Direct bandgap photoluminescence from n-type indirect GaInP alloys Download: 710次
1. INTRODUCTION
Large bandgap alloys have potential use for top junction in tandem solar cells and yellow-green light emitting diodes [1,2]. However, the alloy becomes an indirect band alloy when its band energy is above 2.2 eV. At the direct–indirect band crossover, the Ga content of is within a range between 0.69 and 0.72 [3]. Due to the weak luminescence emission, studies of the optical properties of alloys near the direct–indirect band crossover and in the indirect band region are more challenging than for direct band alloys. In addition, the lack of lattice-matched substrates and ordering effects adds more complexity to the studies.
Previous studies have provided detailed information on the electronic structures of alloys () [4,5]. The experimental results showed that the luminescence efficiency of -based LED degrades significantly for emission wavelengths shorter than 590 nm (2.1 eV) [6]. Despite carrier confinement issues for large bandgap alloys [7,8], a simplistic approach was developed to model the luminescence intensity degradation as a function of the energetic separation between direct and indirect bands [9]. Based on this model, as the Ga content x increases for direct band to approach a larger bandgap energy, the energetic separation between the and the X minima, , becomes smaller. Because the density of states (DOS) of the X valley is about 10 times larger than that of the valley of (), the probability of electrons occupying the indirect band minimum is high. Therefore, the luminescence efficiency drops exponentially for alloys near the direct–indirect crossover, and indirect band alloys have poor luminescence efficiency. Although the indirect band alloys have large direct band energies, they are not suitable for light-emitting applications, as indirect band transitions dominate.
More recently, the concept of “pseudo-direct” bandgap has been developed for Ge by using tensile strain and high n-type doping to promote the direct-band transition [10]. The tensile strain reduces the energetic separation between and L valleys in Ge, and carrier injection into the valley was increased. The high n-type doping increased band filling in the L valleys to promote the thermal excitation of electrons into the valley. As a result, strain and doping effects enhance the direct band transition in Ge. The direct band transition dominates the light emission at room temperature, and Ge-based LEDs and lasers have been demonstrated [10,11]. For the indirect alloys near the direct–indirect band crossover, n-type doping is also expected to enhance the direct band transition due to similar mechanisms.
In this paper, we study the n-type Te doping effects on the photoluminescence (PL) of indirect band alloys. The energy difference between and X valleys of the chosen alloys in this study () is less than 40 meV. Temperature-dependent PL provides the basis for a detailed analysis of the doping effects on the electronic structure of the Te-doped samples. This work is meant to explore the indirect alloys for direct band light emission. This work also can be extended for other III–V alloys [, , , etc.] with indirect bandgap or near the direct-indirect band crossover.
2. EXPERIMENTAL
The epitaxy of films on Si (001) substrates off-cut 6° toward the [110] direction was carried out using an Aixtron Crius metal-organic chemical vapor deposition (MOCVD) reactor, where trimethylgallium (TMGa), trimethylindium (TMIn), arsine (), phosphine (), and diethyltellurium (DETe) were used as precursors. A compositionally graded buffer with a grading rate of was used to reach lattice-matching conditions for the films. The -graded buffer was terminated with a 1 μm thick capping layer, which was closely lattice matched (mismatch ) to the films. Then a chemical mechanical polishing (CMP) process was applied to remove the top 500 nm capping layer and planarize the surface. A 100 nm inter-layer was grown on the CMPed capping layer to initiate the III–V growth, which was also lattice matched to the films. Details of the epitaxy of inter-layer on graded buffer on Si (001) can be found in Ref. [12]. A 400 nm film was grown on at 650°C with a V/III ratio of 200. The effective flow rates of TMGa, TMIn, and were 1.31, 0.69, and 400 sccm, respectively, which were the same for all growth runs of films targeting a Ga content of 0.74. The tellurium was in situ doped during the epitaxy of the films, and the effective flow rate of DETe was adjusted for each sample to introduce different Te-doping concentrations. Capacitance-voltage measurements showed that the active doping concentration, n, increased from to , depending on the DETe flow rate. All samples were pieces of the size of .
X-ray diffraction (XRD) was used to measure the strain status and the compositions of the films. Selective defect etching was used to estimate the etch pit density (EPD) of the samples. The selective etching was carried out by pouring concentrated liquid (85 wt. %) into a glass crucible on a hotplate. The hotplate was then heated up and stabilized at the desired temperature of 250°C for 2 min [13]. The samples were lowered into the hot for 10 s. The EPD was counted using a scanning electron microscope (SEM). Temperature-dependent PL measurements were conducted from 6 to 300 K using a 473 nm solid-state laser for excitation. The laser spot size was , and the excitation power was 50 mW. The PL emission was detected using a Spex 750M spectrometer and a Hamamatsu thermoelectrically cooled GaAs photomultiplier tube (PMT).
3. RESULTS AND DISCUSSION
Our target Ga content was 0.74, and XRD measurements confirmed that there is a small tensile strain (0.1%–0.2%) in the undoped films, and the Ga content was measured to be . However, a composition change was observed for the n-type doped sample with , where the Ga content was reduced to 0.720. The XRD results were consistent with secondary ion mass spectrometry measurements. The composition change was caused by the tellurium dopants acting as surfactant, which changed the gallium and indium incorporation into the crystal lattice [14]. In addition, XRD showed that the Ga content was 0.734 for the sample with a doping level of . This suggests that the composition shift was not reproducible, which cannot be simply resolved by increasing the TMGa flow to compensate for the Ga content reduction. In addition, due to the composition shift, strains varied from to 0.25% for Te-doped samples with no clear correlation with Te-doping concentrations.
In Fig.
Fig. 1. (a) [110]-pole TEM pattern of fully disordered lightly Te-doped sample with . (b) SEM image of the etch pits of film, and .
Temperature-dependent PL is well suited to evaluate materials with direct and indirect bandgap properties. For indirect bandgap semiconductors, the dominant low temperature () PL peak with the highest energy can be attributed to excitonic recombination, while the lower energy peaks are associated with phonon replica and donor-acceptor pair (DAP) recombinations. The excitonic recombination can overlap with band-to-band transitions at higher temperatures, where the band-to-band emission peaks show a redshift with increasing temperature [18]. Direct bandgap semiconductors, however, generally show dominant direct bandgap recombination at low temperatures, as the exciton binding energy is usually quite small [19]. At higher temperatures, the direct bandgap luminescence redshifts and the PL intensity decreases. In the case of an indirect bandgap semiconductor with small difference between direct and indirect bandgap transition, at low temperatures the PL spectrum look like that from an indirect bandgap material while at higher temperatures direct bandgap characteristics dominate [20]. Because the direct bandgap emission is fueled by thermalized electrons, the direct bandgap PL increases with temperature, opposite to the direct bandgap semiconductor PL intensity [21].
According to these rules, the origins of PL emission peaks in indirect samples can be identified. Three Te-doped samples labeled as S01 (), S02 (), and S03 () were selected for detailed analysis of the doping effects on their optical properties. Figure
Fig. 2. Temperature-dependent normalized PL spectra (6–300 K) of Te-doped samples with (a) , (b) , and (c) . The positions of , , and DAP emission peaks from Ref. [3] are labeled, and their positions are indicated by dashed black lines. Solid red lines and arrows indicate the positions of direct band emission, shifting with increasing temperature.
Figure
Fig. 3. Peak positions of Te-doped samples versus temperature. The data points marked with green stars are from Ref. [3]. Dashed lines are fitted to derive the thermal coefficients of the band.
As previously mentioned for S03, the peak broadening and blueshift from 40 K were due to the direct band emission. In Fig.
Direct band materials, such as the () alloy, show a thermal quenching process, and its PL intensity decreases as temperature increases [23]. For the indirect alloy, S01, the thermal quenching effect is not obvious, as shown in Fig.
Fig. 4. Temperature-dependent PL spectra (175–300 K) of Te-doped samples with (a) S01, , (b) S02, , and (c) S03, .
Fig. 5. Arrhenius plot of integrated PL intensity versus temperature for Te-doped samples with (red dots) and (blue diamonds). The activation energies, , were derived from the fitted dashed lines.
Figure
Fig. 6. Integrated PL intensity (black squares) of Te-doped samples and their corresponding peak emission energy (blue circles) versus doping concentration at room temperature. Curve fitting (blue solid line) shows an approximate linear regression of emission energy with increasing doping concentration due to BGN effect.
Figure
However, the PL intensity of Te-doped significantly degrades at . This degradation is likely caused by the formation of inactive complexes related to the Te-doping [28]. As tellurium is a large n-type dopant compared with other dopants (i.e., Si), it may distort the sublattices of III–V compounds and form non-radiative recombination centers at high Te-doping concentration. The cross-section TEM (XTEM) image in Fig.
Fig. 7. (a) XTEM image of Te-doped sample at . (b) Scanning XTEM image at high resolution.
4. CONCLUSIONS
In conclusion, Te-doped indirect bandgap films were deposited on GaAsP/GeSi/Si. The active doping concentration varied from to . The temperature-dependent PL spectra show that the indirect-to-direct band transition occurs between 40 and 100 K, and the direct band emission dominates the room-temperature PL spectra. Due to the BGN effect, the separation energy between and X bands shrinks as Te-doping concentration increases. Because the activation energy for carrier thermalization is decreased as doping concentration increases, the carrier thermalization is significant in n-type doped indirect bandgap samples with and . Therefore, the doping promotes the carrier injection into the valley, which enhances the direct band transition. We show that the integrated PL intensity has been increased by five times for the sample with compared with the lightly doped sample with . The origin of the PL intensity degradation at high doping concentration is not fully understood, but the TEM results can exclude large dopant clusters. In addition, there is a linear relationship between the BGN and the increasing doping concentration, which agrees with the model used for highly n-type doped Ge.
5 Acknowledgment
Acknowledgment. This research was supported by the National Research Foundation Singapore through the Singapore MIT Alliance for Research and Technology’s “Low Energy Electronic Systems (LEES) IRG” research. The authors are grateful for the support provided by the management and technical staff at NTU.
Article Outline
Cong Wang, Bing Wang, Riko I. Made, Soon-Fatt Yoon, Jurgen Michel. Direct bandgap photoluminescence from n-type indirect GaInP alloys[J]. Photonics Research, 2017, 5(3): 03000239.