Gadolinium doped cerium (GDC10) thin film ceramics were prepared by evaporating Gd0.1Ce0.9O1.95 powder,using an e-beam evaporation technique. The initial powder of GDC10, with different surface areas (BET: 6.44,36.2, 201 m2/g), were used to prepare thin ceramic films. The GDC10 thin ceramic films (thickness ~ 1.4 μm)were formed on optical quartz substrate (substrate geometry 2.5 × 2.5 cm).The formed GDC10 thin ceramic films structural properties were studied using X-ray diffraction (XRD) andscanning electron microscopy (SEM). The electrical properties of the samples were investigated with an ACimpedance spectroscopy in the temperature range 473–873 K and frequency range 10−1–106 Hz.The formed GDC10 thin films repeat the crystallographic orientation of evaporated powder and do not de-pend on the initial powder grain size and deposition rate. It was determined that crystallite size has an influ-ence on conductivity and diffusion coefficient. The ionic conductivity and activation energy depends oncrystallite size. The best ionic conductivity and the lowest activation energy are 9.00 · 10−1 S/m and1.13 eV, respectively. The activation energy varies from 1.19 to 1.29 eV, and depends on crystallite size, too.
As an alternative to yttria stabilized zirconia (YSZ), which is usedmost commonly for SOFC (solid oxide fuel cells) electrolyte, gadolin-ium or samarium doped ceria (i.e., GDC and SDC) are applied mostwidely. Ceria-based solid solutions have been acknowledged to bethe most promising electrolytes for use in intermediate-temperature(500–600 °C) fuel cells, since their ionic conductivity is much higherthan yttrium stabilized zirconium (YSZ) [1–3]. The ionic conductivityof ceria-based electrolytes, doped with various dopants (e.g., Ca2+,Sr2+, Y3+, La3+, Gd3+, and Sm3+) at different dopant concentrations,has been investigated extensively [4]. The ionic conductivity dependson doped cations and their concentration.
Solid electrolytes, exhibiting high oxygen ion conductivity, are ofspecial interest for their application, not only in electrochemical de-vices, such as solid oxide fuel cells (SOFCs), but also in oxygen sepa-ration membranes, methane gas conversion reactors, etc. [5]. In thelast two decades, there has been a growing interest in the synthesisof solid electrolyte films by physical vapour deposition (PVD),owing to its strict control of film microstructure, porosity, stoichiom-etry and growth rate during the course of deposition [6]. The deposi-tion parameters influence ionic conductivity and are related to thediffusion coefficient and mobility of oxygen ions. Optimizing resis-tance by the reduction of electrolyte thickness is another importantway to lower the operation temperature. Compared to other vapour
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deposition methods, the e-beam technique has the advantage of ahigh deposition rate and large deposition area [7].
In this work, the relationship between electrical and microstruc-ture growth properties of deposited Gd0.1Ce0.9O1.95 (GDC10) thinfilms was studied. GDC10 powders with different surface area (BET)were used to form thin ceramic films, using the e-beam evaporationmethod.
2. Experimental
The initial powder of GDC10, with different BET (fuel cell mate-rials), was used to forming thin ceramic films. The GDC10 thin ceram-ic films (about ~ 1.4 μm of thickness) deposited evaporating differentinitial powder BET (6.44, 36.2, 201 m2/g) by the electron beam phys-ical vapour deposition method. The GDC10 thin films were formed ondielectric optical quartz (SiO2) substrate (geometry 2.5 × 2.5 cm).The deposition rate was changed from 2 to 16 Å/s, to understand itsinfluence on the thin film growth. The influence of particle size of ini-tial powders, the deposition rate on formed thin ceramic film micro-structure, and electrical properties were studied. The substrate wascleaned in an ultrasonic bath (in pure acetone) before being placedin a vacuum chamber, and its surface was cleaned using radio fre-quency (RF) Ar ions plasma for 10 min before deposition.
The formed GDC10 thin ceramic films were analysed by X-ray dif-fraction (XRD, D8 Discover (Bruker AXS GmbH) standard Bragg-Brentano focusing geometry in a 20–70° range, using the Cu Kα1λ = 0.1540562 nm radiation). The crystallite size of GDC10 thin filmswas estimated using Scherrer's equation [8], using TOPAS software.
42 G. Laukaitis, D. Virbukas / Solid State Ionics 247–248 (2013) 41–47
The crystallite size can be determined using the most intensive peak oraverage crystallite size [9]. The average crystallite size (bd>) was esti-mated by equation evaluating intensity of peaks:
dh i ¼Xn1
I hklð ÞXn
1
I hklð Þ
0BBBB@
1CCCCAd hklð Þ ð1Þ
where b d > is average crystallite size, I(hkl) is reflection intensity of(hkl), d(hkl) is crystallite size of (hkl), n is number of reflection peaks.
A scanning electron microscope (SEM, JSM 5600) was used toinvestigate the microstructure of the GDC10 thin films.
Electrical properties were investigated with ProboStat (NorECsAS) in th temperature range 473–873 K with 20 degrees incrementand in the frequency range 10−1–106 Hz. The Pt electrodes (geome-try 1 × 0.5 cm) were formed before the measurement on electrolyte,and the impedance spectrumwas measured in parallel to coating sur-face (Fig. 1).
3. Experimental result and discussion
The XRD diffraction patterns of the formed GDC10 thin ceramicfilms, at different deposition rates on optical quartz substrate, arepresented in Fig. 2. All visible Bragg peaks correspond to the fluoritestructure (according to DIFFRACplus EVA) (Fig. 2). The XRD peaks ofGDC10 thin ceramics films show the same dominating crystallo-graphic orientation (111). Deposition rate and evaporated powderBET have no influence on it, and it is dominant at all deposition rates.
The crystallite size varies from 5.9 to 11.1 nm when evaporatedpowders are with BET 36.2 and 201 m2/g (Table 1). The depositionrate has a minor influence on crystallite size (it varies from 7.8 to9.5 nm) when is evaporated powder with BET = 6.44 m2/g. It couldbe caused that evaporation conditions are changed using the powderswith different BET. This influences the vapour stream, which changesthe proportion of the atoms and clusters. The deposition rate couldalso influence the crystallite size. As the deposition rate is higher theevaporated particles have higher energies and form larger crystallites[10]. The lattice constants are calculated using TOPAS software (thestrains were eliminated). Nevertheless, the SEM measurements showthat GDC10 thin ceramic films were formed dense, and with homoge-neous structure (Fig. 3). The texture coefficient, as a function of deposi-tion rate, was calculated according [11], and is presented in Fig. 4. Thebiggest influence on the dominant texture coefficient has depositionrate when is evaporated nanometer powder with BET = 201 m2/g(Fig. 4c). Such influence on the texture coefficient is theminor evaporat-ing other submicron powders (BET = 6.44 and 36.2 m2/g) (Fig. 4a andb). It could be seen that the initial evaporated powder BET parameteralso influences the dominant texture coefficient. The texture coefficientT111 is dominant when evaporated powders are with BET = 6.44 and201 m2/g. The dominating texture coefficient starts to be T222 whenevaporated powder BET = 36.2 m2/g (Fig. 4b).
Fig. 1. Sample geometry for two electrode measurements.
Fig. 2. XRD patterns of GDC10 thin ceramic deposited on optical quartz (SiO2) usingdifferent initials powders (different BET): a) 6.44 m2/g; b) 36.2 m2/g; and c) 201 m2/g.
Two semicircles could be seen from the impedance spectroscopymeasurements (Fig. 5): the first one corresponds to Rtot (Rtotal =Rb + Rgb), and the second one (low-frequency arc) corresponds tothe electrode contribution. It is possible to see them because the imped-ance measurement was started from 0.1 Hz. The total resistance (Rtot)
Table 1The crystallite size (d, nm) and lattice constant (a, Å) dependence on different deposition rates of GDC10 with different BET. Thin ceramic films were formed on optical quartzsubstrates.
Deposition rate, (Å/s) Crystallographic orientation bd>, nm a, Å
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and relaxation frequency (fR) at different temperatures were obtainedfrom Cole-Cole plots (Fig. 5). The relaxation frequency was receivedfrom a maximum of Im(Z) at different temperatures. As can be seen,the resistance and relaxation frequency depends on BET. The lowest re-laxation frequency and the resistancewere achieved by evaporating ini-tial powder with BET = 6.44 m2/g (Fig. 5).
The ionic conductivity of the formed GDC10 thin films also dependson crystallite size. It increases when crystallite size is increased, whichcould be seen from the Arrhenius plots of GDC10 (BET 6.44 m2/g) thinceramic films, formed at different deposition rates (Fig. 6). The bestionic conductivity (9.00 · 10−1 S/m) was achieved when depositionrate was 4 Å/s (crystallite size is 9.5 nm) and the lowest conductivity(3.17 · 10−1 S/m), when the deposition rate was 16 Å/s (crystallitesize is 7.9 nm) (Table 2). The same effect was found by Swanson at al.[12] for GDC10 and by Kosacki at al. [13] for GDC20. The values ofionic conductivity for the polycrystalline films are lower than those ofthe single crystal, because total ionic conductivity is the sum of thegrain and the grain boundary conductivities. The grain boundary con-ductivity is smaller compared to the grain conductivity. The grain con-ductivity decreases when crystallinity increases [12]. The average ofcrystallite size is similar (Table 1). The brick-layer model (neglectinglattice strain effect) is commonly employed to analyse the electrical be-haviour of ceramics materials. The true conduction path is lower com-pared to the measured geometric area, in which the electric field isapplied owing to the presence of pores. The electrical resistance of thebulk is expressed by R = (fpor)−1 · ρ · L/A, where fpor is a factor of po-rosity, and ρ is the relative resistivity of the bulk [14]. The densities offormed thin ceramic films depend on the formation parameters. Thecrystallites size has influence on the density of grains and grains bound-ary. The density of grains boundary increases and the conductivity de-creases by reducing the size of the crystallites. The conductivity andactivation energy can be influenced not only by the size of the crystal-lites, but also by the lattice strain,whichmakes ionic diffusionmigrationbarrier affecting oxygen vacancies and chemical bond strength, crystal-lites boundary conditions (boundary density and phase), crystalline andamorphous residual phases in the formed thin film [15,16], andthe elemental composition and amounts of the Ce3+ and Ce4+ valencestate in GDC thin films [17]. Ionic conductivity depends on the chosencrystallographic orientation of crystallites. The same effect was foundby W. Araki and Y. Arai [18] for YSZ electrolytes and L. Yan et al. [19]for La0.7Sr0.3MnO3 electrolytes. The crystallite size of [222] crystallo-graphic orientation and ionic conductivity for micrometrical initials
powders are related. The ionic conductivity is higher than the crystal-lites are large.
The activation energy of total ionic conductivity ΔEa was deter-mined from the slope in the log(σ) vs. 103/T. The activation energydepends on the crystallite size, like the conductivity. It can beexplained that the activation energy decreases by reducing the crys-tallite size [13]. Swanson at al. [12] found that activation energy ofpolycrystalline GDC10 is 0.99 eV and 0.85 eV of the single crystal. Inour case, the activation energy of the formed GDC10 thin films variesfrom 1.19 to 1.29 eV (Table 3). It is slightly higher than that achievedby Swanson et al. [12]. This can be explained by the higher density ofgrain boundary; i.e. increasing the grain boundary density decreasesconductivity and increases activation energy.
The activation energy varies at different deposition rates, and de-creases by reducing crystallite size when using initial powders with(BET 36.2, 201 m2/g). The inverse relationship is obtained when usinginitial (BET 6.44 m2/g). The oxygen ions conductivity in the formed elec-trolyte layers do not only depend on activation energy. The ionic conduc-tivity of GDC10 thin ceramic films is related to the diffusion coefficientand mobility of oxygen ions. The diffusion coefficient of oxygen vacan-cies D is related to the angular frequency ωR by the expression [20–22]:
D ¼ 16γd2ωR; ð2Þ
where D is the oxygen diffusion coefficient, d2 is the mean square jumpdistance between the anion sites in the lattice, γ is correlation factor, andωR angular frequency related to the relaxation frequency fR. For fluoritetape structure, it was derived that d2γ = 0.35a2, where a is lattice con-stant [20–22].
The total ionic conductivity of oxygen ions depends on the mobil-ity of vacancies (μ), the total concentration of Gd cations (N), the va-lence of the oxygen vacancies (z), and the electrical charge (e) [21]:
σ tot ¼ μ⋅N⋅z⋅e; ð3Þ
The diffusion coefficient of oxygen vacancies is proportional to themobility of charge carriers, and can be evaluated form the Nernst–Einstein equation [20]:
D ¼ μkT=ze ¼ kTσ=Nz2e2; ð4Þ
where k is the Boltzmann constant, and T is the absolute temperature.
Fig. 3. SEM images of: a) the surface view; and b) the cross-section of GDC10 thin films formed on optical quartz (SiO2), at 16 Å/s deposition rate, evaporating GDC10 powder withBET = 36.2 m2/g.
44 G. Laukaitis, D. Virbukas / Solid State Ionics 247–248 (2013) 41–47
The fraction of free vacancies tends to increase with temperatureand dopant concentration. The aliovalent cations in the host lattice in-troduce oxygen vacancies, like, for example, the addition of gadoliniumto ceria. This type of process can be illustrated by the following defectequation in Kröger–Vink notation [23]:
Gd2O3⇒2Gd0Ce þ 3OX
O þ V ••O: ð5Þ
Using (2) and (3) equations, the charge carrier concentration (N)can be determined:
N ¼ kTσ=Dz2e2 ð6Þ
The evaluated relaxation frequency (fR), ionic conductivity (σ),diffusion coefficient (D), charge carrier concentration (N) and mobil-ity of oxygen vacancies (μ) of GDC10 thin ceramic films, formed at
Fig. 4. The texture coefficient as a function of deposition rate of formed GDC10 thin filmsby evaporating different powders with different BET: a) 6.44 m2/g; b) 36.2 m2/g; andc) 201 m2/g.
Fig. 5. Cole-Cole plots (measured at 673 K) of formed GDC10 thin films formed by evap-orating powders with different BET at constant deposition rate 2 Å/s (fR – relaxationfrequency).
Fig. 6. Arrhenius plots of GDC10 (BET 6.44 m2/g) thin ceramic films deposited at differ-ent deposition rates.
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different deposition rate, are presented in Table 4. The diffusioncoefficient related to crystallite size and dominating texture coeffi-cient can be seen (Table 1 and Fig. 4). The diffusion coefficient de-creases (at 673 K temperature) when are evaporated submicron
initial powders with BET = 6.44 and 36.2 m2/g. The lowest diffusioncoefficient is 9.12 × 10−15 m2/s, when crystallite size is ~7 nm, andthe highest is 2.72 × 10−14 m2/s when the crystallite size is 9.5 nm.This tendency changes when nanopowder is evaporated (BET =201 m2/g). The highest diffusion coefficient (3.27 × 10−14 m2/s)was achieved when crystallite size was the lowest (7.4 nm). Thiscould mean that the GDC10 thin films (formed evaporatingnanopowder material, BET = 201 m2/g) texture coefficient decreasesto 0.6 (Fig. 4c). Also, it could be influenced by the dominant crystallo-graphic orientation of the crystallites in the formed thin films [24].The diffusion coefficient also relates to conductivity. The oxygen va-cancies occur through grain and grain boundaries. The increasedgrain boundary density decreases ionic conductivity and the diffusioncoefficient (Fig. 7).
The dependence σtot(T) of GDC10 thin ceramic films is caused bythe temperature dependency of oxygen vacancy mobility. The oxygenvacancy mobility depends on temperature and the technological pa-rameters of deposition thin films. The oxygen vacancy mobility in-creases enhancing deposition rate (Fig. 8). The increased depositionrate could increase the energy of evaporated atoms and atom clusters,which could mean that thin films grow with a lower amount ofdefects.
Table 2Total ionic conductivity (σtot, S/m) of GDC10 thin ceramic films formed by evaporating powders with different BET at different deposition rates. The measurements were done at873 K.
46 G. Laukaitis, D. Virbukas / Solid State Ionics 247–248 (2013) 41–47
4. Conclusions
The deposited GDC10 thin films have the same dominant crystallo-graphic orientation (111) and repeat crystallographic orientation as
Table 4The relaxation frequency (fR), ionic conductivity (σ), diffusion coefficient (D), chargecarrier concentration (N) and mobility of oxygen vacancies (μ) of GDC10 thin ceramicfilms, formed at different deposition rates, at 673 K.
Deposition rate, Å/s f, Hz σ, S/m D, m2/s N, m−3 μ, m2/V · s
evaporated powder. The initial powder's BET or deposition rate has noinfluence on it. The crystallite size of formed GDC10 thin films variesfrom 5.8 to 11.1 nm evaporating powders with BET 36.2 and 201 m2/g.The deposition rate has the minor influence on crystallite size (it variesfrom 7.7 to 9.4 nm) using powder with BET = 6.44 m2/g. The deposi-tion rate has the biggest influence on the dominant texture coefficientwhen is evaporated nanometer powder with BET = 201 m2/g. Such in-fluence on the texture coefficient is the minor using other submicronpowders (BET = 6.44 and 36.2 m2/g). The texture coefficient T111 isdominant when are evaporated powders with BET = 6.44 and201 m2/g, but the dominating texture coefficient starts to be T222when is used powder with BET = 36.2 m2/g.
Fig. 7. Diffusion coefficient dependence on temperature of formed GDC10 (BET =6.44 m2/g) thin ceramic films: a) at different deposition rate; and b) by evaporatinginitial powder with different BET.
Fig. 8. The oxygen vacancy mobility dependence on temperature of formed thin ce-ramic films at different deposition rates, by evaporating initial powder with differentBET: a) 6.44 m2/g; b) 36.2 m2/g; and c) 201 m2/g.
47G. Laukaitis, D. Virbukas / Solid State Ionics 247–248 (2013) 41–47
The ionic conductivity and activation energy depends on the crystal-lite size, i.e. as crystallite size decreases ionic conductivity and activation
energy decreases. The ionic conductivity and activation energy are9.00 · 10−1 S/m and 1.13 eV, respectively, at lower crystallite size.The activation energy varies at different deposition rates, and decreaseslessening crystallite size when used initial powders (BET = 36.2 m2/g,201 m2/g). The inverse relationship was obtained when initial powderwas used (BET = 6.44 m2/g). The ionic conductivity and relaxation fre-quency depends on evaporated initial powder's BET. The lowest relaxa-tion frequency and the highest ionic conductivity were achieved byevaporating initial powder with BET = 6.44 m2/g. The diffusion coeffi-cient decreased the reducing crystallite size when submicron initialpowders were evaporated with BET = 6.44 m2/g and 36.2 m2/g. Thelowest diffusion coefficient, 9.12 × 10−15 m2/s, was achieved whencrystallite size was ~7 nm, with the highest 2.72 × 10−14 m2/s whencrystallite size was 9.5 nm. This tendency changed when nanopowderwas evaporated (BET = 201 m2/g); the highest diffusion coefficient(3.27 × 10−14 m2/s) was achieved when the crystallite size was(7.4 nm). The charge carrier concentration depends on the depositionrate and evaporated initials powders BET.
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