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Abstract
A 2.7% tensile strained Ge/SiGe heterostructure nanowire (NW) is in-situ fabricated by a three-dimensional Ge condensation method. The NW metal-semiconductor-metal (MSM) photodetector demonstrates an ultra-broadband detection wavelength of 400-2400 nm, showing a high responsivity of >3.46×102 A/W with a photocurrent gain of >4.32×102 at 1550 nm under −2 V. A high normalized photocurrent to dark current ratio (NPDR) of 1.88×1011 W−1 at 1550 nm under −1 V is achieved. The fully complementary metal-oxide-semiconductor (CMOS) compatible, simple and scalable process suggest that the Ge heterostructure NW is promising for low cost, high performance near-infrared or short wavelength infrared focal plane array applications.
© 2019 Optical Society of America under the terms of the OSA Open Access Publishing Agreement
1. Introduction
Germanium (Ge) is one of the most promising materials for near-infrared (∼1550 nm) photodetectors (PDs) in optoelectronic integrated circuits due to its large absorption coefficient (benefitted from its direct band gap at 0.8 eV), quasi-direct band structure and compatibility with silicon (Si) architecture. The superiority impulses substantial investigations of on-chip Ge PDs featuring high frequency and responsivity in the past decade [1–3]. For example, Si-contacted Ge waveguide p-i-n PD with bandwidth of 67 GHz at −1 V [4] and Ge-on-Si PD without doped Ge or Ge-metal contacts with responsivity of 1.44 A/W [5] at 1550 nm have been demonstrated. And commercial bulk Ge PDs are already available. To realize future integration of the optics with modern nanometer-scaled electronics, nano-scaled PDs are indispensable. In this regard, quasi one-dimensional (1D) nanowires (NWs), which have distinct mechanical, electrical and optical properties from their bulk or film materials have attracted considerable attention. Advanced progresses have been achieved for Ge NW PDs with photoconductive gain of ∼10−1−107 [6–9]. However, due to quantum confinement effect (QCE) in the NWs, the detection cutoff wavelength of NWs would be blue shifted compared to the film and bulk materials (cutoff at ∼1600 nm) resulting in a dramatic drop of responsivity at telecom wavelength (1.3-1.6 µm). To date, few results have been reported for Ge NW PDs with detection wavelength of >1.6 µm [1,10].
By applying tensile strain in Ge NW, QCE can be counteracted since both the direct and indirect band gap of Ge would shrink under tensile strain with a transformation trend from indirect to direct band gap [11]. Many approaches have been proposed to induce tensile strain in Ge [12–18], which initially target at improving the light emitting efficiency of Ge. However, it faces greater challenges to apply tensile strain in Ge NW. With bridge structures or micro-electromechanical-system (MEMS) assisted technique, M. J. Süess et al. [13] and D. Nam et al. [19] induced a 3.1% and 2.77% uniaxial tensile strain in Ge NWs, respectively; By depositing an external Si3N4 stressor layer on Ge NW waveguide, M. de Kersauson et al. [20] obtained a 0.4% biaxial tensile strain in Ge NW. With these strain engineering, the direct band gap of Ge NW is red shifted with superior optical properties. However, post processing is inevitable for these methods rendering the entire fabrication process complicated and lack of productivity. A CMOS compatible, simple and scalable method is desired to fabricate tensile strain Ge NW.
Herein, high tensile strain is in-situ introduced into Ge NW during fabrication of Ge heterostructure NW by a three-dimensional (3D) Ge condensation method. Through a simple modification of SiGe NW width before 3D Ge condensation, a 2.7% tensile strained Ge/SiGe heterojunction NW can be fabricated without any post processing. The highly tensile strained NW metal-semiconductor-metal (MSM) PD presents an ultra-broadband detection wavelength of 400∼2400 nm at room temperature. Under −2 V, the NW PD shows a high responsivity of 3.46×102 A/W with a photocurrent gain of 4.32×102 at 1550 nm. The high responsivity and low dark current of the NW PD result in a high normalized photocurrent to dark current ratio (NPDR) of 1.88×1011 W−1 at 1550 nm under −1 V. The fully CMOS compatible and scalable process suggest a great potential of the NW for low cost, high performance near infrared or short wavelength infrared focal plane array applications.
2. Fabrication of highly tensile strained SiGe/Ge heterostructure NW
A 99-nm Si0.76Ge0.24 layer capped with 5-nm Si on p-type (001) silicon-on-insulator (SOI) substrate was used as initial material, as shown in Fig. 1(a). The thickness and resistivity of top Si for the SOI substrate is 120 nm and ∼10 Ω·cm, respectively. After planar Ge condensation process at 1150°C as reported previously [21,22], a 95-nm Si0.67Ge0.33 layer was obtained on the buried oxide (BOX) layer (Fig. 1(b)). The surface SiO2 generated during planar Ge condensation was then removed by buffered oxide etchant (BOE) to define patterns on SiGe layer. Through e-beam lithography (EBL) and reactive ion etch (RIE) techniques, SiGe-on-insulator NWs were fabricated along [110] orientation on the BOX layer, as displayed in Fig. 1(c). The NW with total length of 45 µm was fixed between two pads with dimension of 50 × 50 µm2. With different widths along the longitudinal direction, the NW can be divided into three parts: two segments with length of 20 µm and width of 677 nm at two sides; one segment with length of 5 µm and width of 200 nm at middle. Ultimately, the patterned sample was loaded into a 2-inch furnace to perform 3D Ge condensation process [23] at 900°C. Detailed condensation recipe is exhibited in Fig. 1(d). The whole condensation process consists 10 cycles of 10-min’s oxidation and subsequent 10-min’s annealing. The oxidation and annealing process were carried out in dry O2 with purity of 99.5% and N2 with purity of 99.999%, respectively.
Fig. 1. (a) Schematic diagram of as grown SGOI sample and (b) SGOI sample after planar Ge condensation process; (c) schematic diagram of SGOI NW structure before 3D Ge condensation process; (d) 3D Ge condensation recipe of SGOI NWs.
Figure 2(a) displays the optical image of the NW structure after 3D Ge condensation process. Figure 2(b) shows the Raman spectra of regions A (pad), B (NW with initial width of 677 nm) and C (NW with initial width of 200 nm) in Fig. 2(a), respectively. The signal was excited by a 488-nm laser then collected by WITec alpha300 confocal micro-Raman system. The spectra from Si0.67Ge0.33 layer and bulk Ge are also presented for comparison. As can be seen, after 3D Ge condensation process, the intensity of Ge-Ge peaks (around 300 cm−1) from the NWs increased significantly compared to that from Si0.67Ge0.33 layer indicating enrichment of Ge content in NWs. For SiGe pad and 677-nm SiGe NW, similar Raman spectra were acquired suggesting that similar Ge content was obtained. For SiGe NW with initial width of 200 nm, Si-Ge peak (around 400 cm−1) disappears after the 3D Ge condensation process manifesting formation of pure Ge NW.
Fig. 2. (a) Optical image of the NW structure after 3D Ge condensation process; (b) Raman spectra of regions A (pad), B (NW with initial width of 677 nm) and C (NW with initial width of 200 nm) in Fig. 2(a), respectively. The spectra from Si0.67Ge0.33 layer and bulk Ge are also presented for comparison.
From the peak positions of Si-Si (at around 500 cm−1) and Si-Ge modes (for SiGe with Ge content of <0.5) [22,24], the Ge content and strain in SiGe Pad are evaluated as 0.49 and −0.19%, respectively. From the integral intensity of Ge-Ge and Si-Ge peaks (for SiGe with Ge content of >0.5) [22,25], the Ge content and strain in SiGe NW with initial width of 677 nm are calculated as 0.52 and 0.19%, respectively. From the difference between peak position of Ge-Ge mode from Ge NW (291.0 cm−1) and that from bulk Ge (301.8 cm−1), a tensile strain of 2.70% in Ge NW can be calculated [26]. For planar sample, based on the thermal expansion coefficient (TEC) difference between Ge [27] (8.51×10−6/°C) and SiO2 [28] (0.50×10−6/°C), the maximum strain that can be accumulated in Ge after cooling from 900 to 20°C is ∼0.70%. It has been reported that the tensile strain of Ge mesa can be enhanced by surrounding SiO2 after Ge condensation process [29]. Thus, we ascribe the large tensile strain in Ge NW to the SiO2-encapsuled structure and small volume of Ge NW. However, quantitative analysis of the strain distribution needs to be further investigated in future.
Figure 3(a) shows the cross-section image of the Ge NW taken from transmission electron microscope (TEM). As can be seen, a 23-nm thick Ge layer with final width of ∼140 nm was encapsulated in SiO2. Due to larger oxidation rate of atoms in the vicinity of the edges and corners, a convex shape was obtained at two sidewalls. The thickness of SiO2 around the NW sidewall ((110) facet) is ∼35 nm, while the thickness of surface oxide varies from ∼35 to ∼45 nm, implying higher oxidation rate on sample surface ((001) facet). Figure 3(b) displays the high resolution TEM (HRTEM) image of region I in Fig. 3(a), from which well-ordered Ge atoms can be observed. It is noticed that a thin interfacial layer is formed on the top surface and sidewall of Ge, which is GeOx produced during Ge condensation process [30]. No dislocation line is found over the whole cross-section structure indicating formation of high-quality Ge NW. By averaging 20 atomic spacing in the HRTEM image, the interplanar spacing of (111) and (001) facets was measured as 0.3310 and 0.5572 nm, respectively. Figure 3(c) shows the selective area electron diffraction (SAED) pattern of the Ge layer. The dot array confirms the good crystal quality of Ge layer. The spots at P, Q and R sites are resulted from electron diffraction from (111), (002) and ($\overline 1 \overline 1 1$) facets, respectively. The length of OP and OQ is 3.128 and 3.586 nm−1, from which an interplanar spacing of 0.3304 and 0.5577 nm can be extracted for (111) and (001) facets, respectively, agreeing well with the results measured from Fig. 3(b). By averaging the results from Figs. 3(b) and 3(c), the interplanar spacing of (111) facet and lattice constant along [001] orientation is 0.3307 ± 0.0003 and 0.5575 ± 0.0003 nm, respectively. A compressive strain of −1.46% along [001] direction and an interplanar spacing of 0.4078 nm for (110) facet can be calculated consequently. Based on the principle of volume invariability of primitive cell under strain, the lattice constant along [100] and [010] orientations is calculated as 0.5768 nm corresponding a tensile strain of 1.96%. Thus, a uniaxial tensile strain of 2.77% is produced along [110] orientation, agreeing well with the Raman result.
Fig. 3. (a) Cross-section TEM image of the Ge NW; (b) HRTEM image of region I in Fig. 3(a); (c) SAED pattern of Ge layer.
3. Highly tensile strained SiGe/Ge heterostructure NW PD
From the above analysis, it was verified that highly tensile strained SiGe/Ge heterojunction NW was formed by 3D Ge condensation technique. With the double heterojunctions, MSM PDs were further fabricated. Figure 4(a) displays the optical image of the fabricated device. Silicon dioxide contact windows with width of 5 µm were opened on SiGe regions astride the Ge NW and 50-nm Al was deposited on SiGe as metal contacts. The distance between the two contacts is 15 µm. Figure 4(b) shows the I-V curve of the MSM PD under dark (solid circle) and illumination of 1550 nm light with a power of 2.75 µW (solid triangle) in semi-log (black) and linear scales (red). In linear scale, the currents increase nonlinearly with applied voltage indicating existence of a Schottky barrier between Al and SiGe. Under −2 V bias, the dark current is only ∼2 nA. Under illumination of 1550 nm light, the current is dramatically enhanced by more than 10 folds under >|±1| V bias, indicating high sensitivity of the PD.
Fig. 4. (a) Optical image of the fabricated tensile strained SiGe/Ge double heterojunction NW PD; (b) I-V curves of the NW PD under dark and illumination of 1550 nm light with a power of 2.75 µW; (c) schematic diagram of the setup for responsivity measurements; (d) response spectra of the NW PD under various bias; (e) schematic diagram of light absorption in the NW PD with different photon energy; (f) dependence of responsivity and NPDR of the NW PD on applied voltage at 950, 1300 and 1550 nm.
We then measured the response spectra of the highly tensile strained NW PD. Tungsten halogen lamp was used as light source (400∼2400 nm) during the measurements. The light was first coupled into a monochrometer to obtain monochromatic light, then coupled into a 600-µm fiber and finally shined normally on the device through a microscope objective. Figure 4(c) shows the schematic diagram of the setup for responsivity measurements. The intensity of monochromatic light was calibrated by a commercial Si (400-1100 nm) and InGaAs (1100∼2400 nm) PD. The spot size of the incident light is a bit larger than 600 µm, while the detection area of our device is less than 8 µm2. Since the detection area of PD is much smaller than the spot size, it is hard to accurately calculate the incident power. For conservative estimation, we consider a Gaussian incident light with a spot size of 600 µm and assume that the light within the most center 8-µm2 region was all received by the PD. Figure 4(d) displays the response spectra of the NW PD under −0.5∼−2.0 V bias. The NW PD shows a strong response over visible and near-infrared wavelength (400-2400 nm) with the following features: a response valley at around 1100 nm, an absorption edge at around 2100 nm and a small response shoulder between 2100 to 2400 nm. The absorption edge at around 2100 nm is determined by the direct band gap of Ge, while the small response shoulder between 2100 to 2400 nm may be determined by the Schottky barrier height (SBH) of SiGe/Al contact due to Fermi-level pinning effect. A similar extended light absorption of up to 1650 nm from Si/metal SBH has been reported [31]. The appearance of response valley at around 1100 nm is resulted from light resonance in the structure and light absorption by Si substrate. As displayed in Fig. 4(e), for photons with energy (hν) of smaller than the band gap of Si (EgSi), since the light is transparent to Si substrate, most of the transmission light from Ge (or SiGe) NW can be reflected back to Ge (or SiGe) NW by SiO2/Si interface leading to a relative high responsivity. As photons energy approximates to the band gap of Si (EgSi), an additional loss of the transmission light due to absorption by Si would be induced. Less photons would be reflected back to Ge (or SiGe) NW giving rise to a decrease of responsivity. However, as photon energy further increases (hν>>EgSi), light absorption in Si substrate can be neglected since the incident photons would be consumed in Ge (or SiGe) due to increase of light absorption coefficient in Ge (or SiGe) (the absorption length of 580 nm light in bulk Ge is less than 24 nm [32]). Thus, a response valley at around 1100 nm can be observed.
The black curves in Fig. 4(f) display the dependence of the responsivity on applied voltage at 950, 1300 and 1550 nm. As reverse bias enlarges, the responsivity of the PD increases dramatically. Under −2.0 V bias, the responsivity at 950, 1300 and 1550 nm is 5.52×102, 2.87×102 and 3.46×102 A/W, respectively. The normalized photocurrent to dark current ratio (NPDR) can be calculated by [33]:
where Iphoto and Idark are photo and dark current, respectively; Pin is incident optical power; R is responsivity. Based on Eq. (1), the NPDR at 950, 1300 and 1550 nm as a function of applied voltage is calculated and displayed as the blue dash curves in Fig. 4(f). Under −1.0 bias, a maximum NPDR of 3.04×1011, 1.45×1011 and 1.88×1011 W−1 at 950, 1300 and 1550 nm is achieved, respectively. The NPDR is significantly higher than those of typical Si (∼10 W−1), Ge (∼102 W−1) and emerging 2D material (∼102 W−1) PDs [34] suggesting promising application potential of the Ge NW PD.For bulk semiconductor PD, the responsivity is generally lower than 1 A/W. The high responsivity of the NW PD indicates existence of high photocurrent gain (G). The photocurrent gain of NW PD can be calculated as:
where q is elementary charge and hν is photon energy of incident light. Based on Eq. (2), under −2 V bias, the photocurrent gain at 950, 1300 and 1550 nm is 4.23×102, 3.01×102 and 4.32×102, respectively. The high photocurrent gain of the heterojunction NW PD indicates prolonged lifetime of the photogenerated carriers since the photocurrent gain is proportional the ratio between the lifetime and the transit time of photogenerated carriers. Due to the large surface-to-volume ratio of NW, the photogenerated carrier dynamics is strongly affected by carrier trapping and/or scattering at the surface localized energy states. It has been reported that the presence of GeOx on Ge can induce a depletion layer on the NW surface with a radial electric field from GeOx to inner NW [9], as shown in Fig. 5(a). The electrons in Ge NW can be effectively trapped in the surface local states by the radial electric field as displayed in Fig. 5(b). This process inhibits carrier recombination and thus prolong the lifetime of photogenerated electrons resulting in a high photocurrent gain of the NW PD.Fig. 5. (a) Schematic diagram of photogeneration process of carriers in Ge NW; (b) schematic diagram of electron trapping process by the surface local states induced by GeOx on Ge NW. (c) schematic diagram of band alignment for the heterojunction NW PD under (I) thermal equilibrium and (II) bias.
Figure 5(c) shows the schematic diagram of band alignment for SiGe/Ge heterojunction NW PD under (I) thermal equilibrium and (II) reverse bias. Silicon, Ge, and SiGe have a similar electron affinity energy (∼4.0 eV). Under tensile strain, their conduction and valance band edges would shift closer to the center of forbidden band [35]. Since the deformation constants of conduction and valence bands of SiGe are smaller than those of Ge [35], with a larger tensile strain in Ge than that in SiGe, a “type I” band alignment is expected for the fabricated SiGe/Ge heterojunction. Under bias, the photogenerated electrons and holes would be driven to the anode and cathode, respectively. The presence of conduction-band discontinuity would further facilitate the trapping of electrons by the surface local states induced by GeOx on the Ge NW.
The above results demonstrate achievement of high-performance Ge heterostructure NW PD at room temperature, which can have many applications, for example, near infrared or short wavelength infrared focal plane array. Traditionally, near-infrared PD arrays are fabricated by HgCdTe or III-V materials [36]. These materials are not compatible with Si architecture. Additionally, for PDs fabricated with these materials, cooling (typically down to 77 K) is always required [37]. The fully CMOS compatible, simple and scalable process and high performance at room temperature suggest that the fabricated Ge heterostructure NW is very promising for near infrared or short wavelength infrared focal plane array applications.
4. Conclusion
In summary, high tensile strain is in-situ introduced into Ge NWs during fabrication of Ge heterostructure NW by a three-dimensional (3D) Ge condensation method. Through simple modification of SiGe NW width before 3D Ge condensation, tensile strain SiGe/Ge heterojunction NWs with tunable Ge content of SiGe can be feasibly fabricated. With the heterojunction NW, metal-semiconductor-metal (MSM) PDs were fabricated. Owing to the 2.7% tensile strain in Ge NW, an ultra-broadband detection wavelength of 400∼2400 nm in Ge NW was observed. Under −2 V, the NW PD presents a high responsivity of 3.46×102 A/W with a photocurrent gain of 4.32×102 at 1550 nm. The high responsivity and low dark current of the NW PD result in a high normalized photocurrent to dark current ratio (NPDR) of 1.88×1011 W−1 at 1550 nm under −1 V. The fully CMOS compatible, simple and scalable process suggest a great potential of the NW for low cost, high performance near infrared or short wavelength infrared focal plane array applications.
Funding
National Key R&D Program of China (2018YFB2200103); National Natural Science Foundation of China (61176092, 61474094).
Disclosures
The authors declare no conflicts of interest.
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