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Nonpolar a-plane MgZnO films were grown on different a-GaN/r-sapphire templates by pulsed laser deposition (PLD), where the growth temperature of GaN buffer layers varied from 700 °C to 1000 °C. High-quality a-plane MgZnO epitaxial film was deposited on the optimized 1000 °C a-GaN/r-sapphire template. Temperature-dependent PL measurements of a-plane MgZnO films reveal an S-type temperature dependence of the excitonic recombination energy. It is resulted that the excitons are localized in alloy-induced potential fluctuations at low temperature and the room-temperature quantum efficiency is calculated to be 9.2%. An involvement of band-tail states in the radiative recombination was considered, and a quantitative description of the blue temperature-induced shift was obtained assuming a Gaussian shape of the band tail.
© 2014 Optical Society of America
1. Introduction
Zinc oxide (ZnO)-based materials have attracted increasing attention for their potential applications in ultraviolet-region optoelectronic devices, especially the UV LEDs and LDs [1–4]. Heterojunctions or quantum wells are the key structures and it is essential to realize the modulation of the bandgap with keeping the lattice constants resembling to each other for constructing various optical devices. A. Ohtomo et al. proposed II-VI oxide semiconductor alloy system, MgZnO, in 1998 to realize the ZnO based heterojunctions [5]. Theoretically the bandgap energy of the MgZnO films is tunable from 3.3 eV to 7.8 eV by artificially controlling the Mg contents in the alloys [5–7]. Furthermore, lattice constants of MgZnO are similar to those of ZnO because the ionic radius of Mg2+ 0.57 Å is basically the same as that of Zn2+ 0.60 Å [5]. Besides composing heterojunctions with ZnO, MgZnO itself is a strong candidate in photodetectors for its high absorption coefficient and abrupt absorption edge [8–12]. However, most of MgZnO films ever reported are c-plane films, which are affected by strong internal electric field introduced by the naturally high piezoelectric and spontaneous polarizations. This built-in electric field leads to the spatial separation of electron and hole wavefunctions in quantum wells, known as the quantum-confined Stack effect (QCSE), which results in the reduction of radiative recombination rate [13,14]. To avoid this adverse effect, it is a promising substitution to grow MgZnO alloys along the nonpolar a-direction [15,16]. In our previous work, we successfully obtained high-quality a-plane ZnO films grown on a-plane GaN [17]. This was derived from that ZnO (a = 3.2498 Å, c = 5.2066 Å) and GaN (a = 3.189 Å, c = 5.185 Å) had the same crystal structure (wurtzite), similar lattice constants (0.4% mismatch for c-axis and 1.9% mismatch for an axis) and a small difference between in-plane linear thermal expansion coefficients (αGaN = 5.59 × 10−6 K−1 and αZnO = 6.51 × 10−6 K−1).
As is known, optical property is one of the most important for optoelectronic materials and devices. However, though considerable efforts have been made to improve the crystallinity of a-plane MgZnO films, few researchers have focused on such properties of a-plane MgZnO films. In this letter, we report on the results of the photoluminescence (PL) characterizations of high-quality a-plane MgZnO layers grown on a-plane GaN. As a result, it is demonstrated that the GaN buffer layer can improve the structural and optical performance of a-plane MgZnO. In addition, a-plane MgZnO films are firstly observed to reveal an S-type temperature dependence of the excitonic recombination energy, suggesting the localization of excitons originating from compositional fluctuation in MgZnO alloys.
2. Experimental details
The a-plane GaN templates were grown directly on r-plane sapphire substrate in a MOCVD reactor [18]. 3-μm-thick a-plane GaN films were deposited on nitridated r-plane sapphire at different temperatures (700 °C, 1000 °C). The 1000 °C a-plane GaN layer is very smooth with typical nonpolar wavy features, while the surface of 700 °C a-plane GaN was covered by small GaN islands, which was the Volmer-Weber three-dimensional growth. The as-grown a-GaN/r-sapphire templates were then transferred into a PLD system to deposit the a-plane MgZnO layer. Prior to deposition, a mixture of ZnO (99.999%), MgO (99.998%) powders had been sintered to form targets with Mg content of 5 at. %. For the deposition of a-plane MgZnO, a KrF (248 nm) excimer laser with a frequency of 5 Hz and 150 mJ pulse energy (2.1 J/cm2 energy density at the target) was used for ablation of this MgZnO ceramic target. Distance between the substrate and the MgZnO target was kept at 50 mm. The chamber was pumped down to 1 × 10−4 Pa by a turbo molecular pump. Then all the MgZnO films were grown at a temperature of 700 °C under the oxygen pressure of 0.1 Pa for 30 min. Under the same conditions a-plane MgZnO grown directly on r-plane sapphire was chosen as the reference. The schematic structures of the samples are shown in Fig. 1.The samples grown on r-sapphire, 700 °C a-GaN and 1000 °C a-GaN were denoted as a, b and c, respectively. The thicknesses of a-plane MgZnO films were determined to be 400 nm with step tests by Profile-system. The surface morphologies of the samples were examined by atomic force microscopy (AFM) in contact mode (Veeco NanoScope MultiMode) and scanning electron microscope (SEM, FEI Nova NanoSEM 450, Holland), crystal structures and in-plane growth orientations were investigated by X-ray diffraction (XRD, PANalytical X’pert PRO MRD, Holland) using Cu Kα1 (k = 1.54056 Å) radiation. Room temperature photoluminescence (PL) spectroscopy was performed with the JY-HR800UV system. The temperature-dependent PL was measured in a flow cryostat over a temperature range from 10K to 300K using the 325 nm line of a He–Cd laser as the excitation source. The PL signals were collected by a charge couple-device (CCD) assembled with a grating monochromator.
Fig. 1 Schematic design of the investigated samples. The a-plane MgZnO films were grown on (a) r-plane sapphire, (b) 700 °C a-GaN/r-sapphire templates, (c) 1000 °C a-GaN/r-sapphire templates using PLD under the same conditions.
3. Results and discussions
3.1 Structural properties
The high-resolution X-ray diffraction (HRXRD) was performed by using the symmetric 2theta/omega (2θ/ω) scanning mode with the X-ray incident plane parallel to the c-axis of the a-plane MgZnO film, and the results are shown in Fig. 2(a).The sharp peaks at 25.5° and 52.6° correspond to and diffraction peaks from r-plane sapphire substrate, respectively. The peaks from a-plane MgZnO and a-plane GaN can be observed at 56.1° and 57.6° for both b and c. No other diffraction peaks are observed for these two samples, indicating that a-plane MgZnO films were solely grown on the a-GaN/r-sapphire templates. Having confirmed the a-plane growth surface, off-axis diffraction peaks were used to determine the in-plane orientation of the MgZnO with respect to the GaN. Phi scanning of -GaN and -MgZnO for the sample c is recorded in Fig. 2(b). The phi scanning curve indicates the following epitaxial relationship: MgZnO||GaN, MgZnO||GaN. While for the sample a no a-plane GaN peak is detected, but the peak at 34.5° from c-plane MgZnO is able to be recognized. It means that the sample a is partly polycrystalline with both and orientations. Furthermore, the FWHM values of the a-plane MgZnO of 2θ/ω scanning are 0.31°, 0.59°, 0.28° for a, b and c, which qualitatively indicates that the crystalline quality of the sample c is the best of these samples.
Fig. 2 (a) HRXRD 2theta/omega scanning curves of the a-plane MgZnO films grown on r-plane sapphire, 700°C a-GaN/r-sapphire templates and 1000°C a-GaN/r-sapphire templates. (b) Phi scanning curve of -GaN and -MgZnO for the sample c.
To further examine the crystalline qualities of the obtained films, the a-plane MgZnO films were characterized by the X-ray rocking curve (XRC). Since materials grown along nonpolar direction typically exhibit strong anisotropic behaviors of crystallinities with respect to in-plane orientations [19], the XRCs were measured as a function of azimuthal angle φ, shown in Fig. 3. From this figure, it is shown that the XRC FWHM values (around 2200 arcsec) of a-MgZnO grown directly on the r-plane sapphire (sample a) are fundamentally the same as what was reported previously [20,21]. While for the the sample c, the FWHM values along the c- and m-directions are 1136 arcsec and 2196 arcsec, respectively. The improvement of crystalline quality may result from the small lattice mismatch between MgZnO and GaN, leading to the reduction of misfit dislocations and stacking faults. However, the ratio of the tilt mosaic along the m-axis to that along the c-axis for sample a is the least of all these samples. As mentioned above, sample a is partly polycrystalline and the -growth may decrease the anisotropy of crystalline features. Meanwhile, the surface of the nonpolar a-GaN is wavy with stripes along the c-direction [18]. Therefore, when grown on this stripe-like surface, the migration of Zn and Mg atoms along c-axis may be enhanced, which will further promote more growth along c-direction. Furthermore, FWHM values of the sample b which is grown on 700 °C-a-GaN are larger than that of the samples a and c. This illustrates that the crystallographic features of the a-plane MgZnO are sensitive to a-plane GaN buffer. The surface coalescence and improvement of crystalline quality of the template layer are essential to the following growth.
Figure 4(a) shows the SEM image of the sample c. Though some small pits present on the surface, the commonly observed streak features along the c-direction on the surface is an obvious symbol for the coalesced nonpolar a-plane MgZnO. The pits are believed to originate from the island growth and coalescing at the early stage of film growth. The surface morphology of wafer grown on 1000 °C-a-GaN was further characterized by AFM, as is shown in Fig. 4(b). The root mean square of surface roughness value over an area of 5 × 5 μm2 is 3.54 nm for the sample c.
The typical EDS spectrum of the sample c was recorded to determine the compositional analyses of Mg and Zn in the samples. Figure 5 shows the typical EDS spectrum for the sample c. Zn, Mg and O are the dominated elements in the deposited layers by PLD. Meanwhile, Ga and N elements from the a-plane GaN buffer layer are also in sight. The Mg and Zn contents are 27.18 ± 2.54 at.% and 4.33 ± 0.81 at.%, respectively, and the ratio of Mg to (Mg + Zn) is 13.74 at.%. As is expected, the Mg content (around 12.5 at.%) is systematically larger than in the targets by a factor of 2.5 [5] [22]. This difference is attributed to the fact that the vapor pressure of ZnO is much larger than that of MgO. Then ZnO can easily desorb from the growing surface and lead to the condensation of MgO on the surface.
3.2 Optical properties
The optical properties were investigated by the room-temperature (RT) PL measurement, which gave insights into the qualities of the a-plane MgZnO epilayers grown on different templates. The RT-PL spectra at room temperature for all samples are shown in Fig. 6.As the penetration depth of the laser light into the MgZnO was approximately 60 nm in the investigated Mg concentration range [23], only a strong emission near the band edge at 3.47 eV is observed for all the three samples. This value corresponds to what was reported previously [5–7]. Meanwhile, it is clear that PL intensity of a-MgZnO deposited on the 1000 °C a-plane GaN buffer is almost twice higher than the one directly grown on the r-plane sapphire and the FWHM of emission of the former is 0.13 eV, smaller than the latter (0.16 eV). Furthermore, the sample b has the lowest PL intensity and greatest FWHM (0.63 eV). This confirms the improvement of crystalline qualities induced by the properly grown a-plane GaN. As shown by the inset of Fig. 6, broadband emissions of all samples at around 2.25 eV mainly due to deep-level impurities can be hardly seen, whose intensities are extremely low under this condition.
Fig. 6 Room-temperature (300K) photoluminescence (PL) spectra recorded for the samples a, b and c. The broadband emissions were magnified and shown in the inset.
To further investigate the excitonic properties, the typical results of temperature-dependent PL measurements for the sample c were recorded in Fig. 7(a).It can be easily observed that nonpolar a-plane MgZnO grown on the well coalescent a-plane GaN reveal an S-type dependence of the excitonic recombination energy. As is shown, first from 12 K to 40 K, PL peaks show a red-shift yielding, higher than the normal energy shrinkage in accordance with Varshni Eq. (1) meV). When the temperature increases from 40 K to 100 K, PL peaks increase gradually, showing an anomalous blue-shift trend. Then, at temperatures > 100 K, PL peaks evolution obeys the normal Varshni formula. S-type temperature dependence has been extensively investigated in (Al,In)GaN systems [24]. It is generally accepted that this anomalous optical transitions is introduced by excitons localized in alloy-induced potential fluctuations [25]. With increasing temperature, they overcome local potential barriers, occupy energetically deeper potential minima, and recombine resulting in a redshift of the luminescence energy. A further increase in the exciton thermal energy however inhibits localization in the deepest potential minima and therefore leads to a blueshift. At higher temperatures, the PL emission follows the temperature dependence of the bandgap [26,27]. The S-type temperature dependence supports the presence of temperature sensitive phonon-assisted hopping of localized excitons, which may have a significant effect on the suppression of nonradiative processes and the performance of optical devices like the InGaN systems [28].
Fig. 7 (a)Temperature dependence of the band edge luminescence of the sample c. (b) Temperature-dependent near-band-edge emission shift of the sample c (black blocks) and the fitting curve of sample c plotted from the band-tail-filling model (red line).
Generally, it is well determined that the PL peak positions conventionally follow the temperature-dependent band-gap shrinking trend given by Varshni equation
where Eg(T) and Eg(0) are the band-gap transition energies at T and 0 K, respectively, while α and β are thermal coefficients adopted from [29] and [30]: α = (8.2 ± 0.3) × 10−4 eV K−1 and β = (700 ± 30) K as used for the A exciton transition of ZnO. However, to quantitatively describe the abnormal temperature-induced blueshift in various alloys system, P. G. Eliseev proposed that the continuous temperature-induced blueshift suggested an involvement of continuum of band-tail states and this shift followed the expressions σ2/kT in the case of the Gaussian profile of the density-of-states (DOS) where σ2 is the Gaussian dispersion, and k is the Boltzmann constant. Then taking the band-tail-filling model into consideration, the Varshni equation transforms into [31]:where σ may represent the degree of the localization of excitons. The fitting curve from the Eq. (2) is plotted in Fig. 7(b), where the fitting parameters α β and σ are 8.4 × 10−4 eV K−1, 700 K and 1.28 × 10−2 eV, respectively. It is resulted that the theoretical curvature agrees well with the experimental values above 50 °C.Figure 8 shows integrated PL intensity as a function of T in the form of an Arrhenius plot. The activation energy required for the nonradiative recombination process in the highest temperature region was estimated by least-squares fitting of the results to the equation
where ε and C are the activation energy and a constant, respectively. The fitting results are shown in the dashed lines, and the most probable value for ε is 15.0 meV. It is known that the activation energy correlated with the depth of the potential minima is an indication for the ability of the suppression of the nonradiative recombination process [32]. This result is close to the value for the nonpolar m-plane MgZnO films [26], while lower than films grown along c-direction [32] at similar Mg content. This may result from the intrinsic differences between the nonpolar and polar materials, such as the different defect types. It can be also found in Fig. 8 that the transition from radiative to non-radiative process occurs at about 70K. It is widely identified that the potential traps in MgZnO alloys are very efficient in inhibiting the non-radiative recombination. The quantum efficiency is calculated by the following formula: IPL(T)/IPL(0K)≈η(T) = τtotal(T)/τr(T) = 1-1/(1 + τn(T)/τr(T)). Then the RT quantum efficiency η(300K) = 9.2%. The result shows that nonradiative recombination is dominant at room temperature and the growth need to be further optimized and it is in progress.4. Conclusions
In summary, we have performed the PLD growth of nonpolar a-plane MgZnO films with high crystallinity on a-GaN/r-sapphire template and investigated the optical properties of such films. This sample has the sole epitaxial orientation, and the FWHM values for MgZnO scanned along c- and m-direction are 1136 arcsec and 2196 arcsec, respectively, indicating a small mosaicity and a low` dislocation density in the MgZnO film. Temperature-dependent PL measurements of a-plane MgZnO films reveal an S-type temperature dependence of the excitonic recombination energy, suggesting the localization of excitons takes place. The band-tail-filling model is involved in the fitting to quantitatively analyze the temperature-induced blueshift of the recombination energy. It is resulted from the Arrhenius plot that the activation energy of the a-plane MgZnO is 15.0 meV and the room-temperature efficiency is calculated to be 9.2%.
Acknowledgments
The authors want to acknowledge Dr. Yulian Li and Prof. Xingjun Wang for the support in the temperature-dependent PL measurements in Key Laboratory of Infrared Imaging Materials and Detectors of Shanghai Institute of Technical Physics of Chinese Academy of Sciences. This work was supported by the National Basic Research Program of China (Grant No. 2012CB619302), The Science and Technology Bureau of Wuhan City (No. 2014010101010003), Natural Science Foundation of Hubei Province (Grant Nos. 2011CDA81), Science Foundation from Hubei Provincial Department of Education (Grant Nos. D20131001), Key Laboratory of infrared imaging materials and detectors of Chinese Academy of Sciences (Grant No. IIMDKFJJ-13-04), and the National Natural Science Foundation of China (Grant No. 10990103, 51002058, 61274010).
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