Samples morphologies and diffraction contrast experiments

All seven specimens exhibited single crystalline structure with the well-defined het- eroepitaxial orientation relationship that was presented in Section 3.1.1. However, the

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samples exhibited differences in the surface morphology, the defect content, and the structural anisotropy as this is expressed by the difference of the (1120) HRXRD RC FWHM along the m- and c- axes. Figure 3.4 shows representative RC FWHM vs. azimuth angle curves for an anisotropic and an almost isotropic sample. The “M-shape” depend- ence corresponds to a maximum along the [1100] azimuth (m-axis) and a minimum along [0001] (c-axis). In Table 3.II the values of the (1120) RC FWHM and the structural anisotropy parameter A are shown. The anisotropy A is defined as the difference in the values along the two in-plane directions, i.e. A = [(m-FWHM) - (c-FWHM)].

Figure 3.4: (1120)HRXRD RC FWHM vs. azimuth angle for two samples, one anisotropic (sam- ple F depicted using squares and a black dashed line,) and one isotropic (sample G depicted us- ing circles and a blue solid line). The variation of the FWHM as a function of the azimuth angle exhibits a strong “M”-shape dependence that corresponds to a maximum along the [1100]azi- muth and a minimum along [0001].

For a correlation of the structural anisotropy to the growth conditions, we point out that layers A, B, and E were grown under nitrogen-rich conditions, since higher growth temperatures lead to gallium desorption despite the increase of the flux ratio, resulting to smaller epilayer thicknesses than samples C, D, F and G that were grown under metal-rich conditions. Furthermore, the structural anisotropy of epilayers A, B, and E was larger as seen in Table 3.II. The epilayers grown under metal-rich conditions exhibited smaller anisotropy except for sample F that was grown on a high temperature (high-T) NL.

Figure 3.5 shows representative 10×10 µm² AFM scans for two anisotropic (A and F) and two isotropic (C and G) samples. The surface grains of the anisotropic sam- ples were larger and more elongated along the c-axis, whereas isotropic samples exhib- ited isotropic surface morphology with smaller, rounded grains. Moreover, the isotropic samples exhibited pits, and sample G exhibited a columnar morphology due to the low growth temperature (low-T). The RMS roughness measured by AFM in all the studied samples ranged between 18-20 nm.

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Table 3.II: (1120) RC FWHM along the c- and m- axes and anisotropy A values.

Sample FWHM c-axis (^o) FWHM m-axis (^o) A

A (G1267) 0.51 1.1 0.59

B (G1117) 0.34 0.71 0.37

C (G1416) 0.20 0.22 0.02

D (G1421) 0.20 0.24 0.04

E (G1553) 0.35 0.57 0.22

F (G1561) 0.26 0.54 0.28

G (G1550) 0.23 0.24 0.01

This morphological behavior can be understood by considering that the nonpolar a-plane surface is inherently anisotropic along the two in-plane directions, as was stat- ed in Section 3.1.1. Density functional theory (DFT) calculations show large anisotropy in the diffusion barriers along and perpendicular to the c-axis.²² On the a-plane surface, the Ga adatoms exhibit smaller diffusion barriers along the c-axis, and thus higher diffu- sion lengths along the c-axis in comparison to the m-axis. On the other hand, a high III/N flux ratio ensures that there is an increased Ga coverage of the surface during growth that reduces the surface anisotropy by reducing the diffusion barrier anisotropy. The reduction of the growth temperature further reduces the anisotropy by decreasing the metal desorption and mobility. The columnar sample G is a limiting case in which the Ga adatom mobility has been reduced. Contrary to Ga-rich, N-rich surface morphologies are not kinetically stabilized on nonpolar surfaces, and the diffusion anisotropy is deter- mined by Ga adatom mobility.²² Hence these calculations are consistent with the ac- quired experimental results regarding surface morphologies. However, in order to un- derstand why the elongated grains are related to the structural anisotropy, TEM obser- vations were employed to correlate the microstructure to the growth conditions.

Figure 3.5: 10×10 µm² AFM scans of samples (a) A: highly anisotropic with elongated grains along the c-axis, (b) C: isotropic with small grains, (c) F: anisotropic with grain elongation due to a high-T NL, and (d) G: isotropic but columnar due to low growth-T.

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Figure 3.6 illustrates BF XTEM images of the same samples as in Figure 3.5. All four images were taken under dynamical two-beam diffraction conditions near the [1100] zone axis. The anisotropic samples (a) and (c), with the elongated grains, exhib- ited flat surfaces, while the isotropic samples appear rough at the nanoscale, which is consistent with the formation and coalescence of small grains. The columnar morpholo- gy of the low-T sample G is depicted in Figure 3.6(d). The values of the rms roughness are also noted for each sample.

Figure 3.6: BF XTEM images taken near the [1100]zone axis under two-beam conditions using g 1120. (a) Sample A (Rq = 12 nm), (b) sample C (Rq = 68 nm), (c) sample F (Rq = 12 nm), (d) sample G (Rq = 32 nm). The XRD RC anisotropic samples exhibit flat surfaces while the isotropic samples have rough surfaces. All samples exhibit a defect rich zone near the interface.

In order to elucidate the influence of defects on the structural anisotropy of the samples, two-beam TEM observations were employed under different diffraction con- trast conditions, in both XTEM and plan-view geometries, using the g.b visibility criteri- on as was explained in detail in Chapter 2.

By using this methodology, it was found that the density of dislocations with c- type Burgers vector components was almost equal to that of the PDs. Figure 3.7 shows representative plan-view TEM images of one anisotropic [Figs. 3.7(a) and 3.7(c)] and one isotropic sample [Figs. 3.7(b) and 3.7(d)] under two-beam conditions using g 1100 and g 0002. The anisotropic samples clearly exhibited higher BSF and PD densities. On the other hand, in the isotropic samples the BSFs were fewer and appeared to terminate more often to PSFs than to PDs. Lattice TDs were observed only in the isotropic case and were rather scarce. From these results we can conclude that the 1/6<2023>PDs are the primary line defects that should be correlated to the structural anisotropy.

The principal difference between nonpolar and polar material is the high density of PDs that are difficult to be distinguished from the lattice TDs in XTEM geometry ob-

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servations. For this reason, plan-view diffraction contrast observations are particularly useful, since the partials are associated with the termination of BSFs.

We also note that all samples exhibited a dislocation rich zone close to the het- eroepitaxial interface, which was more pronounced when an AlN NL was employed.

Such defects were gradually reduced through dislocation annihilation reactions and, by combining plan-view and XTEM observations it was possible to estimate the influence of such reactions on the density of defects that reach the sample’s surface.

The density of TDs with a-type Burgers vector components was deduced by sub- tracting the density of PDs from the overall TD density. The overall TD density is a measure of dislocation density at approximately the middle of epilayer thickness, while the dislocations that reach the surface are in fact fewer. One reason for this is the dislo- cation annihilation reactions and another reason is that surface pits appear as holes in plan-view specimens and so the TDs under the pits cannot be counted.

The results from the defect measurements are summarized in Table 3.III. The isotropic samples exhibited lower dislocation densities whereas the BSF densities did not deviate a lot except in the case of sample A. It is also observed that in the anisotropic samples the PDs dominate or are comparable in density to the lattice TDs, whereas, in isotropic samples, the lattice TDs dominate in the middle of the sample thickness. Sam- ple F, grown under metal-rich conditions but on a high-T NL, also exhibited elevated dis- location densities similar to sample B. Also, the columnar sample G exhibited very few PDs but more lattice TDs at the regions of column coalescence. This dependence is clearly depicted in the graph of the anisotropy A vs. the TD density in Figure 3.8. Sample G has not been included in Figure 3.8 due to its columnar structure.

Table 3.III: Defect densities measured by TEM in XTEM and plan view geometry.

Sample

XTEM geometry Plan-view geometry Overall TD Density

(×10¹⁰ cm^-2)

PD density (×10¹⁰ cm^-2)

BSF Density (×10⁵ cm^-1)

A 27.5 18.4 36.6

B 14.4 6.5 13.4

C 4.6 0.7 6.7

D 6.0 2.5 13.2

E 24.7 15.6 16.0

F 14.3 4.5 13.0

G 10.8 0.6 11.0

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Figure 3.7: Plan-view BF TEM images for (a) anisotropic (sample A), and (b) isotropic (sample C) under two-beam conditions using g 1100 and (c)-(d) corresponding images of (a) and (b) using g 0002. Magenta circles (solid) denote PSFs, yellow (dotted) circles denote PDs, and the blue square denotes a mixed type lattice dislocation with a+c BV vector which is visible under both conditions. The anisotropic sample exhibits higher BSF and PD density, whereas the iso- tropic sample has fewer BSFs that terminate mainly to PSFs.

Figure 3.8: Graph of anisotropy A vs. TD density. Isotropic samples exhibit lower dislocation densities compared to the anisotropic ones.

The difference in the values of the structural anisotropy of samples D and F is remarkable and requires further investigation. From Table 3.III we see that their main difference is the density of lattice TDs, with sample F containing more than twice the TDs of sample D. It is reminded that both samples were grown under the exact same conditions except for the growth temperature of the AlN NL. From the HRTEM images of Figure 3.8, it can be seen that the NL of sample D exhibited a smoother AlN/GaN inter- face when viewed along [0001] (Rq: 1.4 nm) in comparison to sample F (Rq: 2.5 nm). In both samples, Moiré fringes in the vicinity of the AlN/sapphire interface are due to a

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zone of semipolar nanocrystals introduced by the nitridation pre-treatment as de- scribed in Chapter 4 and in Ref. [23]. The NL roughness affects the defect content of the GaN epilayers as it can be deduced from the WBDF XTEM images presented in Figures 3.10(a) and 3.10(b). Both samples exhibit a defect-rich zone near the GaN/AlN interface, but this zone is more pronounced in the case of the anisotropic sample F. This can be attributed to the formation of TD half-loops, such as marked by arrows in the multi- beam TEM image of Figure 3.10(c). The higher defect density of sample F can be at- tributed to the roughness of the NL that leads to nucleation of more lattice TDs. Such TDs emanate in an inclined manner, making it more probable to interact with each oth- er thus forming half-loops.

The NL roughnesses of samples D and F were similar when viewed along[1100] . This difference between the roughness values along [0001] and [1100] probably ex- plains the anisotropic surface morphology of sample F. In particular, the adatom diffu- sion is more easily facilitated along [0001], leading to grain elongation along this direc- tion.

Figure 3.9: HRTEM images along [0001] of the (a) low-T AlN NL (400^oC), and (b) high-T AlN NL (750^oC). The low-T sample has a relatively flat surface when compared to the high-T one. A zone of misoriented nanocrystals near the sapphire substrate is visible through Moiré fringes in both images.

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Figure 3.10: WBDF XTEM images taken off the [1100]zone axis under a g/3g diffraction condi- tion using g 0002 of (a) isotropic sample D, and (b) anisotropic sample F. (c) Higher magnifica- tion multi-beam image of sample F showing the defect-rich zone above the AlN/GaN interface.

The arrows indicate TD half-loops promoted by the inclination of the TDs due to the roughness of the NL.