Luminescence Mechanism in Amorphous Silicon Oxynitride Films: Band Tail Model or N-Si-O Bond Defects Model

Silicon oxynitride films are one kind of important gate dielectric materials for applications in the fabrication of silicon CMOS integrated circuits (ICs), which have been widely and deeply studied. However, with the significant demand of the technologies for Si-based monolithic optoelectronic ICs, the research efforts on the optoelectronic applications of these materials have been continually increasing, particularly in the study of light emission properties and recombination mechanisms. In this paper, we first briefly outline the present photoluminescence (PL) mechanisms in amorphous silicon oxynitride (a-SiO_xN_y) films. Since, the PL properties and recombination processes are affected by both structural disorder and chemical disorder, the PL mechanism has been still unclear and even controversial until now. Among these various PL recombination models, the band tail states and defect state models have gained general consensus. Recently, a N-Si-O bond defect model has been reported, which depends on relative atom concentration of oxygen and nitrogen in the silicon oxynitride materials. It has been revealed that oxygen bonding plays a key role, not only in reducing the structural disorder, but also in creating N-Si-O (N_x) defect states in the band gap. The characteristics of two models, namely band tail and N-Si-O bond defects, have been discussed in detail. Finally, it has been shown that by controlling the chemical composition of these non-stoichiometric silicon oxynitride materials, the optical and electronic properties can be improved.

Introduction

A family of silicon based silicon oxide and nitride is traditional and important electronic and optical films, which have been widely applied in the fields of microelectronic and optoelectronic devices. Among these materials, the structures and properties of amorphous silicon oxynitride (a-SiO_xN_y) films including oxygen-rich silicon oxynitride (O-type a-SiO_xN_y) and nitride-rich silicon oxynitride (N-type a-SiO_xN_y), are very different from those of stoichiometric SiO₂ and Si₃N₄. During the first decade, after the first publications by Fukuda et al. [1] and Hwang et al. [2] in 1990, silicon oxynitride films have been extensively studied as a replacement for the conventional SiO₂ gate dielectric. In contrast to SiO₂, this kind of materials has higher dielectric constants and excellent diffusion-limiting properties, which can therefore suppress the penetration of boron from the poly-Si gate, as well as enhance the reliability [3–5]. On the other hand, incorporation of nitrogen in Si/SiO₂ interfaces can improve both electronic and optical properties [6], which will be discussed in detail in section Role of oxygen bonding and N-Si-O bonding defects.

A large number of techniques have been proposed and evaluated to engineer the concentration and distribution of nitrogen and oxygen in the silicon oxynitride films. Among them, the physical or chemical vapor deposition with various nitrogen based reaction sources, such as N₂O [7–10], NO [11, 12], and NH₃ [13–15] etc., can be used to directly grow a-SiO_xN_y films. Besides, the nitridation treatments [thermal [16] or plasma [17–19]] can also be used to control the nitrogen concentration in the grown SiO₂ films. In the meantime, many works on the characterization of structures, physical and electrical properties of silicon oxynitride films for the microelectronic applications have also been reported [12, 20–22].

On the other hand, the silicon oxynitride films have been employed, not only as an alternative gate dielectric layers for application in microelectronic devices, but also as optical materials in optoelectronic devices, such as ultraviolet transparent waveguide [23–25], microarrays of Si-based light emitters for biosensor [26], CO₂ capturer [27], light emitting devices [28–30], and solar cells [31]. With the prime requirements of the technologies for integrating optoelectronic devices with silicon-based microelectronics, the research efforts on optical properties of silicon oxynitride films have been continually increasing. Among these, most research focused on the study of light emission properties and the recombination mechanism. However, even though the photoluminescence (PL) efficiency of silicon oxynitride is higher than that of other Si-based luminescent materials, it is still lower for practical application requirements. Hence, understanding the PL mechanisms in silicon oxynitride remains the key foundation for fabricating Si-based high-efficiency light emitting devices.

In this paper, we first briefly outline the present PL mechanisms of a-SiO_xN_y films in section PL properties and recombination mechanisms in a-SiOxNy films. Some groups reported that the PL in O-type a-SiO_xN_y was due to radiative recombination between localized band-tail states associated with Si-N bonds which is similar to a-SiN_x [32, 33], while other groups clarified that the complete picture of PL in silicon oxynitride can be explained by the radiative recombination via both band tail states and luminescent defect states [34]. Besides, it was also suggested that the PL might arise from radiative recombination center consisting of the silicon sub-oxide bonding defects [35–37]. Among these various PL models, even though the band tail state and defect state models have gained general consensus, PL mechanisms have still been unclear and even controversial. In section Role of oxygen bonding and N-Si-O bonding defects, we analyzed the role of oxygen bonding being played in silicon oxynitride network, which not only reduced the structural disorder but also created N-Si-O (N_x) defect states in the band gap. Based on this, a N-Si-O bond defect model has been proposed in section Characteristics of band tail states and defect states dominated PL model in a-SiO_xNy, which depends on relative atom concentration of oxygen and nitrogen in the silicon oxynitride materials. In section Improved optical properties in N-type a-SiN_xO_y films, the characteristics of two PL models, i.e., band tail and N-Si-O bond defect, have been discussed in detail. Finally, we showed that by controlling the chemical composition of this non-stoichiometric silicon oxynitride materials the optical and electronic properties can be improved. An n-a-SiN_xO_y/p-Si HJ LED fabricated by the standard CMOS technology has been demonstrated. Finally, the gate controlled Si MOSFET has been emphasized and discussed.

PL Properties and Recombination Mechanisms in a-SiO_xN_y Films

In this section, we briefly outline the present PL mechanisms in a-SiO_xN_y films. It is well-known that, in the non-stoichiometric a-SiO_xN_y materials, the PL properties are affected not only by preparation methods but also by the relative chemical composition of Si, N, and O elements. Over the last two decades, many works on study of PL mechanism were reported, but it has been still unclear and even remained ambivalent up to date. This is probably due to the difficulties of controlling and separating individual contribution made by both structural disorder and chemical disorder which are contained in this kind of material.

In the 1990s, a group at university of North Carolina, Augustine et al. [36], started the study of the visible light emission from thin films containing Si, O, N, and H. The a-SiO_xN_y:H samples were prepared by PECVD using SiH₄ and N₂O at elevated temperatures. The results of X-ray photoelectron spectra (XPS) and Fourier-transform infrared (FT-IR) spectroscopy showed that the chemical composition was dominated by silicon sub-oxide bonding with N atoms serving as a significant impurity. The broad visible PL emission was observed at room temperature. The radiative lifetime was <10 ns and the PL intensity was nearly unchanged from 80 to 300 K. The combined results suggested that N was not the radiative center in the SiO_xN_y matrix, but the silicon sub-oxide bonding acted as a defect state which enabled visible PL emission. Thereafter, they reported [35] the effect of the rapid thermal annealing on PL properties of a-SiO_xN_y:H films. It was suggested that the annealing was characterized by two processes. The first was that the Si-H and N-H bonds were broken and hydrogen effused from hydrogen clustering region, which resulted in an increase of dangling bond density. The second was the local reconstruction of Si-O and Si-N bonds. The results obtained after annealing were in agreement with the above PL characteristics, based on which PL recombination model was schematically shown in Figure 1A. Due to the local chemical bonding disorder, the density of these localized states was not smoothly decaying from the mobility edges of conduction and valence bands, as shown in Figure 1B.

FIGURE 1

Figure 1. (A) Schematic representation of the model of geminate PL recombination through localized defect states within the band tails. Excitation is denoted by EX, radiative recombination by R, and non-radiative recombination at midgap states by NR. (B) Energy vs. density of states in the band gap. Structure is due to localized chemical bonding in the films. The mobility edges are denoted by horizontal dotted lines [Reprinted from Augustine et al. [35], with permission].

Four years later, based on above results, Price, McNeil and Irene et al. [34, 37] further characterized the luminescence center in a-SiO_xN_y films. The experimental results excluded the origin of PL emission from Si dangling bonds, NBOHC [38, 39] and N dangling bonds, as well as silicon nanocrystals. The recombination mechanism of a complete PL spectrum in near-stoichiometric a-SiO_xN_y:H films must be associated with both band tail states and defect states.

At the beginning of the 2000s, Noma et al. [40], reported that the four PL peaks appearing in the range from 2 to 5 eV had been observed in oxygen-rich a-SiO_xN_y:H films, deposited by PECVD using SiH₄, N₂O, and N₂ as reaction gases. The study combining with FTIR, XPS, electron spin resonance (ESR) and scanning electron micrographs (SEM) indicated that the two peaks of 2.7 and 4.4 eV were originated from silicon homo-bonds, and the other broad peaks located around 2.6–2.9 eV were due to the Si-N bonds in the gathering regions of Si-N bonds. Furthermore, since the samples were prepared by PECVD using SiH₄ gas source, a-SiO_xN_y:H samples contained hydrogen. In order to eliminate the effect of hydrogen, the samples were prepared by nitridation of silicon dioxide and, thus, were free of hydrogen [33]. The results of the emission spectrum, excitation spectrum and time-resolved PL (TRPL) spectrum, combined with the theoretical analysis, again confirmed that the origin of the PL with its peak in the range of 2.6–2.9 eV was from the Si-N bonds in a-SiO_xN_y film, and thus irrelevant to hydrogen content.

Soon afterwards, Kato et al. [32, 41, 42] further performed the research on the PL mechanism in a-SiO_xN_y films and a-SiN_x films. The conclusion was drawn that the PL from a-SiO_xN_y films originated from radiative recombination between localized band-tail states associated with Si-N bonds. The experimental evidences were shown in Figures 2, 3. As shown in Figure 2A, the bandwidth of PL (ΔE_PL) caused by the band-tail recombination is proportional to the width of the localized states (E_U), in which the straight line passes the origin of the coordinate axes, indicating that the band-tail states are mostly induced by structural disorder. Figure 2B shows the PL peak energy (E_PL) increases monotonically with an increase in Urbach edge energy E_Uedge in both a-SiO_xN_y and a-SiN_x films, which means that they exhibit the same feature of band-tail recombination mechanism. The time-resolved PL (TRPL) measurements were used to analyze the recombination processes of photo-excited carriers. The results of nanosecond temporal evolution of transient PL showed that the E_PL shifted as a function of specific delay time after excitation. As shown in Figures 3A,B, for both a-SiO_xN_y and a-SiN_x, E_PL first shifts to red, then to blue, and back to red again. It is also a feature of the band-tail recombination mechanism, in which three different effects are included, such as thermalization, coulomb interaction, and extent of localization of carriers, acting together to cause the shift of PL peaks.

FIGURE 2

Figure 2. (A) Relation between the Urbach energy E_U and the FWHM of the PL spectrum ΔE_PL. Closed and open circles are for a-SiO_xN_y:H and a-SiN_x:H, respectively. (B) Correlation between the Urbach edge energy E_edge and the PL peak energy. Closed and open circles are for a-SiO_xN_y:H and a-SiN_x:H, respectively [Reprinted from Kato et al. [32], with permission].

FIGURE 3

Figure 3. Typical time-resolved PL spectra excited by 5.0-eV photons, observed in a-SiO_xN_y:H (A) and a-SiN_x:H (B). Numerals following (i)~(vi) indicate given observation periods [Reprinted from Kato et al. [32], with permission].

In the late 2000s, Tewary et al. [43] reported luminescence from controlling defect and Si nanoparticle in a stoichiometry SiO_1.08N_0.32 films by using CO₂ laser annealing. It was found that there were two distinct peaks located at 570 and 800 nm in the irradiated areas. From the combined study of TEM and PL lifetime, it can be concluded that the visible peak with a short PL lifetime of ~10 ns was related to silicon sub-oxide defects, which was the same as that reported by Augustine et al. [36], while the infrared peak was due to exciton recombination in Si nanoparticles.

Since 2006, Huang et al. [44] have observed strong green light emission from oxidized SiN_x films. The samples were fabricated by PECVD at room temperature followed by thermal or plasma oxidation at 100°C, which is different from a-SiO_xN_y films prepared by PECVD using N₂O and SiH₄ as reaction gas sources. Subsequently, they observed green-yellow electroluminescence from this kind of oxidized a-SiN_x films [29, 45]. At that time, it was suggested that the light emission was due to the radiative recombination through the luminescent center related to Si-O bonds in the Si-N matrix. Based on above investigations, Dong et al. [46] further characterized the luminescent defect states. The results from measurements of FTIR and XPS verified the existence of O-Si-N bonding configuration, which acted as a new luminescent defect state in our a-SiN_xO_y films.

In recent years, Ruggeri et al. [47] reported that the luminescent signal in a-SiO_xN_y layer was affected by silicon chemical environment induced by IR annealing inside the amorphous SiON matrix. Jou et al. [48] reported PL characteristics of a-SiO_xN_y films, which were prepared by reactive sputtering of silicon. The PL spectra with blue and green peaks were observed. The blue and green emissions possibly arose from N related defects and O-Si-N bonding related defects, respectively.

More recently, Zhang et al. [49] and Lin et al. [30] conducted the intensive investigation on PL mechanism in a-SiN_xO_y films. The results of XPS and EPR, as well as PLE measurements, indicated that the oxygen atoms in SiN_x matrix not only reduced the structural disorder but also created N-Si-O (N_x) defect states in the band gap. And then the PL mechanism dominated by the N-Si-O defect states was proposed, which will be discussed in detail in sections Role of oxygen bonding and N-Si-O bonding defects and Characteristics of band tail states and defect states dominated PL model in a-SiO_xNy.

Role of Oxygen Bonding and N-Si-O Bonding Defects

Before discussion about the role of oxygen bonding in a-SiO_xN_y, we would like to very briefly summarize the role of Si/SiO₂ interface region on photoelectronic properties of luminescent silicon. As mentioned above, the structures and properties of Si/SiO₂ interface is a critical factor for influencing the performance of the electronic and optical devices. For the luminescent silicon, like porous silicon, silicon nanocrystals (Si-nc) and silicon quantum dots, the sharp and stable Si/SiO₂ interface can passivate non-radiative defect states on the surface, and consequently provide stable light emission [6]. On the other hand, Wolkin et al. [50] observed that oxidation can also form Si=O double bonds, which act as a luminescent defect states in the Si-nc band gap and pin their emission energy. Moreover, Pavesi et al. [51] reported that radiative luminescent interface states play a key role in the optical gain of Si-nc.

For a-SiO_xN_y films, apart from the preparation methods and conditions, the relative atomic concentration of oxygen and nitrogen O/(O+N) is a crucial factor influencing their chemical bonding configuration and physical properties. For example, when the ratio O/(O+N) > 0.4, the nitride-like structures will be transformed into oxide-like structures and, consequently, the electronic properties will be changed in the silicon oxynitride gate dielectric layer [5]. Thereafter, for simplicity, in this paper the “O-type a-SiO_xN_y” film is used to denote oxygen-rich amorphous oxynitride film, and the “N-type a-SiN_xO_y” film refers to nitrogen-rich amorphous oxynitride film [52].

In the study on optical properties of a-SiO_xN_y films, it has also been found that the PL properties are dependent on the O/(O+N) ratio. As mentioned above, Kato et al. [32], Price et al. [34] and Augustine et al. [36] reported PL from the a-SiO_xN_y films prepared by PECVD with N₂O:SiH₄ reaction gases. From XPS measurements, the O/(O+N) ratio was determined in the range of 0.60–0.90, suggesting these a-SiO_xN_y films can be called O-type a-SiO_xN_y films. They believed that the PL arose either from radiative recombination between localized band-tail states associated with Si-N bonds similar to a-SiN_x films or from both band-tail states and luminescent defect states, which consist of silicon sub-oxide bonding defects. However, in our previous works [44, 46, 49], the a-SiN_xO_y films were prepared by PECVD with NH₃:SiH₄:N₂ reaction gases, and then subsequently oxidized in situ by oxygen plasma [53]. As the O/(O+N) ratio is able to be controlled in the range of 0.05–0.15, and these a-SiN_xO_y films can be called as N-type a-SiN_xO_y films. The origin of the PL was found to be attributed to the radiative recombination through the luminescent N-Si-O bonding defect states, which is different from band-tail states recombination model and will be discussed in section Characteristics of band tail states and defect states dominated PL model in a-SiO_xNy.

In order to explain the role of oxygen bonding in the atomic structure and PL properties of a-SiN_xO_y films, we prepared a series of N-type a-SiN_xO_y samples with different reaction gas ratios R = [NH₃]/[SiH₄] which were listed in Table 1 [49]. It shows that for our samples, the O/(O+N) ratio was in the range of 0.07–0.13. For comparison, three kind of controlled a-SiN_x films were also prepared by PECVD but without oxidation plasma treatment. The optical parameters of both series of a-SiN_xO_y and a-SiN_x samples, such as optical band gap E_opt, Urbach tail edge E_Uedge, Urbach tail width E_U, and PL peak energy E_PL etc., were listed in Table 2. Figure 4A shows the band tail E_Uedge and E_U as a function of R. When the R increases, the E_Uedge and E_U increase monotonically in both a-SiN_xO_y and a-SiN_x samples. This phenomenon resulted from the increase in short range structural disorder with increasing of atomic ratio N/Si in amorphous Si based nitride. Moreover, it is interesting to find that the degree of short range structural disorder in a-SiN_xO_y is smaller than that in a-SiN_x. For example, at R = 2, the band tail width E_U in a-SiN_xO_y film is equal to 0.66 eV, while E_U amounts to 1.5 eV in a-SiN_x film. It can be considered that since O atoms are more electronegative than N atoms, the incorporated O atoms will easily substitute N atoms to form Si-O bonds in the silicon nitride network. On the other hand, according to the chemical bonding principle, oxygen takes two-fold coordination to form Si-O bonds, whereas nitrogen takes three-fold coordination to form Si-N bonds, therefore, the Si-O bonds can relieve the internal stress, which results in smaller network disorder and narrower band tail width (E_U) in our N-type a-SiN_xO_y samples. The data in Figure 4B is reprinted from Kato et al. [32]. They also found that the degree of the structural disorder in a-SiO_xN_y:H was smaller than that in a-SiN_x:H and gave similar explanations.

TABLE 1

Table 1. Binding energy of N 1s and Si 2p in N-Si-O bond configurations, and Si, N, O chemical compositions of a-SiN_xO_y films from XPS spectra.

TABLE 2

Table 2. Summary of the optical parameters and PL peak energies of two series of samples at 9 K.

FIGURE 4

Figure 4. (A) Urbach edge energy E_Uedge (circles) and the Urbach energy E_U (triangles) vs. flow ratio NH₃/SiH₄. Closed and open symbols are for a-SiN_xO_y and a-SiN_x, respectively [After Zhang et al. [49], with permission]. (B) Urbach edge energy E_edge (circles) and the Urbach energy E_U (triangles) vs. abundance ratio. Closed and open symbols are for a-SiO_xN_y:H and a-SiN_x:H, respectively [Reprinted from Kato et al. [32], with permission].

The incorporation (or say doping) of oxygen atoms into amorphous silicon nitride networks not only reduced the band tail structural disorder, but also created N-Si-O bonding defect states in the band gap. We employed XPS and EPR measurements combining with silicon tetrahedral bonding model to verify the existence of N-Si-O bond configuration in our N-type a-SiN_xO_y films. By XPS measurements, the relative atomic concentration of Si, N, and O can be determined by the area integration of each binding energy peak of Si 2p, N 1s, and O 1s, respectively, which were listed in Table 1 [49]. We can see that the relative atom ratio of O and N, i.e., O/(O+N), is <0.1, which means that O atoms serve as impurities in the silicon nitride network, accordingly it can be called as N-type a-SiN_xO_y films.

For the study of bonding configurations in non-stoichiometric a-SiO_xN_y films, several literatures have been published [54–58]. Cova et al. [55] reported that, in the a-SiN_xO_y network, hetero-bonds (Si-N and Si-O bonds) are more preferably formed than homo-bonds (Si-Si, N-N, and O-O bonds), by reason of different bonding energies. That is to say, the bonding energies of homo-bonds [E_b(Si-Si) = 1.8 eV, E_b(N-N) = 1.6 eV, and E_b (O-O) = 1.4 eV] are weaker than those of hetero-bonds [E_b(Si-N) = 3.5 eV and E_b (Si-O) = 3.8 eV]. From XPS measurements, based on analysis of the Si 2p spectra, the results of the binding energy shift, which were induced by the formation of Si-O and Si-N bonds, supported that the N_x-Si-O_y complex was present in the silicon oxynitride network [56]. Besides, if considering the differences of chemical activity, the electronegativity and the competitivity between O and N atoms, the formation of the N_x-Si-O_y bonding configuration is much easier than that of the individual Si-N or Si-O bond [57]. Based on the above analysis, the silicon centered tetrahedral bonding structures with N_x-Si-O_y atomic configuration (x = 1, 2, 3 and y = 1, 2, 3; x + y = 4) have been proposed [55–57].

In our previous work [48, 59], from the results of XPS measurements, the existence of the N_x-Si-O_y bonding configuration was also found in our N-type a-SiN_xO_y films. Figure 5 shows the Si 2p peak position of our N-type a-SiN_xO_y sample with R = 8, 4, and 0.5. It can be found that the dominant Si 2p peak position shifts from 102.5 to 102.0 eV with R increasing from 0.5 to 8, which are just intermediate in the range between 103.35 eV (SiO₂ binding energy) and 101.90 eV (Si₃N₄ binding energy). It can be inferred that multiple bonding configurations, consisting of Si-centered tetrahedral N_x-Si-O_y structures, were present in these samples. In order to quest for the possible multiphase configurations, three phases named N₃-Si-O (102.3 eV), N₂-Si-O₂ (102.65 eV), and N-Si-O₃ (102.95 eV), as well as stoichiometric phases SiO₄ (103.35 eV) and SiN₄ (101.90 eV) [53], were used to fit the dominant Si 2p peak by utilizing Voigt functional form. As shown in Figure 5, the results of good fitting indicated that the above Si, O, N multiphase configurations existed in our N-type a-SiN_xO_y films.

FIGURE 5

Figure 5. XPS spectra of Si 2p peaks for a-SiN_xO_y films with different flow rate ratios, (A) R = 8, (B) R = 4, (C) R = 0.5. The deconvolution is performed on original data denoted by the dotted line [After Dong et al. [59], with permission].

As mentioned above, any type of structure of silicon oxynitride is made up of a center Si atoms, which are tetragonally surrounded by four randomly distributed Si, N and O atoms. The following question is how to verify that the N_x-Si-O_y bonding configurations are a kind of point defect in the a-SiN_xO_y films?

As we know, in a-SiN_x films, the structure of paramagnetic Si dangling bond (K⁰) defect consists of a center Si atom bonded to three N atoms and an unpaired electron (-Si=N). Habraken and Kuiper [5] derived from the EPR measurements that the g value of K⁰ is 2.0022 and its width ΔH_pp amounts to 12.5 G. For the K⁰ defect family with the bonded N atoms replaced by O or Si atoms, the g value of K defect will vary from 2.0022 to 2.0055 (in which 2.0055 is the g value of pure Si dangling bond). Zhang et al. [49] reported that the g values of N-type a-SiN_xO_y samples with R varying from 0.5 to 8 were in the range of 2.0038–2.0026, which were located in the range of 2.0022–2.0055 (g value of K defect family) [5, 60, 61]. It means that the N_x-Si-O_y bonding configurations also acted like point defects in our N-type a-SiN_xO_y samples. For simplicity, in this paper we call this kind of N_x-Si-O_y bonding complex an N_x defect state. The total spin density of N_x defects was calculated to be ~10¹⁷ cm⁻³ by double integration of EPR signals and by correction according to a standard sample a-SiO_0.8 [62]. We also confirmed that the position of the N_x defect states is located in the band gap of around 0.75 eV below the E_Uedge by the absorption and PL measurements.

Characteristics of Band Tail States and Defect States Dominated PL Model in a-SiO_xN_y

According to the typical band-tail states model proposed by Street et al. [63], Austin et al. [64] and Siebert et al. [65], the excited carriers residing in shallow tail states were thermally ionized to give rise to non-radiative decay, but those residing in deeper tail states generated luminescence via the radiative recombination. In general, the PL peak energy E_PL is independent of the excitation energy E_exc when E_exc≥E_opt, while for E_exc<E_opt, the E_PL will red-shift with decreasing E_exc. In our experiments [49], the phenomenon of red-shift of PL peak with decreasing E_exc can be found only in a-SiN_x samples, which was shown in Figure 6B. It can be clearly seen that, when E_exc≥E_opt (3.24 eV), the E_PL kept almost constant at 2.32 eV, while for E_exc<E_opt, the E_PL showed a gradual red-shift down to 2.0 eV. Actually, the trace of this curve is the same as the distribution profile of band tail states. However, for the a-SiN_xO_y sample (R = 1.5, E_opt = 3.5 eV) E_PL is independent of E_exc, regardless of whether E_exc≥E_opt or E_exc<E_opt, which was shown in Figure 6A. As for the a-SiN_xO_y sample with E_opt = 3.5 eV, when E_exc varied from 4.6 to 2.6 eV, E_PL only changed by ~0.05 eV, which was obviously different from that of the band tail defect recombination model. We attributed this PL characteristic to the contribution from N_x defect. On the other hand, the dependence of PL integrated intensity I_PL on the relative value of E_exc and E_opt in a-SiN_x was also different from that in a-SiN_xO_y. As shown in Figure 7, under the condition of E_exc = 4.5 eV and E_exc≥E_opt, the I_PL of the a-SiN_x samples increases with increasing E_opt, while for a-SiN_xO_y samples, the I_PL does not increase monotonically with E_opt, and a peak emerged at E_opt = 3.8 eV, which was probably related to the different densities of luminescent N_x defect states in the a-SiN_xO_y samples with various E_opt.

FIGURE 6

Figure 6. Characteristics of the E_PL dependence on the E_exc at 9K: (A) a-SiN_xO_y (E_opt = 3.50 eV), (B) a-SiN_x (E_opt = 3.24 eV). Inset shows PL spectra under various E_exc at 9K [After Zhang et al. [49], with permission].

FIGURE 7

Figure 7. (A) Integrated PL intensity I_PL dependence on the relative energies of E_exc and E_opt at 9K, under E_exc = 4.5 eV condition [After Zhang et al. [49], with permission]. (B) Schematic diagram of three band energy level structure in our N-type a-SiN_xO_y films. The N-Si-O multi-bonding defect states are located in the band gap below E_Uedge around 0.75 eV [After Lin et al. [69], with permission].

As shown in Figure 2B reported by Kato et al. [32], for both a-SiN_x and a-SiN_xO_y samples, the E_PL increases monotonically with increasing E_Uedge. This relationship is one of the characteristics of PL dominated by band tail states recombination. Herein, we plotted ΔE_stokes (E_Uedge – E_PL) vs. E_Uedge originated from Figure 2B, as shown in the inset of Figure 8. In Figure 8, we also found a similar relationship in our a-SiN_x samples. But for the a-SiN_xO_y samples, ΔE_stokes was independent of E_Uedge and was kept almost constant at around 0.75 eV below the E_Uedge. This phenomenon again shows the different characteristics between band tail state luminescent model and N-Si-O defect state luminescent model.

FIGURE 8

Figure 8. Correlation between the Urbach edge energy E_Uedge and the Stokes shift E_Uedge – E_PL (ΔE_stokes) [After Zhang et al. [49], with permission]. The data in the inset were taken from Kato et al. [32] (with permission). Closed and open symbols are for a-SiN_xO_y and a-SiN_x, respectively.

Moreover, for analyzing the recombination processes of the photo-excited carriers, we also found the different characteristics between O-type a-SiO_xN_y and N-type a-SiN_xO_y samples. Zhang et al. [66] reported the temporal evolution trace of transient PL spectral profiles from nanosecond TRPL measurements. Figures 9A,B record the decay of TRPL intensity as a function of specific delay time at the emission wavelength of 474 nm in a-SiN_xO_y sample, under 8 and 300 K, respectively. The decay of TRPL intensity with specific delay time in the scale of nanosecond after excitation was shown in the Figures 9C,D, and normalized TRPL with SSPL and TIPL was shown in the Figures 9E,F. It is clearly shown that the temporal evolution of transient TRPL spectral profile did not change in both cases. The PL peak energy E_PL was kept almost constant in the time scale of sub-nanoseconds to nanoseconds. This characteristic of transient PL was believed to arise from the behavior of carrier recombination through the defect states, which is obviously different from that of carrier recombination via the band-tail states, as reported by Kato et al. in Figure 3 [32].

FIGURE 9

Figure 9. (A,B) TRPL intensity as a function of decay time and emission wavelength of a-SiN_xO_y films, under (A) 8K and (B) 300K. The color scale represents photon counts. (C,D) Temporal evolution of the transient PL at specific times after excitation, under (C) 8K and (D) 300K. (E,F) Comparison of the SSPL, TIPL, and the normalized TRPL of a-SiN_xO_y films, under (E) 8K and (F) 300K, respectively [After Zhang et al. [66], with permission].

Based on the differing characteristics between band tail and defect models, combined with the results of XPS, EPR with the distinct PL characteristics and absorption spectra [46], we proposed a three-level energy band system to explain the defect states recombination dominated PL model in our N-type a-SiN_xO_y films, which was schematically shown in Figure 7B. The N-Si-O multi-bonding defect states are located in band gap of around 0.75 eV below the E_Uedge. The excited electrons were quickly (~ps) relaxed down to both band tail states and N_x defect states, thereafter recombined with holes via transition between N_x defect states and valence band tail states to give rise to luminescence [49].

Improved Optical Properties in N-type a-SiN_xO_y Films

Despite the benefits of controlling the oxygen content in a-SiN_xO_y films, the optical properties of N-type a-SiN_xO_y films can be improved.

The PL internal quantum efficiency (PL IQE) has reached as high as 60% at the emission wavelength of 470 nm [53]. This higher PL IQE can be explained by the fast radiative recombination rate of excited carriers via the N-Si-O defect states [66]. The faster recombination rate of 10⁸ s⁻¹ in our a-SiN_xO_y films can be comparable to those with direct band gap, e.g., CdSe NCs, reported by Donega et al. [67].

Recently, Valenta et al. [68] employed Si-rich oxynitride (SRON) as well layers to fabricate SRON/SiO₂ superlattice structures. After annealing at 1,150°C for 1 h, the maximum PL external quantum yield (EQY) of nearly 30% in Si nanocrystals has been observed.

Based on the three-level energy band structure, we studied the optical gain property of our N-type a-SiN_xO_y films. The optical gain coefficient G~102 cm⁻¹ has been observed [69]. This kind of amplified emission can also be explained by the carrier recombination radiatively via the N-Si-O bond defect states. As mentioned above, the excited electrons go through relaxation to the N-Si-O defect states very fast (ps), while the radiative recombination lifetime is around ns, which means that the electrons remain in the N-Si-O defect states much longer than relaxation time. Consequently, population inversion of carriers between E_v and N-Si-O defect states can be realized once pumping power exceeds a threshold value.

We also investigated the performance of n-a-SiN_xO_y/p-Si hetero-junction (HJ) LED [30, 69]. The phosphor doped n-a-SiN_xO_y (ρ~2.5 Ω cm, μ~1 cm² V⁻¹ s⁻¹) was deposited on the p-Si substrate to fabricate an ITO/n-a-SiN_xO_y/p-Si HJ LED, which was schematically shown in the inset of Figure 10A. The EL performance was shown in Figure 10B; when the bias voltage reaches ~15 V, the light emission can be observed by naked eye. For comparison, the performance of MIS LED with the structure of undoped a-SiN_xO_y/p-Si was shown in Figure 10C. We can see that the emission threshold voltage of 15 V for MIS LED is higher than that of n-a-SiN_xO_y/p-Si HJ LED. Moreover, the intensity of light output of p-n HJ LED is more than one order of magnitude higher than that of MIS LED under the same forward bias voltage, meanwhile the power efficiency of light output of p-n HJ LED is almost five times higher than that of MIS LED, which was shown in Figure 10D. The reason for the higher EL efficiency is due to the perfect balance between the electrons (n~2.4 × 10¹⁵ cm⁻³) in the n-a-SiN_xO_y layer and the holes injected from p-Si side.

FIGURE 10

Figure 10. (A) Comparison of the PL for n-a-SiN_xO_y film and EL for an n-a-SiN_xO_y/p-Si HJ LED. The inset shows a schematic diagram of n-p HJ LED. (B) EL spectra of n-p HJ LED under different forward bias. (C) EL spectra of ITO/a-SiN_xO_y/p-Si MIS LED under different forward bias. The inset shows the I-V characteristics of MIS LED and p-n HJ LED. (D) The L-V characteristics of MIS LED and p-n HJ LED. The inset shows the light output as a function of the input power for two devices [After Lin et al. [69], with permission].

Finally, for the realization of all-silicon monolithically integrated optoelectronic circuits using standard CMOS processing procedures, the gate controlled Si MOSFET LED should be emphasized. Xu [70, 71] reported that, based on reversed-biased p-n junction of MOSFET operated in the avalanche breakdown mode, the gate controlled Si-LED has been demonstrated. It has been shown that the devices emit light with the wavelength in the range of 100–900 nm and could operate in the GHz range. Since, the device's operation frequency can be modified by gate signals, it can respond to both analog signal transmission and digital signal operation. Lin group reported that Si-based quantum dot (QD) [72] and Si nano-pillar array were used for enhancing the light emission quantum efficiency [73, 74].

On the other hand, direct epitaxial growth of III–V nanostructures on Si substrates is one of the most promising candidates for realizing photonic devices on a Si platform. Liu's group reported [75–77] that, at UCL, high performance InAs/GaAs QD lasers grown on Si substrates have been demonstrated in 2011 and been developing during the last 7 years.

Summary

We reviewed and discussed the present PL mechanisms in the silicon oxynitride films. Apart from preparing methods and conditions, the relative atom concentration ratio of oxygen to nitrogen (O/O+N) is a crucial factor for influencing their chemical bonding configurations and PL mechanisms. For O-type a-SiO_xN_y (0.6 < O/O+N < 0.9), the PL band tail and defect states recombination models have been generally agreed upon. For N-type a-SiN_xO_y (0.05 < O/O+N < 0.2), the N-Si-O bonding defect state recombination model has been proposed. We also analyzed the role of oxygen bonding being played in silicon oxynitride network, which not only reduced the structural disorder but also created N-Si-O (N_x) defect states in the band gap, acting as a new luminescent center. It can be found that, by modifying the chemical composition in these non-stoichiometric silicon oxynitride materials, the optical and electronic properties can be improved. Finally, the n-a-SiN_xO_y/p-Si HJ LED fabricated by the standard CMOS technology has been also demonstrated.

Author Contributions

KC is an advisor and wrote this manuscript. ZL, PZ, and HD carried out experiments and analyzed experimental results. RH did the initial work. XH gave some suggestion and comments on this work.

Funding

The work was supported by National Nature Science Foundation of China Grant No. 11374153 and National Key R&D Program of China Grant No. 2018YFB2200101, and the Scientific Research Program of 333 Project of Jiangsu Province Grant No. BRA2019188.

Conflict of Interest

The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

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