AlN nanoparticle superlattices

ARTICLE IN PRESS

Physica E 35 (2006) 81–87 www.elsevier.com/locate/physe

Growth and visible photoluminescence of SiCxNy/AlN nanoparticle superlattices M. Xua,b,, S. Xub, S.Y. Huangb, J.W. Chaib, V.M. Ngb, J.D. Longb, P. Yangc a

Institute of Solid State Physics and School of Physics and Electronic Egineering, Sichuan Normal University, Chengdu 610068, PR China b Plasma Sources and Applications Center, NIE, Nanyang Technological University 1 Nanyang Walk, 637616, Singapore c Singapore Synchrotron Light Source, National University of Singapore, 5 Research Link, 117603, Singapore Received 1 May 2006; received in revised form 29 May 2006; accepted 29 May 2006 Available online 28 July 2006

Abstract Luminescent SiCxNy/AlN nanoparticle superlattices (NPSLs) were fabricated by sputtering SiC and Al targets alternatively under a mixture gaseous flow of N2, Ar, and H2 by plasma RF magnetron sputtering. It was found that the best periodic structures and greatest content of SiCxNy nanoparticles occur in the case in which the sublayer thickness of SiCxNy and AlN are 42 and 3 nm, respectively. In this case, the NPSLs exhibit strong photoluminescence (PL) either before or after annealing. Annealing at different temperature (Tanneal) revealed that the PL intensity first increases with increasing Tanneal and PL peak is red-shifted when Tanneal4650 1C; the PL intensity reaches a maximum as Tanneal ¼ 1100 1C. A microstructure model that takes into account the formation energy of SiCxNy nanocrystal limited by AlN sublayers was developed to tentatively explain the formation of SiCxNy nanocrystal in SiCxNy/AlN NPSLs. Based on the structural analysis by X-ray diffraction and Fourier-transform infrared spectra, the PL of SiCxNy/AlN NPSLs was discussed. r 2006 Elsevier B.V. All rights reserved. PACS: 71.23.Cq; 78.20.Ci; 81.40E Keywords: SiCxNy nanoparticle film; Superlattices; Photoluminescence

1. Introduction Si-based nanostructures have exhibited their important role in integrated photonics and optoelectronics [1]. Since 1990, much attention has been paid to porous Si structure because of its high luminescence intensity [2]. Unfortunately, its fragility and unstable luminescence will hardly survive modern Si technology. On the other hand, the surfaces of the remaining Si wires and dots are highly reactive. To solve these problems, the research of Si nanostructures was recently focused on the preparation of Si nanocrystals embedded in an oxide host [3]. The control of size, passivation, and density of the nanocrystalline Si Corresponding author. Institute of Solid State Physics and School of Physics and Electronic Egineering, Sichuan Normal University, Chengdu 610068, PR China. E-mail address: [email protected] (M. Xu).

1386-9477/$ - see front matter r 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.physe.2006.05.006

(nc-Si) was explored to achieve a strong photoluminescence (PL) in the visible spectral range. For example, Zacharias and Streitenberger [4] suggested a SiO/SiO2 superlattice approach to well control the nc-Si size and density and hence, improve the PL. Besides the research of nc-Si, the photoluminescent porous SiC [5], a-SiN:H film [6] and SiCN crystallinite [7] have been paid attention in the past years. To grow high-quality SiC film, i.e., cubic silicon carbide (b-SiC), a very high growth temperature is usually required. As a substitute, SiCxNy was synthesized firstly by Chen et al. [7], the formation temperature of which is 800 1C. To date, there has been no detailed work on the luminescence property of SiCxNy film. Recently, most research works have focused on the self-assembly of nanoparticle film because of its simple and cost-effective process. Due to the quantum confinement effects (QCE) at a nanoscale, the nanoparticle films have displayed promising characteristics for the fabrication of

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M. Xu et al. / Physica E 35 (2006) 81–87

light-emitting diodes, non-linear optical devices and electronic devices. Here, we present a SiCxNy/AlN superlattice method for the preparation of SiCxNy nanoparticles, which is expected to enable the control of PL by limiting the nanoparticle size and density. It was found that, when the SiCxNy sublayer is thin, there is a size-dependent shift of the PL position in SiCxNy/AlN nanoparticle superlattices (NPSLs). Strong PL was also observed in our SiCxNy/AlN NPSLs. When the sublayer thickness of SiCxNy and AlN are 40 and 3 nm, respectively, the NPSLs exhibit the best periodic structures and strongest PL whether before or after annealing. Annealing at 1100 1C was found to result in the formation of the greatest number of SiCxNy nanocrystals and strongest PL. Furthermore, it was revealed that the PL intensity first increases with increasing Tanneal and PL peak is red-shifted when Tanneal4650 1C; the PL intensity reaches a maximum as Tanneal ¼ 1100 1C. We suggested a microstructure model that took into account the formation energy of SiCxNy nanocrystal limited by AlN sublayers to tentatively explain the formation of SiCxNy nanocrystal. The PL of SiCxNy/AlN NPSLs was discussed in combination with the results of X-ray diffraction (XRD) and Fourier-transform infrared spectra (FTIR). 2. Experimental details Prior to the growth experiments, Si(1 0 0) substrates used in this study were chemically cleaned before being loaded into the sputtering chamber. A typical base pressure of 3  103 Pa is routinely achieved by a turboelectric pumping system. SiCxNy/AlN NPSLs were prepared by the following procedure. An AlN buffer layer was first grown on Si(1 0 0) wafer by sputtering a pure Al target under the gas flow rate N2:Ar:H2 ¼ 32.1:21.4:3.2 sccm with the power of 300 W at deposition temperature of 350 1C for 30 min. Thereafter, SiCxNy nanoparticle film and amorphous AlN sublayer were alternatively deposited by consecutive sputtering a pure SiC target (Si:C1:1) under the gas flow rate N2:Ar:H2 ¼ 6.4:35.2:9.6 sccm with the power of 500 W and Al target with the same growth condition as that of buffer. The substrate temperature was kept at 400 1C. All the superlattice samples have the same period number of 20. After deposition the samples were annealed at 1100 1C for 20 min under a flowing N2 atmosphere. The superlattice structures were checked by X-ray reflectivity with a step size 0.0051 at Singapore Synchrotron Light Source. The wavelength was set at 1.54 A˚. The width of slits in front of source and detector were set at 0.2 and 0.6 mm, respectively. The cross-section of the samples as-grown and annealed at 1100 1C was observed by a PHILIPS CM300 high-resolution transmission electron microscope (HRTEM). The PL were measured using a 514.5 nm radiation of an Ar+ laser by a Renishaw micro-Raman System 1000 with PL capability at room temperature. The position of the sample was adjusted to avoid the effect of thin-film interference on the PL

spectrum. All the spectra were corrected for system response. XRD with an incident X-ray wavelength of 1.54 A˚ (Cu Ka line) is used to characterize the microstructures of the nanoparticle films with a step size 0.021 at a fixed power of 16 KW. Also, FTIR transmission was adopted to characterise the microstructures of the films in the range of 400–4000 cm-1 with a resolution of 4 cm1 with a Perkin Elmer Spectrum One FTIR spectrometer. 3. Experimental results Fig. 1(a) shows the TEM image of an as-grown SiCxNy/ AlN NPSLs, wherein the bright and dark layers are the SiCxNy and AlN layers, respectively. Good periodic structures are evidenced. There are many amorphous SiCxNy nanoparticles embedded in each SiCxNy sublayer. The composition of SiCxNy sublayer was determined by XPS measurement of a pure SiCxNy film. It was found that the SiCxNy sublayer is composed of 46.4% Si, 10.04% C, 20.36% N, and 23% O. In contrast, there is little O content in the as-grown SiCxNy/AlN NPSLs or the SiCxNy/AlN superlattices annealed not at high temperature, as discussed in the following. The SiCxNy nanoparticles with the typical size of a few nanometers are mainly located at the SiCxNy/ AlN interfaces as shown in Fig. 1(b). For the pure SiCxNy nanoparticle films grown on AlN buffer (not shown here), we observed many amorphous cylinder nanostructures with the size of a few tens of nanometer, indicating that the SiCxNy nanoparticles in the superlattices were indeed formed by the limitation of AlN sublayer. After annealing at 1100 1C, there are nanocrystals o5 nm at the SiCxNy/ AlN interfaces (Fig. 1(c)), in accordance with the XRD results given in the following. X-ray reflectivity (XRR) was applied to investigate the effect of the sublayer thickness on the superlattice structures. Fig. 2 displayed the XRR variation of the superlattices with the SiCxNy sublayer thickness. When the sublayer thickness of SiCxNy and AlN are 42 and 3 nm, respectively, one can see up to three orders of diffraction peak in the XRR curve, indicating a best periodic structure in this case. By theoretically fitting the measured XRR curve (see the inset in Fig. 2), we are able to extract the interface roughness [8]. It was found that, when the sublayer thickness of SiCxNy and AlN are 42 and 3 nm, respectively, the interface roughness reaches a minimum (1 nm), indicative of the smoothest SiCxNy/AlN interface in this case. Furthermore, we checked the effect of the AlN sublayer by changing its thickness from 1.2 to 6 nm. The optimal value of the AlN sublyer thickness was found to be 3 nm. To check the crystal structures of the SiCxNy/AlN NPSLs, XRD measurement was performed. There is no diffraction peak relevant with SiCxNy crystal in the XRD spectra of as-grown samples. After annealing at 1100 1C, the XRD spectra (see Fig. 3) of the SiCxNy/AlN NPSLs exhibit broad diffraction peaks between 341 and 361. Because the AlN buffer under such condition does not give

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Fig. 1. Cross-sectional TEM micrographs of the SiCxNy/AlN NPSLs sample: (a) as-grown SiCxNy/AlN nanoparticle superlattices; (b) SiCxNy/AlN NPSLs annealed at 1100 1C and (c) a SiCxNy nanocrystal formed at the SiCxNy/AlN interface.

Intensity(a.u.)

109

108

106

107 106 105

7

t

10

0.2 =7nm

SiCN

0.4

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θ (degree)

tSiCN=21nm

105 t

t

103

t

105

tSiCN=70 nm

104

tSiCN=42 nm

103

tSiCN=21 nm

102

tSiCN=10.5 nm

4

SiCN

Intensity (a.u.)

10

Counts

107

tSiCN=42nm Square:Experimental Line:Simulated

109

1011

=42nm

=56nm

SiCN

SiCN

tSiCN=7.0 nm

101

=70nm

1

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0.4

0.6

0.8

1.0

1.2

θ (degree)

100 20

25

30

35

40

45

50

2θ θ (degree)

Fig. 2. X-ray reflectivity variation of the SiCxNy/AlN NPSLs with the SiCxNy sublayer thickness. Inset shows the simulated result of the sample with the sublayers of 42 nm SiCxNy and 3 nm AlN, respectively.

Fig. 3. X-ray diffraction spectra of the SiCxNy/AlN NPSLs.

any diffraction information (not shown), the broad peak appearing at around 351 can be identified as the SiCxNy structure [9]. It was further found that, the intensity of the

broad peak reaches a maximum as the sublayer thickness of SiCxNy and AlN are 42 and 3 nm, respectively. Furthermore, we investigated the effect of annealing

ARTICLE IN PRESS M. Xu et al. / Physica E 35 (2006) 81–87

temperature. It was noted that, at low Tanneal, SiCxNy nanocrystal starts to form. The content of SiCxNy nanocrystal increases with the increase of Tanneal and reaches a maximum after annealing at 1100 1C. The further increase in Tanneal will cause the content of SiCxNy nanocrystal to significantly decrease. Shown in Fig. 4(a) are the FTIR spectra of the SiCxNy film and SiCxNy/AlN NPSLs. One can see that, under the same total thickness of SiCxNy, the superlattices exhibit a greater content of SiCxNy nanoparticle in comparison to that of SiCxNy thin film, as indicated by the C–N and Si–N wagging bonds located at 900–1000 cm-1 and CRN stretching bond (2200 cm1). The Si–O bond (1000– 1100 cm1) appears after annealing at 1100 1C. For the superlattices, the content of Si–O bond monotonously

SiCxNy (35min)

60000

[SiCxNy/AlN]20 tSiCxNy=1.5min, tAlN=5min Solid: As-grown Circles: Annealed

50000 Intensity (a.u.)

84

1100°C/20min

40000

30000

20000 As-grown@400°C

10000

0 600

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Emission wavelength (nm)

150 Transmission (a.u.)

Fig. 5. Comparison of the PL spectra of the SiCxNy film and SiCxNy/AlN NPSLs.

C-N Si-N N-Hn

100

C-C

C-N

As-grown SiCN film Annealed SiCN film

Si-O

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As-grown Superlattices Annealed Superlattices

500

1000

(a)

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Wavelength

3000

Transmission (a.u.)

C-N

Si-O

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(cm-1)

N-H

C-C

1200°C 1100°C C-N Si-N

800°C 650°C As-grown

1000 (b)

2000

3000

4000

Wavelength (cm-1)

Fig. 4. (a) FTIR spectra of the SiCxNy film and SiCxNy/AlN NPSLs and (b) variation of the IR absorption of SiCxNy/AlN NPSLs with Tanneal.

increases with increasing Tanneal (see in Fig. 4(b)). As Tanneal increases up to 1200 1C, the C–N and Si–N wagging bonds almost disappear. It should be pointed out that, we observed the obvious C–N and Si–N bonds, but no Si–C bond in the FTIR spectra of SiCxNy/AlN NPSLs. This suggests that the broad peak in the XRD curve should be attributed to the SiCxNy nanocrystal structure, and not SiC nanocrystal. A comparison on the PL of SiCxNy film and SiCxNy/ AlN NPSLs was made in Fig. 5. It can be clearly seen that the PL of SiCxNy/AlN NPSLs is much stronger than that of SiCxNy film whether before or after annealing. Fig. 6(a) displayed the PL results of SiCxNy/AlN NPSLs with different SiCxNy sublayer thickness (tSiCxNy). One can see that, when 10 nmptSiCxNyp70 nm, the PL position is redshifted as tSiCxNy increases (see in Fig. 6(b)). When tSiCxNyp10 nm, the PL peak also exhibits a red-shift. When tSiCxNy is very thick, i.e., tSiCxNyX70 nm, the PL position is independent of tSiCxNy. The PL intensity was found to increase with tSiCxNy and reach a maximum as tSiCxNy ¼ 42 nm. The further increase of tSiCxNy will cause the PL intensity to decrease. To check that the increase of PL intensity is caused by the thickness or superlattice effect, we also fabricated a 10-bilayer superlattice sample with the SiCN sublayer thickness of 42 nm. In comparison to that with 20 bilayers and SiCN sublayer thickness of 21 nm, the PL of the former is much stronger (not shown), indicative of the significant influence of superlattice effect on the PL as the SiCN sublayer thickness is below 42 nm. Annealing at 1100 1C can strongly enhance the PL intensity, but the trend of the variation of the PL intensity versus tSiCxNy does not change, as shown in Fig. 6(c). In Fig. 6(d), we showed the variation of the PL with Tanneal. It

ARTICLE IN PRESS M. Xu et al. / Physica E 35 (2006) 81–87

tAlN=3.0nm

Intensity (a.u.)

tSiCN=7nm tSiCN=10.5nm tSiCN=21nm tSiCN=42nm tSiCN=56nm tSiCN=70nm tSiCN=875nm

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PL position (eV)

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Emission wavelength (nm)

Fig. 6. (a) PL spectra of the SiCxNy/AlN NPSLs with different SiCxNy sublayer thickness (tSiCxNy), (b) variation of the PL position of the SiCxNy/AlN NPSLs with tSiCxNy, (c) PL spectra of the as-grown and annealed SiCxNy/AlN NPSLs with tSiCxNy and (d) PL spectra of the SiCxNy/AlN NPSLs with Tanneal.

was seen that the PL intensity first increases but PL peak does not shift at low temperature; when Tanneal4650 1C, the PL increases again in intensity and is accompanied by a red-shift of peak position; the PL intensity reaches a maximum when Tanneal ¼ 1100 1C; a further increase in Tanneal degrades the PL of SiCxNy/AlN NPSLs. 4. Discussions The superlattice method to limit the Si nanocrystal size was first explored by Zacharias and Streitenberger [4]. Using an empirical model that takes into account the different interfacial energies and materials, they successfully explained the exponential scaling of crystallization temperature with the layer thickness of amorphous Si

sublayer and revealed a critical thickness for amorphous Si sublayer below which no crystallization can occur for the Si/SiO2 system. In the model suggested by Zacharias and Streitenberger [4], they assumed that the crystallization nucleus is symmetrically embedded in the amorphous material between the oxide interfaces and is cylindrical in shape. In our SiCxNy/AlN NPSLs, we assume that the crystallization nucleus is on the surface of amorphous AlN sublayer because of their close lattice constants and thermal expansion coefficients at 300 K favoring of the formation of SiCxNy nanocrystal on the AlN surface. The TEM results have supported this issue. On the other hand, we still assume that the SiCxNy nanocrystal is cylindrical in shape. Fig. 7 gives a schematic diagram of the SiCxNy nanocrystal formed on the AlN surface.

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r

SiCxNy h

AlN sublayer Fig. 7. Schematic diagram of the SiCxNy nanocrystal formed on the AlN surface.

Based on these assumptions above, the Gibbs free energies for the cylindrical nuclei and amorphous cylinder with radius r and height h can be expressed as G c ¼ pr2 hG vc þ pr2 gA2c þ ðpr2 þ 2prhÞgac ,

(1)

G a ¼ pr2 hG va þ pr2 gAa ,

(2)

where Gvc and Gva are the Gibbs free energies per unit volume of the bulk crystalline and amorphous phase, respectively; gAc (or gAa) is the interface free energy per unit area between amorphous AlN and crystalline (or amorphous phase); gac is the interface free energy per unit area between amorphous and crystalline SiCxNy. As a subsequence, the nucleation energy barrier can be given by the difference of the Gibbs free energies DG ¼ G c  G a ¼ pr2 hDG v þ pr2 DgAac þ ðpr2 þ 2prhÞgac , (3) where DGv ¼ G va  G vc 40 and DgAac ¼ gAc  gAa . The nucleation barrier of SiCxNy nanocrystal can be given by the maximum of DG. By solving the equations qDG=qr ¼ 0 and qDG=qh ¼ 0, one can obtain the critical radius and cylinder height of SiCxNy nanocrystal r ¼

2gac , DG v

(4)

h ¼

2ðgac þ DgAac Þ . DG v

(5)

Then the nucleation barrier is given by DG  ¼

4pg2ac ðgac þ DgAac Þ . DG 2v

(6)

If the SiCxNy nanocrystal is formed directly from the amorphous SiCxNy network, the nucleation barrier is DG  ¼

8pg3ac . DG 2v

(7)

It can be seen from the Eq. (5), when the SiCxNy sublayer thickness is too thin, no crystallization can occur for the SiCxNy/AlN system. As such, for the SiCxNy/AlN superlattices with tSiCxNy ¼ 7 nm, no SiCxNy nanocrystal can form even after high-temperature annealing. The values of gac, DgAac and DGv are not known yet, so h cannot be given by calculation. Still, from our experimental

results, h can be estimated to be at 7–10 nm. For the formation temperature of SiCxNy nanocrystal, it can be estimated from kTDG when the amorphous-to-crystalline transition is induced by a thermally activated process. Chandru [10] reported that the AlN crystal can be formed at 350 1C by RF magnetron sputtering. For SiCxNy, the formation temperature of SiCxNy crystalline is up to 800 1C. Since there are close lattice constants and thermal expansion coefficients between AlN and 6H–SiC, it is most probably that there are close lattice constants and thermal expansion coefficients between AlN and SiCxNy. Reasonably, DgAac should be smaller than gac. So, the crystallization of SiCxNy sublayer thicker than 10 nm on the AlN surface occurs at a lower temperature (i.e., 650 1C for the SiCxNy/AlN superlattice system). It is noted that the model of Zacharias and Streitenberger describing the crystallization of ultrathin amorphous Si layers embedded in thermally stable SiO2 layers is only valid for thin films of a thickness below 7 nm. In our samples, the SiCxNy nanocrystals are formed on the AlN surface, but the other surfaces of the SiCxNy nanoparticles is very likely adjacent to the SiO2 network from the FTIR results, which works as an energy barrier against the growth of the SiCxNy nanocrystal. So, the size of the SiCxNy nanocrystal embedded in the SiCxNy/AlN superlattice is very small (o5 nm). Furthermore, since there is a critical radius of SiCxNy nanocrystal, it is reasonable that the smooth interface is in more favor of the crystallization of SiCxNy on the AlN surface. In Fig. 2, we have shown that the smoothest SiCxNy/AlN interface can be achieved when the sublayer thickness of SiCxNy and AlN are 42 and 3 nm, respectively. As such, it is not surprising that there exist more SiCxNy nuclei in the SiCxNy sublayers. As a result, the most SiCxNy nanocrystal content was observed after annealing at 1100 1C, as shown in Fig. 3. Next, we turn to the PL of SiCxNy/AlN NPSLs. In general, the PL intensity I is given by [11] t I / sf N, (8) tR where s, f, t, tR, and N are the excitation cross-section, photon flux, lifetime, radiative lifetime, and total number of emitting centers, respectively. It is believed that the passivation of nanoparticles can effectively enlarge the excitation cross-section and increase the lifetime. Furthermore, the emitting center should be relevant with the SiCxNy nuclei or nanocrystal for the as-grown SiCxNy/AlN NPSLs. After annealing, the Si oxides will also affect the PL in SiCxNy/AlN NPSLs. By considering so, the PL of SiCxNy/AlN NPSLs can be understood as follows. (1) In comparison to the pure SiCxNy film, the SiCxNy/ AlN NPSLs with the same total thickness of SiCxNy have exhibited the greater content of SiCxNy nanoparticle (SiCxNy nuclei or nanocrystal), as revealed by the FTIR spectra. That is, the SiCxNy/AlN NPSLs have more emitting center. Since the growth conditions of

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SiCxNy layer are the same between the pure SiCxNy film and SiCxNy/AlN superlattices, the s, t and tR are quite similar. It is reasonable that the PL intensity is mainly dependent on the content of SiCxNy nanoparticle. So, the stronger PL was observed in the SiCxNy/ AlN NPSLs. (2) The red-shift of the PL position of SiCxNy/AlN NPSLs at tSiCxNyp70 nm can be attributed to the QCE of SiCxNy nanoparticle. In Fig. 2(b), the SiCxNy nanoparticle size was shown to be a few nanometer at tSiCxNy ¼ 42 nm. The shorter growth time of SiCxNy sublayer, the smaller SiCxNy nanoparticle size, which causes the blue-shift of the PL position. (3) A maximum of PL intensity observed as tSiCxNy ¼ 42 nm can be also accounted for in combination with the content of nanoparticle. In this case, the annealed SiCxNy/AlN NPSLs have exhibited the most content of SiCxNy nanocrystal, as shown in Fig. 3. It can be deduced that the as-grown SiCxNy/AlN NPSLs should have the most number of SiCxNy nuclei. Hence, a maximum of PL intensity was observed as tSiCxNy ¼ 42 nm and tAlN ¼ 3 nm, whether before or after annealing. (4) The effect of annealing on the PL of SiCxNy/AlN NPSLs is well explained in terms of the PL model suggested by Qin et al. [12]. In this model, Qin et al., suggested that there were two competitive processes, namely, the quantum confinement (QC) process and the quantum confinement-luminescence center (QCLC) process, in the PL of the nanoscale Si/Si oxide systems. In the case that the most probable size of Si nanoparticle is smaller than a critical one, the QC process dominates the PL. When the most probable size of Si nanoparticle is larger than a critical one, the QCLC process dominates the PL. In our SiCxNy/AlN NPSLs, the PL most probably originates from SiCxNy nanoparticle (QC center, PL  630 nm) and Si oxide (QCLC, PL  680 nm). When Tanneal is low (i.e., 650 1C or lower), the oxidation is not obvious (see Fig. 4(b)). As such, the QC process dominates the PL, which exhibits the shorter PL wavelength. As Tanneal increases up to 800 1C or higher, the oxidation becomes significant (see Fig. 4(b)) and SiCxNy nanocrystal increases. As result, the PL is dominated by Si oxide and hence, red-shifted. It is noted that, the content of SiCxNy nanoparticle is remarkably diminished and PL intensity strongly decreases as Tanneal increases up to 1200 1C. This suggests that process b suggested by Qin [12] should play a key role in the PL in SiCxNy/AlN NPSLs. That is, in the annealed SiCxNy/AlN NPSLs, photoexcitation occurs in a SiCxNy nanoparticle (nanocrystal or nanocrystal nuclei), while photoexcited electrons and holes in the SiCxNy nanoparticle transfer in the luminescence centers in the SiOx layer surrounding the SiCxNy nanoparticle and radiatively recombine there. After annealing at 1200 1C, there are numerous Si oxides but the SiCxNy nanoparticles become few, so that the PL intensity strongly decreases.

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5. Conclusions SiCxNy/AlN NPSLs were grown by sputtering SiC and Al targets alternatively by plasma RF magnetron sputtering. When the sublayer thickness of SiCxNy and AlN are 42 and 3 nm, respectively, the best periodic structures, greatest content of SiCxNy nanoparticles and strongest PL were observed. The PL intensity first increases with increasing Tanneal and it is red-shifted when Tanneal4650 1C; a maximum of the PL intensity occurs in the case of Tanneal ¼ 1100 1C; a further increase of Tanneal will degrade the PL. Taking into account the formation energy of SiCxNy nanocrystal limited by AlN sublayers, we developed a microstructure model to tentatively explain the formation of SiCxNy nanocrystal in SiCxNy/AlN NPSLs. Furthermore, the PL of SiCxNy/AlN NPSLs was explained in terms of the structural analysis by XRD and FTIR. Acknowledgments This work is supported by A*STAR, Singapore (Project No. 012-101-0024) and partly performed at SSLS under NUS core support C-380-003-003-001, A*STAR/MOE RP 3979908M and A*STAR 12 105 0038 Grants. The support by Education Bureau of Sichuan Province, PR China (Project No. 2005A092) is also acknowledged. References [1] O. Bisi, S.U. Campisano, L. Pavesi, F. Priolo (Eds.), Silicon Based Microphotonics: from Basic to Applications, IOS Press, Amsterdam, 1999. [2] L.T. Canham, Appl. Phys. Lett. 57 (1990) 1045. [3] H. Takagi, H. Ogawa, Y. Yanazaki, T. Nakagiri, Appl. Phys. Lett. 56 (1990) 2397; T.S. Iwayama, S. Nakao, K. Saitoh, Appl. Phys. Lett. 65 (1994) 1814; K.D. Hirschman, L. Tsybeskov, S.P. Duttagupta, P.M. Fauchet, Nature 384 (1996) 338; S. Hayashi, K. Yamamoto, J. Lumin. 70 (1996) 352; L. Pavesi, L.D. Negro, C. Mazzoleni, G. Franzo, J.P. Prolo, Nature 408 (2000) 440. [4] M. Zacharias, P. Streitenberger, Phys. Rev. B 62 (2000) 8391; M. Zacharias, J. Heitmann, R. Scholz, U. Kahler, M. Schmidt, J. Blasing, Appl. Phys. Lett. 80 (2002) 661. [5] L.S. Liao, X.M. Bao, Z.F. Yang, N.B. Min, Appl. Phys. Lett. 66 (1995) 2382. [6] M. Molinari, H. Rinnert, M. Vergnat, Appl. Phys. Lett. 77 (2000) 3499 ibid 79 (2001) 2172. [7] L.C. Chen, C.K. Chen, S.L. Wei, D.M. Bhusari, K.H. Chen, Y.F. Chen, Y.C. Jong, Y.S. Huang, Appl. Phys. Lett. 72 (1998) 2463. [8] H. Ueda, O. Kitakami, Y. Shimada, Y. Goto, M. Yamamoto, Jpn. J. Appl. Phys. 33 (1994) 6173. [9] V.M. Ng, M. Xu, S.Y. Huang, J.D. Long, S. Xu, in: Proceedings of the 7th APCPST and 17th SPSM (30P-17, June 29–July 2, 2004, Fukuoka, Japan) and Thin Solid Films 506–507 (2006) 283–287. [10] M. Chandru, Thesis for the Bachelor Degree of Science with Honours in Physics, Nanyang Technological University (NIE), 2003. [11] A.R. Wilkinson, R.G. Elliman, Appl. Phys. Lett. 83 (2003) 5512. [12] G.G. Qin, Mater. Res. Bull. 33 (1998) 1857; G.G. Qin, Y.J. Li, Phys. Rev. B 68 (2003) 085309.