Eu3+ doped tin dioxide (SnO2) thin films are deposited by the sol-gel-dip-coating process on top of GaAs films, which is deposited by resistive evaporation on glass substrate. This heterojunction assembly leads to interesting luminescent emission from the rare-earth ion, unlike the SnO2 deposition directly on a glass substrate, where the Eu3+ transitions are absent. In the heterojunction, the Eu3+ transitions are clearly identified and are similar to emission from samples in the form of pressed powder (pellets), thermally treated at much higher temperatures. However, in the form of films, the Eu emission comes along a broad band, located at higher energy compared to Eu3+ transitions. This broad band is blue shifted as the thermal annealing temperature as well as the crystallite size increase. Although the size of nanocrystallites points toward quantum confinement, another cause of the detected broad band is more feasible: the electron transfer between oxygen vacancies, originated from the disorder in the material, and trivalent rare-earth ions, which present acceptor-like character in this matrix. This electron transfer may relax for higher temperatures in the case of pellets, and the broad band is eliminated.
© 2014 Optical Society of America
Rare-earth doped semiconductors have increased interest due to applications in optoelectronic, such as displays and optical communication devices . Doping wide bandgap semiconductors with rare-earth ions reduces the thermal quenching effects , allowing luminescent emission with high quantum efficiency from the rare-earth ion. Tin dioxide (SnO2) is a wide bandgap semiconductor (energy about 3.6–4.0 eV ) with above 90% transparency in the visible range [3, 4]. Sol-gel deposited thin films are composed of small crystallites, generally of nanoscopic dimensions, which may influence the emission spectra, because it is well-known that particle size influences the optical properties of materials. Rare-earth based luminescent nanocomposites are widely investigated for the production of high definition images, where the image resolution of the cathodic-ray tube is determined by the particle size . Photoluminescence (PL) properties of Eu-doped SnO2 xerogels have been investigated under different lasers excitation , which allowed distinguishing two Eu3+ families in SnO2. In the first of these families the PL measurements show that Eu3+ ions are located in symmetric centers, substitutional to Sn4+ in the lattice, and in the other family, the PL measurements allow identifying that rare-earth impurity gets preferentially located at asymmetric sites, the particles surface. In the latter case the location is related to the segregation caused by excess of doping. Excitation with 266 nm is energetically higher than the SnO2 bandgap, promoting matrix band-to-band excitation and an efficient energy transfer to symmetric located Eu3+, substitutional to Sn4+. This happens for samples with doping level until 0.5 at.%. Above this doping level, energy transfer to ions located close to the particles surface, with lower symmetry, becomes operative. This behavior may be also related to the particles size, which decreases by increasing the Eu concentration, as deduced from X-ray diffraction data. The location of Eu3+ ion at surface sites of particles would explain the preferential excitation of low symmetry sites with 266 nm (band-to-band excitation). Temperature dependent PL allows identifying a sharp growing integrated intensity curve from 10 to 240 K, with a slight decrease on the intensity from 240 to 300 K. Then, the maximum emission is about 240 K  and considerable luminescence intensity is observed even at room temperature. This observation points out for technological applications, for instance as electroluminescence devices operating at room temperature.
Along with the Eu3+ transitions, some transitions from the SnO2 matrix itself have been identified. Kim and associates  detected a UV broad band peaking about 400 nm, which was attributed to oxygen vacancies in SnO2. Another crucial point concerning the emission from a matrix with nanosized building blocks refers to the possibility of quantum confinement. The emission is highly dependent on the crystallite size, similar to quantum dots. As its size decreases the emission spectra shifts to high energies and the oscillator strength is concentrated into just a few transitions . The quantum confinement arises as a result of changes in the density of electronic states. For a bulk system the energy and the crystal momentum can both be precisely defined, unlike the position. For a particle localized in a quantum dot or, by analogy, inside a nanocrystallite, the uncertainty in the position decreases, so that the momentum is no longer well defined. Considering that a series of nearby transition in the bulk occur at slightly different energies and momentum, it may be overlapped by quantum confinement in the crystallite into a single and intense transition. Then, with size reduction the emission shifts to higher energies [8–10]. Fine-tuning the luminescence emission is possible by controlling the nanocrystallite size of silicon in silicon nitride samples, where the size corresponding to red, green and blue are 4.6, 3.1 and 2.7 nm, respectively .
On the other hand, luminescence from thin films is not usual as from xerogels. The main difficulty is probably related to the sample dimensions, where the number of emission centers in the laser beam path leads to a low efficiency. There are a few publications reporting luminescence on SnO2 films. Eu-doped SnO2 thin films deposited by rf magnetron sputtering showed an emission peak at 588 nm, with a maximum intensity for 1.0at% doping of Eu doping and a 1200°C firing temperature . Eu-doping and its luminescence has been reported in other sort of thin films [13–17]. In general, the two main peaks are assigned to the 5D0 → 7F1 and 5D0 → 7F2 transitions of the Eu3+, whose relative intensity depends on the doping concentration. These emission lines are often generated for some sort of energy transfer. Besides, the high annealing temperature, above 1000°C, helps to split the emission lines.
SnO2/GaAs heterojunctions thin films have been successfully obtained, giving birth to smooth interface and improved electrical properties [18,19]. In this paper, it is shown photoluminescence from Eu3+ transitions in SnO2 films, deposited on top of GaAs as a heterojunction. The Eu3+ transitions are not identified in the PL of SnO2:Eu thin films deposited on top of glass substrates, without the GaAs layer, annealed at similar temperatures as the heterojunction samples. A broad band, not related to Eu3+, is also observed in the PL spectra of thin films, which is blue shifted as the thermal annealing temperature and, thus, the crystallite size increase, being destroyed for pressed powder, annealed to high temperature.
2. Experimental procedure
GaAs thin films are deposited by the resistive evaporation technique, in a BOC Edwards evaporator system Model AUTO 500. GaAs tiny pieces were placed on a tungsten boat inside a low pressure chamber (10−5 mbar). This technique is basically the sublimation of the boat stuff by controlling of electrical current through the metallic boat. It leads to evaporation of GaAs due to the low pressure inside the chamber, and a beam of the material in gaseous phase is directed to the substrate, located in the up part of the evaporator system. The rotation of the substrate holder (round plate) assures homogeneous composition and thickness of deposited films. After evaporation, films where submitted to a thermal annealing at 150°C by 30 minutes in an EDGCON 3P oven.
The deposition of Eu3+ - doped SnO2 thin film layer was described in details elsewhere . The deposition takes place in air atmosphere (room temperature), and after each layer, samples are dried in air by 20 minutes and treated at 200°C by 10 minutes in the same oven used for GaAs annealing. This procedure was repeated 10 times. The final annealing was at 200°C by 1 hour or 400°C by 20 minutes for different samples. In the case of the SnO2 film sample without GaAs, deposited directly on glass substrate, the final annealing was 500°C by 1 hour. The xerogels (powder) SnO2:2%Eu and SnO2:4%Eu were obtained by evaporating all the solvent. Then the powder was treated at 1000°C by 1hour. The remaining powder was pressed down to pellet format with 5000kgf/cm2 by 3 minutes, and treated again at 1000°C by 20 minutes. In order to understand the rule of thermal annealing for films, a SnO2:2%Eu film was deposited directly on quartz substrate, and the final annealing was 1000°C by 1hour.
X-ray diffraction patterns are obtained using a Rigaku diffractometer, model D/MAX 2100PC operated with a Cu Kα radiation source (1.5405 Å) and a Ni filter for reducing the undesirable Kβ radiation. The data are collected with 0.02° of step and scanning rate of 1°/minute in the 2theta mode for films (fixed incident angle of 1.5°) in the angular region of 20° ≤ 2θ ≤ 80°.
For the photoluminescence (PL) measurements in the UV-Vis range, samples were excited with a modulated 350 nm line of a Kripton (Kr+) laser and the signal was detected by a R955 PMT from Hamamatsu and a SR530 Lock-in Amplifier from Stanford Research System. A single configuration monochromator was used for scanning the PL spectra.
Morphological characterization as well as the compositional analysis were carried out with a scanning electron microscopy (SEM) system (FEI model Quanta 200 or Carl Zeiss model LS15), equipped with a energy dispersive X-ray (EDX) microanalysis system (Oxford model Inca 250P20). The samples were attached to a sample holder (stub) by using conducting silver paint. After the paint had dried, metallization was carried out with gold (Au) deposition in an Edwards Scancoat Six sputter coater system with a deposition time of about 3 minutes. Electron were accelerated under 25.0 kV for EDX analysis.
3. Results and discussion
Figure 1 shows a schematic diagram of samples investigated in this paper, where the GaAs film is deposited on a soda-lime glass substrate and the Eu-doped SnO2 film is deposited on top of the GaAs film. The bottom inset, in the left, represents the band energy diagram of the heterojunction, showing a possible formation of a two-dimensional electron gas (2DEG) at the interface. The inset at the upper part, in the left side yet, is a SEM image of the GaAs film surface, previous to the deposition of SnO2. The upper inset in the right side shows the SEM image of the SnO2:2at%Eu, deposited on top of GaAs. As can be seen the homogeneous distribution of nanometric particles at GaAs interface gives rise to a surface with more separated and larger particles. A similar heterojunction, but deposited in the inverse order of layers, gave rise to interesting interface properties [18,19], allowing the possible generation of a 2DEG, and showed the quality of samples deposited by this combination of techniques.
Figure 2 shows the luminescence spectra for the heterojunction samples GaAs/SnO2:2at%Eu with distinct thermal annealing (T.A.) of the SnO2 film, as described in the previous section: at 200°C by 1 hour, and at 400°C by 20 minutes. A third PL spectra is shown, for a SnO2:2%Eu film with thermal annealing at 500°C by 1 hour.
The heterojunction luminescence spectra, with distinct thermal annealing, show a broad band as well as well-defined peaks. This broad band has its maximum about 550 nm for the sample treated at 200°C whereas the peak of the broad band is blue-shifted for the sample treated at 400°C, exhibiting a peak about 470 nm. Besides, the higher annealing temperature induces higher intensity for the whole spectra. The well defined peaks are coincident with Eu3+ transitions: 5D0→7F1 (about 596 nm), 5D0→7F2 (about 618 nm) and 5D0→7F4 (about 704 nm). On the other hand, the emission spectra of SnO2:2at%Eu film, not coupled to a GaAs layer on a heterojunction, and treated at 500°C, presents a more intense broad band, but does not allow the identification of the Eu3+ transitions. It is interesting to mention that a series of 4 heterojunction samples treated from 300 to 500°C for 1 hour (final annealing) yield PL (not shown) with decreasing Eu3+ transition intensity and increasing broad band intensity, which is slightly blue-shifted.
Figure 3 shows X-ray difratograms for these three samples. They exhibit the noisy and diffuse profile typical of materials with nano-crystallized domain, where the building blocks are nanocrystallites. The labeled peaks correspond to the planes of rutile structure of tin dioxide (JCPDS- 41-1445) and planes of GaAs (JCPDS-80-0016). Characteristic peaks of SnO2 are seen at 33.7° and 51.5° corresponding to (101) and (211) crystallographic planes, and characteristic peaks of GaAs at 45.4° and 53.7°, related to (220) and (311) planes. At 27.1° there is probably an overlapping of two peaks: one related to plane (100), characteristic of SnO2, and another peak corresponding to plane (111) of GaAs.
As seen in Fig. 2, the heterojunction thermally treated at 200°C by 1 hour presents the broad emission peak centered at lower energy (higher wavelength), whereas for the heterojunction treated at 400°C the broad band peak shifts to a higher energy and the SnO2:2at%Eu film show the highest energy position of the broad band peak. Taking the (101) direction in the curves of Fig. 3, which is the most intense peak related to SnO2, and using the Scherrer equation , the evaluated crystallite size varies from 7.3 nm for the GaAs/SnO2:2%Eu treated at 200°C to 10.9 nm for the heterojunction treated at 400°C and 16.2 nm for the SnO2:2%Eu sample, treated at 500°C. Considering the broad band in PL emission shown in Fig. 2, the result is rather distinct from what we would expect from bands generated from quantum confinement, since the smaller crystallites would force overlap of nearby-lying bulk transitions into a single and intense transition, where size reduction shifts the emission to higher energies .
SnO2 nanocrystalline thin films grown on p-type InP (100)  has yield an evident broad band with peak at 3.13 eV. However, no precise explanation was found, since it could be due to all the luminescent centers, such as nanocrystals and defects in the SnO2 film. Considering that the peak position does not change with temperature, there is a series of possible reasons for the luminescence such as defects or nanocrystallline grains (quantum confinement) or to defect levels associated with oxygen vacancies or interstitials tin atoms, resulting from the nanosized dimensions of the building blocks in the SnO2 film . PL spectra have also been reported with intense UV-violet peak at 390 nm (about 3.18 eV) and a shoulder at 430 nm (about 2.88 eV) along with a broad peak at 520 nm (about 2.38 eV) . In polycrystalline oxides, oxygen vacancies are known to be the most common defect and, besides, the diffused oxygen recombines with oxygen vacancies, decreasing the oxygen vacancies concentration. It has been reported a UV-violet peak at room atmosphere after the sample be annealed , which means that the decrease of the number of oxygen vacancies does not lead to the decrease of the intensity of the UV-violet peak. It shows that the oxygen vacancies alone do not play an important role in the origin of the UV-violet peak. The main origin of the UV-violet luminescence band is ascribed to the electron transition between the donor level and the acceptor level formed by Sb ions . It is much more common to a trivalent rare-earth ion, such as Eu3+, to form an acceptor level in SnO2, than a Sb ion, whose predominant state is pentavalent [22,23]. Besides, rare-earths do not present a pentavalent ion. Then, it is expected that the broad band seen in Fig. 2 has contribution related to the electron transition between the donor level, formed by oxygen vacancies, and the acceptor level formed by Eu3+ ions. The dominance of the 3+ oxidation state of Eu also leads to high charge compensation as determined by electrical characterization measurements .
Figure 4 shows the emission spectra for pressed powder (pellets) treated at 1000°C, with two distinct Eu doping concentrations: 2at%, identical to the investigated films, and a higher concentration 4at%. The inset in Fig. 4 is the X- ray diffraction pattern of SnO2:4at%Eu powder.
The pellets do not present a broad band, and the Eu3+ transition are clearly identified at 557, 565 and 572 nm corresponding to the transition 5D1→7F2, 583 nm corresponding to 5D1→7F3, 593, 598 and 604 nm (5D0→7F1), 614, 618 and 625 nm (5D0→7F2), 633, 638 e 657 nm (5D0→7F3), and 716, 721 and 727 nm (5D0→7F6). Peaks related to 5D0→7F1 transition are the most intense. Although it is not shown here, some peaks were identified more clearly by magnification of the spectra region corresponding to low intensity transitions. It is evident that the more doped sample show more intense peaks, but all of the peak position are preserved. Another interesting feature is that excitation with the line 350 nm from the Kr+ laser leads to the 5Do→7F1 transition as the most intense, unlike the PL for the films shown in Fig. 2 where the transition 5Do→7F2 turned out as the most intense. It is also different from excitation with the fourth harmonic of a Nd:YAG laser (266 nm) for the same sort of high doping samples, where the transition 5Do→7F2 also turned out as the most intense . The asymmetric ratio, defined by the ratio between the lines (5D0 → 7F2) / (5D0 → 7F1), is very low in the case of pressed SnO2 powder shown in Fig. 4, indicating dominant emission from excited ion located at symmetric sites, for example Eu3+ substitutional to Sn4+ in the SnO2 lattice [25–27]. It is rather unexpected for 2 at% Eu3+ sample where the low solubility (0.05–0.06 at.% ) yields a very high Eu concentration at asymmetric sites, for instance, at particles surface. It may be explained by a rather efficient excitation of symmetric sites provided by the 350 nm line of the Kr+ laser, which is much closer to the SnO2 bandgap, compared to the 266 nm line, being responsible for more efficient energy transfer from SnO2 band-to-band transition to substitutional Eu3+ ions. The broad band is no longer seen in this case. The annealing must be providing the diffusion of adsorbed oxygen species which recombines with oxygen vacancies, decreasing the number of available vacancies to electron transition between the donor level and the acceptor level formed by Eu3+ ions, in good agreement with Wang and associates . Besides, the consequence of overall lattice reorganization provided by the thermal annealing at longer time and higher temperature certainly is decreasing the number of lattice defects, decreasing the concentration of the non-intentionally introduced luminescent centers, as predicted by Kim and associates .
In order to fully understand the rule of PL emission for this sort of sample, a film was deposited on quartz substrate and thermally treated at 1000°C by 1 hour, an identical annealing condition of the powders. The soda-lime glass substrate would not support such a high temperature. Results of PL and X-ray diffraction are shown in Fig. 5.The Eu3+ transitions show up and the broad band is not destroyed in this case, moreover, it becomes even higher as compared to PL curves shown in Fig. 2. Then, one may conclude that the broad band is a characteristic of films. Besides, the Eu3+ dominant band changes, because, like the powder, the 5Do→7F1 transition is the most intense, unlike the PL for the heterostructure films deposited at low temperature (Fig. 2) where the transition 5Do→7F2 turned out as the most intense. Looking carefully at Fig. 2 it seems that the transition 5Do→7F2 intensity decreases for higher temperature an then, it may be postulated that it changes at some temperature for the 5Do→7F1 transition, which is the dominant for the SnO2:2at%Eu film, deposited on quartz substrate. The existence of a temperature where the relative dominance of each of these Eu3+ transition changes from 5Do→7F2 to 5Do→7F1 could be accomplished with higher number of samples, and could be done in the future. The broad band peaks at 445 nm, preserving the tendency for the blue shift observed for the heterostructures PLs. Concerning the diffraction patterns of samples treated at the same temperature (1000°C) it can be observed that for powders, the peak related to (110) plane is the most intense whereas for the film the most intense is related to (101) plane. This change in the arrangement is probably consistent with the broad band appearing in the film and not in the powder.
Table 1 shows the crystallite size evaluated by the Scherrer equation  based on the diffraction data of for the direction (101) (most intense for all the SnO2 films) shown in Figs. 3 and 5, beside the calculated crystallite size from the diffraction data of the 2at%Eu doped SnO2 powder. As can be seen in Table 1, as the thermal annealing temperature increases, the crystallite size also increases. The crystallite size varies from 7.3 nm for the GaAs/SnO2:2%Eu treated at 200°C to 22.1 nm for the SnO2:2%Eu powder, treated 1000°C by 1 hour. Comparing the spectra shown in Figs. 2, 4 and 5 with the crystallite size evaluated in the direction (101), it is possible to identify a crystallite size increase with the broad band peak position shifting to higher energy (lower wavelength) in the case of films, but the broad band vanishes for the highest crystallite size, which was found for the SnO2:2%Eu powder. Although the small size of the building blocks (crystallites) points towards a quantum confinement, the expected behavior for the confinement is that smaller crystallites would force the emission to higher energies , in opposition to what was observed here. Then, the most plausible explanation is that the decreasing number of available vacancies to electron transition between the donor level and the acceptor level formed by Eu3+ ions, makes the broad band disappears for the case of the powder, even though the film deposited on quartz does not respect this rule.
The luminescence peaks shown here can be separated in two distinct phenomena: (1) Eu3+ transitions which are observed only in heterojunction samples, when annealed at low temperature, (2) the broad band, which is observed only in films. These situations are sketched in Fig. 6.Although the crystallite size has influence in both types of emissions, the influence is rather different. Figure 6(a) shows the differences on the dominant Eu3+ transition depending on the location of the ion. If the crystallite size is small, the ion is located preferentially at crystallite boundary layer, as shown at left in the top figure, and the PL emission is as seen at right, for the heterojunction thermally treated at 200°C/1 hour, with the 5Do→7F2 transition as dominant. When the crystallite grows, the Eu3+ ions become located inside the crystallite, substitutional to Sn4+, as seen at left of the bottom of Fig. 6(a). In this case, the PL structure is as seen at right, obtained for SnO2 film deposited on quartz, with the 5Do→7F1 transition as dominant. The high annealing temperature also allows splitting the emission lines . Figure 6(b) shows that as the crystallite size decreases, the bandgap energy becomes wider . However, it increases the disorder at neighborhood of defects, creating a wider distribution of defect levels inside the bandgap. Considering that the broad band may be generated for the electron transfer between oxygen vacancies (donors) and Eu3+ levels (acceptors) , the larger distribution of levels caused by the more disordered neighborhood of both sort of defects, may lead to a lower energy transition for smaller crystallites, even though the bandgap is larger, as seen in Fig. 6(b).
In order to explain the low temperature PL data of heterostructure films (Fig. 2), the samples were submitted to an Energy Dispersive X-Ray (EDX) analysis throughout the sample surface. This EDX-scanning on the surface was carried out considering that the large particles observed in the sample surface (inset of Fig. 1) suggested the possible existence of Eu-concentrated regions. It is a fair explanation for the PL data, because the Eu-concentrated regions could be selectively excited by the Kr+ laser, emitting the observed bands. It also would be in good agreement with 5Do→7F2 transition as the most intense for the heterostructures, since this transition is characteristic of Eu3+ ions located in rather asymmetric sites. Figure 7 is representative of a detailed SEM analysis carried out for the sample GaAs/SnO2:2at%Eu, annealed at 400°C by 20 minutes. Table 2 shows the relative composition of Sn and Eu, measured in the marked regions of Fig. 7, obtained by the EDX analysis. The most external particle, which will be called as “balls” here, may be regions where the Eu ions are concentrated. The EDX analysis, shown in Table 2, displays indeed variation of up to 700% in the relative Eu/Sn composition, among measured sites. Although it not possible to ensure that electrons interacted only with the elements within the demarked area under this magnification and so high voltage, the balls seem to be very important in this variation, for instance regions 17, 18 and 19, which are exactly around the balls, present high Eu concentration. There are exceptions, as region 24, where the concentration of Eu is the highest and Region 20, on the “ball”, where the concentration of Eu is the lowest. In the former case the demarked area is located in the neighborhood of several balls and there may be buried balls under the chosen area as the figure shows irregularities at the surface of this region. However, this is not the formal rule, because there are a few balls not presenting a very high Eu concentration.
Figure 8 shows the relative Eu/Sn concentration along a selected line crossing several balls. Although the Eu3+ - local concentration peaks do not happen for the centers of all of the balls, it is clear that the presence of a ball has a relation to the Eu distribution. Besides, the dominance of the Eu transition of asymmetric sites is in very good agreement with the disorder of the surface.
Figure 9 shows SEM on the SnO2 film (deposited directly on glass substrate) surface, to compare with the SEM of the heterojunction surface, along with SEM of the heterojunction GaAs/SnO2:2%Eu thermally annealed at 200°C by 1 hour. The SEM picture of the SnO2:2%Eu thin film (Fig. 9(a)) shows a more uniform surface when compared to the SEM picture of the heterojunction GaAs/SnO2:2%Eu annealed at 400°C by 20 minutes (Fig. 7), which are in very good agreement with the SEM of the heterojunction GaAs/SnO2:2%Eu thermally annealed at 200°C by 1hour (Fig. 9(c)). In this heterojunction, the regions called here as “balls” are clearly seen again in the SEM picture, as prominent particles in the surface. Figure 9 also brings the Electron Backscatter Diffraction (EBSD) done in the same SEM equipment. The great novelty in these results are the possibility of identification of atoms with high atomic number, such as Eu or Sn (atomic number 63 e 50, respectively), which are seen as clearer spots in the EBSD image, whereas O, Ga and As (atomic number 8, 31 and 33, respectively), are darker. Since Eu and Sn have close atomic numbers, it is not possible to identify with precision where each of these ions is preferentially located in the clear spots. However, comparing Figs. 9(b) and 9(d), it is easily seen that the regions around the balls are differentially populated comparing with the rest of the surface. It reinforces the hypothesis that the existence of these ball regions provides some sort of Eu3+ concentrated regions in the heterojunction surface, which are responsible for the luminescence from the Eu3+ ions, seen in the heterojunction and not in the SnO2 film annealed at lower temperatures. As confirmed in Figs. 9(b) and 9(d), there are not such “balls” in the SEM image of the SnO2 film deposited on glass substrate, neither concentrated regions of heavier atoms in the EBSD image.
The luminescence from Eu-doped oxide, in the form of thin films is a rather convenient way, providing accessibility to construction and operation of luminescent and electroluminescent devices, rather than powders. The combination with GaAs in the form of heterojunction films provides an interesting emission, not seen from Eu-doped films deposited in an isolated way directly on glass substrates, when thermally annealed at lower temperatures. Although, the transition was also observed for a SnO2 film deposited on quartz substrate, it was needed a much higher annealing temperature, and a much more expensive substrate, making the process of producing photoluminescent films unviable in this case.
Doping SnO2 thin films with Eu3+ leads to interesting luminescent emission from the rare-earth ion when the doped tin dioxide film is grown on the top of a GaAs layer, unlike the SnO2 deposition directly on a glass substrate. Eu transitions are clearly identified when the samples are in the form of pressed powder (pellets), thermally treated at much higher temperatures. However, in the form of films, the Eu emission comes along a broad band which is present in all the films and blue shifts as the thermal annealing temperature as well as the crystallite size increase. Although quantum confinement seems to be an explanation, this evidence is in opposition to what would be expected, where the energies and intensity should rise for smaller nanocrystallites. Then, the most feasible contribution for the detected broad band is related to electron transfer from oxygen vacancies and acceptor Eu3+ ions, or nanocrystalline defects, originated from the disorder in the material. These phenomena relax for annealing at higher temperatures and longer times, in the case of pellets, and eliminate the broad band, but it does not happen in the case of films.
The great relevance of this result is the possibility of obtaining luminescence from films, a much more convenient way of making light emitting devices, which can be enhanced by the making of a heterojunction with GaAs.
The authors would like to thank the financial support of the Brazilian agencies: CNPq and FAPESP. XRD was performed at Multiuser Laboratory at UNESP/DF Campus Bauru.
References and links
1. S. Coffa, G. Franzò, F. Priolo, A. Polman, and R. Serna, “Temperature dependence and quenching processes of the intra-4f luminescence of Er in crystalline Si,” Phys. Rev. B Condens. Matter 49(23), 16313–16320 (1994). [CrossRef] [PubMed]
2. M. Ishii, S. Komuro, and T. Morikawa, “Study on atomic coordination around Er doped into anatase and rutile TiO2:Er-O clustering dependent on the host crystal phase,” J. Appl. Phys. 94(6), 3823 (2003). [CrossRef]
3. S. C. Ray, M. K. Karanjai, and D. Dasgupta, “Tin dioxide based transparent semiconducting films deposited by the dip-coating technique,” Surf. Coat. Tech. 102(1), 73–80 (1998). [CrossRef]
4. E. Dien, J. M. Laurent, and A. Smith, “Comparison of optical and electrical characteristics of SnO2-based thin films deposited by pyrosol from different tin precursors,” J. Eur. Ceram. Soc. 19(6–7), 787–789 (1999). [CrossRef]
5. H. Peng, H. Song, B. Chen, J. Wang, S. Lu, X. Kong, and J. Zhang, “Temperature dependence of luminescent spectra and dynamics in nanocrystallineY2O3:Eu3+,” J. Chem. Phys. 118(7), 3277–3282 (2003). [CrossRef]
6. E. A. Morais, L. V. A. Scalvi, A. Tabata, J. B. B. De Oliveira, and S. J. L. Ribeiro, “Photoluminescence of Eu3+ ion in SnO2 obtained by sol–gel,” J. Mater. Sci. 43(1), 345–349 (2008). [CrossRef]
7. T. W. Kim, D. U. Lee, and Y. S. Yoon, “Microstructural, electrical, and optical properties of SnO2 nanocrystalline thin films grown on InP (100) substrates for applications as gas sensor devices,” J. Appl. Phys. 88(6), 3759–3761 (2000). [CrossRef]
8. A. P. Alivisatos, “Semiconductor Clusters, Nanocrystals and Quantum Dots,” Science 271(5251), 933–937 (1996). [CrossRef]
9. R. S. D. Bella and K. Navaneethakrishnan, “Donor binding energies and spin–orbit coupling in a spherical quantum dot,” Solid State Commun. 130(11), 773–776 (2004). [CrossRef]
10. G. Ledoux, J. Gong, F. Huisken, O. Guilois, and C. Reynaud, “Photoluminescence of size-separated silicon nanocrystals: Confirmation of quantum confinement,” Appl. Phys. Lett. 80(25), 4834–4836 (2002). [CrossRef]
11. T. Y. Kim, N. M. Park, K. H. Kim, G. Y. Sung, Y.-W. Ok, T.-Y. Seong, and C.-J. Choi, “Quantum confinement effect of silicon nanocrystals in situ grown in silicon nitride films,” Appl. Phys. Lett. 85(22), 5355–5357 (2004). [CrossRef]
12. D. H. Park, Y. H. Cho, Y. R. Do, and B. T. Ahn, “Characterization of Eu-Doped SnO2 Thin Films Deposited by Radio-Frequency Sputtering for a Transparent Conductive Phosphor Layer,” J. Electrochem. Soc. 153(4), H63–H67 (2006). [CrossRef]
13. H. Guo, Y. Zhu, S. Qiu, J. A. Lercher, and H. Zhang, “Coordination Modulation Induced Synthesis of Nanoscale Eu1-xTbx-Metal-Organic Frameworks For Luminescent Thin Films,” Adv. Mater. 22(37), 4190–4192 (2010). [CrossRef] [PubMed]
14. C. M. Leroy, T. Cardinal, V. Jubera, C. Aymonier, M. Treguer-Delapierre, C. Boissière, D. Grosso, C. Sanchez, B. Viana, and F. Pellé, “Luminescence properties of ZrO2 mesoporous thin films doped with Eu3+ and Agn,” Micro. Mes. Mat. 170, 123–130 (2013). [CrossRef]
15. J. W. E. Wiegman and E. van der Kolk, “Building integrated thin film luminescent solar concentrators: Detailed efficiency characterization and light transport modeling,” Sol. Energy Mater. Sol. Cells 103, 41–47 (2012). [CrossRef]
16. M. Yu, J. Lin, J. Fu, H. J. Zhang, and Y. C. Han, “Sol–gel synthesis and photoluminescent properties of LaPO4:A (A - Eu3+, Ce3+, Tb3+) nanocrystalline thin films,” J. Mater. Chem. 13(6), 1413–1419 (2003). [CrossRef]
17. K. L. Frindell, M. H. Bartl, M. R. Robinson, G. C. Bazan, A. Popitsch, and G. D. Stucky, “Visible and near-IR luminescence via energy transfer in rare earth doped mesoporous titania thin films with nanocrystalline walls,” J. Solid State Chem. 172(1), 81–88 (2003). [CrossRef]
18. T. F. Pineiz, L. V. A. Scalvi, M. J. Saeki, and E. A. Morais, “Interface Formation and Electrical Transport in SnO2:Eu3+/GaAs Heterojunction Deposited by Sol–Gel Dip-Coating and Resistive Evaporation,” J. Electron. Mater. 39(8), 1170–1176 (2010). [CrossRef]
19. T. F. Pineiz, E. A. de Morais, L. V. A. Scalvi, and C. F. Bueno, “Interface formation of nanostructured heterojunction SnO2:Eu/GaAs and electronic transport properties,” Appl. Surf. Sci. 267, 200–205 (2013). [CrossRef]
20. B. D. Cullity and R. Stock, Elements of X-Ray Diffraction (Prentice Hall, 2001).
21. Y. Wang, J. Ma, F. Si, X. Yu, and H. Ma, “Structural and photoluminescence characters of SnO2:Sb films deposited by RF magnetron sputtering,” J. Lumin. 114(1), 71–76 (2005). [CrossRef]
22. V. Geraldo, V. Briois, L. V. A. Scalvi, and C. V. Santilli, “EXAFS investigation on Sb incorporation effects to electrical transport in SnO2 thin films deposited by sol–gel,” J. Eur. Ceram. Soc. 27(13-15), 4265–4628 (2007). [CrossRef]
23. V. Geraldo, V. Briois, L. V. A. Scalvi, and C. V. Santilli, “Structural Characterization of Nanocrystalline Sb-Doped SnO2 Xerogels by Multiedge X-ray Absorption Spectroscopy,” J. Phys. Chem. C 114(45), 19206–19213 (2010). [CrossRef]
24. E. A. Morais, L. V. A. Scalvi, A. A. Cavalheiro, A. Tabata, and J. B. B. Oliveira, “Rare earth centers properties and electron trapping in SnO2 thin films produced by sol–gel route,” J. Non-Cryst. Solids 354(42-44), 4840–4845 (2008). [CrossRef]
25. G. E. S. Brito, S. J. L. Ribeiro, V. Briois, J. Dexpert-Ghys, C. V. Santilli, and S. H. Pulcinelli, “Short Range Order Evolution in the Preparation of SnO2 Based Materials,” J. Sol-Gel Sci. Technol. 8(1–3), 261–268 (1997). [CrossRef]
26. S. J. L. Ribeiro, S. H. Pulcinelli, and C. V. Santilli, “SnO2:Eu nanocrystallites in SnO2 monolithic xerogels,” Chem. Phys. Lett. 190(1–2), 64–66 (1992). [CrossRef]
27. E. A. Morais, S. J. L. Ribeiro, L. V. A. Scalvi, C. V. Santilli, L. O. Ruggiero, S. H. Pulcinelli, and Y. Messaddeq, “Optical characteristics of Er –Yb doped SnO xerogels,” J. Alloy. Comp. 344(1-2), 217–220 (2002). [CrossRef]
28. T. Matsuoka, T. Tohda, and T. Nitta, “The Low-Energy-Electron (LEE) Excitation of SnO2:Eu Powder Phosphor; Fundamental Characteristics,” J. Electrochem. Soc. 130(2), 417–423 (1983). [CrossRef]
29. R. E. Marotti, P. Giorgi, G. Machado, and E. A. Dalchiele, “Crystallite size dependence of bandgap energy for electrodeposited ZnO grown at different temperatures,” Sol. Energy Mater. Sol. Cells 90(15), 2356–2361 (2006). [CrossRef]