Si bi-layer containing Ge QDs

Journal of Luminescence 154 (2014) 51–57

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Optical and structural investigations of self-assembled Ge/Si bi-layer containing Ge QDs Alireza Samavati a, Z. Othaman a,n, S.K. Ghoshal b, M.R. Dousti b a b

Ibn Sina Institute for Fundamental Science Studies, Universiti Teknologi Malaysia, Skudai 81310, Johor, Malaysia Advanced Optical Material Research Group, Department of Physics, Faculty of Science, Universiti Teknologi Malaysia, 81310 UTM Skudai, Johor, Malaysia

art ic l e i nf o

a b s t r a c t

Article history: Received 7 October 2013 Received in revised form 31 March 2014 Accepted 1 April 2014 Available online 12 April 2014

We report the influence of Si spacer thickness variation (10–40 nm) on structural and optical properties of Ge quantum dots (QDs) in Ge/Si(1 0 0) bi-layer grown by radio frequency magnetron sputtering. AFM images reveal the spacer dependent width, height, root mean square roughness and number density of QDs vary in the range of
12–25 nm,
2–6 nm,
1.95–1.05 nm and
0.55  1011–2.1  1011 cm  2, respectively. XRD patterns exhibit the presence of poly-oriented structures of Ge with preferred growth along (1 1 1) direction accompanied by a reduction in strain from 4.9% to 1.2% (estimated from Williamson–Hall plot) due to bi-layering. The room temperature luminescence displays strong blue– violet peak associated with a blue shift as much as 0.05 eV upon increasing the thickness of Si spacer. This shift is attributed to the quantum size effect, the material intermixing and the strain mediation. Raman spectra for both mono and bi-layer samples show intense Ge–Ge optical phonon mode that is shifted towards higher frequency. Furthermore, the first order features of Raman spectra affirm the occurrence of interfacial intermixing and phase formation during deposition. The excellent features of the results suggest that our systematic method may constitute a basis for the tunable growth of Ge QDs suitable in nanophotonics. & 2014 Elsevier B.V. All rights reserved.

Keywords: Germanium quantum dots Intermixing Photoluminescence Raman spectroscopy

1. Introduction Growth and optical characterization of low-dimensional semiconductor nanostructures with high-density and narrow size distribution is challenging for device fabrication [1,2]. The spontaneous formation of coherent three-dimensional (3D) Ge QDs on Si substrate due to lattice mismatch of semiconductor heteroepitaxy follows the Stranski–Krastanow (SK) growth mode. About 4.2% lattice mismatch between Ge layer and Si substrate leads to a linear increase in the film elastic energy that drives a layer-toisland transition above a critical thickness. However, the mechanism of bi-layer growth process is far from being understood. Multilayers containing Ge quantum dots (QDs) have been fabricated by several groups using rf magnetron sputtering [3], molecular beam epitaxy [4], ion implantation [5], dc sputtering in a reactive oxygen environment [6] and chemical vapor deposition with subsequent thermal annealing to induce film crystallization [7]. Among all these growth techniques, rf magnetron sputtering is usually one step method [8], comprising of thin film deposition

n

Corresponding author. Tel.: þ 60 7 553 4024; fax: þ60 7 556 6162. E-mail addresses: [email protected] (A. Samavati), [email protected] (Z. Othaman). http://dx.doi.org/10.1016/j.jlumin.2014.04.003 0022-2313/& 2014 Elsevier B.V. All rights reserved.

and post-growth thermal annealing treatment, which is compatible with the conventional integrated circuit fabrication processes. The strong room temperature luminescence from Ge QDs became promising for optoelectronics applications. Recently, the improvement in growth techniques for self-assembled Ge/Si QDs of high quality open up new avenues in nanophotonics [9]. The Ge/Si interface being a hetero-structure exhibits a staggered (type-II) band gap alignment. The momentum-conservation in Si-based materials with indirect nature of band gap requires phonon assisted radiative recombination processes. It is believed that the enhanced no-phonon luminescence observed in Ge/Si QDs is caused by the exciton localization at the Ge/Si interface [10,11]. Because the QDs are not fully strained, some strain is elastically transferred into the surrounding Si. This locally strained Si can confine the electrons in the vicinity of the Ge QDs by producing local potential minima in the conduction band and thereby provides a possible pathway to relax the requirement of momentumconservation [11,12]. The minimization of global elastic energy for stacked Ge/Si QDs requires the strain transfer from the buried Ge dots into the Si spacer in which the amount of spacer thickness is determinant [13,14]. Additionally, the strain relaxation is also achievable by material intermixing, which in turn, dramatically modifies the QD potential and the dot/spacer interfacial behavior. Therefore, the thickness of Si space layer significantly influences the structural

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and the optical properties of stacked Ge QDs. To our knowledge, this is the first time that the effect of Si spacer thickness on surface evolution, size, and spectral features of Ge/Si bi-layer is studied and compared to its mono-layer counterpart by means of atomic force microscopy (AFM), X-ray diffraction (XRD), photoluminescence (PL) and Raman spectroscopy.

2. Experimental procedure The radio frequency magnetron sputtering (13.56 MHz) method is used to grow Ge/Si mono and bi-layer structures consist of Ge layer thickness of 10 nm and Si space layer thickness of 10, 20, 30 and 40 nm at substrate temperature of 400 1C. This temperature is found to be optimum for growing high quality Ge QDs as reported elsewhere [15]. Polycrystalline Si and Ge targets (99.999% purity) are used for deposition. Silicon wafers (p-type) with (1 0 0) surface orientation is used as substrates. The schematic representation of the mono and bi-layered structures in addition to their band structure are depicted in Fig. 1. The reflectometry spectroscopy (Filmetrics) is used to measure the thickness. The thickness of layers is controlled by adjusting the rf sputtering power and the time. The nominal rate of deposition used for Ge and Si are 6 and 4 nm/min, respectively. The growth morphology of samples is characterized using AFM (SPI3800) built by Seiko Instrument Inc. (SII) and XRD (Bruker D8 Advance Diffractometer) using Cu-Kα1 radiations (1.540 Å) at 40 kV and 100 mA. The scanning range of 2θ is from 0 to 601. A slow speed of scanning
1.21/min with a resolution of 0.0111 is employed. Raman spectroscopy is performed using spectrum GX (NIR, FTRaman) system with an Nd crystal laser source having a spot size of 1 mm. The room temperature PL measurement (Perkin Elmer Ls 55 Luminescence Spectrometer) is carried out using a xenon flash lamp under different excitation wavelength.

estimated width distribution for the bi-layer structure is found to be narrower (smaller FWHM) than the mono-layer one. Similar trends are evidenced in the variation of QDs heights. The observed enhanced surface homogeneity for the bi-layer samples with increasing Si spacer is in agreement with Miura et al. [16]. This enhancement is attributed to the increased evenness of the strain distribution in the upper layer as predicted by the elastic continuum model [17]. However, the decrement in vertical correlation of QDs with the increase of Si spacer is decided by the attenuated strainfield coupling as follows. A rapid coarsening of the QDs in the subsequent layers with increasing spacer is responsible for the decrease in mean size of dots. The relaxation of these QDs is partly due to the dislocations on the dot edge which originates from the large accumulated strain in the spacer. The Williamson–Hall plot also confirms higher relaxation of strain in the bi-layered ([Ge/Si (10 nm)]  2) structure compare to the mono-layer one. The root mean square (rms) roughness is a measure of surface morphology and sample quality. AFM is used to measure the variation of rms roughness, number density and the average size of Ge QDs as summarized in Table 1. The rms roughness continuously decreased as the thickness of Si space layer is increased. The Si/Ge/Si sample in the absence of any Ge QDs on the surface shows smooth variation of rms roughness and is quite robust. The roughness fluctuation and the variation in width distribution follow the same trend. Moreover, the number density of monolayer sample is higher than the bi-layer one. The number density of bi-layer samples increases monotonically from 0.55  1011 cm  2 to 2.1  1011 cm  2 as the Si spacer thickness increase from 10 nm to 40 nm respectively. The XRD spectra of Ge/Si, Si/Ge/Si and [Ge/Si(10 nm)]  2 are presented in Fig. 5. Mono-layered structure display two prominent peaks attributed to preferred facets of Ge(1 1 1) and Ge(2 2 0) accompanied by a sharp peak associated with Si substrate [18,19]. The Ge peaks remain unaffected as Si is deposited on the top of Ge QDs. However, the absence of crystalline peak from Si substrate in

3. Results and discussion The energy dispersive X-ray (EDX) spectra of the bi-layer sample [Ge/Si(10 nm)]  2 as shown in Fig. 2 clearly reveal the presence of Ge, Si, O, C and Au elements. The appearance of Ge peaks exhibits the Ge QDs, which is in consistent with the AFM results. The occurrence of carbon (C) signal is due to the supportive carbon tape attached to the sample holder, while the oxygen (O) peak originates from the surface passivation of the dangling bonds and subsequent formation of SiO2 and GeOx defects. The Au peak appears from the Au coating on the sample used also for field emission scanning electron microscopy which is not shown here. Fig. 3 shows the 3D AFM images for mono-layered Ge/Si (Fig. 3a), sandwiched layer Si/Ge/Si (Fig. 3b), and bi-layered structures with different Si spacer thickness such as [Ge/Si(10 nm)]  2 (Fig. 3c), [Ge/ Si(20 nm)]  2 (Fig. 3d), [Ge/Si(30 nm)]  2 (Fig. 3e) and [Ge/Si (40 nm)]  2 (Fig. 3f). The corresponding width distributions are displayed in Fig. 4. Following SK growth mode, during the sputtering, the Ge add atoms diffuse at surface and produce new nucleation centers to form dots due to entropy maximization. The

Fig. 2. The EDX spectra of [Ge/Si(10 nm)]  2.

Fig. 1. Schematics of the sample structure displaying structural features of mono-layer (a), bi-layer (b) and energy band diagram (c).

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Fig. 3. The AFM micrograph of Ge/Si ((a) and (g)), Si/Ge/Si ((b) and (h)), [Ge/Si]  2 having 10 nm ((c) and (i)), 20 nm ((d) and (j)), 30 nm ((e) and (k)) and 40 nm ((f) and (l)) Si space layer thickness.

the sandwiched sample may be due to the strong scattering from Si deposited layer. Moreover, bi-layering through continuous Ge deposition re-structured the entire XRD pattern by forming polyoriented structures of the Ge at (0 0 4) and (3 1 1) planes accompanied by (1 1 1) and (2 2 0) planes. The appearance of sharp Ge (0 0 4) peak with strong intensity may be due to the formation of bulk Ge structure in the sub-layer Ge caused by heat treatment during deposition. The intensity of preferably orientated Ge(1 1 1) gradually enhances while the signal of Ge(2 2 0) diminishes. The presence of GeO2 and SiO2 peaks are attributed to the surface passivation of oxygen arises due to atmospheric exposure during sample transfer or deposition. Finally, the oxygen may diffuse and oxide the sub-layers to form GeO2 and SiO2. The strain field between Ge nanocrystallites and the surrounding matrix alters the lattice parameters by changing the interspacing of the planes. The broadening of the XRD peaks is majorly attributed to the presence of various crystallite sizes, microstrain and instrumental effects. The broadening caused by the effect of strain can be determined by Williamson–Hall plot (βcosθ versus sinθ) which is depicted in Fig. 6. Here, β is the line broadening at the half of the maximum intensity (FWHM). A linear relation between βcosθ and sinθ suggests the contribution of strain in the broadening of the XRD peak. On the other hand, if the contribution in broadening is appeared from the size variation only then βcosθ would have remained constant [20]. Consequently, the observed shift in the XRD spectra for Ge(1 1 1) peak corresponding to the bilayer sample is due to the lowering of the strain [21] as shown in Williamson–Hall plot.

The Raman spectra of mono-layer and bi-layer samples are shown in Fig. 7 in which the peak of Ge–Ge optical phonon for bilayer sample is marked as circle in the inset. The vertical solid line at 300 cm  1 indicates the optical phonon position corresponding to the bulk crystalline Ge. The peaks of Ge–Ge optical phonon related to the mono and bi-layer structures are slightly shifted towards the higher frequencies (marked by dotted line) with respect to their bulk value. This shift is ascribed to the competitive effect of phonon confinement and strain. The optical phonon branches of the bulk Ge are known to be quadratic and nearly flat at the Brillouin zone center (k
0) as described by simple linear chain model. Restricted phonons in a nanocrystal are equal to those vibrations in an infinite crystal whose wave vector is given by mπ/d with m an integer and d the diameter of the nanocrystal [22]. Therefore, the spatial limitations of QDs result a shift of optical phonons towards lower frequency. In the case of dots with very small heights
1.5–2 nm, the effect of phonon confinement that results a shift
2 cm  1 is observed [23]. It is the effect of strain, which plays a predominant role in shifting the peak towards the higher frequency. Interestingly, the lattice mismatch of Si and Ge leads to a compressive strain on the dots in the lateral directions that induces a Ge–Ge mode shift towards the higher frequency. The Ge–Ge optical phonon frequency induced by a biaxial strain in Ge can be written as ω ¼ ω0 þ ð1=2ω0 Þ ½pεzz þ qðεxx þ εyy Þ where ω0 ¼ 0.564  1014 s  1 is the frequency of the Ge zone center LO phonon; p¼  4.7  1027 s  2 and q¼  6.16  1027 s  2 are the Ge deformation potentials; &xx ¼  0.042 and &yy are the biaxial strain. Here, &zz ¼  (2C12)xx/C11, &xx ¼ &yy, C11 ¼1288 kbar,

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750

3000

Average Height~ 2nm

Counts

2000

Counts

Counts

500

Average Height~ 6nm

250

500 Average Height~ 6nm

250

1000

0

0 5

10

15

20

25

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Normalized width (nm)

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5

10

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Normalized width (nm)

2250 1500

Counts

Counts

1500 1000 Average Height~ 5nm

Average Height~ 3nm

750

500

FWHM

0

0 5

10

15

20

25

5

30

10

15

20

25

30

Normalized width (nm)

Normalized width (nm)

Fig. 4. The size distribution of Ge/Si (a), [Ge/Si]  2 having 10 nm (b), 20 nm (c), 30 nm (d) and 40 nm (e) Si space layer thickness.

Table 1 The rms roughness, number density and average size of Ge QDs for different samples. Rms roughness (nm)

Ge/Si Si/Ge/Si [Ge/Si(10 nm)]  2 [Ge/Si(20 nm)]  2 [Ge/Si(30 nm)]  2 [Ge/Si(40 nm)]  2

2.15 0.42 1.95 1.77 1.21 1.05

Number density (  1011 cm  2)

3 – 0.55 0.70 1.35 2.10

Average size (nm) Width

Height

8 – 25 25 20 16

2 6 6 5 3

C12 ¼482.5 kbar are elastic coefficients [24]. The calculated value of ω is 317.4 cm  1 for fully strained pure Ge on Si and the corresponding experimental value is 302 cm  1 (Fig. 7). This difference in the theoretical and experimental values suggests that the dots in the samples may not be completely strained as assumed in the theory, which is indicated by the Williamson–Hall plot. The strain relaxation of QDs may arise due to the microscopic mechanisms such as the presence of threading dislocations and the Ge–Si atomic intermixing which results in a weak shoulder at higher energy. The signature of the intermixing is evidenced as Ge–Si peak in the full-scale Raman spectra of [Ge/Si(10 nm)]  2 depicted in the inset of Fig. 7. The intensity of peaks depends on the relative number of corresponding bonds, therefore, the degree of the interface intermixing can be determined by the integrated peak intensity ratio IGe–Ge/IGe–Si. The relationship for SiGe alloy and Ge QDs is presented by Mooney et al. [25] as I Ge  Ge x ffi 3:2 2ð1  xÞ I Ge  Si

where x is the average Ge concentration which is estimated to be 0.46. Furthermore, the intermixing of Ge and Si is greatly enhanced by the heat treatment during deposition. This in turn, offers alternative pathways to the strain relaxation that results the reduction of mismatch energy at the phase boundary as shown in Fig. 6. The formation of QDs during the deposition of subsequent Ge layer is strongly influenced by (1) the compressive strain induced by the buried Ge wetting layer and (2) the inhomogeneous strain field due to the partially strain relaxed dots. The above mechanisms cause improved size-uniformity and lateral ordering of the nanodots in the topmost layer even when the first layer is randomly nucleated [26,27]. Fig. 8 represent the normalized PL spectra of prepared samples under
5.18 eV excitation energy. No significant PL signal is observed from Si/Ge/Si sandwiched layer. For mono-layer Ge/Si QDs, a PL band centered at 3.28 eV is observed in conformity with Jie et al. and Mestanza et al. [28,29]. The origin of this peak can be explained by the Ge-QDs/SiO2 or GeOx interface model depicted in Fig. 9 in which energy band values are adapted from Cohen and Chelikowsky [30]. If the confinement within the potential well is sufficiently deep, confined electrons and holes in the Ge dots favors radiative recombination. The radiative recombination requires the energy
3.28 eV. However, to confine electrons with energies
2.5 eV above the conduction band edge the potential barrier at the Ge-dot/Si interface is not sufficient for such recombination. Conversely, such electrons with higher energies can only exist in the structural defects states of SiO2 or GeOx. This fact suggests that the radiative recombination does not involve electrons from the conduction band of the Ge or Si, rather it involves localized electrons in radiative defect states at the Ge-dot/SiO2 or GeOx interfaces. Min et al. [31] and Takeoka et al. [32] proposed a similar mechanism for radiative recombination on the ion implanted Ge in SiO2 and the rf co-sputtered Ge and SiO2 nanostructures, respectively.

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Ge (220)

55

Ge (220)

12

10

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14

16 18 20 2θ (degree)

40

22

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24

60

Intensity (a.u.)

Intensity (a.u.)

Intensity (a.u.)

Ge (111)

Si substrate

Intensity (a.u.)

Ge (111)

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Si substrate

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Si substrate

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GeO2 (311)

40

30

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Ge (311) GeO2 (012)

GeO2 (011)

GeO2 (am)

SiO2 (am)

Ge (220)

Intensity (a.u.)

Ge (004)

Ge (111)

2θ (degree)

12

60

12

14

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24

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70

2θ (degree) Fig. 5. XRD spectra from Ge/Si (a), Si/Ge/Si (b) and [Ge/Si(10 nm)]  2 (c).

0.4

ε= 4.9% ε= 4.1% ε=1.2% Ge/Si Si/Ge/Si [Ge/Si(10 nm)]×2

0.1 0.0

Ge-Ge

Si-Si

Ge-Si

0.2

0.9

1.2

sinθθ/λ //λ λ Fig. 6. Williamson–Hall plot of Ge/Si, Si/Ge/Si and [Ge/Si(10 nm)]  2.

The variation of PL intensity and peak energy as a function of Si spacer thickness is depicted in Fig. 10. The PL intensity shows continuous decrement going from mono to bi-layer Ge/Si and the same trend persists as the thickness of Si space layer is increased. This is related to the decrease in the concentration of photocarriers, their aggregations, and the energy transfer to the down layer in particular. The observed red shift
0.08 eV in the PL peak for bi-layer sample (Ge/Si(10 nm)  2) compare to mono-layer one is related to the increase of average width and height of QDs. There is no significant change in the band position of bi-layer samples when the Si spacer thickness increases from 10 nm to 20 nm since the diameter of dots varies negligibly. The blue shift
0.05 eV is observed due to the continual increase of the Si spacer thickness from 20 nm to 40 nm. Such

Intensity (a.u.)

β cosθ /λ

0.3

Ge-Ge

Ge-O

302 cm-1 300

400

500

600

[Ge/Si(10 nm)]××22 Ge/Si

Bulk Ge 280

300

320

Raman shift (cm-1) Fig. 7. Raman spectra of Ge/Si and [Ge/Si(10 nm)]  2 for the Ge–Ge peak. The inset shows the full-scale spectra of the [Ge/Si(10 nm)]  2.

shift is related to the decrease of the width and height of QDs as well as the modifications in strain and intermixing. Moreover, the degree of intermixing depends on the accumulated strain energy.

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Excitation (Ε) = 5.18 eV

Intensity (a.u.)

Red shift 0.08 eV 3.28 eV

Ge\Si

Intensity (a.u.)

3.20 eV

E x 5.30eV

3.23 eV

2

3

4

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6

Energy (eV)

[Ge\Si(10 nm)]××22 [Ge\Si(20 nm)]××22 [Ge\Si(30 nm)]××2 ×2 2 [Ge\Si(40 nm)]××2 ×2 2

Si\Ge\Si Blue shift 0.05 eV

2.8

3.2

3.6

4.0

E x 5.18eV

Intensity (a.u.)

3.25 eV

E x 5.05eV

Energy (eV) Fig. 8. Room temperature PL spectra.

Stokes Shift 2-3.5 eV Ec (Si)

E x 4.86 eV

electron

1.17 eV

0.66 eV hole

0.86 eV

Ev (Si)

2

Growth direction

SiO2 or GeOx

Ge QDs

Defect Fig. 9. Schematic band structure of Ge-dot/SiO2 or GeOx. Dotted lines indicate the confined state energy levels in the Ge QDs. The energy of 1.17 eV and 0.66 eV correspond to the band gap of bulk Si and Ge respectively.

3.30

12 11

×10

3.28 3.26

Εex=5.18 eV

8

3.24 7 6

Intensity energy

3.22

Peak energy (eV)

PL intensity (a.u.)

Ge\Si monolayer

10 9

5 3.20

4 10

0.51 eV

20

30

40

Si space layer thickness (nm)

3 4 Photon energy (eV)

5

Fig. 11. Photoluminescence spectra of [Ge/Si(10 nm]  2 under different excitation energies (Ex) at room temperature: 5.30 eV (a), 5.18 eV (b), 5.05 eV (c) and 4.86 eV (d). The peak energies are estimated by using Gaussian functions and the dotted lines are Gaussian profiles. The photon energy of the excitation laser is shown by arrows.

dots. The strain transfer mechanism tends to lower the emission energy of QDs. On the other hand, the deformation potential of the Si band (which is about three times higher than the Ge band [13]) favors the material intermixing that tend to increase the QD emission energy due to the increased Si content in the QDs. Following Kanemitsu et al. [33] and Okamoto et al. [34] the resonant-excitation energy dependence of the PL (Fig. 11) is carried out on [Ge/Si(10 nm)]  2 sample to demonstrate the origin of PL bands due to the defect states of Ge-dots/SiO2 or GeOx. When the excitation photon energy is obove 5.18 eV, the PL peak energy nearly remains constant around 3.20 eV. Generally, if the luminescence mechanism is dominated by the excitonic recombination related to the quantized states in Ge nanocrystalite, the Stokes shift under 4.86 eV excitation should be close to excitons binding energy. In contrast, the observed Stoke shift
0.51 eV seems to be too large compare to the excitons binding energy of Ge nanocrystals. Thus, the origin of luminescence is mainly associated with the radiative recombination through localized states at or in the proximity of the Ge nanocrystallite-SiO2\GeOx interfacial defect states.

Fig. 10. PL intensity (left) and band gap energy (right) versus Si space layer thickness. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

4. Conclusion

Reduction of the spacer thickness accumulates more strain energy that cause stronger Ge–Si intermixing. Although, reducing the spacer thickness may provide better confinement for the electron in the Si conduction band due to the elastic strain transfer effect, the predominant role played by material intermixing tends to counteract this effect and degrades the interface sharpness of the

We investigate the mechanism of bi-layer growth processes prerequisite for determining multi-layered structure and its impact on surface and optical properties of Ge QDs at room temperature. Modification of the optical response of Ge QDs by controlling the Si spacer thickness, improving the uniformity of dots through bilayering and tuning their size distribution are reported. The width

A. Samavati et al. / Journal of Luminescence 154 (2014) 51–57

and height of QDs is substantially reduced with the increase of Si spacer thickness. The process of bi-layering and the thickness variation of Si spacer strongly influence the rms roughness and the number density. The bi-layering favors the growth of polyoriented structures of Ge with preferred orientation along (1 1 1) direction. The reduction of strain value due to the conversion of mono into bi-layered is quantified using Williamson–Hall plot. The occurrence of asymmetric PL band centered at 3.20 eV for bi-layer sample is attributed to the radiative recombination that involves electrons localized in the defect states at the Ge-dot/ SiO2 or GeOx interfaces. Subsequently, the shift in the intense PL peak is achieved by changing the thickness of the Si spacer. Intense Ge–Ge optical phonon mode accompanied by a shift is evidenced in the Raman spectra. These blue shifts in luminescence and Raman band are ascribed to the effects of quantum confinement, strain and interfacial intermixing. We demonstrate a tunable bi-layering of heterostructure via careful manipulation of the Si spacer which may significantly improve the growth kinetics of Ge QDs decisive for widespread applications. Acknowledgments Universiti Teknologi Malaysia through Vote Q.J130000.2526.02H94 and Ibnu Sina Institute for Fundamental Science Study. Moreover, Dr. Samavati is thankful to RMC for postdoctoral research grants. References [1] A.A. Shklyaev, M. Ichikawa, Surf. Sci. 514 (2002) 19. [2] S.K. Ray, K. Das, Opt. Mater. 27 (2005) 948. [3] S. Huang, Z. Xia, H. Xiao, J. Zheng, Y. Xie, G. Xie, Surf. Coat. Technol. 204 (5) (2009) 558. [4] S.V. Kondratenko, O.V. Vakulenko, Y.N. Kozyrev, M.Y. Rubezhanska, A. G. Naumovets, A.S. Nikolenko, V.S. Lysenko, V.V. Strelchuk, C. Teichert, J. Mater. Sci. 46 (17) (2011) 5737. [5] W. Skorupa, L. Rebohle, T. Gebel, Appl. Phys. A 76 (7) (2003) 1049. [6] M. Zschintzsch, N.M. Jeutter, J. Von Borany, M. Krause, A. Mücklich, J. Appl. Phys. 107 (3) (2010) 034306. [7] D.C. Paine, C. Caragianis, T.Y. Kim, Y. Shigesato, T. Ishahara, Appl. Phys. Lett. 62 (22) (1993) 2842.

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