Enhanced direct bandgap emission in germanium ... - OSA Publishing

Enhanced direct bandgap emission in germanium by micromechanical strain engineering Peng Huei Lim, Sungbong Park, Yasuhiko Ishikawa, and Kazumi Wada* Department of Materials Engineering, the University of Tokyo, 7-3-1, Hongo Bunkyo,Tokyo 113-8656, Japan *[email protected]

Abstract: We propose a new class of optoelectronic devices in which the optical properties of the active material is enhanced by strain generated from micromechanical structures. As a concrete example, we modeled the emission efficiency of strained germanium supported by a cantilever-like platform. Our simulations indicate that net optical gain is obtainable even in indirect germanium under a substrate biaxial tensile strain of about 1.5% with an electron-hole injection concentration of 9x1018 cm−3 while direct bandgap germanium becomes available at a strain of 2%. A large wavelength tuning span of 400 nm in the mid-IR range also opens up the possibility of a tunable on-chip germanium biomedical light source. ©2009 Optical Society of America OCIS codes: (130.3120) Integrated optics devices; (130.0250) Optoelectronics

References and links 1. 2. 3. 4. 5. 6. 7.

8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18.

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Received 6 May 2009; revised 4 Aug 2009; accepted 5 Aug 2009; published 31 Aug 2009

31 August 2009 / Vol. 17, No. 18 / OPTICS EXPRESS 16358

19. T. Namazu, Y. Isono, and T. Tanaka, “Evaluation of size effect on mechanical properties of single crystal silicon by nanoscale bending test using AFM,” J. Micromech. Sys. 9(4), 450–459 (2000). 20. T. Alan, A. T. Zehnder, D. Sengupta, and M. A. Hines, “Methyl monolayers improve the fracture strength and durability of silicon nanobeams,” Appl. Phys. Lett. 89(23), 231905 (2006). 21. K.-S. Chen, A. Ayon, and S. M. Spearing, “Controlling and Testing the Fracture Strength of Silicon on the Mesoscale,” J. Am. Ceram. Soc. 83(6), 1476–1484 (2000). 22. D. C. Houghton, G. C. Aers, S. Yang, E. Wang, and N. L. Rowell, “Type I band alignment in Si1-xGex/Si(001) quantum wells: photoluminescence under applied [110] and [100] Uniaxial Stress,” Phys. Rev. Lett. 75(5), 866–869 (1995). 23. P. G. Evans, D. S. Tinberg, M. M. Roberts, M. G. Lagally, Y. Xiao, B. Lai, and Z. Cai, “Germanium hut nanostressors on freestanding thin silicon membranes,” Appl. Phys. Lett. 87(7), 073112 (2005). 24. A. E. Franke, and M. J. Heck, T.-J. King, R. T. Howe, “Polycrystalline silicon-germanium films for integrated microsystems,” J. Micromech. Sys. 12(2), 160–171 (2003). 25. A. E. Franke, D. Bilic, D. T. Chang, P. T. Jones, T.-J. King, R. T. Howe, and G. C. Johnson, “Post-CMOS integration of germanium microstructures,” MEMS '99. Twelfth IEEE International Conference on, 630–637 (1999). 26. P. T. Jones, G. C. Johnson, and R. T. Howe, “The fracture strength of polycrystalline silicon,” Proceedings of MRS 1998 Spring Meeting, April 13–17, San Francisco, 197–202 (1998). 27. N. Lobontiu and E. Garcia, Mechanics of Microelectromechanical Systems, (Kluwer Academic Publishers, 2005). 28. F. Zhang, V. H. Crespi, and P. Zhang, “Prediction that uniaxial tension along produces a direct band gap in germanium,” Phys. Rev. Lett. 102(15), 156401 (2009). 29. S. Dhar, “Analytical Mobility Modeling for Strained Silicon-Based Devices,” Dissertation, Vienna University of Technology (2007). 30. O. Madelung ed., Semiconductors – Basic Data, (Springer, Berlin, 1996). 31. C. G. Van de Walle, “Band lineups and deformation potentials in the model-solid theory,” Phys. Rev. B 39(3), 1871–1883 (1989). 32. S. L. Chuang, Physics of Optoelectronic Devices, (John Wiley & Sons, New York, 1995). 33. T. C. Chong, and C. G. Fonstad, “Theoretical Gain of Strained-Layer Semiconductor Lasers in the Large Strain Regime,” IEEE J. Quantum Electron. 25(2), 171–178 (1989). 34. M. M. Rieger, and P. Vogl, “Electronic-band parameters in strained Si1-xGex alloys on Si1-yGey substrates,” Phys. Rev. B 48(19), 14276–14287 (1993). 35. F. Schaeffler, “High-mobility Si and Ge structures,” Semicond. Sci. Technol. 12(12), 1515–1549 (1997). 36. D. J. Paul, “Si/SiGe heterostructures: from material and physics to devices and circuits,” Semicond. Sci. Technol. 19(10), R75–R108 (2004). 37. J. Liu, Massachusetts Institute of Technology, Cambridge, MA 02139 (personal communication, 2008).

1. Introduction Crystalline germanium (Ge)-based [1–6] emitters are the subject of growing interest because of their CMOS compatibility and their C band direct bandgap emission spectrum. One major disadvantage of Ge as an emission medium is its indirect nature, which limits direct optical transitions by way of electronic scattering into the indirect valleys. Such losses can be suppressed using heavy n-doping [2,5,6] and the emission rate enhanced by the Purcell effect in highly confined optical cavities [4]. The small 0.14 eV separation between the direct Γ and the indirect L valley electrons can be further minimized in tensile-strained Ge grown on germaniumsilicon-tin (GeSiSn) alloys [1]. A biaxial tensile strain of approximately 2% [7] is thereby sufficient to transform Ge into a direct bandgap material and an efficient luminescence source. Despite the potential these methods hold, some issues remain unresolved. Firstly, Purcell effect does not directly address the large intervalley electronic scattering which is the most serious hurdle for efficient Ge light emission. In the case of the n-doped Ge light emitter, the near ohmic carrier concentration in the active layer causes unwanted current flow during electrical operation. Forward biasing of the device is needed to introduce minority holes into the active medium but will simultaneously inject significant number of electrons from the electronrich active layer into neighbouring non-active regions where their energy will be dissipated. For Sn-based group IV alloys, buffer deposition is constrained by the thermal budget of downstream CMOS processes, some of which, such as dopant annealing, require temperatures of about 800°C. As a comparison, liquid phases in Ge1-xSnx mixtures exist at equilibrium for temperatures greater than 232°C for x>1.2% and x is restricted to below 0.7% for solid alloys at a temperature of 800°C [8]. There is therefore sufficient motivation to explore alternatives for enhancing Ge light emission. Strain engineering of Ge appears promising since the direct to indirect bandedge offset of Ge can be significantly reduced or even eliminated under tension. Johansson et al. [9] first reported large uniaxial tensile strains of almost 6% in Si micromechanical cantilevers and the #111065 - $15.00 USD

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use of mechanically applied strain to enhance transistor performance is well-known [10]. We believe that applying a similar approach in Si/Ge photonics will reap similar benefits. The electroabsorption contrast of a Ge-based Franz-Keldysh modulator is limited by the L valley optical absorption [11], which is reducible by tensile strain. The linear electro-optical effect has also been observed in strained Si-based photonic crystals where the inversion symmetry is broken [12]. Further, Liu et al. [2] estimated that at a biaxial tensile strain of 2%, the Ge direct bandgap will narrow down to 0.5 eV or a wavelength of 2.5 µm. Such small bandgaps corresponds to the mid-IR range around 3 µm, which is the “fingerprint” spectrum region in biochemical sensing. While tensile strained Ge is advantageous for many applications, present techniques to realize this are limited. Ultrathin layers of compressive Ge have been pseudomorphically grown on Si whose lattice mismatch is as large as 4.2%. Compression is however, unsuitable for light emission work since the absorptive indirect band now penetrates deeper into the direct bandgap of the material. By exploiting the difference in thermal expansion coefficients of Si and Ge, Ishikawa et al. [13] achieved a 0.2% biaxial tensile strain in bulk Ge on Si deposited by a 2-step chemical vapor deposion (CVD) process [14]. Post-growth silicidation on the Si wafer backside can increase biaxial tensile strain to 0.24% through substrate warping [15] while the larger 1.2% biaxial tensile strain on free standing SiGe/Si sandwich structures supported by SiO2 pedestals exists in the Si layer [16]. Thus, strain obtainable by these methods remains insufficient to create a direct bandgap in Ge. In this paper, we analyze the properties of biaxial tensile strained Ge deposited on deformable micromechanical system (MEMS) Si structures and show that significant optical gains are obtainable in Ge under realistic conditions in MEMS devices. 2. Micromechanical strain generation The theoretical fracture strengths of Si and Ge are large: for the crystalline direction, Ruoff [17] has computed the maximum Si and Ge uniaxial tensile strain and stress to be 20.6% and 21.4 GPa as well as 18.3% and 14.7 GPa respectively, while Roundy et al. [18] has calculated the corresponding tensile strain to be 17% at 22 GPa stress for Si and 20% at 14 GPa stress for Ge. In practice, fracture limits are significantly lower and highly dependent on sample size and crystalline quality. Namazu et al. [19] noted that the average bending strength of 17.53 GPa for submicron Si beams is 38 times larger than that of millimeter-sized samples. Recently, Alan et al. [20] also achieved high tensile stresses of 18.2 GPa in Si using cantilevers with surface passivation by methyl monolayers. These results indicate that micromechanical structures are a feasible method to approach the theoretical strain limits predicted in [17] and [18]. Legend (GPa) 5 -1.1

z

y

4

x 3 2 0

-2.1 -3.1 -4.1 -5.1

Fig. 1. in-plane stress distribution of our proposed strain generation structure as calculated by FEA. The x, y and z axes are the crystallographic axes , and respectively. The distribution is obtained by a 90° rotation about the axis. Positive stress values are tensile.

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Biaxial Si fracture stresses was investigated by Chen et al. [21] in 10 mm2 -oriented samples 230 nm to 500 nm thick. A maximum Weibull strength of 4.6 GPa was found, which, based on their assumed isotropic Young's modulus of 170 GPa and Poisson's constant of 0.1, would yield a biaxial tensile strain of 2.44%. We note that this strain is sufficient to transform Ge into a direct bandgap material. Some preliminary experiments on strained Ge structures on deformed substrates have already been reported. Houghton et al. [22] observed type I alignment of SiGe/Si quantum wells with the application of uniaxial stress on their samples and Evans et al. deposited Ge hut nanostructures on thin Si membranes by molecular beam epitaxy and created significant bending under the Ge mesas [23]. However, this growth method does not lend itself to blanket deposition of bulk Ge on Si. We are unaware of any published work on crystalline Ge-based MEMS structures but low temperature deposited poly-Ge is a widely used sacrificial material in CMOS-based MEMS due to its high etching selectivity with respect to poly-SiGe and Si [24]. Franke et al. [25] measured the average uniaxial fracture strain in poly-Ge to be 1.7% for micron-scale structures and a corresponding poly-Si value of 1.5% is also obtained via the same method [26]. The value for poly-Si is still an order of magnitude smaller than the best reported experimental results for crystalline Si [20] and the theoretical values [17], [18]. Based on the above considerations, we shall assume that maximum biaxial tensile strain of crystalline Ge is much larger than that of poly-Ge as well and exceeds the 2% required for a direct bandgap transformation. Figure 1 depicts an embodiment of our proposed strain-generating structure supporting a Gebased light emitter. A square 80 µm by 80 µm thin Si platform is suspended using 4 Si cantilevers which are themselves attached to the main Si chip. The entire Si structure forms a cross with each arm extending 100 µm from the platform and having a uniform thickness of 10µm. A Ge ring or disk of radius 10 µm and thickness 400 nm is sited at the center of the platform. Such Ge rings can either be fabricated by selective CVD growth [14] or by postgrowth patterning. A rib-ring can be an alternative structure when the channel ring is not strained as efficiently as we expected. To exert biaxial tensile stress on the Ge emitter in the cross structure, a vertical force of 330 mN is applied at the bottom of the platform center, either using an external L-shaped hook constituting part of the external chip packaging or a MEMS-based electrostatic actuator with the silicon membrane. Locating the force applicator off-chip makes this technique amenable even to processes without the addition of MEMS capabilities. On integrated MEMS chips, on-chip movement control has the added advantage that the strain and the device parameters are tunable on the fly. Exerting upwards pressure on the platform contorts the supporting cantilevers which, in tandem, generate a biaxial in-plane stress on the structure center. As in the case of a vertically loaded MEMS cantilever [27], the arching of the stressed structure creates a linear strain field which changes from compressive at the bottom surface to tensile at the top. A simple cantilever-type mechanism is sufficient when a uniaxial tensile strain is required. In reference [28], Feng Zhang et al. predicted that uniaxial tension along Ge results in a direct bandgap at a tensile stress of 6.5 GPa or a tensile strain of 4.2%. Attaining such high uniaxial strains typically involves complex selective growth processes, Ge nanowires as suggested by the authors or the optimization of the geometries and orientations of quantum dots. With cantilever structures [20], we believe that comparable uniaxial tensile strains are likewise achievable in bulk Ge. The strain field calculated by finite element analysis (FEA) of our proposed device is shown in Fig. 1. For simplicity, we neglect the small Ge ring structure and assume that it will take on the substrate curvature and strain given the relative thicker height and substantially larger volume of the substrate. The strain distribution within the Ge layer is then linearly extrapolated from that of the Si substrate. A more exact simulation to account for the strain variation near the Ge ring will be the subject of further study. We further simplify our analysis by approximating crystalline Si as isotropic with a Young’s modulus constant of 170 GPa and a Poisson coefficient of 0.1. Chen et al. [21] found that this assumption does not substantially alter their FEA results as compared to a full fledged simulation with cubic elastic constants for Si. #111065 - $15.00 USD

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The maximum uniaxial stress magnitude of 5 GPa is located at the outer ends of the cantilever support and is comparable to that of similar-sized cantilevers reported in [9]. The largest biaxial tensile stress, also about 5 GPa for both in-plane crystallographic axes, is located at the platform center where the Ge ring is located. This amounts to an in-plane biaxial tensile strain of 2.77% (and a corresponding out-of-plane compression strain of 2.14% and 2.08% in the silicon and germanium layers respectively) using the elastic constants given in the next section and an in-plane biaxial tensile strain of 2.64% using the isotropic assumption of the FEA simulations here. The other components of the surface stress tensor are much smaller at Ge ring. By the same axes convention in Fig. 1, the surface vertical stress σzz is negligible while the shear stress σxy and the sum total of σyz and σzx amount to 2.5% of the biaxial stress, σxx or σyy along the Ge emitter ring. 3. Electrooptical properties of strained Ge The normally degenerate light hole (lh) and heavy hole (hh) bands of unstrained Ge are now lifted, resulting in a lh valence bandedge. Simultaneously, the energetic offset from the conductance Γ to L valleys, ∆EΓL is reduced. All bandedges, move towards the bandgap center leading to a smaller direct Γ-lh interband transition energy EgΓlh. In our simulation model, we assume that the Ge layer takes on the curvature and hence the strain of the underlying Si substrate. The Si strain tensor, and by extension the Ge strain tensor, can be calculated from the Si surface stress tensor by the generalized Hookes’ Law [29] with the appropriate elastic stiffness constants for Si, c11 = 1.66 Mbar, c12 = 0.639 Mbar and c44 = 0.796 Mbar and for Ge, c11 = 1.29 Mbar, c12 = 0.48 Mbar and c44 = 0.68 Mbar [30]. Following the discussion in reference [31], the change in the Γ and L bandedges due to strain, ∆ECΓ and ∆ECL respectively are derived from the strain tensor by Eq. (7) and Eq. (17) of reference [31]. The Ge deformation potential constants (in the notation as the aforementioned reference) are taken to be adirC = −8.24 eV, aindirC = −1.54 eV and ΞLU = 15.13 eV [31]. The corresponding behavior of hole states is more complex due to their close energetic and momentum proximity at the Γ point. Bandmixing occurs in the presence of shear strain and the new hole eigenstates are taken as a linear sum of the unstrained states by diagonalizing the Pikus-Bir Hamiltonian [32]. Since the shear components in our system are small, we shall treat the valence band carriers as hh-like or lh-like and ignore the intermixing and anti-crossings of the hole states. As with the case for the conduction band calculations, the relevant Ge parameters to determine the lh and hh bandedges, b = −2.55 eV, d = −5.5 eV, ∆0 = 0.3 eV are extracted from reference [31]. Although shear strain lifts the degeneracy of the L band, we assume the energy of the least energetic L valley as that for all L valleys. This simplification will require a larger applied tensile strain and is therefore a more stringent criterion for a given decrease in the indirect L to direct Γ separation. In accordance with the preceding FEA analysis, the shear components of the surface Si stress tensor are taken to be 2.5% of the in-plane biaxial hydrostatic stress value and their signs are chosen to give the least energetic L valley in the Ge layer using the proceeding method while the vertical stress components are taken to be zero. The strain tensor is calculated from the stress tensor using Hook’s law. At the Si/Ge interface, all Si strain components except for the vertical strain, εzz are transferred over to the Ge layer and the latter is subsequently calculated from the in-plane strain assuming zero vertical stress via Hook's law. We approximate the linear Ge tensile strain field in the cantilever structure by discretizing it into 10 sublayers parallel to the Si base. For each strained sublayer, the in-plane strain changes linearly as the same rate as the Si base and the energy shifts in the density of states (DOS) of the L, Γ, hh and lh bands are calculated in accordance with the deformation potential theory. To determine the gain, we assume dynamic equilibrium in the Ge layer and a constant number of injected electrons and holes, whose distributions are regulated by their respective pseudo-Fermi levels. For given hole and electron populations, global quasi-Fermi levels common across all sublayers were computed numerically from standard Fermi statistics. The strain-dependent change in both the conductance Γ [33] and L valleys [34], [35] are assumed small and neglected in the following computations. Biaxial tensile strain lifts the lh and hh #111065 - $15.00 USD

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degeneracy and introduces anisotropy in their bandedges. We approximate the modified lh and hh masses masses parallel and perpendicular to the strain plane as mlh///mo = 1/(γ1 − f+γ2), mlh,perpendicular/mo = 1/(γ1 + 2f+γ2) mhh///mo = 1/(γ1 + γ2), mhh,perpendicular/mo = 1/(γ1 − 2γ2)

( (

2 x 1 + 32 x − 1 + 1 + 2 x + 9 x 2 where f + = 3 4

( x −1 +

1 + 2x + 9x

2

2

) + x −1 +

)) + 6x

(1)

2

is a correction factor for the 2

1 + 2 x + 9 x − 3x

2

spin-orbit split-off band interaction and x = Qε/∆0, whereby Qε = −1/2b(εxx + εyy − 2εzz) [32]. The unstrained Ge Luttinger parameters are given by γ1 = 13.4, γ2 = 4.25 and γ3 = 5.69 [36]. The use of Eq. (1) assumes that the optical transitions of interest take place near the Γ point where the kinetic term E(k) is smaller than the lh-hh strain energy splitting (2δES) [33]. We do not expect any gain in the low strain regime where 2δES is small and the approximation less precise but the latter becomes more accurate towards the device operating tensile biaxial strain of about 2%. We assume parabolic DOS for all bands, where the L and Γ band effective masses, mL and mΓ, are given by 0.22 m0 and 0.038 m0 respectively [30], where m0 is the electron mass. Their hh and lh counterparts, mx are approximated by the DOS masses given by (mx//2mx,perpendicular)1/3 where x is lh or hh respectively. The assumption of parabolic bands here neglects the anisotropy and non-parabolicity induced by strain, which will affect the correctness of physical quantities such as the DOS or the inversion population threshold. More exact modeling of the curvature of the strained carrier dispersion will be the subject of future work. At the same time, the unstrained material parameters in our model, such as the valence band Luttinger parameters as well as the conductance masses should be replaced by their strained equivalents if they become experimentally available, since the parabolic model cannot determine these parameters as a function of strain. Tensile strain in Ge lifts the lh band above that of the hh band, resulting in stronger conduction band to lh transitions where TM-polarised (polarisation perpendicular to biaxial strain field) photons predominate [33]. The TM gain expression for the dichroic Ge medium can be derived from Eqs. (15) and (20) of reference [33] assuming negligible intraband relaxation time and parabolic dispersions for both the conductance and lh bands. Equations (14) and (15) in reference [33] also assume that ∆0 is large compared to δES. At a biaxial tensile strain of 2% in our Ge modeling, δES is approximately 0.11 eV compared to ∆0 = 0.3 eV. To further improve the accuracy of our modeling, the spin-orbit split-off state would have to be included in the density matrix formulation of [33]. Based on these assumptions, we arrive at the following expression,

g (E photon ) = (mr / mr − 0.2% strain )

3/ 2

( ( (E

× fc

mr mc

ABlh E photon − EgΓlh / E photon

) ( (E

photon − E gΓlh ) − f v

mr mlh

))

photon − E gΓlh )

(2)

where Ephoton is the photon energy, EgΓlh is the lh direct bandgap and (fc − fv) is the population inversion factor, which is a function of the carrier masses and Ephoton − EgΓlh. mr is the reduced mass given by (mlhmΓ)/(mlh + mΓ). The empirical constant A = 2.01x104 (eV)1/2cm−1 is derived from top down absorption experiments [2], [37] of 0.2% biaxial tensile strained Ge where the illuminating light polarization is parallel to the biaxial strain plane. We derive the reduced mass at 0.2% biaxial tensile strain, mr-0.2%strain based on the DOS mlh by Eq. (1) even though they give an overestimation of the actual values at low lh-hh strain separation 2δES. The computed gain value will consequently yield a conservative underestimation of the gain at the operating strain because of the smaller mr/mr-0.2%strain ratio. Blh = 0.54 is a correction factor reflecting the relative hh and lh absorption contributions as well as the TM/TE interaction strengths of the lh band. It is calculated using the unstrained Ge masses (which again leads to a conservative low estimation of the lh gain due to the small unstrained lh DOS) as well as the dipole moments given in Eqs. (14) #111065 - $15.00 USD

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and (15) of reference [33]. hh do not interact with TM photons and we neglect their influence in Eq. (2). Finally, the modal gain of the Ge-based waveguide can be calculated from the average of the gain in the active layer weighted by the local field intensity. As a simplification, we assume that the mode overlaps uniformly with all parts of the strained germanium layer and that the net gain is simply the average of all sublayer gain coefficients. 4. Discussion

The relative shifts in Fig. 2(a) of the conductance and valence bands with increasing in-plane biaxial tensile strain result in several consequences for light emission. Firstly, the conductance band Γ to L valley separation, ∆EΓL decreases and eventually disappears under a tensile strain of 2% when Ge becomes a direct material. Secondly, the valence band degeneracy is lifted and the lh optical transitions dominate that of the hh due to the smaller bandgap associated with the lh band. This allows for a more rapid change of the pseudo-Fermi levels towards inversion because of the smaller effective DOS. lh also have a stronger interaction with TM photons than TE photons and for subsequent discussion, we focus only on TM emission. hh, on the other hand does not contribute to TM gain at all. EcГ ΔEГL

(c) 3100 18

-3

-0.1 0.5 -0.2

0.4

Wavelength (nm)

0.6

Energy (eV)

0.7 0

0

18

-3

5x10 cm

-1000

-2000

-3000 18

2.0

Biaxial Tensile Strain (%)

-4000 0.45

0.55

2700

2500

2500

2000

2300

1500

2100

1000

1900

500

1700

0

1500

0.3 1.0

3000

max gain incl. free carrier loss

-3

1x10 cm

-0.3

max interband gain

2900

9x10 cm

0.0

3500 lambda @ max gain incl free carrier loss

Towards top surface

1000

0.8

0.1

Energy (eV)

(b)

0.9

Evlh EgГlh

Gain (/cm)

0.2

Gain (/cm)

(a)

0.65

Energy (eV)

0.75

-500

0.00

0.50

1.00

1.50

2.00

2.50

Biaxial Tensile Strain (%)

Fig. 2. (a) Strain dependencies of Ge direct conductance bandedge, the lh valence bandedge, the conductance band Γ to L valley separation as well as the lh direct bandgap, ECΓ, EVlh, ∆EΓL = ECΓECL and EgΓlh respectively with the biaxial tensile strain, where ECL is the L valley bandedge. ECΓ and EVlh are plotted with respect to their unstrained positions. (b) Interband gain profiles of Ge at 1.75% substrate surface biaxial tensile strain and various carrier injection concentrations. The colored lines show the average biaxial tensile strain within the active layer for each concentration. For the concentration 9x1018 cm−3, the gray lines represent the height-dependent biaxial tensile strain within each strain sublayer of our model. (c) Wavelength and gain variation of Ge with biaxial tensile strain, assuming an injection concentration of 9x1018 cm−3.

Figure 2(b) shows typical interband gain profiles of the 1.75% strained Ge layer based on our model at injection levels from 1 to 9x1018 cm−3. The larger Ge biaxial tensile strain further away from the Si substrate and the consequently smaller lh bandgap induces e-h pair migration to the Ge top surface, generating a stronger inversion and larger gain there. Since the active Ge layer is undoped, minority carrier injection from the adjacent p-doped regions is not suppressed unlike the case for the highly n-doped active layer proposed in reference [2]. In conjunction with SiGe cladding layers, a type I structure is also attainable for electron confinement. Figure 2(c) plots the maximum Ge optical gain coefficient with and without free carrier absorption (FCA) as well as the wavelength at maximum net gain against the biaxial tensile strain component at the same 9x1018 cm−3 injection concentration used in reference [2]. The gradual movement of the Γ point towards the bandgap center at higher tensile stress lowers the required electron concentration at inversion. Beyond a biaxial tensile strain of 1.5%, FCA is compensated by the interband gain even while Ge remains an indirect material. By varying the applied biaxial tensile strain between 1.5% and 2%, the peak wavelength is also tunable within a range of 400 nm up to a maximum wavelength of 2.6 µm. The net gain of direct bandgap Ge at 2% biaxial tensile strain can reach 1077 cm−1, compared to 400 cm−1 predicted in n-doped Ge

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[2]. However, in the latter case, a value of Blh of 0.318 [37] is applied, assuming unstrained masses as well as polarisation independent and identical average dipole moments for both hh and lh. Substituting for this value of Blh yields a weaker gain of 325 cm−1 in our case. 5. Conclusion

We modeled Ge emission characteristics under biaxial tensile strain generated by micromechanical structures. Our simulations indicate that transparency is attainable even in indirect Ge at a substrate biaxial tensile strain of 1.5% assuming a carrier injection concentration of 9x1018 cm−3. Larger biaxial tensile strains greater than 2% will transform the Ge into a direct bandgap material, resulting in an efficient mid-IR light emitter. A strain-tunable wavelength range of 400 nm also opens up the possibility of on-chip tunable light sources. Our technique can be further combined with n-doping of the Ge layer. Jifeng et al. [2] estimated that the net optical gain increases by a factor of 10 in 0.25% biaxial tensile strained Ge compared to unstrained Ge. By means of our proposed technique, we can substantially increase the gain of ndoped Ge on CMOS fabricated by deposition or bonding processes of Ge on wafer substrates, where the built-in strain is low.

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