metals, 1971. 3: p. 475–495.
69. H.W. Tawancy, N.M.Abbas, and T.N. Ryns-Jones, The role of platinum in aluminide
coatings, 1991, Surface Coatings Technology. 49: p. 1–7.
70. J.R. Davis, Handbook of Thermal Spray Technology. 2004: Materials Park, OH: The
ASM Thermal Spray Society.
71. D.H.Boone, Physical vapour deposition processes. Materials Science and Technology,
1986. 2 p. 220–224.
72. A. R. Nicholl, H.Gruner, G. Wuest and S. Keller, Future developments in plasma spray
coating. Materials Science and Technology, 1986. 2: p. 214–219.
121
73. A.R. Nicholl and G.Wahl, Oxidation and high temperature corrosion behaviour of
modified MCrAlY cast materials-Superalloys ed. W.B.K. R. H. Bricknell, M. Gell, C. S.
Kortovich and J. F. Radavich, eds. 1984: Warrendale, PA: The Metallurgical Society of
AIME.
74. H. Zhou, F.Li, B. He, J. Wang, and B.D. Sun, Air Plasma Sprayed Thermal Barrier
Coatings on Titanium Alloy Substrates. Surf. Coat. Technol, 2007. 201: p. 7360-7367.
75. T. Lauwagie, K.Lambrinou, S. Patsias, W. Heylen, and J. Vleugels, Resonant-Based
Identification of the Elastic Properties of Layered Materials: Application to Air-Plasma
Sprayed Thermal Barrier Coatings. NDT&E Int., 2008. 41: p. 88-97.
76. D.R. Clarke, Materials selection guidelines for low thermal conductivity thermal barrier
coatings. Surface and Coatings Technology, 2003. 163–164: p. 67–74.
77. B. A. Pint, I.G.Wright, W.Y. Lee, Y. Zhang, K. Prussner and K.B. Alexander, Substrate
and bondcoat compositions: factors affecting alumina scale adhesio, . Materials Science
and Engineering, 1998. A245: p. 201–211.
78. J.S. Wang, and A.G. Evans, Effects of strain cycling on buckling, cracking and spalling
of a thermally grown alumina on a nickel-based bondcoat. Acta Materialia, 1999. 47: p.
699–710.
79. W.D. Kingery, Thermal conductivity: temperature dependence of conductivity for single-
phase ceramics. Journal of the American Ceramic Society, 1955. 38 p. 251–255.
80. D.R. Clarke. and C.G. Levi, Materials design for the next generation of thermal barrier
coatings. Annual Review of Materials Research, 2003. 33 p. 383–417.
81. S. Stecura, Advanced thermal barrier system bond coatings for use on nickel-, cobalt- and
iron-base alloy substrates. Thin Solid Films, 1986. 136: p. 241–256.
82. S. Alperine, and L. Lelait, Microstructural investigations of plasma-sprayed Yttria
partially stabilised zirconia TBC. Transactions ASME: Journal of Engineering formGas
Turbines and Flows, 1994. 116 p. 258–265.
83. R. A. Miller, J.L. Smialek and R.G.Garlik, Phase stability in plasma-sprayed partially
stabilised zirconia-yttria, . Science and Technology of Zirconia, Advances in Ceramics.
Vol. 3 1981: Columbus, OH: American Ceramic Society.
84. J.A. Thompson, and T.W. Clyne, The effect of heat treatment on the stiffness of zirconia
topcoats in plasma-sprayed TBCs. Acta Materialia, 2001. 49 p. 1565–1575.
85. J.R.Brandon, and R. Taylor, The phase stability of zirconia-based thermal barrier
coatings: 1. Zirconia-yttria alloys. Surface Coatings Technology, 1991. 46: p. 75–90.
86. R.L. Jones, and D.Mess, Improved tetragonal phase stability at 1400 ◦C with scandia,
yttria-stabilised zirconia. Surface and Coatings Technology, 1996. 86–87: p. 94–101.
87. M. Leoni, R.L.Jones, and P. Scardi, Phase stability of scandia-yttria-stabilised zirconia
TBCs. Surface and Coatings Technology 1998. 108–109: p. 107–113.
88. K. Matsumoto, Y.Itoh, and T.Kameda, EB-PVD process and thermal properties of
hafnia-based thermal barrier coating. Science and Technology of Advanced Materials,
2003. 4: p. 153–158.
122
89. R. Vassen, X.Q.Cao, F. Tietz, D. Basu and D. Stover, Zirconates as new materials for
thermal barrier coatings. Journal of the American Ceramic Society, 2000. 83: p. 2023–
2028.
90. B. Saruhan, P.Francois, K. Kritscher and U. Schulz, EB-PVD processing of pyrochlore-
structured La2Zr2O7-based TBCs. Surface and Coatings Technology, 2004. 182: p. 175–
183.
91. J.Wu, Low thermal conductivity rare-earth zirconate for potential thermal-barrier coating
applications. Journal of the American Ceramic Society, 2002. 85: p. 3031–3035.
92. R. Gadow and M. Lischka, Lanthanum hexaaluminate – novel thermal barrier coatings
for gas turbine applications – materials and process development. Surface and Coatings
Technology, 2002. 151–152: p. 392–399.
93. C.G. Levi, Emerging materials and processes for thermal barrier systems. Current
Opinion in Solid State and Materials Science, 2004. 8 p. 77–91.
94. T. Beck, M.Schweda and L.Singheiser, Procedia Engineering, 2013 55 p. 191 – 198.
95. C. Zhu, P.Lu, A. Javed, G.Y. Liang, P. Xiao, Surface & Coatings Technology, 2012. 206:
p. 3739–3746.
96. K.H.Stern, Metallurgical and Ceramic Protective Coatings. 1996,Chapman & Hall, 26
Boundary Row London,UK.
97. A. Rauf, Q.Yu, L. Jin and C. Zhou, Scripta Materialia 2012. 66 p. 109–112.
98. A.C.Fox and T.W. Clyne, Surface and Coatings Technology, 2004. 184 p. 311-321.
99. H. H. Guo, S.Gong, Beijing University of Aeronatical and Ausronautics: China.
100. S. Mihm, T.Duda, H. Gruner, G. Thomas, and B. Dzur, Journal of Thermal Spray
Technology, 2012. 21: p. 400-408.
101. A. C. Léger, J.Wigren and M.O. Hansson, Surface and Coatings Technology 1998. 108-
109 p. 86-92.
102. W. J. Quadakkers, V.Shemet, D. Sebold, R. Anton, E. Wessel and L. Singheiser, Surface
and Coatings Technology, 2005. 199 p. 77-82.
103. H. M. Tawancy, N.M.Abbas and T.N. Rhys-Jones, Surface and Coatings Technology,
1991. 49(1-3) p. 1.
104. A.G. Evans, D.R.Mumm, J.W. Hutchinson, G.H. Meier, F.S. Pettit, Mechanisms
controlling the durability of thermal barrier coatings. Prog. Mater. Sci., 2001. 46: p. 505-
553.
105. S. Das, S.Datta, D. Basu and G. C. Das, Ceramics International, 2009. 35 p. 1413-1406.
106. D. Seo, K.Ogawa, Y. Nakao, H. Miura and T. Shoji ,Surface and Coatings Technology,
2009. 203 p. 1979-7983.
107. C. Zhou, J.Yu, S. Gong and H. Xu, materials Science and Engineering, 2003. A 348 p.
327-332.
123
108. W. D. Kingery, H.K.Bowen and D.R.Uhlmann, Introduction to Ceramics. Second Edition
ed. 1976: John Wiley & Sons.
109. Y. M. Chiang, D.P.Birnie and W.D.Kingrey, Physical Ceramics:Principles for Ceramic
Science and Engineering. 1997: John Wiley & Sons Inc.
110. M.W. Barsoum, Fundamentals of Ceramics. 1997: The McGraw-Hill Companies, Inc.
111. R. A. Pidani, R.S. Razavi., R. Mozafarinia, H. Jamali, Optics and Lasers in Engineering,
2012. 50: p. 780–786.
112. J. M. Drexler, C.H.Chen, A. D. Gledhill, K. Shinoda, S. Sampath, N. P. Padture, Surface
& Coatings Technology, 2012. 206: p. 3911–3916.
113. A. Guignard, G.Mauer, R. Vaßen, and D. Sto¨ver, Journal of Thermal Spray Technology,
2012. 21: p. 416-424.
114. P.S. Anderson, Surface & Coatings Technology, 2006. 185: p. 106-119.
115. A. Stuke, H.Kassner, J. L. Marque´s, R. Vassen, and D. Sto¨ver, International Journal of
Applied Ceramic Technology 2012: p. 1–14.
116. S. Guo, Y.Tanaka and Y.Kagawa, Journal of the European Ceramic Society, 2007. 27: p.
3425-3431.
117. J.R. Nicholls, Surface and Coatings Technology, 2002. 151-152 p. 383-391.
118. J.y. Kwon, Surface & Coatings Technology, 2006. 201: p. 3483-3490.
119. Y. Zhao, A.Shinmi, X. Zhao, P.J. Withers, S. Van Boxel, N. Markocsan, P. Nylen, P.
Xiao, Surface & Coatings Technology 2012. 206 p. 4922–4929.
120. G. Mauer, M.O.Jarligo, D. E. Mack, and R. Vaßen, Journal of Thermal Spray
Technology, 2013. 22: p. 646-658.
121. S. Wei, W.F.chi, F. Qun-bo, M. Zhuang, Surface & Coatings Technology, 2013. 217: p.
39–45.
122. J. Malzbender, R.W.Steinbrech, Surface & Coatings Technology, 2012. 209: p. 97–102.
123. R.B. Heimann, Plasma Spray Coating. 1996: VCH Verlagsgesellschaft mbH, Fedral
Republic of Germany.
124. D. Matejka, B.Benko, Plasma Spraying of Metallic and Ceramic Materials. 1989: John
Wiley & Sons Ltd.
125. W. Huang, X.Fan, Y. Zhao, X. Zhou, X. Meng, Y. Wang, B. Zou, X. Cao and Z. Wang, .
Surface & Coatings Technology 2012. 207 p. 421–429.
126. Y. Tan, V.Srinivasan, T. Nakamura, S. Sampath, P. Bertrand, and G. Bertrand, Journal of
Thermal Spray Technology, 2012. 21: p. 950-962.
127. M. G. Hocking, V.Vasantasree and P.S.Sidnye, Metallic and Ceramic Coating:
Production High Temperature properties and Applications. 1989: Longman Scientific and
Technical.
128. E. Lang, Coating for the High Temperature Applications. 1983: Applied Science
Publisher.
124
129. I.J. McColm, Ceramic Science for Materials Technologists. Leonard Hill.
130. A.Nusair.Khan, and J. Lu, Journal of Materials Processing Technology, 2009. 209 p.
2508-2524.
131. P. Fauchais, Journal of Thermal Spray Technology, 2001. 10: p. 44-66.
132. E. Lugscheider, C.Barimani, P. Eckert and U. Eritt, Computational materials Science,
1996. 71 p. 109-114.
133. I. V. Kostikov, A.I.Shestrin, Plazmennyie pokrytiia. 1978, Moscow.
134. J.H.Zaat, Ann. Rev. Mater. Sci., 1983. 13: p. 9-42.
135. W. Huang, Y.Zaho, X. Fan, X. Meng, Y. Wang, X. Cai, X. Cao, and Z. Wang, Journal of
Thermal Spray Technology, 2013. 22: p. 918-925.
136. J.F.Shackelford and R.H Doremus, Ceramic and Glass Materials Structure, Properties
and Processing. 2008: Springer Science+Business Media, LLC.
137. X. Zhang, M.Watanabe, S. Kuroda, Engineering Fracture Mechanics, 2013. 110: p. 314–
327.
138. B.Sudhangshu, High Temperature Coatings,. 2007: Burlington, MA, Elsevier Science &
Technology Books.
139. J.L Smialek, Advances in the oxidation resistance of high-temperature turbine materials.
Surf. Interface Anal., 2001. 31 (7) p. 582–592.
140. R.A.Rapp, Chemistry and electrochemistry of the hot corrosion of metals. Corrosion
1986. 42 (10): p. 568–577.
141. R.A Rapp, Hot corrosion of materials: a fl uxing mechanism. Corros. Sci., 2002. 44 (2):
p. 209–221.
142. F.S. Pettit, c.s.giggens, Hot corrosion, in: C.T. Sims. 1987, New York: Superalloys II,
Wiley.
143. G.W Goward, Low-temperature hot corrosion in gas turbines: A review of causes and
coatings therefor, J. Eng. Gas Turbine Power. Tran ASME, 1986. 108: p. 421–425.
144. R.A. Rapp, Y.S.Zhang, Hot corrosion of materials: fundamental studies. JOM. 46: p. 47-
55.
145. T.S. Sidhu, S.Prakash, R.D. Agrawal, Hot corrosion and performance of nickel based
coatings. Curr. Sci. , 2006. 90 (1): p. 41–47.
146. P.Hancock, Vanadic and chloride attack of superalloys. Mater. Sci. Technol, 1987. 3: p.
536–544.
147. N. Eliaz, G.Shemesh, r.m. latanision, Hot corrosion in gas turbine components. Eng. Fail.
Anal. , 2002. 9 (1): p. 31–43.
148. J.Stringer, Hot Corrosion in Gas Turbines. 1972, Metals and Ceramics Information
Center, MCIC Report New York. p. 72–08.
149. K.L. Luthra, D.A. Shores, Fundamentals of High-Temperature Corrosion. 1992,
Academic Press: New York.
125
150. N.S. Bornstein, M.A. Decrescente,, Formation and reactivity thermodynamics of sodium
sulfate with gas turbine alloys. Corrosion 1968, 24: p. 127–133.
151. J.A. Goebel, F.S.Pettit., G.W. Goward, Mechanisms for the hot corrosion of nickel-base
alloys, Metall. Trans. , 1973. 4 p. 261-278.
152. R.L. Jones, The development of hot-corrosion-resistant zirconia thermal barrier coatings.
Mater. at High Temp. , 1991. 9: p. 228–236.
153. H.E. Eaton, N.S.Bronstein., J.T. Demasi-marcin, The effects of environmental
contaminants on industrial gas turbine thermal barrier coatings. Proc. Int. Gas Turbine
Aeroengine Congress and Exhibition, ASME, New York, 1996: p. Paper 96-GT-283.
154. M.F. Trubelja, D.M.Nissley, N.S. Bornstein, J.T.D. Marcin, Pratt & Whitney thermal
barrier coating development, Proc. Advan Turbine Systems Annu Program Rev. Meet,
U.S. Department of Energy, Washington, DC (1997).
155. C. Leyens, U.Schulz, B.A. pint, I.G. Wright, Infl uence of electron beam physical vapor
deposited thermal barrier coating microstructure on thermal barrier coating system
performance under cyclic oxidation conditions. Surf. Coat.Technol. , 1999. 120–121: p.
68–76.
156. F.C. Toriz, A.B.Thakkar, S.K. Gupta, Thermal barrier coating for jet engines,
Proceedings of the International Gas Turbine Aeroengine Congress and A. , Paper 88-
GT-279 (1988).
157. M.P. Borom, C.A.Johnson, L.A. Peluso, Role of environment deposits and operating
surface temperature in spallation of air plasma sprayed thermal barrier coatings. Surf.
Coat. Technol., 1996. 86–87(116–126).
158. R.A. Rapp, K.S.Goto, Hot corrosion of metals by molten salts. 1981: Molten Salts, I,
Pennington, NJ, Electrochemical Society.
159. A.S. Khanna, S.K.Jha, Degradation of materials under hot corrosion conditions. Trans.
Indian Inst. Met., 1988. 51 (5) p. 279–290.
160. N.S. Bornstein, M.A.Decrescente, H.A. Roth, The relationship between relative oxide ion
content of Na2SO4, the presence of liquid metal oxides and sulfi dation attack. Metall.
Trans., 1973. 4: p. 1799–1810.
161. Y.S. Hwang, R.A.Rapp, Thermochemistry and solubilities of oxides in sodium sulfate-
vanadate solutions. Corrosion 1989. 45 (11): p. 933–937.
162. C.G. Zhou, Y.X.Song, C.L. wang, H.B. Xu, Cyclic-oxidation behavior of thermal-barrier
coatings exposed to NaCl vapor. Oxid. Met. , 2008. 69 (1-2).
163. X. Chen, Y.Zhao., L. Gu, B. Zou, B. Zou, and X. Cao, Hot Corrosion Behaviour of
Plasma Sprayed YSZ/LaMgAl11O19 Composite Coatings in Molten Sulfate-Vanadate
Salt. Corros.Sci., 2011. 53: p. 2335-2343.
164. A.Nusair Khan, J.Lu, Thermal cyclic behavior of air plasma sprayed thermal barrier
coatings sprayed on stainless steel substrates. Surface and Coatings Technology, 2007.
201: p. 4653-4658.
126
165. C.S. Ramachandran , V.B.alasubramanian, P.V. Ananthapadmanabhan, On the cyclic hot
corrosion behaviour of atmospheric plasma sprayed Lanthanum Zirconate based coatings
in contact with a mixture of sodium sulphate and vanadate salts: A comparison with the
traditional Ysz duplex and NiCrAlY coated samples. Vacuum, 2013. 97: p. 81-95.
166. Saremi Mohson, Afrasiabi Abbas, Kobayashi Akira, Bondcoat oxidation and hot
corrosion behavior of plasma sprayed YSZ coating on Ni superalloy. Transactions of
Joining and Welding Research Institute, 2007. 36(1): p. 41-45.
167. N.Toshio, I.Takeshi, N.Takumi, S.Yoshimitsu, T.K.Zaini, H.Shigenari, Advanced
Coatings on High Temperature Applications. Materials Science Forum, 2006. 522-523: p.
1-14.
168. C. H Lee, H.K.Kim, H. S. Choi, H. S. Ahn, Phase transformation and bondcoat oxidation
behavior of plasma-sprayed zirconia thermal barrier coating. Surface and Coatings
Technology, 2000. 124: p. 1.
169. Diffusion Processes in Advanced Technology Materials, E.D.G., William Andrew, Inc.
and Springer-Verlag GmbH & Co. KG, USA, (2005)
170. A.Keyvani, M.Sarimi, M. Heydarzadeh Sohi, An investigation on oxidation, hot
corrosion and mechanical properties of plasma-sprayed conventional and nanostructured
YSZ coatings. Surface and Coatings Technology, 2011. 206: p. 208-216.
171. M.H. Habibi, Li Wang., S.M.Guo, Evolution of hot corrosion resistance of YSZ,
Gd2Zr2O7, and Gd2Zr2O7 + YSZ composite thermal barrier coatings in Na2SO4 +
V2O5 at 1050 ◦C. Journal of European Ceramic Society, 2012. 32: p. 1635-1642.
172. Powder diffraction files, J.-I.C.f.D.D., USA (1999).
173. J. W. Edington, Electron Diffraction in the Electron Microscopy, Monographs in
Practical Electron Microscopy in Materials Science, The Macmillan press Ltd, London,
(1975) page-80.
174. Smithels Metals Reference Book, Eds. E.A.Brands, G.B.Brook, 7th Edition, Butterworth
Heinemann, USA, 1999, p 27-4.
175. L.A.Bursill, B.G.Hyde, Crystallographic shear in the higher titanium oxides: structure,
texture, mechanisms and thermodynamics. Prog. Solid State Chemistry, 1972. 7: p. 177.
176. G.J.Wood, L.A.Bursill,. The Formation Energy of Crystallographic Shear Planes in
TinO{2n - 1}. Proc. R. Soc. Lond., 1981. A375: p. 105.
177. E. Cho, S.Han, H. S.Ahn, K. R. Lee, S. K. Kim, and C. S. Hwang, First-principles study
of point defects in rutile TiO2−x. Phys. Rev. B, 2006. 73: p. 193-202.
178. Ed. J.P.Carmo and J.E.Ribeiro, New Advances in Vehiculaar Technology and
Automotive Engineering, Croatia: In Tech, (2012), p.133.
179. S. Somiya, S.Hirano, S. Kamiya, Phase relations of the Cr2O3-TiO2 system. J. Solid
State Chem., 1978. 25: p. 273-284.
180. C. Winde, Deposition and Characterization of Cr-film on TiO2 (110) surfaces.,
Doctorate dissertation, Max-Planck Institut für Metallforschung, Stuttgart, Germany,
2002.
127
181. M. Harju, E.Levanen, T. Mantyla Wetting behaviour of plasma sprayed oxide coatings.
Appl. Surf. Sci., 2006. 252: p. 8514-8520.
182. ASTM E2860 (Standard Test Method for Residual Stress Measurement by X-Ray
Diffraction)
183. ASTM E837 (Standard Test Method for Determining Residual Stresses by the Hole-
Drilling Strain-Gage Method)
184. C. A. Walker and J. McKelvie, The Measurement Of Residual Stress By A Moire Fringe
Method, SPIE 0136, 1st European Conf. on Optics Applied to Metrology, 302 (April 18,
1978); doi:10.1117/12.956177.
185. André Mézin, Audrey Hemel, Measuring coating internal stresses by the curvature
method applied to a beveled sample. Surface and coating technology, 2014. 258: p. 1128-
1136.
186. T.W. Clyne, Residual stresses in the surface coatings and their effect on interfacial
debonding. Key Eng. Mater., 1996. 116-117: p. 307-330.
187. Thermal Barrier Coatings, Ed. Hubin Xu, Hongbo Guo, Woodhead publishing limited,
2011.
188. Imran Nazir Qureshi, M.Shahid, A. Nusair Khan, and Yaseer A. Durrani, Journal of
Thermal Spray Technology,, 2015. 24(8) p. 1520-1528
189. I. Gurrappa , A.Sambasiva Rao, Surface and coating technology, 2006. 201: p. 3016-
3029.
190. Z. Yu, K.P.Dharmasena, D.D Hass, H.N.G. Wadley, Surface and Coating Technology,
2006. 201 p. 2326-2334.
191. K.G.S. Thomas, M.Hertter, Surf. Coat. Technol., 1999. 120-121: p. 84-88.
192. W.J. Brindley, J. Thermal Spray Technol. , 1997. 6 p. 85.
193. R.A. Miller, J. Am. Ceram. Soc., 1984. 67 p. 517.
194. L. Lelait, S.Alpirine, R. Mevre, J. Mater. Sci. Lett., 1992. 27 p. 5.
195. E.Y. Lee, R.R.Biederman., R.D. Sisson, Mater. Sci. Eng. A, 1989. 121: p. 467.
196. D.R. Mumm, A.G.Evans, I.T. Spitsberg, Acta Mater. , 2001. 49 p. 2329-2340.
197. Y.H. Sohn, J.H.Kim, E.H. Jordan, M. Gell, Surf. Coat. Technol. , 2001. 146-147: p. 70-
78.
198. E. Tzimas, H. Mu¨ llejans, S.D. Peteves, J. Bressers, W. Stamm, Acta Mater., 2000. 48:
p. 4699-4707.
199. D.R. Mumm, G.A.Evans, Key Eng. Mater., 2001. 197: p. 199-250.
200. A. Rabiet, A.G.Evans, Acta Mater., 2000. 48 p. 3963.
![Corrosion Protection.ppt [Read-Only]swan/courses/53086/Corrosion_Protection.pdfMethods of Corrosion Control Barrier Protection Provided by a protective coating that acts as a barrier](https://static.fdocuments.in/doc/165x107/610108f5c9335c35fd4589d5/corrosion-read-only-swancourses53086corrosionprotectionpdf-methods-of-corrosion.jpg)
