GeSn lasers covering a wide wavelength range thanks to uniaxial

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GeSn lasers covering a wide wavelength range thanks to uniaxial tensile strain Jérémie Chrétien, Nicolas Pauc, Francesco Armand Pilon, Mathieu Bertrand, Quang-Minh Thai, Lara Casiez, Nicolas Bernier, Hugo Dansas, Patrice Gergaud, Eric Delamadeleine, Rami Khazaka, Hans Sigg, Jerome Faist, Alexei Chelnokov, Vincent Reboud, Jean-Michel Hartmann, and Vincent Calvo ACS Photonics, Just Accepted Manuscript • DOI: 10.1021/acsphotonics.9b00712 • Publication Date (Web): 27 Aug 2019 Downloaded from pubs.acs.org on August 30, 2019

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GeSn lasers covering a wide wavelength range thanks to uniaxial tensile strain Jérémie Chrétien1, Nicolas Pauc1, Francesco Armand Pilon3, 4, Mathieu Bertrand2, Quang-Minh Thai1, Lara Casiez2, Nicolas Bernier2, Hugo Dansas2, Patrice Gergaud2, Eric Delamadeleine1, Rami Khazaka2, Hans Sigg3, Jerome Faist4, Alexei Chelnokov2, Vincent Reboud2, Jean-Michel Hartmann2, Vincent Calvo1 1 Univ. Grenoble Alpes, CEA, IRIG-DEPHY, F-38000, Grenoble, France 2 Univ. Grenoble Alpes, CEA, LETI, F-38000, Grenoble, France 3 Laboratory for Micro- and Nanotechnology, Paul Scherrer Institut, 5232 Villigen PSI, Switzerland 4 Institute for Quantum Electronics, ETH Zürich, 8093 Zürich, Switzerland KEYWORDS. Germanium Tin laser, strain, tunability, group IV, nanophotonics

ABSTRACT. Silicon photonics continues to progress tremendously, both in near-infrared datacom/telecoms and in mid-IR optical sensing, despite the fact a monolithically integrated group-IV semiconductor laser is still missing. GeSn alloys are one of the most promising candidate materials to realize such devices, as robust lasing under optical pumping was demonstrated by several groups up to mild cryogenic temperatures. Ideally, the integrated lasers should be tunable by design over a wide spectral range, offering a versatility which is required for optical sensing devices. We present here an innovative approach taking advantage of local strain management in

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the semiconductor laser’s active zone. Arrays of differently strained Fabry-Pérot GeSn microlasers were fabricated side-by-side on the very same chip after blanket epitaxy on a Ge-buffered Silicon-On-Insulator substrate. Thanks to the local strain design, laser emission over a very large wavelength range under optical pumping, with laser lines peaking from 3.1 up to 4.6 µm at 25K, and with thresholds lower than 10 kW.cm-2. Laser operation persists up to 273 K, i.e. very close to room temperature. This strategy, implemented on group IV semiconductors, opens up a new route to control the emission properties of micro-lasers integrated on a chip over significant photon energy windows representing a significant step forward in the integration and miniaturization of light sources emitting at a process defined wavelength.

Strain has for long been used in III-V direct band gap structures to modify the laser emission wavelength and lower the lasing threshold. Fabrication of the active stacks occurs through a layerby-layer epitaxial growth onto a host substrate, for instance GaAs or InP. Control of strain in the active layers is then restricted by notions of critical thickness for plastic relaxation (for individual layers such as quantum wells or for whole stacks). The operating regime of strained semiconductor lasers is defined by the heterostructures used and emission wavelength tunability is, generally, small. More recently, strain induction from relaxed “strain reservoirs” into adjacent constrictions (such as bridges) has proven to be a very efficient tool to control the strain state in “bulk” semiconductor devices. In particular, this strategy was successfully used on Germanium On Insulator (GeOI) films [1], with very high uniaxial [2-6] or biaxial strains [6-8] achieved in several µm long suspended bridges or crosses with sub-micrometer cross-sections. Interest in doing so is driven by the peculiar feature of strained Ge, which transits from an indirect to a direct fundamental band gap as the tensile strain increases, but also by the significant strained volumes, able of sustaining optical modes. This potentially leads to optical resonators built on “bulk”, highly


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strained materials [9-11]. However, the Ge band gap is close to a direct one only at the highest achievable strains. In parallel, with the recent advent of high quality Ge1-xSnx alloys, laser operation in groupIV semiconductors is now routinely reported in the litterature [12-21]. Indeed, this semiconductor is direct for low tin contents and residual compressive strains. In as-grown stacks, the emission wavelength is mainly controlled by the tin content, and to a lesser extent by the residual compressive strain in the active layers (as in other semiconductor lasers) [22, 23]. Here we explore a new route which combines direct group-IV semiconductor with strain redistribution, tunable by a simple membrane design, to control the semiconductor band gap [24, 25] and finally the emission wavelength. In this paper, we describe the fabrication and optical characteristics of tensilely strained and free standing GeSn heterostructure microbridges placed or not between two broadband mirrors offering optical feedback. Strain values were extracted by comparing experimental Photoluminescence (PL) data with a simple deformation potential model. Lasing was achieved in strained optical cavities, pumped either with a pulsed 1064 or 2650 nm laser, from 25 K up to 273 K, the highest GeSn lasing temperature to date, to the best of our knowledge.


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Figure 1. Structure and device description. (A) Schematic showing the principle of strain induction in micro-bridges after under-etching of the arms for tensile (𝜀 > 0) or compressively (𝜀 < 0) prestrained layers. (B) Cross-sectional schematics of the GeSn heterostructure with targeted tin contents and thicknesses. (C) Cross-sectional TEM image with EDX elemental Sn mapping. (D) Strain maps using nano-beam PED. The scale uses the Ge buffer lattice constant as a reference. (E) Schematic view of an optically pumped corner-cube cavity GeSn device with two differently strained neighboring cavities. (F) SEM image of a 150 µm arm length device. The mechanism behind the formation of tensile strained Ge bridges, as reported in Ref. (1), relies on the use of Ge stretching arms on both sides of a central, tiny Ge bridge. Indeed, there is a slight amount of tensile strain (~0.2 %) in Ge grown on Si or in GeOI films (owing to disprepancies in terms of thermal dilation coefficients between Ge and Si, which come into play when cooling down the Ge Strain-Relaxed Buffer on Si after epitaxial growth). Arms release and contraction induced by 𝑆𝑖𝑂2 underetching will longitudinally stretch the bridge (see Fig. 1 A, top). In opposition to Ge epilayers, GeSn alloys are radically different since the residual strain in these


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layers is compressive, as Sn is a much bigger atom than Ge. Releasing a GeSn structure with the same architecture than in [1] would induce bridge squeezing by the arms and structure bowing, which would be harmful from an optical point of view, as the bandgap would become less direct (see Fig. 1 A, bottom). Fig. 1 B shows the starting stack of our samples. A 5 µm thick Ge strain-relaxed buffer was first of all grown in two steps onto a thinned-down Silicon On Insulator (SOI) substrate (the interest of using a SOI substrate instead of a Si substrate will be clear in the following). We grew on top a step-graded GeSn epixatial stack with 8%-10%-13%-16%-13% Sn concentration steps to mitigate Sn segregation and confine misfit dislocations at Sn concentration steps, reducing thereby the threading dislocation density resulting from the high lattice mismatch [26, 27]. The whole heterostructure was encapsulated by two SiGeSn electronic barriers to minimize surface recombinaison and improve confinement of electronic carriers in the optically active GeSn 16% layer. The recipe used to grow that stack was the very same used in Ref. [18] for our stateof-the-art suspended micro-disk laser lasing up to 230 K (see S1 for growth details). The Transmission Electron Microscopy (TEM) image shown in Fig. 1C shows that the GeSn 16% active layer is 415 nm thick. Strain maps acquired thanks to nanobeam Precession Electron Diffraction (PED) are shown in Fig. 1 D. The residual biaxial compressive strain (from PED) in the GeSn 16% active layer is -0.66%, with a 70 % macroscopic degree of strain relaxation (see S2). This peculiar stack enabled us to adapt the structure shown in Fig. 1 A to have, in the end tensile strain in constrictions made from an GeSn epilayer which was initially compressively strained. To do so, we took advantage of the 5 µm thick underlying Ge buffer, which was slightly tensilestrained as it was grown on a thinned-down SOI substrate (same levels of strain in that Ge-On-


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SOI stack than in GeOI substrates). The structure basically consists in a tensile strained central GeSn resonator, suspended in the air, which is obtained after an under-etching of the SOI’s Buried Oxide (BOX) and thus a release of the thick Ge stretching arms. Fig. 1 E shows a schematic of a strained device and Fig. 1 F is a Scanning Electron Microscopy (SEM) view of the final device, where we underlined with a dashed line the under-etching front in the Ge layer. Under-etching was stopped below the GeSn pad and close to its longitudinal bondaries. This results in (i) an anchoring of the GeSn bridge to the Ge arms on both sides of these lines and (ii) an injection of tensile strain in the GeSn bridge (from Ge arm contraction). Such a process results in a local inversion of strain without any need for external stressors or flip chip approaches. The detailed fabrication process is described in S3. GeSn bridges are 8 µm long, 1.5 µm wide and 1 µm thick, while the arm width is 40 µm. Strain tuning in the bridge area is then controlled by changing the total device length L. The oscillator is described in Refs. [28,29]; it is made of a bridge cavity with cornercube mirrors at each extremity. Each mirror is built from two face to face and intersecting parabolic GeSn-air dioptres. Finite Difference Time Domain (FDTD) simulations give a modal reflectivity around 93 % (see S4) for the fundamental mode in the guide at 3.5 µm, corresponding to approximately 35 cm-1 mirror losses. This high degree of reflectivity arises from the total internal reflection mechanism of the incoming wave at the air-semiconductor interface and ensures a robust insensitivity of the mirror reflectivity to wavelength changes. An amazingly broad range of high reflectivity (>91%) is indeed obtained from simulations over the 2.5 µm to the 5 µm wavelength range (see Fig. S6). This design is particularly relevant and adapted to strain-tuned lasers owing to its ease of integration, as no adapted design is necessary to fit the emission wavelength.


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Figure 2. CW characterisation. Photoluminescence spectra under 1047 nm CW optical pumping of GeSn heterostructures for different arm lengths without (A) and with mirrors (B) at 25K for a 𝑃 = 1.15𝑘𝑊.𝑐𝑚 ―2 pumping power. Spectra coming from a full stack (i.e. without any processing, referred to as “Blanket Layer”) and a fully relaxed structure (i.e. a broken cavity) are plotted at the bottom. Arrows point out the line cutoff used for the estimate of bandgap shift. The inset is an enlarged view of the spectrum showing cavity modes.

Figure 2 (A) and (B) show PL spectra obtained on structures without or with cornercube mirrors for increasing arm lengths (see S5 for PL setup details). They were recorded at 25 K with a 1047 nm Continuous Wave (CW) pump laser (𝑃 = 1.15𝑘𝑊.𝑐𝑚 ―2). In both cases, two main


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contributions can clearly be seen (the bump at ~4200 nm corresponds to the CO2 absorption lines). The first one is fixed at 0.42 eV, whatever the device length. This line peaks at the same energy than the PL signal coming from broken bridges on the same chip, i.e. made of free standing and relaxed material. This contribution is assigned to light emission from the unstrained parts of the samples, located outside the 8 µm long waveguide, but inside the 12.5 µm pumping spot. This is in agreement with the expected emission from a relaxed GeSn 16% layer as already shown in [17, 30] at 25 K, indicating that recombination takes place in the highest Sn content layer. Note that the signal from the epilayer (referred to as “Blanket Layer” in Fig. 2 A) is slightly blue shifted compared with the signal from the broken bridges (called “Relaxed” in Fig. 2 A), due to the slight residual compressive strain in the GeSn stack after epitaxy. The other line redshifts as the arm length increases, in line with a strain induced band gap shrinking, as observed for Ge bridges in [2] and [3]. This qualitatively shows that our design is efficient to invert and control strain in the central zone of the resonator and potentially achieve gain curve energetic shifting. Adding mirrors (see Fig. 2 B) to this structure leads to similar trends. However, one can note a set of distinct features compared to the mirror-free devices: (i) due to mirror induced mechanical softening of the arm ends, longer arms are necessary to have similar strain levels in the waveguide (see for instance the L=150 µm devices, peaking above (resp. below) the CO2 absorption line for mirror (resp. mirror-less) structures); (ii) a peak broadening is observed in the redshifted emission of optical cavities; it is likely due to a tensile gradient in the mirror region, as a result of the mechanical softening of the arm evidenced in (i) and (iii) emerging modes are visible and start to modulate the signal from the bridges, unambiguously showing longitudinal modes of oscillation. No functional devices with L>175 µm (resp. 250 µm) were found for mirror-less (resp. mirror)


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samples, as longer structures presented cracks in high shear strain regions, i.e. close to the waveguide end or mirror edges. An in-depth quantitative study of the strain distribution and intensity would require a dedicated study and falls outside the scope of this paper. Indeed, Raman spectroscopy, when used with its conventional visible wavelength pump, gives a picture of the crystal deformation within the first tens of nm below the surface, which is of little use to determine the strain state in the buried GeSn 16 % active layer. Instead, we will use here a simple model aiming at relating a spectroscopic shift with the uniaxial strain in the bridge. It is based on the deformation potential theory with scaled deformation potentials. For uniaxial strain, the energetic shift of conduction and valence (heavy HH and light LH holes) band edges can be determined analytically using the following formula [31, 32]: 𝛿𝐸𝑐 = 𝑎𝑐,𝛤𝑇𝑟(𝜺) (1) 𝛿𝐸𝑣,𝐻𝐻 = 𝑎𝑣𝑇𝑟(𝜺) + 𝛿𝐸𝑣,𝐿𝐻 = 𝑎𝑣𝑇𝑟(𝜺) ―

Δ0 6

+

𝛿𝐸100 4

Δ0 3

1

― 2𝛿𝐸100 (2)

1

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+ 2 Δ20 + Δ0𝛿𝐸100 + 4𝛿𝐸2100 (3)

where 𝜺 is the strain tensor, depending only on the longitudinal deformation 𝜀0, and 𝛿𝐸100 𝑆12

= 2𝑏(1 ― 𝑆11)𝜀0. Parameter values are given in S6 and Table S8. Tensile strain 𝜀0 along lifts the valence band degeneracy at the Γ point and results in lower energy (resp. higher energy) Γ-HH (resp. Γ-LH) band gaps. Band to band recombination gives rise to photon emission starting from the semiconductor Band Gap edge “Renormalized” (i. e. reduced) by an energetic amount ΔEBGR depending on the free carrier concentration. Since data in Figs. 2 A and B are collected for constant free carrier generation (and thus carrier concentration), and by neglecting the effect of strain induced band distortions on ΔEBGR, we will arbitrarily take the low energy cutoff at 20 % of peak maximum on the redshifted


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spectra as a criterion to point out the energetic shift with respect to the relaxed sample (black arrows in Fig. 2). Two close components (splitting of the order of 50 meV) can be clearly distinguished in low strain samples (L=50 µm for mirror less samples and L=75 µm for resonators). Assuming that, upon strain induction, the strain change is the same in the 16% and 13% layers, and neglecting the very small differences between the deformation potentials of these layers, we will consider the band offsets as being constant between the well and the barriers as the strain level is changed in the stack. As seen in Fig. 2 A and B, when the tensile strain increases, the redshifted PL spectra become purer, suggesting that emission in low strain samples does not originate from the 13% barriers. Otherwise, a two lines spectrum would still be present for higher strains. Instead, and taking equations (1) to (3), we find a strain induced HH to LH gap splitting of the order of 50 meV for the L=50 µm mirror less sample. This is a value similar to the experimental peak splitting, suggesting that these two lines arise from recombination on HH and LH states in the 16% tin content layer due to hole band filling of the HH to LH gap. Increasing the strain splits even more these gaps and shrinks hole population in the LH band, favoring vertical transitions with the HH band only.


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Figure 3. Strain analysis. Line red-shifting from CW photoluminescence measurements and estimated longitudinal strain in the bridge associated to this redshift for HH band edge as a function of device length L.

Figure 3 summarizes the estimated band gap shifts and associated strain using equations (1) to (3) in the waveguide as a function of L, based on the experimental energetic shift given by the relative positions of the arrows (see in Fig. 2) between released and strained bridge spectra. Compared with the band gap of the released bridge, there is a 40% relative shrink of the band gap in the most strained bridge. As seen above, resonators are less strained than mirror–free structures for the same arm geometry, with typically a 20-30 % strain loss for high strains in cavities. The estimated breaking limit is about 2.2 % for this set of samples.


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Cavities were also optically pumped at 25 K by an Optical Parametric Oscillator (OPO) laser emitting at 2650 nm, with a repetition rate of 82 MHz and a pulse duration of 0.1 ns. The OPO laser source enabled to optically pump GeSn closer to its band gap, helping in minimizing heating in these free-standing structures, while yielding transient and high carrier injection levels, in opposition to CW pumping. Fig. 4 A shows a Light in-Light out curve measured at 25K, on a L=150µm device. A slope change is observed, characteristic of laser emission, with a threshold as low as 9 kW/cm-2 and an intensity dynamic range reaching 4 decades. Thresholds measured on other devices (having different L) span the 9 to 12 kW.cm-2 range at this temperature (see S7). These thresholds are very small for bulk GeSn lasers and compete with thresholds obtained on GeSn/SiGeSn quantum well lasers, i. e. 35 𝑘𝑊.𝑐𝑚 ―2 at 20K with a 1064 nm pumping wavelength [21]. Finding similar thresholds for different strain levels can be surprising at first glance since the directness D (𝐷 = 𝐸L𝑔 ― 𝐸Γ𝑔) is expected to increase with the tensile uniaxial strain, removing free electrons from the L valleys and reducing material losses. Indeed, and taking the values found for Ge [3], we find 𝛿𝐷 (𝑚𝑒𝑉) ≈ 31 𝜀0(%) as strained induced D change for the uniaxial case. Considering that this estimate does not deviate too much for the alloy under study, we find 𝛿𝐷 ≈ 60 𝑚𝑒𝑉 as gain in D for the highest strains estimated in our study. This has to be compared with the “compositional” directness in the unstrained material, calculated to be approximately 150 meV-190 meV at 16 % [17, 32], which offers a very high energy barrier to confine electrons in the Γ valley even without “strain assistance”. This situation would not clearly hold for lower tin content samples, where the strain is expected to have a deep impact on the electron distribution in the L and Γ conduction valleys. No clear relationship (see S7) between these low lasing thresholds and applied tensile strain can be made, as many parameters come into play such as strain changes,


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slight modifications in the free carrier absorption, mode confinement, mirror reflectivity, or variability in the device geometry and roughness due to process conditions.

Figure 4. Lasing in GeSn cornercube cavities. (A) Light in-Light out curve at 25 K for a L=150 µm cavity with a lasing threshold at 9 kW.cm-2; (B) Spectra at the same temperature (25 K) and laser pumping conditions (178 𝑘𝑊.𝑐𝑚 ―2, pumping wavelength 2650 nm) for different length cavities.


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Fig. 4 B shows the entire collection of normalized laser spectra recorded on different length cavities on the very same chip, with therefore the same tin content in the GeSn active layer. Laser emission is observed for all resonators. It covers the 3.1 to 4.6 µm range, with a Dirac comb like emission pattern. This is in agreement with the strain induced band gap shrinking observed under CW excitation. It is assigned to recombination in the lowest energy states in the GeSn 16% layer. Compared with spectra recorded in the CW regime, the Signal to Noise ratio is greatly improved because of the high signal dynamic range in the laser emission regime. The Free Spectral Range is constant and measured to be 6.25 meV (or 1.51 THz). This is the signature of Fabry Pérot type longitudinal resonances having different mode indexes, within the same fundamental mode, as no other wavelength shifted Dirac comb can be detected. Taking as an effective group index ng=3.8 in the whole cavity (see supporting information), we find a round trip cavity length of 53 µm, in reasonable agreement with the geometric mirror to mirror spacing of 20 µm measured between the intersecting points of the two parabola. Note that a thermal redshift of the resonance wavelength was only observed for the highest pumping powers (not shown).


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Figure 5. (A) Mode competition near threshold for a L=150 µm device at 130 K ; (B) LinLout curve of this bridge at the same temperature showing a threshold of 46kW.𝑐𝑚 ―2 Pumping is performed with the OPO laser at λp=2650 nm.

Figure 5 A shows photoluminescence spectra obtained at 130 K on the 150 µm long bridge, in the threshold region. A series of sharp peaks emerges at low pumping power, with a relative peak weight change as the pumping power becomes stronger. This behavior is characteristic of mode competition and reveals laser emission at energies given by the strain shifted band gap of the semiconductor. The associated Light in- Light out curve is shown in Fig. 5 B, giving a threshold of 46 kW.𝑐𝑚 ―2, which is higher than the typical thresholds of 10 kW.cm-2 found at 25 K (Fig. 4A). Increasing the sample temperature up to 260 K leads to a narrowing of the total emission window, as shown in Fig. 6 A. PL spectra were obtained at this temperature with a Nd:YAG laser emitting at 1064 nm, with a repetition rate of 50 kHz and a pulse duration of 0.6 ns with a 20 µm spot size. This configuration provides a much smaller duty cycle than with the OPO. The mean power transmitted to the oscillator and corollary emission flux are then much smaller; it is nevertheless a good configuration in order to maintain high electron hole densities and limit sample heating despite the shorter excitation wavelength.


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Figure 6 (A) Spectra at 260 K and 3 MW.cm-2 for the same cavities showing a narrowing of the emission window down to 281 nm; (B) Laser spectrum recorded at 273 K (power density 4 MW.cm-2) on a L=75 µm cavity, together with the Light in-Light out curve.


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The arm stretching force is fixed not only by design rules, but also depends on the material thermal expansion coefficients which come into play during any temperature change. A temperature increase from 25 K up to 260 K results in Ge arm dilatation, reducing the stretching force exerted on the cavity and the induced tensile strain. A 281 nm shift in the laser emission (at a power density of 3 MW.cm-2) between the L=50 and L=250 µm devices is obtained at 260K, showing that laser tunability is still significant close to room temperature, despite a significant emission narrowing. This extremely high tunability of our group IV laser compares well with its III-V or II-VI compounds counterparts, with emission windows of 60-202 nm for the same blanket layers [34-37]. Clear lasing operation is shown at 273 K, i. e. very close to room temperature, on a L=75µm resonator, as shown in Fig. 6 B. Spectra below and above threshold are shown in Fig. 6 B, together with the Light in-Light out curve giving a~2 MW.cm-2 threshold. This is the highest operation temperature reported so far in any GeSn lasers [38]. The limitation of the operating temperature might have different origins such as thermal activation of traps within the optically active medium itself. Compared with the low temperature case, the loss of D due to lower strains at elevated temperatures might, together with thermal energy, result in the injection of a small fraction of the electrons in L states, contributing to the killing of lasing at high temperatures. Variability in the device geometry or other losses such as free carrier absorption or a smaller mirror reflectivity (as discussed above for the low temperature thresholds) could explain why the longest arm device does not have the highest lasing temperature. In conclusion, we have demonstrated laser operation over a very high wavelength window. The emission wavelength was controlled at will using strain management in the semiconductor active zone and simple design rules, on a pure group IV semiconductor platform. The influence of


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strain on laser performances has still to be investigated in details. Nevertheless, the highest lasing temperature was 273 K, which is to date the highest operation temperature in any GeSn semiconductor lasers, despite a not fully optimized thermal management in our free standing structure. At cryogenic temperatures we report extremely low thresholds for a GeSn laser (in the 9 to 12 kW.cm-2 range), and a very high wavelength tunability for a semiconductor laser fabricated with the same starting material. Such an approach could be implemented on other semiconducting stacks, opening the way to widely tunable lasers in the visible or infrared range.

ASSOCIATED CONTENT Supporting Information. Materials growth, Determination of the residual compressive strain in GeSn layers, Device fabrication, Finite Difference Time Domain simulation of mirror reflectivity, Photoluminescence, Elastic constants and deformation potentials for strain computation, Summary of Lin – Lout curves for different arm length long devices, Parameters table. (PDF) AUTHOR INFORMATION Corresponding Author *E-mail : [email protected] *E-mail : [email protected] ACKNOWLEDGMENT This work was supported by Elegante ANR project and CEA Carnot project. REFERENCES


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