Considerable attention has been focused on InGaN green LDs during the past few years due to the demand for green laser diodes in pico-projectors and laser displays. Since the first breakthrough of InGaN-based green LDs was achieved by Osram Corp. in 2009 [1], green InGaN LDs with emission wavelength above 500 nm grown on c-plane [1–4], (11-22) plane [5], and (20-21) plane [6–9] have been realized. However, the threshold current density of InGaN-based green LDs is still high. The reasons for high threshold current density for green LDs mainly include increased defects and polarization field in InGaN QWs as nearly two-fold higher indium (In) composition is required in green LDs compared to blue LDs. In this article, by observing the morphology evolution of green InGaN/GaN QW and studying the CL property, we investigate indium-segregation-related defects that are formed at green InGaN/GaN QW interfaces. Meanwhile, we also propose the approach and suggest the mechanism to remove them for green InGaN/GaN QW grown on both GaN templates and free-standing GaN substrates. By engineering the interface of green InGaN/GaN QWs, we have achieved green LD with low threshold current density of 1.85 kA cm⁻². The output power of the green LD is 58 mW at a current density of 6 kA cm⁻² under continuous-wave operation at room temperature.

The epitaxial growth of samples in this experiment was carried out on a commercial Aixtron 6 × 2 in. Close Coupled Showerhead (CCS) metal-organic chemical vapor deposition (MOCVD) reactor. The green LD structures consist of a 2 μm n-GaN layer, an n-AlGaN cladding layer, an InGaN lower waveguide (WG) layer, 2 pairs of InGaN/GaN QWs as the active region, an upper WG layer, a p-AlGaN electron blocking layer (EBL), a p-AlGaN/GaN (superlattices) SLs cladding layer, and a heavily doped p + -GaN contact layer.

Figure 1 shows the morphology evolution of single green InGaN/GaN QW grown on GaN/sapphire templates. The GaN/sapphire template has feature of atomic terraces and steps since it was grown by a step-flow growth mode at 1000 °C, as shown in Fig. 1(a). In order to incorporate enough In composition into InGaN QW to obtain green emission, the growth temperature was lowered to around 700 °C to grow InGaN QW. As shown in Fig. 1(b), green InGaN QW features two dimensional nano-size islands, which are formed along the atomic terraces underneath. The islands are only one to two monolayers high, as shown by an atomic force microscope (AFM) line profile in Fig. 1(c). This morphology is typical for green InGaN QW samples grown on GaN/sapphire templates under widely varied conditions in our experiments. Similar morphology of green InGaN QW samples has been reported [10,11]. We suggest this morphology is caused by the large strain existing in InGaN QW and related In surface segregation occurring during InGaN QW growth. Size difference between In and Ga as well as the strength difference of In-N and Ga-N bonds result into In segregation [12,13]. Figures 1(d) and 1(e) show the AFM images of samples with various thickness of GaN layers which were grown following the InGaN QW at the same growth condition with TMIn turned off (named as GaN cap layer afterward). The morphology of GaN cap layer with 1 nm thickness is rough, and atomic steps almost disappear. Moreover, trench defects appear in the samples with GaN cap layer thickness of 3 nm and larger, as shown in Fig. 1(e). A correlation between trench defects and In-rich clusters has been reported [14,15]. It is suggested that both the island morphology of InGaN QW and the trench defects result from severe In segregation. It should be pointed out that this morphology is not due to low growth temperature used for the GaN cap layer. As a comparison, the morphology of GaN cap layer grown directly on GaN/sapphire template without InGaN QW underneath at the same growth condition is shown in Fig. 1(f). It shows step-flow morphology. We suggest that the 2D island morphology and In clusters on the top of green InGaN QW hinder the diffusion of adatoms during the GaN cap layer growth, which results into rough surface and trench defects. The excess In clusters may also cause thermal degradation of InGaN/GaN QWs as reported in our previous study [16]. These defects greatly lower the emission efficiency of green InGaN/GaN QWs, which is especially a challenge for green LD growth since a much larger thermal budget imposed to InGaN QWs during p-AlGaN cladding layer growth.

 

Fig. 1 AFM images of (a) GaN/Sapphire template, (b) green InGaN QW. (c) is the AFM line profile of green InGaN QW along the white line shown in (b). AFM images of (d) 1 nm GaN cap layer with green InGaN QW underneath, (e) 4 nm GaN cap layer with green InGaN QW underneath, (f) 1 nm GaN cap layer directly grown on GaN/Sapphire template.

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Since severe In segregation and resulting trench defects significantly decrease the quantum efficiency of InGaN/GaN QWs, an interface treatment approach that can remove the excess In clusters existing at the InGaN QW top surface is necessary. We used thermal annealing to remove the excess In clusters. The InGaN QW annealed by ramping the temperature from 700 °C to 850 °C in 240 s and holding for 30 s. However, this annealing caused the decomposion of In-N bonds and thus reduce In composition of InGaN QW due to relatively weak In-N bonds, which results into large shortening of InGaN QW emission wavelength, as shown in Fig. 2(a). In order to prevent reduction of In composition in InGaN QW, an thin GaN cap layer was grown following InGaN QW at the same temperature, and its thickness should be large enough to prevent InGaN QW decomposition while be kept below which trench defects appear. GaN cap layer with thickness of 1.2 nm and 1.8 nm were investigated. After 1.2 nm or 1.8 nm GaN cap layer was grown, the InGaN QW with GaN cap was annealed by ramping the temperature from 700 °C to 850 °C in 240 s and holding for 30 s, and then GaN QB layer with thickness of 15 nm was grown at the same temperature of 850 °C. Figure 2(b) shows the AFM images of 15 nm QB layers grown on GaN/sapphire template. Atomic step and terrace can be clearly seen, indicating a step flow growth. Except for the v-pits, which is related to threading dislocations, no newly formed defects such as trench defects can be seen in the QB layer, which assure that InGaN QW active region with high quality be obtained. Therefore, an interface engineering and subsequent QB growth at raised temperature is effective to suppress defects and essential to obtain high quality InGaN active region. As shown in Fig. 2(c), the electroluminescence (EL) full-width-at-half-maximams (FWHMs) of LD structures with different GaN cap layer thickness are plotted as a function of EL wavelengths. Each data point represents an LD structure grown on GaN/sapphire template, while the dashed lines plotted along the lower limit of the EL FWHMs are guide lines for the eyes. It shows a tendency that the EL FWHMs of both groups of samples increased with emission wavelength, which is an indication of enhanced potential inhomogeneity as In composition in the InGaN QWs increases. However, it is noted that EL FWHMs of LDs with 1.2 nm cap layer are larger than that of LDs with 1.8 nm cap layer, and they show a more pronounced increase with increasing wavelength. Moreover, the wavelength of InGaN QW with 1.8 nm GaN cap layer grown at the same temperature is longer than that of InGaN QW with 1.2 nm GaN cap layer, as shown in Fig. 2(d). As we mentioned previously, the GaN cap layers act to protect InGaN QW layers from decomposition during temperature ramping up. Therefore, we believe that the GaN cap layers with nominal thickness of 1.2 nm may be not thick enough to protect InGaN QW layers, which results into additional fluctuation of InGaN QW layer thickness and In composition caused by In desorption. It is expected that this effect increases as the In composition of InGaN QW layers increases. As a result, the EL FWHMs of LDs with 1.2 nm GaN cap layer show a pronounced increase with wavelength. On the other hand, the guide line for the lower limit of EL FWHMs of LD structures with 1.8 nm GaN cap layer has a smaller slope, which suggests an GaN cap layer with nominal thickness of 1.8 nm can protect the InGaN QW layers from temperature ramping.

 

Fig. 2 (a) PL spectra of green InGaN/GaN QWs with and without GaN cap layer, (b) AFM images of 15 nm QB layers grown on GaN/sapphire template, (c) EL FWHMs dependent on emission energy of LD groups with 1.2 nm and 1.8 nm GaN cap layer, (d) PL spectra of green InGaN/GaN QW with 1.2 nm and 1.8 nm GaN cap layer.

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We then study the mechanism how excess In clusters can be removed while InGaN QW is covered by a GaN cap layer. Figures 3(a) and 3(b) show the AFM and the secondary electron microscope (SEM) images of 1.8 nm GaN cap layer with InGaN QW underneath upon thermal annealing. It should be pointed out the AFM and the SEM images are not the same region since they are separate observations. It is noted that both images show voids at the surface. We speculate that these void regions correspond to where In clusters exist prior to thermal annealing. In clusters were removed upon thermal annealing and voids were left at the surface. In order to clarify this assumption, we did CL measurement on the same region as the SEM observation. The CL measurement system is integrated in the SEM equipment, and can be done at the same time as the SEM observation. Monochromatic CL images at wavelength of 520 nm and 365 nm are shown in Figs. 3(c) and 3(d). The 520 nm emission comes from InGaN QW, while the 365 nm emission comes from the GaN/sapphire template since CL excitation depth is about 1 μm, which is deep into the GaN/sapphire template. The dark spots in the monochromatic CL image indicate no emission at specific wavelength, which could result from non-radiative recombination or inhomogeneous composition distribution of the emission medium. The dark spots in the 365 nm CL image as shown in Fig. 3(d) should correspond to dislocations, which are well-known non-radiative recombination centers. Since these dislocations extend to the InGaN QW and the surface, dark spots should also appear at the same region in the 520 nm CL image, which is the case as shown Fig. 3(c). In fact, the density of dark spots in the 365 nm CL image agrees with the density of dislocations which is about 3.7 × 10⁸ cm⁻² in this experiment. Dark spot density in the 520 nm CL image is 5.2 × 10⁸ cm⁻², which is higher than that in the 365 nm CL image. As shown in Fig. 3(c), the excess dark spots are marked by red circles, which also correspond to void regions in Fig. 3(b). These voids are not related to dislocations and are thus bright in the 365 nm CL image. In segregation to dislocations and thus formation of V-shape pits has been reported [17–21]. In clusters at dislocation regions are readily to be removed upon thermal annealing since they are exposed at the surface due to the existence of v-pits. In clusters at regions other than dislocations and their removal have not yet been reported. Our observations give evidence that In clusters exist not only at dislocation regions but also at dislocation-free regions, and they can be removed by thermal annealing too. We suggest the removal mechanism of In clusters formed at dislocation-free regions as following. The In clusters are metallic In and do not incorporate into III-nitride crystal, which hinders adatoms to incorporate on top of them during subsequent GaN cap layer growth. In fact, high density of pits can be observed in Fig. 1(d). Even that In clusters could be covered by a thin GaN cap layer, these regions are expected to have poor crystalline quality and thus to readily decompose upon thermal annealing. Therefore, in both cases In clusters not related to dislocations can be removed upon thermal annealing.

 

Fig. 3 AFM, SEM and CL images of annealed GaN cap layer with InGaN quantum well underneath grown on GaN/sapphire template.

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In order to study the In segregation and removal process irrelevant to dislocations, we also prepared sample of 1.8 nm GaN cap layer with InGaN QW underneath on free-standing GaN substrate upon thermal annealing. The dislocation density of the GaN substrate used in this experiment is less than 5 × 10⁶ cm⁻² which means it is dislocation free over a 5 μm by 5 μm observation area in average. As shown in Fig. 4(a) and 4(b), there exists high density of voids both in the AFM and the SEM images, which are certainly irrelevant to dislocations given the density. For the SEM and the CL images taken at the same region, these voids correspond to dark spots one by one in monochromatic CL image at 520 nm, as shown in Fig. 4(c). Therefore, these observations confirm our assumption that In clusters exist not only at dislocation regions but also at dislocation-free regions, and they can be removed by thermal annealing.

 

Fig. 4 AFM, SEM and CL images of annealed GaN cap layer with InGaN quantum well underneath grown on free-standing GaN substrate.

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We then grew green LD structure on free-standing GaN substrate with interface engineered InGaN/GaN QW active region, and the green LD structure was fabricated into ridge waveguide LD chips by conventional lithography and lift-off technique with ridge width of 10 μm and cavity length of 800 μm. The cavities and mirror facets were formed by cleaving, and the rear and front facets were coated with dielectric films. The measurements of the laser diodes were done under continuous-wave operation at room temperature. Figure 5(a) shows the emission spectra at various current density below and above lasing threshold. The spectrum FWHM at 0.025 kA cm⁻² is 123 meV, indicating sharp interface and small In composition fluctuation. The FWHM decreases quickly as injected current density increases, which leads to low threshold current density of 1.85 kA cm⁻². The lasing wavelength is 508 nm. Figure 5(b) shows the power-current-voltage curve. The voltage at threshold current is 4.3 V, which is quite low and is attributed to the suppressed carbon impurity in AlGaN: Mg cladding layer as reported in our previous study [22]. The output power of the green LD is 58 mW at a current density of 6 kA cm⁻².

 

Fig. 5 (a) Emission spectra at different current density, (b) Power-current-voltage curve of a typical green laser diode under continuous-wave operation at room temperature.

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In conclusion, by observing the morphology evolution of green InGaN/GaN QW and studying the CL property, we found In cluster defects exist not only at dislocation regions but also at dislocation-free regions, and thermal annealing at the interface can remove the In cluster defects. By engineering the interface of green InGaN/GaN QWs, we have achieved green LDs with low threshold current density of 1.85 kA cm⁻². The output power of the green LD is 58 mW at a current density of 6 kA cm⁻² under continuous-wave operation at room temperature.