thermal barrier coatings_chapter2

8/7/2019 Thermal Barrier Coatings_chapter2

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Chapter 2

Thermal Barrier Coatings

2.1 Overview

By attaching an adherent layer of a low thermal conductivity material to the surface of a

internally cooled gas turbine blade, a temperature drop can be induced across the thickness

of the layer, Fig. 2.1. This results in a reduction in the metal temperature of the component

to which it is applied. Using this approach temperature drops of up to 170oC at the metal

surface have been estimated for 150 m thick yttria stabilized zirconia coatings[19]. This

temperature drop reduces the (thermally activated) oxidation rate of the bond coat applied

to metal components, and so delays failure by oxidation. It also retards the onset of ther-

mally induced failure mechanisms (i.e. thermal fatigue) that contribute to component

durability and life. It is important to note that coatings of this type are currently used only

for component life extension at current operating temperatures. They are not used to

increase the operating temperature of the engine. However, the development of a prime

reliant TBC system, for which the probability of failure is sufficiently low, would allow

these coatings to be used to increase the engine operating temperature and lead to signifi-

cant improvements in engine performance.


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Chapter 2. Thermal Barrier Coatings 7

Figure 2.1 A schematic illustration of a modern thermal barrier coating system consisting

of a thermally insulating thermal barrier coating, a thermally grown oxide

(TGO) and an aluminum rich bond coat. The temperature gradient during

engine operation is overlaid.

Turbineblade

Hotgases

Interiorcoolant

gas

Interiorcoolant

gas

Thermally grownoxide

Insulativethermal

barrier

coatingOxidation

resistant

bond coat

Nickelsuperalloy

turbine blade

Temperatu

re

150mm

Cooling

gas

temperature

170oC

Enginegas

temperature


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2.2 TBC Systems

Modern TBCs are required to not only limit heat transfer through the coating but to also

protect engine components from oxidation and hot corrosion. No single coating

composition appears able to satisfy these multifunctional requirements. As a result, a

coating system has evolved. Research in the last 20 years has led to a preferred coating

system consisting of three separate layers[20] to achieve long term effectiveness in the

high temperature, oxidative and corrosive use environment for which they are intended to

function, Fig. 2.1.

First, a thermally protective TBC layer with a low thermal conductivity is required to

maximize the thermal drop across the thickness of the coating. This coating is likely to

have a thermal expansion coefficient that differs from the component to which it is

applied. This layer should therefore have a high in-plane compliance to accommodate the

thermal expansion mismatch between the TBC and the underlying nickel superalloy

component. In addition, it must be able to retain this property and its low thermal

conductivity during prolonged environmental exposure. A porous, columnar, 100-200 m

thick, yttria stabilized zirconia (YSZ) layer is currently preferred for this function[21].This layer may be applied using either air plasma spray (APS)[15] or electron beam

physical vapor deposition (EB-PVD)[16].

Second, an oxidation and hot corrosion resistant layer is required to protect the

underlaying turbine blade from environmental degradation. This layer is required to

remain relatively stress free and stable during long term exposure and remain adherent to

the substrate to avoid premature failure of the TBC system. It is important that it also

provide an adherent surface for the TBC top coat. Normally, the thin (< 1 m), protectivealuminum rich oxide which is thermally grown upon the bondcoat is utilized for this

purpose[22]. Since the aluminum content of modern nickel based superalloy is not

typically high enough to form a fully protective alumina scale, an aluminum rich layer


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(bond coat) is applied onto which the thermally grown oxide may form[23]. A ~50m

thick layer of either a low sulphur platinum aluminide[24] or MCrAlY (where M is Ni or

Co)[25] is utilized for this purpose. Either low pressure plasma spray (LPPS)[22] or pack

cementation[26] are used to apply the bond coat.

In addition, these layers are desired to be thin and low density to limit the centrifugal load

on rotating engine components and have good thermal and mechanical compatibility. The

focus of the work in this dissertation is on the insulative YSZ top layer where improved

coating morphologies are desired to improve TBC performance.

2.3 Materials Selection

Yttria stabilized zirconia has become the preferred TBC layer material for gas turbine

engine applications because of its low thermal conductivity, , and its relatively high

(compared to many other ceramics) thermal expansion coefficient, Fig. 2.2[27]. This

reduces the thermal expansion mismatch with the high thermal expansion coefficient met-

als to which it is applied. It also has good erosion resistance which is important because of

the entrainment of high velocity particles in the engine gases[1]. The low thermal conduc-tivity of bulk YSZ results from the low intrinsic thermal conductivity of zirconia (reported

to be between 2.5 and 4.0 depending on the phase, porosity and temperature[28]) and

phonon scattering defects introduced by the addition of yttria[29]. These defects are intro-

duced because yttria additions require the creation of O2- vacancies to maintain the electri-

cal neutrality of the ionic lattice. Since both the yttrium solutes and the O2- vacancies are

effective phonon scattering sites the thermal conductivity is decreased as the yttria content

is increased. In practice, a yttria concentration in the range of 6 to 8 wt.% is generally used

since this composition maximizes spallation life due to the formation of the metastable t

phase[30], Fig. 2.3. This phase yields a complex microstructure (containing twins and

antiphase boundaries) which resist crack propagation and transformation into the mono-

clinic phase (with an attendant 4% volume change) upon cooling. YSZ has a room temper-


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ature, grain size dependent, thermal conductivity of 2.2-2.6 W/mK in the densest form.

Adding porosity further reduces and can improve the in-plane compliance[31].

Figure 2.2 Plot of thermal conductivity vs. thermal expansion coefficient. Materials forthe TBC layer are desired to have a thermal expansion coefficient close to that

of nickel based superalloys and the lowest possible value. The nickel alloys

used for turbine blades have a CTE from 14.0 to 16.0 x 10-6 K-1. YSZ has a

CTE = 9.0 x 10-6 K-1.

nickel alloys

Al2O3MgO

SiO2

mullites sialons

BeO

ZrO2

Thermal Conductivity (W/mK)

ThermalExpansionCoefficient(10-6/K)

1.0 10.0 100.00.1

1.0

10.0

100.0

1000.0

YSZ


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Figure 2.3 Yttria - Zirconia phase diagram. Note that the shaded region indicates the

region where the formation of the metastable t phase occurs upon cooling.

Compositions which result in the monoclinic + cubic phase at room

temperature (i.e. 3 to 18 mol% YO1.5) are termed partially stabilized zirconia

and compositions which result in solely the cubic phase at room temperature

(> 18 mol% YO1.5) are termed fully stabilized zirconia.

T + F

T

cubic (F)

liquid

L+F

M M + F

T'

3000

2500

2000

1500

1000

500

Tem

perature(

oC)

YO1.5

(mol%)

20151050

Y2O3(wt%)

6 8 10 12

M = monoclinicT = tetragonal


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2.4 Property Profiles

TBC coating systems must possess a combination of properties to be effective. These

include a low thermal conductivity, high resistance to spallation, good erosion resistance,

phase stability and pore morphological stability. For aircraft turbine applications the spal-

lation resistance and the thermal conductivity of the coating system are the most critical to

performance[32]. The thermal conductivity is strongly dependent on the volume fraction

and morphology of the porosity found in this layer. The spallation resistance, however, is

dependent on the mechanical properties of all three layers [33,34]. For example the TBC

top layer must have a high in-plane compliance to minimize the coefficient of thermal

expansion (CTE) mismatch stress[35] between the top TBC layer and the underlying

superalloy substrate.

Even when highly compliant TBC top layers are deposited, spallation failure can still

occur. Such failures have been observed to initiate either within the TBC layer, at the

TBC/TGO interface or at the TGO/bond coat interface[36]. One contributing factor is the

development of large stresses in the TGO layer. Clarke and Christensen[37,38] have mea-

sured ambient temperature residual compressive stresses of 3 to 4 GPa in the TGO layer of

TBC systems. This stress has been linked to the CTE mismatch between the TGO layer

and the substrate/bond coat and to growth stresses in the TGO. Evans et al.[39,40] have

analyzed the thermomechanical stresses in these systems and shown that they can lead to

the initiation of cracks at the TGO/bond coat interface. Out-of-plane tensile stresses result-

ing from undulations or morphological defects that form on an otherwise smooth sur-

face[39], ratcheting effects caused by cyclic plasticity in the substrate[40], TGO/bond coat

interface embrittlement (due to sulphur impurities)[40,41] and sintering induced increases

in the TBC in-plane compliance are all thought to play a role in the spallation failure of

TBC systems[42,43]. Recent work also suggests that the TGO undulations can result in


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the formation of cracks the TBC layer[44]. Control of these thermally induced failure

mechanisms is clearly a critical issue for the development of more durable TBC systems.

Increasing the thermal resistance of the TBC layer is expected to reduce the growth rate of

the TGO layer and slow the rate of ratcheting by reducing the temperature below it.

However, the performance of a TBC coatings system depends on the top layer morphol-

ogy as well. A manifest of this is seen when TBC top coats are deposited using APS or

EB-PVD. Coatings produced by APS have a thermal conductivity in the range of 0.8 - 1.0

W/mK at 25oC[14,45]. This is significantly lower than the 1.5 - 1.9 W/mK reported for

EB-PVD coatings at 25oC[14,46] and as a result, the APS coatings provides superior ther-

mal protection. However, the spallation resistance of these layers is less than that of EB-

PVD TBC layers (8 to 10 times shorter spallation lifetimes)[47]. This arises because of the

superior in-plane compliance of the EB-PVD coating. As a result, EB-PVD TBC layers

are preferred for aerospace gas turbine applications[22].

2.5 TBC Pore Morphology

The thermal and mechanical property differences of YSZ coatings synthesized by the two

processing routes result from differences in the morphology of the porosity present within

the TBC layer, Fig. 2.4. In APS layers, inter-splat pores result from the impingement of

molten droplets onto a substrate. These pores are roughly aligned parallel to the substrate

surface and are accompanied by micro-cracks and fine grain boundaries, Fig. 2.4(a).

Brindley[48] has shown, Fig. 2.5(a), that the thermal conductivity of YSZ coatings will

decrease as the pore volume fraction increases. In this case, the pores provide a high

impedance to heat flow through the thickness of the coating resulting in a TBC layer with

low thermal conductivity. The spallation life of these coatings is believed to be governed

by a combination of the disc-like coating defects and the significant coefficient of ther-

mal expansion (CTE) mismatch between the TBC layer and the underlying substrate.


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Since the CTE of YSZ is approximately 5x10-6oC-1 smaller than that of the nickel super-

alloys to which it is applied, significant strains are formed during thermal cycling which

leads to the initiation of cracks in the TBC layer. These cracks eventually result in spalla-

tion of the coating[50]. Such failures limit the use of these coatings to applications where

only moderate thermal cycling is experienced (e.g. land base power generation tur-

bines)[51].

In contrast, the TBC layers produced by EB-PVD have a columnar microstructure with

elongated intercolumnar pores that become predominantly aligned perpendicular to the

Figure 2.4 Schematic illustrations of the pore morphology of (a) a plasma spray

deposited YSZ showing its coarse, disc-like pores aligned parallel to the

substrate surface and (b) an electron beam physical vapor deposited YSZcoating with elongated pores aligned perpendicular to the substrate surface.

TBC

A) Plasma Spray

B) EB-PVD

Bond coat

Bond coat

TGO

TGO

Thermal flux

Low lateral compliance

Highthermalconductivity

(k= 1.5-1.9 W/mK)

Lowthermal

conductivity(k= 0.8-1.0 W/mK)

High lateral compliance

TBC


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plane of the coating as its thickness increases[52]. A finer distribution of intracolumnar

pores also exists. The elongated intercolumnar pores increase the compliance of the coat-

ing in the plane of the substrate, and leads to the improved spallation lifetimes of the TBC

systems.

For the deposition temperatures typically employed to obtain good coating adhesion (i.e.

T/Tm ~ 0.47 where T is the substrate temperature and Tm is the melting point of the depos-

ited material), substrate rotation is required during EB-PVD to obtain sufficient inter- and

intracolumnar porosity. This rotation causes flux shadowing and a varying deposition rate.

The resulting competitive growth process leads to tapered, poorly bonded columns aligned

perpendicular to the substrate surface and the formation of the finer intracolumnar pores,

Fig.2.4(b). The large, through thickness intercolumnar pores are not effective at reducing

heat transfer through the thickness of the coating. However, the fine intracolumnar pores

contribute a moderate reduction in the thermal conductivity as they are generally inclined

to the heat flow. Even so, EB-PVD coatings still have a considerably higher thermal con-

ductivity than their APS counterparts[14]. These results indicate that the morphology of

the porosity in the YSZ coating will strongly effect its thermal conductivity, see Fig.

2.5(b)[49].

One approach to improve TBC system performance is to optimize the pore morphologies

in order to reduce the TBCs thermal conductivity while still retaining a high in-plane

compliance. Lower thermal conductivities lead to temperature reductions at the

TGO/bond coat interface which slows the rate of the thermally induced failure mecha-

nisms. For example, lower temperatures in the bond coat and TGO layers reduce the CTE

mismatch strain in the TGO layer, slow the growth rate of the TGO layer and retard impu-

rity diffusion within the bond coat. Alternatively, lower thermal conductivity TBC layers

might allow designers to reduce the TBC thickness thereby decreasing the significant cen-

trifugal load that the mass of the TBCs imposes on the rotating turbine engine compo-

nents[11].


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Figure 2.5 The effect of pore volume fraction and morphology on the thermal

conductivity of a given material. In a) the experimentally determined thermal

conductivity reduction is shown for plasma sprayed coatings having a

increasing amount of porosity. In b) the calculated reduction in the thermal

conductivity is shown for a material in which elongated cracks with different

orientations to the heat flux have been introduced.

0 . 0

1 . 0

2 . 0

3 . 0

T

h

e

r

m

a

l

C

o

n

d

u

c

t

i

v

i

t

y

(

W

/

m

-

K

)

1 0 . 0 2 0 . 0

3 0 . 0

4 0 . 0 5 0 . 0

P o r e V o l u m e F r a c t i o n ( % )

p l a m s a s p r a y e d

Y S Z c o a t i n g s

Bond coat

TGOTBC

A)

3.0

2.0

1.0

0.0

1.0 5.04.03.02.0

pores alignedperpendicular to

heat flow (k ~ k )o

q = 32.7o

q = 45.0o

q = 60.0o

Pore Volume Fraction (%)

ThermalConductivity

(W/mK)

Thermal flux

pores aligned

parallel toheat flux

Thermal flux

B)


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2.6 Advanced Processing Approaches for TBCs

In order to develop tailored pore morphologies, a processing approach must be identi-

fied or developed in which pore evolution can be manipulated and controlled. Several

alternative methods have been proposed to deposit YSZ coatings. These include sputtering

[53] chemical vapor deposition (CVD)[54], and sol-gel based approaches[55]. In sputter-

ing, energetic particles are employed to remove (sputter) atoms from a target. The sput-

tered atoms then deposit on a substrate to grow a film. In CVD, a volatile compound of the

material to be deposited is chemically reacted with other gases producing a nonvolatile

solid that deposits atomistically on a substrate. Sol-gel based coating approaches incorpo-

rate solvent based chemical solutions to transport coating materials onto a substrate.

All of these approaches, however, have significant drawbacks for TBC application. Sput-

tering results in low deposition rates making it undesirable for the economical production

of porous YSZ layers. CVD and sol-gel approaches also deposit material at a low rate and

typically require the use of dangerous (and expensive) precursor materials. In addition,

none of these approaches allows any means for precisely controlling and manipulating the

pore morphology within the coating. As a result, processing approaches which are supe-

rior to EB-PVD for the high rate production of controllable, porous coatings do not appear

to currently exist. Thus, improvements to electron beam based deposition approaches

which facilitate the economical deposition of controllable, porous coating morphologies

appear to be the most suitable option.

The morphology of vapor deposited coatings is generally controlled by multiple process-

ing variables including: adatom kinetic energy[56,57], adatom angle of incidence[58,59],

substrate temperature[60,61], deposition rate[60,62], the presence and nature of the sur-

rounding gas[60,63,64], elemental compositions of the adatoms[65], substrate rough-


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ness[66] and the relative amount of chamber ionization/plasma generation[67]. It has been

experimentally shown by several researchers that porous, columnar morphologies are typ-

ically associated with low adatom kinetic energy, large angles of incidence, low substrate

temperatures, high deposition rates, rough substrates and high chamber pres-

sures[56,60,66,68,69].

These parameters can facilitate coating porosity in many ways including: reductions in the

adatom surface mobility[70], flux shadowing[71] and vapor phase cluster deposition[72].

For example, the effect of substrate temperature and chamber pressure can be observed

from empirical structure zone models (SZM), Fig. 2.6, in which the results of an experi-

mental study of sputtered materials revealed systematic changes in the coating morphol-

ogy as these coating parameters were altered[65,73]. Note that the porous structures are

indicated in the high chamber pressure / low substrate temperature region. Flux shadowing

effects can result in porosity when oblique atom arrivals are shadowed by surface asper-

ities creating local variations of the vapor flux, Fig. 2.7. This leads to a reduced growth

rate in flux depleted regions of the substrate. As growth progresses, these regions lead to

pore formation in the coating.

Unfortunately, significant control over such variables can often be limited by other engi-

neering aspects of coating design. For example, TBC top layers deposited by EB-PVD are

deposited on relatively rough substrates at high substrate temperatures (T/Tm = 0.47) in

order to achieve proper coating adhesion, at low chamber pressures to obtain proper work-

ing condition for EB-gun operation and moderate deposition rates due to the poor materi-

als utilization efficiency of the EB-PVD process and the desire to minimize coating

defects resulting from enhanced evaporation rates. The adatom energy is generally rela-

tively low in these systems (~0.2 eV) and not easily variable. Despite this low energy, the

processing conditions in conventional EB-PVD processes do not result in a YSZ layer


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with the desired porous, columnar morphology unless the component is rotated around

one axis[74]. This rotation alters the angle of incidence of the adatom during deposition

(from -90 to 90o) resulting in significant shadowing effects and a varying deposition rate.

A columnar morphology, as described, results with the column morphology strongly

effected by the rotation speed.

The engineering constraints on depositing coatings of this type using EB-PVD therefore

can limit the means available to introduce porosity in the coating. However, an electron

beam processing approach which allowed for coatings to be deposited at higher chamber

pressures, higher deposition rate or had an additional means of manipulating the angle of

incidence distribution at the substrate would appear to be useful for this application.

Figure 2.6 Structure zone model (SZM) showing the results of an experimental study of

sputtered materials (e.g. Ti, Fe, Cr, Cu, Mo and Al) showing how substrate

temperature and chamber pressure similarly effect the morphology of a range

of coating materials.


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2.7 Summary

Investigation of the current state-of-the-art in TBC technology has indicated that opportu-

nities exist to significantly improve upon modern TBC systems. One such opportunity is

the tailoring of pore morphologies to improve the insulative properties of the coatings. In

doing so an increase in the spallation life of the coating may also be expected provided the

in-plane compliance is not compromised. To accomplish this, however, it appears that

improved TBC deposition approaches must be developed which exhibit improved control

over the coating morphology.

Figure 2.7 Schematic illustration showing how an oblique atom arrival will have its

substrate impingement point altered by the presence of surface asperities.

Note that the peak of the asperity will have an enhanced vapor flux as a result.

shadowinginduced

depositionlocation

unshadowedlocation

vaporatom

surfaceasperity

substrate

enhancedflux

region

depletedflux

region

Flux Shadowing

thermal barrier coatings_chapter2

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