**4. Results and discussion**

Let us then discuss the surface structure of the substrate after thermal cleaning under *Profile a* and before LT‐AlN BL growth under *Profile b*. (For this case, the substrate was immediately cooled down, removed from the reactor, and analyzed.) As confirmed by AFM measurements, the thermal annealing under *Profile a* produced a parallel step‐and‐terrace surface structure with monolayer steps (**Figure 5(a)**). The average terrace‐width and step height are ∼125 and ∼0.21 nm, respectively, as estimated from AFM line scan profile. Hence, the substrate's offcut angle was estimated to be ∼0.11°, which is comparable to the expected off‐cut angle of 0.15°. Moreover, because the ML step height corresponds to 1.299 nm *co*/6 = 0.22 nm, where *co* is the unit cell of sapphire having six ML steps of oxygen layers along the *c* axis [27], this confirms that the steps are correlated with the periodicity of oxygen and the interaction between oxygen atoms of successive layers is stronger than Al atoms of successive layers [26, 28].

Formation Mechanism and Elimination of Small‐Angle Grain Boundary in AlN Grown on (0001) Sapphire Substrate http://dx.doi.org/10.5772/66177 49

∼15‐nm‐thick LT‐AlN BL. Then the temperature was increased to *T*<sup>g</sup> = 1285°C for HT‐AlN growth. In the experiment, all temperature readings are from those indicated by the thermocouple placed near the substrate. Note that although both profiles incorporate LT‐AlN BL, the timing, hence, the substrate's surface structure at which the LT‐AlN BL is introduced is quite different, which is crucial forthe formation/elimination of small‐angle grain boundary. For analyses, atomic force microscopy (AFM) measurements were conducted to study the surface morphologies both of the substrate's surface and AlN epilayer, while XRD and transmission electron microscope (TEM) measurements were conducted to study the structural qualities and to assess the SAGB. CL measurements were conducted to study the

**Figure 4.** Temperature profiles for two‐step growth of AlN on sapphire substrate. *Profile a* incorporates thermal clean‐

Let us then discuss the surface structure of the substrate after thermal cleaning under *Profile a* and before LT‐AlN BL growth under *Profile b*. (For this case, the substrate was immediately cooled down, removed from the reactor, and analyzed.) As confirmed by AFM measurements, the thermal annealing under *Profile a* produced a parallel step‐and‐terrace surface structure with monolayer steps (**Figure 5(a)**). The average terrace‐width and step height are ∼125 and ∼0.21 nm, respectively, as estimated from AFM line scan profile. Hence, the substrate's offcut angle was estimated to be ∼0.11°, which is comparable to the expected off‐cut angle of 0.15°. Moreover, because the ML step height corresponds to 1.299 nm *co*/6 = 0.22 nm, where *co* is the unit cell of sapphire having six ML steps of oxygen layers along the *c* axis [27], this confirms that the steps are correlated with the periodicity of oxygen and the interaction between oxygen

atoms of successive layers is stronger than Al atoms of successive layers [26, 28].

optical properties of AlN.

48 Study of Grain Boundary Character

ing, while *Profile b* incorporates no thermal cleaning.

**4. Results and discussion**

**Figure 5.** AFM surface morphologies of sapphire substrate after (a) thermal cleaning under *Profile a*, and before LT‐AlN BL growth at (b) 1100°C and (c) 800°C under *Profile b*. The corresponding (d) HT‐AlN growth for thermally cleaned substrate and under *Profile b* using (e) 1100°C and (f) 800°C LT‐AlN BL *T*g.

On the other hand, for sapphire substrate under *Profile b*, the formation of surface steps are undefined at BL *T*g =   800°C (**Figure 5(c)**). However, two monolayer steps can be defined when the BL *T*g is increased to 1100°C (**Figure 5(b)**). This indicates that with increasing BL *T*g, the surface structure changes from "rough" to smooth having two ML steps. Therefore, different surface structures of sapphire substrate are formed depending on these two experimental conditions. It is therefore interesting to find out its effect on the structural quality of the subsequently grown AlN epilayer.

The corresponding surface morphologies for AlN grown under *Profile a* (**Figure 5(d)**) and *Profile b* at LT‐AlN BL *T*g =1100°C (**Figure 5(e)**) and 800°C (**Figure 5(f)**) are also shown in **Figure 5**. It can be seen that both profiles show an AlN with atomically smooth surfaces, as evidenced by their root‐mean‐square (RMS) roughness values. Moreover, a step‐and‐terrace surface morphology structure which replicates that of the substrate is exhibited for AlN grown under *Profile a*, while a meandering surface morphology is observed for AlN grown under *Profile b* using LT‐AlN BL *T*g =800°C (**Figure 5(f)**). Furthermore, these meandering steps begin to align along a certain direction as LT‐AlN BL *T*g is increased to 1100°C (**Figure 5(e)**). Clearly, these differences in AlN morphologies are most likely influenced by the surface structure of the substrate prior to AlN growth.

**Figure 6** shows the asymmetric (10‐12) *φ*‐scan of AlN epilayers grown under *Profile a* (**Figure 6(a)**) and *Profile b* using LT‐AlN BL *T*g =1100°C (**Figure 6(b)**). The (11‐23) *φ*‐scan of the substrate is also shown. It is well known that the in‐plane epitaxial relationship between AlN and sapphire is AlN<10‐10>||α‐Al2O3<11‐20> [8]. Moreover, a closer look at AlN (10‐12) diffraction shows two peaks under *Profile a* (**Figure 6(a)**) with separation *Δθ* ∼3.72°, while only a single diffraction peak is observed under *Profile b*. This indicates the presence of a special grain boundary under *Profile a*, where the two AlN grains have a particular in‐plane misor‐ ientation relationship, while it is successfully suppressed under *Profile b*. Furthermore, because sapphire (11‐23) diffraction does not exhibit two peaks, the special grain boundary is confirmed to exist only in AlN epilayer.

**Figure 6.** XRD *φ*‐scans of AlN grown sapphire substrate under (a) *Profile a* and (b) *Profile b* using LT‐AlN BL *T*g =  1100°C.

To further confirm the existence of the special grain boundary in AlN under *Profile a*, a plan‐ view bright‐field TEM micrograph and the corresponding selected area electron diffraction pattern (SAEDP) were taken under [0001] zone axis, as shown respectively in **Figures 7(a)** and **(b)**. A periodic bright and dark contrast of two AlN grains is observed. The AlN grain width is found identical to the step width of thermally cleaned sapphire substrate, implying that the origin of the grain boundary is related to the substrate's surface structure. As the AlN is grown onto sapphire substrate with either *A* or *B* oxygen stacking, the characteristic of that surface is also carried into AlN [26]. And as supported by XRD results and because only a slight misorientation relative to one another exists between these two AlN grains, this grain boundary is confirmed to be a small‐angle grain boundary [25, 26]. Furthermore, due to the arrays of edge dislocations that exist at the boundary, this type of special boundary is called pure low‐angle tilt SAGB. The spacing *D* between adjacent edge dislocation array can also be estimated using the formula = /sin /sin, where **b** is the in‐plane burger's vector (= 0.3112 nm) and Δ*θ* is the misorientation angle (Δ*θ* ∼ 3.72°) obtained from XRD measurement. Hence, the spacing between edge dislocations is estimated to be ∼4.75 nm. The SAEDP also supports the observation of SAGB, as seen from the double diffraction spots (denoted by arrow marks in **Figure 7(b)**). On the other hand, no special grain boundary is observed for AlN grown under *Profile b* (not shown). This suggests that the buffer layer technique is effective for suppressing the SAGB.

**Figure 6** shows the asymmetric (10‐12) *φ*‐scan of AlN epilayers grown under *Profile a* (**Figure 6(a)**) and *Profile b* using LT‐AlN BL *T*g =1100°C (**Figure 6(b)**). The (11‐23) *φ*‐scan of the substrate is also shown. It is well known that the in‐plane epitaxial relationship between AlN and sapphire is AlN<10‐10>||α‐Al2O3<11‐20> [8]. Moreover, a closer look at AlN (10‐12) diffraction shows two peaks under *Profile a* (**Figure 6(a)**) with separation *Δθ* ∼3.72°, while only a single diffraction peak is observed under *Profile b*. This indicates the presence of a special grain boundary under *Profile a*, where the two AlN grains have a particular in‐plane misor‐ ientation relationship, while it is successfully suppressed under *Profile b*. Furthermore, because sapphire (11‐23) diffraction does not exhibit two peaks, the special grain boundary is confirmed

**Figure 6.** XRD *φ*‐scans of AlN grown sapphire substrate under (a) *Profile a* and (b) *Profile b* using LT‐AlN BL *T*g = 

To further confirm the existence of the special grain boundary in AlN under *Profile a*, a plan‐ view bright‐field TEM micrograph and the corresponding selected area electron diffraction pattern (SAEDP) were taken under [0001] zone axis, as shown respectively in **Figures 7(a)** and **(b)**. A periodic bright and dark contrast of two AlN grains is observed. The AlN grain width is found identical to the step width of thermally cleaned sapphire substrate, implying that the origin of the grain boundary is related to the substrate's surface structure. As the AlN is grown onto sapphire substrate with either *A* or *B* oxygen stacking, the characteristic of that surface is also carried into AlN [26]. And as supported by XRD results and because only a slight misorientation relative to one another exists between these two AlN grains, this grain boundary is confirmed to be a small‐angle grain boundary [25, 26]. Furthermore, due to the arrays of edge dislocations that exist at the boundary, this type of special boundary is called pure low‐angle tilt SAGB. The spacing *D* between adjacent edge dislocation array can also be estimated using the formula = /sin /sin, where **b** is the in‐plane burger's vector

to exist only in AlN epilayer.

50 Study of Grain Boundary Character

1100°C.

**Figure 7.** Plan‐view TEM bright‐field image of AlN grown under (a) *Profile a*. (b) Corresponding selected‐area electron diffraction pattern of AlN in (a).

**Figure 8.** (a) LT‐AlN BL *T*g as a function of symmetric (0002) and asymmetric (10‐12) *ω*‐scans. (b) AFM surface mor‐ phology of AlN grown at optimum LT‐AlN BL *T*g (1050°C).

Then the LT‐AlN BL *T*<sup>g</sup> was optimized to improve the mosaicity of the epitaxial film using the current growth condition. The result of the XRD symmetric (0002) and asymmetric (10‐ 12) *ω*‐scans are shown in **Figure 8**. While the XRD linewidth of the tilt component corresponding to the symmetric (0002) becomes narrow with increasing temperature, the twist component corresponding to the asymmetric (10‐12) initially becomes narrow and then broadens with increasing temperature. Hence, the narrowest linewidth for each component does not coincide with each other, and the reason is still unknown at this time. But by balancing these components, it can be deduced that the optimum LT BL temperature is *T*<sup>g</sup> ∼1050°C, where the XRD linewidths are ∼66 and ∼1443 arcsec, respectively, for (0002) and (10‐12) *ω*‐scans. The symmetric component is comparable to that grown under *Profile a* (FWHM is ∼64 arcsec), suggesting that both have highly‐oriented films along the *c*‐axis growth direction. On the other hand, the asymmetric (10‐12) linewidth is wider than that grown under *profile a* (FWHM is 1145 arcsec). **Figure 8(b)** shows the AFM surface morphology of AlN grown under the optimum LT‐AlN BL *T*<sup>g</sup> of 1050°C. Although nanopits have been reduced compared with other BL *T*<sup>g</sup> (**Figures 5(e)**, **(f)**), it is believed that the quality of the film can be further improved especially the twist mosaicity upon optimizing the buffer layer thickness or other growth parameters.

To demonstrate the effect of eliminating the SAGB on the optical properties of AlN epilayer, we obtained the CL spectra for both profiles, as shown in **Figure 9**. The CL measurements were acquired at 93 K under 10 kV and 0.1 μA emission condition (spot size is ∼1 μm). CL peaks are assigned to free and bound excitonic emissions, including the LO phonon replicas, as shown in the inset figure [14]. Moreover, the emission intensity of AlN under *Profile b* using the optimum LT‐AlN BL *T*g of 1050°C is approximately more than two times higher than that under *Profile a*. This result is attributed to the higher probability of radiative recombination of electron and hole pairs due to the elimination small‐angle grain boundary. The differences in the CL peak position also suggest their different residual strain, where the tensile‐strained AlN under *Profile b* has smaller bandgap energy (∼6.011 eV) than that of an almost relaxed AlN (∼6.031 eV) under *Profile a*. The slow relaxation experienced by AlN under *Profile b* is most likely due to the reduced generation of dislocations upon the suppression of SAGB. As SAGB is a type of an edge dislocation, eliminating it is expected to enhance the optical properties of AlN.

**Figure 9.** CL spectra of AlN grown under (a) *Profile a* and (b) *Profile b* using optimum LT‐AlN BL *T*g = 1050°C.

Hence, the origin of SAGB can be ascribed to the surface structure of sapphire substrate. It was shown that as long as the surface has ML steps, even the introduction of LT‐AlN BL would be ineffective in eliminating the SAGB. Hence, there is a maximum allowable BL *T*g at which the substrate's surface does not transform into monolayer steps yet. In this study, we consider that the maximum BL *T*g would be around 1100°C. The result also implies that there is a proper timing to introduce the LT‐AlN BL [21]. This tendency is also observed in direct HT‐AlN growth using different growth methods [16, 21, 25], suggesting the wide observation of this phenomenon. It is noteworthy that even during temperature ramp‐up, because the optimal *T*<sup>g</sup> for HT‐AlN growth is in the vicinity around *T*g = 1200°C, it is likely that the surface could also change to periodic (ML) steps, as pointed out [25]. This is also supported in the present study based on the evolution of surface structure in both profiles*.* Therefore, the LT‐AlN BL would be necessary in order to prevent the surface from transforming into structures with monolayer steps. Conversely, with the application of LT‐AlN BL, the two ML steps observed using LT‐ AlN BL *T*g = 1100°C is kept, hence, preserving their heteroepitaxial relationship (as expected, 2 ML steps do not form SAGB). For BL *T*g = 800°C, it is believed that the "weak" heteroepitaxial relationship due to the undefined (rough) surface plays a role in circumventing the SAGB.

(10‐12) *ω*‐scans. The symmetric component is comparable to that grown under *Profile a* (FWHM is ∼64 arcsec), suggesting that both have highly‐oriented films along the *c*‐axis growth direction. On the other hand, the asymmetric (10‐12) linewidth is wider than that grown under *profile a* (FWHM is 1145 arcsec). **Figure 8(b)** shows the AFM surface morphology of AlN grown under the optimum LT‐AlN BL *T*<sup>g</sup> of 1050°C. Although nanopits have been reduced compared with other BL *T*<sup>g</sup> (**Figures 5(e)**, **(f)**), it is believed that the quality of the film can be further improved especially the twist mosaicity upon optimizing the buffer layer

To demonstrate the effect of eliminating the SAGB on the optical properties of AlN epilayer, we obtained the CL spectra for both profiles, as shown in **Figure 9**. The CL measurements were acquired at 93 K under 10 kV and 0.1 μA emission condition (spot size is ∼1 μm). CL peaks are assigned to free and bound excitonic emissions, including the LO phonon replicas, as shown in the inset figure [14]. Moreover, the emission intensity of AlN under *Profile b* using the optimum LT‐AlN BL *T*g of 1050°C is approximately more than two times higher than that under *Profile a*. This result is attributed to the higher probability of radiative recombination of electron and hole pairs due to the elimination small‐angle grain boundary. The differences in the CL peak position also suggest their different residual strain, where the tensile‐strained AlN under *Profile b* has smaller bandgap energy (∼6.011 eV) than that of an almost relaxed AlN (∼6.031 eV) under *Profile a*. The slow relaxation experienced by AlN under *Profile b* is most likely due to the reduced generation of dislocations upon the suppression of SAGB. As SAGB is a type of an edge dislocation, eliminating it is expected to enhance the optical properties of

**Figure 9.** CL spectra of AlN grown under (a) *Profile a* and (b) *Profile b* using optimum LT‐AlN BL *T*g = 1050°C.

thickness or other growth parameters.

52 Study of Grain Boundary Character

AlN.

On the other hand, the SAGB can also be prevented even when using thermally cleaned substrate. In fact, a lot of AlN growth optimizations have been performed under such condition and yet, there was no observation of SAGB or hardly any mention if at all. The reason could be due to the weakening of the epitaxial relationship between the sapphire and AlN after introducing some extrinsic factors such as nitridation before LT‐AlN BL or AlN seeding layer [11, 14, 26]. This could be the reason why the growth interruption by V/III ratio resulted in not only dislocation bending or coalescence but also the elimination of domain structure [18, 24]. Moreover, the insertion of an intermediate layer after LT‐AlN BL could have also eliminated the rotation domain [30]. In addition, the no observation of rotation domain on thermally cleaned substrate would also depend on thermal cleaning temperature or temperature ramp‐ up. In their study, substrate surface after thermal cleaning at 1145°C prior to LT‐AlN BL growth could have 2 ML step structure, thus preventing the SAGB [23]. Also, the initial supply of TMA (TMA preflow) during the pulse growth could have replaced the oxygen‐terminated surface into Al‐terminated one [25]. As mentioned earlier, the sapphire's surface with Al termination is expected to have no rotation domain.

Finally, **Table 1** shows the comparison between several techniques for eliminating the SAGB performed on as‐received and nitrided sapphire substrates. Note that techniques included in the table are only those that directly discussed the rotation domain. Although a high‐quality AlN could be obtained from nitrided sapphire substrate, residual SAGB still exists, albeit coalesced with other domain upon nitrogen radical or annealing treatment [19, 20]. Moreover, the step bunching of sapphire substrate into *n* × 2 ML, where *n* is an integer, would require a critical off‐cut angle during *ex situ* surface treatment [26]. On the other hand, the LT‐AlN BL technique would eliminate these additional processes owing to its in situ treatment. It is noteworthy that although rotation domain phenomenon was not observed in previous studies using LT‐AlN BL, we believe that whether intentionally or unintentionally, the elimination of rotation domain is one of its underlying purposes. For example, Xi et al. adopted this approach, and a high‐quality and atomically smooth AlN epilayer was obtained [22]. The same is true in the work of Zhang et al.; however, they introduced the PALE approach to enhance the migration of Al adatoms during the HT‐AlN growth [15]. But Hu et al. modified this method by performing nitridation pretreatment under H2/NH3 ambient prior to LT‐AlN BL growth. However, they ascribed the elimination of rotation domain to the different strain relaxation mechanisms induced by lattice mismatch [21]. In addition, nitridation pretreatment may induce rough AlN surface. Therefore, LT‐AlN BL is a promising technique for eliminating the rotation domain. It is noteworthy that LT BL has been used as a standard technique for the growth of GaN, resulting in crack‐free and high‐quality epilayer [31, 32]. However, to the best of our knowledge, this is the maiden report that clarifies the role of LT‐AlN BL for obtaining high‐quality AlN without SAGB.


**Table 1.** Comparison among several techniques reported for eliminating SAGB in AlN, where "+" sign shows the satisfied property, and the "‐" sign shows the unsatisfied property.
