We report on tensile-strained Ge/Si0.11Ge0.89 quantum-well (QW) metal-semiconductor-metal (MSM) photodetectors on Si substrates. A tensile strain of 0.21% is introduced into the Ge wells by growing the QW stack on in-situ annealed Ge-on-Si virtual substrates (VS). The optical characterization of Ge/Si0.11Ge0.89 QW MSM photodetectors indicates that the optical response increases to a wavelength of 1.5 μm or higher owing to the strain-induced direct bandgap shrinkage. Analysis of the band structure by using a k · p model suggests that by optimizing the tensile strain and Ge well width, tensile-strained Ge/SiGe QW photodetectors can be designed to cover the telecommunication C-band and beyond for optical telecommunications and on-chip interconnection.
© 2016 Optical Society of America
Ge/SiGe quantum well (QW) systems have recently attracted increasing attention for Si-based photonic devices in terms of the development of optical telecommunications and on-chip interconnections [1, 2]. Such systems are fully compatible with standard complementary metal-oxide-semiconductor (CMOS) processes, allowing for monolithic integration with Si electronics. The Ge/SiGe QW systems have type-I alignment at the Γ-point and modified density-of-states by quantum confinement, leading to strong and fast direct interband transitions that are compatible with III–V counterparts [3,4]. This has inspired the development of Si-based Ge/SiGe QW photonic devices, including optical modulators [1–4], light emitters [5–7], and photodetectors [8–10]. However, the Ge/SiGe QW photonic devices are limited by the relatively larger direct bandgap energy of ∼0.88 eV, and can only operate in the telecommunications E-band wavelength range of 1410–1460 nm; they therefore hardly reach the commercially important telecommunications C-band (1530–1565 nm). Therefore there is a significant demand to shift the wavelength range of operation into the C-band for use in practical applications. To achieve this goal, it is necessary to shrink the direct bandgap to redshift the absorption edge. Adding Sn, another group-IV element, into the active layer has been proven effective for extending the absorption edge into longer wavelengths, and the photodetection range of GeSn photodetectors has been extended into the mid-infrared (MIR) region [11–19]. In addition to Sn-alloying, another intriguing approach is to introduce tensile strain into the Ge well [20, 21]; this involves growing tensile-strained Ge/SiGe QW structures. Recently, a few attempts have been made realizing tensile-strained Ge QW structures on Si substrates using post-growth rapid thermal annealing [22–24]. In such attempts, the bandgap shrinkage caused by tensile strain has been verified by photoluminescence and electroluminescence experiments. However, very little work has been reported on tensile-strained Ge/SiGe QW photonic devices.
In this paper, we investigate tensile-strained Ge/SiGe QW metal-semiconductor-metal (MSM) photodetectors on Si substrates. Tensile strain is introduced to the Ge wells by growing the Ge/SiGe QW stack on in-situ annealed Ge-on-Si virtual substrates (VSs). The structural properties of the grown Ge/SiGe QW are systematically characterized by different experimental techniques including cross-sectional transmission electron microscopy (XTEM) and Raman microscopy. Optical responsivity measurements of the fabricated Ge/SiGe QW photodetectors show that the photodetection cutoff wavelength extends beyond 1.5 μm. Lastly, band structure of the tensile-strained Ge/SiGe QW structure is also analyzed in order to investigate the impact of tensile strain.
2. Sample growth and characterization
The samples used in this study were grown on n-type Si (001) wafers with a resistivity of 1–10 Ω·cm using solid-source molecular beam epitaxy (MBE) at a base pressure of less than 2 × 10−10 torr. The epitaxy process began with the growth of a high-quality Ge VS on a Si wafer, which was realized by a two-step growth technique [25, 26]. The Ge VS consists of a 25 nm thick Si buffer grown at 700 °C, a 50 nm thick Si buffer grown at 350 °C, a 60 nm thick Ge seed layer grown at 350 °C, and a 60 nm thick Ge buffer layer grown at 500 °C. After the growth of the Ge VS, an in-situ annealing was then performed at 800 °C for 5 min. This annealing step introduces a tensile strain in the Ge VS during the cooling process from the elevated growth temperature to room temperature owing to the mismatch between the thermal expansion coefficients of Ge and Si . Subsequently, the growth temperature was decreased to 500 °C for the growth of five pairs of Ge/SiGe QWs. The entire structure was finally capped by a 3 nm thick Si layer. Although the samples were unintentionally doped, there is a p-type background doping concentration of ∼ 1 × 1015 cm−3 in the layers.
The microstructures and layer thicknesses of the samples were probed using XTEM (Model JEOL-2100F). Figure 1 shows the XTEM image of the Ge/SiGe QW sample. First, a large density of misfit dislocations is observed at the Ge/Si interface because of the 4.2% lattice mismatch between the Ge and Si layers. (The bulk lattice constants of Si and Ge are and , respectively .) The misfit dislocations indicate that the Ge VS is strain-relaxed. For the Ge/SiGe QW stack, homogenous periodicity and smooth barrier-well interfaces are observed. From the XTEM image, the thicknesses of the Ge wells and SiGe barriers are determined to be 7 nm and 10 nm, respectively. Additionally, dislocation lines were not observed near the Ge/SiGe interface in the high-resolution TEM (HRTEM) image, indicating that the Ge/SiGe QW is coherently strained to the underlying Ge VS.
The composition and strain status of the layers of the Ge/Si1−xGex QW structure were determined using Raman microscopy using a XploRA Raman microscope with a non-polarized 532 nm laser source. The excitation light source was focused onto a spot size of ∼ 1 μm, and the spectra were recorded in a backscattering geometry at room temperature, so that the Raman shift of the LO phonon was measured. Figure 2 shows the measured Raman shift of the Ge/Si1−xGex QW sample compared to that of a bulk Ge reference sample. For bulk Ge, a peak is observed at 300 cm−1, which is associated with the Ge-Ge LO mode, as indicated by the solid gray line. For the Ge/Si1−xGex QWs, a main peak accompanied by a shoulder at lower frequencies was observed. The former is assigned to the Ge-Ge LO mode from the Ge layers, whereas the later is attributed to that of the Si1−xGex layers. These peaks were fitted using Gaussian curves. The results are depicted in Fig. 2, which contains two peaks located at 299.05 cm−1 and 293.34 cm−1; from this we can deduce the frequency shifts of ΔωGe = −0.95 cm−1 and ΔωSiGe = −6.66 cm−1. Moreover, the strain value and composition in the layers can be deduced from the Raman shifts. For biaxially stressed Si1−xGex alloys, the strain- and composition-dependent frequency shift of the Ge-Ge LO mode is expressed as follows :Eq. (1), indicating that the Ge layers were tensily strained. Correspondingly, the in-plane lattice constant of the Ge layers is . For the SiGe barriers, the in-plane lattice constant can be expressed as 30]. Because of the pseudomorphic growth conditions of the Ge/Si1−xGex QW stack, . Given the frequency shift of ΔωSiGe = −6.66 cm−1 and using Eqs. (2)–(4), the Ge composition in the Si1−xGex barrier is determined to be x = 0.89 with a in-plane strain of .
3. Device fabrication and characterization
The Ge/Si0.11Ge0.89 QW MSM photodetectors were fabricated using CMOS-compatible processing. The mesas were defined using optical lithography and reactive ion etching (RIE) techniques upon the Si substrate. A 300 nm thick SiO2 passivation layer was then deposited by plasma-enhanced chemical vapor deposition (PECVD). To expose the surface of Ge/Si0.11Ge0.89 QW for metal contact, oxide windows were then opened using the RIE technique. 250 nm thick Ni contact pads were deposited by sputtering and patterned using a photoresist lift-off technique. A cross-sectional schematic diagram of the fabricated photodetector device is displayed in Fig. 3(a), and a scanning-electron-microscopy (SEM) image is shown in Fig. 3(b).
Figure 4 shows the measured dark current-voltage (I–V) characteristics curve of the fabricated Ge/Si0.11Ge0.89 QW MSM photodetectors at room temperature using a Keithley 2400 source meter. The dark current increases with increasing bias voltage at a low forward bias voltage, followed by a sharp increase in current at higher voltages. This nonlinear behavior indicates a clear rectifying behavior, suggesting that the Ge/Si0.11Ge0.89 QW is Schottky-type contact with Ni electrodes. Under reverse bias, the I–V curve displays similar behaviors, indicating a symmetric dark I–V characteristics attributed to the symmetric MSM structure. By fitting the linear part of the dark I–V curve under forward bias, as shown in the inset of Fig. 4, we obtain the turn-on voltage of Von = 0.62 V for the Ge/Si0.11Ge0.89 QW MSM photodetectors. A dark current of 6 μA is obtained before the Schottky diode is turned on.
Optical characterization of the Ge/Si0.11Ge0.89 QW MSM photodetectors was performed at room temperature using a quartz-tungsten-halogen broad-band light source, which was filtered using a 1200 nm long-pass filter, dispersed by a monochromator with a grating of 600 lines/mm blazed at a wavelength of 1.6 μm, chopped at 200 Hz, and normally incident onto the active area of the samples. The photodetector device was connected in series with a 50 Ω resistor, across which the modulated voltage was measured through a lock-in amplifier to obtain the photocurrent used to determine the optical responsivity (i.e., the ratio of the collected photocurrent to the power of the incident light, denoted by R.) Figure 5(a) shows the optical responsivity spectrum at room temperature obtained from the tensile-strained Ge/Si0.11Ge0.89 QW MSM photodetector. The optical responsivity gradually decreases with increasing wavelength, and reaches a cutoff near ∼ 1540 nm, as indicated by the solid arrow. Compared with the photodetection cutoff wavelength of ∼ 1410 – 1460 nm in conventional Ge/Si0.15Ge0.85 QW photodetectors [8–10], this tensile-strained Ge/Si0.11Ge0.89 QW photodetector exhibits an extended optical response in the infrared region covering the telecommunications O-, E-, S-, and C-bands. In addition, it is also noted that the photodetection range of this tensile-strained Ge/Si0.11Ge0.89 QW photodetector is comparable with Ge-on-Si photodetectors [31–33], but the performance is still interior to commercially available InGaAs photodetectors currently used in telecommunications. The extended optical response is mainly attributed to the direct bandgap energy shrinkage caused by the tensile strain (discussed later). Beyond the cutoff wavelength, the optical response becomes relatively weak, which can be attributed to the optical absorption by indirect L-valley interband absorption. To confirm the direct bandgap energy, the absorption coefficient (α) was extracted from the responsivity spectrum using34]. The extracted absorption coefficient is displayed in Fig. 5(b). It can be clearly seen that the absorption coefficient decreases with increasing wavelength, which is followed by a sharp onset at ∼ 1490 nm, as indicated by the arrow; this corresponds to the direct-gap absorption edge. Near the absorption edge, the lowest direct transition energy (Ed) can be extracted using the Tauc equation : Fig. 5(b). By fitting the near-edge absorption using Eq. (5), we obtain Ed = 0.847 eV.
4. Band structure analysis
To gain a deeper understanding of the strain effect on the Ge/Si0.11Ge0.89 QW photodetector, the strained electronic band structure was calculated and compared with that of a conventional strain-compensated Ge/Si0.15Ge0.85 QW grown on a strain-relaxed Si0.1Ge0.9 buffer [1,2]. The strained band structures were calculated using deformation potential theory , and the subband states were calculated using a multi-band k · p model [5, 28] with the parameters taken from . Figure 6 shows the calculated band structures and confined subband states for the conventional strain-compensated Ge/Si0.15Ge0.85 QW structure and the tensile-strained Ge/Si0.11Ge0.89 QW structure. For the conventional strain-compensated Ge/Si0.15Ge0.85 QW, as shown in Fig. 6(a), the structure exhibits a crucial type-I alignment with pronounced band offsets to localized electrons and holes in the Ge well region for efficient direct transitions. The 0.4% compressive strain in the Ge well increases the direct bandgap from 0.8 eV to 0.826 eV. In addition, the compressive strain pushes the heavy-hole (HH) band above the light-hole (LH) band, so the valence band is HH-like. The Γ-conduction band profile confines two electrons states in the QW (cΓ1 and cΓ2), leveling at 0.886 eV and 1.084 eV, respectively. For the valence band, two confined HH states (HH1 at −0.02 eV and HH2 at −0.074 eV) and one LH state (LH1 at −0.052 eV) are found. As a result, the lowest direct interband transition is the HH1-cΓ1 transition with a transition energy of Ed=0.905 eV, which defines the position of the absorption edge. On the other hand, for the tensile-strained Ge/Si0.11Ge0.89 QW structure, as shown in Fig. 6(b), the tensile strain effect significantly impacts the band structure. First, the 0.21% tensile strain lowers the direct bandgap energy of the Ge layer from 0.8 eV to 0.762 eV. For the conduction band, the profile shows that the Γ-valley conduction band minimum is located in the Ge layer with a barrier height of 0.205 eV, confining one electron state (cΓ1 at 0.812 eV). For the valence band, the tensile strain shifts the LH band above the HH band, both in the Ge and Si0.11Ge0.89 layers. The LH band profile indicates that the Ge layer serves as the barrier layer, where no hole is confined. For the HH band, the energy minimum is located in the Ge layer with a barrier height of 0.087 eV. In this way, holes are confined in the Ge layer in the HH1 and HH2 states situated at −0.042 meV and −0.097 meV, respectively. The band structure analysis indicates that the band alignment of the Γ-conduction band and the HH band is type-I, whereas that of the Γ-conduction band and the LH band is type-II. As a result, only interband transitions from the HH subbands to the Γ-valley conduction subband are possible, and the lowest direct interband transition is the HH1-cΓ1 transition. The calculated HH1-cΓ1 energy is Ed=0.857 eV, which agrees well with the experimental result of 0.847 eV. These results confirm that the optical absorption in the tensile-strained QW structure can be attributed to the HH1-cΓ1 direct interband transition. Furthermore, the comparison between the band structures of the conventional strain-compensated Ge/Si0.15Ge0.85 QW structure and the tensile-strained Ge/Si0.11Ge0.89 QW structure demonstrates that the use of tensile strain can effectively offset the transition energies in the Ge/SiGe QW structure, thereby redshifting the absorption edge in the infrared region.
Through the increased tensile strain in the Ge/SiGe QW structure (which shrinks the bandgap energy), the absorption edge can be redshifted in order to broaden the optical response of Ge/SiGe QW photodetectors. In addition to tensile strain, another approach to extending the photodetection range into longer wavelengths is to increase the well width. The combination of these two approaches has the potential to further extend the photodetection range towards longer wavelengths in order to cover the telecommunications C-band and beyond for practical applications. We therefore studied the effects of tensile strain in the Ge well and Ge well width on the lowest transition energy for the pseudomorphic Ge/Si0.11Ge0.89 QW structure. We first calculated the various band edges for the Ge wells and Si0.89Ge0.11 barrier as a function of tensile strain in the Ge well; the results are depicted in Fig. 7(a). For the conduction band, the Γ-conduction edges shift downward rapidly with increasing tensile strain. For the valence band, the LH bands are lifted higher than the HH bands, and their energy separation becomes larger with increasing tensile strain. As a result, the direct bandgap energies decrease with increasing tensile strain. Figure 7(b) depicts the barrier heights for the various bands in the Ge/Si0.89Ge0.11 QW structure as a function of tensile strain in the Ge well. The Γ-valley conduction band and HH band exhibit positive barrier heights, which are solely dependent on the tensile strain in the Ge well, indicating a type-I band alignment. In addition, the barrier heights are much larger than 1.5 kBT at 300 K; hence carriers can be effectively confined in the Ge well. On the other hand, the barrier height for the LH band is negative, so no holes are confined, as discussed above. As a result, the HH1-cΓ1 transition remains the lowest direct transition in the structure that designates the absorption edge. Figure 7(c) shows the calculated HH1-cΓ1 transition energy for the pseudomorphic Ge/Si0.11Ge0.89 QW structure as a function of tensile strain in the Ge well and the Ge well width, where the Si0.89Ge0.11 barrier width is set to 10 nm. As the tensile strain in the Ge well increases, the HH1-cΓ1 transition energy decreases owing to the bandgap reduction. With an increase of the Ge well width, the HH1-cΓ1 transition energy also decreases owing to the reduced quantized energies of carriers. However, increasing the Ge well width also weaken the quantum confinement effect, and the material eventually becomes more bulk-like. To maintain a proper quantum confinement, a well width of smaller than 15 nm is usually favorable. In this situation, a tensile strain of >0.274 % is required to shift the HH1-cΓ1 transition energy into the C-band window. A further increase in the tensile strain can lower the HH1-cΓ1 transition energy and redshift the absorption edge toward longer wavelengths in order to fully cover the C-band (and even the entirety of telecommunications windows) for important applications in telecommunications and chip-scale optical interconnections.
In summary, we have demonstrated tensile-strained Ge/Si0.11Ge0.89 QW photdetectors on Si substrates with extended infrared response. The structural properties were investigated using TEM and Raman microscopy, and revealed a tensile strain of 0.21% in the Ge wells. The optical response of tensile-strained Ge/Si0.11Ge0.89 QW MSM photodetector exhibits an extended optical response beyond a wavelength of 1.5 μm owing to the bandgap shrinkage caused by the tensile strain, as confirmed by the band structure analysis using the multi-band k · p method. The optical response can be further extended to longer wavelengths by optimizing the tensile strain and well width in the QW structure in order to enhance device performance at the important telecommunications C-band (and beyond). Those results represent a step toward realizing efficient Ge/SiGe QW photodetectors on Si substrates for telecommunications and on-chip optical connection applications.
Ministry of Science and Technology of Taiwan (MOST) (MOST 101-2112-M-002-015-MY3).
The authors would like to thank Dr. H. Li at National Taiwan University for the technical assistance with XTEM experiments.
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