Surface and Coatings Technology 108–109 (1998) 114–120
`
`Sintering and creep behavior of plasma-sprayed zirconia- and hafnia-based
`thermal barrier coatings
`
`*
`Dongming Zhu , Robert A. Miller
`National Aeronautics and Space Administration, Lewis Research Center, Cleveland, OH 44135,USA
`
`Abstract
`
`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
`ZrO –8wt%Y O , ZrO –25wt%CeO –2.5wt%Y O , ZrO –6w%NiO–9wt%Y O , ZrO –6wt%Sc O –2wt%Y O
`and HfO –
`2
`2
`3
`2
`2
`2
`3
`2
`2
`3
`2
`2
`3
`2
`2
`3
`27wt%Y O coating materials were investigated using dilatometry. It was found that the HfO –Y O and baseline ZrO –Y O exhibited
`2
`3
`2
`2
`3
`2
`2
`3
`the best sintering resistance, while the NiO-doped ZrO –Y O showed the highest shrinkage strain rates during the tests. Higher
`2
`2
`3
`shrinkage strain rates of the coating materials were also observed when the specimens were tested in Ar15%H as compared to in air.
`2
`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.
`1998 Elsevier Science
`S.A. All rights reserved.
`
`Keywords: Thermal barrier coatings; Ceramic sintering and creep; Defect structure; Dilatometry
`
`1. 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 coat-
`ing 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–segmen-
`tation-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
`
`*Corresponding author. Ohio Aerospace Institute, NASA Lewis Re-
`search Center. Tel.: 11-216-4335422; fax: 11-216-4335544; e-mail:
`Dongming.Zhu@lerc.nasa.gov
`
`been made in characterizing the ceramic coating sintering
`and creep behavior at high temperature and under tempera-
`ture 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 ZrO –8wt%Y O at
`room
`2
`2
`3
`temperature has also been observed at a tensile stress of
`7.4 MPa [15]. A mechanism-based model has been pro-
`posed to describe the densification and deformation occur-
`ring in thermal barrier coatings at temperature by taking
`into account the thermally and stress-activated diffusion,
`and the mechanical compacting processes [13]. The dop-
`ants 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.
`
`0257-8972/98/$ – see front matter
`PII: S0257-8972( 98 )00669-0
`
`1998 Elsevier Science S.A. All rights reserved.
`
` 1
`
`UTC 2008
`General Electric v. United Technologies
`IPR2016-01289
`
`(cid:211)
`(cid:211)
`

`
`D. Zhu, R.A. Miller / Surface and Coatings Technology 108–109(1998)114–120
`
`115
`
`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) ZrO –8wt%Y O , a NASA–Lewis Research
`2
`2
`3
`Center
`reference
`(or baseline) material;
`(b) ZrO –
`2
`25wt%CeO –2.5wt%Y O ,
`a
`commercially
`available
`2
`2
`3
`coating material developed for hot corrosion resistance; (c)
`HfO –27wt%Y O , a potential new coating material
`2
`2
`3
`developed at NASA for high temperature stability [16]; (d)
`ZrO –6wt%NiO–9wt%Y O , a NiO-doped ZrO –Y O
`2
`2
`3
`2
`2
`3
`coating material reported to suppress the tetragonal–mono-
`clinic phase transformation [17]; (e) ZrO –6wt%Sc O –
`2
`2
`3
`2wt%Y O , developed for improved hot corrosion resist-
`2
`3
`ance [18]. The coating sintering and creep mechanisms and
`dopant effect on coating sintering rates are discussed based
`on experimental observations and possible defect reactions.
`
`approximately 450 g using a spring device) on the speci-
`men, a uniaxial stress of 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 tempera-
`tures, all specimens were heated at a rate of 58C /min and
`held at the given test temperature for 15 h, and then cooled
`down at a rate of 58C/min to room temperature. Specimen
`shrinkage during the heating /cooling cycles was continu-
`ously 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 di-
`latometer sintering tests were also examined by X-ray
`diffractometry with Cu Ka radiation.
`
`2. Experimental materials and methods
`
`The five ceramic coating materials mentioned above,
`ZrO –8wt%Y O ,
`ZrO –25wt%CeO –2.5wt%Y O ,
`2
`2
`3
`2
`2
`2
`3
`ZrO –6wt%NiO–9wt%Y O ,
`ZrO –6wt%Sc O –
`2
`2
`3
`2
`2
`3
`2wt%Y O and HfO –27wt%Y O , were chosen for this
`2
`3
`2
`2
`3
`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 ZrO –25CeO –2.5wt%Y O which was spray dried
`2
`2
`2
`3
`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 mm 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 6008C for 6 h 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 Ar15%H within the temperature range of 900–
`2
`TM
`14008C, using a UNITHERM
`high temperature di-
`latometer system shown in Fig. 1. Since the push rod in the
`dilatometer exerts a certain amount of force (measured at
`
`Fig. 1. Schematic diagram showing the ceramic sintering experiment
`using dilatometry.
`
`3. Experimental results
`
`Fig. 2 shows thermal expansion (shrinkage) results for
`the coating materials during the sintering experiments at
`various temperatures measured by the dilatometry tech-
`nique. Sintering shrinkage was observed for all materials
`when the specimens were held at temperature for 15 h. The
`shrinkage strains increased with increasing temperature. It
`can be seen that the HfO –27wt%Y O showed the best
`2
`2
`3
`sintering resistance. In contrast, CeO -, Sc O -, and NiO-
`2
`2
`3
`doped ZrO –Y O materials exhibited significant sintering
`2
`2
`3
`shrinkage. Below the temperature of 9008C, no significant
`shrinkage strains were detected for the given test time. Fig.
`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 Fig. 4a,b,
`faster shrinkage rates were observed initially, however,
`relatively constant rates were observed for longer sintering
`times. At 14008C as shown in Fig. 4c, the ‘steady-state’
`sintering rates for ZrO –Y O , ZrO –CeO –Y O , ZrO –
`2
`2
`3
`2
`2
`2
`3
`2
`Sc O –Y O , ZrO –NiO–Y O , and HfO –Y O are
`2
`3
`2
`3
`2
`2
`3
`2
`2
`3
`28
`28
`28
`29
`2.6310 , 3.8310 , 8.5310
`and 6.4310
`/s, respec-
`tively. Fig. 4d shows that
`for ZrO –NiO–Y O ,
`the
`2
`2
`3
`second cycle resulted in further shrinkage of the specimen
`at 12008C.
`Fig. 5 shows the sintering shrinkage kinetics of plasma-
`sprayed ZrO –8wt%Y O at 12008C in air and in Ar1
`2
`2
`3
`5%H . It can be seen that when the specimen was tested in
`2
`a reducing atmosphere, a faster sintering shrinkage rate
`was observed. In addition, the ceramic coating material
`turned black after this Ar1H treatment. Increased sinter-
`2
`ing rates and darkened color were observed for all other
`materials under the reduced oxygen partial pressure con-
`dition. This may imply that the more defective structure of
`the materials due to Ar1H sintering would increase the
`2
`
` 2
`
`

`
`116
`
`D. Zhu, R.A. Miller / Surface and Coatings Technology 108–109(1998)114–120
`
`Fig. 2. Thermal expansion and sintering shrinkage response for the
`coating materials during the dilatometry sintering experiments at various
`temperatures.
`(a)
`ZrO –8wt%Y O ;
`(b)
`ZrO –25wt%CeO –
`2
`2
`3
`2
`2
`2.5wt%Y O ; (c) ZrO –6wt%Sc O –2wt%Y O ; (d) ZrO –6wt%NiO–
`2
`3
`2
`2
`3
`2
`3
`2
`9wt%Y O ; (e) HfO –27wt%Y O .
`2
`3
`2
`2
`3
`
`Fig. 4. Sintering behavior of the ceramic materials at the isothermal
`stages. (a,b) The sintering strains as a function of time and temperature
`for ZrO –8wt%Y O and ZrO –25wt%CeO –2.5wt%Y O , respective-
`2
`2
`3
`2
`2
`2
`3
`ly; (c) steady-state creep rates for the ceramic materials at 14008C; (d)
`sintering shrinkage of ZrO –6wt%NiO–9wt%Y O at 12008C under two
`2
`2
`3
`temperature cycles.
`
`Fig. 3. Total sintering shrinkage strains for the coating materials at the
`15-h isothermal sintering stages as a function of temperature.
`
`Fig. 5. Sintering shrinkage kinetics of plasma-sprayed ZrO –8wt%Y O
`2
`2
`at 12008C in air and in Ar15%H .2
`
`3
`
`minority defect transport especially at the internal surfaces
`and grain boundaries, thus resulting in a faster sintering
`rate.
`Fig. 6 illustrates some examples of the X-ray diffraction
`
` 3
`
`

`
`D. Zhu, R.A. Miller / Surface and Coatings Technology 108–109(1998)114–120
`
`117
`
`phase structures. As shown in Fig. 6, the majority phase in
`ZrO –NiO–Y O was the cubic c phase,
`instead of
`2
`2
`3
`tetragonal phase in the baseline material. However, the
`monoclinic m phase was also present
`in this material.
`Because of the limited solubility of NiO in ZrO –Y O
`2
`2
`3
`(about 3 mol% at 16008C [17]), NiO phase was observed
`in the as-sprayed and air-sintered specimens. In the Ar1H2
`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. Fig.
`7 shows micrographs of ceramic surfaces of the ZrO –2
`8wt%Y O coating material before and after
`the di-
`2
`3
`latometry sintering at 12008C. It can be seen that sintering
`which occurred could result
`in microcrack healing and
`material densification, accompanying with substantial grain
`growth in some regions.
`
`Fig. 6. X-ray diffraction spectra of the plasma-sprayed ceramic coating
`materials. (a) Diffraction spectra of ZrO –Y O , ZrO –CeO –Y O ,
`2
`2
`3
`2
`2
`2
`3
`ZrO –Sc O –Y O and HfO –Y O after 12008C sintering in air; (b)
`2
`2
`3
`2
`3
`2
`2
`3
`diffraction spectra of ZrO –NiO–Y O under the as-sprayed condition,
`2
`2
`3
`and after 12008C sintering in air and 12008C sintering in Ar15%H .2
`
`spectra for ZrO –Y O , ZrO –CeO –Y O , HfO –Y O
`2
`2
`3
`2
`2
`2
`3
`2
`2
`3
`and ZrO –Sc O –Y O after 15 h sintering at 12008C in
`2
`2
`3
`2
`3
`air. From X-ray diffraction experiments, it was found that
`the baseline ZrO –8wt%Y O primarily consisted of
`2
`2
`3
`tetragonal
`t9 phase. The CeO -doped ZrO –Y O also
`2
`2
`2
`3
`showed significant amount of
`t9 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, HfO –2
`27wt%Y O had a fully stabilized cubic c phase. No
`2
`3
`appreciable monoclinic phase was observed in these three
`materials. Heat treatments related to the sintering experi-
`ments under various temperature and oxygen pressure
`conditions did not measurably alter the phase structures of
`these materials. The as-sprayed ZrO –Sc O –Y O ma-
`2
`2
`3
`2
`3
`terial showed tetragonal t9 phase and an increased amount
`of the monoclinic m phase. The monoclinic phase in the
`Sc O -doped materials increased after the sintering tests.
`2
`3
`The NiO-doped ZrO –Y O showed more complex
`2
`2
`3
`
`Fig. 7. Surface SEM micrographs of the ZrO –8wt%Y O ceramic
`2
`2
`3
`coating material before and after the dilatometry sintering at 12008C for
`15 h. (a) Before the sintering test; (b) after the sintering test.
`
` 4
`
`

`
`118
`
`D. Zhu, R.A. Miller / Surface and Coatings Technology 108–109(1998)114–120
`
`(1)
`
`(2)
`
`(3)
`
`(4)
`
`..
`3
`Y O 5 2Y 1 3O 1 V
`9
`O
`O
`Zr
`2
`3
`
`..
`3
`O 5V 1 2e9 1 O ( g)
`O
`O
`2
`2
`(for majority defects) and
`....
`3
`3
`2O Zr 5 Zr 1 4e9 1 O ( g)
`O
`Zr
`2
`i
`
`1]
`
`1
`]
`2
`
`O ( g) 5V 999 1 4h
`9
`Zr
`2
`
`.
`
`(for minority defects)
`
`(5)
`
`In the extrinsic region, the majority defect oxygen vacancy
`..
`concentration [V ] is determined by the dopant yttria
`O
`concentration [Y ], which follows the electroneutrality
`9
`Zr
`condition
`..
`[Y ] 5 2[V ].
`9
`Zr
`O
`At lower oxygen partial pressures in the intrinsic region
`where the electron conductivity becomes important, oxy-
`gen vacancies can be further introduced according to Eq.
`(2), that is
`
`..
`2
`[V ]n 5 K p
`..
`O
`V
`O
`
`21 / 2
`O
`
`2
`
`exp 2
`
`S
`
`D]]
`
`(6)
`
`(7)
`
`(8)
`
`DH ..
`VO
`RT
`where n is electron concentration, K is reaction constant,
`..V O
`DH is the enthalpy of formation of oxygen vacancies, R
`..V O
`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 Eq. (3)
`DH ....
`Zr
`RT
`is the enthalpy of
`is a constant, DH
`where K
`Zr
`Zr
`formation of zirconium interstitials. In this intrinsic region,
`the electroneutrality can be expressed as
`..
`....
`n 5 2[V ] 1 4[Zr
`O
`i
`
`[Zr
`
`....
`i
`
`4
`]n 5 K
`
`p
`
`21
`O
`
`2
`
`Zr
`
`....
`i
`
`S
`
`exp 2
`
`D]]
`
`i
`
`.
`
`....
`i
`
`....
`i
`
`].
`
`the oxygen vacancy and
`By combining Eqs. (6)–(8),
`zirconia interstitial concentrations can be written as
`DH ..VO
`3RT
`
`..
`[V ] 5 (K / 4)
`..
`O
`V
`O
`
`1 / 3 21 / 6
`p
`O
`
`2
`
`S
`
`exp 2
`
`D]]
`
`(9a)
`
`[Zr
`
`....
`i
`
`] 5 (K
`
`Zr
`
`....
`i
`
`S
`
`4 / 3 21 / 3
`/(2K )
`)p
`..
`V
`O
`O
`2
`3DH 2 4DH
`..
`....
`V
`Zr
`O
`i
`]]]]]
`3RT
`
`D
`
`.
`
`exp 2
`(when [V ] 4 [Zr
`
`..
`O
`
`..
`[V ] 5 (K /(4K
`..
`O
`V
`O
`
`S
`
`exp 2
`
`2 / 5 21 / 10
`)p
`Zr
`O
`2
`5DH 2 2DH
`..
`....
`V
`Zr
`O
`i
`]]]]]
`5RT
`
`....
`i
`
`])
`
`)
`
`....
`i
`
`....
`i
`
`[Zr
`] 5 (K
`/256)
`(when [V ] < [Zr
`
`Zr
`
`....
`i
`
`..
`O
`
`1 / 5 21 / 5
`p
`O
`
`2
`
`])
`
`....
`i
`
`exp 2
`
`D
`S
`
`(9b)
`
`(10a)
`
`(10b)
`
`D]]
`
`DH ....
`Zr i
`5RT
`
`.
`
`4. 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 mechani-
`cal creep tests [1,12,20] and the laser 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]. Fig. 8 illustrates the creep rates of
`the plasma-sprayed ZrO –8wt%Y O as a function of
`2
`2
`3
`stress and temperature determined by the laser sintering
`technique [13]. 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 phe-
`nomena also suggest
`the complex diffusion occurring
`during the dilatometer sintering test.
`Creep deformation of ceramic coating materials requires
`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 Kroger–
`Vink notation [22] as
`
`Fig. 8. The coating creep rates of plasma-sprayed ZrO –8wt%Y O ,
`2
`2
`3
`determined by laser high heat flux sintering/creep technique, as a
`function of stress and temperature.
`
` 5
`
`

`
`D. Zhu, R.A. Miller / Surface and Coatings Technology 108–109(1998)114–120
`
`119
`
`In the very high oxygen pressure region where the
`zirconium vacancies are predominant, the electroneutrality
`condition can be written as
`
`p 5 4[V 999].
`9
`Zr
`
`(11)
`
`where p is electron hole concentration. The zirconium
`vacancy concentration can be obtained from Eq. (4) as
`
`[V 999] 5 (K
`9
`Zr
`V 9999
`Zr
`
`/256)
`
`1 / 15 1 / 10
`p
`O
`
`2
`
`S
`
`exp 2
`
`D]]]
`
`DHV 9999
`Zr
`5RT
`
`.
`
`(12)
`
`is the enthalpy of
`is a constant, DH
`where K
`V 9999
`V 9999
`Zr
`Zr
`¨
`formation of zirconium vacancies. A Kroger–Vink dia-
`gram is constructed based on these defect reactions and Eq.
`(5), Eqs. (9a) and (9b), Eqs. (10a) and (10b) and Eq. (12),
`as shown in Fig. 9.
`The increased sintering rate of ZrO –8wt%Y O at the
`2
`2
`3
`reducing Ar15%H atmosphere is probably related to the
`2
`defect structure change in the oxide. From the proposed
`¨
`Kroger–Vink diagram shown in Fig. 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 intersti-
`tials can even become the dominant defect type. Therefore,
`it is possible that
`the highly defective oxide structures
`under low oxygen pressures facilitate 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 ZrO –25wt%CeO –2.5wt%Y O ma-
`2
`2
`2
`3
`terial under relatively moderate reducing conditions, fur-
`ther 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
`
`¨
`Fig. 9. Proposed pseudo-Kroger–Vink diagram illustrating the possible
`majority and minority defects in ZrO –8wt%Y O .
`2
`2
`3
`
`there is a close relationship between the oxide chemical
`and phase stability and the sintering rate. Hafnia-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 a surprise that
`the HfO –2
`27wt%Y O exhibited the lowest sintering rates. On the
`2
`3
`other hand, the CeO -doped ZrO exhibited large electron
`2
`2
`contributions at even moderate temperatures and oxygen
`activities [25] . As shown in Fig. 9, the increased region of
`electron conductivity implies an extended metal cation
`interstitial region,
`in which the cation interstitial con-
`centration 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 Sc O -doped ZrO –
`2
`3
`2
`Y O , will have a similar effect on the metal cation
`2
`3
`diffusion. For the NiO-doped ZrO –Y O , the observed
`2
`2
`3
`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.
`
`5. Conclusions
`
`the
`(1) Sintering shrinkage strains were observed at
`isothermal stage for all ceramic coating materials tested in
`the
`dilatometer
`sintering
`experiments. The HfO –
`2
`27wt%Y O and baseline ZrO –8wt%Y O exhibited the
`2
`3
`2
`2
`3
`best sintering resistance, and NiO-doped ZrO –Y O
`2
`2
`3
`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 Ar1
`5%H as compared to those tested in air. This phenomenon
`2
`was attributed to a proposed enhanced metal cation inter-
`stitial 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 resist-
`ance. Insufficient doping and dopant-segregation-induced
`depletion will facilitate the sintering process.
`
`Acknowledgements
`
`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.
`
` 6
`
`

`
`120
`
`D. Zhu, R.A. Miller / Surface and Coatings Technology 108–109(1998)114–120
`
`References
`
`in: J.D.
`[1] R.F. Firestone, W.R. Logan, J.W. Adams, R.C.J. Bill,
`Buckley, C.M. Packer, J.J. Gebhardt
`(Eds.), The 6th Annual
`Conference on Composites and Advanced Ceramic Materials, The
`American Ceramic Society, Columbus, OH 43214, 1982, p. 758.
`[2] R.F. Firestone, W.R. Logan, J.W. Adams, NASA CR-167868,
`November, 1982.
`[3] H.E. Eaton, R.C. Novak, Surf. Coat. Technol. 32 (1987) 227.
`[4] T.A. Cruse, S.E. Stewart, M. Ortiz (Eds.), in: The Gas Turbine and
`Aeroengine Congress and Exposition, The American Society of
`Mechanical Engineers, New York, NY 10017, 1988, p. 1.
`[5] S.M. Meier, D.M. Nissely, K.D. Sheffler, NASA CR-182230, July
`1991.
`in: T.W. Tong (Ed.),
`[6] H.E. Eaton, J.R. Linsey, R.B. Dinwiddie,
`Thermal Conductivity, Lancaster, PA, 1994, p. 289.
`[7] J.G. Goedjen, W.J. Brindley, R.A. Miller,
`in: C.C. Berndt, S.
`Sampath (Eds.), Advances in Thermal Spray Science and Technolo-
`gy, ASM International, Materials Park, OH, 1995, p. 73.
`[8] R.B. Dinwiddie, S.C. Beecher, W. Porter, A.B. Nagaraj (Eds.), The
`International Gas Turbine and Aeroengine Congress and Exhibition,
`The American Society of Mechanical Engineers, New York, NY
`10017, 1996, p. 1.
`[9] D. Zhu, R.A. Miller, Surf. Coat. Technol. 94–95 (1997) 94.
`[10] D. Zhu, R.A. Miller, Mater. Sci. Eng. A245 (1998) 212.
`[11] K.F. Wesling, D.F. Socie, B. Beardsley, J. Am. Ceramic Soc. 77
`(1991) 1863.
`
`[12] G. Thurn, G.A. Schneider, F. Aldinger, Mater. Sci. Eng. A233
`(1997) 176.
`[13] D. Zhu, R.A. Miller, NASA TM-113169, Army Research Labora-
`tory Report ARL-TR-1565, November 1997.
`[14] T.A. Cruse, B.P. Johnsen, A. Nagy, J. Thermal Spray Technol. 6
`(1997) 57.
`[15] S.R. Choi, D. Zhu, R.A. Miller (Eds.), in: International Symposium
`on Advanced Synthesis and Processing, The 22nd Annual Confer-
`ence on Composites and Advanced Ceramic Materials, 1998.
`[16] R.A. Miller, G.W. Leissler, NASA TP-3296, March 1993.
`[17] S. Chen, P. Shen, Mater. Sci. Eng. A123 (1990) 145.
`[18] R.L. Jones, R.F. Reidy, D. Mess, Surf. Coat. Technol. 82 (1996) 72.
`[19] D. Zhu, 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, W.J. Brindley, J. Eng. Mater.
`Technol. 117 (1995) 305.
`[21] H. Solmon, J. Chaumont, C. Dolin, C. Monty, in: T.O. Mason, J.L.
`Routbort (Eds.), Ceramic Transactions—Point Defects and Related
`Properties of Ceramics, The American Ceramic Society, Westerville,
`OH, 1991, p. 175.
`¨
`[22] F.A. Kroger, The Chemistry of Imperfect Crystals, North-Holland,
`Amsterdam, 1964.
`[23] P. Kofstad, Nonstoichiometry, Diffusion, and Electrical Conductivity
`in Binary Metal Oxides, Robert E. Krieger, Malabar, FL, 1983.
`[24] J. Thornton, A. Majumdar, G. McAdam, Surf. Coat. Technol. 94–95
`(1997) 112.
`[25] R.F. Reidy, G. Simkovich, Solid State Ionics 62 (1993) 85.
`
` 7