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Oxygen vacancy stabilized zirconia (OVSZ); a joint experimental and theoretical study
Scripta Materialia 124 (2016) 26–29
Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Regular Article
Oxygen vacancy stabilized zirconia (OVSZ); a joint experimental and theoretical study Mohsin Raza a,⁎, David Cornil b, Jérôme Cornil b, Stéphane Lucas c, Rony Snyders a,d, Stéphanos Konstantinidis a a
Chimie des Interactions Plasma-Surface (ChIPS), University of Mons, 23 Place du Parc, 7000 Mons, Belgium Service de Chimie des Matériaux Nouveaux, University of Mons, 23 Place du Parc, 7000 Mons, Belgium c Research center for the Physics of Matter and Radiation (PMR-LARN), University of Namur, B-5000 Namur, Belgium d Materia Nova Research Center, Parc Initialis, B-7000 Mons, Belgium b
a r t i c l e
i n f o
Article history: Received 3 May 2016 Received in revised form 16 June 2016 Accepted 20 June 2016 Available online 5 July 2016 Keywords: DFT calculations Zirconia (ZrO2) Oxygen vacancy Phase transformation Phase stabilization
a b s t r a c t Understanding the phase formation in zirconia (ZrO2) has triggered a great debate over the last couple of decades, with several mechanisms proposed so far. In the present letter, we demonstrate by well-optimized experimental measurements supported by Density Functional Theory (DFT) calculations that only O vacancies allow for the stabilization of the cubic (c) phase at room temperature. These vacancies distort the zirconia lattice, forcing the crystal to arrange itself in a high symmetric c structure. © 2016 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
The specific crystal structure of any material is a decisive factor for controlling its properties [1,2] and leads sometime to the discovery of new functionalities [3,4]. In this respect, zirconia (ZrO2) is a material which exists in three crystallographic phases under atmospheric pressure: (i) the monoclinic phase (m, space group P21/c) stable up to ~ 1205 °C; (ii) the tetragonal phase (t, space group P42/nmc) appears from ~ 1205 °C to 2377 °C; and finally (iii) the cubic phase (c, space group Fm-3m) from 2377 °C to 2710 °C (melting temperature) [5]. Because of ZrO2 superior chemical stability, high hardness [6], high dielectric constant [7] and prominent optical properties [8], ZrO2 films have been exploited for a broad range of applications e.g., medical applications [9,10], wear resistant coatings [11], oxygen detectors [12,13] and for thermal barrier coatings (TBC) [14,15]. However, for pure zirconia, it is not possible to exploit most of the above mentioned applications as this is restricted by the change in volume of the zirconia-based components (~5 vol.%) due to the phase transformation upon heating and cooling of the device, which ultimately leads to the deterioration of the device components [16,17]. Therefore, the stabilization of the high temperature c-phase at room temperature is of paramount importance. This has been achieved for decades by doping of cations of lower ⁎ Corresponding author. E-mail addresses: [email protected] (M. Raza), [email protected] (S. Konstantinidis).
http://dx.doi.org/10.1016/j.scriptamat.2016.06.025 1359-6462/© 2016 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
valence than Zr in the zirconia lattice (e.g. Y or Mg) [18,19]. By adding around 12 mol% of yttria (Y2O3), the c-phase of zirconia is found to be stabilized at room temperature and is known as yttria-stabilized zirconia (YSZ) [19]. In YSZ, zirconium (Zr4+) is replaced by yttrium (Y3+) so that to maintain the charge neutrality, for two substituting yttrium cations, one oxygen vacancy is created. This makes YSZ not only useful for TBC but also as an electrolyte in solid oxide fuel cells (SOFC) [20– 22] and in oxygen sensors [13] because of its very good ionic conductivity [23]. However, it has also been found that the doping by aliovalent cations leads to the perturbation of the periodic potential of the oxideion array, which results in higher energy barrier for O2– ions during their diffusion to a vacant site in the solid as compared to intrinsic vacancy-doped oxides [24]. Therefore, to stabilize high temperature cphases of zirconia at room temperature without any doping of yttria, an intense research has been developed during the last one and a half decade using various synthesis techniques. The stabilization procedure has been related to the grain size, energy input during growth, stresses in the film and O vacancies/N atom incorporation in the zirconia lattice [16,25–34]. However, a consensus over what drives the phase formation in zirconia has not been reached so far. In this letter, we demonstrate that O vacancies are the sole responsible for the stabilization of the high temperature c-phase of zirconia at room temperature. To achieve this, we coupled cold plasma-based reactive magnetron sputtering experiments to Ab Initio Density Functional
M. Raza et al. / Scripta Materialia 124 (2016) 26–29
Fig. 1. Influence of O vacancies on the phase constitution. Calculations were performed at the DFT level using SIESTA code [50] including periodic boundary conditions. The functional used was the PBE [51] version initially developed by Perdew, Burke and Ernzerhof with a double zeta basis plus polarization orbitals (DZP). The optimization of the structure and the calculation of the electronic structure were performed using a mesh cut-off of 190 Ry and a (2x2x2) k-point mesh. Calculations were performed at 0 K using 96 atoms in the unit cell (32 Zr and 64 O) and O vacancies were inserted randomly in the zirconia lattice by removing O atoms.
Theory (DFT) calculations. The impact of randomly introduced O vacancies on the energetics of ZrO2-x cell in the t- and c-phases is shown in Fig. 1 based on the DFT calculations (details of the DFT calculations can be found in supplementary information). The later reveal that, for N3 at.% of O vacancies, the c-phase is thermodynamically the most stable. Similar theoretical results have also been reported, via a self-consistent tight-binding model, by Fabris et al. [35]. In their publication, these authors suggest that the stabilization of the t- and c-phase of zirconia can be achieved (only in theory) solely by incorporating O vacancies in the zirconia lattice. To overcome the limitation placed by Fabris et al. [35] and to assess experimentally how the incorporation of O vacancies influence the zirconia phase formation/stabilization, we designed experiments using dc reactive magnetron sputtering (dc-RMS) in such a way that we were not only able to incorporate O vacancies in the zirconia lattice as the
27
film grows, but also to vary their concentration. For this purpose, a voltage feedback control unit [36–39], an auxiliary device (Speedflo mini from Gencoa UK) which provides a rapid control over oxygen partial pressure using target voltage as a feedback signal, was used to work inside the so-called metal-to-compound transition zone [40] of the sputtered zirconium target. The transition zone is an experimental working window in between the metallic and oxidized modes of the sputtered target where under-stoichiometric films are grown, as shown in Fig. 2(a). Using this well-optimized synthesis setup, 100 ± 10 nm thick ZrO2-x films were deposited on Si (100) single crystal substrates at 50, 65, 75, 80, and at 100% signal set point value of the feedback unit (0% represents the metallic mode setpoint while 100% represents the peak of the transition zone) as well as in the oxidized mode in a high vacuum chamber with a base pressure b 4 × 10−4 Pa. During the depositions, non-intentionally heated substrates were placed at a distance of 6.5 cm from a 5 cm in diameter purity (99.97%) Zr target. The sputter target (i.e., the cathode of the system) was fed by 200 mA of discharge current by an Advanced Energy MDX 500 dc power supply equipped with an arc suppressor (Sparkle from Advanced Energy). In order to reach a precise control over the sputter target and therefore film-chemistry, 18O2 (purity 97.1%) was injected at the target during the film growth and was regulated by the voltage feedback device whereas Ar was injected away from the target by using a conventional mass flow controller. During the depositions, working pressure was kept constant at 1.33 Pa in each case. The chemical composition of the films deposited in such conditions is shown in Fig. 2(b). It is observed that the films deposited inside the transition zone at 50, 65, 75, 80 and 100% signal are indeed under-stoichiometric and thus contain 32, 20, 16, 6, and 3 at.% of O vacancies, respectively. In such under-stoichiometric zirconia, the local Zr charge state may vary from Zr+ to Zr4+ depending on the surrounding atoms and density of vacancies [41–45]. Interestingly, the quantum-chemical calculations show that the net charge on zirconium cation varies from 2.54 | e| for ZrO2 down to 1.87 | e | for 15% of O vacancy. These values are in good agreement with those reported in [42], which highlight a variation from 2.57 | e| for ZrO2 down to 2.02 | e| for Zr2O3. According to the DFT calculations presented in Fig. 1, such amount of O vacancies should induce the formation of the c-phase. In contrast, the film deposited in the oxidized mode is stoichiometric (ZrO2). The observed decrease in O vacancy concentration with the increase in signal (i.e., 18O2 partial pressure) is due to the increased oxide compound formation on the target surface.
Fig. 2. (a) Target voltage curve of Zr target as a function of O2 flow shows the transition zone and working points inside the transition zone where the ZrO2-x films were deposited. (b) The concentration of Zr in the samples was probed by RBS. The incident energy of the alpha particles was 2 MeV, and the beam impinged the sample surface at normal incidence. Backscattered particles were collected at 165° in a PIPS detector. Spectra were analyzed with the SIMNRA software assuming Rutherford backscattering cross-section. Specific 18O depth profiles were determined using the resonant nuclear reaction 18O(p,α)15N at 151 keV. Samples were tilted at 30° with respect to the incident beam and the alpha particles were collected in large area PIPS detectors facing the sample surface and parallel to it. The incident energy was varied from 145 to 200 keV. Depth profiles were deconvoluted in order to take into account the energy straggling of the beam and then quantified with the help of a Si18O2 standard produced by thermal oxidation in a pure 18O atmosphere.
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M. Raza et al. / Scripta Materialia 124 (2016) 26–29
Film crystallinity was checked by grazing incidence X-ray diffraction (GIXRD) and the resulting diffractograms were compared to ICDD PDF cards representing patterns of three polymorphs of ZrO2: monoclinic (PDF# 37-1484), tetragonal (PDF# 81-1544) and cubic (PDF# 491642). The peak positions of the films deposited inside the transition zone matched very well those corresponding to the tetragonal and/or cubic phases and were hard to distinguish. To unambiguously identify the phase, we compared the theoretical diffractograms of the t- and cphases (obtained from the DFT optimized structure) with the diffractogram of one of our film grown at 100% sensor signal i.e. film with the lowest concentration of O vacancies i.e. 3 at.%, as shown in Fig. 3. Furthermore, Lamas et al. [46] has pointed out the splitting of the (400) peak of the c-phase of ZrO2 into the (004) and (400) peaks in the t-phase of ZrO2, with more than one degree of separation in their XRD spectra. In Fig. 3, the same peak splitting can be clearly seen for the theoretical diffractogram of the t-phase while we did not observe any splitting of the (400) peak of the film grown at 100% sensor signal, thus confirming that the oxygen vacancy-doped films (ie. deposited inside the transition zone) belong to the cubic zirconia. The diffractograms of the DFT structures obtained for the various percentages considered in O vacancies are shown in Fig. 4(a) together with the diffractograms of the as-deposited films recorded using an incidence angle of 0.5° shown in Fig. 4(b). The remarkable agreement between theory and experiments demonstrates that the oxygen-vacancy doped films (synthesized inside the transition zone) are of pure cphase while the film changes to phase pure monoclinic when deposited in the oxidized mode, i.e. when containing no O vacancies. However, in DFT based diffractograms Fig. 4(a), it is observed that some new peaks start to appear with the increase in O vacancies. The later could be an artifact coming from the large number of O vacancies in the cell. On the other hand, it is also possible that these peaks are not detected in the experimental diffractograms because of the low signal-to-noise ratio. We also found a very good agreement of our DFT calculations for the transition enthalpy of ZrO2 from m to c phase, which is 14.45 kJ/mol in our case and 14.26 kJ/mol measured by X. Luo et al. [47]. It has been also reported that incorporating O vacancies induces lattice distortions and provoke a contraction of the zirconia lattice, which may result in similar zirconia lattice as in cubic zirconia due to the coulomb forces between O vacancy-Zr and O vacancy-O atoms [48,49]. Furthermore, Fabris et al. [35] have also shown in their work that having a low content of O vacancies, i.e. 1 at.% (= 3.2 mol% Y2O3), leads to the tetragonal distortion,
Fig. 3. Comparison of the diffractograms of the cubic and tetragonal phases obtained at the DFT level with the experimental diffractogram of the film grown at 100% sensor signal inside the transition zone. The inset data shows that there is no splitting of the experimental c(400) peak, as it is observed for the tetragonal (004) and (400) peak.
Fig. 4. GIXRD spectra of the (a) structures resulting from the DFT calculations and (b) deposited films, as a function of concentration of O vacancies. Crystallinity of the deposited films was analyzed by GIXRD using PANalytical Empyrean with a Cu-Kα radiation source. The diffractograms were recorded with a step size of 0.07° using an incidence angle of 0.5° at 40 mA, 45 kV of generator settings.
while having a high content of O vacancies, i.e. 4 at.% (= 14.4 mol% Y2O3) resulting in every oxygen atom to be a neighbor of a vacant site or at least four of its six neighboring oxygen atoms, leads to the cubic structure. For these reasons, we believe that the only possible atomistic mechanism behind stabilization of c-phase of zirconia is lattice distortion caused by O vacancy incorporation in the zirconia lattice, ultimately forcing the O and Zr atoms to arrange themselves in the high symmetry cubic crystal. In conclusion, based on our DFT calculations and their remarkable agreement with our experimental data, we conclude that incorporating (as low as 3 at.%) O vacancies is the sole mechanism responsible of promoting zirconia to the high-temperature c-phase, at room temperature. Moreover, development of such stable phase pure cubic zirconia at room temperature by implementing such an optimized and controlled method might not only provide an alternate of YSZ but could also be of great fundamental and technical interest for other oxides (e.g. HfO2, TiO2 and Bi2O3). This work is supported by the Belgian Government through the “Pôle d'Attraction Inter universitaire” (PAI, “Plasma-Surface Interaction”, Ψ). S. Konstantinidis is a Research Associate and J. Cornil a Research Director of the FNRS (Fonds National de la Recherche Scientifique), Belgium.
M. Raza et al. / Scripta Materialia 124 (2016) 26–29
Appendix A. Supplementary data Supplementary data to this article can be found online at http://dx. doi.org/10.1016/j.scriptamat.2016.06.025. References
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Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Regular Article
Oxygen vacancy stabilized zirconia (OVSZ); a joint experimental and theoretical study Mohsin Raza a,⁎, David Cornil b, Jérôme Cornil b, Stéphane Lucas c, Rony Snyders a,d, Stéphanos Konstantinidis a a
Chimie des Interactions Plasma-Surface (ChIPS), University of Mons, 23 Place du Parc, 7000 Mons, Belgium Service de Chimie des Matériaux Nouveaux, University of Mons, 23 Place du Parc, 7000 Mons, Belgium c Research center for the Physics of Matter and Radiation (PMR-LARN), University of Namur, B-5000 Namur, Belgium d Materia Nova Research Center, Parc Initialis, B-7000 Mons, Belgium b
a r t i c l e
i n f o
Article history: Received 3 May 2016 Received in revised form 16 June 2016 Accepted 20 June 2016 Available online 5 July 2016 Keywords: DFT calculations Zirconia (ZrO2) Oxygen vacancy Phase transformation Phase stabilization
a b s t r a c t Understanding the phase formation in zirconia (ZrO2) has triggered a great debate over the last couple of decades, with several mechanisms proposed so far. In the present letter, we demonstrate by well-optimized experimental measurements supported by Density Functional Theory (DFT) calculations that only O vacancies allow for the stabilization of the cubic (c) phase at room temperature. These vacancies distort the zirconia lattice, forcing the crystal to arrange itself in a high symmetric c structure. © 2016 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
The specific crystal structure of any material is a decisive factor for controlling its properties [1,2] and leads sometime to the discovery of new functionalities [3,4]. In this respect, zirconia (ZrO2) is a material which exists in three crystallographic phases under atmospheric pressure: (i) the monoclinic phase (m, space group P21/c) stable up to ~ 1205 °C; (ii) the tetragonal phase (t, space group P42/nmc) appears from ~ 1205 °C to 2377 °C; and finally (iii) the cubic phase (c, space group Fm-3m) from 2377 °C to 2710 °C (melting temperature) [5]. Because of ZrO2 superior chemical stability, high hardness [6], high dielectric constant [7] and prominent optical properties [8], ZrO2 films have been exploited for a broad range of applications e.g., medical applications [9,10], wear resistant coatings [11], oxygen detectors [12,13] and for thermal barrier coatings (TBC) [14,15]. However, for pure zirconia, it is not possible to exploit most of the above mentioned applications as this is restricted by the change in volume of the zirconia-based components (~5 vol.%) due to the phase transformation upon heating and cooling of the device, which ultimately leads to the deterioration of the device components [16,17]. Therefore, the stabilization of the high temperature c-phase at room temperature is of paramount importance. This has been achieved for decades by doping of cations of lower ⁎ Corresponding author. E-mail addresses: [email protected] (M. Raza), [email protected] (S. Konstantinidis).
http://dx.doi.org/10.1016/j.scriptamat.2016.06.025 1359-6462/© 2016 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
valence than Zr in the zirconia lattice (e.g. Y or Mg) [18,19]. By adding around 12 mol% of yttria (Y2O3), the c-phase of zirconia is found to be stabilized at room temperature and is known as yttria-stabilized zirconia (YSZ) [19]. In YSZ, zirconium (Zr4+) is replaced by yttrium (Y3+) so that to maintain the charge neutrality, for two substituting yttrium cations, one oxygen vacancy is created. This makes YSZ not only useful for TBC but also as an electrolyte in solid oxide fuel cells (SOFC) [20– 22] and in oxygen sensors [13] because of its very good ionic conductivity [23]. However, it has also been found that the doping by aliovalent cations leads to the perturbation of the periodic potential of the oxideion array, which results in higher energy barrier for O2– ions during their diffusion to a vacant site in the solid as compared to intrinsic vacancy-doped oxides [24]. Therefore, to stabilize high temperature cphases of zirconia at room temperature without any doping of yttria, an intense research has been developed during the last one and a half decade using various synthesis techniques. The stabilization procedure has been related to the grain size, energy input during growth, stresses in the film and O vacancies/N atom incorporation in the zirconia lattice [16,25–34]. However, a consensus over what drives the phase formation in zirconia has not been reached so far. In this letter, we demonstrate that O vacancies are the sole responsible for the stabilization of the high temperature c-phase of zirconia at room temperature. To achieve this, we coupled cold plasma-based reactive magnetron sputtering experiments to Ab Initio Density Functional
M. Raza et al. / Scripta Materialia 124 (2016) 26–29
Fig. 1. Influence of O vacancies on the phase constitution. Calculations were performed at the DFT level using SIESTA code [50] including periodic boundary conditions. The functional used was the PBE [51] version initially developed by Perdew, Burke and Ernzerhof with a double zeta basis plus polarization orbitals (DZP). The optimization of the structure and the calculation of the electronic structure were performed using a mesh cut-off of 190 Ry and a (2x2x2) k-point mesh. Calculations were performed at 0 K using 96 atoms in the unit cell (32 Zr and 64 O) and O vacancies were inserted randomly in the zirconia lattice by removing O atoms.
Theory (DFT) calculations. The impact of randomly introduced O vacancies on the energetics of ZrO2-x cell in the t- and c-phases is shown in Fig. 1 based on the DFT calculations (details of the DFT calculations can be found in supplementary information). The later reveal that, for N3 at.% of O vacancies, the c-phase is thermodynamically the most stable. Similar theoretical results have also been reported, via a self-consistent tight-binding model, by Fabris et al. [35]. In their publication, these authors suggest that the stabilization of the t- and c-phase of zirconia can be achieved (only in theory) solely by incorporating O vacancies in the zirconia lattice. To overcome the limitation placed by Fabris et al. [35] and to assess experimentally how the incorporation of O vacancies influence the zirconia phase formation/stabilization, we designed experiments using dc reactive magnetron sputtering (dc-RMS) in such a way that we were not only able to incorporate O vacancies in the zirconia lattice as the
27
film grows, but also to vary their concentration. For this purpose, a voltage feedback control unit [36–39], an auxiliary device (Speedflo mini from Gencoa UK) which provides a rapid control over oxygen partial pressure using target voltage as a feedback signal, was used to work inside the so-called metal-to-compound transition zone [40] of the sputtered zirconium target. The transition zone is an experimental working window in between the metallic and oxidized modes of the sputtered target where under-stoichiometric films are grown, as shown in Fig. 2(a). Using this well-optimized synthesis setup, 100 ± 10 nm thick ZrO2-x films were deposited on Si (100) single crystal substrates at 50, 65, 75, 80, and at 100% signal set point value of the feedback unit (0% represents the metallic mode setpoint while 100% represents the peak of the transition zone) as well as in the oxidized mode in a high vacuum chamber with a base pressure b 4 × 10−4 Pa. During the depositions, non-intentionally heated substrates were placed at a distance of 6.5 cm from a 5 cm in diameter purity (99.97%) Zr target. The sputter target (i.e., the cathode of the system) was fed by 200 mA of discharge current by an Advanced Energy MDX 500 dc power supply equipped with an arc suppressor (Sparkle from Advanced Energy). In order to reach a precise control over the sputter target and therefore film-chemistry, 18O2 (purity 97.1%) was injected at the target during the film growth and was regulated by the voltage feedback device whereas Ar was injected away from the target by using a conventional mass flow controller. During the depositions, working pressure was kept constant at 1.33 Pa in each case. The chemical composition of the films deposited in such conditions is shown in Fig. 2(b). It is observed that the films deposited inside the transition zone at 50, 65, 75, 80 and 100% signal are indeed under-stoichiometric and thus contain 32, 20, 16, 6, and 3 at.% of O vacancies, respectively. In such under-stoichiometric zirconia, the local Zr charge state may vary from Zr+ to Zr4+ depending on the surrounding atoms and density of vacancies [41–45]. Interestingly, the quantum-chemical calculations show that the net charge on zirconium cation varies from 2.54 | e| for ZrO2 down to 1.87 | e | for 15% of O vacancy. These values are in good agreement with those reported in [42], which highlight a variation from 2.57 | e| for ZrO2 down to 2.02 | e| for Zr2O3. According to the DFT calculations presented in Fig. 1, such amount of O vacancies should induce the formation of the c-phase. In contrast, the film deposited in the oxidized mode is stoichiometric (ZrO2). The observed decrease in O vacancy concentration with the increase in signal (i.e., 18O2 partial pressure) is due to the increased oxide compound formation on the target surface.
Fig. 2. (a) Target voltage curve of Zr target as a function of O2 flow shows the transition zone and working points inside the transition zone where the ZrO2-x films were deposited. (b) The concentration of Zr in the samples was probed by RBS. The incident energy of the alpha particles was 2 MeV, and the beam impinged the sample surface at normal incidence. Backscattered particles were collected at 165° in a PIPS detector. Spectra were analyzed with the SIMNRA software assuming Rutherford backscattering cross-section. Specific 18O depth profiles were determined using the resonant nuclear reaction 18O(p,α)15N at 151 keV. Samples were tilted at 30° with respect to the incident beam and the alpha particles were collected in large area PIPS detectors facing the sample surface and parallel to it. The incident energy was varied from 145 to 200 keV. Depth profiles were deconvoluted in order to take into account the energy straggling of the beam and then quantified with the help of a Si18O2 standard produced by thermal oxidation in a pure 18O atmosphere.
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Film crystallinity was checked by grazing incidence X-ray diffraction (GIXRD) and the resulting diffractograms were compared to ICDD PDF cards representing patterns of three polymorphs of ZrO2: monoclinic (PDF# 37-1484), tetragonal (PDF# 81-1544) and cubic (PDF# 491642). The peak positions of the films deposited inside the transition zone matched very well those corresponding to the tetragonal and/or cubic phases and were hard to distinguish. To unambiguously identify the phase, we compared the theoretical diffractograms of the t- and cphases (obtained from the DFT optimized structure) with the diffractogram of one of our film grown at 100% sensor signal i.e. film with the lowest concentration of O vacancies i.e. 3 at.%, as shown in Fig. 3. Furthermore, Lamas et al. [46] has pointed out the splitting of the (400) peak of the c-phase of ZrO2 into the (004) and (400) peaks in the t-phase of ZrO2, with more than one degree of separation in their XRD spectra. In Fig. 3, the same peak splitting can be clearly seen for the theoretical diffractogram of the t-phase while we did not observe any splitting of the (400) peak of the film grown at 100% sensor signal, thus confirming that the oxygen vacancy-doped films (ie. deposited inside the transition zone) belong to the cubic zirconia. The diffractograms of the DFT structures obtained for the various percentages considered in O vacancies are shown in Fig. 4(a) together with the diffractograms of the as-deposited films recorded using an incidence angle of 0.5° shown in Fig. 4(b). The remarkable agreement between theory and experiments demonstrates that the oxygen-vacancy doped films (synthesized inside the transition zone) are of pure cphase while the film changes to phase pure monoclinic when deposited in the oxidized mode, i.e. when containing no O vacancies. However, in DFT based diffractograms Fig. 4(a), it is observed that some new peaks start to appear with the increase in O vacancies. The later could be an artifact coming from the large number of O vacancies in the cell. On the other hand, it is also possible that these peaks are not detected in the experimental diffractograms because of the low signal-to-noise ratio. We also found a very good agreement of our DFT calculations for the transition enthalpy of ZrO2 from m to c phase, which is 14.45 kJ/mol in our case and 14.26 kJ/mol measured by X. Luo et al. [47]. It has been also reported that incorporating O vacancies induces lattice distortions and provoke a contraction of the zirconia lattice, which may result in similar zirconia lattice as in cubic zirconia due to the coulomb forces between O vacancy-Zr and O vacancy-O atoms [48,49]. Furthermore, Fabris et al. [35] have also shown in their work that having a low content of O vacancies, i.e. 1 at.% (= 3.2 mol% Y2O3), leads to the tetragonal distortion,
Fig. 3. Comparison of the diffractograms of the cubic and tetragonal phases obtained at the DFT level with the experimental diffractogram of the film grown at 100% sensor signal inside the transition zone. The inset data shows that there is no splitting of the experimental c(400) peak, as it is observed for the tetragonal (004) and (400) peak.
Fig. 4. GIXRD spectra of the (a) structures resulting from the DFT calculations and (b) deposited films, as a function of concentration of O vacancies. Crystallinity of the deposited films was analyzed by GIXRD using PANalytical Empyrean with a Cu-Kα radiation source. The diffractograms were recorded with a step size of 0.07° using an incidence angle of 0.5° at 40 mA, 45 kV of generator settings.
while having a high content of O vacancies, i.e. 4 at.% (= 14.4 mol% Y2O3) resulting in every oxygen atom to be a neighbor of a vacant site or at least four of its six neighboring oxygen atoms, leads to the cubic structure. For these reasons, we believe that the only possible atomistic mechanism behind stabilization of c-phase of zirconia is lattice distortion caused by O vacancy incorporation in the zirconia lattice, ultimately forcing the O and Zr atoms to arrange themselves in the high symmetry cubic crystal. In conclusion, based on our DFT calculations and their remarkable agreement with our experimental data, we conclude that incorporating (as low as 3 at.%) O vacancies is the sole mechanism responsible of promoting zirconia to the high-temperature c-phase, at room temperature. Moreover, development of such stable phase pure cubic zirconia at room temperature by implementing such an optimized and controlled method might not only provide an alternate of YSZ but could also be of great fundamental and technical interest for other oxides (e.g. HfO2, TiO2 and Bi2O3). This work is supported by the Belgian Government through the “Pôle d'Attraction Inter universitaire” (PAI, “Plasma-Surface Interaction”, Ψ). S. Konstantinidis is a Research Associate and J. Cornil a Research Director of the FNRS (Fonds National de la Recherche Scientifique), Belgium.
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Appendix A. Supplementary data Supplementary data to this article can be found online at http://dx. doi.org/10.1016/j.scriptamat.2016.06.025. References
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