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Available online at www.sciencedirect.com

Journal of the European Ceramic Society 34 (2014) 55–61

Preparation and thermophysical properties of fluorite-typesamarium–dysprosium–cerium oxides

Zhang Hongsong a,∗, Yan Shuqing b, Chen Xiaoge c

a Department of Mechanical Engineering, Henan Institute of Engineering, Zhengzhou 450007, Chinab School of Mathematics and Information Science, North China University of Water Conservancy and Hydroelectric, Zhengzhou 450007, China

c Department of Construction Engineering, Henan Institute of Engineering, Zhengzhou 450007, China

Received 11 May 2013; received in revised form 9 July 2013; accepted 19 July 2013Available online 15 August 2013

bstract

n the present study, (Sm1−xDyx)2Ce2O7 solid solutions were synthesized by solid reaction at 1600 ◦C for 10 h in air. The phase structure, micro-orphology and thermophysical properties of (Sm1−xDyx)2Ce2O7 oxides were examined. XRD results indicated that pure (Sm1−xDyx)2Ce2O7

xides with fluorite structure were prepared. SEM revealed that their microstructures were very dense and there were no other phases among thearticles. The thermal conductivity and thermal expansion coefficient of the ceramics remarkably decreased through Dy-doping. Their thermalxpansion coefficients were higher than that of YSZ, and their thermal conductivities were much lower than that of 8YSZ. Their excellent

hermophysical properties imply that these solid solutions are potential materials for the ceramics layer in thermal barrier coatings.

2013 Elsevier Ltd. All rights reserved.

eywords: Thermal barrier coatings; Rare earth cerium oxides; Thermophysical property; Doping; Solid reaction

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. Introduction

Thermal barrier coatings (TBCs) are widely used in gas tur-ine engines and other components which must endure highemperature.1–3 Generally, a typical TBC consists of a ther-

ally insulating ceramic top layer and intermediate oxidationesistant metallic bond-coat layer. The main purpose of thisual layer system is to provide thermal protection of the super-lloy components, thus improving performance or lifetime.4

ow, the demand for improved performance in high-temperatureechanical systems has led to increasingly harsh operating envi-

onments. Further improvements in gas turbine performanceill require even higher thermal efficiencies, longer operating

4,5

ifetimes, and reduced emissions. Currently, yttria stabilizedirconia (YSZ), especially zirconia containing 8 wt.% yttriaoatings are still the mostly used TBCs on rotating parts in the

∗ Corresponding author. Tel.: +86 371 62508765.E-mail addresses: zhs761128@163.com,

hsandchen@126.com (Z. Hongsong).

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b

955-2219/$ – see front matter © 2013 Elsevier Ltd. All rights reserved.ttp://dx.doi.org/10.1016/j.jeurceramsoc.2013.07.017

urbine, showing an amazing balance of all required properties.6

owever, the major disadvantage of YSZ is the limited opera-ion temperature of 1473 K for the long-term application. Atigher temperatures, the t’-tetragonal phase transforms into theetragonal and the cubic (t + c) phases. During cooling, the t-hase will further transform into the monoclinic (m) phase,esulting in the volume increase and leading to the formation ofracks in the coating.7 Moreover, the sintering-induced volumehrinkages would degrade the columnar structure of EB-PVDoatings and increase the elasticity modulus and, as a result,educe the favorable strain tolerance of the coating.8 In ordero overcome the shortcomings and meet to the ambitious devel-pment goal of TBCs, the most feasible and economic methods to further reduce the thermal conductivity of new ceramicBCs on the premise of the usage of advanced superalloys andooling technique. This would have a dual benefit of increasedngine efficiency and increased reliability. With this objective,

ecent development efforts have focused toward identifying neweramic materials for TBCs.9

The selection of TBC material is restricted by someasic requirements, such as: (1) high melting point, (2) no

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6 Z. Hongsong et al. / Journal of the Eu

hase transformation between room temperature and operatingemperature, (3) low thermal conductivity, (4) chemical inert-ess, (5) thermal expansion match with the metallic substrate,6) good adherence to the metallic substrate, and (7) low sinter-ng rate of the porous microstructure.10,11 Therefore, the numberf materials that can be used as TBCs is very limited. In recentears, the rare earth oxides with type of Ln2B2O7 (Ln = rarearth elements, B = Zr, Hf, Hf or Ce), have been developed fordvanced turbine engines, which are intended to operate at tem-eratures as high as possible.12 The rare earth zirconates witheneral formula Ln2Zr2O7 with pyrochlore structure or defectuorite-type structure have been regarded as the most poten-

ial ceramic materials for thermal barrier coatings. The thermalonductivities of Ln2Zr2O7 (Ln = La, Nd, Sm, Eu, Gd, Dy, etc.)eramic materials are in the range from 1.1 to 1.2 W/m K, whichre much lower than that of YSZ. Because of their promisinghermophysical properties, efforts have been made to investi-ate the co-doped Ln2Zr2O7 ceramics with one or more metalxides in recent years.13–16 However, their relative low ther-al expansion coefficients (CETs) can result in high thermal

tresses in TBC applications, which is very harmful for TBC’serformance.

The rare earth hafnate with type of Ln2Hf2O7 show typicalyrochlore structure.17 However, their thermophysical proper-ies were mainly studied by numerical simulation, further reportn the thermophysical property of Ln2Hf2O7 series was rela-ively less.18,19 Thermophysical properties reported by Q. Z.ue revealed that thermal conductivities of Ln2Sn2O7 seriesere higher than that of 8YSZ, and their thermal expansion

oefficients were lower than that of 8YSZ.20 These results con-rary to the requirements of thermal barrier coatings.

In recent years, rare earth cerium oxides with type ofn2Ce2O7 (Ln = rare earth elements) have attracted consider-ble attention. H. Dai et al.21 reported that thermal expansionoefficient of Nd2Ce2O7 was higher than that of YSZ and evenore interesting was the thermal expansion coefficient change

s a function of temperature paralleling that of the superalloyond coat. Moreover, the thermal conductivity of Nd2Ce2O7as 30% lower than that of YSZ. X.Q. Cao et al.22 has proposeda2Ce2O7 as a new TBC material that has a low thermal conduc-

ivity and a large thermal expansion coefficient close to that ofhe bond coat. After long-term annealing at 1400 ◦C, La2Ce2O7as still stable without phase transformation. S. J. Patwe

t al.23 reported that thermal expansion coefficient of Gd2Ce2O7as higher than that of Gd2Zr2O7 and YSZ. In our labora-

ory, thermo physical properties of several rare earth ceriumxides, such as Sm2Ce2O7, Yb2Ce2O7, Er2Ce2O7, Dy2Ce2O7,2Ce2O7 and doped-La2Ce2O7, have been investigated.24–28.esults indicated that their thermal conductivities were 25–40%

ower than that of YSZ, and their thermal expansion coeffi-ient were much higher than that of Ln2Zr2O7 and YSZ. Inhese reported rare earth cerium oxides, Sm2Ce2O7 has rela-ively lower thermal conductivity and higher thermal expansion

oefficient. It was claimed that the substitution on Ln site byther cations in Ln2B2O7 ceramics led to a low thermal con-uctivity as contrasted with unitary oxides, which also haveypical influence on thermal expansion coefficient. However,

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n Ceramic Society 34 (2014) 55–61

o data on phase structure and thermophysical properties ofamarium–dysprosium cerium oxides system have been reportedp to now. In the present study, (Sm1−xDyx)2Ce2O7 ceram-cs were prepared by pressureless sintering. The microstructurend thermophysical properties of different (Sm1−xDyx)2Ce2O7eramics were examined.

. Experiment

Samarium oxide, dysprosium oxide and cerium oxide pow-ers (Rare-Chem Hi-Tech Co., Ltd., Huizhou, China, purity99.9%) were chosen as raw materials. These rare-earth oxide

owders were heat-treated at 800 ◦C for 2 h before weighting.Sm1−xDyx)2Ce2O7 (x = 0, 0.05, 0.15) oxides were preparedy solid-state reaction. For each composition, the weightedowders were mixed by planetary milling with zirconia ballsn isopropyl alcohol for 6 h. Subsequently, the dried powder

ixtures were then sieved for granulation and compacted into disk form under uniaxial pressure of 50 MPa followed byold isostatic pressing with 150 MPa. Finally, the bulks wereressureless-sintered at 1600 ◦C for 10 h in air.

The phase structures of sintered bulk materials were char-cterized by X-ray diffraction (XRD, X’Pert PRD MPD Theetherlands) with Cu K� radiation at a scan rate of 4◦/min.canning electron microscopy (SEM, Model Hitachi S-4800,apan) was used to observe the microstructure of bulk ceramics.he specimens were polished with 1 �m diamond paste, and

hen thermally etched at 1500 ◦C for 2 h in air for SEM observa-ions. Qualitative X-ray element analysis of various phases wasarried out using SEM equipped with energy dispersive spec-roscopy (EDS). For heat-treated samples, the actual densitiesere measured by using Archimedes method with an immersionedium of deionized water. The lattice parameters of developed

hases were calculated from the XRD results.The thermal diffusivity (λ) of the sintered samples was mea-

ured using laser-flash method (Model Anter FlashLineTM3000,SA) from 200 ◦C to 1000 ◦C in an argon gas atmosphere.ylindrical disk-shaped sample was about 12.7 mm in diam-ter and about 1.2 mm in thickness. Each sample was groundo that both surfaces were coplanar. In order to avoid any trans-ission of the laser beam through the samples, both the front

nd back faces of the samples were coated with a thin layerf graphite. Appropriate corrections were made in the thermaliffusivity calculations to account for the presence of these lay-rs. The thermal diffusivity measurements were made at 200 ◦Cntervals from 200 ◦C to 1000 ◦C. The thermal diffusivity mea-urement of these samples was carried out three times at eachemperature. The specific heat capacities (Cp) were determineds a function of temperature from the chemical compositions ofSm1−xDyx)2Ce2O7 and the heat capacity data of the constituentlement (O, Dy, Sm and Ce) obtained from reference.29,30 inonjunction with the Neumann–Kopp rule. The thermal con-uctivity (k) of the specimen was calculated by Eq. (1) with

pecific heat capacity (Cp), density (ρ) and thermal diffusivityλ) as follows.

= λ · ρ · CP (1)

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psa(Aeu9niples are very clean, no other inter-phases or un-reacted oxides

Fig. 1. XRD patters of (Sm1−xDyx)2Ce2O7 bulk ceramics.

ecause the sintered specimen was not full dense, the measuredhermal conductivity was modified for the actual value k0 usingq. (2), where φ is the fractional porosity and the coefficient/3 is used to eliminate the effect of porosity on actual thermalonductivity.22

k

k0= 1 − 4

3φ (2)

The linear thermal expansion coefficient (CTE) of theintered samples was determined with a high-temperatureilatometer (Model NeTZSCH DIL 402 C/7, Germany). Theize of sample was approximately 25 mm × 3 mm × 4 mm. Dataor precise calculation of thermal expansion coefficient wereeasured in the temperature of ambient and 1000 ◦C at a heating

ate of 5 ◦C/min in argon atmosphere, and they were correctedsing the known thermal expansion of a certified standard alu-ina.

. Results and discussion

.1. XRD

In this work, the phase structure of these three samples wasnalyzed by XRD at room temperature, as shown in Fig. 1, itan be noted that all patterns are coincident with the standardpectrum of CeO2, and no evidence of pyrochlore phase cane found, which means that pure (Sm1−xDyx)2Ce2O7 ceramicsith defect structure are successfully synthesized. The peaks ofm2Ce2O7 at 2θ = 28.3◦, 32.79◦, 47.22◦, 55.9◦, 58.7◦, 75.9◦ cor-esponding to the (1 1 1), (2 0 0), (2 2 0), (3 1 1), (2 2 2) and (3 1 1)eflections of fluorite structure, respectively. The absence ofiffraction peaks around 36.9◦ and 44.5◦ are always considereds the characteristic of the fluorite structure distinguished fromhe pyrochlore structure. As shown in Fig. 1, with incorporationf smaller Dy3+ cations partially instead of larger Sm3+ cations,

he (Sm0.95Dy0.05)2Ce2O7 and (Sm0.9Dy0.1)2Ce2O7 ceramicslso exhibit a defect-fluorite structure. Also, the lattice param-ters of (Sm1−xDyx)2Ce2O7 (x = 0, 0.05 and 0.15) ceramics

ett

n Ceramic Society 34 (2014) 55–61 57

ere calculated from XRD patterns. The value of Sm2Ce2O7s 0.549 nm which is higher than that of (Sm0.95Dy0.05)2 Ce2O70.5488 nm) and (Sm0.85Dy0.15)2Ce2O7 (0.5482 nm). This factan be attributed to the difference of ionic radii of Sm andy. The decreasing lattice parameter of the (Sm1−xDyx)2Ce2O7

eramics indicate that Dy3+ ions have successfully entered intohe lattice of Sm2Ce2O7. However, there are a few of broadeaks in XRD pattern of (Sm0.85Dy0.15)2Ce2O7, such as peakst 2θ values of about 56.04◦(3 1 1), 58.78◦(2 2 2), 69.11◦(4 0 0),6.26◦(3 3 1), 78.57◦(4 2 0). The broad peaks always indicatesittle grain size of products according to Scherrer’s equation,hich also show that the crystal growth rare at these peaks is slow

ompared to that of peaks at 2θ = 28.27◦(1 1 1) and 32.78◦(2 0 0).his phenomenon maybe attributed to the uneven temperatureeld in sintering furnace, and the concrete reason will be dis-ussed in the future work.

In addition, in the Ln2Ce2O7 system, the phase structure isainly governed by the relative ionic radius ratio of (rA/rB).30–33

he pyrochlore oxides were found to be stable when the ionicadius ratio (rA/rB) lies in range 1.46–1.78.33 The stability ofefect fluorite structure in rare earth cerium at an atmosphericressure is limited to the range of r(Ln3+)/r(Ce4+) < 1.46. Above.78, there is a transition to a monoclinic phase with La2Ti2O7-ype structure.34 The ionic radius of Ce4+ is 0.97 A, however, theonic radius of Sm3+ and Dy3+ are 1.08 A and 1.027 A, respec-ively. The average ionic radius r(Lnaver) at the Ln-sites in theSm1−xDyx)2Ce2O7 system is calculated from the ionic radiusf the component ions and the chemical composition using theollowing equation.21

(Lnaver) = xr(Dy3+) + (1 − x)r(Sm3+) (3)

For the samarium–dysprosium cerium oxides in this inves-igation, the calculated values of (rA/rB) for Sm2Ce2O7,Sm0.95Dy0.05)2Ce2O7 and (Sm0.85Dy0.15)2Ce2O7 are 1.11,.109 and 1.104, respectively, which is clearly lower than 1.46.herefore, these samples are expected to crystallize in the defectuorite structure. This result is coincident with the conclusionrom the XRD analysis.

.2. SEM

The typical microstructure of the synthesized products is dis-layed in Fig. 2. As can be seen from Fig. 2(a–c), the grainize of Sm2Zr2O7 and (Sm0.95Dy0.05)2Ce2O7 is homogeneous,nd some fine grains can be found in the microstructure ofSm0.85Dy0.15)2Ce2O7, which is consistent with XRD result.lthough these products have dense microstructure, some appar-

nt pores are still found. Their relative densities measured bysing Archimedes method in sequence are 98.2%, 96.2% and4.1%, respectively. Fig. 2(d–f) show high magnification scan-ing electron micrograph of (Sm1−xDyx)2Ce2O7 ceramics. Its can be observed that the grain boundaries in these sam-

xist in the interfaces. The chemical compositions of the syn-hesized oxides were determined using EDS. Table 1 showshe results of chemical compositions for (Sm1−xDyx)2Ce2O7

58 Z. Hongsong et al. / Journal of the European Ceramic Society 34 (2014) 55–61

Fig. 2. Microstructure of synthesized oxides (a) Sm2Ce2O7, (b) (Sm0.95Dy0.05)2Ce2O7, (c) (Sm0.85Dy0.15)2Ce2O7, (d) Sm2Ce2O7, (e) (Sm0.95Dy0.05)2Ce2O7, and(f) (Sm0.85Dy0.15)2Ce2O7.

TA

E

S((

able 1tomic percentage of (La1−xDyx)2Ce2O7 ceramics.

lement atomic% Sm Dy Ce O

m2Ce2O7 17.8 0 17.6 64.6Sm0.95Dy0.05)2Ce2O7 15.7 1.2 17.9 65.2Sm0.85Dy0.15)2Ce2O7 13.5 3.2 18.4 64.9

cet

3

(rc

eramics. According to EDS analysis, the mole ratio of differ-nt elements in (Sm1−xDyx)2Ce2O7 ceramics is very close toheir stoichiometery.

.3. Thermal conductivity

The calculated specific heat capacities ofSm1−xDyx)2Ce2O7 ceramics based on Neumann–Koppule at different temperatures are presented in Fig. 3. The spe-ific heat capacities of (Sm1−xDyx)2Ce2O7 ceramics increase

Z. Hongsong et al. / Journal of the European Ceramic Society 34 (2014) 55–61 59

Fw

w1

C

C

fftiihc

Fp

Fc

2dTwelstdtsc

ig. 3. Variation of specific heat capacities of (Sm1−xDyx)2Ce2O7 ceramicsith temperature.

ith the increasing temperature in the range from ambient to000 ◦C, which can be fitted using the following equations.

p(Sm1.9Dy0.1Ce2O7) = 0.37034 + 0.0001 × T

− 56.85978 × T−2 (4)

p(Sm1.7Dy0.3Ce2O7) = 0.36921 + 0.0001 × T

− 56.68542 × T−2 (5)

The thermal diffusivity of (Sm1−xDyx)2Ce2O7 ceramics as aunction of temperature is shown in Fig. 4. The thermal dif-usivity values in Fig. 4 are the arithmetic means of everyhree measurements of identical ceramic materials. The errors derived from the mean standard deviation, and the error bars

n Fig. 4 are smaller than the symbols which are not plotted,ere. Clearly, the thermal diffusivity of (Sm1−xDyx)2Ce2O7eramics decreases rapidly with increasing temperature from

ig. 4. Thermal diffusivities of (Sm1−xDyx)2Ce2O7 ceramics for different tem-eratures.

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ig. 5. Temperature dependence of thermal conductivity of (Sm1−xDyx)2Ce2O7

eramics.

00 ◦C to 1000 ◦C, which suggests a dominant phonon con-uction behavior in most inorganic non-metallic materials.35

hermal diffusivity of (Sm1−xDyx)2Ce2O7 ceramics decreasesith increasing Dy2O3 content at identical temperature lev-

ls. The decrease in thermal diffusivity at identical temperatureevels is assigned to the lattice distortion due to Dy3+ cationsubstitution for Sm3+ cations in the structure. Cations substi-ution causes lattice distortion, increases phonon scattering andecreases phonon mean free path, according to Eq. (6).21 Onhe other hand, the phonon mean free path is proportional to thequare of atomic weight difference between the solute and hostations, according to the following Eq. (7)21:

1

l= 2ca3ω4

πυ4 J2γ2(

R

R

)2

(6)

1

lp= a3

4πv4 ω4c

(R

R

)2

(7)

here, a3 is the volume per atom, υ the transverse wave speed, the phonon frequency, c the concentration per atom, J theonstant, � the Gr uneisen parameter, M and R the average massnd ionic radius of the host atom, respectively, M and Rhe differences of masses and ionic radius between the sub-tituted and the substituting atoms, respectively. Because thetomic weights of Sm and Dy are 150.4 and 162.5, respectively.he effective ionic radius of Sm3+ ion and Dy3+ ion are 1.079 And 1.027 A, the effective phonon scattering by Dy3+ is signifi-antly higher than that of Sm3+, which contributes to the lowerhermal conductivity.

According to Eqs. (1) and (2), the corrected thermal conduc-ivities of (Sm1−xDyx)2Ce2O7 ceramics are plotted in Fig. 5.he values in Fig. 5 are corrected to 100% theoretical den-ity according to Eq. (2). The error bars are omitted in Fig. 5or the reason that they are smaller than the symbols. Fig. 5

hows that the thermal conductivity gradually decrease withncreasing temperature up to 1000 ◦C, which is attributed tohe lattice thermal conduction. From Fig. 5, doping of dyspro-ium oxide clearly reduces thermal conductivities of these rare

60 Z. Hongsong et al. / Journal of the Europea

Fig. 6. Thermal expansion coefficient of (Sm1−xDyx)2Ce2O7 ceramics.

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10. Ramachandran CS, Balasubramanian V, Ananthapadmanabhan PV,Viswabaskaran V. Influence of the intermixed interfacial layers on the ther-

arth cerium oxides over the entire temperature range. The ther-al conductivities of (Sm1−xDyx)2Ce2O7 system in the current

nvestigation were located within the range 1.38–1.69 W/m K at000 ◦C, which are obviously lower than those of fully dense

wt.% YSZ (3.0 at room temperature to 2.3 W/m.K at 700 ◦Ceported by Wu et al.)36 Thus (Sm1−xDyx)2Ce2O7 ceramics cane explored as potential candidates for thermal barrier coatingspplications.

.4. Thermal expansion coefficient

The thermal expansion coefficients of (Sm1−xDyx)2Ce2O7eramics are presented in Fig. 6. It can be seen that the ther-al expansion coefficients of these ceramics increase with

ncreasing temperature. The essence of the thermal expansionoefficients of solid materials increasing is that the averageistance between particles among the lattice increases. Withncreasing of temperature, the crystal lattice vibration of solid

aterials is intensified, which resulted in the increase of thehermal expansion coefficients.

Fig. 6 also indicates that the thermal expansion ofSm1−xDyx)2Ce2O7 ceramics decreased with increasing Dy2O3ontent. The minimum value around 11.49 × 10−6/K waschieved for the sample with x = 0.15 at 1000 ◦C. This result cane attributed to the lower ionic radius of Dy3+ (1.027 A) than thatf Sm3+ (1.079 A), which leads to the crystal lattice vibrationeduced, hence (Sm1−xDyx)2Ce2O7 ceramics have decreasinghermal expansion coefficients with increasing content ofy2O3.16 The average values of (La1−xGdx)2Ce2O7 ceram-

cs in the temperature range between 500 ◦C and 1000 ◦C are1.24 × 10−6/K, 11.1 × 10−6/K and 11.02 × 10−6/K, respec-ively, which are significantly greater than that of YSZ. Theirigher thermal expansion coefficients are beneficial to reducehe residual thermal stresses due to thermal expansion mismatch

t the interface between a top ceramic layer and a metal bondingoating.4

n Ceramic Society 34 (2014) 55–61

. Conclusions

1) The (Sm1−xDyx)2Ce2O7 ceramics with single fluorite struc-ture were successfully prepared by solid reaction methodat 1600 ◦C for 10 h in air using Sm2O3, Dy2O3 and CeO2as raw materials. Their relative densities are higher than94%, and there are no other inter-phases or unreacted oxidesexisted in the boundaries between grains.

2) The thermal conductivities of (Sm1−xDyx)2Ce2O7 ceramicsgradually decrease with raising temperature. Their thermalconductivities are located in the range 1.38–1.69 W/m K at1000 ◦C, which are obviously lower than those of fully dense7 wt.% YSZ.

3) Thermal expansion coefficients of (Sm1−xDyx)2Ce2O7oxides increase with increasing temperature. The averagevalues of (La1−xGdx)2Ce2O7 ceramics in the temperaturerange between 500 ◦C and 1000 ◦C are 11.24 × 10−6/K,11.1 × 10−6/K and 11.02 × 10−6/K, respectively, whichare significantly greater than that of YSZ. At identi-cal temperature levels, thermal expansion coefficient of(Sm1−xDyx)2Ce2O7 decreases with increasing Dy2O3 con-tent.

cknowledgement

Great thanks to the financial support of Doctor Research FundD2007012) in Henan Institute of Engineering.

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