GaAs quantum dots

GaAs quantum dots

Journal of Luminescence 153 (2014) 109–117 Contents lists available at ScienceDirect Journal of Luminescence journal homepage: www.elsevier.com/loca...

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Journal of Luminescence 153 (2014) 109–117

Contents lists available at ScienceDirect

Journal of Luminescence journal homepage: www.elsevier.com/locate/jlumin

H  ion implantation induced ten-fold increase of photoluminescence efficiency in single layer InAs/GaAs quantum dots R. Sreekumar a, A. Mandal a, S. Chakrabarti a,n, S.K. Gupta b a b

Centre for Nanoelectronics, Department of Electrical Engineering, Indian Institute of Technology Bombay, Mumbai 400076, Maharashtra, India Nuclear Physics Division, Bhabha Atomic Research Centre, Mumbai 400085, Maharashtra, India

art ic l e i nf o

a b s t r a c t

Article history: Received 19 September 2013 Received in revised form 5 March 2014 Accepted 8 March 2014 Available online 16 March 2014

We demonstrate a ten-fold increase in photoluminescence (PL) efficiency from 50 keV H  ion-implanted InAs/GaAs quantum dots (QDs) at a temperature of 8 K and/or 145 K. Enhancement occurred without post-annealing treatment. PL efficiency increased with increasing implantation fluence from 6  1012 ions/cm2 up to an optimum value of 2.4  1013 ions/cm2, beyond which PL efficiency decreased drastically (up to a fluence of 2.4  1015 ions/cm2). Passivation of non-radiative recombination centres (due to direct interaction of H  ions with lattice defects) and de-excitation of photo-generated carriers to QDs through quantum mechanical tunnelling via H  ion-induced defects (e-traps) that are created near the QD–cap layer interface, resulted in PL efficiency enhancement. Shallow e-traps with activation energy  90 meV and 30 meV created near the conduction band of GaAs cap layer for the samples implanted with H  of fluence 6  1012 and 2.4  1013 ions/cm2 respectively are identified using low temperature PL study. Contribution of de-trapped electrons from the e-traps to the QDs enhanced the PL efficiency at 145 K. Cross-section transmission electron microscopy and X-ray diffraction study revealed that the structural damage created by H  ions at the high fluence level of 2.4  1015 ions/cm2, caused the degradation in PL efficiency. & 2014 Elsevier B.V. All rights reserved.

Keywords: PL efficiency InAs/GaAs quantum dots Ion implantation Defect passivation

1. Introduction Devices containing quantum dots (QDs) are emerging as an alternate platform for quantum well (QW) based devices. High efficiency InAs/GaAs QD lasers, photodetectors that show better yield than QW devices, have been demonstrated recently [1,2]. The advantage of these III–V compound, semiconductor-based QD devices is their low threshold current density [3], and the ability to tune their emission wavelength from 1100 nm to 1500 nm, that made them use in infrared detectors [4], focal plane arrays [5] and in the telecom sector [6]. Recent studies have revealed that the tolerance of QD-based devices toward energetic radiation is higher than that of QW devices [7–9], which opened up new applications for such devices in outer space vehicles and satellites, where these devices are exposed to energetic radiations. Much research is being carried out to improve the characteristics of QDs to allow them to reach their theoretically expected device efficiency [10]. One of the major characteristics of III–V compound, semiconductorbased QDs is their photoluminescence (PL) efficiency, which can be advantageous for realising highly efficient, low-threshold QD lasers. To improve the PL efficiency in InAs/GaAs QDs, different techniques

n

Corresponding author. Tel.: +91 22 2576 7421; fax: +91 22 25723704. E-mail address: [email protected] (S. Chakrabarti).

http://dx.doi.org/10.1016/j.jlumin.2014.03.016 0022-2313/& 2014 Elsevier B.V. All rights reserved.

including deuteration [11], hydrogen plasma treatment [12,13], nitrogen treatment of QDs before capping [14], stacking of multilayer QDs [15], QDs grown in dot-in-the-well (DWELL) structures [16] and rapid thermal annealing [17] have been used. Most of the techniques described above result in undesirable spectral emission shifts that are due to the intermixing of atoms between the QDs and the barrier material. Hence, a technique that does not promote inter-diffusion and undesirable spectral shifts would benefit QD-based devices, such as vertical cavity surface emitting lasers (VCSELs) and resonant cavity light emitting diodes (with fewer QDs in the active region) that demand high-yield PL and low spectral emission tolerance. Ion beams have been used to modify the structural and optoelectronic properties of materials for several decades [18– 20]. Precision control over the size and depth of damage or modification created by ions is well reported, and is utilised in the semiconductor industry. Extensive studies have been carried out to understand the effect of ion beams in insulators, metals and semiconductors. However, few reports have been published on the effect of ion beams in QDs. Sreekumar et al. [21] studied the effect of heavy S  ion implantation on InAs/GaAs QDs and demonstrated the degradation of luminescence properties induced by the structural damage created by heavy sulphur ions. In contrast, recently Sreekumar et al. [22] demonstrated 7 times enhancement in PL efficiency in single layer InGaAs/GaAs QDs induced by proton

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irradiation (3–5 MeV) without any post annealing treatment following the report by Leon et al. where they demonstrated slight enhancement of PL emission at 5 K using 1.5 MeV proton irradiation on single layer InGaAs/GaAs QDs [7,23]. Later, Lu and coworkers demonstrated enhancement of PL efficiency in multilayer InAs QDs on proton implantation (50–70 keV) followed by rapid thermal annealing at 700 1C [24,25]. In this study, we report an unusual response and enhanced PL efficiency (10 fold at 145 K, with a spectral shift of 15 meV) from H  ion implanted single layer InAs/GaAs QDs without a post-annealing treatment. To explain the observed PL enhancement we performed structural (X-ray diffraction and cross-sectional transmission electron microscopy), optical (low temperature PL) analysis and proposed a plausible model to explain the PL mechanism taking place in the InAs/GaAs QD hetero-structure.

2. Experimental Single layer InAs/GaAs QDs were grown over a semi-insulating GaAs substrate in Stranski–Krastanov growth mode using the solid source molecular beam epitaxy method. After desorbing the protective oxide layer, a 0.5 mm intrinsic GaAs buffer layer was grown at 590 1C. Subsequently, the temperature was brought down to 500 1C and a thin intrinsic GaAs layer of 0.1 μm was deposited. Thereafter, 2.7 ML of InAs was deposited (500 1C), which formed the wetting layer and gave rise to self-assembled QDs of density 1.84  1010/cm2 (deduced from atomic force microscopy – Veeco Digital NanoScope IV). The QDs were further capped (capping/ barrier layer) by a 0.1 μm intrinsic GaAs layer, over which surface QDs were grown under similar growth conditions (to determine the density and the morphology of the QDs). The deposition rates for GaAs and InAs were maintained at  0.72 mm/h and  0.2 ML/s, respectively. Fig. 1 shows the schematic of the epitaxial grown InAs/ GaAs QD hetero-structure on the semi-insulating GaAs substrate. The InAs/GaAs QDs were implanted with H  ions of energy 50 keV and fluence varying in the range of 6  1012–2.4  1015 ions/cm2 (at room temperature). To minimise the channelling effect, implantation were performed at 101. Samples were removed from the implantation chamber and underwent atomic force microscopy (AFM) study (Veeco Digital NanoScope IV) at atmospheric pressure. Structural characterisation was obtained by using X-ray diffraction (XRD) analysis (Philips PANalytical X'Pert Pro X-ray diffractometer) and cross sectional transmission electron microscopy (XTEM; Philips EM420). For XTEM measurements, samples were prepared using conventional mechanical polishing and ion milling techniques. The XTEM micrographs were recorded under an acceleration voltage of 200 kV and with a lattice resolution of 0.3 nm. The PL measurements were carried out over a range from 8 K to room temperature by

Fig. 1. Schematic of the InAs/GaAs hetero-structure under study.

Fig. 2. PL spectra recorded at 145 K from as-prepared and H  ion implanted samples at various fluence levels recorded at a laser excitation power of 51.6 W/ cm2.

exciting the samples with a diode laser operating in continuouswave mode at a wavelength of 405 nm. Power dependent PL spectra were recorded by varying the excitation density over the range 6.5– 51.6 W/cm2. PL emission was dispersed by a 0.75 mm triple-grating monochromator and detected using liquid nitrogen cooled 0.2 μm InGaAs detector array. Electronic and nuclear energy losses of 50 keV H  ions in InAs/GaAs systems are  13.67 eV/Å and 0.05 eV/Å respectively, while the range is about  0.36 mm (determined using transport of ions in matter (TRIM) calculations) [26]. TRIM-based simulation (depth-wise) was performed in a layer sequence to study the damage created in the system under study.

3. Results and discussion Fig. 2 depicts the PL spectra (at 145 K) for as-prepared and from GaAs/(InAs/GaAs) QD hetero-structures implanted with 50 keV H  ions at fluences ranging from 6  1012 to 2.4  1015 ions/cm2 (laser excitation density of 51.6 W/cm2). The as-prepared sample exhibited PL emissions centred at 1191 nm and 1122 nm, corresponding to the ground state and first excited state, respectively. No emission from wetting layer is observed from the samples under study. A clear enhancement in PL emission is observed on H  ion implantation up to an optimum fluence of 2.4  1013 ions/cm2. Samples implanted with a fluence of 6  1012 ions/cm2 exhibited a 5-fold increase in PL efficiency (ratio of overall integrated PL intensity of implanted sample PLI to that in the as-prepared sample PLas), whereas samples implanted with 2.4  1013 and 7.2  1013 ions/cm2 exhibited PL efficiency enhancement levels of 10-times and 3.5-times, respectively. The increasing behaviour of PL efficiency with implantation fluence indicates that H  ion implantation could passivate non-radiative recombination defects that were created during growth of the hetero-structure, up to a certain fluence level (2  1013 ions/cm2). The optimum fluence of 2.4  1013 ions/cm2 observed to enhance the PL efficiency in the present study could be explained as: the defects (interstitial) in the system created during the growth process might have been in the order of 2.4  1013 ions/cm2; thus, the direct interaction of H  ions with these defects imparted energy to the interstitial atoms in order to occupy the original/respective lattice positions in the lattice. By increasing the implantation fluence higher than 2.4  1013 ions/cm2, passivation of the defects is complete and initialised the creation of additional defects (displacement/vacancies) that can trap photo-generated carriers in the system. This behaviour is clearly evident in the PL analysis (Fig. 2). Lu et al. [24] observed a similar type of enhancement in PL efficiency up to

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Fig. 3. PL spectra recorded at 5 K from the samples implanted with H  ions with laser excitation power (a) 51.6 W/cm2 and (b) 5 W/cm2.

80-fold induced by proton implantation (followed by rapid thermal annealing) from a multi-layer InAs/GaAs QD structure. They attributed the PL enhancement as a combined effect of defect passivation and increased capture rate of photo-generated carriers from GaAs cap layer/barrier layer due to proton implantation. In addition to the increase in PL efficiency, H  ion implantation resulted in a shift in PL emission to a shorter wavelength; i.e., the ground state emission showed a blue shift from 1191 nm to 1175.45 nm. It is interesting to note that the sample which showed maximum PL enhancement only showed a blue shift in PL emission about  14 meV (fluence – 2.4  1013 ions/cm2). The blue shift could be due to one or two phenomena: inter-diffusion of atoms across the InAs/GaAs QD-GaAs cap layer interface [25] or due to the compressive stress induced by the cap layer on the InAs/GaAs QDs [27]. Fig. 3a depicts the PL spectra recorded at 8 K from samples implanted with H  ions of various fluences (laser excitation density of 51.6 W/cm2). An enhancement of PL efficiency with H  ion implantation fluence is observed as obtained at 145 K. However the enhancement of PL efficiency for the sample implanted with optimum fluence of H  ions was about 2-times with respect to the as prepared sample. To extract information on non-radiative recombination centres, it is a good practice to record PL spectra at a low laser excitation density; a density at which the non-radiative recombination centres are more pronounced as lesser number of carriers are photo-generated in the system. To demonstrate the same, PL spectra were recorded at an excitation density of 5 W/cm2 (at 8 K) from the as-prepared and H  ion implanted samples (Fig. 3b). This excitation density was selected so that the as-prepared sample just started exhibiting first excited state emission. A similar trend of increasing PL efficiency with increasing implantation fluence was detected, even at a low excitation density. The overall PL efficiency increased with fluence and at the optimum fluence of 2.4  1013 ions/cm2 reached a level that was  9 times greater than that of the as-prepared sample; PL efficiency reduced on further increases in fluence. This experiment clearly showed that H  ion implantation can passivate nonradiative recombination centres in InAs/GaAs systems. We conducted excitation density dependent PL study of the obtained samples at 8 K by varying the excitation density from 6.5 W/cm2 to  51.6 W/cm2. Fig. 4 shows the plot of integrated PL intensity versus laser excitation density for the as-prepared and samples implanted with various fluences of 50 keV H  ions. Each plot is normalised to its respective integrated PL intensity, recorded at an excitation density of 6.5 W/cm2. The rate of increase in integrated PL intensity with excitation density for the samples implanted with 6  1012 and 2.4  1013 ions/cm2 fluence were higher than that of the as-prepared sample (Fig. 4). Whereas on increasing

Fig. 4. Variation of integrated PL intensity at 8 K with laser excitation density.

the fluence to 7.2  1013 ions/cm2, the rate of increase in PL intensity decreased, reaching a value comparable to that of the as-prepared sample. With an implantation fluence of 2.4  1015 ions/cm2, the rate of increase in PL intensity was markedly reduced. The observed phenomenon is explained as follows: barrier layer plays a vital role in the PL properties of InAs/GaAs QD hetero-structure. When the hetero-structure is optically excited, carriers are generated not only in the QDs, but also in the barrier layer, which serves as a reservoir of photo-generated carriers. These photo-generated carriers in the GaAs barrier layer are subsequently diffused towards a small capture volume in and around the QDs from which these carriers are captured into the QDs to undergo radiative recombination/relaxation to the ground state. Defects created in the barrier layer during the growth process or that are intentionally created can modify the carrier dynamics in the barrier layer which in effect alter the PL efficiency of the QDs. In our case, we observed an enhanced rate of PL efficiency in the samples implanted with fluence 6  1012 and 2.4  1013 ions/cm2 with respect to the as-prepared sample. Implantation with H  ions passivated the defects (non-radiative recombination centers) in the barrier layer and/or near the barrier layer/QD interface. This increased the effective carrier concentration in the barrier layer thereby enhancing the carrier capture rate into the QDs to engage in radiative recombination. Passivation of defects in effect enhanced the non-radiative recombination time, which also favoured the survival of the photo-generated carriers. Popescu et al. [28] showed the relationship between th PL quantum yield and the effect of non-radiative recombination rate (Rnr) in the wetting layer and the thermal escape rate ðRD-W Þ of carriers from

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dots to wetting layer in a InAs/GaAs QD hetero-structure considering an ‘ambipolar’ approach, which is given by, PL quantum yield;

Rr N o 1 ¼ 1 þ ðRnr =C W-D Þð1 þðRD-W =Rr ÞÞ Ge

ð1Þ

where Rr is the radiative recombination rate in the QDs (electron– hole recombination), No the electron population in the QDs and Ge the number of carriers captured per unit time by the wetting layer from the barrier layer. C W-D is the capture rate of carriers in to QDs from the wetting layer. A decrease in Rnr as a result of passivation of non-radiative recombination centres/defects by H  implantation can in effect enhance the overall PL yield or the radiative recombination in the QDs. Passivation of defects induced by proton implantation and subsequent enhancement of non-radiative recombination time and carrier capture rate by InAs QDs has been reported by Lu et al. [24]. Increasing the H  ion fluence beyond the optimum value created additional structural damages/defects in the barrier layer/QD which captured the photo-generated carriers from the barrier layer (and QDs) even though the carrier population in the heterostructure were un-changed. This reduced the non-radiative recombination time and effective carrier concentration in the barrier layer, which resulted in the slower rate of PL intensity from samples implanted with higher fluence. Temperature dependent PL measurements of the samples were carried out from 8 K to room temperature with a laser excitation density of 51.6 W/cm2. Arrhenius plots for the low temperature PL measurement data (Fig. 5) clearly show an increase in integrated PL intensity with temperature, reaching a maximum at 145 K for all the implanted samples (except sample implanted with 2.4  1015 ions/cm2) whereas a decrease in PL integrated intensity is shown in the as-prepared sample. Activation energies Ea1 and Ea2 of the samples were determined by measuring the slope of the straight line regions in the Arrhenius plots of the PL data for temperature ranges 50–150 K (Ea1) and 150–270 K (Ea2). As-prepared samples exhibited activation energies of Ea1 ¼3 meV and Ea2 ¼312 meV. We attribute the activation energy Ea1 ¼3 meV is due to the defects that are created during the growth process near the cap layer–QD interface, that can act as non-radiative recombination centers, whereas Ea2 ¼312 meV as to the thermal escape of carriers from QDs to the wetting layer. On implantation with H  ions at a fluence 6  1012 ions/cm2, defects with activation energy Ea1 ¼3 meV were eliminated (non-radiative recombination centres) and the samples exhibited activation energies Ea1 ¼91.6 meV and Ea2 ¼312 meV. Increase of PL intensity on H 

ion implantation (6  1012 ions/cm2) supports the fact of elimination defect level with Ea1 ¼3 meV. Samples implanted with 2.4  1013 ions/cm2 showed Ea1 ¼30.5 meV and Ea2 ¼ 315 meV. We attribute the activation energies Ea1 of 30.5 and 91.6 meV are due to As vacancies (VAs) created by H  implantation in the GaAs cap layer for samples implanted at fluences of 2.4  1013 ions/cm2 and 6  1012 ions/cm2, respectively. The activation energy Ea1 ¼ 30.5 and 90 meV obtained from the H  implanted sample is in accord with the defect levels of As vacancies (VAs) observed in GaAs [29–32]. These defects VAs acts as e-traps and contributed the trapped electrons to the conduction band (of GaAs) on thermal activation [30]. Corbel et al. [29] proposed defect levels 30 meV and 90 meV situated near the conduction band of GaAs corresponds to two charge state transitions of As vacancies from V2As to   VAs and VAs to V0As respectively. We attribute this e-trap as level-A in the conduction band of GaAs cap layer which is formed due to QDs with special conduction band structure at the InAs/GaAs interface induced by local potential deformation as a result of H  ion implantation. On increase of implantation fluence to 7.2  1013 ions/cm2 resulted in decrease of activation energy Ea1 ¼ 12 meV (defect level-A) and Ea2 ¼272 meV. Similarly an increase in fluence to 2.4  1014 and 2.4  1015 ions/cm2, resulted in further decreases of activation energy Ea2 to 229 and 138 meV, respectively (Ea1 ¼2 meV for both samples). Correspondingly, a shift in the offset of the exponential PL quenching to a lower temperature is observed for the samples implanted with 2.4  1014 and 2.4  1015 ions/cm2, whereas for samples implanted with H  ions up to a fluence of 2.4  1013 ions/cm2 exhibited a shift in PL quenching towards higher temperatures (Fig. 5). The shifting of the offset of the exponential PL quenching towards higher temperature is an indication of defect annihilation in and around QDs and at the QD–cap layer interface [28,33]. It is important to note that utmost care was taken during the low temperature PL measurements to keep the laser excitation power constant for the entire set of samples under study; thereby taking into account findings in works by Le Ru et al. [33] and Patane et al. [34] in which the authors claim that the activation energy deduced from Arrhenius plot is strongly dependent on the excitation power. The probable reason for reduction in Ea1 ¼ 2 meV might be the depopulation of carriers from QDs, through thermally assisted tunnelling, to the defects introduced by H  ion implantation with higher fluence [35]. The temperature dependent increase in integrated PL intensity in the range 50–140 K for the samples implanted with H  ions can be explained as follows: on photo excitation carriers are not only excited into the QDs, barrier layer/wetting layer but also trapped by the level-A. The carrier diffusion time ‘τ’ from the actual photoexcitation position in the GaAs barrier layer towards the QDs, which are separated by a distance L is given by [36] τ¼

L2 Da

where Da is the ambipolar diffusion co-efficient defined as   μ μ 2K B T Da ¼ e h e μe þμh

Fig. 5. Arrhenius plots of as-prepared samples and samples implanted with 50 keV H  ions of fluences 2.4  1013 ions/cm2 and 2.4  1015 ions/cm2 (Io is the integrated PL intensity recorded at 8 K).

ð2Þ

ð3Þ

where me(h) is the electron (hole) mobility, KB is the Boltzmann constant, e the electron charge and T is the temperature. On increasing the temperature, the carriers trapped in the level-A gets sufficient energy to reach conduction band of the barrier level. As the carriers are re-emitted to the GaAs barrier layer, the effective carrier concentration/density in the barrier layer increases thereby, enhancing the effective carrier diffusion length. This delays the non-radiative recombination time and results in the slowing down of the capture time as the carriers initially trapped in the level-A starts to move towards the QD capture

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volume, attaining a maximum occupation at a large time delay. As the temperature is increased further, the level-A no longer able to trap the carriers, causing all the carriers in the barrier layer to diffuse towards QD region. The effective carrier diffusion length and carrier concentration/density is much larger than that at low temperature and hence leads to faster capture into the QDs, resulting in high PL intensity around 140 K. In addition, another phenomenon that can increase the PL intensity with temperature has been proposed by Popescu et al. [28]. They suggested that the interface between the dots and the barrier layer will not be compositionally abrupt, which develops a strain field between them during the growth process. This strain field creates a potential barrier between the interfaces which block the photocarriers to capture by the dots at low temperature. As the temperature increases the carriers acquires energy to overcome the energy barrier created at the barrier layer–QD interface, which promote the capture rate by the QDs, thereby increasing the PL efficiency [28]. Whereas the decrease in PL intensity observed at temperature higher than 145 K is due the combined effect of (i) reduced capture rate by the QDs; (ii) thermal activation of carriers from the QDs to the GaAs barrier and (iii) non-radiative recombination within the QDs. At higher temperatures the level-A (e-trap) cannot trap electrons and contribute extra carriers to the conduction band of the barrier layer, as the density of the trapped carriers are very much less than that at 10 K (lower temperatures). Furthermore the non-radiative recombination of the thermally escaped carriers from the QDs in the barrier layer and within QDs results in reduction of carriers in the QDs to engage in radiative recombination. On the other hand, the enhancement in PL efficiency at 8 K is attributed to the tunnelling of trapped electrons from the e-traps (level-A) to the QDs. The sample implanted with H  with a fluence of 2.4  1013 ions/cm2 which has an Ea1 ¼30.5 meV exhibited high PL efficiency as compared with the sample that is implanted with a fluence of 6  1012 ions/cm2 (Ea1 ¼90 meV). It is obvious that the potential of the e-trap with Ea1 ¼30.5 meV will be lower than that with Ea1 ¼90 meV and hence the probability of tunnelling of electrons from e-traps having lower potential (close to the QDs) will be higher to the QDs, that resulted in higher PL efficiency. Fig. 6 shows the XRD pattern of samples implanted with H  ions at fluences of 6  1012–2.4  1015 ions/cm2. XRD peaks centred at 2θ¼ 66.12 and 31.69 along the (4 0 0) and (2 0 0) planes, respectively, are from the GaAs substrate, while that at 2θ¼ 72.14

is from the GaAs cap layer [37]. The reflection from InAs QDs along (4 0 0) plane could also be identified. Implantation with H  ions at fluences of 6  1012 and 2.4  1013 ions/cm2 resulted in improvement in x-ray diffraction from InAs QDs (a decrease in full width at half maximum (FWHM) was observed, Fig. 7) as the H  ions passivate the defects from the InAs/GaAs QDs created during the growth process [38] whereas the XRD from the GaAs cap layer was degraded (an increase in FWHM, Fig. 7). Similarly the x-ray diffraction from (2 2 2) plane of GaAs also degraded on H  implantation. A further increase in fluence to 7.2  1013 ions/cm2 reduced the reflections from InAs/GaAs QDs and GaAs cap layer and at a fluence of 2.4  1015 ions/cm2 resulted in elimination of diffraction peaks from the GaAs cap layer and from the QDs, indicating a reduction in crystalline quality of the hetero-structure upon H  ion implantation. The variation in FWHM with fluence clearly indicates the introduction of structural damages in the hetero-structure at high implantation fluence. At a fluence level lower than a critical value, implantation creates point defects, which can be As and Ga vacancies, interstitials, antisites and/or complexes of these defects. Overlapping of these defect complexes can result in an amorphous layer formation [39]. With an increase in implantation fluence, the total energy deposited by the H  ions also increases; thus the overlap of these defects increases. In the present study, the XRD peak in the GaAs cap layer disappeared at a fluence of 2.4  1014 ions/cm2 (Fig. 6e), and a subsequent increase of fluence to 2.4  1015 ions/cm2 resulted in the absence of reflection from GaAs buffer layer (Fig. 6f). Perhaps at a lower H  ion fluence, overlapping of defects was initiated at the surface layer, and, on a subsequent increase of fluence, the overlapping was extended to a depth in and around the InAs QDs, where the dots are embedded. There are reports on defects induced by protons in GaAs system [40,41] and even amorphization of an InAs/GaAs system using 200 keV argon, 300 keV selenium, zinc, cadmium and 280 keV nitrogen ions [39]. Bench et al. [42] demonstrated that amorphous zones can be created in GaAs system with a total nuclear energy density of 1.72  1019 keV/cm3 and 1.13  1019 keV/ cm3 using 50 keV xenon and krypton ions, respectively, at room temperature (fluence of 6  1011 ions/cm2). In our study, the total nuclear energy deposited by 50 keV H  ions at a fluence of 2.4  1015 ions/cm2 was  1.2  1019 keV/cm3, which is within the range of critical nuclear energy density required for the amorphization of InAs/GaAs system [42]. Of particular note here is that the XRD pattern still exhibits a diffraction pattern from the GaAs substrate. This suggests that the amorphous layer is only created to a depth of 360 nm (calculated using TRIM) from the surface of the hetero-structure. Fig. 8 depict the X-TEM images of surface QDs recorded from the sample implanted with various fluences of 50 keV H  ions. As the implantation fluence increases the height of the surface QDs

Fig. 6. X-ray diffraction pattern of as-prepared and samples implanted with H  ions at various fluences.

Fig. 7. Variation of FWHM of X-ray reflections from InAs/GaAs QDs and the GaAs cap layer with different implantation fluence of 50 keV H  ions.

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Fig. 8. Cross sectional TEM images of surface quantum dots (a) as-prepared, sample implanted with H  ions of fluence (b) 6  1012 ions/cm2, (c) 2.4  1013 ions/cm2 and (d) 2.4  1015 ions/cm2.

decreases and lateral size increases a bit with respect to the as-prepared sample. This is probably due to the inter-mixing taken between the surface QDs and the GaAs cap-layer interface. A similar behaviour is observed in the case of embedded dots on H  ion implantation. From the XRD analysis it is observed that the H  ion implantation modifies the structural/crystalline property of the QDs and cap layer. The modification made by the H  ions in the cap layer could exert stress on the QDs. Stress induced by the cap layer and the inter-diffusion taking place between the QD-cap layer interfaces resulted in the blue shift of the PL emission on ion implantation. The embedded dots were of regular shape (Fig. 9a and b) in the case of as-prepared sample whereas on implantation with H  ions, the regular shape is lost as a result of inter-diffusion took place between the QDs-cap layer interface and strain induced by the cap layer on the QDs (Fig. 9c and d). The signatures of the damages created by H  ion implantation at high fluences are observed in the form of fractured QD and deformation of the QD shape. An obvious strain contrast could be visible near the QDs. Moreover, creation of damages in the QDs and cap layer also reflected in the PL emission as the PL efficiency drastically reduced on H  ion implantation at higher fluences. XRD analysis showed improvement of crystalline quality of the QDs (reduction in FWHM of the XRD peak up to an optimum fluence), however from TEM

analysis no passivation of defects are visible in the range of magnification of the TEM images recorded from the samples here studied. At higher fluence strain contrast/damages created by H  ion implantation could be seen in the cap layer and near the QDs. Moreover, at a higher fluence of 2.4  1015 ions/cm2, one could see that the shape of the dots is completely lost as it diffused within the barrier layer (Fig. 9e and f). This could be one of the reasons for the reduction of PL intensity (reduction in number of dots), which is also reflected in the XRD spectra of the sample implanted with H  ions of fluence 2.4  1015 ions/cm2 (Fig. 6f). Fig. 10 depicts a depth-wise TRIM simulation performed in layer sequence to study the damage created in the GaAs/(InAs/GaAs) hetero-structure. Fig. 10 shows clear evidence that maximum damage was created at the GaAs cap (layer 2) and near the cap layer–InAs QD interface; damages were also created in the GaAs buffer layer (layer 4). Difference in the lattice constants of InAs QD and GaAs cap layer could exert compressive stress on InAs QDs which can modify the energy gap of the QDs. For simplicity we consider the dots are of spherical shape. Within the dot, the strain component and pressure can be expressed as,  err ¼ 

 P 3λ þ2μ

ð4Þ

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Surface dots

Embedded dots Embedded dots

Damaged zone

Damaged dots

Damaged zone/ strain contrast

Strain contrast or damage

Damaged dots

Damaged dots Fig. 9. TEM images of (a and b) as-prepared samples and sample implanted with 50 keV H  ions at fluences of (c and d) 2.4  1013 ions/cm2 and (e and f) 2.4  1015 ions/cm2.



  4ð3λ þ 2μÞðα  1Þ ð4α þ 8Þ

ð5Þ

where λ, μ are the Lame's constant and α is the ratio of lattice constants of InAs and GaAs. Lattice constant of InAs QDs and the GaAs cap layer is determined from the XRD analysis and the ratio α is calculated for the samples implanted with various fluence of H  ions. Then the change in conduction and valence band edges are δEc and δEv respectively, where δEc;v ¼ ac;v ð3err ÞðeV Þ

ð6Þ

where ac and av are the deformation potentials of InAs dots (ac ¼  5.08 eV, av ¼ 1 eV) [43]. Applying these values in Eq. (6), we obtain an energy gap change (δE¼δEc  δEv) of 0.545 eV in the case of as-prepared sample. For the sample implanted with H  ions of fluence 6  1012, 2.4  1013 and 7.2  1013 ions/cm2, the change in energy gap were δE¼0.559, 0.561 and 0.5585 eV respectively. So the shift in emission energy (blue shift) deduced from the Eq. (4) is about 14.23, 16.25 and 13.46 meV respectively for the samples implanted with H  ions of fluence 6  1012, 2.4  1013 and 7.2  1013 ions/cm2, whereas, the blue shift recorded from the PL spectra (at 8 K) is about 15.52, 17.01 and 18.27 meV. However the

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blue shift calculated using Eq. (6) is lower, which suggests that the inter-diffusion of atoms across the QD–cap layer interface also contributing for blue shift. Fig. 11a and b shows a pictorial representation of the proposed PL mechanism taking place in the samples under study before and after implantation respectively. Photo-generated carriers are excited from the valence band of the InAs/GaAs QDs, wetting layer and the GaAs cap layer to the conduction band by laser excitation (Process A). The photo-generated carriers from the cap layer, having sufficient energy to overcome the potential barrier created due to the strain at the QD–cap layer interface, were first captured by the excited state of the QDs (Process C) through the wetting layer (Process B) and then relaxed to the ground state (Process D) to engage in radiative recombination (processes E and F). Non-radiative recombination centres present at the QD–cap layer interface (ENR1 level) and in QDs (ENR2 level) destroy the photogenerated carriers from the GaAs cap layer and the QDs, respectively, via non-radiative recombination (processes I and J). This reduces the overall PL emission. As evident from the low excitation density (5 W/cm2) PL measurements, the H  implantation favoured the passivation of the non-radiative recombination centres (ENR1 and ENR2 levels). This allows a greater number of photo-generated carriers to survive to engage in radiative recombination. The H  ion

Fig. 10. TRIM simulation (performed in layer sequence) showing the damages/ vacancies produced in the GaAs/(InAs/GaAs) hetero-structure under 50 keV H  ion implantation. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

implantation resulted in formation of shallow e-traps at the QD–cap layer interface and/or in the GaAs cap layer (level A), which is evident from the damage region near the QD–cap layer interface XTEM image (Fig. 9). Photo-generated carriers are excited from the valence band to the level-A (Process K) and/or can be trapped from cap layer, and are then relaxed to the QDs via quantum mechanical tunnelling (Process T). The potential of the e-traps created via H  implantations are lower than that of the non-radiative recombination centers and hence the probability of tunnelling from these e-traps to the QDs is high. Ultimately, both the passivation of the non-radiative recombination centres and the de-excitation of the trapped electrons to QDs result in enhancement of PL efficiency. At elevated temperatures, thermalisation of trapped electrons (de-trapping) from level A to the conduction band of the GaAs cap layer occurs, which facilities an increase in carriers in the GaAs cap layer. These carriers are captured and undergo radiative recombination from the InAs QDs. This process is responsible for the increase in PL efficiency at elevated temperature. Similar implantation studies were performed using 50–80 keV protons on multi-layer InAs/GaAs QDs by Yalin et al. [25] and Lu et al. [24] in which they observed an enhancement in PL efficiency only after a rapid thermal annealing treatment at temperature 550–700 1C. Moreover the behaviour of PL emission only after proton implantation is not discussed. The significant contribution of the present work compared with others [24,25] is that the enhancement in PL efficiency is attained without any post annealing treatment and without any considerable blue shift in PL emission [45,46]. This is of potential importance in the perspective of industrial scale production of devices. The disadvantages of rapid thermal annealing treatment are: (i) strongly affect the emission efficiency [44] due to the creation of non-radiative centers, (ii) blue shift in PL emission caused by the intermixing at QD–cap layer interface and (iii) excessive loss of arsenic from the hetero-structure. The blue shift of about 40–100 meV was observed in the case of sample annealed at temperatures 550 and 700 1C [25]. On the other hand, hydrogen plasma treatment has been used to enhance PL efficiency from InAs/GaAs QDs without any post annealing treatment and without considerable shift in PL emission wavelength [12,13]. However, the hydrogen plasma treatment was done (using Kaufman source) at a temperature of  300 1C (for 1 h), and it was observed that the benefit of hydrogen plasma treatment was lost on post annealing higher than 300 1C for 5 min [13]. This limits the hydrogen plasma treatment technique, if the temperatures for device processing

Fig. 11. Band diagram shows a proposed model to explain the PL mechanism takes place in the hetero-structure (a) before and (b) after H  ion implantation, and (c) the mechanism taking place under an elevated temperature condition.

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are high. Moreover, the control over the implantation depth of hydrogen implantation is also limited. The present study reveals the scope of using an ion beam technique to improve the optoelectronic properties of InAs/GaAs QDs without the creation of a considerable spectral shift. Such a technique could be employed to improve the performance of VCSELs and resonant cavity LEDs that have small active QD regions (typically 20 μm diameter) and demand high yield PL and low spectral emission tolerance. One of the advantages of this ion beam technique is the ability to control depth and size of modifications to the material's properties. This is important as techniques like rapid thermal annealing can result in undesirable modifications to device structures. As a future work, we are extending this H  ion implantation work to QD-based device structures. 4. Conclusions Here, we demonstrate an enhancement of PL efficiency up to 10 fold in a single layer GaAs/(InAs/GaAs) QD hetero-structure at 8 K and/or at 145 K. This efficiency is achieved in the absence of postannealing treatment by using 50 keV H  ion implantation. The optimum fluence at which maximum PL efficiency was attained using 50 keV H  ions was 2.4  1013 ions/cm2. Introduction of structural damage on increasing the fluence beyond an optimum value resulted in degradation of the PL efficiency. The increase in PL efficiency is attributed: (i) to the passivation of non-radiative recombination centres present at the GaAs cap layer–QD interface, wetting layer and in the QDs; and (ii) to de-excitation of carriers from the GaAs cap layer to the QDs via H  ion implanted induced defects through quantum mechanical tunnelling. Interaction of H  ions with the defects is seen as the most probable reason for the passivation of non-radiative recombination centres in the system. The shift in the onset of PL quenching (Arrhenius plot) towards a higher temperature and the higher PL efficiency exhibited by the H  implanted sample at a lower excitation density, supported the concept of passivation of defects in and around the QDs. Detrapping of carriers from H  ion implantation induced shallow defects via thermal activation to the conduction band; subsequent capture by QDs to engage in radiative recombination resulted in the increase of PL efficiency at 145 K. Acknowledgements The authors are grateful to Arindam Basu, S P Sarode and N B V Subrahmanyam, LEAF facility, BARC for the co-operation during the beam time. We acknowledge DST, Govt. of India for financial support. Partial funding was also obtained from MCIT, Govt. of India through the Centre of Excellence in Nanoelectronics, IIT Bombay and from MNRE, India through NCPRE, IIT Bombay. The Central SPM Facility of IIT Bombay was used for the AFM study. In addition, the European Commission is acknowledged for partial funding through Contract SES6-CT-2003-502620 (FULLSPECTRUM). References [1] A. Salhi, L. Fortunato, L. Martiradonna, M.T. Todaro, R. Cingolani, A. Passaseo, M. De Vittorio, Semicond. Sci. Technol. 22 (2007) 396. [2] S.Y. Lin, Y.J. Tsai, S.C. Lee, Jpn. J. Appl. Phys. 40 (2001) L1290. [3] P.G. Eliseev, H. Li, A. Stintz, G.T. Liu, T.C. Newell, K.J. Malloy, L.F. Lester, Appl. Phys. Lett. 77 (2000) 262.

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