In-Situ Annealing and Hydrogen Irradiation of Defect-Enhanced Germanium Quantum Dot Light Sources on Silicon

: While light-emitting nanostructures composed of group-IV materials fulﬁl the mandatory compatibility with CMOS-fabrication methods, factors such as the structural stability of the nanostructures upon thermal annealing, and the ensuing photoluminescence (PL) emission properties, are of key relevance. In addition, the possibility of improving the PL e ﬃ ciency by suitable post-growth treatments, such as hydrogen irradiation, is important too. We address these issues for self-assembled Ge quantum dots (QDs) that are co-implanted with Ge ions during their epitaxial growth. The presence of defects introduced by the impinging Ge ions results in pronounced PL-emission at telecom wavelengths up to room temperature (RT) and above. This approach allows us to overcome the severe limitations of light generation in the indirect-band-gap group-IV materials. By performing in-situ annealing, we demonstrate a high PL-stability of the defect-enhanced QD (DEQD) system against thermal treatment up to 600 ◦ C for at least 2 h, even though the Ge QDs are structurally a ﬀ ected by Si / Ge intermixing via bulk di ﬀ usion. The latter, in turn, allows for emission tuning of the DEQDs over the entire telecom wavelength range from 1.3 µ m to 1.55 µ m. Two quenching mechanisms for light-emission are discussed; ﬁrst, luminescence quenching at high PL recording temperatures, associated with the thermal escape of holes to the surrounding wetting layer; and second, annealing-induced PL-quenching at annealing temperatures > 650 ◦ C, which is associated with a migration of the defect complex out of the QD. We show that low-energy ex-situ proton irradiation into the Si matrix further improves the light emission properties of the DEQDs, whereas proton irradiation-related optically active G-centers do not a ﬀ ect the room temperature luminescence properties of DEQDs. modelling: T.F. and J.S., draft preparation: M.B., L.S., and J.A.; writing—review and editing: L.S., J.A., A.P., J.H., H.G., T.F., F.S., and M.B.; supervision: and


Introduction
Practical, Si-compatible, monolithically-integrated lasers are the main missing ingredient towards the realization of efficient and low-cost integrated silicon photonics circuits that are designated to uplift fields of on-chip data communication [1][2][3] and on-chip sensing [4]. There has been substantial

Materials and Methods
The DEQD samples were grown in a SIVA-45 molecular beam epitaxy (MBE) system (RIBER, Paris, France) on high-resistivity (>5000 Ωcm) Si(001) substrates. After a standard substrate cleaning process, including Piranha solution and RCA cleaning [24], the samples were dipped for 30 s in diluted hydrofluoric acid (HF 1%), to remove the native oxide before their introduction into the load lock vacuum chamber. The samples were degassed at 720 • C for 15 min before the growth of a crystalline Si buffer layer of 50 nm thickness at a growth temperature (T G ) of 500 • C. Note that this high degassing temperature is not mandatory for DEQD formation, and that, in principle, any T G higher than~400 • C can be employed, above which the hydrogen termination of the HF-dipped Si surfaces starts to desorb [25]. Thus, the degassing process can be readily adapted to ensure CMOS-compatibility.
For the self-organized, strain-driven growth of QDs, 7.3 Å of Ge were deposited at T G = 500 • C. This leads to the formation of QDs in "hut-cluster" shape [26,27] on a Ge-rich wetting layer (WL) [28]. Hut clusters are elongated pyramidal structures terminated by {105}-oriented crystal facets [26]. For the given growth conditions, the typical QD height and density are~1.4 nm to 2 nm, and 2 × 10 10 cm −2 , respectively. During the entire QD deposition process, Ge ions with a dose of about 10 4 µm −2 and an energy of 1.5 keV were impinging onto the sample surface. This low energy leads to implantation depths of roughly 1-2 nm for the positively charged Ge ions which, in turn, corresponds to the average height of the Ge hut clusters. Finally, the implanted Ge QDs were capped with a Si matrix of 70 nm thickness, which was deposited during a ramp-up of T G from 475 • C to 500 • C. Hereafter, the samples were annealed in-situ for 2 h at six different temperatures T A , ranging from 500 • C up to 675 • C. In addition, a reference sample was grown with the same layer sequence, but without the annealing step.
Each sample was cut into four parts. Hydrogen passivation experiments with a low-energy (100 eV) Kaufman source [29][30][31][32] were performed on three of these. To enhance hydrogen diffusion, the sample temperature during hydrogen treatment was always kept at 300 • C. The three pieces from each sample were exposed to proton irradiation doses of 5 × 10 17 cm −2 , 1 × 10 18 cm −2 , and 2.5 × 10 18 cm −2 , respectively.
For the spectroscopic investigation, the samples were placed in a helium flow cryostat with a heatable cold finger that allows sample temperatures (T PL ) between 4 K and 350 K. The samples were excited using a fiber-coupled continuous-wave diode laser with an emission wavelength of 442 nm, and maximum output power of up to 3 mW. A microscope objective with a numerical aperture of 0.7 focused the laser to the sample surface, resulting in a laser spot diameter of~2 µm. The resulting PL signal was collected by the same objective, fiber-coupled to a grating spectrometer and detected by a liquid nitrogen-cooled InGaAs line detector (Princeton Instruments OMA V, Trenton, NJ, USA), with a cut-off wavelength of 1580 nm.

DEQD Nanostructures
The present understanding of the formation process of DEQDs and the reason for their superior optical properties are as follows: The implantation of Ge ions of low energies (<2 keV) during the QD growth leads to a partial amorphization of the QD's crystal lattice. The subsequent solid-phase recrystallization at growth temperature has to remain incomplete, since an additional Ge atom was brought into a nanostructure which had perfect crystallinity before the ion implantation process. According to theoretical modelling, recrystallizing the structure results in the formation of a split-[110] self-interstitial [18] surrounded by a distorted Ge QD crystal lattice region consisting of about 45 atoms. These results suggest that it is this interstitial defect and the surrounding distorted crystal that are responsible for the enhanced light-emission properties of DEQDs at room temperature. The defect system introduces electronic states at the Γ-point of Ge, from where electrons directly recombine with holes confined by the large band offsets of the Si-Ge-Si QD heterostructure. Electrons tunnel from the surrounding Si matrix into these states, and this tunneling process is independent of temperature.
Therefore, we previously assumed [18] that the only mechanism leading to an eventual PL quenching at temperatures above~300 K is caused by QD ionization, i.e., thermal escape of holes from the QD's valence band potential well, induced by the band offsets of the Si-Ge-Si heterostructure to the Si matrix. However, no proof has been given so far for the assumption that hole-escape to the Si matrix is indeed the major mechanism for PL decay at high temperatures. As will be shown, this picture is indeed not fully adequate.

In-Situ Annealing of DEQDs
To study the influence of thermal hole-escape on the PL decay and the activation energy barrier for defect migration E A(migr) , the same growth conditions for the Si matrix and the DEQDs were applied for all in-situ annealed samples. The annealing time was 2 h in all cases, and the only parameter varied Crystals 2020, 10, 351 4 of 14 was the annealing temperature T A , which was systematically changed within the range from 500 • C to 675 • C. Figure 1 depicts PL spectra for the annealed samples, as well as for a pristine reference sample (grey) recorded at temperatures of T PL = 10 K (Figure 1a) and T PL = 300 K (Figure 1b). For low T PL , and low T A of 500 • C, we observe that sample annealing for 2 h leads to an increase of the integrated PL intensity, I PL , of the DEQD signal, but does not significantly change the spectral shape of the PL signal. Increasing T A results initially in a further improvement of I PL , followed by saturation in the T A range from 600 • C to 650 • C. However, a further slight increase of T A from 650 • C to 675 • C leads to almost complete quenching of the PL intensity.

In-Situ Annealing of DEQDs
To study the influence of thermal hole-escape on the PL decay and the activation energy barrier for defect migration EA(migr), the same growth conditions for the Si matrix and the DEQDs were applied for all in-situ annealed samples. The annealing time was 2 h in all cases, and the only parameter varied was the annealing temperature TA, which was systematically changed within the range from 500 °C to 675 °C. Figure 1 depicts PL spectra for the annealed samples, as well as for a pristine reference sample (grey) recorded at temperatures of TPL = 10 K ( Figure 1a) and TPL = 300 K (Figure 1b). For low TPL, and low TA of 500 °C, we observe that sample annealing for 2 h leads to an increase of the integrated PL intensity, IPL, of the DEQD signal, but does not significantly change the spectral shape of the PL signal. Increasing TA results initially in a further improvement of IPL, followed by saturation in the TA range from 600 °C to 650 °C. However, a further slight increase of TA from 650 °C to 675 °C leads to almost complete quenching of the PL intensity.
The second obvious result of Figure 1a is a strong blueshift of the PL peak wavelengths (energies) from ~1520 nm (0.816 eV) for TA = 500 °C to 1330 nm (0.93 eV) for TA = 650 °C. At room temperature (TPL = 300 K, Figure 1b), this blueshift is present as well. The peak energy shift of the PL signal with increasing TA is plotted in Figure 2b, for TPL = 10 K and TPL = 300 K. As already seen in the PL spectra in Figure 1, the PL shifts for both TPL monotonously to higher energies with increasing TA. Note that all peak maxima of the spectra recorded at TPL = 300 K are also offset to the blue by 9.8-17.1 meV, as compared to the maxima observed at TPL = 10 K.
In contrast to the spectra recorded at TPL = 10 K, the PL signal measured at TPL = 300 K shows the onset of quenching already at TA ≥ 600 °C. The obtained values for IPL versus TA are plotted in Figure  2a for TPL = 10 K and TPL = 300 K, respectively. In a previous publication [18] we have argued, just on the basis of room-temperature PL measurements, that the decay of the DEQD PL with high TA is likely related to thermally-induced migration of the defect complex from the QD into the Si matrix. There, we extracted an activation energy EA(migr) of 3.4 eV for this process, which is indicated in Figure   Figure 1. Photoluminescence spectra of defect-enhanced quantum dots (DEQDs) samples recorded at T PL = 10 K (a) and room temperature (b). The DEQD samples were in-situ annealed for 2 h at temperatures T A of 500 • C (violet), 550 • C (dark blue), 600 • C (light blue), 625 • C (green), 650 • C (orange), and 675 • C (red). The spectrum of the pristine reference sample is depicted in grey.
The second obvious result of Figure 1a is a strong blueshift of the PL peak wavelengths (energies) from~1520 nm (0.816 eV) for T A = 500 • C to 1330 nm (0.93 eV) for T A = 650 • C. At room temperature (T PL = 300 K, Figure 1b), this blueshift is present as well. The peak energy shift of the PL signal with increasing T A is plotted in Figure 2b, for T PL = 10 K and T PL = 300 K. As already seen in the PL spectra in Figure 1, the PL shifts for both T PL monotonously to higher energies with increasing T A . Note that all peak maxima of the spectra recorded at T PL = 300 K are also offset to the blue by 9.8-17.1 meV, as compared to the maxima observed at T PL = 10 K.
In contrast to the spectra recorded at T PL = 10 K, the PL signal measured at T PL = 300 K shows the onset of quenching already at T A ≥ 600 • C. The obtained values for I PL versus T A are plotted in Figure 2a for T PL = 10 K and T PL = 300 K, respectively. In a previous publication [18] we have argued, just on the basis of room-temperature PL measurements, that the decay of the DEQD PL with high T A is likely related to thermally-induced migration of the defect complex from the QD into the Si matrix. There, we extracted an activation energy E A(migr) of 3.4 eV for this process, which is indicated in Figure 2a by the Arrhenius fit to I PL (T A ). For the data recorded at TPL = 10 K, the same Arrhenius fit cannot be performed, as the fit does not converge. However, it is evident already, from the comparison between the low and high TPL spectra in Figure 1, that a significant difference in EA(migr) has to exist. In the next paragraph, we argue that the EA(migr) value extracted from the PL intensities at TPL = 300 K presents a lower limit to describe the migration of the defect out of the DEQD structure. In addition, the PL signal is affected by spurious PL quenching, due to a reduction of the hole binding energy that is caused by intermixing of the Si/Ge/Si interfaces upon annealing. Still, the observed large lower limit of EA(migr) = 3.4 eV rationalizes the resilience against thermal budgets of the DEQDs, that allows, for example, DEQD annealing at 600 °C for 2 h without loss of PL yield. Moreover, it was shown recently that DEQDs withstand thermal annealing, even for temperatures up to 800 °C, during 3-20 ms long flash-lamp annealing processes that are suitable for dopant activation [33], again demonstrating the robustness against thermal treatment of this nanosystem.
To gain further insight into the optical response of DEQDs, we recorded PL spectra in the TPL range between 10 K and 300 K, as shown in Figure 3. The quenching of IPL with increasing TPL beyond 140 K shows Arrhenius-type behavior. From these curves, we extracted activation energies of EA(ref) = 255 meV, EA(600) = 190 meV, EA(625) = 157 meV, and EA(650) = 124 meV for the reference sample and the samples annealed at TA = 600 °C, 625 °C, and 650 °C, respectively.
Since these activation energies for PL quenching decrease very strongly with increasing annealing temperature TA, it is immediately clear that the activation energy for defect migration EA(migr) cannot be extracted from the single experiment at TPL = 300 K reported in [18]. Evidently, quenching of the PL intensity at TPL = 300 K is affected by both defect migration and hole escape. For the data recorded at T PL = 10 K, the same Arrhenius fit cannot be performed, as the fit does not converge. However, it is evident already, from the comparison between the low and high T PL spectra in Figure 1, that a significant difference in E A(migr) has to exist. In the next paragraph, we argue that the E A(migr) value extracted from the PL intensities at T PL = 300 K presents a lower limit to describe the migration of the defect out of the DEQD structure. In addition, the PL signal is affected by spurious PL quenching, due to a reduction of the hole binding energy that is caused by intermixing of the Si/Ge/Si interfaces upon annealing. Still, the observed large lower limit of E A(migr) = 3.4 eV rationalizes the resilience against thermal budgets of the DEQDs, that allows, for example, DEQD annealing at 600 • C for 2 h without loss of PL yield. Moreover, it was shown recently that DEQDs withstand thermal annealing, even for temperatures up to 800 • C, during 3-20 ms long flash-lamp annealing processes that are suitable for dopant activation [33], again demonstrating the robustness against thermal treatment of this nanosystem.
To gain further insight into the optical response of DEQDs, we recorded PL spectra in the T PL range between 10 K and 300 K, as shown in Figure 3. The quenching of I PL with increasing T PL beyond 140 K shows Arrhenius-type behavior. From these curves, we extracted activation energies of E A(ref) = 255 meV, E A(600) = 190 meV, E A(625) = 157 meV, and E A(650) = 124 meV for the reference sample and the samples annealed at T A = 600 • C, 625 • C, and 650 • C, respectively.
Since these activation energies for PL quenching decrease very strongly with increasing annealing temperature T A , it is immediately clear that the activation energy for defect migration E A(migr) cannot be extracted from the single experiment at T PL = 300 K reported in [18]. Evidently, quenching of the PL intensity at T PL = 300 K is affected by both defect migration and hole escape.   To assess the structural changes induced by two hours of annealing at T A , we started from the experimental values of the PL peak positions at T PL = 10 K for different T A . For the non-annealed sample (grey spectrum in Figure 1a), we assume for simplicity a box-like Ge composition profile with 100% Ge content along the growth direction for both the QD and the WL. This should be reasonably close to reality, since at the low employed T G = 500 • C, trench-formation concomitant with Si-Ge intermixing is suppressed [34]. Slight differences in the Ge content, e.g., between 90% and 100%, are in general difficult to determine experimentally. Also, asymmetric smearing out of the initial Ge box profile induced by surface segregation can be assumed at the employed capping layer growth temperature of 475 • C [35]. However, this effect dominates the optical properties only for much thinner nanostructures such as very thin wetting layers (<3 ML thickness), as only the topmost monolayers of the nanostructure are affected by surface segregation at this growth temperature [35]. We justify the approximation of the box-like profiles a posteriori, as the bulk-diffusion induced blurring of the Ge composition is affecting the QD and WL potential much more significantly.
For the different T A , we assume a degradation of the Ge profiles through bulk diffusion according to the model given in [36], which was already successfully applied for Ge/Si nanostructures and heterointerfaces [37]. Also, we assume that the initial state of the PL transition, i.e., the split-[110]-interstitial in the QDs, virtually does not shift in energy as a consequence of Ge diffusion. This assumption appears to be reasonable, as this defect probes the local Ge environment. Especially in the center of the QD, where the defect is most effective, this concentration does not decay below~80% of its initial value and, thus the corresponding modifications of the defect's energy levels are expected to be small, and were neglected. The height-to-base-length aspect ratio of our hut clusters is <0.1. Therefore, we modelled the DEQDs also as quantum wells, both for solving the Ge bulk diffusion equation, and for calculating the ground state energy of the holes confined to the resulting SiGe layer with locally varying Ge concentration. For the energy level calculations, we used the 6-band k·p envelope function approach described in [38], with the band parameters given in [39]. The heavy-hole ground state (HH 1 ) energy in the DEQDs and the surrounding WL after Ge bulk diffusion were calculated as a function of the Ge diffusion length L D , assuming that the Ge diffusion constant is independent of the Ge concentration for the limit of high Ge concentrations as used in this work. A comparison of the observed DEQD-PL shifts for different T A to the calculated HH 1 energy shifts allows us to determine for each T A a value for L D for which the calculated HH 1 energy shift and the measured PL shift coincide. For the different T A used in this work, the resulting Ge profiles and the HH 1 energy levels for the QD and WL potentials are shown in Figure 4a    In addition to the shift of the PL peak as a function of T A , our model also predicts a value for the ionization energy of the holes in the HH 1 DEQD state that dominates PL quenching via the thermal hole emission from the QD to the higher energy states in the WL, as schematically indicated in Figure 4c,d. For the L D values and the corresponding shapes of the confinement potentials that correctly reproduce the observed PL shifts, the extracted hole ionization energies also agree well with the experimentally observed activation energies for PL quenching. Thus, we are able to describe two independent experimental quantities (PL-and E A -shifts with T A ) with one parameter, L D . These findings justify a posteriori our model assumptions.
From the shift of the PL transition energy with T A and the correlated diffusion lengths L D = (D·τ) 0.5 , where D is the diffusion constant and τ the total annealing time (7200 s), we can additionally determine the activation energy E A,diff for bulk diffusion for Ge in Si. From D(T) = D 0· exp(−E A,diff /kT) we extract an activation energy of E A,diff = 1.4 eV and the diffusion constant at infinite temperature, D 0 = 1.7 × 10 −12 cm 2 /s. The value for E A,diff is surprisingly low compared to the reported values in the literature for coherent, i.e., defect-free, bulk crystals, which are typically about~3 eV [40,41]. Diffusion of highly strained Ge in Si was, however, not investigated thus far. Here, we argue that it is the highly strained region around the split-[110] self-interstitial [18] that leads to the significant lowering of E A,diff . Furthermore, the disorder in the crystal, induced by the Ge ion implantation, might enhance vacancy formation, the main driving factor for bulk diffusion [42].

H-Irradiation Treatment of DEQDs
Hydrogen passivation is an indispensable tool for improving the electrical and optical quality of silicon in applications ranging from solar cells to CMOS technology to emission improvement of GeSi nanostructures [23,43,44]. DEQDs-the optically active material in this work-are surrounded by crystalline Si; the Si float-zone substrate, and the epitaxial Si buffer and capping layers.
Process steps, such as sample cleaning before growth, imperfect Ge ion implantation through the QD layer into the Si buffer layer, or the growth of the Si epilayers itself might lead to the introduction of point defects, which might not be detectable using methods such as transmission electron microscopy (TEM). Furthermore, it is known that the presence of larger crystal defects, such as stacking faults or misfit dislocations, leads to significant PL quenching from SiGe QD emitters [45], which makes the interpretation of the observed changes of the PL yield with respect to the performed in-situ/ex-situ treatments impossible. However, such extended defects are not observed in our samples, as evidenced by the TEM micrographs in Figure 5. Nevertheless, point defects might still limit the PL yield of DEQDs by the non-intentional introduction of non-radiative recombination paths in the Si matrix. Hydrogen passivation is an indispensable tool for improving the electrical and optical quality of silicon in applications ranging from solar cells to CMOS technology to emission improvement of GeSi nanostructures [23,43,44]. DEQDs-the optically active material in this work-are surrounded by crystalline Si; the Si float-zone substrate, and the epitaxial Si buffer and capping layers. Process steps, such as sample cleaning before growth, imperfect Ge ion implantation through the QD layer into the Si buffer layer, or the growth of the Si epilayers itself might lead to the introduction of point defects, which might not be detectable using methods such as transmission electron microscopy (TEM). Furthermore, it is known that the presence of larger crystal defects, such as stacking faults or misfit dislocations, leads to significant PL quenching from SiGe QD emitters [45], which makes the interpretation of the observed changes of the PL yield with respect to the performed in-situ/ex-situ treatments impossible. However, such extended defects are not observed in our samples, as evidenced by the TEM micrographs in Figure 5. Nevertheless, point defects might still limit the PL yield of DEQDs by the non-intentional introduction of non-radiative recombination paths in the Si matrix. In Figure 6 we present the PL-response of a series of samples annealed at T A = 600 • C and recorded at T PL = 300 K (Figure 6a) and T PL = 10 K (Figure 6b). Besides an untreated reference sample, the samples underwent an additional low-energy proton irradiation treatment with various irradiation doses (5 × 10 17 cm −2 , 1 × 10 18 cm −2 , and 5 × 10 18 cm −2 ). For these samples and hydrogenated samples that were previously annealed at other T A , the integrated PL intensity is given in Appendix A.
as stacking faults or misfit dislocations, leads to significant PL quenching from SiGe QD emitters [45], which makes the interpretation of the observed changes of the PL yield with respect to the performed in-situ/ex-situ treatments impossible. However, such extended defects are not observed in our samples, as evidenced by the TEM micrographs in Figure 5. Nevertheless, point defects might still limit the PL yield of DEQDs by the non-intentional introduction of non-radiative recombination paths in the Si matrix. In Figure 6 we present the PL-response of a series of samples annealed at TA = 600 °C and recorded at TPL = 300 K ( Figure 6a) and TPL = 10 K (Figure 6b). Besides an untreated reference sample,   (Figure 6, green and blue curves) can be increased by 48% and 89%, respectively, as compared to the untreated reference sample ( Figure 6, black curve). In contrast, the highest investigated H-implantation dose leads to a 63% decrease of I PL . At T PL = 300 K, the shape of all PL spectra is similar, except for a slight increase of the full width at half maximum of ≈29 meV for some of the H-treated samples. The PL response recorded at T PL = 10 K, however, reveals that H-implantation causes changes in the carrier recombination paths, and thus in the shape of the PL spectra. An additional peak at wavelengths between 1200 nm and 1300 nm appears that is absent in the DEQD reference sample without low-energy proton irradiation and has not been observed in DEQD samples before [10,[18][19][20][21]. This additional peak exhibits a systematic shift to higher energies with increasing proton irradiation dose. It can be assigned to so-called G-centers that are ascribed in the literature to the formation of Si-C vacancy clusters in Si [46].
The PL response from the low-energy proton-irradiated sample with a dose of 1 × 10 18 cm −2 is shown in Figure 7 (black curve) over a T PL range from 20 K to 300 K. These spectra are compared to the ones of the not-H-treated DEQD sample (blue spectrum in Figure 7). Additionally, we compare the signal to another reference sample that contains only the Si layers, but no Ge DEQD layer. The PL response of this Si sample, which underwent the same proton irradiation treatment as the DEQD sample, is depicted in red in Figure 7. For both H-treated samples, the peak intensity associated with the G-center defect quenches already at T PL <120 K. For T PL >120 K, the spectral shapes of the hydrogen-treated DEQD sample and DEQD-reference coincide.
The corresponding temperature dependence of I PL for DEQDs, proton-irradiated DEQDs, and proton-irradiated Si are plotted in Figure 8. As mentioned before, E A(600) amounts to 190 meV. This value slightly improves to 210 meV for the DEQD sample treated with hydrogen. For the proton-irradiated Si epilayer, a low activation energy of 40 meV was found for the G-center related PL signal. The results of Figures 7 and 8 indicate that the PL signal associated with the G-center is caused by structural changes in the Si capping layer upon low-energy proton irradiation, whereas the PL-results confirm that the DEQD layer is structurally not affected by H-treatment. Moreover, the light-emission improvement observed for the H-treated samples at T PL = 300 K proves that H-treatment is a valid tool to be combined with group IV nanostructures, since non-radiative recombination in the Si cladding layers can be suppressed.

Conclusions
We demonstrate that DEQD light-emitters can withstand a high thermal budget of at least up to 600 °C over 2 h, which makes them compatible with front-end integration in CMOS processes. Two main reasons for the quenching of the PL emission from DEQDs have been identified. First, at high

Conclusions
We demonstrate that DEQD light-emitters can withstand a high thermal budget of at least up to 600 °C over 2 h, which makes them compatible with front-end integration in CMOS processes. Two main reasons for the quenching of the PL emission from DEQDs have been identified. First, at high

Conclusions
We demonstrate that DEQD light-emitters can withstand a high thermal budget of at least up to 600 • C over 2 h, which makes them compatible with front-end integration in CMOS processes. Two main reasons for the quenching of the PL emission from DEQDs have been identified. First, at high luminescence temperatures, quantum dot ionization occurs by the escape of heavy holes from the deep valence band potentials to the higher energy states in the surrounding wetting layer (up to 250 meV). Second, at high annealing temperatures (T A >650 • C), defect migration out of the QD occurs, which quenches the PL intensity at all T PL . The latter is associated with a surprisingly high activation energy, larger than 3.4 eV. We showed that annealing at moderate temperatures of 500 • C to 550 • C, as well as low-energy ex-situ proton irradiation, are advantageous technology steps to improve the optical properties of DEQDs by the curing and passivation of point defects in the Si matrix surrounding the DEQDs. Figure A1 depicts a summary of the integrated PL intensities for the reference sample and the annealed samples at TA, as well as the samples that underwent low-energy proton irradiation treatment. Figure A1. Integrated PL intensity for samples treated with different TA and hydrogen curing conditions. Red and black data points were extracted from measurements at TPL = 300 K and 10 K, respectively. Open symbols correspond to IPL from the DEQDs only, while full symbols indicate IPL from both DEQD emission and the light emission originating from the low energy proton irradiation. Figure A1. Integrated PL intensity for samples treated with different T A and hydrogen curing conditions. Red and black data points were extracted from measurements at T PL = 300 K and 10 K, respectively. Open symbols correspond to I PL from the DEQDs only, while full symbols indicate I PL from both DEQD emission and the light emission originating from the low energy proton irradiation.