Understanding the local structure of Eu³⁺- and Y³⁺-stabilized zirconia: insights from luminescence and X-ray absorption spectroscopic investigations

Zirconia (ZrO₂) is a material with exceptional properties such as a very high melting point of 2715 °C and the capability of forming solid solutions with various isovalent and aliovalent cations [1]. The incorporation of under- or oversized cations leads to a phase transformation from the thermodynamically stable phase, i.e., the monoclinic phase under ambient conditions, to its high temperature polymorphs, the tetragonal or cubic structures, depending on the type and amount of dopant [2‐4]. In case of isovalent substitution of Zr⁴⁺, the mismatch of the ionic radii of the host versus dopant cation causes a distortion of the crystal lattice, leading to the aforementioned phase transition. When subvalent doping is performed, charge compensation takes place via the introduction of oxygen vacancies into the system. This is assumed to have a strong stabilizing effect on the t- and c-ZrO₂ phases [4, 5]. The incorporation of supervalent dopants is also feasible; however, less is known about these structures [3]. Due to this unique set of properties, the material can be tailored to meet the requirements of the targeted applications. The three most common stabilizers are ceria (CeO₂), yttria (Y₂O₃) and magnesia (MgO). CeO₂-doped ZrO₂ is used in three-way catalysts of internal combustion engines [6]. Y₂O₃-stabilized zirconia (YSZ) exhibits very good oxygen conducting properties, resulting from the oxygen vacancy formation in ZrO₂. At the same time, YSZ is acting as an electrical insulator. In combination with the very high temperature stability, this makes it an ideal electrode material in the field of solid oxide fuel cells (SOFCs). Besides the use of YSZ as cathode material due to the high oxygen mobility [7‐9], it also finds application as basis for the anode material in SOFCs, where a metallic anode material is deposited on the YSZ [10‐12]. YSZ is further used as sensor material for reducing gases, such as CO, H₂ and hydrocarbons but also for H₂S and NO_x [13]. MgO-stabilized ZrO₂ is an important construction material due its exceptional strength and fracture toughness. The fracture toughness rests upon the martensitic transformation of metastable t-ZrO₂ to m-ZrO₂, which is accompanied by a volumetric expansion. This expansion resulting from the diffusionless transformation, induced by the stress field around a crack in a ZrO₂ ceramic material, leads to local compressive stress which counterweighs the driving force of the crack propagation [1, 14‐16]. Besides these three very commonly studied materials, applications for various other stabilizing cations can be found. Ti⁴⁺-, Al³⁺- or K⁺-stabilized ZrO₂ can be used as catalysts in the biofuel production [17]. Lanthanide-stabilized zirconia finds usage, especially in the field of optoelectronics [18‐29]. Further research of lanthanide-doped ZrO₂ has been performed in the field of oxygen ion conductors [30], basic material properties [31‐33], as well as in fundamental research. Zirconia is also studied in nuclear research in relation to the transmutation of actinides (An) [34‐37], and the production of inert matrix fuels [38, 39]. Furthermore, ZrO₂ is an important solid phase at several stages of the nuclear fuel cycle. It occurs as a corrosion product at the fuel rod cladding surface where it may come into contact with dissolved radionuclides in repositories for nuclear waste storage, and it has been considered a potential host matrix for the immobilization of high-level radioactive waste streams containing especially the transuranium elements (TRUs) Pu, Am and Cm [40‐42]. The most stable oxidation state of the latter two TRUs is +3, while Pu³⁺ can be stabilized under anoxic, mildly reducing conditions. For applications as nuclear materials, the crystalline host structure must be flexible enough to accommodate the rather large actinide cations. Further, structural properties such as a large overall grain size and low porosity, which reduce the reactive surface area and subsequently increase their corrosion resistance in contact with water, are important.

Due to the large amount of applications for ZrO₂ doped with trivalent rare earth elements (REE), special interest exists in the structural properties of these solid phases. The formation of oxygen vacancies in the lattice as a result of REE³⁺ incorporation leads to a deviation between the long-range order in the bulk and the short-range order of the local structure. Despite the increase in the bulk symmetry from m- to t- to c-ZrO₂ with increasing doping fractions, the local symmetry is heavily distorted in these stabilized t- and c-ZrO₂ phases. A change in the coordination number from 8 to 7 occurs, especially for the Zr⁴⁺ but also for the trivalent dopant depending on the overall concentration of oxygen vacancies in the crystal structure [4]. The reduction in the Zr–O coordination number is assumed to result in the stabilization of the high-temperature phases, due to a reduction of the stress around the small Zr⁴⁺ cation. In addition, distortion of the crystal structure takes place due to the size mismatch between the small Zr⁴⁺ and the much larger REE³⁺ cations.

The structural properties of zirconia solid solutions have been investigated in numerous studies using various experimental and computational methods. Several extended X-ray absorption fine structure (EXAFS) studies of doped zirconia phases have been conducted with various trivalent dopants such as Fe³⁺, Ga³⁺, Ge³⁺, Nb³⁺, Ce³⁺, Gd³⁺ [4, 43], Er³⁺ [44] and Am³⁺ [36]. YSZ has received especially large attention in this field [45‐49]. The EXAFS studies have shown that the direct coordination environment around the dopant stays very constant, independent of the ZrO₂ crystal phase. This has partly been explained by the location of the oxygen vacancies in the crystal structure and their preference for the Zr⁴⁺ cation rather than the dopant cation [36, 45]. Laser-induced lanthanide luminescence spectroscopy, especially of Eu³⁺, has further been employed to study the local structure of zirconia solid solutions [1, 21, 30‐32, 50‐55]. Here, differing luminescence behavior is observed for m-ZrO₂ as compared to the high-temperature polymorphs (t- and c-ZrO₂). However, no significant change in the luminescence of Eu³⁺ doped t-ZrO₂ and c-ZrO₂ can be observed, especially when using the classical nonselective excitation method, which results in co-excitation all Eu³⁺ environments in the crystal structure. When using site-selective excitation of the ⁵D₀ ← ⁷F₀ or the ⁵D₁ ← ⁷F₀ transition, multiple Eu³⁺ species have been observed in several studies [30‐32, 50‐52, 55]. However, contradicting conclusions have been drawn on the nature of the individual species, and their assignment to specific luminescence emission spectra. Various computational studies have been performed to investigate structural properties of doped zirconia phases [56‐61]. Special interest in experimental and computational studies has been devoted to the oxygen diffusion in YSZ [62‐65]. It was found that the doping fraction and the location of oxygen vacancies in the lattice are crucial for the oxygen mobility in these phases [58, 62].

Despite the existing knowledge of t- and c-ZrO₂ phases stabilized with trivalent cations, knowledge gaps still exist regarding the relationship between the bulk structure and order/disorder phenomena occurring around the dopant cation. To our knowledge, no direct comparison of changes occurring in the dopant (Ln³⁺) versus Zr⁴⁺ environments has been attempted, which could provide a deeper insight into how similar or dissimilar structural changes around the host and dopant cations are for various dopant concentrations and how they influence the overall symmetry of the bulk structure in comparison with order/disorder phenomena on the dopant level.

Thus, in the present work, we report on first combined luminescence spectroscopic (TRLFS) and EXAFS investigations of Eu³⁺ or Y³⁺/Eu³⁺ co-doped tetragonal and cubic zirconia polymorphs. Powder X-ray diffraction (PXRD) has been used to study the bulk behavior of doped zirconia phases in dependence of the doping percentage. Scanning electron microscopy (SEM) and transmission electron microscopy (TEM) were employed as imaging method for macroscopic investigations. Site-selective Eu³⁺ TRLFS at 10 K was used to study the local dopant environment independently from the bulk. Additionally, EXAFS investigations at the yttrium and zirconium K-edges were performed on Y³⁺/Eu³⁺ co-doped zirconia samples for complementary information on the local structure of the dopant and the host cations. To bridge the gap between the bulk structural and dopant-specific investigations, additional X-ray pair distribution function (XPDF) analyses were conducted for Eu³⁺ doped zirconia samples. XPDF is a complementary method to study local structures and disorder phenomena in crystalline as well as amorphous materials, capable of providing information of the average, periodic structure of a material as well as potential local deviations from this average structure. While many studies can be found where the PDFs of zirconia systems were obtained based on computational approaches [63, 66, 67], the availability of PDFs based on experimental data [i.e., X-ray diffraction (XPDF) or neutron diffraction (NPDF)] is rather limited. To the best of our knowledge, only one study can be found where the XPDF of a monoclinic zirconia sample is presented [68]. Therefore, we report on the first XPDF studies on stabilized ZrO₂ polymorphs in the present work.

With the help of PXRD, it can be seen that the cubic zirconia phase dominates the phase composition of all samples with some remaining tetragonal zirconia in the intermediate doping range (below 25 mol%). With the combination of PXRD and XPDF, the formation of a secondary phase could be excluded. As briefly mentioned before, the crystallites of the Eu³⁺ doped samples are approximately three times smaller than the Y³⁺/Eu³⁺ co-doped ones for comparable dopant concentrations. This is likely an effect of the cation radius, which is larger for Eu³⁺ in comparison with Y³⁺, resulting in a larger strain and stronger lattice distortion in the samples with a direct influence on the maximum crystallite size. This may be of importance for applications where the crystallite size directly influences the material performance, such as its optical properties [96, 97], oxygen ion conductivity [98] or dissolution rate and degradation [99‐103]. Despite that our combined PXRD and XPDF analyses showed a linear increase in the lattice parameters with increasing dopant concentration, indicative of solid solution formation, several of our luminescence spectroscopic observations point toward the presence of multiple distinct dopant environments in the solid structure, i.e., Eu³⁺ sites with differing coordination environments. Arguments for this can be found in the shifting of the excitation peak at very high doping, as well as the differing emission spectra and lifetimes for differing excitation energies in the individual samples. The presence of multiple Eu³⁺ environments in stabilized ZrO₂ solids has been reported in multiple studies. Montini et al. [55] have studied Eu³⁺/Ce⁴⁺ co-doped zirconia phases and found clear evidence for the presence of three nonequivalent Eu³⁺ environments in the stabilized host structures based on the collected Eu³⁺ excitation spectra. The authors attributed the more redshifted species with peak positions of 579.5 nm and 580.5 nm to two nonequivalent incorporation species due to the presence of different sites in the solid matrix. A blueshifted component with a peak position around 578.5 nm was shown to arise from Eu³⁺ incorporation into a superficial site. In studies by Yugami et al. [104] and Borik et al. [51, 52], three to four different Eu³⁺ environments in Y³⁺/Eu³⁺ co-doped polycrystalline samples or single crystals of zirconia, respectively, were observed. The authors attributed the specific emission signals to the formation of oxygen vacancies in the crystal structure and the subsequent change in the oxygen coordination number around the Eu³⁺ cation depending on the location [nearest-neighbor (NN) or next-nearest-neighbor (NNN) positions] of the oxygen vacancies.

Thus, in analogy with these studies, our excitation profiles were fitted using multiple Gaussian peaks. The best fit was achieved using three Gaussian peaks to fit the overall excitation spectra (see Fig. 10 for two examples and Figure S8 for all fits). The peak positions for the samples from 11 to 24 mol% Eu³⁺ and 16 to 26 mol% Y³⁺ doping were first fitted by varying all parameters. Thereafter, the obtained average values for the peak positions of the present species, namely 578.1 ± 0.5 nm (species 1), 579.0 ± 0.4 nm (species 2) and 579.7 ± 0.5 nm (species 3), were fixed to avoid overparametrization. All samples could be fitted with these peak positions. The results of the fitting of all excitation spectra are summarized in Table 3.

Some systematic trends can be deduced from the obtained Gaussian fits of the excitation profiles. A systematic decrease in the full width at half maximum (FWHM) can be observed from species 1 to species 3. This is in line with the recorded emission spectra at excitation wavelengths of 577.2 nm (predominantly exciting species 1), 579.3 nm (predominantly exciting species 2), and 580.0 nm (predominantly exciting species 3), where emission peak narrowing could be recorded from species 1 to species 3. In addition, the overall magnitude of the crystal field splitting of the ⁷F₁ band decreases in the same order pointing toward a lowering of the crystal-field perturbation. Thus, there seems to be a systematic increase in the local order in these three Eu³⁺ environments (species 1 → species 3).

In both the Eu³⁺ doped and Y³⁺/Eu³⁺ co-doped samples, species 1 increases significantly for samples of high doping where a clear shift of the excitation peak was observed (see Fig. 5), while species 2 varies slightly in the range of 30–45% for Eu³⁺ doping and 51–55% for Y³⁺/Eu³⁺ co-doped samples. Species 3 decreases in both cases with increasing doping fraction to about 10% for the highest doping fractions used here.

The obtained peak position for species 1 of 578.1 nm agrees rather well with the peak position of 578.5 nm obtained by Montini et al. [55] for Eu³⁺ incorporation at the zirconia surface. In the zirconia single crystals studied by Borik et al. [51, 52], such a blueshifted ⁷F₀ peak position was not observed. However, given the much smaller surface area of a single crystal compared to nanoparticle powder, no significant contribution of a surface-associated Eu³⁺ species to the luminescence is expected. The increasing trend of species 1 with increasing doping percentage follows the trend of decreasing crystallite size seen in the PXRD bulk studies (see Table 1). This is visualized in Figure S9. Thus, it is likely that the Eu³⁺ dopant is located on a superficial site, which will increase in abundance as the crystallite size decreases. This hypothesis is further supported by the larger abundance of this species 1 in the Eu³⁺ doped ZrO₂ solid phases, which are smaller in crystallite size, in comparison to the Y³⁺/Eu³⁺ co-doped ones, which yield larger crystallites. Due to lattice relaxations at a superficial site, a higher degree of freedom and therefore a greater variety of coordination spheres are possible, resulting in the rather large FWHM in comparison with the two other species.

The excitation peak positions of species 2 (579.0 nm) and species 3 (579.7 nm) are very similar to those identified by Borik et al. [52] at 578.9 nm and 579.3 nm for Y³⁺/Eu³⁺ co-doped ZrO₂. Excitation spectra with similar excitation peak positions can also be deduced from the data in Yugami et al.; however, exact peak positions for the extracted species were not given. In Borik et al., the abundance of the latter species was seen to increase with increasing doping, leading to the conclusion that the species results from a Eu³⁺ environment with one oxygen vacancy in the first coordination sphere. Due to a subsequent decreasing trend of the excitation peak at 578.9 nm, this Eu³⁺ environment was assigned to an eightfold coordinated Eu³⁺ site with oxygen vacancies located only at NNN positions. Exactly the opposite trend was reported in Yugami et al. for the investigated polycrystalline Y³⁺/Eu³⁺ co-doped samples. Based on their data, the excitation peak at > 579 nm, decreasing in abundance with increasing dopant concentration, was attributed to the pristine Eu³⁺ environment, coordinating to eight oxygen atoms in the crystal structure. Based on our results presented in Table 3, a trend similar to what was observed in Yugami et al. can be deduced for species 3, where an increasing dopant substitution in the ZrO₂ solid phases results in a decrease in its fraction. The percentage of this species is very similar in both the Eu³⁺ and Y³⁺/Eu³⁺ co-doped sample compositions, i.e., the amount of this species seems to be related to the overall dopant concentration. In order to visualize the trends of the various Eu³⁺ species derived in the present study, the mole percentages of all Eu³⁺ species and the number of oxygen vacancies have been plotted as a function of the Y³⁺/Eu³⁺ co-doping in Fig. 11, left, and Eu³⁺ + Y³⁺ doping fractions in Fig. 11, right.

A line for the distribution of oxygen vacancies around the Eu³⁺ ion (orange line), assuming a random (nonpreferential) distribution of the vacancies in the crystal lattice, has been included in the figures. For more details of these plots, the reader is referred to the SI (Equations S1–S3 and Figure S10). In both the pure Eu³⁺ doping row and the Y³⁺/Eu³⁺ co-doping row, species 2 can be seen to correlate well with the overall number of oxygen vacancies; especially, in the Y³⁺/Eu³⁺ samples (Fig. 11, right), the amount of this species and the number of oxygen vacancies are almost identical. In the case of solely Eu³⁺ doped ZrO₂ (Fig. 11, left), an increase in the amount of species 2 from 11 to 16 mol% doping can be seen, after which a plateau is obtained. In this plateau region, a strong increase in the surface-associated species (species 1) occurs, which may explain the lower amount of species 2 in these Eu³⁺ doped samples in comparison with the Y³⁺/Eu³⁺ co-doped ones. In general, however, this species clearly follows the trend of the increasing amount of oxygen vacancies in the host lattice with increasing M³⁺ doping, which implies that this Eu³⁺ environment is associated with an oxygen vacancy in the first coordination sphere, i.e., in NN position.

Species 3 on the other hand is rather low in abundance over the whole investigated dopant concentration range, and the overall amount of the species is rather constant or slightly declining. As no correlation between the number of oxygen vacancies and species 3 can be seen, it is not reasonable to assume that this species is Eu³⁺ associated with one oxygen vacancy in the crystal structure as found in Borik et al. [52] but rather a Eu³⁺ species surrounded by eight oxygen atoms in the crystal lattice, in agreement with the assignment in Yugami et al. [104]. Reasons for these contradicting observations could be the very different sample types used in the different studies. While Borik et al. used single crystals of Eu³⁺ doped YSZ in their luminescence spectroscopic investigations, polycrystalline powder samples have been used in the current study and in the study by Yugami et al. These are not in thermodynamic equilibrium due to the slow cation mobility in these phases. Another reason for the dissimilar results could be the different excitation wavelengths used to excite the Eu³⁺ species. In Borik et al., excitation was performed via the ⁵D₁ ← ⁷F₀ transition with a constant wavelength of 532 nm. In Yugami et al. and the present work, site-selective ⁵D₀ ← ⁷F₀ excitation spectra were collected by varying the excitation wavelength. This way, different crystal field perturbations of the various Eu³⁺ environments are taken into account, i.e., preferential excitation of a species with a matching energy gap (energy of incident laser light versus energy of the ⁵D₁ ← ⁷F₀ or ⁵D₀ ← ⁷F₀ transition) does not take place.

Interestingly, for dopant percentages of 11–16 mol% (Eu³⁺ doped samples) and of 13–26 mol% (Y³⁺/Eu³⁺ co-doped samples), the amount of Eu³⁺ species with the oxygen vacancy in NN position (species 2) is larger than the statistical distribution of vacancies between Zr and Eu. This implies that oxygen vacancies have a slight tendency for the dopant Eu³⁺ in our samples. Even though both Yugami et al. and Borik et al. reported on the relative amounts of Eu³⁺ species with one or two vacancies in NN or NNN positions in their ZrO₂ host phases, our study is the first one to report on the preferential location of oxygen vacancies for Eu³⁺doped ZrO₂ samples. Therefore, a direct comparison with other experimental studies using Eu³⁺ as dopant cation cannot be made. There are nonetheless both experimental [105] and computational [58] studies supporting the results obtained in the present study, where a preferential location of oxygen vacancies in the crystal lattice has been found in the direct coordination of other oversized trivalent dopants rather than Zr⁴⁺. These studies, however, are outnumbered by both experimental results and computational evidence for the preferential location of oxygen vacancies around Zr⁴⁺. This vacancy formation around the Zr host is explained to promote the stabilization of the tetragonal and cubic zirconia phases at ambient conditions as the reduction of the Zr–O coordination from eight to seven reduces the stress around the rather small host cation. We therefore have reason to believe that the sample synthesis procedure in terms of calcination time and temperature has a strong influence on the location of oxygen vacancies in the crystal lattice. As already discussed in connection to the phase composition of our synthetic zirconia samples, our samples are not in thermodynamic equilibrium due to the very slow cation diffusion in ZrO₂ at 1000 °C, which, however, is a rather typical calcination temperature in the synthesis of ZrO₂ [1, 26, 106, 107]. Although the anion mobility in stabilized zirconia is higher, the same assumption can be applied for the oxygen distribution in the crystal lattice. When the dopant is introduced into the zirconia solid phase, vacancy formation occurs next to the dopant due to the charge mismatch (M³⁺ ↔ M⁴⁺) [58]. For short synthesis times, especially when combined with rather moderate temperatures, oxygen migration through the lattice may not reach a steady state which subsequently results in an unexpected distribution of oxygen vacancies in the crystal lattice. Thus, we assume that an increase in the calcination time and calcination temperature would drive the vacancies toward the host cations.

Finally, having assigned the three Eu³⁺ environments in the ZrO₂ crystal structures, an attempt was made to designate the measured luminescence lifetimes to these Eu³⁺ species. To do so, an excitation spectrum of Y³⁺/Eu³⁺ co-doped ZrO₂ with an overall dopant concentration of 18 mol% was measured after a delay time of 8 ms and compared to the excitation spectrum recorded 1 µs after the laser pulse (Figure S11). Based on the recorded lifetimes of τ₁ = 1030 ± 310, τ₂ = 2950 ± 250 and τ₃ = 6560 ± 360 µs, the excitation spectrum after a delay of 8 ms should show no presence of the species with the shortest lifetime and approximately 7% and 30% of the species with lifetimes of 2950 µs and 6560 µs, respectively. However, as evident from Figure S11, only a shoulder on the blue side, corresponding to species 1, is completely absent from the delayed excitation spectrum, while the main peak has remained rather unchanged. Based on this, the short lifetime of 1030 ± 310 µs can indeed be assigned to the surface-associated Eu³⁺ species. This lifetime is rather short for a fully incorporated Eu³⁺ ion. In fact, the Horrocks equation [Eq. (1)] would predict partial hydration of Eu³⁺ with on average 0.4 H₂O molecules coordinating to the cation. Species 2 and species 3, on the other hand, seem to have rather similar lifetimes based on the very constant excitation peak shape independent of delay. This is likely to be related to some energy transfer mechanism either from Eu³⁺ to Eu³⁺ centers or, more likely, between Eu³⁺ centers and defect electrons present in the lattice, which hampers the assignment of individual Eu³⁺ environments in the crystal structure based on their luminescence decay properties.

Our EXAFS investigations of ZrO₂-doped samples with dopant concentrations ranging from 16 to 26 mol% show almost identical Zr–O/Zr-cation or Y–O/Y-cation distances of 2.147 ± 0.005/3.553 ± 0.004 Å and 2.317 ± 0.003/3.607 ± 0.003 Å, respectively, over the entire compositional range. Furthermore, from EXAFS as well as XPDF, it can be excluded that the first Zr–O shell is split into 2 × 4 subshells in these samples. It is striking that despite a rather large difference in the Zr–O and Y–O distance (~ 8%), the Zr-cation and Y-cation distances are very similar (difference of ~ 1.5%). The changes introduced by the oversized dopant are therefore mostly limited to the first coordination sphere. This observation was explained by Li et al. [43] with the deformability of oxygen ligands as well as with the continuum elasticity theory, where a quadratic decay of the radial displacement is predicted.

Nevertheless, the different species observed with TRLFS cannot be observed with EXAFS as both the Zr and Y environments stay constant over a large doping range. As EXAFS yields information of the average Zr or Y environment in the sample, the distinction of nonequivalent species in complex samples such as the investigated M³⁺ doped ZrO₂ ones is hardly possible. In our sample rows, multiple structural changes take place simultaneously. With increasing doping, oxygen vacancies are formed in the lattice, decreasing the overall M–O CN. However, at the same time, the crystallite size decreases with increasing doping leading to a higher abundance of superficial sites. The changes in the fractions of the three species observed with TRLFS are rather small, further hindering the observation of a trend in the EXAFS fits; especially, the structural difference between species 2 and 3 is vastly defined by a change in the M–O coordination number. However, the determination of the coordination number with EXAFS inhibits a rather larger error and the coordination number in ZrO₂ does not seem to be captured properly by EXAFS fitting. This can be seen in our attempts to fit the coordination numbers (see Table S3), as well as in the literature, where the coordination number of doped ZrO₂ is typically fixed to a certain value.

The three different species observed in our studies consist of a subset of a large multitude of slightly differing coordination environments themselves. This becomes apparent when considering that in a 3 × 3 × 3 supercell of ZrO₂ doped with 16 mol% Y³⁺, a number of 9.3·10²³ possible theoretical constellations of Zr⁴⁺ and Y³⁺ cations, O²⁻ anions and oxygen vacancies exist. Such multiple, similar species can lead to destructive interference of the EXAFS signal of the individual Eu³⁺ environments, causing a cancelation of their signal. Therefore, the differences of the samples observed with TRLFS are not accessible with our EXAFS investigations.

The PDF analysis based on X-ray diffraction data is not affected by self-quenching effects. Nevertheless, similar atomic distances are derived for the M–O and M–M distances and no clear trend of the atomic distances with changing doping fractions can be derived. This supports the assumption that the small changes in the system resulting from multiple species cannot be readily resolved. A self-quenching effect in the EXAFS studies could, nevertheless, be present, which could explain the differing values obtained for the second M–O distances with EXAFS and XPDF. The large number of oxygen atoms in this shell could make the self-quenching of parts of the signal rather likely. These distances, however, also inherit a rather large error margin due to the high degeneracy and a strong overlap of PDFs at this distance.

An alternative explanation for the observed differences in the luminescence behavior of the individual species, but not in the structural studies with EXAFS and XPDF could be that all three species do in fact have very similar abs-O or abs-Zr distances and that solely the change in the coordination number and the local symmetry is responsible for the differences observed in the luminescence spectroscopic investigations.

In the present study, we have combined bulk structural investigations with spectroscopic studies of Eu³⁺- or Y³⁺/Eu³⁺ co-doped tetragonal and cubic zirconia polymorphs. Our bulk structural characterizations show that the dopant is homogenously distributed in the ZrO₂ host structure, thereby stabilizing the tetragonal and cubic polymorphs with little to no changes in the phase composition when using Y³⁺ versus Eu³⁺ in the stabilization process. The incorporation of Y³⁺ and/or Eu³⁺ results predominantly in the cubic phase at doping percentages of 16 mol% and above with fractions of the tetragonal phase present until approximately 19 mol%.

Our EXAFS and XPDF results imply that even though the abs-O distances for the host (Zr) and dopant (Y or Eu) differ by about 8% or 11%, the overall average coordination environment seems to remain unchanged over the studied doping range. Further insights from luminescence spectroscopic studies, i.e., site-selective luminescence spectroscopic investigations of Eu³⁺ at temperatures below 10 K, reveal the presence of nonequivalent Eu³⁺ environments in the ZrO₂ solid structures. These Eu³⁺ environments are assumed to arise from Eu³⁺ incorporation at a superficial site (species 1, λ_ex = 578.1 nm) and incorporation on two bulk sites (species 2, λ_ex = 579.0 nm and species 3, λ_ex = 579.7 nm) differing in the location of the oxygen vacancies with respect to the dopant cation.

The results from our study clearly show that structural modifications, such as the phase transformation from t-ZrO₂ to c-ZrO₂ which can be described with the help of various ideal crystal structures on the bulk level, do not accurately capture the structural changes and/or displacements occurring around the incorporated dopant. A mechanistic description of such local changes requires spectroscopic techniques such as TRLFS capable of distinguishing species or environments with very subtle differences for an adequate understanding of how bulk structural changes in terms of crystallite size or oxygen vacancy formation influence the local coordination environment of an incorporated cation. With the TRLFS method, we could show that the applied synthesis procedure results in a preferential placement of oxygen vacancies in the direct coordination environment of the Eu³⁺ dopant (species 2). The location of vacancies in the crystal lattice has been shown to have direct implications for the oxygen mobility in stabilized ZrO₂ and will play a role in applications such as SOFCs, where a high oxygen mobility is a prerequisite. Further, our combined bulk structural and luminescence spectroscopic results showed that the decreasing crystallite size renders Eu³⁺ at superficial sites in the zirconia matrix. The association of the dopant surface could have direct implications for the corrosion resistance and performance of crystalline zirconia ceramics used for the immobilization of actinides from nuclear fuel-related waste streams.

This study shows the importance of combining the strength of multiple techniques to understand complex systems like the solid solutions of zirconia and thereby enable an advanced understanding of the changes taking place in the material on the molecular level.