1. Introduction

The realization of photonic interconnects requires the development of new efficient and reliable on-chip optical device elements, including light sources, modulators, amplifiers, buffers, switches, and detectors. The transition from electronic to photonic interconnects will be cost effective only if it is performed with mature CMOS technology. This requires silicon compatibility of all the integrated optical components. A significant amount of work has already been done and different approaches to the realization of each of the mentioned components have been attempted: microdetectors based on III-V structures evanescently coupled to a silicon waveguide [1], Ge-on-Si [2, 3], and ion-implanted silicon [4] have been demonstrated; modulators based on the free carrier plasma dispersion effect [5] and electro-absorption effect [6–8] have been realized; and hybrid evanescent amplifiers [9], and switches [10] have been shown. The realization of a silicon-based light source is considered to be one of the most challenging tasks. Due to the indirect silicon band-gap, the emission efficiency of silicon is low and, even at liquid helium temperatures does not exceed ~10-4% [11]. Consequently, alternative approaches are being considered, such as silicon Raman lasers [12], hybrid evanescent lasers [13], devices based on germanium-on-silicon [14], silicon-nanocrystal doped SiO₂ [15, 16], silicon nanocrystal doped silicon nitride [17], erbium-doped Si-rich silicon nitride [18], erbium doped porous silicon [19–22], and erbium-doped silicon-rich SiO₂ [23, 24].

One of the possible solutions for the realization of silicon based light source is to use the emission from Er-silicates. Er silicates (Er₂SiO₅ and Er₂Si₂O₇) have been attracting considerable attention as Er based materials for small size and high optical gain light source in silicon photonics integration, because they contain a higher Er density of 10²² cm⁻³ than Er doped Si-based materials [25, 26]. H. Omi et al. have studied the mixture formation of Er_xYb_2-xSi₂O₇ and Er_xYb_2-xO₃ on silicon for broadening the C-band in an optical amplifier [27]. However, such a high concentration of Er results in up-conversion due to closely neighboring Er ions which limits the Er luminescence. Therefore, it is necessary to characterize and control the distance between Er ions in such Er silicates. An effective strategy for reducing this up-conversion is to incorporate yttrium (Y) cations into the structure, where they substitute Er ions in the silicate lattice and prevent neighboring Er ions from causing up-conversion due to the similar ionic radius between Y and Er [28, 29].

Scandium ions (Sc³⁺), on the other hand, are small (ionic radius = 0.75 Å) than erbium (Er³⁺) (ionic radius = 0.881 Å). Generally, this can result in enhancing the crystal field strength for Er doped silicates and oxides [30, 31]. In fact, Fornasiero et al. synthesized single crystal of Er doped Sc silicates using the Czochralski technique with the idea that Sc³⁺ ions would increase the Stark-splitting of the thermally populated erbium ground state as well as that of other electronic energy levels of the silicates and thereby reduce re-absorption losses [31]. A. Najar et al. has fabricated Er_xSc_2-xSi₂O₇ compound after annealing at 1250 °C and estimated the erbium diffusion coefficient to be 1x10⁻¹⁵ cm²/s with a peak emission at 1537 nm [32]. However, the effect of thermal annealing and Sc ions on the structural and optical properties of Er-Sc silicates with different compositions have not been sufficiently characterized yet, compared with those of Er-Y or Er-Yb silicates.

In this work, we synthesized a polycrystalline Er-Sc silicate and discilicate compounds in which Er and Sc are homogeneously distributed using RF-sputtering with multilayer Er₂O₃, Sc₂O₃, and SiO₂ deposited on SiO₂/Si (100) substrate and thermal annealing at high temperature. The concentration profiles of Er³⁺ and Sc³⁺ and the morphology structures were investigated by RBS, TEM-SAED and X-ray. The diagram energy of Er³⁺ ions in Er_xSc_2-xSi₂O₇ phase is presented based in the optical properties studies coupled with the morphology structure. The excitation cross-section will be determined and discussed.

2. Experimental conditions

Er-Sc multilayer thin films were grown by RF-sputtering by alternating 15 nm thick layers of Er₂O₃, and Sc₂O₃ separated by 15 nm thick SiO₂ layer. These layers were deposited on 50 nm thick Er₂O₃ on SiO₂ (1.3 µm)/Si (100) substrate at room temperature. After deposition, the samples were annealed in O₂ at 900 °C, 1000 °C, 1100 °C, or 1250 °C for 1 h. The concentration of different chemical elements in the films was determined by Rutherford backscattering spectroscopy (RBS). He⁺⁺ ions delivered by a 2.275 MeV Van de Graf accelerator with the normal detector angle of 160° was used. Spectra are fitted by applying a theoretical layer model and iteratively adjusted elemental concentrations and thickness until good agreement is found between the theoretical and the experimental spectra. The samples were analyzed from transmission electron microscopy (TEM) images obtained at 200 KeV, and by synchrotron grazing incidence X-ray diffraction (GIXD) experiments performed on the as-grown and annealed samples at the BL24 in SPring-8 using an X-ray wavelength of 1.24 Å and an incidence angle of 1.0◦. Photoluminescence (PL) and photoluminescence excitation (PLE) spectra were obtained by using two excitation sources at 980 and 1527.6 nm. The excitation laser was focused to a spot with a diameter of about 15 µm and an incident angle of 45 degrees through an objective lens. The luminescence from the sample was collected perpendicularly by a different objective lens with a numerical aperture of 0.40. The PL spectra were detected using a 0.5-m spectrometer and cooled InGaAs detector. Time-resolved PL was measured by using a pulsed laser formed by an acousto-optic modulator and a streak camera (Hamamatsu C11293S). The spectral and temporal resolutions of this system are 500 µeV and 20 ps (for a 1 ns time range), respectively.

3. Results and discussion

To determine the concentration distribution of chemical elements in the samples, we performed RBS measurements. RBS spectra in Fig. 1 show the profiles of Er, Sc, Si, and O atom distributions after deposition and annealing at 1100 and 1250 °C. As shown in Fig. 1, RBS signal intensities changed with increasing annealing temperature, suggesting that the concentration changes for Er³⁺ and Sc³⁺. Note that for the as-grown sample, the Er distribution is formed by peaks corresponding to Er₂O₃ layers. Whereas, the intensity decrease for sample annealed at 1100 °C and erbium diffuse in-depth. The concentration of Si and O has an atomic ratio of Si to O of about 1:4, indicating that the films are mainly formed by two phases Er_xSc_2-xSiO₅ and Er_xSc_2-xSi₂O₇.

 

Fig. 1 RBS signal of as-grown and after annealing at 1100 and 1250 °C for 1 h in O₂.

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In contrast, at 1250 °C erbium starts to penetrate the silicate matrix and becomes more homogeneous. From the RBS analysis, the concentration of Si and O, has an atomic ratio of 2:7, indicating that the films are mainly composed of the Er_xSc_2-xSi₂O₇ phase. The RBS analysis indicates erbium and scandium doping concentrations close to [Er] = 12.7x10²⁰, and [Sc] = 4.1 x10²⁰ at/cm³ at 1100 °C and [Er] = 6.7x10²¹, and [Sc] = 1.4x10²¹ at/cm³ at 1250 °C.

To determine the microscopic structures after deposition, we performed TEM analysis of the cross-section coupled with selected area electron diffraction (SAED) images of the samples after deposition and annealing at 1100 °C and 1250 °C. Figure 2(a) shows a schematic cross-section of the different layers. The cross-sectional image in Fig. 2(b) obtained after deposition shows different layers of Er₂O₃, Sc₂O₃, and SiO₂ with a total deposition thickness of around 109 nm. The inset SAED image from the Er₂O₃ layer at the bottom shows multicrystalline rings. The interplanar spacings d are about 1.29, 1.32 and 1.52 Å, corresponding respectively to the (203), (440), and (20-3) planes for Er₂Si₂O₇, and with d of 1.32 and 1.52 Å corresponding to the (800) and (444) planes for Er₂O₃.

 

Fig. 2 (a) The fabricated structure. (b)-(c) and (d) cross-sectional TEM images after deposition and after annealing at 1100 and 1250 °C, respectively, with SAED images in insets.

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For the sample annealed at 1100 °C, the TEM image in Fig. 2(c) shows two layers separated by some voids. The appearance of these voids may be due to the diffusion of SiO₂ layer. The interplanar spacings are about 1.7, 2.83 and 1.88 Å, corresponding respectively to the (233), (−402), and (−523) planes for Er₂SiO₅, and the (132), (130), and (22-2) planes for Er₂Si₂O₇. These results show the mixture formation of Er_xSc_2-xSiO₅, Er_xSc_2-xSi₂O₇ phases at 1100 °C.

After thermal annealing at 1250 °C, we formed a unique layer with an average thickness of 102 nm as shown in Fig. 2(d). Not that the size of the voids becomes very small at 1250 °C compared to 1100 °C. The SAED images show a single crystal compound. The interplanar spacings are 1.30, 1.54 and 2.61 Å, corresponding respectively to the (203), (33-2), and (220) planes for Er₂Si₂O₇, and Er₂SiO₅ phase with (−442) plane for 1.54 Å spacing. TEM images show for both annealing temperatures, the existence of Er_xSc_2-xSiO₅ and Er_xSc_2-xSi₂O₇ phases in the structures.

To confirm the dominant phase after annealing, we performed GIXD experiments on the as-grown and annealed samples. GIXD profiles of the crystalline structure after deposition and annealing are shown in Fig. 3. The GIXD profile of the sample after deposition shows the presence of Er₂O₃, Sc₂O₃, Er₂SiO₅, Sc₂SiO₅, and Er₂Si₂O₇ in the films. After annealing at 1100 °C, the main intense peaks are close to Er₂SiO₅ and Sc₂SiO₅ with two peaks also close to Er₂Si₂O₇. This result confirms the TEM and RBS measurements showing the presence of both Er_xSc_2-xSiO₅ and Er_xSc_2-xSi₂O₇ phases. The size of the polycrystalline Er_xSc_2-xSiO₅ is around 39 nm at 1100 °C according to the analysis of the diffraction peak width, whereas is equal to 50 nm for Er_xSc_2-xSi₂O₇ polycrystalline.

 

Fig. 3 GIXD profiles obtained from samples after deposition and after annealing at 1100 and 1250 °C for 1 h in O₂, with the Joint Commission for Powder Diffraction Standards (JCPDS) corresponding to the different compound.

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After annealing at 1200 °C (data not shown here) and 1250 °C, we have only Er₂Si₂O₇ because of diffusion of Er and Sc into different layers and the formation of new polycrystalline mixed compounds assigned to Er_xSc_2-xSi₂O₇, which is the dominant phase. The size of the polycrystalline Er_xSc_2-xSi₂O₇ is around 35 nm at 1250 °C according to the analysis of the diffraction peak width.

To demonstrate a relation between composition stoichiometry and optical properties, we analyzed the annealed samples by PL. Figure 4 shows PL spectra of samples annealed between 900 and 1250 °C using λ_ex = 980 nm at room temperature. For annealing temperature equal to 900 and 1100 °C, the spectra exhibit several emission peaks, corresponding to the ⁴I_13/2 to ⁴I_15/2 transition of Er³⁺ ions. The main intense peak is observed at 1528 nm and the second intense peak is at 1537 nm at 1100 °C. The PL spectrum with a single main intense peak at 1528 nm is similar to the PL spectrum of the Er silicates [26]. The structure of the Er silicates has been confirmed to comprise Er₂SiO₅ crystalline films with a highly ordered layer structure [26, 33]. In our case, this peak corresponds to the Er_xSc_2-xSiO₅ phase and confirms the previous observations at 1100 °C.

 

Fig. 4 PL spectra at room temperature obtained from the sample annealed from 900 to 1250 °C with λ_ex = 980 nm.

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However, the samples annealed at 1200 °C and 1250 °C exhibit PL spectra with a different shape compared to 900 and 1100 °C, with the main intense peak at 1537 nm with narrower linewidth, indicating that a relatively homogeneous crystalline environment for Er³⁺ ions, and with sub-peaks at 1546.2 and 1551 nm. The PL peak intensity of annealed samples reaches the maximum at 1250 °C. The peak at 1537 nm appears as the main peak for the films at 1200 and 1250 °C, while the 1528 nm main PL peak typical of Er_xSc_2-xSiO₅ is observed for the films with low intensity. It is considered that the 1537 nm peak come from another Er silicate structure formed by the reaction of the Er–O and Si–O precursors with the SiO₂ inter-layers. We consider that high temperature annealing at 1200 and 1250 °C for 1 h enhances the reaction of the Er–O and Si–O precursors with the SiO₂, converting the most of Er_xSc_2-xSiO₅ to Er_xSc_2-xSi₂O₇. The surface morphology of the films has observed by AFM, which revealed that the surface is still flat, smooth, and noncracking after annealing up to 1250 °C [32].

To determine the Stark energy levels of the Er_xSc_2-xSi₂O₇ phase, PLE and PL tests were conducted at 4 K in Fig. 5. Figure 5(a) shows a PLE color plot under ⁴I_13/2 manifold excitation. The dashed lines in this figure indicate the assigned energy levels of the ⁴I_13/2 manifold measured for Er_xSc_2-xSi₂O₇ in the polycrystalline phase. The PLE spectrum of the sample annealed at 1250 °C in Fig. 5(b) shows peaks at 1512, 1519, and 1529 nm corresponding to the direct absorption mechanism with an FWHM of 2, 2.5, and 2.25 nm, respectively. The spectral linewidths of the PL emissions are 1.1, 1.76, and 1.37 nm, which correspond respectively to the emission peaks at 1537, 1546, and 1551 nm.

 

Fig. 5 (a) PLE color plot measured at 4 K obtained from the sample annealed at 1250 °C. (b) PLE and PL spectra at 4K.

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From the collected spectra, we obtained the energy level schematic in Fig. 6 corresponding to transition between the ⁴I_15/2 and ⁴I_13/2 manifolds of Er in Er_xSc_2-xSi₂O₇ for the sample annealed at 1250 °C. The excited state ⁴I_13/2 level of Er³⁺ is composed of four levels instead of seven levels, and the ground state ⁴I_15/2 level is composed of two levels instead of eight levels. The six component peaks corresponding to Y₂-Z₁, Y₃-Z₃, Y₁-Z₁, Y₃-Z₁, Y₁-Z₂, and Y₁-Z₃ transitions are presented in Fig. 6.

 

Fig. 6 Schematic illustration of main optical transitions of Er³⁺ in Er_xSc_2-xSi₂O₇. The observed absorption and emission transitions are indicated in red and blue, respectively.

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These level energies of Er_xSc_2-xSi₂O₇ are different from Er:Sc₂O₃ and Er in scandium silicate [34, 35]. J. H. Shin et al. have shown a peak emission at 1529 nm with an FWHM of 11 nm for Er_xY_2-xSiO₅ annealed at 1200 °C using an excitation wavelength of 488 nm [36, 37]. In addition, M. Miritello et al. obtained a peak emission at 1535 nm for α-(Yb_1-xEr_x)₂Si₂O₇ with a 37 nm FWHM using 532 nm excitation wavelength after annealing at 1200 °C [34]. Furthermore, the energy levels were different from Er₂O₃ with strong luminescence from Y₂-Z₁ and Y₁-Z₁ transitions in C_3i site observed at wavelengths of 1544.7 and 1548.2 nm [38].

The split of Stark levels are due to the smaller Sc³⁺ radius’s enhancing the crystal field strength, which affects the luminescence properties with a smaller FWHM compared to the effect of Y ions. However, Er can be substituted with Y ions in the silicate phase, but Sc has additional effect due to it is smaller radius compared to Er and Y.

The scalar crystal-field strength is defined by following Eq. (1) [32, 39]:

where $B_q^k$represents the non-zero crystal-filed parameters that contain the radially dependant parts (k and q) of the one-electron crystal-field. $B_q^k$ characterizes the interaction between ligands and the central ions and includes the radial integral of the wavefunction. The values of the crystal field strength N_V have been calculated for Er³⁺ doped into different matrices Sc₂O₃ (5300 cm⁻¹), Y₂O₃ (2700 cm⁻¹), and Er₂O₃ (2200 cm⁻¹) [40, 41], and they show clearly that the crystal-field strength increases regularly with decreasing ionic radius of the RE host cation, which generates the Stark-splitting energy levels. The presence of peaks at 1542, 1546, and 1551 nm is due to the presence of Sc ions, which increase the crystal field strength and thereby enhance the Stark-splitting of the thermally populated Er energy levels (⁴I_15/2 and ⁴I_13/2 levels) as well as that of the other electronic energy levels.

Figure 7 shows the variation of FWHM at 1537 nm from 4 K to room temperature of Er_xSc_2-xSi₂O₇. FWHM increases from 1.1 to 2.3 nm with increasing of temperature. The PL spectra are plotted as a function of temperature in the inset in Fig. 7. As can be seen, the Er³⁺ PL intensity becomes stronger at low temperature without shifting. The absence of shoulders peak emission for the main peak at 1537 nm like Er doped Si nanocluster that can be inducing larger inhomogeneous broadening of Er³⁺:⁴I_13/2-⁴I_15/2 transition [42].

 

Fig. 7 Variation of FWHM at 1537 nm as a function of temperature. Inset shows temperature dependent PL spectra of Er_xSc_2-xSi₂O₇ for λ_ex = 1527.6 nm.

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The results of integrated PL intensity at 1537 nm as a function of excitation power are shown in Fig. 8. The PL intensity increases linearly and then exhibits a sublinear increase associated with the upconversion phenomena where two Er³⁺ ions both excited at the ⁴I_13/2 level interact, with one being de-excited to the ⁴I_15/2 ground state and the other being resonantly excited to the ⁴I_9/2 level [43]. At low pump powers, however, a nonradiative concentration quenching process may occur. In this case the excitation is resonantly transferred from one excited Er³⁺ ion to a nearby Er³⁺ ion in the ground state and hence it travels along the sample. In order to determine the excitation cross-section, it can be shown that the pump intensity dependence of the PL from erbium is given by the following Eq. (2) [22]:

 

Fig. 8 Integrated PL intensity as a function of excitation power at 1527.6 nm.

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where I_PLmax is the maximum PL intensity, h is Plank’s constant, c is the speed of light, p_ex is the pump power, and τ the lifetime at room temperature, which is equal to 12 μs fitted with single exponential for the excitation wavelength at λ_ex = 1527.6 nm. Fitting the experimental data of this equation, an estimate of the excitation cross-section σ_ex can be obtained. The value of σ_ex is 3.0x10⁻²⁰ cm² at λ_ex = 1527.6 nm. This is one order of magnitude higher than the σ_ex of 1.4 10⁻²¹ cm² at λ_ex = 980 nm reported by us for the same structure [44]. The difference between both values is due to the excitation wavelength (resonant excitation at 1527.6 nm) and the life time depending on excitation wavelength. In addition, σ_ex at λ_ex = 980 nm is on the same order as the excitation cross-section in α-(Y_1-xEr_x)₂Si₂O₇ of 2.0x10⁻²¹ cm². This is means that Sc plays the same role as Y in the structure - an inactive element and acts as a disperser of Er ions – in addition, it presents a higher crystal field strength that enhances the Stark-splitting of Er energy levels. On the other hand, σ_ex in Er_xSc_2-xSi₂O₇ (1.4 10⁻²¹ cm²) is smaller than that in α-(Yb_1-xEr_x)₂Si₂O₇, 2.0x10⁻²⁰ cm² at 980 nm, because of the energy transfer from Yb to Er ions [34].

The key advantage of Er_xSc_2-xSi₂O₇ is that present high excitation cross-section, small FWHM at 1537 nm, and high crystal field strength due to Sc ions that enhances the Stark-splitting of Er. Er_xSc_2-xSi₂O₇ compound is promising material by combining with photonic crystal structures to build light sources or on-chip optical amplifiers.

4. Conclusion

In summary, polycrystalline Er_xSc_2-xSiO₅ and Er_xSc_2-xSi₂O₇ compounds were fabricated using RF-sputtering by alternating Er₂O₃, Sc₂O₃ layers separated by a SiO₂ layer and annealing in O₂ gas. We demonstrated that stoichiometric Er-Sc silicate films and crystalline structures can be controlled by changing the annealing temperature and we identified the stoichiometry. The Er-Sc silicate layer shows a sharp emission peak at room temperature centered at 1528 and 1537 nm corresponding respectively to Er_xSc_2-xSiO₅ and Er_xSc_2-xSi₂O₇ phases. The PL efficiency in Er_xSc_2-xSi₂O₇ phase is due to the decrease of nonradiative transition rates.

The narrow FWHM of Er_xSc_2-xSi₂O₇ is due to the small ionic radii of Sc³⁺ which enhance the crystal field strength. The excitation cross-section is 3.0x10⁻²⁰ cm² at λ_ex = 1527.6 nm. The narrow emission is very promising for photonic crystal light emitting devices because the extraction efficiency can be increased with a pronounced narrowing of the emission.