Summary…

In summary, the effects of δ-TMIn-flow process with an initial fTMIn of 400 sccm during the well layer growth on the optical properties of InGaN/GaN MQWs were investigated. The HRXRD θ-2θ spectra and HRTEM images indicate the good layer periodicity and the structural quality of the InGaN/GaN MQW. And in both sample A and B, there exists the In-rich clusters in the InGaN/GaN MQW layers whether the δ-TMIn flow or not. But we can observe from PL spectra, the PL peak energy was different at 10 K even though the same composition extracted from XRD measurement. From the FWHM result of PL measurement, In-rich clusters were more uniform in size of sample as compared to sample A. And according to the PL and PLE measurement result, the larger values of σ, Ea and Stokes’ shift in sample B indicate that the δ-TMIn flow resulted in the increase the composition fluctuation in InGaN MQW region and shows the stronger carrier localization effect. And the light output of the GaN LEDs with the δ-TMIn-flow process is increased up to 24% without obvious deterioration of interfacial abruptness.

Figure 3.1 The structure of InGaN/GaN MQW.

Figure 3.2 Schematic diagrams of fTMIn variation over time in InGaN QWs for sample A

(a) (b)

(c) (d)

(a) (b)

(c) (d)

Figure 3.3 Cross-sectional TEM image of the InGaN/GaN MQWs of (a) sample A and (b) sample B. High magnification TEM image of the InGaN/GaN MQWs of (c) sample A and (d) sample B.

Figure 3.4 HRXRD spectra for (0004) reflection from the InGaN/GaN MQW structure of (a)

Figure 3.5 The normalized PL emission spectra of samples A (solid line) and B (solid circle) at 10K.

Figure 3.6 Temperature-dependent PL spectra of (a) sample A and (b) sample B.

Figure 3.7 The diagram of peak energy versus temperature of (a) sample A and (b) sample B.

Figure 3.8 Normalized PL intensity as a function of T^-1 for (a) sample A and (b) sample B.

Figure 3.10 PL and PLE spectra of the InGaN/GaN MQWs of (a) sample A and (b) sample B at 10K.

Fig 3.11 L-I characteristics for the LEDs of samples A and B

Chapter 4 Optical Characteristic of GaN-Quantum-dots grown on AlN Nanoholes 4.1 Introduction

The GaN-based materials generally contain many threading dislocations that reduce the quantum efficiency of LEDs and can be detrimental for the lifetime of LDs. To reduce this dislocation density is extremely hard due to large lattice mismatch between GaN and Sapphire (~13%). It had been theoretically predicted and verified that quantum dots (QDs) in the active layer lead to improved optical properties such as low and temperature independent threshold current [50].

The epitaxial growth by MBE and MOVPE was essentially a non-equilibrium process.

However, it was very useful to categorize it into three different modes as in the equilibrium theory. As schematically shown in Figure 4.1, Frank-Van der Merwe mode represented a layer-by-layer or 2D growth. Volmer-Weber mode corresponds to island or 3D growth.

Stranski-Krastanow (S-K) mode was 2D growth of a few monolayers, called as wetting layer, followed by 3D island formation. The last mode was the one most relevant to the growth of semiconductor QDs. Several approaches had been investigated for fabricating GaN QDs, the most common method was using the S-K growth mode, which the 2-D wetting layer transformed into 3-D islands due to 2.5% lattice mismatch between AlN and GaN [15, 51].

However, in the strain-driven self-assembled QD growth process, nonuniformity in the wetting layer gave rise to QD nucleation. The nucleation sites were weakly linked to surface steps resulted in non-uniform distribution; the QD nucleation sites were thus distributed randomly on the growth surface. The random nucleation resulted in a nonuniform QD size distribution and a broadened inhomogeneous linewidth.

There was another way to fabricate GaN QDs so-called anti-surfactant method. GaN QDs were fabricated on AlGaN surfaces, which were grown SiC substrates by MOVPE. The AlGaN surface was treated by tetraethyl-silicon (TESi) prior to the deposition of GaN. By adding a small amount of Si atoms on the AlGaN surface, the growth of GaN was changed from step-flow to 3D mode. And the most different between SK mode and anti-surfactant was the existence of wetting layer.

Moreover, recently some groups fabricated QDs structures by using self-assmebled nanoholes. Schuler et. al. [52] had found that in situ AsBr3 etching of a thin GaAs films in a molecular beam epitaxy (MBE) chamber can result in an array of small dips on a nanometer scale. These small dips were vertically aligned due to enhanced etching at locally strain areas.

The hole can be overgrowth with InAs such that an atomic flay surface was recovered.

Further InAs deposited on the filled-hole layer formed into pairs of self-assembled QDs.

In my thesis, the optical properties of GaN QDs structures with utilizing self-assembled nanoholes on AlN surface were investigated. We grew a flat AlN thin film and applied an in-situ etching process with H2 gas at high temperature for AlN surface to form nanoholes structures. Subsequently, these holes were filled with GaN.

4.2 Sample Structure

The epitaxial growth of GaN QDs on c-plane (0001) sapphire (Al2O3) substrates were performed by a EMCORE D-75 MOCVD system. The precursors of Ga, Al, and N were trimethylgallium (TMGa), trimethylaluminum (TMAl), and ammonia NH3. And Hydrogen (H2) and Nitorgen (N2) were used as carrier gas when growing the whole structure. A thermal cleaning process was carried out at 1080 ℃ for 10 min in a stream of hydrogen ambient before the growth of epitaxial layers. Afer depositing of a 30-nm thick GaN nucleation layer at 530 ℃, heated up to 1045 ℃ for growth of a 1-µm thick GaN layer. After completion of this GaN thickness of the 10-pairs of AlN/GaN structure was deposited. And the sample structure was illustrated in Figure 4.2.

4.3 Material Properties Analysis

Figure 4.3 was the surface image of AlN layers which grew in N2 ambient (Figure 4.3 (a)) and in H2 ambient (Figure 4.3(b)). From atomic force microscopy (AFM) analysis, the sample surface which grown in N2 ambient was flat (Ra=1.4 nm). In contrast with N2 ambient condition, the surface of sample grew in H2 ambient had many nanoholes on surface and the depths and width are 40 nm and 167 nm, respectively. We considered the discrepancy between Figure 4.3(a) and 4.3(b) as the thermal decomposition activation energies of GaN-based materials were different in H2 and N2 ambient, that H2 dissociation rate played a critical role in the decomposition rate [53]. As shown in Figure 4.3, the results indicated that H2 produced thermally unstable centers that were decomposed out of the growing layer, which also presented by K. W. Hipps et al. [54].

Figure 4.4(a) showed cross-section bright field TEM image obtained from the GaN/AlN QDs structure. And Figure 4.4(b) was the magnified bright field TEM images of GaN QDs on AlN layer. The periodic stacked layers of GaN QDs separated by AlN spacer were fabricated.

The GaN layer and AlN layer can be easily distinguished by fluctuation of different composition. The structure consisted of 10-stackes of GaN QDs with 100-nm thick AlN

4.4 Micro-Photoluminescence Measurement

We carried out the temperature dependent µ-PL measurements to investigate the emission properties and the internal quantum efficiency of the GaN QDs. Figure 4.5 shows µ-PL spectra of GaN QDs using the 325 nm excitation of He-Cd laser to selective excitation at temperature range from 80K to 300K. The µ-PL peak of this structure was observed around 3.464 eV at 300K. Figure 4.6 shows the fitting of the µ-PL spectrum with two Gaussian lineshape with GaN QDs and GaN bulk at temperature 80 K. Compare with GaN bulk structure, the GaN QDs ground state was blue shift 63 meV..

We extracted all GaN QDs spectrum from 80 to 300K and plotted in Figure 4.7. And the PL emission peak energy does not change much with this temperature range. The energy gap shrinkage was just about 35 meV in the QDs structures as shown in Figure 4.8 which is the PL peak energies versus temperature diagram. Generally, band gap energies of semiconductors decrease with increasing temperature following by Varshni empirical relation [41]

where T temperature in Kelvin, Eg(0) the band gap at 0 K, and α and β known as Varshni’s fitting parameters. And we also plotted bulk GaN according to Varshni empirical equation with dotted line in Figure 4.8. And from Figure 4.8, GaN QDs structure was less sensitive to temperature as compared with bulk GaN.

An Arrhenius plot of the integrated luminescence intensity of the QD emission was shown in Figure 4.9. The experimental data was fitted with well-known thermal activation relation [47]

where C the constant, Ea the activation energy and kB the Bolzmann’s constant. The thermal activation energy Ea can be obtained from the data. The thermal activation energy for the GaN island structure was 54 meV. Due to the small value of experimentally determined activation energy, in contrast with the relatively large thermal ionization barriers predicted for this structure. It can be postulated that the decrease of the GaN QDs PL intensity was due to the thermal activation of the non-radiative recombination center as often observed in many semiconductors, rather than the evaporation of the carriers from the GaN QDs to AlN [55]. In this process, deep level in the AlN barriers were thermally ionized, creating non-radiative

recombination centers of QD carriers.

Figure 4.10 shows the FWHM change with temperature from 80 K to 300 K. There was a narrowing of FWHM with increasing temperature to 100K. This unusual variation of FWHM can be explained as follows. When the thermal energy became comparable to the excition binding energy in the quantum dots structure, free carriers generated by exciton were able to jump to state of lower energy due to thermally assisted hopping. Under these conditions, the recombination of e-h pairs will be more frequent with predominant sizes. This was why the FWHM reduced as the temperature increase. At very high temperatures, of course the exciton was thermalized and emission took place from the entire range of excitation energies in the GaN QDs, so the FWHM increased [56-59].

4.5 Summary

In summary, we performed the structural and optical studies on inverted pyramid-shaped GaN QDs with dimensions of 40/40 nm (length/depth). The µ-PL measurements of these GaN QDs were performed over a temperature range from 80 to 300 K. Comparing with GaN bulk structure, the ground state of GaN QDs was blueshifted by 63 meV. PL emission peak energy does not change much with temperature, the energy gap shrinkage is just about 35 meV in the QD structures compared with 50-60 meV in GaN bulk materials. Finally we observed the narrowing of full-width at half maximum (FWHM) with increasing temperature to 100 K, this phenomenon can be attributed to carrier redistribution of different GaN QD sizes.

Figure 4.1 Schematic diagram of the three possible growth modes: Frank-van der Merwe, Volmer-Weber, and Stranski-Kranstanow.

Figure 4.2 Sample Structure of GaN QDs grown on AlN nanoholes.

Figure 4.3 AFM image of AlN grown under (a) Nitrogen ambient and (b) Hydrogen ambient

Figure 4.4 (a) Cross-section bright field TEM image obtained from the GaN/AlN QDs structure. (b) magnified bright field TEM images of GaN QDs on AlN layer.

Figure 4.5 Temperature dependent µ-PL spectra of GaN QDs.

Figure 4.6 µ-PL spectra at 80K. And the original µ-PL spectra fitted by two Gaussian peak with GaN QDs (353 nm) and GaN bulk (360.2 nm).

Figure 4.7 Gaussian fit of GaN QDs signal extracted from Figure 4.4 from 80 to 300K.

Figure 4.8 Peak position vs temperature. The dotted line is calculated according to Varshni law using parameter of GaN.

Figure 4.9 Temperature dependent of integrated PL intensity

Figure 4.10 FWHM change with temperature from 80 K to 300 K

Chapter 5 Conclusion 5.1 Conclusion

In this dissertation, we have studied the GaN-based quantum confined structure for two kinds of structures.

First, the effects of δ-TMIn-flow process with an initial fTMIn of 400 sccm during the well layer growth on the optical properties of InGaN/GaN MQWs were investigated. The HRXRD θ-2θ spectra and HRTEM images indicate the good layer periodicity and the structural quality of the InGaN/GaN MQW. And in both sample A and B, there exists the In-rich clusters in the InGaN/GaN MQW layers whether the δ-TMIn flow or not. But we can observe from PL spectra, the PL peak energy was different at 10 K even though the same composition extracted from XRD measurement. From the FWHM result of PL measurement, In-rich clusters were more uniform in size of sample as compared to sample A. And according to the PL and PLE measurement result, the larger values of σ, Ea and Stokes’ shift in sample B indicate that the δ-TMIn flow resulted in the increase the composition fluctuation in InGaN MQW region and shows the stronger carrier localization effect. And the light output of the GaN LEDs with the δ-TMIn-flow process is increased up to 24% without obvious deterioration of interfacial abruptness.

Second, we performed the structural and optical studies on inverted pyramid-shaped GaN QDs with dimensions of 40/40 nm (length/depth). The µ-PL measurements of these GaN QDs were performed over a temperature range from 80 to 300 K. Comparing with GaN bulk structure, the ground state of GaN QDs was blueshifted by 63 meV. PL emission peak energy does not change much with temperature, the energy gap shrinkage is just about 35 meV in the QD structures compared with 50-60 meV in GaN bulk materials. Finally we observed the narrowing of full-width at half maximum (FWHM) with increasing temperature to 100 K, this phenomenon can be attributed to carrier redistribution of different GaN QD sizes

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