For high mobilities in 2DEGs, the interface quality in heterostructures is of immense importance. The growth under indium bilayer stabilized conditions showed a spiral growth mode. However, another step had to be applied in order to achieve smooth InGaN surfaces.
All samples grown in the In bilayer stabilized regime showed a degraded surface after indium desorption. The observation of pits with a diameter of 2 nm on InGaN (0001) surfaces for indium rich growth conditions has been reported before [64]. The impact on the surface morphology observed in this work is however much stronger. An example is shown inFigure 6.7 a) exhibiting cracks and pits up to 3 nm deep. These defects act as scattering centers for electron transport, making the observed morphologies fatal for the mobility in 2DEGs.
Some observations suggested that the degradation of the surface is a post growth process. Some structures showed spiral hillocks with atomically flat terraces on an otherwise degraded surface (see Figure 6.7 b)). On one samples a missing link was found (see Figure 6.7 c)). This spiral hillock showed atomically smooth terraces at the base changing to pitted terraces further up. If the pits would form during growth, this should not be observed.
This suggests that the surface degrades after growth.
Possible processes are oxidation or rough crystallization of a metallic ad-layer during cooldown. This was studied by treating the samples with HCL, which removes metals, and with KOH, which etches oxides. No change in the morphology was found. As nitrides withstand both treatments, a reaction with nitrogen or surface segregation are a plausible explanation for the degradation.
z-scale: 6.0 nm z-scale: 14 nm z-scale: 9.0 nm
a) b) c)
Figure 6.7: Atomic force micrograph of the InGaN (0001) surface showing a) a degraded surface, b) a spiral on an otherwise degraded surface and c) missing link - a spiral which is smooth at the base but degraded at the top. The observations suggest a post growth degradation.
z-scale: 7.0 nm z-scale: 4.0 nm z-scale: 4.5 nm
a) b) c)
Figure 6.8: Atomic force micrographs of the InGaN (0001) surface a) after indium desorption under UHV conditions, b) no indium desorption and c) indium desorption under excess gallium, which resulted in the best morphology.
For b) and c) the metals were removed by HCL after growth.
This presumption was investigated by applying three different treatments after InGaN growth: 1) indium desorption without supplying nitrogen to the growth chamber (UHV:pbase= 5×10
-10
mbar) 2) no indium desorption with excess In during cool down, 3) excess Ga during In desorption and cool down. The results are shown inFigure 6.8. The morphology for UHV desorption showed only a minor improvement compared to regular desorption.
The other two steps, however, resulted in pit free surfaces. The degradation therefore occurs if indium desorbs directly from the surface into an ambient nitrogen atmosphere leaving a bare InGaN surface. The best morphology was obtained for excess Ga.
By applying excess Ga before InGaN growth, the number of threading dislo-cations was significantly reduced. This is shown in Figure 6.9 where two samples with and without excess Ga before InGaN growth are compared.
The black spots originate from edge dislocations penetrating the surface.
z-scale: 3.0 nm z-scale: 4.0 nm
a) b)
Figure 6.9: Atomic force micrographs of the InGaN (0001) surface with excess Ga a) only after InGaN growth and b) before and after InGaN growth. Excess In before InGaN growth was applied in a). The number of dislocations (black spots in a)) was reduced by excess Ga before InGaN growth.
z-scale: 5.5 nm z-scale: 6.0 nm
a) b)
Figure 6.10: Atomic force micrographs of the InGaN (0001) surface of GaN/InGaN/-GaN heterostructures with a) complete desorption before GaN/InGaN/-GaN cap layer growth and b) only indium desorption with excess Ga remaining on the surface before GaN cap layer growth. A lower defect density is achieved for no metal desorption before cap layer growth.
This observation cannot be explained by surface segregation. In this process, incorporated In atoms leave an InGaN layer at high temperatures due to the weak In-N bond. However, the dislocation reduction was observed when excess Ga was applied before InGaN growth, when no InGaN has grown yet. Surface segregation during the beginning of InGaN growth can also be excluded. Because the excess Ga reduced the dislocation generation and the surface degradation, it is reasonable to assume that both processes have the same origin. The sample shown inFigure 6.9 a) was, however, with excess In before InGaN growth. Excess In also reduced the surface degradation after growth, and should thus also suppress the dislocation generation if surface segregation has a negative effect on InGaN growth. Both processes were only observed if thin In adlayers were in direct contact with the ambient nitrogen, and it is therefore assumed that their origin is the reaction of indium with nitrogen.
The impact of excess Ga on the growth of GaN caplayers was also studied.
Figure 6.10 a)shows the surface of a GaN/InGaN/GaN heterostructure. For this sample, all metals desorped before GaN caplayer growth.Figure 6.10
Figure 6.11
RHEED-[110] after indium desorption from the de-graded InGaN (0001) surface.
A
√ 3x
√
3R30°is clearly ob-served.
b)shows the surface of a similar structure, where only indium desorped after InGaN growth, leaving excess Ga on the surface. Atomically flat terraces are observed for both samples. However, the desorption of all metals results in many defects. The optimum growth of InGaN heterostructures was therefore carried out under indium bilayer stabilized conditions with excess Ga during growth interruptions before and after InGaN growth.
During the optimization process presented in the previous paragraphs it was found, that a
√ 3x
√
3R30° RHEED reconstruction coincided with the surface degradation. This reconstruction, which is shown in Figure 6.11, always appeared after indium desorption if no excess Ga was applied. No reconstruction was observed after metal desorption, if first indium desorpted and then gallium. This compares well with the observation of the reconstruc-tion reported in the literature [64]. The reconstrucreconstruc-tion forms under nitrogen atmosphere on GaN (0001) or InGaN (0001) surfaces covered with 1/3 ML of indium. It is not observed after indium desorption if In is deposited onto a gallium adlayer. This supports the presumption that the degradation arises from a reaction of a thin indium layer with the ambient nitrogen. The authors also analyzed the growth depending on Ga and In supply. Above a certain indium flux they obtained smooth surfaces. During growth with this flux, they observed a
√ 3x
√
3R30° reconstruction by RHEED. The reconstruction was therefore linked to the growth of smooth layers. In this work the re-construction rather indicated the degradation of a surface that was smooth beforehand.
The requirement of excess Ga during the entire growth process complicates
Ga In Ga
surface coverage
GaN InGaN
growth
Ga In
supply N
59 60 61 62 63 64 65 66
RHEED intensity (arb. units)
growth time (min)
Figure 6.12: RHEED intensity during InGaN growth with excess Ga before and after InGaN growth. The supply, the growing material and the surface coverage are schematically indicated. The effective growth time is much shorter than the In+Ga+N supply. InGaN grows only after all excess Ga incorporated, which is observed by a minimum in the RHEED intensity.
the control of the InGaN thickness. InGaN grows only if there is no Ga on the surface, because Ga is preferably incorporated over In. The beginning of InGaN growth can however be monitored by RHEED. The RHEED intensity is reduced by the presence of liquid metallic layers. The intensity drop is material dependent since the scattering cross section of Ga is higher compared to In. The RHEED intensity is therefore different for surfaces cover with Ga or In or both.
A typical RHEED intensity signal during growth is shown in Figure 6.12.
At the beginning, the surface is covered with gallium. If the In,Ga and N shutters are opened, the intensity drops. During this time In accumulates and excess Ga incorporates because Ga/N < 1. At some point, no Ga is left on the surface and InGaN starts growing under a thick In wetting layer.
This is indicated by another decrease in the RHEED intensity.