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Abstract
We have investigated the structural evolution of Si-rich Er silicate films. Er silicate layers embedded with amorphous silicon (a-Si) clusters are formed upon the annealing above 1000 °C of Er-doped Si-rich SiO2 films with high Er concentrations. It is found that the crystallization of Er silicates annealed at higher temperatures leads to the disappearance of Si-NCs formed at 900 °C. On the other hand, the interface between Er silicate and embedded a-Si clusters increases the nucleation barrier for Si-NCs in these clusters. A two-step annealing process is utilized to obtain a fine structure and a good crystal quality of Er silicates at the same time, which will benefit the sensitized photoluminescence of Er silicates.
© 2019 Optical Society of America under the terms of the OSA Open Access Publishing Agreement
1. Introduction
Developing an efficient Si-based light source has become the most urgent requirement in Si photonics due to the poor light emission efficiency of bulk Si. [1,2] Among various approaches, Er-doped Si-based materials have attracted much attention because Er ions can emit at 1.54 µm due to electron transition from 4I13/2 to 4I15/2 in the 4f electronic shell, which corresponds to a minimum in the loss spectrum of silica optical fibers. [3–10] However, the Er content in most solid hosts is limited due to its low solid solubility, which results in a very small optical gain. [11,12] On the other hand, Er compounds, such as Er2O3, Er2SiO5, Er2Si2O7 and Er3(SiO4)2Cl, can have Er concentrations as high as 1022cm−3. [13–15] It has been reported that a net gain exceeding 100 dB cm−1 is achieved in erbium chloride silicate nanowires. [15] In addition to Er contents, the sharp absorption lines and the small absorption cross section of Er ions also limits the optical gain. In general, this problem could be solved by introducing sensitizers into host materials. The sensitizers, such as Si nanocrystals (Si-NCs), luminescent centers and amorphous Si (a-Si) nanoclusters, can be efficiently excited by a wide range of wavelengths and transfer energy to nearby Er ions, [16–20] thus increasing the effective excitation cross section of Er ions by three orders of magnitude. [21] Therefore, the incorporation of Er compounds and sensitizers and the efficient energy transfer between them is the key to reaching high optical gain in Er-doped Si-based materials. Miritello et al [13] have reported sensitized photoluminescence (PL) from Er ions in a Er-Si-O/Si multilayer structure. Yin et al [22] have achieved PL and electroluminescence (EL) from amorphous Er silicates coexisting with a-Si excess in Si-rich Er silicate (SRES) films. Shen et al [23] synthesized Er silicates on porous Si framework and obtained Er PL via the coupling between porous Si and Er ions. These researches have achieved sensitized PL from Er ions in amorphous or crystallized Er silicates. However, the crystallization of Er silicates is very important because not all Er ions are optically active and the defect concentration is high in amorphous Er silicates. [13,24,25] In addition, the distributions of sensitizers and Er silicate crystallites require further improvement to obtain high coupling efficiency between Er silicates and sensitizers.
In this paper, we report on an Er silicate film embedded with small a-Si clusters. The relationship among the evolution of Si-NCs, a-Si clusters and Er silicates during annealing is investigated. A theoretical model that takes into account the interface energies, the sizes of a-Si clusters and the nucleation barrier for Si-NCs is proposed to explain the absence of Si-NCs in the Si-rich Er silicate film. A two-step annealing process is utilized in order to modify the film structure. This is the first paper in a series of two. The second one will focus on the sensitized photoluminescence of Er silicates.
2. Experiment
Er-Si-O films with thickness of ∼550 nm were deposited on Si (100) substrates by reactive magnetron co-sputtering using Si and Er targets in a reactive atmosphere of Ar and O2 mixed gas. The radio frequency (rf) powers for Si and Er targets are set to 130 W and 30 W, respectively. The depositions were made on a rotating substrate heated at 300 °C. After deposition, the films were thermally treated at different temperatures in flowing nitrogen in a tubular furnace.
Chemical composition of the Er-Si-O film was measured by Rutherford backscattering spectrometry (RBS) using a 2.02 MeV 4He ion beam at a scattering angle of 160°. Crystalline structures were characterized by X-ray diffraction (XRD) analyses using Cu Kα radiation with the grazing angle of 0.5° or 1°. The structural characteristics were studied by transmission electron microscopy (TEM) including high resolution TEM (HRTEM) and the distributions of elements were studied using TEM equipped with X-ray energy dispersive spectroscopy (XEDS).
3. Results and discussion
The average chemical composition of the film measured by RBS is estimated to be Si, 38 at%; O, 48 at%; and Er, 14 at%. The Si excess of the film compared with Er2Si2O7 is calculated to be 111% (NSi-ex = (NEr-Si-O − NEr2Si2O7)/NEr2Si2O7, where NEr-Si-O and NEr2Si2O7 stand for the atomic percentage of Si atoms in the Er-Si-O film and that in Er2Si2O7, respectively). The as-deposited films were treated with different annealing processes. Figure 1 presents the XRD spectra of the Er-Si-O films annealed at different temperatures. The spectra of the films annealed at 900 °C or 950 °C show no diffraction peaks, demonstrating the amorphous nature of the films. Crystalline structures start to appear at temperatures above 1000 °C, indicating that the crystallization temperature of Er silicate lies between 950 °C and 1000 °C. The XRD peaks do not correspond to any powder diffraction files of Er2SiO5 and Er2Si2O7 in the Joint Committee for Powder Diffraction Standards (JCPDS) reference database, but it is reasonable to use y-Y2Si2O7 and α-Tm2Si2O7 as references for y-Er2Si2O7 and α-Er2Si2O7 due to the similar ionic radius of Er3+, Y3+, and Tm3+. [26] It can be seen that the films annealed at 1000 °C for 30 min, 1000°C for 30 min followed by 1100°C for 30 min, and 1000°C for 30min followed by 1200°C for 1 min are mainly composed of y-Er2Si2O7 while the film annealed at 1000°C for 30 min followed by 1200 °C for 30 min is composed of α-Er2Si2O7. This change of Er silicate polymorph is consistent with the previous demonstration that higher temperature annealing transforms y-Er2Si2O7 into α-Er2Si2O7. [27]
Figure 2(a) shows the cross-sectional TEM image of the Er-Si-O film annealed at 1000 °C for 30 min. It is observed that an embedded structure is formed. The dark area corresponds to Er silicates. The HRTEM image and the corresponding FFT (fast Fourier Transformation) pattern of the indicated dark area is presented in Fig. 2(b). The lattice spacing is 3.05 Å, which corresponds to the (021) plane of y-Er2Si2O7. The bright area is mostly composed of Si atoms according to the large Si excess in the Er-Si-O film. Figures 2(c) and (d) show the HAADF (high-angle annular dark field) image of the film and the corresponding distribution Er, Si, O atoms, which confirms the above statement. It is worth noting that no crystalline structure is observed in the bright area in Fig. 2(a). In other words, a self-organized a-Si-embedded Er silicate film is formed. However, Si-NCs should be able to form under normal circumstances according to the available Si amount and the annealing temperature of 1000 °C. To investigate the origin of the absence of Si-NCs, TEM images of the Er-Si-O films annealed at different temperatures are analyzed. Figure 3(a) shows the HR-TEM image of the Er-Si-O film annealed at 900 °C for 30 min. Er silicates have not crystallized at this temperature while many Si-NCs have formed. The Si-NCs are marked with white circles. An enlarged image of the area indicated by the dashed red circle is shown in the inset of Fig. 3(a). The lattice spacing is calculated to be 3.14 Å, which corresponds to the (111) plane of Si. Figure 3(b) shows the HR-TEM image of the film which is further annealed at 950 °C following the annealing at 900 °C. There is still no crystallized Er silicates in this sample. From a comparison between Figs. 3(a) and (b), it is observed that the distribution of Er ions in the 900 °C-annealed film is relatively uniform, but the annealing at 950 °C has led to the formation of a network of Er-rich areas. At the same time, the Si-NCs formed at 900 °C have completely disappeared. The network structure of this sample is similar to that of the film annealed at 1000 °C except that the film is mainly composed of Er-rich areas and amorphous silicon clusters. Figures 3(c) and (d) present the cross-sectional TEM image and the HR-TEM image of the indicated dark area of the film annealed at 900 °C followed by 1000 °C. When the annealing temperature of the second step increases to 1000 °C, Er silicates have crystallized and still no Si-NCs are observed. The FFT pattern of the HR-TEM image is shown in the inset of Fig. 3(d). The lattice spacing is calculated to be 5.23 Å corresponding to the (110) plane of y-Er2Si2O7. The structure size, the structure distribution, and the Er silicate polymorph of the multi-step annealed samples are all the same with the sample annealed at 1000 °C directly, which demonstrates that the pre-annealing at 900 °C cannot influence the final structure of the films.
Fig. 2. TEM images of the Er-Si-O film annealed at 1000 °C for 30 min. (a) Cross-sectional TEM image. (b) HR-TEM image of the indicated area with the corresponding FFT pattern as inset. (c) HAADF image and (d) the corresponding XEDS map of Er, Si, O atoms.
Fig. 3. TEM images of the Er-Si-O films annealed at different temperatures. (a) HR-TEM image of the film annealed at 900°C for 30 min. Inset is an enlarged image of the indicated area. (b) HR-TEM image of the film annealed at 900°C for 30 min followed by 950 °C for 30 min. (c) Cross-sectional TEM image and (d) HR-TEM image of the indicated area of the film annealed at 900°C for 30 min followed by 1000 °C for 30 min. (e) Cross-sectional TEM image and (f) HR-TEM image of the indicated area of the film annealed at 1000°C for 30 min followed by 1200 °C for 30 min. Insets of (d) and (f) are the corresponding FFT patterns.
Based on the above results, we can conclude that the evolution of the film structure during annealing is as follows. When the annealing temperature is around 900 °C, a large number of Si-NCs are formed. As the annealing temperature rises, Er silicate tends to crystallize. However, the large difference in cell parameters between Er2Si2O7 and Si results in an excess term related to strain in the expression of the nucleation barrier for heterogeneous nucleation of Er2Si2O7 on Si-NCs, which significantly increases the nucleation barrier. Besides, due to the high Er concentrations of the film, the amount of Er silicates in the final structure is very large, as shown in Fig. 3(c), so massive nucleation of Er silicates need to occur in the film and the existence of the Si-NCs greatly hinders this process. On the other hand, the sizes of Si-NCs are very small in the film, so the energy barrier for the dissolution of the Si-NCs is not very high. Therefore, Si-NCs tend to dissolve to reduce the energy barrier for the formation of the Er silicate network structure. The dissolution of Si-NCs occurs simultaneously with the accumulation of Er atoms and the ejection of excess Si atoms from the Er-rich areas, finally forming the structure shown in Fig. 3(b). When the annealing temperature increases to about 1000 °C, Er silicates start to crystallize in the Er-rich areas, which leads to the formation of the a-Si-embedded Er silicate layer.
After the crystallization of Er silicates, a large number of a-Si clusters exist in the film. In general, amorphous Si clusters should have crystallized at 1000 °C. [28,29] However, neither prolonging the duration time of the annealing at 1000 °C to 1 h nor adding a second annealing step of 1200 °C for 30 min can lead to the formation of Si-NCs. Figure 3(e) shows the cross-sectional TEM image of the Er-Si-O film annealed at 1000 °C for 30 min followed by 1200 °C for 30 min. Only the dark areas, which represent Er silicates, are found to be crystallized. The HR-TEM image of the marked dark area and the corresponding FFT pattern are shown in Fig. 3(f). The indicated lattice spacing is 4.45 Å corresponding to the (1-11) plane of α-Er2Si2O7. As shown in Fig. 2(a) and Figs. 3(c) and (e), the sizes of the embedded a-Si clusters are mostly in the range of 10-25 nm. Under the circumstance of such a small size, the nucleation of Si-NCs in the a-Si clusters will be greatly influenced by the interface of the surrounded crystalline Er silicates. This is similar to the nanocrystal growth in ultrathin a-Si films with an oxide interface on both sides of the thin Si layer. [30–32] It is reported that an amorphous SiO2 interface on both sides of a Si layer with the thickness less than 50 nm will not result in a homogeneous and uninfluenced nucleation of Si within the layer and the crystallization temperature will increase. [31] We assume that the a-Si cluster is spherical in shape and the nucleus is located at the center of the cluster and is also spherical in shape. The model is illustrated in Fig. 4. The material s represents the Er silicate, material a the a-Si, and material c the crystalline nucleus.
Fig. 4. Model of a spherical shaped crystalline nucleus embedded in an a-Si cluster with Er silicate interfaces.
We define ${\gamma _{ac}}$, ${\gamma _{sc}}$, and ${\gamma _{sa}}$ as the interfacial free energies per unit area between the a-Si (a) and the crystalline nucleus (c), between the Er silicate (s) and the crystalline nucleus (c), and between the Er silicate (s) and the a-Si (a), respectively. However, for the structure considered in Fig. 4, the interface between material s and material c is not well defined if the distance l is very small. When l = 0, a sharp interface between materials s and c with the interface energy ${\gamma _{sc}}$ is formed. When l → ∞, the material s and material c are separated by two well defined interfaces (the interface between materials a and c with the interface energy ${\gamma _{ac}}$ and the interface between materials s and a with the interface energy ${\gamma _{sa}}$). Considering the interaction between these two interfaces for small l, an effective interface energy $\gamma _{sc}^{eff}$ is defined, which lies between the two cases for l = 0 and l → ∞. Therefore, we write
where K is an effective order parameter which increases with the decrease of l and is unity for l = 0 (Er silicate/crystalline nucleus interface) and zero for l → ∞ (a-Si/crystalline nucleus interface). Similarly, for a spherical particle of the same size with the crystalline nucleus but in amorphous phase a, the effective interface energy $\gamma _{sa}^{eff}$ is written as Given the above assumptions, the change in Gibbs free energy when the spherical shaped nucleus with radius r in Fig. 5 is formed can be written asFig. 5. (a) Cross-sectional TEM image of the Er-Si-O film annealed at 1200°C for 1 min. Inset is the HR-TEM image of the indicated bright area. (b) HR-TEM image of the indicated dark area in (a). Inset is the corresponding FFT pattern.
The nucleation barrier is given by
In order to make the calculation simpler, we choose to ignore the fact that $\Delta {\gamma _{eff}}$ depends on the radius r of the nucleus. Then the critical radius of the nucleus derived from the Eq. (5) is Inserting the Eq. (6) into the Eq. (3), we obtain the nucleation barrier For l → ∞, namely, $K$=0, $\Delta {\gamma _{eff}}$=${\gamma _{ac}}$. Therefore, the bulk nucleation barrier $\Delta {G_b}^\ast $ is Then the ratio of $\Delta {G^\ast }$ to $\Delta {G_b}^\ast $ is equal to $\Delta \gamma _{eff}^3/\gamma _{ac}^3$.For thin Si layers capped with amorphous SiO2 in a superlattice structure, the relation of the interface energy of oxide/crystalline nucleus interface $\gamma _{oc}^{}$, the interface energy of oxide/a-Si interface $\gamma _{oa}^{}$, and $\gamma _{ac}^{}$ is given by [31]
In our model, the difference between ${\gamma _{sc}}$ and ${\gamma _{sa}}$ should be even larger due to the crystalline nature of Er silicates and the large difference in cell parameters between Er2Si2O7 and Si. Therefore, we obtain andIf the Er-Si-O film was annealed at 1200 °C directly, the film structure is quite different from the structure shown in Fig. 3(e). Figure 5 shows the cross-sectional TEM image and the HR-TEM image of the Er-Si-O film annealed at 1200 °C for 1 min. The average size of the a-Si clusters is much larger than that of the film annealed at 1000 °C followed by 1200 °C. In some large a-Si clusters, Si-NCs are observed. The inset of Fig. 5(a) shows the HR-TEM image of the indicated bright area. The lattice spacing is calculated to be 3.14 Å, which corresponds to the (111) plane of Si. The presence of Si-NCs in large a-Si clusters and the absence of Si-NCs in small a-Si clusters (Fig. 3(e)) confirms the above statement that the influence of interface on the crystallization of a-Si clusters increases greatly with the decrease of cluster size. Figure 5(b) shows the HR-TEM image and the corresponding FFT pattern of the indicated dark area in Fig. 5(a). The lattice spacing is 6.63 Å which corresponds to the (100) plane of α-Er2Si2O7. It is worth noting that some of the dark areas are found to be not crystallized due to the short annealing time.
From a comparison between Fig. 3(e) and Fig. 5(a), it is concluded that the pre-annealing at 1000 °C has great influences on the final structure of Er-Si-O films. The sizes of the a-Si clusters and the Er silicate crystallites of the film annealed at 1000 °C followed by 1200 °C are almost the same with the film annealed at 1000 °C alone and are much smaller than those of the film annealed at 1200 °C directly, demonstrating that the fine structure formed at 1000°C is maintained after the annealing at 1200°C. In order to optimize the sensitized PL of Er silicates, the structure of the a-Si-embedded Er silicate film should be as fine as possible, which demands that the film should be annealed at lower temperatures. However, the crystal quality of Er silicates is also very important, and low temperature annealing cannot effectively reduce the defect concentrations. Therefore, using a two-step annealing process is very important because the fine structure formed at the first step can be maintained and the second step at higher temperatures can improve the crystal quality of Er silicates. Although the annealing temperature of the first step should be very low, the crystallization of Er silicates during the first step is essential because further diffusion of Er ions during the second step cannot be restrained without the crystallization of Er silicates. As shown in Fig. 3(c), the film annealed at 900°C followed by 1000°C has a structure that is the same with the film directly annealed at 1000°C because 900 °C is not high enough for the crystallization of Er silicates. Therefore, the annealing temperature of the first step should be as low as possible because the Er silicates can crystallize to achieve a fine film structure which can be maintained after the second step higher temperature annealing.
4. Conclusion
In conclusion, we have fabricated a-Si embedded Er silicate layers in Er-doped Si-rich SiO2 films with Er concentrations as high as 14 at%. Si-NCs form at low annealing temperatures and disappear upon the crystallization of Er silicates. The interface between Er silicates and small size a-Si clusters restrains the crystallization of Si-NCs in these clusters by increasing the nucleation barrier. The utilization of a two-step annealing process solves the contradiction between obtaining a fine structure and a high crystal quality, which has laid a foundation for efficient sensitized photoluminescence of Er silicates.
Funding
National Key R&D Program of China (2018YFB2200102); National Natural Science Foundation of China (61874095).
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