Photoluminescence and carrier transport mechanisms of silicon-rich silicon nitride light emitting device

Applied Surface Science 351 (2015) 1053–1059

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Photoluminescence and carrier transport mechanisms of silicon-rich silicon nitride light emitting device Wugang Liao a , Xiangbin Zeng a,∗ , Wei Yao b , Xixing Wen a a b

School of Optical and Electronic Information, Huazhong University of Science and Technology, Wuhan, Hubei 430074, China Shenzhen Institute of Huazhong University of Science and Technology, Shenzhen 518000, China

a r t i c l e

i n f o

Article history: Received 10 December 2014 Received in revised form 3 June 2015 Accepted 9 June 2015 Available online 18 June 2015 Keywords: Silicon quantum dots Silicon nitride Light emitting diode Photoluminescence

a b s t r a c t Silicon-rich silicon nitride (SRSN) films were prepared on p-type silicon substrates using plasmaenhanced chemical vapor deposition (PECVD). Small size (∼3 nm) amorphous silicon quantum dots (a-Si QDs) were obtained after 1100 ◦ C annealing. Two different excitation sources, namely 325 nm and 532 nm lasers, were introduced to investigate the photoluminescence (PL) properties. The PL bands pumped by 325 nm laser at ∼2.90 eV and ∼1.80 eV were contributed to the radiative centers from N dangling bonds (DBs), while the dominant PL bands at 2.10 eV were ascribed to the instinct PL centers in the nitride matrix. However, PL emissions from band tail luminescence and quantum confined effect (QCE) in a-Si QDs were found under the excitation of 532 nm laser. Light emitting diode (LED) with ITO/SiNx/p-Si/Al structure was fabricated. Intensely red light emission was observed by naked eyes at room temperature under forward 20 V. Three different carrier transport mechanisms, namely Poole–Frenkel (P–F) tunneling, Fowler–Nordheim (F–N) tunneling and space charge limited current (SCLC), were found to fit different electric field regions. These results help to understand the PL mechanisms and to optimize the fabrication of a-Si QD LED. © 2015 Elsevier B.V. All rights reserved.

1. Introduction Enormous researches on Si-based light source have been performed due to its potential applications in all-Si based optoelectronics integrated-circuits (IC) [1–5]. Si quantum dots (Si QDs) are expected to overcome the shortcoming of poor optical properties in bulk Si [6,7]. Si-rich nitride (SRN) [8,9] and Si-rich oxide (SRO) [10–12] films have been generally applied as the surrounding matrix. However, a high turn-on voltage is required to inject the carriers into Si QDs through the SRO due to its relatively high barrier height between Si and SRO [13,14]. Moreover, the existence of Si O double bond at the surface of Si QDs would induce new energy levels which enter into the band gap of Si QDs, resulting in an untunable photoluminescence (PL) emission when the size of Si QDs is below 3 nm [15]. SRN as an alternative matrix has some advantages due to its lower potential barrier for carrier injection, better electrical stability with direct impact on the device operation lifetimes compared with that of SRO [3]. By now, PL and electroluminescence (EL) covering visible zone of SRN containing Si QDs were achieved by changing the size of Si QDs [16]. However, the origin of PL from

∗ Corresponding author. Tel.: +86 027 87544760; fax: +86 027 87544760. E-mail address: [email protected] (X. Zeng). http://dx.doi.org/10.1016/j.apsusc.2015.06.050 0169-4332/© 2015 Elsevier B.V. All rights reserved.

SRN films is still unclear. The dominated PL can mainly be ascribed to the quantum confinement effect (QCE) [17–19], radiative defects [20], and band tail states [21]. What’s worthy noting is that different radiative mechanisms of PL emission were found to be selectively excited under different excitation energies [22]. In addition, the conduction mechanism of the SiNx LED is also still controversial [1,20]. As for device applications, it is essential to understand the carrier transport mechanisms inside the Si QDs systems. In this paper, we investigated the carrier transport mechanisms from SiNx based LED, the origin of room temperature PL emission under different excitation power and the temperature-dependent PL emission from 77 K to 296 K. 2. Experiment Silicon-rich silicon nitride (SRSN) films were deposited using hydrogen-diluted SiH4 (5%) and pure NH3 on p-type (1 0 0) crystalline silicon wafers (thickness: 280 ␮m, resistivity: 1–5  cm) by plasma enhanced chemical vapor deposition (PECVD). The native surface oxide was removed from the silicon wafers by diluted HF acid before deposition. The radio frequency of the PECVD system was 13.56 MHz. The work pressure in the chamber was maintained at 100 Pa and the substrate temperature was kept at 200 ◦ C. The flow rates of SiH4 and NH3 were kept at 45 sccm and

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Fig. 1. (a) HRTEM images, and (b) Raman spectra of the annealed sample at 1100 ◦ C, (c) size distribution histogram of Si QDs.

15 sccm, respectively. The time of deposition was 10 min and the thickness of as-deposited sample was controlled at 100 nm estimated by scanning electron microscope (SEM: Nova nanoSEM). For post annealing, the samples were annealed at 950 ◦ C and 1100 ◦ C in nitrogen for 30 min. High-resolution transmission electron microscopy (HRTEM JEM-2100F) was used to obtain the information about the presence of Si QDs. Room temperature PL (HORIBA Jobin-Yvon LabRAM spectrometer HR 800 UV) has been excited at high energy with the 325 nm line of a He–Cd laser source, and at low energy with the 532 nm laser line. PL emission from 77 K to room temperature in range of 1.4–3.5 eV has also been performed using the 325 nm line of a He–Cd laser. For forming a SiNx LED, the ITO film with a dimensional thickness of 100 nm was deposited by RF sputtering and patterned with 1 mm diameter as cathode. Al anode was evaporated on the rear of the p-type substrates. Finally, a thermal annealing at 400 ◦ C for 10 min was performed in nitrogen to improve the contact property. The EL spectrum was obtained using a charge-coupled device. Current–Voltage (I–V) measurements were carried out by Keithley 4200SCS semiconductor parameter analyzer. 3. Results and discussion The cross-sectional HRTEM was performed to evident the existence of Si QDs in the SRSN sample. As shown in Fig. 1(a), a large number of Si QDs which appear as dark spots are dispersed uniformly in the surrounding nitride matrix for the annealed sample at 1100 ◦ C. However, no Si QDs was found in the as-deposited sample and the annealed sample at 950 ◦ C. The average size of Si QDs is estimated to be ∼3 nm from its histogram in Fig. 1(c). Please note that all the Si QDs with a density of ∼1.5 × 1012 /cm2 were still in the amorphous state though annealed at a relatively high temperature, probably due to the quite small size (∼3 nm) which is hard to be crystallized [23]. The broad band located at ∼480 cm−1 of Raman spectra in Fig. 1(b) can also be attributed to the amorphous Si clusters. The limitation of further growth of Si QDs is due to the formed densificated and near-stoichiometric matrix after 1100 ◦ C annealing [24]. Meanwhile, the ratio of N/Si was calculated to be

∼0.97 by X-ray photoelectron spectroscopy (XPS) analysis, which would essentially lead to the small size of Si nano-particles. In addition, possible nitridation process [25] would even induce the size decrease under 1100 ◦ C annealing in nitrogen. We investigated the PL spectra of the as-deposited and annealed samples pumped by the 325 nm line of a He–Cd laser at 50 mW/cm2 . As displayed in Fig. 2(a), the asymmetric PL exhibits little change in shape, accompanied with an enhancement of the luminescence intensity with the increase of annealing temperature to 1100 ◦ C. We deconvoluted the PL spectra into three parts with the best Gaussian fits. As shown in Fig. 2(b), it is obvious that all the PL spectra are composed of three fixed PL bands at ∼1.80, ∼2.10 and ∼2.90 eV, respectively. Considering the absence of Si QDs in the as-deposited and 950 ◦ C annealed samples, we can exclude the role of QCE by amorphous Si QDs. The weak PL bands at ∼2.90 eV (3.0–0.1 eV) can be contributed to the radiative center from N DBs, which is in accordance with previous works [26]. The low energy band at ∼1.80 eV can be ascribed to radiative recombination between the N4 + and N2 0 levels [27]. What’s worthy noting is that no PL band at ∼2.40 eV from Si DBs was found, which may due to the good passivation of Si QDs by N atoms. The dominant PL band at 2.10 eV with a negligible variation of full width at half maximum (FWHM) 0.55 ± 0.05 eV is similar to the 2.15 eV value of Si3 N4 reference sample [28]. Meanwhile, the intensity of dominant PL is enhanced with an increase of annealing temperature. It can be explained by the enrichment of the silicon nitride phase occurring with the segregation of Si QDs after high temperature annealing [24]. Near-stoichiometric nitride should play an important role in passivation of non-radiative defects. Hence, it is reasonable to contribute the dominant PL bands at 2.10 eV pumped by 325 nm laser to the instinct PL centers in the Si3 N4 matrix. Furthermore, the optical absorbance measurements of all films were performed to study the optical band gaps of the SRSN films. As displayed in Fig. 3, the data were plotted as the well-known Tauc plots, (˛h)1/2 as a function of h, where ˛, h and  represent the absorption coefficient, Planck’s constant and the frequency of the incident light, respectively [29]. Two straight lines from the linear region for each sample can be extrapolated to estimate the

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Fig. 2. (a) Room-temperature PL emission spectra of as-deposited, 950 ◦ C and 1100 ◦ C annealed samples, (b) the experimental PL spectra have been deconvoluted into three Guassians.

absorption edges. The energies of the absorption edge 2 (Eedg2 ) can be ascribed to the optical band gap energy of the SRSN matrix. As shown in Fig. 3, the band gap of the as-deposited films is ∼4.0 eV. After annealing process, the band gap decreases to 3.50 eV at 950 ◦ C and then increases a little to 3.55 eV at 1100 ◦ C. It is well known that hydrogen in a-Si related alloy material would broaden the bandgap [30]. Therefore, the decrease of band-gap can be explained by the effusion of hydrogen from the SRN films after high temperature annealing. The energies of the absorption edges 1 (Eedg1 ), according to Rodriguez et al. [31], can be attributed to the band tail gap in the SiNx matrix. Furthermore, the Eedg1 was calculated to be 2.18, 2.38, and 2.35 eV for the as-deposited, 950 ◦ C and 1100 ◦ C annealed samples, respectively. As displayed in Fig. 3, band tails due to structure disorder in SRSN films play a significant role to the absorption spectra. The existence of CB and VB energy band tails would result in the sub-gap optical absorption. Due to this sub-gap absorption, it can be explained that the PL of as-deposited sample with a relatively low value of Eedg1 was excited by photons with energy of 3.81 eV even though its optical gap is as high as 4.0 eV.

To gain more insights on the light-emitting mechanisms, the characteristic of temperature-dependent PL of the sample annealed at 1100 ◦ C had also been investigated using a He–Cd laser (325 nm) as excitation source. The excitation power density is kept at 50 mW/cm2 . On one hand, the shape and peak positions of PL spectra were almost independent of the temperature. Even the corresponding PL bands obtained after deconvolution of each PL spectrum remained unchanged with different temperature. These suggest that the PL emission cannot originate from QCE in amorphous Si QDs [32], which accords with the room temperature PL analysis. On the other hand, as shown in Fig. 4, the PL intensity shows the typical temperature-dependent characteristics. As displayed in the inset of Fig. 4, the integrated PL intensity slightly increases with increasing temperature up to 100 K from 77 K, and then decreases monotonically with further increasing temperature to 296 K. In general, the integrated PL intensity increases with decreasing temperature due to a decrease of temperature-activated non-radiative recombination process and an increase of excitonic transition [33,34]. According to Rinnert et al. [33], the weak increase

Fig. 3. Tauc plots of as-deposited sample, annealed samples at 950 ◦ C and 1100 ◦ C.

Fig. 4. Temperature-dependent PL spectra of the annealed sample at 1100 ◦ C.

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Fig. 5. (a) Room-temperature PL emission spectra of as-deposited, 950 ◦ C and 1100 ◦ C annealed samples excited by 532 nm He–Ne laser, normalized PL spectra of samples (b) as-deposited and (c) annealed at 1100 ◦ C for various excitation intensities of 532 nm laser of power 50 mW/cm2 .

of PL integrated intensity from 77 K to 100 K can be explained by the saturation effect which is due to an increase in the PL radiative lifetime at low temperature. Furthermore, the temperature dependence of the integrated PL intensity can be described by a simple model based on thermal ionization from localized radiative traps [35]. According to this model, the TD-PL intensity (IPL (T)) at a given temperature (T) can be expressed as Eq. (1). IPL (T ) =  × IPL (T0 ) =

IPL (T0 ) 1 + ˇexp(−EA /kB T )

(1)

In the set of Fig. 4, the dotted line is fitted curve by Eq. (1) with EA = 96.6 meV and ˇ = 70, which are in agreement with previous work by Deshpande et al. [36]. In addition, the PL IQE (300 K) for the sample annealed at 1100 ◦ C was estimated to be 37% based on the work of Hartel et al. [37]. A 532 nm laser as the excitation source was also used to investigate the room temperature PL emission of the as-deposited and the annealed samples. As can be shown in Fig. 5(a), the shape and peak positions of PL spectrum of the sample annealed at 950 ◦ C remained unchanged compared with the as-deposited sample. But its intensity reached the most. When the annealing temperature increased to 1100 ◦ C, the PL peak shifted to a lower energy at ∼1.80 eV. To elucidate the PL mechanisms, we measured the PL emission spectra under various intensities at 532 nm excitation. As shown in Fig. 5(b), the dominant PL peak of the as-deposited sample broadens and

shifts to the low energetic transitions with decreasing excitation intensity. Meanwhile, one broad peak appears at the shoulder of PL spectra when the intensity of excitation source is lower than 50%. This behavior can be attributed to the band tail luminescence [38] which assumes that excited carriers thermalize towards low energetic transitions. A red-shift is expected when the amount of excited carriers decreases with decreasing excitation intensity. The annealed sample at 950 ◦ C shows the similar behavior with the asdeposited sample. Considering the absence of Si QDs in both the annealed sample at 950 ◦ C and as-deposited sample, both the PL emissions can be ascribed to band tail luminescence of SRN films. It is worth noting that there is a relatively large absorption emission Stokes shift (Eedg1 –EPL ∼ 0.4–0.6 eV) for as-deposited sample and annealed sample at 950 ◦ C. The enhancement of luminescence intensity for the annealed sample at 950 ◦ C, comparing with the as-deposited sample, can be attributed to the better passivation of non-radiative defects in the SRN film after annealing. However, the PL emission of the sample annealed at 1100 ◦ C shows a totally different evolution with decreasing excitation intensity. As displayed in Fig. 5(c), the shape and peak positions of normalized PL spectra are almost the same even when the excitation intensity is reduced to 10%. Thus, the contribution from band tail states in the PL spectra of the annealed sample at 1100 ◦ C can be excluded. PL emission from Si QDs may arise because different radiative mechanisms are selectively excited upon usage of

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Fig. 7. P–F tunneling, F–N tunneling and SCLC fitting of J–V data.

Fig. 6. I–V curve of the ITO/SiNx/p-Si/Al light-emitting device at room temperature.

different excitation energies [22]. The average size of amorphous Si QDs is calculated to be 3 nm in the annealed sample at 1100 ◦ C according to the empirical equation E(PL) = 1.56 + 2.40/a2 [6], where a is the average size of Si QDs. The good correlation between the Si QDs size observed by HRTEM and that calculated by QCE model demonstrates that the PL of the annealed samples at 1100 ◦ C is dominated by QCE in amorphous Si QDs. Actually, direct band to band recombination may be the dominant process because the number of disorder-induced tail states is strongly lowered in amorphous Si QDs [39]. The ITO/SiNx/p-Si/Al light-emitting device was fabricated by sputtering ITO films on the SRN and thermal evaporation Al on the backside of Si substrate. The current–voltage (I–V) characteristic of SiNx sample annealed at 1100 ◦ C for 30 min is shown in Fig. 6. A forward dc bias voltage from −20 V up to 40 V was applied to the bottom electrode. The breakdown voltage of our LED is higher than 70 V whose electrical field is 7 MV/cm. Under the reverse bias, the current is negligible. Actually, the SRN-c-Si interface behaves like a Schottky diode, where the SRN film plays the role of a metal [40]. Thus the reverse bias current is limited by the diode. Under forward bias, the majority of voltage is dropped across the SRN film. It is obvious that the current increases very slowly at the beginning and then goes up dramatically when the applied voltage is larger than 5 V. As shown in the inset of Fig. 6, the red light emission can be observed with naked eyes in the dark room when the forward dc bias voltage increases to 20 V. The electroluminescence (EL) spectra shows the main emission peak at ∼1.85 eV and a shoulder peak at ∼2.1 eV when measured at 3 MV/cm. Under the dc voltage driving, the electrons and holes are injected from ITO electrode and p-Si substrate, respectively. The wavelength of red emission is consistent with the PL emission at ∼1.80 eV, which also elucidates the recombination of electron–hole pairs in amorphous Si QDs. It was found that light emission can be only observed under the forward bias voltage, which illustrates that the p-type silicon substrate can offer little electrons under reverse bias. With further increase of forward voltage, the intensity of EL increases while the shape of that keeps unchanged, which indicates that injected holes and electrons of higher energy into Si QDs results in higher recombination ratio under relatively higher applied voltage. Since the band gap of amorphous Si QDs is fixed, the EL emission should be scarcely affected. In order to investigate the carrier transport mechanisms in our Si QD LEDs, as shown in Fig. 7, we measure the current density–voltage (J–V) of our Si QDs LED. Usually, P–F tunneling is dominated for SiNx-based LED under a low electrical field [20].

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The injection of carriers is thermally activated in P–F transport mechanism. Under this bias condition, the substrate undergoes an accumulation of holes. The analytic relationship of P–F tunneling conduction mechanism is described as equation (2) [20,41]:



JPF = CEexp −

q(f −



qE/εr ε0 )



(2)

KT

where q is elementary charge, E is the electric field, f is the energy barrier height for P–F emission, εr is the relative static permittivity, ε0 is the permittivity in free space, K is the Boltzmann constant, and T is the temperature. As shown in the inset of Fig. 7, this P–F transport of carriers has a linear relationship of ln(J/E) and E1/2 when the applied voltage is less than 5 V corresponding to an electrical field of 0.5 MV/cm. This indicates that the carrier transport process is governed by the P–F tunneling in this regime. To study the transport mechanism beyond 5 V, the F–N tunneling effect based on the ln(J/E2 )–1/E relation is used to analyze the carrier transport mechanism and fit the experimental data. The F–N tunneling is described as equation (3) [14]: JFN

A 2 = E exp 4B



−

3/2

2BB 3E



,B=

4



2m∗ q

h

(3)

where A and B are the constants, m* is the effective of electron mass, and B is the effective barrier height for carrier injection in F–N model. As shown in the inset of Fig. 7, the applied voltage to initiate F–N tunneling is ranging from 5.5 V to 15.5 V. The F–N model is based on the hot electron tunneling through a triangle-shaped energy barrier formed in the dielectric layer due to high applied electric field [20]. Neither P–F nor F–N tunneling mechanisms suits the experimental data with higher electric field above 1.55 MV/cm. However, as shown in the inset of Fig. 7, the current above a voltage value 12.5 V increased along a straight line in the log–log plot, corresponding to an I ∝ Vm dependence. This characteristic at 295 K can be explained using the space charge limited current (SCLC) model with an exponential density of traps. To further confirm the carrier transport mechanism under high bias voltage, we have also investigated the temperature-dependent I–V behaviors in logarithmic scales ranging from 295 K to 398 K. As displayed in Fig. 8, the experiment data ranging from 295 K to 398 K are well fitted by the SCLC model at the high bias. In addition, the curves of higher temperature show smaller slopes than those of lower temperature. This behavior is also consistent with the SCLC theory. Therefore,

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technology project of Ministry of Education of the People’s Republic of China (Grant No. 62501040202). We also acknowledge the Analytical and Testing Center of Huazhong University of Science and Technology for the PL and UV–vis absorption spectra measurements and Optoelectronic Micro/nano Fabrication Faculty of Wuhan National Laboratory for Optoelectronics for the I–V measurements. References

Fig. 8. I–V curves for the ITO/SiNx/p-Si/Al structure at different temperature from 295 K to 398 K.

the SCLC model could be the predominant conduction mechanism when the electric field is above 1.25 MV/cm. From the analysis of above, linear fittings in different regions indicate different carrier transport mechanisms under different electrical fields. P–F tunneling and SCLC mechanisms dominant the carrier transport process when the applied electrical field is under 0.5 MV/cm and above 1.25 MV/cm. The F–N tunneling mechanisms would contribute the carrier transport under the moderate electric field. 4. Conclusions In conclusion, SRSN thin films were prepared by PECVD and treated at 950 ◦ C and 1100 ◦ C in N2 flow. Dense a-Si QDs with small size (∼3 nm) were precipitated after 1100 ◦ C annealing. Three pinned PL peaks at 1.80, 2.10 and 2.90 eV can be obtained by deconvolution of PL spectra for all films pumped by the 325 nm (3.81 eV) laser. Both the weak PL band at ∼2.90 eV and the band at ∼1.80 eV could be contributed to the radiative centers from N DBs. The dominant PL band at 2.10 eV was ascribed to the instinct PL centers in the Si3 N4 matrix. The characteristic of temperature-dependent PL of annealed sample at 1100 ◦ C has also been investigated using a 325 nm laser as the excitation source. The shape and peak positions of PL spectra were almost independent of temperature. These also suggested that the PL emission excited by 325 nm laser did not originate from the QCE in amorphous Si QDs. However, the PL emission pumped by 532 nm (2.33 eV) laser exhibited a totally different behavior. Both the PL emissions of as-deposited and annealed sample at 950 ◦ C at ∼1.80 eV were ascribed to band tail luminescence of SRN films, while the PL emission of 1100 ◦ C annealed sample was ascribed to QCE in a-Si QDs. LED with metal-nitride-semiconductor (MNS) structure was fabricated. Intense red light emission was observed by naked eyes at room temperature under forward 20 V and corresponding EL was also measured. We investigated the carrier transport mechanisms for SiNx-LED under different electrical fields. The P–F tunneling and SCLC mechanism dominant the carrier transport process when the applied electrical field is under 0.5 MV/cm and above 1.25 MV/cm, respectively. The carrier transport process under the moderate electric field can be contributed to the F–N tunneling. Acknowledgments This work was supported by the National Natural Science Foundation of China under No. 51472096, Central University Basic Research Funds (Grant No. 2014NY004) and the supporting

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