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GaAs quantum wells grown by molecular beam epitaxy
Optical Materials 64 (2017) 361e365
Contents lists available at ScienceDirect
Optical Materials journal homepage: www.elsevier.com/locate/optmat
Very high dose electron irradiation effects on photoluminescence from GaInNAs/GaAs quantum wells grown by molecular beam epitaxy lt¸a t¸eanu b, S.I. Spa ^nulescu b, E. Arola c, d E.-M. Pavelescu a, b, *, N. Ba a
National Institute for Research and Development in Microtechnologies, Erou Iancu Nicolae 126A, 077190, Bucharest, Romania la ras¸ilor 169, 030615, Bucharest, Romania Faculty of Exact Sciences and Engineering, Hyperion University, Calea Ca c Department of Physics, Tampere University of Technology, FIN-33101, Tampere, Finland d Optoelectronics Research Centre, Tampere University of Technology, FIN-33101, Tampere, Finland b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 13 July 2016 Received in revised form 19 October 2016 Accepted 6 December 2016
The effects of 7 MeV electron irradiation at very high doses of 2 1017 and 1:5 1018 electrons=cm2 and subsequent rapid thermal annealing on photoluminescence from a strain-compensated GaInAsN/GaAsN/ GaAs quantum well structure are investigated. A large additional blueshift of photoluminescence has been observed from the lower-dose irradiated sample as compared to the non-irradiated one when annealed after the irradiation. This additional blueshift will become considerably reduced by an ageing effect, which occurs already at room temperature. The mechanism causing the additional blueshift of photoluminescence and its reduction is qualitatively assigned to metastable complex defects promoted by electron irradiation in the nitrogen containing layers. No such additional blueshift of photoluminescence under the thermal treatment has been observed in the higher-dose irradiated sample. © 2016 Elsevier B.V. All rights reserved.
Keywords: Electron irradiation Photoluminescence Dilute nitrides
1. Introduction The electronic band structure properties of Ga1x Inx As1y Ny =GaAs have made this semiconductor alloy very attractive for a potential use in 1.3 mm GaInAsN/GaAs heterojunction based quantum well lasers and solar cells [1,2]. Unfortunately, GaInAsN is a polymorphic and metastable alloy containing a large number of point defects, notably, nitrogen interstitials and gallium vacancies [3]. The defects act as non-radiative recombination centers which deteriorate the optical and electrical properties of the alloy. The density of defects can be reduced by annealing, indicating that the non-radiative defects exhibit low activation energies. However, annealing induces a large undesired blueshift (BS) of the GaInAsN alloy band gap [4]. The BS is also facilitated by the presence of point defects [5]. Electron irradiation of semiconductors is known to generate point defects [6]. Applying electron irradiation to GaInAsN is of interest because it provides information about the nature and formation process of the defects. It has been shown that the electron irradiation effects on GaInAsN and its heterostructures strongly
* Corresponding author. National Institute for Research and Development in Microtechnologies, Erou Iancu Nicolae 126A, 077190, Bucharest, Romania. E-mail address: [email protected] (E.-M. Pavelescu). http://dx.doi.org/10.1016/j.optmat.2016.12.007 0925-3467/© 2016 Elsevier B.V. All rights reserved.
depend on the magnitude of radiation dose. Thus, at relatively low doses within the range of 1013 1015 electrons=cm2 it has been observed a direct enhancement in luminescence efficiency of GaInNAs-containing heterostructures [7,8]. Relatively medium doses around 1016 electrons=cm2 have been seen to promote an enhancement in luminescence efficiency of GaInNAs layers and quantum wells when annealed upon irradiation, in spite of an initial deterioration of the photoluminescence spectra of the samples upon irradiation [9,10]. Very recently, higher doses around 1017 electrons=cm2 were shown to promote unstable defects, as reflected by the instability of photoluminescence, in GaInAsN/GaAs quantum wells when annealed upon irradiation [11]. The influence of electron irradiation at high doses beyond 1017 electrons=cm2 on GaInAsN-containing heterostructures has not been studied so far. Such a study would be useful for many applications because electron irradiation can create similar defects as are being introduced during the processing of GaInAsN-based devices, e.g. during ion implantation. Furthermore, such defects are created, for instance, in high-efficiency tandem solar cells in a radiation environment in space. Therefore, the electron irradiation experiments could be a vital method in estimating the irradiation stability of the dilute nitrideearsenides and other semiconductor compounds used in devices operating in space. In fact, irradiation with high-energy electrons are commonly employed to assess the end-of-life parameter of solar cells [12].
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Fig. 1 shows our experimental results on photoluminescence. The PL spectra of the three, initially identical and non-electronirradiated, strain-compensated GaInAsN/GaAsN/GaAs epitaxial quantum well (QW) structures, denoted as S, S17, and S18. Consequently, two of these samples, namely the samples S17 and S18 are being exposed to 7 MeV electron irradiation, and the corresponding PL spectra has also been plotted in Fig. 1 (for further details of the
samples and electron irradiation characteristics, see Section 2). It can be seen from Fig. 1 that electron irradiation systematically deteriorates the PL intensity with the increasing dose, initially before the electron irradiation all the three samples exhibiting similar peak wavelengths, intensities and linewidths (FWHMs) within the uncertainty limits of the PL setup. Furthermore, one can observe from Fig. 1 that the PL spectra from the S17 sample is clearly red-shifted by 5 nm relative to the PL spectra from the sample S, whereas the PL spectral peak from the sample S18 is nearly unaffected by the higher-dose irradiation as can be judged from the slight redshift (RS) by only 2 nm, similar to the wavelength resolution of our PL mapper. Obviously, the electron irradiation promoted changes in the S17 sample either in the nitride GaInAsN and GaAsN layers or in the (Al) GaAs layers through which the photons passed in the PL experiments. The observed redshift suggests that the changes in lattice structure have taken place preferentially in the nitride layers. The interpretation behind this redshift could be understood if electron irradiation helps nitrogen atoms to become incorporated into substitutional sites of the As sublattice [14,15]. After the electron irradiation, all three samples were annealed at 650 C for 5 min (RTA1 treatment). Fig. 2(a) shows the PL spectra from these annealed samples. It can be seen that annealing has promoted a large blueshift (BS) z80 nm in the PL spectra for the sample S17 but only 21 and 22 nm blueshift in the PL spectra for the samples S and S18, respectively. Besides the large additional BS (around 60 nm) observed, it is interesting to notice from Fig. 2(a) that annealing has caused the PL intensity from the S17 sample to become significantly improved, by a factor of 20 as compared to the PL intensity from the same non-annealed electron irradiated sample in Fig. 1. Fig. 2(a) further shows that the PL intensity from the annealed S17 sample is about two times larger than the PL intensity from the non-irradiated annealed sample S. Furthermore, it is noticeable that the intensity of the PL spectra from the nonannealed S18 sample in Fig. 1 has also remarkably improved, by a factor of 10, upon annealing [see Fig. 2(a)]. Interestingly, Fig. 2(a) also shows that the PL spectral profile from the annealed electron irradiated S18 sample follows closely that measured from the nonirradiated annealed sample S whose intensity has undergone only 5 times enhancement when compared to its non-annealed case in Fig. 1. Concerning the linewidths, Fig. 2(a) shows that while the FWHM values of the PL spectral peaks from the S and S18 samples
Fig. 1. 300-K PL's plot recorded from samples S, S17 and S18 before and after irradiation.
Fig. 2. 300-K PL's plot recorded from samples S, S17, and S18 upon (a) RTA1 and (b) 4 days at 20 C subsequent to the RTA1 treatment.
In this work we have investigated the effects of 7 MeV electron irradiation at high doses of 2 1017 and 1:5 1018 electrons=cm2 on photoluminescence from a strain-compensated GaInAsN/ GaAsN/GaAs quantum well structure. 2. Experimental Our sample consisted of an epitaxial structure with a 7 nm thick Ga0:63 In0:37 As0:99 N0:01 layer/20 nm thick GaAs0:99 N0:01 layer/ 130 nm thick GaAs layer, surrounded by a 100 nm thick Al0:3 Ga0:7 As cladding layer on both sides of the structure and capped with a 10 nm GaAs layer. The structure was deposited onto n-type GaAs (100) substrates using a molecular beam epitaxy (MBE) system equipped with a nitrogen rf-plasma source. The N-containing layers were grown at the temperature of 430 C and the (Al)GaAs layers were grown at the temperature of 580 C. Upon growth, three pieces of identical size (3 mm 3 mm 350 mm) were cut from the same region of the wafer. Two of these samples, labeled S17 and S18, were subject to irradiation at room temperature with 7 MeV electrons and electron beam doses of 2 1017 and 1:5 1018 electrons=cm2 , respectively. The details of the electron irradiation are described elsewhere [13]. The irradiation electronbeam current was set to 5 mA corresponding to the electron radiation flux density of 3:125 109 electrons=cm2 s. The third sample, labeled S, was not irradiated. The photoluminescence (PL) spectra was recorded at room temperature using an automated PL mapper equipped with a 532 nm Nd-YAG laser and an InGaAs array detector. The spectral resolution of the PL mapper in the near infrared region of interest (around 1300 nm) is around 2 nm (1 mm slit width, 0:3 m monochromator, 150 grooves=mm grating blazed at 1250 nm). Rapid thermal annealing (RTA) was performed at 650 C in dry nitrogen atmosphere in steps of 5 min. 3. Results and discussion
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have decreased by 4 and 3 meV, respectively, the FWHM value of the PL spectral peak from the S17 sample has suffered a noticeable enhancement of 12 meV, making the differences in linewidths very noticeable. We briefly discuss in the following the ageing phenomenon as seen in the PL spectra from various GaInAsN/GaAsN/GaAs based quantum well structures. When Fig. 2(b) is compared with Fig. 2(a) it is noticeable that the annealed S17 sample exhibits a rapid ageing effect in terms of the redshifted PL spectra and reduced value of the FWHM of the PL peak, while the PL spectral features from the annealed S and S18 samples have remained unchanged. When the samples were kept at 20 C for four days, the differences in the wavelength values (DlðS; S17Þ≡lðS17Þ lðSÞ) and FWHM values (DFWHMðS; S17Þ≡FWHMðS17Þ FWHMðSÞ) of the PL peaks between the S17 and S samples have considerably become reduced in terms of their absolute values to about 23 nm and 6 meV, respectively. This reduction in jDlðS; S17Þj and jDFWHMðS; S17Þj values is due to a significant PL redshift and decrease in the FWHM value experienced by the S17 sample during the ageing process as can be deduced from Fig. 2(a) and (b). Interestingly, this ageing behavior of the PL spectral peak position of the S17 sample can also be seen in similar samples (not shown here for brevity) whose N-containing layers were grown both at a higher, 450 C temperature [11], and at a lower 410 C temperature. This clearly indicates that the electron irradiationrelated defects are, on one hand, unstable with very low activation potential barriers and, on the other hand, are remarkably easily created by electron irradiation irrespective of the growth temperature. This conclusion should be valid over the most commonly used range of growth temperatures (410 450 C) at which the growth of dilute nitride materials takes place by the MBE. Fig. 3(a) shows that when the second RTA treatment was performed (RTA2, 650 C=5 min) a significant additional BS (around 75 nm) was again seen for the S17 sample as compared to the S sample taking into account the large blueshift of 82 nm generated in the case of the S17 sample and the blueshifts for the S and S18 samples of only 6 and 7 nm, respectively, relative to the PL peak positions of the RTA1 treated samples after 4 days of ageing at 20 C [cf. Fig. 2(b)]. Furthermore, when comparing Fig. 3(a) with Fig. 2(b) we notice that the change in the FWHM was þ5 meV for S17 sample and around 1 meV for the S18 sample. Fig. 3(b) shows the PL spectra measured for the samples S, S17,
and S18 after the third RTA treatment (RTA3, 650 C=5 min) was performed very soon after the RTA2 treatment. Another surprising and remarkable redshift of 93 nm and the decrease in the FWHM of the PL spectral peak (DFWHM ¼ 10 meV) were observed in the case of the S17 sample. Interestingly, the RTA3 treatment of the S17 sample returned its PL spectral peak wavelength and the FWHM values to nearly the same peak wavelength and FWHM values measured from the RTA3-treated S sample (in other words, the significant RTA2-induced additional BS of the S17 sample as compared to the S sample almost vanished after the RTA3). During the RTA3 treatment, the PL peaks of the S and S18 samples blueshifted only marginally, 2 3 nm, and the FWHM remained essentially unchanged. It appears from Figs. 3 and 2 that the RTA2 and RTA3 treatments activated more pronouncedly the same microscopic mechanisms as did the RTA1 treatment and the ageing process at the temperature of 20 C, respectively. In particular, we notice that the blueshift after RTA2 treatment is noticeably larger than the blueshift after the RTA1 treatment. Furthermore, it is interesting to see that the RTA3 treatment, which was carried out soon after the RTA2 treatment, caused such amount of redshift and decrease in the FWHM capable of making the PL spectra of the S17 sample nearly identical to the PL spectra of the S sample upon the same RTA3 treatment. It is worth mentioning that the PL spectra of a sample, similar to the S17 sample, whose N-containing layers were grown at 450 C became almost completely identical to that of the corresponding nonirradiated sample after either the ageing process at room temperature for 4 days or after the RTA3 treatment performed soon after [11] the RTA2. In Fig. 4 are depicted high resolution XRD curves recorded from the samples S, S17 and S18 after irradiation (black lines) and after the RTA3 stage (gray lines). As compared to the corresponding nonannealed sample, the GaInNAs quantum well of each RTA3-treated sample experienced a very tiny decrease in strain whose maximum value was observed for S17 and amounted only 0.011% (corresponding to 24 arcsecs left shift of the XRD feature related to the GaInNAs well). At the same time, no change can be observed in the shape and position of the XRD features related to the GaNAs barriers. This is to say that strain relaxation by means of either misfit dislocations (which would lead to red shift) or elemental
Fig. 3. 300-K PL's plot recorded from samples S, S17 and S18 upon (a) RTA2 and (b) RTA3.
Fig. 4. High resolution x-ray diffraction (XRD) curves recorded from S, S17 and S18 after irradiation (black lines) and RTA3 treatment (gray lines).
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interdiffusion (which would lead to blue shift) across quantum well and barrier interfaces was negligible and cannot account alone for any of the significant blue shifts as well as additional blue shifts and subsequent red shifts encountered in our samples. Hence, the remarkable blue shifts and the additional blue shifts/red shifts observed must be assigned to other causes. Thus, the blue shift seen for the samples S and S18 upon annealing as well as for the sample S17 after the RTA3 stage is mainly due to formation of IneN bonds on annealing at the expense of Ga-N ones, promoted during growth. However, this change on annealing in the atomic composition of the N-centered tetrahedral N Gam In4m clusters (0 m 4) from a Ga-rich composition to an In-rich composition (which generates blue shift) should be excluded as the cause for the additional BS/subsequent RS observed because a change back of the N atom nearest neighborhood from the N-rich to the Ga-rich one (which generates red shift) is very improbable in the equilibrium theory of the alloying formation [4]. The mechanisms responsible for the additional blueshift and subsequent redshift observed in the PL spectra from the S17 sample relative to the PL spectra from the S sample should be assigned to the presence of additional particular type of defects created/influenced by electron irradiation. The fact that the spectral line shifts were always accompanied with substantial changes in the FWHM suggests to us that the defects are mainly in the quantum well and its barriers region, which accommodates the biggest part of the carriers' wave functions involved in the PL process. We further assume that electron irradiation yields metastable defect centers. Metastable centers are known to exist in electron irradiated GaAs [6], lattice-matched Ga1x Inx Ny As1y (xz2:8y) on GaAs [16], and were theoretically predicted to occur in ternary GaAsN alloys [17]. Let us assume that electron irradiation initially produced complex neutral or charged defect centers in the quantum well of the pristine sample. It is then conceivable that heating the sample at 650 C (under the RTA1 or RTA2 treatments) could influence the atomic and electronic structures of these defects and the surrounding host material, for example, by moving nearby nitrogen atoms from their initial substitutional sites to energetically more favorable defect sites [17,18]. Our first-principles studies in the framework of the density functional theory (DFT) [17] shows that when a nitrogen atom in the dilute GaAs1x Nx ternary alloy moves from a substitutional group V site into any of the stable or metastable interstitial sites defined by the nearest Ga or As tetrahedron, then the alloy bandgap will increase causing the blueshift in the PL spectra. Due to the isoelectronic nature of the Ga and In atoms it is then obvious that in the present case of the Ga1x Inx As1y Ny quaternary alloy in the quantum well, the nitrogen atom moving from its subsitutional group V site into the nearest interstitial site defined by a Ga tetrahedron, As tetrahedron, or Gam In4m (0 m 4) tetrahedron, the bandgap will similarly get increased, causing the blueshift in our observed PL spectra. Furthermore, our calculations on the GaAs1x Nx ternary alloy show that, out of many possible interstitial states capable of causing the blueshift, it is the single N impurity low-energy eGN type interstitial defect defined in Ref. [18] which will cause the largest blueshift. Therefore, we tentatively propose that the same type of single N impurity interstitial defect (eGN) is responsible for the blueshift observed in the Ga0:63 In0:37 As0:99 N0:01 =GaAs quantum well, rather than certain GaIn atomic configurations around the substitutional N in the Ga0:63 In0:37 As0:99 N0:01 alloy [5] as the potential-energy transition barrier between these configurational states is obviously quite large. We have further experimental evidence to exclude the possibility of having a substitutional N atom surrounded by disordered Ga-In atomic configurations (N@Ga In) to explain the blueshift.
Namely, as it is well known from the theory of disordered alloys, this substitutional disorder between the Ga and In atoms in the group III sublattice of the Ga0:63 In0:37 As0:99 N0:01 alloy would enhance scattering of carriers involved in the photoluminescence process, therefore decreasing their overall lifetime and consequently broadening the linewidth of the PL spectral transitions. However, this disorder related broadening is not clearly visible in our PL spectra, therefore excluding the possibility that the observed blueshift would be, at least majorly, due to the N@Ga In type defects. It is obvious that large scale ab initio molecular dynamics calculations would be required in order to quantitatively differentiate between these two defect possibilities (N interstitial vs. N substitutional cases) responsible for the blueshift observed in the Ga0:63 In0:37 As0:99 N0:01 =GaAs quantum well. Since we know that the activation energy of the defects in the S17 sample is very low, as evidenced by the instability of this sample during the ageing process at the room temperature of 20 C [cf. Fig. 2(b)], the nitrogen atoms are expected to be only loosely bound to the obvious interstitial defect sites which can then readily hop back to substitutional sites by slight thermal activation. This process would give rise to the redshift. Interestingly, this could explain why the sizes of the redshift and blueshift observed in the photoluminescence spectra are almost the same (jRSjzjBSj) when the complete cycle of activation and deactivation of the electronirradiated defects is carried through. This is the case with the RTA2/RTA3 treatments and most likely also with the previous RTA1/ageing at 20 C processes if the ageing duration would have been prolonged beyond 4 days or the environmental temperature would have been higher than 20 C. Finally we notice that by increasing the electron irradiation dose up to one order of magnitude larger up to 1.5 1018 electrons/cm2 produced additional defects or modified those already produced at lower doses during irradiation process till 1018 electrons/cm2 such that the resulting photoluminescence spectral behavior was essentially different from that observed in the sample S17 irradiated with the dose of 2 1017 electrons=cm2 . The additional defects may interact with the metastable defects produced at lower doses up to 1017 electrons/cm2 or the amount of the metastable defects are so high as they strongly interact each other (they could agglomerate forming clusters) in such way that the potential barrier between their ground and metastable state is so high as they cannot longer undergo a change in their energy states neither in short time by thermal annealing at elevated temperatures nor in longer time by ageing at room temperature. Thus, the formation of IneN bonds remains the only major mechanism for annealinginduced blue shift in the sample S18, explaining why the magnitude of annealing-induced blue shift observed in the sample S18 is similar to that seen for S despite irradiation. Regarding the different behavior of photoluminescence with the applied dose of electron irradiation, it is worth mentioning our recent observation that a 7 MeV electron irradiation on GaInNAs thin layers up to the dose of 1015 electrons=cm2 produced a photoluminescence enhancement whereas one order of magnitude higher dose induced a deterioration of luminescence efficiency of the investigated GaInNAs layers [8]. 4. Summary and conclusions It has been found in our studies that 7 MeV electron irradiation with doses of 2 1017 and 1:5 1018 electrons=cm2 will cause atomic structural damage in the GaInNAs/GaNAs/GaAs quantum wells, probably creating complex defect centers. When the lowerdose sample (exposed to the dose of 2 1017 electrons=cm2 ) was annealed at 650 C, a pronounced additional blueshift of PL was observed, which made the total blueshift much larger than the blue
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shift of a non-irradiated sample. The irradiation-created defects turned out to be unstable, making the additional blue shift remarkably reduce or almost vanish (by a subsequent red shift of photoluminescence), a phenomenon which occurred even at room temperature. The additional blueshift observed in this sample can qualitatively be explained by assuming that there is a hopping mechanism for nitrogen atoms to move from their initial substitutional sites to energetically more favorable defect sites. On the other hand, hopping of nitrogen atoms in the reverse order, from the defect sites to the substitutional sites, will cause the subsequent redshift of the PL spectra. Due to the instability of this defect state under thermal activation, the nitrogen atom could return back to its initial site, restoring the PL peak and line shape very close to their initial values. No such PL instability upon thermal annealing was observed in the case of the higher-dose irradiated sample (exposed to the dose of 1:5 1018 electrons=cm2 ), probably due to inactivation of the irradiation-induced unstable defects by interaction with each other following their increased concentration or with additional defects generated at the higher irradiation dose. Acknowledgments The research leading to these results has received funding from the EEA Financial Mechanism 2009e2014 under the project contract no 23SEE/30.06.2014.
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References [1] W. Li, T. Jouhti, C.S. Peng, J. Konttinen, P. Laukkanen, E.M. Pavelescu, M. Dumitrescu, M. Pessa, Appl. Phys. Lett. 79 (2001) 3386. [2] B. Royall, N. Balkan, S. Mazzucato, H. Khalil, M. Hugues, J.S. Roberts, Phys. Stat. Sol. 248 (2011) 1191. [3] W. Li, M. Pessa, T. Ahlgren, J. Dekker, Appl. Phys. Lett. 79 (2001) 1094. [4] K. Kim, A. Zunger, Phys. Rev. Lett. 86 (2001) 2609. [5] P.J. Klar, H. Gruning, J. Koch, S. Schafer, K. Volz, M. Stolz, W. Heimbrodt, A.M. Saadi, A. Lindsay, E.P. O'Reilly, Phys. Rev. B 64 (2001) 121203. venard, J. Appl. Phys. 64 (1988) R65. [6] J.C. Bourgoin, H.J. Bardeleben, D. Stie t¸eanu, T. Jouhti, V. Cimpoca, M. Pessa, [7] E.-M. Pavelescu, A. Gheorghiu, N. B alt¸a in: Semiconductor Conference, CAS 2004 Proceedings (IEEE, Piscataway, NJ, 2004), vol. 1, 2004, p. 220. [8] E.-M. Pavelescu, R. Kudrawiec, J. Puustinen, A. Tukiainen, M. Guina, J. Luminesc 136 (2013) 347. lt¸ [9] E.-M. Pavelescu, R. Kudrawiec, N. Ba at¸eanu, S. Sp^ anulescu, M. Dumitrescu, M. Guina, Semicond. Sci. Technol. 28 (2013) 025020. [10] E.-M. Pavelescu, A. Gheorghiu, R. Kudrawiec, M. Dumitrescu, A. Tukiainen, T. Jouhti, T. Hakkarainen, N. Tkachenko, V.D.S. Dhaka, H. Lemmetyinen, J. Andrzejewski, J. Misiewicz, M. Pessa, Appl. Phys. Lett. 85 (2004) 6158. [11] E.-M. Pavelescu, M. Dumitrescu, M. Guina, J. Luminesc 154 (2014) 584. [12] F. Dimroth, Phys. Stat. Sol. C 3 (2006) 373. [13] E.-M. Pavelescu, S. Antohe, V. Ruxandra, S. Sp^ anulescu, V. Cimpoca, M. Pessa, in: Semiconductor Conference, 2001, CAS 2001 Proceedings (IEEE, Piscataway, NJ, 2001), vol. 2, 2001, p. 319. [14] X. Yang, J.B. Heroux, M.J. Jurkovic, W.I. Wang, J. Vac. Sci. Technol. B 17 (1999) 1144. [15] S. Francoeur, G. Sivaraman, Y. Qiu, S. Nikishin, H. Temkin, Appl. Phys. Lett. 72 (1998) 1857. [16] J.Z. Li, J.Y. Lin, H.X. Jiang, J.F. Geisz, S.R. Kurtz, Appl. Phys. Lett. 75 (1999) 1899. [17] E. Arola, J. Ojanen, H.-P. Komsa, T.T. Rantala, Phys. Rev. B 72 (2005) 045222. [18] W.K. Loke, S.F. Yoon, S.Z. Wang, T.K. Ng, W.J. Fan, J. Appl. Phys. 91 (2002) 4900.
Contents lists available at ScienceDirect
Optical Materials journal homepage: www.elsevier.com/locate/optmat
Very high dose electron irradiation effects on photoluminescence from GaInNAs/GaAs quantum wells grown by molecular beam epitaxy lt¸a t¸eanu b, S.I. Spa ^nulescu b, E. Arola c, d E.-M. Pavelescu a, b, *, N. Ba a
National Institute for Research and Development in Microtechnologies, Erou Iancu Nicolae 126A, 077190, Bucharest, Romania la ras¸ilor 169, 030615, Bucharest, Romania Faculty of Exact Sciences and Engineering, Hyperion University, Calea Ca c Department of Physics, Tampere University of Technology, FIN-33101, Tampere, Finland d Optoelectronics Research Centre, Tampere University of Technology, FIN-33101, Tampere, Finland b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 13 July 2016 Received in revised form 19 October 2016 Accepted 6 December 2016
The effects of 7 MeV electron irradiation at very high doses of 2 1017 and 1:5 1018 electrons=cm2 and subsequent rapid thermal annealing on photoluminescence from a strain-compensated GaInAsN/GaAsN/ GaAs quantum well structure are investigated. A large additional blueshift of photoluminescence has been observed from the lower-dose irradiated sample as compared to the non-irradiated one when annealed after the irradiation. This additional blueshift will become considerably reduced by an ageing effect, which occurs already at room temperature. The mechanism causing the additional blueshift of photoluminescence and its reduction is qualitatively assigned to metastable complex defects promoted by electron irradiation in the nitrogen containing layers. No such additional blueshift of photoluminescence under the thermal treatment has been observed in the higher-dose irradiated sample. © 2016 Elsevier B.V. All rights reserved.
Keywords: Electron irradiation Photoluminescence Dilute nitrides
1. Introduction The electronic band structure properties of Ga1x Inx As1y Ny =GaAs have made this semiconductor alloy very attractive for a potential use in 1.3 mm GaInAsN/GaAs heterojunction based quantum well lasers and solar cells [1,2]. Unfortunately, GaInAsN is a polymorphic and metastable alloy containing a large number of point defects, notably, nitrogen interstitials and gallium vacancies [3]. The defects act as non-radiative recombination centers which deteriorate the optical and electrical properties of the alloy. The density of defects can be reduced by annealing, indicating that the non-radiative defects exhibit low activation energies. However, annealing induces a large undesired blueshift (BS) of the GaInAsN alloy band gap [4]. The BS is also facilitated by the presence of point defects [5]. Electron irradiation of semiconductors is known to generate point defects [6]. Applying electron irradiation to GaInAsN is of interest because it provides information about the nature and formation process of the defects. It has been shown that the electron irradiation effects on GaInAsN and its heterostructures strongly
* Corresponding author. National Institute for Research and Development in Microtechnologies, Erou Iancu Nicolae 126A, 077190, Bucharest, Romania. E-mail address: [email protected] (E.-M. Pavelescu). http://dx.doi.org/10.1016/j.optmat.2016.12.007 0925-3467/© 2016 Elsevier B.V. All rights reserved.
depend on the magnitude of radiation dose. Thus, at relatively low doses within the range of 1013 1015 electrons=cm2 it has been observed a direct enhancement in luminescence efficiency of GaInNAs-containing heterostructures [7,8]. Relatively medium doses around 1016 electrons=cm2 have been seen to promote an enhancement in luminescence efficiency of GaInNAs layers and quantum wells when annealed upon irradiation, in spite of an initial deterioration of the photoluminescence spectra of the samples upon irradiation [9,10]. Very recently, higher doses around 1017 electrons=cm2 were shown to promote unstable defects, as reflected by the instability of photoluminescence, in GaInAsN/GaAs quantum wells when annealed upon irradiation [11]. The influence of electron irradiation at high doses beyond 1017 electrons=cm2 on GaInAsN-containing heterostructures has not been studied so far. Such a study would be useful for many applications because electron irradiation can create similar defects as are being introduced during the processing of GaInAsN-based devices, e.g. during ion implantation. Furthermore, such defects are created, for instance, in high-efficiency tandem solar cells in a radiation environment in space. Therefore, the electron irradiation experiments could be a vital method in estimating the irradiation stability of the dilute nitrideearsenides and other semiconductor compounds used in devices operating in space. In fact, irradiation with high-energy electrons are commonly employed to assess the end-of-life parameter of solar cells [12].
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Fig. 1 shows our experimental results on photoluminescence. The PL spectra of the three, initially identical and non-electronirradiated, strain-compensated GaInAsN/GaAsN/GaAs epitaxial quantum well (QW) structures, denoted as S, S17, and S18. Consequently, two of these samples, namely the samples S17 and S18 are being exposed to 7 MeV electron irradiation, and the corresponding PL spectra has also been plotted in Fig. 1 (for further details of the
samples and electron irradiation characteristics, see Section 2). It can be seen from Fig. 1 that electron irradiation systematically deteriorates the PL intensity with the increasing dose, initially before the electron irradiation all the three samples exhibiting similar peak wavelengths, intensities and linewidths (FWHMs) within the uncertainty limits of the PL setup. Furthermore, one can observe from Fig. 1 that the PL spectra from the S17 sample is clearly red-shifted by 5 nm relative to the PL spectra from the sample S, whereas the PL spectral peak from the sample S18 is nearly unaffected by the higher-dose irradiation as can be judged from the slight redshift (RS) by only 2 nm, similar to the wavelength resolution of our PL mapper. Obviously, the electron irradiation promoted changes in the S17 sample either in the nitride GaInAsN and GaAsN layers or in the (Al) GaAs layers through which the photons passed in the PL experiments. The observed redshift suggests that the changes in lattice structure have taken place preferentially in the nitride layers. The interpretation behind this redshift could be understood if electron irradiation helps nitrogen atoms to become incorporated into substitutional sites of the As sublattice [14,15]. After the electron irradiation, all three samples were annealed at 650 C for 5 min (RTA1 treatment). Fig. 2(a) shows the PL spectra from these annealed samples. It can be seen that annealing has promoted a large blueshift (BS) z80 nm in the PL spectra for the sample S17 but only 21 and 22 nm blueshift in the PL spectra for the samples S and S18, respectively. Besides the large additional BS (around 60 nm) observed, it is interesting to notice from Fig. 2(a) that annealing has caused the PL intensity from the S17 sample to become significantly improved, by a factor of 20 as compared to the PL intensity from the same non-annealed electron irradiated sample in Fig. 1. Fig. 2(a) further shows that the PL intensity from the annealed S17 sample is about two times larger than the PL intensity from the non-irradiated annealed sample S. Furthermore, it is noticeable that the intensity of the PL spectra from the nonannealed S18 sample in Fig. 1 has also remarkably improved, by a factor of 10, upon annealing [see Fig. 2(a)]. Interestingly, Fig. 2(a) also shows that the PL spectral profile from the annealed electron irradiated S18 sample follows closely that measured from the nonirradiated annealed sample S whose intensity has undergone only 5 times enhancement when compared to its non-annealed case in Fig. 1. Concerning the linewidths, Fig. 2(a) shows that while the FWHM values of the PL spectral peaks from the S and S18 samples
Fig. 1. 300-K PL's plot recorded from samples S, S17 and S18 before and after irradiation.
Fig. 2. 300-K PL's plot recorded from samples S, S17, and S18 upon (a) RTA1 and (b) 4 days at 20 C subsequent to the RTA1 treatment.
In this work we have investigated the effects of 7 MeV electron irradiation at high doses of 2 1017 and 1:5 1018 electrons=cm2 on photoluminescence from a strain-compensated GaInAsN/ GaAsN/GaAs quantum well structure. 2. Experimental Our sample consisted of an epitaxial structure with a 7 nm thick Ga0:63 In0:37 As0:99 N0:01 layer/20 nm thick GaAs0:99 N0:01 layer/ 130 nm thick GaAs layer, surrounded by a 100 nm thick Al0:3 Ga0:7 As cladding layer on both sides of the structure and capped with a 10 nm GaAs layer. The structure was deposited onto n-type GaAs (100) substrates using a molecular beam epitaxy (MBE) system equipped with a nitrogen rf-plasma source. The N-containing layers were grown at the temperature of 430 C and the (Al)GaAs layers were grown at the temperature of 580 C. Upon growth, three pieces of identical size (3 mm 3 mm 350 mm) were cut from the same region of the wafer. Two of these samples, labeled S17 and S18, were subject to irradiation at room temperature with 7 MeV electrons and electron beam doses of 2 1017 and 1:5 1018 electrons=cm2 , respectively. The details of the electron irradiation are described elsewhere [13]. The irradiation electronbeam current was set to 5 mA corresponding to the electron radiation flux density of 3:125 109 electrons=cm2 s. The third sample, labeled S, was not irradiated. The photoluminescence (PL) spectra was recorded at room temperature using an automated PL mapper equipped with a 532 nm Nd-YAG laser and an InGaAs array detector. The spectral resolution of the PL mapper in the near infrared region of interest (around 1300 nm) is around 2 nm (1 mm slit width, 0:3 m monochromator, 150 grooves=mm grating blazed at 1250 nm). Rapid thermal annealing (RTA) was performed at 650 C in dry nitrogen atmosphere in steps of 5 min. 3. Results and discussion
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have decreased by 4 and 3 meV, respectively, the FWHM value of the PL spectral peak from the S17 sample has suffered a noticeable enhancement of 12 meV, making the differences in linewidths very noticeable. We briefly discuss in the following the ageing phenomenon as seen in the PL spectra from various GaInAsN/GaAsN/GaAs based quantum well structures. When Fig. 2(b) is compared with Fig. 2(a) it is noticeable that the annealed S17 sample exhibits a rapid ageing effect in terms of the redshifted PL spectra and reduced value of the FWHM of the PL peak, while the PL spectral features from the annealed S and S18 samples have remained unchanged. When the samples were kept at 20 C for four days, the differences in the wavelength values (DlðS; S17Þ≡lðS17Þ lðSÞ) and FWHM values (DFWHMðS; S17Þ≡FWHMðS17Þ FWHMðSÞ) of the PL peaks between the S17 and S samples have considerably become reduced in terms of their absolute values to about 23 nm and 6 meV, respectively. This reduction in jDlðS; S17Þj and jDFWHMðS; S17Þj values is due to a significant PL redshift and decrease in the FWHM value experienced by the S17 sample during the ageing process as can be deduced from Fig. 2(a) and (b). Interestingly, this ageing behavior of the PL spectral peak position of the S17 sample can also be seen in similar samples (not shown here for brevity) whose N-containing layers were grown both at a higher, 450 C temperature [11], and at a lower 410 C temperature. This clearly indicates that the electron irradiationrelated defects are, on one hand, unstable with very low activation potential barriers and, on the other hand, are remarkably easily created by electron irradiation irrespective of the growth temperature. This conclusion should be valid over the most commonly used range of growth temperatures (410 450 C) at which the growth of dilute nitride materials takes place by the MBE. Fig. 3(a) shows that when the second RTA treatment was performed (RTA2, 650 C=5 min) a significant additional BS (around 75 nm) was again seen for the S17 sample as compared to the S sample taking into account the large blueshift of 82 nm generated in the case of the S17 sample and the blueshifts for the S and S18 samples of only 6 and 7 nm, respectively, relative to the PL peak positions of the RTA1 treated samples after 4 days of ageing at 20 C [cf. Fig. 2(b)]. Furthermore, when comparing Fig. 3(a) with Fig. 2(b) we notice that the change in the FWHM was þ5 meV for S17 sample and around 1 meV for the S18 sample. Fig. 3(b) shows the PL spectra measured for the samples S, S17,
and S18 after the third RTA treatment (RTA3, 650 C=5 min) was performed very soon after the RTA2 treatment. Another surprising and remarkable redshift of 93 nm and the decrease in the FWHM of the PL spectral peak (DFWHM ¼ 10 meV) were observed in the case of the S17 sample. Interestingly, the RTA3 treatment of the S17 sample returned its PL spectral peak wavelength and the FWHM values to nearly the same peak wavelength and FWHM values measured from the RTA3-treated S sample (in other words, the significant RTA2-induced additional BS of the S17 sample as compared to the S sample almost vanished after the RTA3). During the RTA3 treatment, the PL peaks of the S and S18 samples blueshifted only marginally, 2 3 nm, and the FWHM remained essentially unchanged. It appears from Figs. 3 and 2 that the RTA2 and RTA3 treatments activated more pronouncedly the same microscopic mechanisms as did the RTA1 treatment and the ageing process at the temperature of 20 C, respectively. In particular, we notice that the blueshift after RTA2 treatment is noticeably larger than the blueshift after the RTA1 treatment. Furthermore, it is interesting to see that the RTA3 treatment, which was carried out soon after the RTA2 treatment, caused such amount of redshift and decrease in the FWHM capable of making the PL spectra of the S17 sample nearly identical to the PL spectra of the S sample upon the same RTA3 treatment. It is worth mentioning that the PL spectra of a sample, similar to the S17 sample, whose N-containing layers were grown at 450 C became almost completely identical to that of the corresponding nonirradiated sample after either the ageing process at room temperature for 4 days or after the RTA3 treatment performed soon after [11] the RTA2. In Fig. 4 are depicted high resolution XRD curves recorded from the samples S, S17 and S18 after irradiation (black lines) and after the RTA3 stage (gray lines). As compared to the corresponding nonannealed sample, the GaInNAs quantum well of each RTA3-treated sample experienced a very tiny decrease in strain whose maximum value was observed for S17 and amounted only 0.011% (corresponding to 24 arcsecs left shift of the XRD feature related to the GaInNAs well). At the same time, no change can be observed in the shape and position of the XRD features related to the GaNAs barriers. This is to say that strain relaxation by means of either misfit dislocations (which would lead to red shift) or elemental
Fig. 3. 300-K PL's plot recorded from samples S, S17 and S18 upon (a) RTA2 and (b) RTA3.
Fig. 4. High resolution x-ray diffraction (XRD) curves recorded from S, S17 and S18 after irradiation (black lines) and RTA3 treatment (gray lines).
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interdiffusion (which would lead to blue shift) across quantum well and barrier interfaces was negligible and cannot account alone for any of the significant blue shifts as well as additional blue shifts and subsequent red shifts encountered in our samples. Hence, the remarkable blue shifts and the additional blue shifts/red shifts observed must be assigned to other causes. Thus, the blue shift seen for the samples S and S18 upon annealing as well as for the sample S17 after the RTA3 stage is mainly due to formation of IneN bonds on annealing at the expense of Ga-N ones, promoted during growth. However, this change on annealing in the atomic composition of the N-centered tetrahedral N Gam In4m clusters (0 m 4) from a Ga-rich composition to an In-rich composition (which generates blue shift) should be excluded as the cause for the additional BS/subsequent RS observed because a change back of the N atom nearest neighborhood from the N-rich to the Ga-rich one (which generates red shift) is very improbable in the equilibrium theory of the alloying formation [4]. The mechanisms responsible for the additional blueshift and subsequent redshift observed in the PL spectra from the S17 sample relative to the PL spectra from the S sample should be assigned to the presence of additional particular type of defects created/influenced by electron irradiation. The fact that the spectral line shifts were always accompanied with substantial changes in the FWHM suggests to us that the defects are mainly in the quantum well and its barriers region, which accommodates the biggest part of the carriers' wave functions involved in the PL process. We further assume that electron irradiation yields metastable defect centers. Metastable centers are known to exist in electron irradiated GaAs [6], lattice-matched Ga1x Inx Ny As1y (xz2:8y) on GaAs [16], and were theoretically predicted to occur in ternary GaAsN alloys [17]. Let us assume that electron irradiation initially produced complex neutral or charged defect centers in the quantum well of the pristine sample. It is then conceivable that heating the sample at 650 C (under the RTA1 or RTA2 treatments) could influence the atomic and electronic structures of these defects and the surrounding host material, for example, by moving nearby nitrogen atoms from their initial substitutional sites to energetically more favorable defect sites [17,18]. Our first-principles studies in the framework of the density functional theory (DFT) [17] shows that when a nitrogen atom in the dilute GaAs1x Nx ternary alloy moves from a substitutional group V site into any of the stable or metastable interstitial sites defined by the nearest Ga or As tetrahedron, then the alloy bandgap will increase causing the blueshift in the PL spectra. Due to the isoelectronic nature of the Ga and In atoms it is then obvious that in the present case of the Ga1x Inx As1y Ny quaternary alloy in the quantum well, the nitrogen atom moving from its subsitutional group V site into the nearest interstitial site defined by a Ga tetrahedron, As tetrahedron, or Gam In4m (0 m 4) tetrahedron, the bandgap will similarly get increased, causing the blueshift in our observed PL spectra. Furthermore, our calculations on the GaAs1x Nx ternary alloy show that, out of many possible interstitial states capable of causing the blueshift, it is the single N impurity low-energy eGN type interstitial defect defined in Ref. [18] which will cause the largest blueshift. Therefore, we tentatively propose that the same type of single N impurity interstitial defect (eGN) is responsible for the blueshift observed in the Ga0:63 In0:37 As0:99 N0:01 =GaAs quantum well, rather than certain GaIn atomic configurations around the substitutional N in the Ga0:63 In0:37 As0:99 N0:01 alloy [5] as the potential-energy transition barrier between these configurational states is obviously quite large. We have further experimental evidence to exclude the possibility of having a substitutional N atom surrounded by disordered Ga-In atomic configurations (N@Ga In) to explain the blueshift.
Namely, as it is well known from the theory of disordered alloys, this substitutional disorder between the Ga and In atoms in the group III sublattice of the Ga0:63 In0:37 As0:99 N0:01 alloy would enhance scattering of carriers involved in the photoluminescence process, therefore decreasing their overall lifetime and consequently broadening the linewidth of the PL spectral transitions. However, this disorder related broadening is not clearly visible in our PL spectra, therefore excluding the possibility that the observed blueshift would be, at least majorly, due to the N@Ga In type defects. It is obvious that large scale ab initio molecular dynamics calculations would be required in order to quantitatively differentiate between these two defect possibilities (N interstitial vs. N substitutional cases) responsible for the blueshift observed in the Ga0:63 In0:37 As0:99 N0:01 =GaAs quantum well. Since we know that the activation energy of the defects in the S17 sample is very low, as evidenced by the instability of this sample during the ageing process at the room temperature of 20 C [cf. Fig. 2(b)], the nitrogen atoms are expected to be only loosely bound to the obvious interstitial defect sites which can then readily hop back to substitutional sites by slight thermal activation. This process would give rise to the redshift. Interestingly, this could explain why the sizes of the redshift and blueshift observed in the photoluminescence spectra are almost the same (jRSjzjBSj) when the complete cycle of activation and deactivation of the electronirradiated defects is carried through. This is the case with the RTA2/RTA3 treatments and most likely also with the previous RTA1/ageing at 20 C processes if the ageing duration would have been prolonged beyond 4 days or the environmental temperature would have been higher than 20 C. Finally we notice that by increasing the electron irradiation dose up to one order of magnitude larger up to 1.5 1018 electrons/cm2 produced additional defects or modified those already produced at lower doses during irradiation process till 1018 electrons/cm2 such that the resulting photoluminescence spectral behavior was essentially different from that observed in the sample S17 irradiated with the dose of 2 1017 electrons=cm2 . The additional defects may interact with the metastable defects produced at lower doses up to 1017 electrons/cm2 or the amount of the metastable defects are so high as they strongly interact each other (they could agglomerate forming clusters) in such way that the potential barrier between their ground and metastable state is so high as they cannot longer undergo a change in their energy states neither in short time by thermal annealing at elevated temperatures nor in longer time by ageing at room temperature. Thus, the formation of IneN bonds remains the only major mechanism for annealinginduced blue shift in the sample S18, explaining why the magnitude of annealing-induced blue shift observed in the sample S18 is similar to that seen for S despite irradiation. Regarding the different behavior of photoluminescence with the applied dose of electron irradiation, it is worth mentioning our recent observation that a 7 MeV electron irradiation on GaInNAs thin layers up to the dose of 1015 electrons=cm2 produced a photoluminescence enhancement whereas one order of magnitude higher dose induced a deterioration of luminescence efficiency of the investigated GaInNAs layers [8]. 4. Summary and conclusions It has been found in our studies that 7 MeV electron irradiation with doses of 2 1017 and 1:5 1018 electrons=cm2 will cause atomic structural damage in the GaInNAs/GaNAs/GaAs quantum wells, probably creating complex defect centers. When the lowerdose sample (exposed to the dose of 2 1017 electrons=cm2 ) was annealed at 650 C, a pronounced additional blueshift of PL was observed, which made the total blueshift much larger than the blue
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shift of a non-irradiated sample. The irradiation-created defects turned out to be unstable, making the additional blue shift remarkably reduce or almost vanish (by a subsequent red shift of photoluminescence), a phenomenon which occurred even at room temperature. The additional blueshift observed in this sample can qualitatively be explained by assuming that there is a hopping mechanism for nitrogen atoms to move from their initial substitutional sites to energetically more favorable defect sites. On the other hand, hopping of nitrogen atoms in the reverse order, from the defect sites to the substitutional sites, will cause the subsequent redshift of the PL spectra. Due to the instability of this defect state under thermal activation, the nitrogen atom could return back to its initial site, restoring the PL peak and line shape very close to their initial values. No such PL instability upon thermal annealing was observed in the case of the higher-dose irradiated sample (exposed to the dose of 1:5 1018 electrons=cm2 ), probably due to inactivation of the irradiation-induced unstable defects by interaction with each other following their increased concentration or with additional defects generated at the higher irradiation dose. Acknowledgments The research leading to these results has received funding from the EEA Financial Mechanism 2009e2014 under the project contract no 23SEE/30.06.2014.
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References [1] W. Li, T. Jouhti, C.S. Peng, J. Konttinen, P. Laukkanen, E.M. Pavelescu, M. Dumitrescu, M. Pessa, Appl. Phys. Lett. 79 (2001) 3386. [2] B. Royall, N. Balkan, S. Mazzucato, H. Khalil, M. Hugues, J.S. Roberts, Phys. Stat. Sol. 248 (2011) 1191. [3] W. Li, M. Pessa, T. Ahlgren, J. Dekker, Appl. Phys. Lett. 79 (2001) 1094. [4] K. Kim, A. Zunger, Phys. Rev. Lett. 86 (2001) 2609. [5] P.J. Klar, H. Gruning, J. Koch, S. Schafer, K. Volz, M. Stolz, W. Heimbrodt, A.M. Saadi, A. Lindsay, E.P. O'Reilly, Phys. Rev. B 64 (2001) 121203. venard, J. Appl. Phys. 64 (1988) R65. [6] J.C. Bourgoin, H.J. Bardeleben, D. Stie t¸eanu, T. Jouhti, V. Cimpoca, M. Pessa, [7] E.-M. Pavelescu, A. Gheorghiu, N. B alt¸a in: Semiconductor Conference, CAS 2004 Proceedings (IEEE, Piscataway, NJ, 2004), vol. 1, 2004, p. 220. [8] E.-M. Pavelescu, R. Kudrawiec, J. Puustinen, A. Tukiainen, M. Guina, J. Luminesc 136 (2013) 347. lt¸ [9] E.-M. Pavelescu, R. Kudrawiec, N. Ba at¸eanu, S. Sp^ anulescu, M. Dumitrescu, M. Guina, Semicond. Sci. Technol. 28 (2013) 025020. [10] E.-M. Pavelescu, A. Gheorghiu, R. Kudrawiec, M. Dumitrescu, A. Tukiainen, T. Jouhti, T. Hakkarainen, N. Tkachenko, V.D.S. Dhaka, H. Lemmetyinen, J. Andrzejewski, J. Misiewicz, M. Pessa, Appl. Phys. Lett. 85 (2004) 6158. [11] E.-M. Pavelescu, M. Dumitrescu, M. Guina, J. Luminesc 154 (2014) 584. [12] F. Dimroth, Phys. Stat. Sol. C 3 (2006) 373. [13] E.-M. Pavelescu, S. Antohe, V. Ruxandra, S. Sp^ anulescu, V. Cimpoca, M. Pessa, in: Semiconductor Conference, 2001, CAS 2001 Proceedings (IEEE, Piscataway, NJ, 2001), vol. 2, 2001, p. 319. [14] X. Yang, J.B. Heroux, M.J. Jurkovic, W.I. Wang, J. Vac. Sci. Technol. B 17 (1999) 1144. [15] S. Francoeur, G. Sivaraman, Y. Qiu, S. Nikishin, H. Temkin, Appl. Phys. Lett. 72 (1998) 1857. [16] J.Z. Li, J.Y. Lin, H.X. Jiang, J.F. Geisz, S.R. Kurtz, Appl. Phys. Lett. 75 (1999) 1899. [17] E. Arola, J. Ojanen, H.-P. Komsa, T.T. Rantala, Phys. Rev. B 72 (2005) 045222. [18] W.K. Loke, S.F. Yoon, S.Z. Wang, T.K. Ng, W.J. Fan, J. Appl. Phys. 91 (2002) 4900.