Structural and optical characteristics of Ge<sub>1−x</sub>Snx/Ge superlattices grown on Ge-buffered Si(001) wafers

Jia-Zhi Chen; H. Li; H. H. Cheng; Guo-En Chang

doi:10.1364/OME.4.001178

1. Introduction

Group-IV elements of silicon (Si) and germanium (Ge) are the building elements for electronic devices. In moving to photonic devices, owing to the indirectness of its energy band, the performance is less pronounced. In a recent development, tin (Sn)-another group IV element-was employed in the growth of IV-IV compounds. The incorporation of Sn into Ge modulates the energy band of the host materials. To a certain Sn composition and above, the energy band of the host materials transforms from indirect to direct bandgap which has been demonstrated by several different types of measurements [1–5]. From these investigations, the critical Sn composition is found to be around 6% to 10%. This leads to the development of photonic devices made of thick GeSn films, such as infrared light-emitting diodes and photodetectors [5–8]. In light of these progresses, laser designs with active layers employing GeSn quantum wells (QWs) have been proposed and analyzed [9, 10]. Compared to the bulk counterpart, GeSn QWs provide more efficient carrier utilization due to the quantum confinement effect, showing promises to achieve lasing action at a reasonable threshold carrier density. Thus Sn-based low-dimensional quantum structures are highly desired for the development of new Si-based photonic devices.

Previously, few attempts have been made to grow Ge_1−xSn_x/Ge superlattices (SLs) or quantum wells [5, 11, 12]. However, no clear quantum confinement was obtained. Here, we demonstrate the growth and characterization of low-dimensional Ge_1−xSn_x/Ge SLs grown on Si wafers via a high-quality Ge virtual substrate (VS). Ge_1−xSn_x/Ge SLs with Sn compositions up to the threshold value that affords a direct bandgap are achieved by a special low temperature growth technique using molecular beam epitaxy (MBE). Quantum confinement is found as revealed by room-temperature absorption measurements, in a good agreement with our k.p analysis. Those results show that Ge_1−xSn_x/Ge low-dimensional structures can be achieved, opening new opportunities for Si photonics.

2. Experimental details

The samples were grown on Si (001) wafers with a resistivity of 1–10 Ω · cm by solid-source MBE at a base pressure less than 2 × 10⁻¹⁰ torr. Deposition was performed by using an e-gun to evaporate the target source (Si and Ge). Sn is grown using an effusion cell equipped with dual heating filaments placed at the bottom and top of the crucible to prevent Sn condensation at the orifice area of the crucible. Prior to being loaded into the chamber, Si substrates were degreased, dipped into HF solution, and rinsed in de-ionized water. To remove the oxide layer at the surface, in-situ surface preparation was carried out using thermal oxide adsorption at 900 °C, assisted by a low Si flux. The process continued for about 4–5 minutes until a sharp 2 × 1 reflection high energy electron diffraction (RHEED) pattern was observed. During the crystal growth, in-situ RHEED was employed to monitor the growth. The sample consists of: (a) a Ge buffer layer of 100 nm grown at low temperature of 350 °C (LT) followed by in-situ annealing at 800 °C for 5 minutes, (b) a Ge layer of 100 nm grown at high temperature of 550 °C (HT), (c) a thick Ge buffer layer of 100 nm grown at low temperature of 350 °C, (d) 20-period Ge_1−xSn_x/Ge SL with a nominal layer thickness of 6.5/2.5 nm grown at 150 °C, and (e) a Si cap layer of 3 nm grown at 150 °C. The 200 nm Ge layer (a), (b) serves as a high-quality VS for the subsequent growth of the Ge_1−xSn_x/Ge SL. This gives a smaller lattice mismatch than that between GeSn and Si to improve the material quality. Layer (c) acts as a buffer layer which can trap the dislocations caused by the lattice misfits between the Ge VS and the subsequently grown GeSn layers [13]. Five samples were grown with different Sn compositions ranging from 1.6% to 6.96% determined by x-ray measurements, which will be discussed later. A reference sample of Ge VS with the same design but without the Ge_1−xSn_x/Ge SL was also grown for the absorption experiment. Between the growth of different layers, the growth was interrupted by closing the shutter in front of the sources for changing the growth temperature. The GeSn layers were grown at a temperature below the melting point of Sn (230 °C) to suppress the Sn segregation at the surface [14].

Different experimental techniques were employed to characterize the samples. Cross-sectional transmission electron microscope (XTEM) experiments were conducted to probe microstructure of the samples. High-resolution X-ray diffraction in both (004) ω-2θ scanning and (224) reciprocal space mapping (HRXRD-RSM) was used to measure the lattice constants of the layers in the film [8, 15]. ((004) measurement gives the lattice constant in the growth direction (a^⊥) while (224) measurement yields the lattice constant in the plane perpendicular to growth direction (a^‖).) These data were acquired with a Bede D1 system. Fourier transform infrared spectroscopy (FTIR) was employed to measure the absorbance spectra of the samples for determining their optical bandgap. Data are acquired using a liquid-nitrogen cooled MCT photodetector operated at wavelength range of 1.43 μm to 22 μm. The absorption signal was enhanced using a total internal reflection technique. (The sample is polished in a parallelogram shape. Incident light is guided to one side of the polished surface and exits from the other side of the sample to the detector after multiple reflection through the epilayer.)

3. Results and discussion

We first present the microstructure of the film. A typical cross-sectional image (sample N2) is shown in Fig. 1. Two characteristics are observed. Firstly, clear misfit dislocations are observed at the Ge/Si interface attributed to the 4.2% lattice mismatch between the Ge and Si layers. (The bulk lattice constants of Si, Ge, and α-Sn are $a_{0}^{Si} = 5.4307 \overset{˚}{A}$ , $a_{0}^{Ge} = 5.6575 \overset{˚}{A}$ , and $a_{0}^{Sn} = 6.4892 \overset{˚}{A}$ , respectively [9,16].) Those misfit dislocations indicate that the Ge buffer layer is strain-relaxed. Secondly, for the Ge_1−xSn_x/Ge heterostructure, the interface is flat with a good periodicity. Dislocation lines are not observed at the Ge_1−xSn_x/Ge interface indicating the layers are coherently strained and the strain values will be discussed later. From the TEM image, the thickness of the Ge_1−xSn_x/Ge layer is determined and summarized in Table 1.

Fig. 1 XTEM image of sample N2, showing smooth and flat interfaces between the Ge and GeSn layers, as well as the good periodic structure. The inset shows the Ge/Si interface of the sample where a high density of misfit dislocations is visible, indicating the Ge VS is strain-relaxed.

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Table 1. Summary of the layer thicknesses, out-of-plane lattice constant, Sn composition, and direct bandgap energy of the strained-layer Ge_1−xSn_x/Ge SLs. The direct bandgap energies extracted from the absorbance measurements match the lowest direct transition energies obtained from the PL measurements reported in Ref. [18].

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Next, we move to the determination of Sn composition and strain in the layers. From the HRXRD-RSM experiments [8, 15], the in-plane and out-of-plane lattice constants of the Ge and GeSn layers are deduced. The results are summarized in Table 1. For the Ge layers, a^‖ is 5.65 Å which agrees well with the bulk Ge lattice constant reflecting that the Ge layers are fully strain-relaxed. For the Ge_1−xSn_x layers, the same value of a^‖ is obtained, showing that the Ge_1−xSn_x layers are coherently strained. (For the other samples, the same result is obtained.)

Knowing a^‖ and a^⊥, the unstrained lattice constant of the Ge_1−xSn_x layers can be deduced by

a_{0} = \frac{a^{⊥} + 2 \frac{C_{12}}{C_{11}} a^{‖}}{1 + 2 \frac{C_{12}}{C_{11}}},

where C₁₁ and C₁₂ are the elastic constants. For Ge_1−xSn_x alloys, their ratio is C₁₂/C₁₁ = 0.3738 + 0.1676x − 0.0296x² [16]. By comparing the deduced unstrained lattice constant from the x-ray experiments to the composition-dependent lattice constant of unstrained Ge_1−xSn_x alloys, the Sn compositions of the Ge_1−xSn_x layers are determined and the results are listed in Table I. (In making the comparison, the composition-dependent lattice constant of unstrained Ge_1−xSn_x alloys is

a_{0}^{GeSn} = a_{0}^{Ge} (1 - x) + a_{0}^{Sn} x + θ^{GeSn} x (1 - x)

where θ^GeSn = − 0.066 Å is the bowing parameter [16].) As described above, since the Ge_1−xSn_x layers are in-plane lattice-matched to the Ge VS, they are subjected to an in-plane biaxially compressive stress. The corresponding in-plane strain is

ε_{x x} = ε_{y y} = (a_{0}^{Ge} - a_{0}^{GeSn}) / a_{0}^{GeSn}

. The strain in the growth direction can be evaluated by ε_zz = −2(C₁₂/C₁₁)ε_xx. Because the Ge layers are strain-relaxed, the SL stack overall is asymmetrically strained. As more Ge_1−xSn_x/Ge periods are added to this stack, the build-up of net strain across the stack will impose an upper limit upon the total number of layers that can be used in the SL.

The measured absorbance spectra of the samples together with the Ge VS reference are plotted in Fig. 2(a). For the Ge VS, the absorbance increases rapidly at the onset which represents its direct bandgap. Near the absorbance edge, the direct bandgap energy ( $E_{g}^{Γ}$ ) can be determined using the Tauc equation [17]

α h ν = A {(h ν - E_{g}^{Γ})}^{1 / 2},

where α is the absorption coefficient, A is an energy-independent constant, h is the Planck constant, and ν is the frequency of light. A plot of (αhν)² versus hν is shown in Fig. 2(b). By numerically fitting the band-edge absorbance curve using Eq. (2), we obtain an energy of 0.789 eV which is in good agreement with the direct bandgap of 0.79 eV for bulk Ge [19].

Fig. 2 (a) Room-temperature absorbance spectra of the Ge_1−xSn_x/Ge SL samples, compared to that of the Ge VS. (b) Tauc plot for the determination of direct bandgap. The dashed lines show fits using the Tauc equation.

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With the incorporation of Sn, the absorbance edge shifts to lower photon energy with increasing Sn composition. Using the same fitting procedure, we obtain the direct bandgap energies for the Ge_1−xSn_x/Ge SLs, which are consistent with lowest transition energies extracted from the photoluminescence (PL) experiments [18]. The results are summarized in Table I. The direct bandgap energy corresponds to the transition energy between the confined states of the Γ-valence and conduction bands which we will discussed later. Furthermore, in contrast to the smooth absorption curve of the Ge VS, it is noted that for sample N6 the absorbance spectrum exhibits clear step-like structures as indicated by the arrows. Those structures are distinguished features of two-dimensional quantum structures with step-like density-of-states.

Having determined the Sn composition and direct bandgap, we move to the composition-dependent bandgap of the Ge_1−xSn_x/Ge SLs. To quantitatively interpret the transition energies of the Ge_1−xSn_x/Ge SLs, the effects of strain and quantum confinement are analyzed. We start with the unstrained bandgap of Ge_1−xSn_x alloys

E_{g} ({Ge}_{1 - x} {Sn}_{x}) = (1 - x) E_{g} (Ge) + x E_{g} (Sn) - b x (1 - x),

where E_g(Ge) and E_g(Sn) denote the bulk bandgap energies of Ge and α-Sn, respectively, and b is the bowing parameter. For the Γ- and L-valley bandgaps, the parameters used in the calculation are listed in Table 2. The unstrained direct bandgap of Ge_1−xSn_x alloys as a function of Sn composition is depicted as the solid line in Fig. 3(a), compared with the experimental data of the Ge_1−xSn_x/Ge SLs determined from the absorption experiments. Obviously, the measured direct bandgap energies are larger than those of unstrained bulk Ge_1−xSn_x alloys. This is mainly attributed to the compressive strain and the quantum confinement effect that shift the bandgap energy. To distinguish their individual effects, we first calculate the strain-induced bandgap energy shifts. The strain-induced energy shifts for the conduction and valence bands are calculated using the deformation potential theory [20]

Δ E_{c}^{Γ} = a_{c} (ε_{x x} + ε_{y y} + ε_{z z})

Δ E_{c}^{Γ} = a_{L} (ε_{x x} + ε_{y y} + ε_{z z})

Δ E_{HH} = - P_{ε} - Q_{ε}

Δ E_{LH} = - P_{ε} + Q_{ε} + \frac{1}{2} [\sqrt{{(Δ + Q_{ε})}^{2} + 8 Q_{ε}^{2}} - (Δ + Q_{ε})]

P_{ε} = - a_{v} (ε_{x x} + ε_{y y} + ε_{z z})

Q_{ε} = - \frac{b_{v}}{2} (ε_{x x} + ε_{y y} - 2 ε_{z z})

where a_c (a_v) is the hydrostatic deformation potential of the Γ-conduction (valence) band, a_L is the deformation potential of the L-conduction band, b_v is the shear deformation potential, and Δ is the spin-orbit split-off energy. The various bands are lined up using the model-solid theory [9]. By adding the strain-induced energy shifts to the unstrained bandgap (Eq. (3)) and taking parameters from Ref. [9], we obtain the band edges of the various bands for strained Ge_1−xSn_x alloys as a function of Sn composition. For pseudomorphic Ge_1−xSn_x alloys on Ge, the energy positions of the conduction and valence band edges as a function of Sn composition (in units of eV) can be expressed as

E_{c}^{Γ} (x) = 0.787 - 1.46 x + 2.31 x^{2}

,

E_{c}^{L} (x) = 0.664 - 0.91 x + 1.01 x^{2}

, E_HH(x) = 1.7x − 0.26x², and E_LH(x) = 0.15x + 3.76x². The results are depicted in Fig. 3(b). For the conduction bands, as the Sn composition increases, both the Γ- and L-conduction bands shift downward. Most importantly, their energy difference decreases as the Sn content increases, leading to an important indirect-to-direct bandgap transition at a Sn composition of 10.9%. This crossover suggests that pseudomorphic GeSn alloys on Ge can become a direct-bandgap material for efficient light emitters. For the valence band, the heave-hole (HH) band and light-hole (LH) band are lifted with increasing Sn composition. In particular, the compressive strain pushes the HH band above the LH band, making the top valence band HH-like. As a result, the direct transition from the HH band to the Γ-conduction band is the lowest direct transition which defines the strained direct bandgap

E_{g}^{s} (x) = E_{c}^{Γ} (x) - E_{HH} (x)

. The calculated results are depicted as the dashed line in Fig. 3(a), compared with the experimental data from Ref. [5,20]. A good agreement is found between our calculations and the experimental results, showing the validity of our calculations. However,

E_{g}^{s} (x)

still lies below the observed values of direct bandgap energy for the Ge_1−xSn_x/Ge SLs, and the energy difference becomes larger with increasing Sn composition. This discrepancy may derive from the quantum confinement effect in the Ge_1−xSn_x/Ge SL that increases the transition energy.

Table 2. Γ- and L-valley bandgap parameters for Ge_1−xSn_x alloys (Refs. [4, 9]).

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Fig. 3 (a) Direct bandgap energies of unstrained Ge_1−xSn_x alloys, pseudomorphic Ge_1−xSn_x alloys on Ge, and pseudomorphic Ge_1−xSn_x/Ge SLs on Ge. (b) Calculated band edges of the various bands for pseudomorphic Ge_1−xSn_x alloys on Ge as a function of Sn composition. The energy zero is chosen at the valence band maximum of unstrained Ge. The indirect-to-direct bandgap transition occurs at x=10.9%.

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When bringing relaxed Ge and fully-strained Ge_1−xSn_x on Ge (VS) to form a heterostructure, both conduction and valence bands exhibit band offsets. Since the strained bandgap of the Ge_1−xSn_x layer is smaller than that of the Ge layer, the Ge_1−xSn_x/Ge heterostructure has a type-I band alignment where the Ge_1−xSn_x layer serves as the quantum well. From Fig. 3(b), the band offsets become larger with increasing Sn composition. Once the well width is sufficiently thin, quantum confinement effects become apparent and can significantly increase the bandgap energy (blueshift). To confirm the quantum confinement in the Ge_1−xSn_x/Ge SLs, the quantized energies and subband structures are calculated using the multi-band k.p method [21, 22] with the parameters taken from Ref. [9]. The calculated bandgap energies for the Ge_1−xSn_x/Ge SLs, that is, the transition energy from the first HH subband to the first Γ-conduction subband (HH1-cΓ1), are depicted in Fig. 3(a). A good agreement between our theoretical calculations and experimental data is obtained. This suggests that the blueshift of direct bandgap can be quantitatively explained by the quantum confinement effects at the Γ point. (In performing the modeling, it is noted that different sets of input parameters [23,24] used in the calculation may cause a change in the results. However, considering the Sn composition is not large in this study, there is no significant variation for the calculated bandgap energies due to the variation of input parameters.) Based on the results of blue-shifted transition energies and observed step-like features in the absorbance spectra of the Ge_1−xSn_x/Ge SLs, we thus confirm the room-temperature quantum confinement effect in the Ge_1−xSn_x/Ge SLs.

4. Conclusion

In summary, low dimensional heterostructure of newly developed group-IV material system is investigated. Ge_1−xSn_x/Ge superlattices are demonstrated by a low temperature growth technique using molecular beam epitaxy. The structural and optical properties are investigated by different experimental techniques. As the Sn content is increased, the absorbance edge significantly redshifts and step-like features are observed at room temperature. The band structure is analyzed using a multi-band k.p method to identify the bandgap energy shifts by strain and quantum confinement effects of the superlattices. The analysis shows that the HH1-cΓ1 transition is the lowest transition, and the calculated transition energy agrees with the experimental results of direct bandgap determined from the absorption experiments, confirming the quantum confinement in the Ge_1−xSn_x/Ge superlattices. Those results show that Ge_1−xSn_x/Ge heterostructures represent a new type of group-IV quantum structure for the development of high-performance Si-based photonic devices.

Acknowledgments

This work was financially supported by the Ministry of Science and Technology of Taiwan under Grant Nos. NSC 102-2221-E-194-053-MY3 and NSC 101-2112-M-002-015-MY3.

References and links

1. G. He and H. A. Atwater, “Interband transitions in Sn_1−xGe_x alloys,” Phys. Rev. Lett. 79, 1937–1940 (1997). [CrossRef]

2. V. R. D’Costa, C. S. Cook, A. G. Birdwell, C. L. Littler, M. Canonico, S. Zollner, J. Kouvetakis, and J. Menéndez, “Optical critical points of thin-film Ge_1−ySn_y alloys: A comparative Ge_1−ySn_y/Ge_1−xSi_x study,” Phys. Rev. B 73, 125207 (2006). [CrossRef]

3. R. Roucka, J. Mathews, R. T. Beeler, J. Tolle, J. Kouvetakis, and J. Menéndez, “Direct gap electroluminescence from Si/Ge_1−ySn_y p-i-n heterostructure diodes,” Appl. Phys. Lett. 98, 061109 (2011). [CrossRef]

4. H. Lin, R. Chen, W. Lu, Y. Huo, T. I. Kamins, and J. S. Harris, “Investigation of the direct band gaps in Ge_1−xSn_x alloys with strain control by photoreflectance spectroscopy,” Appl. Phys. Lett. 100, 102109 (2012). [CrossRef]

5. A. Gassenq, F. Gencarelli, J. V. Campenhout, Y. Shimura, R. Loo, G. Narcy, B. Vincent, and G. Roelkens, “GeSn/Ge heterostructure short-wave infrared photodetectors on silicon,” Opt. Express 20, 27297–27303 (2012). [CrossRef] [PubMed]

6. J. Mathews, R. Roucka, J. Xie, S.-Q. Yu, J. Menéndez, and J. Kouvetakis, “Extended performance GeSn/Si(100) p-i-n photodetectors for full spectral range telecommunication applications,” Appl. Phys. Lett. 95, 133506 (2009). [CrossRef]

7. M. Oehme, M. Schmid, M. Kaschel, M. Gollhofer, D. Widmann, E. Kasper, and J. Schulze, “GeSn p-i-n detectors integrated on Si with up to 4 % Sn,” Appl. Phys. Lett. 101, 141110 (2012). [CrossRef]

8. H. H. Tseng, K. Y. Wu, H. Li, V. Mashanov, H. H. Cheng, G. Sun, and R. A. Soref, “Mid-infrared electroluminescence from a Ge/Ge_0.922Sn_0.078/Ge double heterostructure p-i-n diode on a Si substrate,” Appl. Phys. Lett. 102, 182106 (2013). [CrossRef]

9. G.-E. Chang, S.-W. Chang, and S.-L. Chuang, “Strain-balanced Ge_1−zSn_z/Si_1−x−yGe_xSn_y multiple-quantum-well lasers,” IEEE J. Quantum Electron. , 46, 1813–1820 (2010). [CrossRef]

10. G. Sun, R. A. Soref, and H. H. Cheng, “Design of an electrically pumped SiGeSn/GeSn/SiGeSn double-heterostructure midinfrared laser,” J. Appl. Phys. 108, 033107 (2010). [CrossRef]

11. O. Gurdal, M. Hasan, M. R. Sardela, J. E. Greene, H. H. Radamson, J. E. Sundgren, and G. V. Hansson, “Growth of metastable Ge_1−xSn_x strained layer superlattices on Ge(001)2×1 by temperature-modulated molecular beam epitaxy,” Appl. Phys. Lett. 67, 956–958 (1995). [CrossRef]

12. A. Tonkikh, C. Eisenschmidt, V. Talalaev, N. Zakharov, J. Schilling, G. Schmidt, and P. Werner, “Pseudomorphic GeSn/Ge(001) quantum wells: Examining indirect band gap bowing,” Appl. Phys. Lett. 103, 032106 (2013). [CrossRef]

13. I. S. Yu, T. H. Wu, K. Y. Wu, H. H. Cheng, V. Mashanov, A. Nikiforov, O. Pchelyakov, and X. S. Wu, “Investigation of Ge_1−xSn_x/Ge with high Sn composition grown at low-temperature,” AIP Advances 1, 042118 (2011). [CrossRef]

14. E. Kasper, J. Werner, M. Oehme, S. Escoubas, N. Burle, and J. Schulze, “Growth of silicon based germanium tin alloys,” Thin Solid Films 520, 3195– 3200 (2012). [CrossRef]

15. H. Li, Y. X. Cui, K. Y. Wu, W. K. Tseng, H. H. Cheng, and H. Chen, “Strain relaxation and Sn segregation in GeSn epilayers under thermal treatment,” Appl. Phys. Lett. 102, 251907 (2013). [CrossRef]

16. R. Beeler, R. Roucka, A. V. G. Chizmeshya, J. Kouvetakis, and J. Menéndez, “Nonlinear structure-composition relationships in the Ge_1−ySn_y/Si(100) (y< 0.15) system,” Phys. Rev. B 84, 035204 (2011). [CrossRef]

17. J. Tauc, Optical Properties of Solids (North Holland, 1969).

18. G. E. Chang, W. Y. Hsieh, J. Z. Chen, and H. H. Cheng, “Quantum-confined photoluminescence from Ge_1−xSn_x/Ge superlattices on Ge-buffered Si(001) substrates,” Opt. Lett. 38, 3485–3487 (2013). [CrossRef] [PubMed]

19. H.-S. Lan, S.-T. Chan, T.-H. Cheng, C.-Y. Chen, S.-R. Jan, and C. W. Liu, “Biaxial tensile strain effects on photoluminescence of different orientated Ge wafers,” Appl. Phys. Lett. 98, 101106 (2011). [CrossRef]

20. J. P. Gupta, N. Bhargava, S. Kim, T. Adam, and J. Kolodzey, “Infrared electroluminescence from GeSn hetero-junction diodes grown by molecular beam epitaxy,” Appl. Phys. Lett. 102, 251117 (2013). [CrossRef]

21. S. L. Chuang, Physics of Photonic Devices (Wiley, New York, 2009), 2nd edition

22. S.-W. Chang and S.-L. Chuang, “Theory of optical gain of Ge-Si_xGe_y Sn_1−x−y quantum-well lasers,” IEEE J. Quantum Electron. 43, 249–256 (2007). [CrossRef]

23. V. D’Costa, Y.-Y. Fang, J. Tolle, J. Kouvetakis, and J. Menendez, “Ternary GeSiSn alloys: New opportunities for strain and band gap engineering using group-IV semiconductors,” Thin Solid Films 518, 2531–2537 (2010). [CrossRef]

24. P. Moontragoon, R. A. Soref, and Z. Ikonic, “The direct and indirect bandgaps of unstrained Si_xGe_1−x−ySn_y and their photonic device applications,” J. Appl. Phys. 112, 073106 (2012). [CrossRef]

Sample	t_GeSn (nm)	t_Ge (nm)	a^⊥(GeSn) (Å)	Sn content x (%)	$E_{g}^{Γ}$ (meV)	$E_{g}^{Γ}$ from PL(meV)
N1	6.0	2.5	5.6762	1.60	739	735
N2	6.5	2.3	5.6856	2.39	728	726
N3	6.4	2.2	5.7006	3.65	705	702
N4	6.2	2.2	5.7270	5.83	645	651
N5	6.1	2.2	5.7410	6.96	628	634

	E_g(Ge) (eV)	E_g(Sn) (eV)	b (eV)
Γ-valley	0.787	−0.413	2.42
L-valley	0.664	0.092	0.89

Sample	t_GeSn (nm)	t_Ge (nm)	a^⊥(GeSn) (Å)	Sn content x (%)	$E_{g}^{Γ}$ (meV)	$E_{g}^{Γ}$ from PL(meV)
N1	6.0	2.5	5.6762	1.60	739	735
N2	6.5	2.3	5.6856	2.39	728	726
N3	6.4	2.2	5.7006	3.65	705	702
N4	6.2	2.2	5.7270	5.83	645	651
N5	6.1	2.2	5.7410	6.96	628	634

	E_g(Ge) (eV)	E_g(Sn) (eV)	b (eV)
Γ-valley	0.787	−0.413	2.42
L-valley	0.664	0.092	0.89

Structural and optical characteristics of Ge_1−xSn _x /Ge superlattices grown on Ge-buffered Si(001) wafers

Abstract

1. Introduction

2. Experimental details

3. Results and discussion

4. Conclusion

Acknowledgments

References and links

Cited By

Figures (3)

Tables (2)

Equations (9)

Optical Materials Express