NASA/TM--1998-208406

Sintering and CreepZirconia and Hafnia Based

Barrier Coatings

Behavior of Plasma-SprayedThermal

Dongming Zhu

Ohio Aerospace Institute, Cleveland, Ohio

Robert A. Miller

Lewis Research Center, Cleveland, Ohio

Prepared for the

25th International Conference on Metallurgical Coatings and Thin Films

sponsored by the American Vacuum Society

San Diego, California, April 27--May 1, 1998

National Aeronautics and

Space Administration

Lewis Research Center

July 1998

https://ntrs.nasa.gov/search.jsp?R=19980211578 2020-03-18T11:16:10+00:00Z

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Acknowledgments

The authors are grateful to George W. Leissler for his assistance in the preparation of plasma-sprayed ceramic

coating specimens, and to Ralph G. Garlick for performing X-ray diffraction experiments.

NASA Center for Aerospace Information7121 Standard Drive

Hanover, MD 21076Price Code: A03

Available from

National Technical Information Service

5287 Port Royal RoadSpringfield, VA 22100

Price Code: A03

SINTERING AND CREEP BEHAVIOR OF PLASMA-SPRAYED

ZIRCONIA AND HAFNIA BASED THERMAL BARRIER COATINGS

Dongming Zhu _"and Robert A. Miller

"_Ohio Aerospace Institute

National Aeronautics and Space Administration

Lewis Research Center, Cleveland, OH 44135

ABSTRACT

i

The sintering and creep of plasma-sprayed ceramic thermal barrier coatings under high

temperature conditions are complex phenomena. Changes in thermomechanical and

thermophysical properties and in the stress response of these coating systems as a result of the

sintering and creep processes are detrimental to coating thermal fatigue resistance and performance.

In this paper, the sintering characteristics of ZrO2-8wt%Y203, ZrO2-25wt%CeO2-2.5wt%Y203,

ZrO2-6w%NiO-9wt%Y203, ZrO2-6wt%Sc20322wt%Y203 and HfO2-27wt%Y203 coating

materials were investigated using dilatometry. It was found that the HfO2-Y203 and baseline

ZrO2-Y203 exhibited the best sintering resistance, while the NiO-doped ZrO2-Y203 showed the

highest shrinkage strain rates during the tests. Higher shrinkage strain rates of the coating, materials

were also observed when the :specimens were tested in Ar+5%H2 as compared to in air. This

phenomenon was attributed to an enhanced metal cation interstitial diffusion mechanism under the

reducing conditions. It is proposed that increased chemical stability of coating materials will

improve the material sintering resistance.

I. INTRODUCTION

Plasma-sprayed ceramic thermal barrier coatings are being developed for advanced gas

turbine and diesel engine applications to improve engine reliability and efficiency. Since these

coatings are experiencing severe thermomechanical cycling during engine operation, it is especially

challenging to develop coating systems with high reliability and durability. In particular, ceramic

Paper presented in The 25th International Conference on Metallurgical Coatings and Thin Films, San Diego,California, April 27-May 1, 1998. To be published in Surface and Coatings Technology.NASA/TM--1998-208406 1

coating sintering and creep at high temperature are among the most important issues for the

development of advanced thermal barrier coatings, as has been recognized by many investigators

[1-10] . The ceramic sintering and creep at high temperature can result in coating shrinkage and

through-thickness cracking during cooling, thereby further accelerating the coating failure process.

Sintering-segmentation-enhanced delamination can be an important failure mechanism for a

thermal barrier coating system, due to stress concentration from the through-thickness cracks, and

increased coating elastic modulus from the sintering densification process. The increase in coating

thermal conductivity is also detrimental to coating performance. Research efforts involving

various techniques have also been made in characterizing the ceramic coating sintering and creep

behavior at high temperature and under temperature gradients simulating those encountered in the

engine [1, 2, 7, 11-13]

The sintering and creep of plasma-sprayed, porous and microcracked ceramic thermal

barrier coatings are complex phenomena. The early work by Firestone et al [1, 2] indicated that the

ceramic creep appeared to be a thermally activated process, with the ceramic splat-sliding being an

important creep deformation mechanism. More recently, it has been reported that the ceramic

thermal barrier coatings can sinter and creep significantly under compressive stress states at

relatively low temperatures [11, 13, 14] . The "creep" of plasma-sprayed ZrO2-8wt%Y203 at room

temperature has also been observed at a tensile stress of 7.4 MPa [15]. A mechanism-based model

has been proposed to describe the densification and deformation occurring in thermal barrier

coatings at temperature by taking into account the thermally and stress activated diffusion, and the

mechanical compacting processes [13] The dopants in the ceramic coatings can significantly

modify the point defect and microstructures in the bulk, at splat-grain boundaries and microcrack

surfaces of the materials, thereby can significantly affect these sintering and creep processes. A

better understanding of the dopant effects will help to develop future advanced, sintering/creep

resistant "superalloy-type" ceramic coatings.

The purpose of this paper is to investigate sintering kinetics of several zirconia and hafnia

based ceramic coating materials. The ceramic materials investigated include: (a) ZrO2-8wt%Y203,

a NASA-Lewis Research Center reference (or baseline) material; (b) ZrO2-25wt%CeO2-

2.5wt%Y203, a commercially available coating material developed for hot corrosion resistance;

(c) HfO2-27wt%Y203, a potential new coating material developed at NASA for high temperature

stability [16] ., (d) ZrO2-6wt%NiO-9wt%Y203, a NiO-doped ZrO2-Y203 coating material

reported to suppress the tetragonal-monoclinic phase transformation [17] ; (e) ZrO2-6wt%Sc203-

NASA/TM 1998-208406 2

................................• .........: • ,: _,_:_•_:,:_:_::__:__::_:_:_ __,_:_::_< _•__:: ___:_:_•_ ::_::_i_i?::_:::_:i_:_!_?:_z_i_ _:! ::::__?_::_!_!i_!_i/!!3!:_!_3i:ii!i!if!!!iii:i!ii̧i!!!i_ii!i!i!!!!_i:_iii_!_!!i_!iii:i!_}_iii_ii_iii_i_ii_i_iiiii!iii_i_i_iiiTi!i_iii!_iiiiiiiiiiii!iiiiiii!iiiTi_iiiiii_ii_iiii_

2wt%Y203, developed for improved hot corrosion resistance [18] . The coating sintering and creep

mechanisms and dopant effect on coating sintering rates are discussed based on experimental

observations and possible defect reactions.

II. EXPERIMENTAL MATERIALS AND METHODS

The five ceramic coating materials mentioned above, ZrO2-8wt%Y203, ZrO2-

25wt%CeO2-2.5wt%Y203, ZrO2-6wt%NiO-9wt%Y203, ZrO2-6wt%Sc203-2wt%Y203 and

HfO2-27wt%Y203, were chosen for this study. The actual compositions of these materials were

close to the their nominal compositions. Each of the above materials was prepared by sintering and

crushing except for ZrO2-25CeO2-2.5wt%Y203 which was spray dried and plasma spheroidized.

A single set of standard plasma-spray parameters was Used for each material. The powders with an

average particle size of 60 ]xm of these coating materials were first plasma-sprayed onto 3 mm

diameter graphite cylindrical bars, using the plasma spray conditions described previously [19]. The

coating thickness was about 0.76 mm, and porosity was about 10%. The graphite bars were then

slowly burnt off at 600°C for 6 hours in a furnace in air. The hollow ceramic cylinders were cut

into 25.4 mm dilatometer specimens.

Ceramic sintering experiments were carried out in air and in Ar+5%H2 within the

temperature range of 900 to 1400°C, using a UNITHERM TM high temperature dilatometer system

shown in Figure 1. Since the push rod in the dilatometer exerts a certain amount of force

(measured at approximately 450g using a spring device) on the specimen, a uniaxial stress of

eramac coating material

Pt reference Thermocouple

_V =g_Water outlet

aSS2D --_ "-'--_-Gas outlet

JSJ_ G_:I:::2_,41:_ Water inlet

as inlet

Fig. 1 Schematic diagram showing the ceramic sintering experiment

using dilatomentry.

NASA/TM--1998-208406 3

approximately 0.5 MPa was acting on the specimen during the entire sintering test. Therefore, this

experiment can also be considered a low constant-stress creep test for the ceramic materials.

During the sintering/creep experiments at various test temperatures, all specimens were heated in at

rate of 5°C/minute and held at the given test temperature for 15 hours, and then cooled down at a

rate of 5°C/minute to room temperature. Specimen shrinkage during the heating/cooling cycles was

continuously recorded in a computer system. Surface morphology changes of the specimens due

to the sintering process were examined using a scanning electron microscope (SEM). Phase

structures of the specimens before and after dilatometer sintering tests were also examined by

X-ray diffractometry with Cu K a radiation.

III. EXPERIMENTAL RESULTS

Figure 2 shows thermal expansion (shrinkage) results for the coating materials during the

sintering experiments at various temperatures measured by the dilatometry technique. Sintering

shrinkage were observed for all materials when the specimens were held at temperature for 15

hours. The shrinkage strains increased with increasing temperature. It can be seen that the HfO2-

27wt% Y20 3 showed the best sintering resistance. In contrast, CeO2- , Sc203, and NiO-doped

ZrO2-Y203 materials exhibit significant sintering shrinkage. Below the temperature of 900°C, no

significant shrinkage strains were detected for the given test time. Figure 3 illustrates the sintering

shrinkage strains occurring at the isothermal sintering stages as a function of temperature.

The sintering rates of the ceramic materials at the isothermal stages change with time,

especially at the early sintering time period. As shown in examples in Figure 4 (a) and (b), faster

shrinkage rates were observed initially, however, relatively constant rates were observed for longer

sintering times. At 1400°C as shown in Figure 4 (c), the "steady state" sintering rates for ZrO2-

Y203, ZrO2-CeO2-Y203, ZrO2-Sc203-Y203, ZrO2-NiO-Y203, and HfO2-Y203 are

2.6×10 -8/sec., 3.8x10 -8/sec. 4.2x10 -8/sec., 8.5×10 -8/sec. and 6.4x10 -9/sec., respectively.

Figure 4 (d) shows that for ZrO2-NiO-Y203, the second cycle resulted in further shrinkage of the

specimen at 1200°C.

Figure 5 shows the sintering shrinkage kinetics of plasma-sprayed ZrO2-8wt%Y203 at

1200°C in air and in Ar+5%H2. It can be seen that when the specimen was tested in a reducing

atmosphere, a faster sintering shrinkage rate was observed. In addition, the ceramic coating

material turned black after this Ar+H2 treatment. Increased sintering rates and darkened color were

observed for all other materials under the reduced oxygen partial pressure condition. This may

imply that the more defective structure of the materials due to Ar+H2 sintering would increase the

NASA/TM--1998-208406 4

Fig. 2

1.5

_" 1.0

0.5

"_ 0.0

.= -0.5

r_

1.5

1.0

0.5

0.0

-0,5

1.5

1,0

0.5

0.0

-0.5

1.5

1.0

0.5

0.0

-0.5

1.5

1.0

0.5

0.0

-0.5

-1.0

.'.. : ..... n:.,. n,..i.., n,,, n., i.t.u ,, ,a. ,j n... u, ,, io.. u.., i,,, n...n. In

.-q.,,

• .I; -.l... l...n.., i...i...i...n...l.. ,i...i. _. n "

. .. I...I... n...i...u.., n... n... l...l...i _..i...n...l. 1.1... I...

0 200 400 600 800 1000 1200 1400 1600Temperature, °C

Thermal expansion and sintering shrinkage response for the coating materials during the

dilatometry sintering experiments at various temperatures. (a) ZrO2-8wt%Y203;

(b) ZrO2-25wt%CeO2-2.5wt%Y203; (c) ZrO2-6wt%Sc203-2wt%Y203; (d) ZrO2-

6wt%NiO-9wt%Y203; (e) HfO2-27wt%Y203,

NASA/TM--1998-208406 5

¢D

¢D

r_

¢D_0

r_

Fig. 3

0.80

0.70

0.60

0.50:1

0.40

0.30

0.20

0.10

0.00

.... I .... I .... I .... I .... 1''''/3 ....

in air

ZrO -Y 02 : 2:3

-----o--- ZrOa-CeO.Y 20

_ _ ZrO2-NiO-Y203 / F

/ //

i i i II i i i i

800 900 1000 1100 1200 1300 1400 1500

Temperature, °C

Total sintering shrinkage strains for the coating materials at the

15 hour isothermal sintering stages as a function of temperature.

minority defect transport especially at the internal surfaces and grain boundaries, thus resulting in a

faster sintering rate.

Figure 6 illustrates some examples of the X-ray _action spectra for ZrO2-Y203, ZrO2-

CeO2-Y203, HfO2-Y203 and ZrO2-Sc203-Y203 after 15 hours sintering at 1200°C in air. From

X-ray diffraction experiments, it was found that the baseline ZrO2-8wt%Y203 primarily consisted

of tetragonal t' phase. The CeO2-doped ZrO2-Y203 also showed significant amount of t' phase,

however, the possibility that the cubic c phase might also be increased as compared to the baseline

material requires further study. Due to the high concentration of yttria dopant, HfO2-27wt%Y203

had a fully stabilized cubic c phase. No appreciable monoclinic phase was observed in these three

materials. Heat treatments related to the sintering experiments under various temperature and

oxygen pressure conditions did not measurably alter the phase structures of these materials. The as-

sprayed ZrO2-Sc203-Y203 material showed tetragonal t' phase and an increased amount of the

monoclinic m phase. The monoclinic phase in the Sc203-doped materials increased after the

sintering tests.

NASA/TM--1998-208406 6

1.5=_ 1.4

"a1.3

"_ 1.2

1.1O_

1.0

0.9

.4

.4

E

"_ 1.31.2

-a1.11.00.9

0.8

- d '_'_x._- _-_ .... _ 900 C-- ___22 _- ----o--- 1000ooC

_ _// _ _10o c -- /[// _ 1200°C: J/ .... //_. _ 1300°C

::,, __ 7_14°°_5 10 15 20 25 30

: : Timg, hours(a)

1.4

1.3 oC

1.2 _- ...... _ i IO0:C///........ _ _. _// + 1200 C

1.1 _- F- ___i:::_-_-_- ::\/// _ 1300oC

,.o_ _ . , -_ -+-,4ooo_0.9 ............... ,

0 5 10 :15 20 25 30

,_m(e_)_ours1.6

1.4 ) ' _ ...... i400°C ' 6.4.10-9/s-_..._ '

1.2 ....

1.o _

0.8

0.6

0.4

0'2 . .

0 5 10Time ' holr 5 20 25:(c)

1.6i i . , i i

/ 1 ooo zr02 -NiO'Y203 i

10 20 30 40 '5Time, hours

(d)

150010005000

150010005000

150010005000

150010005000

o

o_

©

Fig. 4 Sintering behavior of the ceramic materials at the isothermal stages. (a) and (b) The

sintering strains as a function of time and temperature for ZrO2-8wt%Y203 and ZrO2-

25wt%CeO2-2.5wt%Y203, respectively; (c) Steady state creep rates for the ceramic

materials at 1400°C; (d) Sintering shrinkage of ZrO2-6wt%NiO-9wt%Y203 at 1200°C

under two temperature cycles.

NASAfrM--1998-208406 7

..................... ..................__ _ _ _ __ __ : ,, _:_:_ :_ ___ :_ i _i_i__!i_?i_:i_i!_ _i_!i_:i____ iii!_i;!_iiii!_!!_i!!_ii!!_!i_ii!!!ii!!!!!_i!i!iiiii_!iii!i!!i_!!i_!i_!i!_!_i!_!_i_i_i_ill_ii!!_!_i!i!i_i!i_i_!_!!ii_!i_iii_!i!_i_!_!iiiii!iiiii!iii_i_iii!i!iii!iiiiiiiii_iiiJii_iiii!iii_iiiiiiiii

1.60

1.50

=_ 1.40

1.30

_ 1.20

1.10

t I I

T= 1200 °C

Air

Ar+5%H 2

I

1.00

0:0 5.0 10.0 15'0 20.0 25.0

Time, hours

1500O

1000

500

o

Fig. 5 Sintering shrinkage kinetics of plasma-sprayed ZrO2-8wt%Y203 at

1200°C in air and in Ar+5%H2.

The NiO-doped ZrO2-Y203 showed more complex phase structures. As shown in Figure

6, The majority phase in ZrO2-NiO-Y203 was the cubic c phase, instead of tetragonal phase in

the baseline material. However, the monoclinic rn phase was also present in this material. Because

of the limited solubility of NiO in ZrO2-Y203 (about 3 mol% at 1600°C [17] ), NiO phase was

observed in the as-sprayed and air-sintered specimens. In the Ar+H2 sintered specimens, however,

a Ni phase was present because of the reduction of NiO.

Surface microstructure changes were also observed after the sintering experiments. Certain

regions showed more noticeable sintering densification and grain growth as compared to other

regions, indicating there were some heterogeneities in the observed sintering phenomena. Figure 7

shows micrographs of ceramic surfaces of the ZrO2-8wt%Y203 coating material before and after

the dilatometry sintering at 1200°C. It can be seen that sintering which occurred could result in

microcrack healing and material densification, accompanying with substantial grain growth in

some regions.

NASAfrM--1998-208406 8

Z>

I_D_Doo

_aOoo4a

C_

Ca _

o _.

-_._o_°

O__o

= g

_-__

_ ° 0 o0

n o _ .

_" 0 _'t..,a

o ._

Relative intensity

tO

-m_11_) m(110)_ t - I_m_111) _1) 1

-mff_(111) c(lll) _m(m(ill) c(ill) f-m(ill) c(lll)

_::==m(200) _ c(200) _ c(200)

NiO(lll) ,NiO(11I)

_g_'.

NiO(200) "NiO(200) Ni(H 1}

¢:D _ _ _ " _c(220)" _ to t Ni(200)

, o :: o _ "

_:_ (311) _" _1!) _ii) "?J

O ('_c(222) + ...... _:>c(222) f _= >¢(222): 8

0

tO0

t.a.)0

g4_

-.1

Relative intensity :

I:::' " ' I ' ' '

C t(11t) _t(£: t(l I1) _ i_i e(I11 _ c(iil), ::: ' ( ) : :_ :

• t(O02) 02) ,,m(002,020) c;20:" !_t(20(J) r -_t(200) _c(200) _ I u)

l_ _,m_il_,_*i_ Ige "o lb

0 1% i I'_f_ _ ,_o

.-_____c ( 2 2 O), _-_.c ( 2 2 _),

_(202)/" _(202)

_t(l13) _-_ t(! l;_}

_c,t(31 i) Uc_t(3l 1)

c,t(222) f c,t(222)

m(022)--_._m(220), _- C(220).

_c(220)

m(300.202)

m(0t 3_221)

_c_311)

,e(222) 1 c(222i !

" i lil• ®iii_iiii!.

_i_iiiiiii!!_i_i_!!iiiiiiiiiiiil__i!_i!iiiiiiiiiiiii_i_i_i_i_i_i_i_i_iiiiiiiiiii<i_!_J__

(a)

: ........ '_!_ :iii: _i:_i_: ' i:i.:.:i. ...... _: :.

__ii!_:::": "_N "'::::::: :!?iii' ..5 ..::i?

• .:.....::_$_:i:!_:$ii_:i:_. "i:i::$::" .:4:4"..

....." !<::,iiiiiiii;ii!iii#, "

(b)

Fig. 7 Surface micrographs of the ZrO2-8wt%Y203 ceramic

coating material before and after the dilatometry sintering

at 1200°C for 15 hours. (a) Before the sintering test;

(b) After the sintering test.

IV. DISCUSSION

The sintering and low-stress creep characteristics of the ceramic coating materials,

determined by the dilatometer technique, are similar to the creep behavior of plasma-sprayed

coatings obtained from high temperature mechanical creep tests [1, 12, 20] and the laser

NASA/TM--1998 -208406 10

sintering/creep test [13] The fast initial creep rate and low creep activation energy have been

attributed to mechanical sliding, fast surface and grain boundary diffusion, and temperature and

stress gradient enhanced transport in the porous and weak ceramic coatings [13] . Figure 8 iUustrates

the creep rates of the plasma-sprayed ZrO2-8wt%Y203 as a function of stress and temperature

determined by the laser sintering technique [131. It can be seen that with higher compressive stresses

acting on the coating, a long primary creep stage and substantial sintering/creep rates can be

observed at much lower temperatures. In the high temperature, low stress sintering/creep tests by

the dilatometer technique, mechanical sliding becomes less predominant, and a nearly "steady-

state" creep region has been reached in a relatively short period of time. Diffusion-related processes

become more important mechanisms for the low stress sintering and creep deformation. The

observed grain growth phenomena also suggest the complex diffusion occurring during the

dilatometer sintering test.

10 -6

I I I t I, .Stress

[_. 50 MPa10 -7 _,_,_,"... - 1O0 MPa

_,_.._ .... ......... 200 MPa

....£-- _ ----..: ______10-8 'ZLZ

10-9 ..... ..... .... . .....

10 -10 , , , I , , , I , , , I , , , I , , ,

0.0 20.0 80.0 100.0

200°C

000°C

800°C

40.0 60.0

Time, hours

Fig. 8 The coating creep rates of plasma-sprayed ZrO2-8wt%Y203, determined

by laser high heat flux sintering/creep technique, as a function of stress

and temperature.

NASA/TM--1998-208406 11

Creep deformation of ceramic coating materials require diffusion of the cations and anions

in these materials. The creep rate in ceramics is therefore determined by the diffusion of the slowest

species, diffusing along the fastest path. In yttria-stabilized zirconia, the majority defect types are

oxygen vacancies and yttrium aliovalent dopants at normal cation sites. The possible minority

defects are zirconia interstitials, zirconium vacancies, and yttrium interstitials, and the zirconium

and yttrium cation transport is confirmed to be the slowest process in yttria-stabilized single

crystals [21] . The defect reactions in the yttria-stabilized zirconia can be written according to

Kr6ger-Vink notation [22] as

and

Y203=2Yzr +30_ + V_

O_9= VO +2e' +lo2(g )2

(1)

(2)

(for majority defects)

20 9+Zr r = Zri""+4e' +O2(g)

02(g) = Vzr +4h"

In the extrinsic region, the majority defect oxygen vacancy concentration Iv b]

determined by the dopant yttria concentration [Y'zr ], which follows the electroneutrality condition

At lower oxygen partial pressures in the intrinsic region where the electron conductivity

becomes important, oxygen vacancies can be further introduced according to Equation (2), that is

-1/2 ( M-/Vo[V'o]n2=KVoPO2 exp. RT (6)

where n is electron concentration, Kvo is reaction constant, AHvo is the enthalpy of formation of

oxygen vacancies, R and T are gas constant and temperature, respectively. Metal interstitials can

be an important defect type in the oxygen deficient oxide [23] , and the zirconium interstitial

concentration can be obtained from Equation (3)

zSd-lzri I[Zri'"]n4=Kzri""P_)_exp RT (7)

where Kzri... is a constant, M-Izri... is the enthalpy of formation of zirconium interstitials. In this

(3)

(4)

(for minority defects)

is

NASA/TM--1998-208406 12

intrinsic region, the electroneutrality can be expressed as

: n=2[Vo]+4[Zri'"].: (8)

By combining Equations (6)-(8), the oxygen vacancy and zirconia interstitial concentrations can be

written as

P02 exp AHv°3RT

4/3 -1/3I r'I ex(i

,2/5) -1/10 exp([Vo]?(Kvo/(4Kzr}) JPo 2

[Zri'"]=(Kzri... /25611/5po_15expl

3AHzr}... -4AHvo

3RT

5AHvo - 2AHzri...

5RT

AH zri"5RT )"

(9a)

(9b)

(when[v'ol>>[Zri'"])

(10a)

(10b)

(when[Vo]<<[Zri])

In the very high oxygen pressure region where the zirconium vacancies are predominant,

the electroneutrality condition can be written as

P= 4[Vz;] (11)

where p is electron hole concentration. The zirconium vacancy concentration can be obtained from

Equation (4) as

[Vz'_]=(Kvz'_ /25611/SplO/12Oexp AHVz_5RT (12)

where Kvz _ is a constant, AHvz r is the enthalpy of formation of zirconium vacancies. A Kr6ger-

Vink diagram is constructed based on these defect reactions and Equations (5), (9), (10) and (12),

as shown in Figure 9.

The increased sintering rate of ZrO2-8wt% Y203 at the reducing Ar+5%H2 atmosphere is

probably related to the defect structure change in the oxide. From the proposed Kr6ger-Vink

diagram shown in Figure 9, it can be seen that both concentrations of oxygen vacancies and

zirconia interstitials increase with reducing partial pressure of oxygen, especially in the low oxygen

activity region, At extremely low oxygen pressures, the metal cation interstitials can even become

the dominant defect type. Therefore, it is possible that the highly defective oxide structures under

NASA/TM--1998 -208406 13

_ 02 i • •• • /

- _, i: . ,

[ 1,,.l _. i _ t .I

_: 1/4o _ Vr;."] n ,o: " "02'_ __:P02

I n+p

Fig. 9

L°g [POE1

Proposed Pseudo-Kr6ger-Vink diagram illustrating the possible

majority and minority defects in ZrO2-8wt%Y203.

low oxygen pressures facilitates the metal cation interstitial formation, thus resulting in faster metal

cation diffusion and the increased sintering rate. Thornton et al. [24] have also observed enhanced

cerium migration and segregation in the ZrO2-25wt%CeO2-2.5wt%Y203 material under relatively

moderate reducing conditions, further confirming the increased cation mobility in more oxygen-

deficient oxide under the low oxygen activity conditions.

Since the ceramic sintering requires the transport of the minority cations, the stability of the

ceramic materials (both dopants and base materials) will have influence on the sintering behavior.

The present study has shown that there is a close relationship between the oxide chemical and

phase stability and the sintering rate. Halfnia-based oxides have higher chemical stability, and lower

oxygen partial pressures for the transition of ionic conductivity to electronic conductivity, as

compared to zirconia-based oxides, therefore it is not surprised that the HfO2-27wt%Y203

exhibited the lowest sintering rates. On the other hand, the CeO2-doped ZrO2 exhibited large

electron contributions at even moderate temperatures and oxygen activities [25]. As shown in

Figure 9, the increased region of electron conductivity implies an extended metal cation interstitial

NASA/TM-- 1998 -208406 14

region, in which the cation interstitial concentration is increased with decreasing oxygen partial

pressure, and thus resulting in possible enhanced metal cation diffusion with reducing oxygen

activity. Insufficient doping, as possibly occurred for Sc203-doped ZrO2-Y203, will have a

similar effect on the metal cation diffusion. For the NiO-doped ZrO2-Y203, the observed high

sintering rates may also be related to NiO segregation at the grain boundaries, which may act as a

sintering agent. At lower oxygen partial pressures, the NiO reduction to metallic Ni, as observed in

this experiment, can further enhance the sintering process. It is suggested that the chemical and

phase stability of both the base oxides and dopant oxides is critical to the sintering and creep

behavior of the ceramic materials.

V. CONCLUSIONS

1. Sintering shrinkage strains were observed at the isothermal stage for all ceramic coating

materials tested in the dilatometer sintering experiments. The HfO2-27wt%Y203 and baseline

ZrO2-8wt%Y203 exhibited the best sintering resistance, and NiO-doped ZrO2-Y203 showed

the highest shrinkage strain rates during the tests.

2. The higher shrinkage strain rates of the coating materials were observed for the specimens

tested in Ar+5%H2 as compared to those tested in air. This phenomenon was attributed to a

proposed enhanced metal cation interstitial diffusion mechanism under the reducing conditions.

3. There was a close relationship between the observed sintering behavior and chemical and phase

stability of the coating materials. Increased chemical stability of base oxides and dopants seems

to improve materials phase stability at high temperature, and sintering/creep resistance.

Insufficient doping and dopant-segregation-induced depletion will facilitate the sintering

process.

[1]

[21

[31

[4]

REFERENCES

R. F. Firestone, W. R. Logan, J. W. Adams and R. C. J. Bill, in J. D. Buckley, C. M.

Packer and J. J. Gebhardt (eds.), The 6th Annual Conference on Composites and Advanced

Ceramic Materials, The American Ceramic Society, Inc., Columbus, OH 43214, 1982,

p. 758.

R. F. Firestone, W. R. Logan and J.W. Adams, NASA CR-167868, November, 1982.

H. E. Eaton and R. C. Novak, Surface and Coatings Technology, 32 (1987) 227.

T. A. Cruse, S. E. Stewart and M. Ortiz, in (eds.), The Gas Turbine and Aeroengine

Congress and Exposition, The American Society of Mechanical Engineers, New York,

New York 10017, 1988, p. 1.

NASA/TM--1998 -208406 15

[5] S.M. Meier, D. M. Nissely and K. D. Sheffier, NASA CR-!82230, July 1991.

[6] H.E. Eaton, J. R. Linsey and R. B. Dinwiddie, in T. W. Tong (eds.), Thermal

Conductivity, Lancaster, Pennsylvania, 1994, p. 289.

[7] J.G. Goedjen, W. J. Brindley and R. A. Miller, in C. C. Berndt and S. Sampath (eds.),

Advances in Thermal Spray Science and Technology, ASM International, Materials Park,

Ohio, 1995, p. 73.

[8] R.B. Dinwiddie, S. C. Beecher, W. Porter, D. and A. B. Nagaraj, in (eds.), The

International Gas Turbine and Aeroengine Congress and Exhibition, The American Society

of Mechanical Engineers, New York, New York 10017, 1996, p. 1.

[9] D. Zhu and R. A. Miller, Surface and Coatings Technology, 94-95 (1997) 94.

[10] D. Zhu and R. A. Miller, in W. J. Brindley (eds.), The 3rd Thermal Barrier Coating

Workshop, sponsored by the TBC Interagency Coordination Committee NASA Lewis

Research Center, Cleveland, Ohio, 1997, p. 139. Materials Science and Engineering, A245

(1998) 212.

[11 ] K.F. Wesling, D. F. Socie and B. Beardsley, Journal of the American Ceramic Society, 77

(1991) 1863.

[12] G. Thurn, G. A. Schneider and F. Aldinger, Materials Science and Engineering, A233

(1997) 176.

[13] D. Zhu and R. A. Miller, NASA TM-113169, Army Research Laboratory Report ARL-

TR-1565, November 1997.

[14] T.A. Cruse, B. P. Johnsen and A. Nagy, Journal of Thermal Spray Technology, 6 (1997)

57.

[15] S.R. Choi, D. Zhu and R. A. Miller, in (eds.), International Symposium on Advanced

Synthesis and Processing, The 22nd Annual Conference on Composites and Advanced

Ceramic Materials, 1998.

[16] R.A. Miller and G. W. Leissler, NASA TP-3296, March 1993.

[17] S. Chen and P. Shen, Materials Science and Engineering, A123 (1990) 145.

[18] R.L. Jones, R. F. Reidy and D. Mess, Surface and Coatings Technology, 82 (1996) 72.

[19] D. Zhu and R. A. Miller, NASA Technical Paper TP-3676, Army Research Laboratory

Technical Report ARL-TR-1341, May 1997.

[20] B.P. Johnsen, T. A. Cruse, R. A. Miller and W. J. Brindley, Journal of Engineering

Materials and Technology, 117 (1995) 305.

NASA/TM--1998 -208406 16

[21] H. Solmon, J. Chaumont, C. Dolin and C. Monty, in T. O. Mason and J. L. Routbort

(eds.), Ceramic Transactions - Point Defects and Related Properties of Ceramics The

American Ceramic Society, Inc, Westerville, Ohio, 1991, p. 175.

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Amsterdam, 1964.

[23] P. Kofstad, Nonstoichiometry, Diffusion, and Electrical Conductivi_ty in Binary_ Metal

Oxides, Robert E. Krieger Publishing Company, Malabar, Florida, 1983.

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(1997) 112.

[25] R.F. Reidy and G. Simkovich, Solid State Ionics, 62 (1993) 85.

NASA/TM--1998 -208406 17

REPORT DOCUMENTATION PAGE Form ApprovedOMB No. 0704-0188

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1. AGENCY USE ONLY (Leave blank) 2. REPORT DATE 3. REPORT TYPE AND DA_.=5 COVERED

July 1998 Technical Memorandum

4. TITLE AND iSU_ t i tLE 5. FUNDING NUMBERS

Sintering and Creep Behavior of Plasma-Sprayed Zirconia and Hafnia Based

Thermal Barrier Coatings

6. AUTHOR(S) '

Dongming Zhu and Robert A. Miller

7. PERFORMING ORGANIZATION NAME(S) AND ADDRESS(ES)

National Aeronautics and Space Administration

Lewis Research Center

Cleveland, Ohio 44135 -3191

9. SPONSORING/MONITORING AGENCY NAME(S) AND ADDRESS(ES)

National Aeronautics and Space Administration

Washington, DC 20546-0001

WU-505-23-2U-00

8. PERFORMING ORGANIZATIONREPORT NUMBER

E-11239

10. SPONSORING/MONITORINGAGENCY REPORTNUMBER

11. SUPPLEMENTARY NO_,'-S

NASA TM--1998-208406

Prepared for the 25th International Conference on Metallurgical Coatings and Thin Films sponsored by the American

Vacuum Society, San Diego, California, April 27--May 1, 1998. Dongming Zhu, Ohio Aerospace Institute, 22800 Cedar

Point Rd., Cleveland, Ohio 44142, and Robert A. Miller, NASA Lewis Research Center. Responsible person, DongmingZhu, organization code 5160, (216) 433-5422.

12b. DISTRIBUTION CODE12a. Di-_¥HIBUTION/AVAILABILITY STATEMENT

13.

Unclassified - Unlimited

Subject Category: 27 Distribution: Nonstandard

This publication is available from the NASA Center for AeroSpace Information, (301) 621-0390.

AB_THACT (Maximum 200 words)

The sintering and creep of plasma-sprayed ceramic thermal barrier coatings under high temperature conditions are complex

phenomena. Changes in thermomechanical and thermophysical properties and in the stress response of these coating sys-

tems as a result of the sintering and creep processes are detrimental to coating thermal fatigue resistance and performance.

In this paper, the sintering characteristics of ZrO2-8wt%Y203, ZrO2_25wt%CeO2_2.5wt%Y203 ' ZrO2_6w%NiO_

9wt%Y203, ZrO2-6wt%Sc203-2wt%Y203 and HfO2-27wt%Y203 coating materials were investigated using dilatometry.

It was found that the HfO2-Y203 and baseline ZrO2-Y203 exhibited the best sintering resistance, while the NiO-doped

ZrO2-Y203 showed the highest shrinkage strain rates during the tests. Higher shrinkage strain rates of the coating materials

were also observed when the specimens were tested in Ar+5%H2 as compared to in air. This phenomenon was attributed to

an enhanced metal cation interstitial diffusion mechanism under the reducing conditions. It is proposed that increased

chemical stability of coating materials will improve the material sintering resistance.

14. SUBJECT I=HMS

Thermal barrier coatings; Ceramic sintering and creep; Defect structure; Dilatometry

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