1. Introduction

Rare-earth doped semiconductors have increased interest due to applications in optoelectronic, such as displays and optical communication devices [1]. Doping wide bandgap semiconductors with rare-earth ions reduces the thermal quenching effects [2], 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 [3]) 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 [5]. Photoluminescence (PL) properties of Eu-doped SnO₂ xerogels have been investigated under different lasers excitation [6], 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 [6] 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 [7] 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 [8]. 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 [11].

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 [12]. 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 [6]. 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 K_α radiation source (1.5405 Å) and a Ni filter for reducing the undesirable K_β 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.

 

Fig. 1 Schematic diagram of heterojunction sample. Bottom inset: band energy diagram of the heterostructure interface. Upper left inset: SEM image of the GaAs film surface. Upper right inset: SEM image of the GaAs/SnO₂:2at%Eu film surface.

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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.

 

Fig. 2 Luminescence spectra of GaAs/SnO₂:2at%Eu thermally annealed at 200°C by 1 hour, and at 400°C by 20 minutes, as well as SnO₂:2at%Eu film treated at 500°C by 1 hour. Excitation: line 350 nm of the Kr⁺ laser.

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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.

 

Fig. 3 X-ray diffraction patterns of GaAs/SnO₂:2at%Eu thermally annealed at 200°C by 1 hour, and at 400°C by 20 minutes, and SnO₂:2at%Eu film treated at 500°C by 1 hour.

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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 [20], 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 [8].

SnO₂ nanocrystalline thin films grown on p-type InP (100) [7] 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 [7]. 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) [21]. 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 [21], 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 [21]. 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 [24].

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.

 

Fig. 4 Luminescence spectra of SnO₂:2at%Eu and SnO₂:4at%Eu xerogels treated at 1000°C. Excitation: line 350 nm of the Kr⁺ laser. Inset: X- ray diffractogram of SnO₂:4at%Eu powder (previous of being pressed down to pellet format). Blue lines are from the file JCPDS- 41-1445.

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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 ⁵D_o→⁷F₁ transition as the most intense, unlike the PL for the films shown in Fig. 2 where the transition ⁵D_o→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 ⁵D_o→⁷F₂ also turned out as the most intense [6]. 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.% [28]) 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 [21]. 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 [7].

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 ⁵D_o→⁷F₁ transition is the most intense, unlike the PL for the heterostructure films deposited at low temperature (Fig. 2) where the transition ⁵D_o→⁷F₂ turned out as the most intense. Looking carefully at Fig. 2 it seems that the transition ⁵D_o→⁷F₂ intensity decreases for higher temperature an then, it may be postulated that it changes at some temperature for the ⁵D_o→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 ⁵D_o→⁷F₂ to ⁵D_o→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.

 

Fig. 5 (a) Photoluminescence of a Eu 2at%-doped SnO₂ thin film deposited on quartz substrate and thermally annealed at 1000°C by 1 hour. Excitation: line 350 nm of the Kr⁺ laser. (b) X-ray diffraction pattern for this film. SnO₂ planes are from the file JCPDS- 41-1445.

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Table 1 shows the crystallite size evaluated by the Scherrer equation [20] 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 [8], 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.

Table 1. – Crystallite size for the thin film samples and powder, evaluated by the Scherrer equation

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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 ⁵D_o→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 ⁵D_o→⁷F₁ transition as dominant. The high annealing temperature also allows splitting the emission lines [16]. Figure 6(b) shows that as the crystallite size decreases, the bandgap energy becomes wider [29]. 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) [21], 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).

 

Fig. 6 (a) Eu³⁺ ion located preferentially at crystallites surface (top) and the PL structure is as shown at right, for the heterojunction treated at 200°C/1 hour; or Eu³⁺ ion is located prefentially in the middle of the larger crystallite (bottom), and the PL is as obtained for the SnO₂ film on quartz (right). (b) at left, it is the energy band diagram for the heterojunction thermally annealed at 200°C/1hour, and at right, it is the energy band diagram for SnO₂ film on quartz.

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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 ⁵D_o→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.

 

Fig. 7 SEM image (magnitude 10000x) for the sample GaAs/SnO₂:2%Eu annealed at 400°C by 20 minutes, showing the marked regions where the EDX analysis was carried out.

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Table 2. Relative composition of Sn and Eu obtained by EDX with magnitude 10000x, for the sample GaAs/SnO₂:2%Eu annealed at 400°C by 20 minutes.

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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.

 

Fig. 8 Relative Eu/Sn composition obtained by scanning EDX-signal along the yellow line drawn in SEM image shown in the inset, for the heterostructure GaAs/SnO₂:2%Eu annealed at 400°C by 20 minutes.

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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.

 

Fig. 9 SEM (magnitude 10000x) of: (a) SnO₂:2%Eu and (c) heterojunction GaAs/SnO₂:2%Eu annealed at de 200°C / 1hour. EBSD carried out in the SEM microscope: (b) SnO₂:2%Eu thin film and (d) heterojunction GaAs/SnO₂:2%Eu thermally annealed at de 200°C / 1hour

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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.

4. Conclusions

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.