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We report the nanostructure study results, mainly based on plan-view transmission electron microscopy (TEM) on the coalescence process during the overgrowth by metalorganic chemical vapor deposition of GaN nanocolumns grown by molecular beam epitaxy. In cross-section scanning electron microscopy images, one can observe a two-stage coalescence overgrowth process. First, a group of nearby nanocolumns is merged into a thicker column. One of the possible merging processes is the growth of a bridging domain between two columns for their connection. The thicker columns are then developed into horn-shaped structures for the second-stage coalescence. Because different columns may have different crystal orientations, stacking faults can be formed for implementing the coalescence between two domains. Such stacking faults around the boundaries of merged domains represent one of the major kinds of defect after the threading dislocation density is reduced based on the nanocolumn growth technique.
©2013 Optical Society of America
1. Introduction
Although free-standing GaN substrate has been available, currently sapphire is still the most widely used substrate for nitride compound growth. However, the large lattice mismatch between GaN and sapphire leads to a high threading dislocation (TD) density (109-1010 cm−2) in a GaN thin film. Several techniques have been introduced for reducing the TD density, including epitaxial lateral overgrowth [1–6] and patterned substrate growth [7–15]. The former type of technique can result in significant reduction of TD density to the level of 105-106 cm−2 if the patterned overgrowth is repeated by shifting the pattern [16]. Also, the later types of technique can also lead to a TD density around 105-108 cm−2. However, most of the aforementioned techniques involve in complicated process procedures. For fabricating GaN template of a low TD density, the coalescence overgrowth of GaN nanocolumns has become quite attractive because the lateral strains of nanocolumns can be relaxed to prevent the formation of TDs [17–26]. Almost TD-free GaN nanocolumns, grown with molecular beam epitaxy (MBE), can be quite straight, parallel, and vertical with respect to the substrate. Because the procedures for fabricating optoelectronics or electronics devices based on the column structures are normally complicated, the coalescence overgrowth of such GaN nanocolumns is crucial to provide us with a high-quality GaN template for practical device fabrication. Coalescence overgrowth of GaN nanocolumns, including those containing InGaN/GaN quantum wells, with MBE has been reported [27,28]. However, the detailed studies, particularly the coalescing mechanisms and the nanostructures formed during the overgrowth process, have not been discussed yet. Such a study is useful for understanding the overgrowth process and optimizing the growth conditions of the columns and the overgrowth layer. In addition to the GaN nano-columns-based approach, significant advances have been reported for addressing high performance mid / deep UV sources by using AlGaN-based [29–32] and AlInN-based [33,34] material systems.
In this paper, we use the techniques of scanning electron microscopy (SEM) and plan-view transmission electron microscopy (TEM) to study the nanostructures of the overgrown samples of GaN nanocolumns. The nanocolumns are grown with MBE and the coalescence overgrowth is performed with metalorganic chemical vapor deposition (MOCVD). From the microscopy images, one can observe the coalescence process of nanocolumns. Also, it is found that stacking faults are formed for connecting two domains of atomic misalignment when they contact each other for coalescence. Section 2 of this paper shows the sample growth conditions and the measurement parameters of SEM and TEM. In section 3, we report the plan-view TEM results and discuss the results. Further discussions are made in section 4. Finally, conclusions are drawn in section 5.
2. Sample preparation and measurement conditions
The plasma-assisted MBE-grown nanocolumns were grown on Si (111) substrate. Before growth, the substrate was cleaned with a dip in 10% aqueous HF solution, followed by rinsing in deionized water, before blowing dry. The substrate was heated to approximately 710 °C for the AlN layer. The active nitrogen flux was provided by an EPI Unibulb plasma source operated at 330W and 2.0 sccm in nitrogen flow rate. Substrate temperatures were measured using an optical pyrometer. Prior to igniting the nitrogen plasma, a small amount of aluminum was deposited onto the silicon surface to prevent the formation of silicon nitride. The AlN nucleation layer was first grown at a growth rate of 0.04 nm/sec before the deposition of GaN nano-columns. After the AlN nucleation layer was deposited, the substrate temperature was raised to 810 °C for growing GaN nanocolumns at a growth rate of 0.07 nm/sec. The nanocolumn cross-sectional size is about 100 nm with the column density estimated to be 109 cm−2. The height of the nanocolumns is about 2 μm. The MOCVD-overgrown GaN layers of about 700 nm in thickness were deposited at 800 and 900 °C in a Veeco P75 system for sample I and sample II, respectively. The growth pressure was 200 torr and the rotation speed was 1500 rpm. The V/III ratio was 2600 and the flow rate of NH3 was 1000 sccm. The growth pressure, the rotation speed, and the total gas flow rate were the same as those of standard GaN thin film growth. The estimated growth rate was about 0.29 nm/sec.
The TEM investigations were performed using a Philips Tecnai F30 field emission electron microscope with an accelerating voltage of 300 kV and a probe forming lens of Cs = 1.2 mm. In this study, we used the (0001) zone-axis beam to take high-resolution TEM images. Also, we used Gatan Digital Micrograph to process the electron beam diffraction pattern. The scanning electron microscopy (SEM) measurement was performed with a JEOL JSM 6700F system. In this system, we used a cool-field emission gun and a scanning probe form lens of Cs = 1.0 mm. The operation acceleration voltage is 15 kV.
Figure 1 shows the plan-view (a) and cross-section (b) SEM images of sample I. Here, one can see that domain structures exist on the surface with the domain size in the scale of a couple μm. In Fig. 1 (b), the blue/green arrow indicates the MOCVD-overgrown GaN layers and the plasma-assisted MBE-grown nanocolumns layer, respectively. One can see basically there are three layers, as divided by the two marked arrows. The top layer of about 700 nm in thickness corresponds to the overgrown layer. The middle layer shows thicker columns, which result from the merge of a few nanocolumns during the overgrowth process. The bottom layer, which demonstrates denser distribution of nanocolumn, results from the fill-in of GaN between nanocolumns during the overgrowth stage [35,36]. The lateral dimension of the original nanocolumns is around 100 nm. In the overgrown layer, one can see the growth of horn-shape structures to gradually achieve coalescence of different columns. It has been shown that by overgrowing a thicker layer (up to 2 μm in thickness), the domains shown in Fig. 1(a) become connected. To observe the merging process of nanocolumns, we prepared plan-view TEM samples with a double-side ion-milling procedure to obtain a thin layer for TEM observation around the circled portion of Fig. 1(b).
3. Plan-view TEM results
Figure 2(a) shows a plan-view TEM image demonstrating the connection of two hexagonal domains. The domain size is in the range of a couple hundred nm. The two domains correspond to two nearby nanocolumns. The circled area is magnified to give Fig. 2(b). Here, one can see the essentially continuous atomic arrangement across the boundary between the two domains. However, one can observe a region of stacking fault, as circled. It is believed that when the two domains, which are usually not perfectly aligned in atomic arrangement, are to be coalesced, stacking faults can be formed around the boundaries. The results in Fig. 2 also demonstrate that the crystal orientations among nanocolumns in proximity can be quite similar. In spite of the imperfect crystal alignments between two nanocolumns, they can still be coalesced through the formation of stacking faults.
Fig. 2 (a) A plan-view TEM image showing the coalesced connection between two hexagon domains of sample I; (b) A magnified TEM image showing the essentially continuous crystalline orientation across the coalescence boundary.
Figure 3(a) shows a plan-view TEM image around the junction of two domains of another similar sample (sample II). Parts of the circled region are magnified to give Figs. 3(b) and (c). Here, a line of stacking fault extending for a few hundred nm can be observed. Such an extended range of stacking fault can be used to explain the low cathodoluminescence intensity near the boundary of two coalesced domains [35,36]. Around the boundaries, the existence of defects in the form of stacking fault leads to low emission efficiency.
Fig. 3 (a): A plane-view TEM image of sample II showing the coalescence boundary of two domains; (b) and (c): Magnified TEM images showing the stacking fault structures for implementing the coalescence.
Figure 4 shows the connection between two hexagon domains and one relatively smaller transition region in sample I. The small transition region is supposed to be added to the sample during the overgrowth stage. We are interested in the conditions of coalescence between the added region and the nanocolumns on both sides. The portions indicated by the circle and the square are magnified to give Figs. 5 and 6, respectively. In the upper portion of Fig. 5, the middle lighter image portion corresponds to the boundary between the large and the small domains. Three regions denoted by A-C are assigned by the circles for Fourier transform analysis. The atomic diffraction patterns after Fourier transforms of the three regions are demonstrated in the lower portion of Fig. 5. Here, one can clearly see the three almost identical hexagonal patterns showing that the coalescence between the large and the small regions is quite successful.
Fig. 4 A plan-view TEM image showing the connection between two large hexagon domains and a small bridging domain of sample I. The areas within the circle and square are to be further analyzed.
Fig. 5 A magnified TEM image showing the area of the circle in Fig. 4. Three circled areas, labeled as A, B, and C, are used for diffraction pattern analysis to give the results at the lower portion of this figure for demonstrating the same crystalline orientation in the two coalesced domains and the boundary
Fig. 6 A magnified TEM image showing the region of the square in Fig. 4. Two square regions, labeled as D and E, are used for diffraction pattern analysis to give the results at the lower portion of the figure for demonstrating the different crystalline orientations in the two domains.
However, the coalescence between the middle small region and the large region on the left in Fig. 4 is not so successful, as indicated by Fig. 6. Here, we can first observe a clear gap between the two regions. The atomic diffraction patterns in the two regions, denoted by D and E, are shown in the lower portion of Fig. 6, in which quite different crystal orientations between the two regions can be seen. One can estimate that the orientations in the c-plane of the two regions differ by around 15 degrees. Based on the results in Figs. 5 and 6, one can speculate that the middle small region in Fig. 4 stems from the hexagon domain on the right such that they have the same crystal orientations. However, because the large crystal orientation difference between the two hexagon domains, the coalescence between the small bridging domain and the left hexagon domain becomes difficult.
Figure 7 shows another plan-view TEM image of sample I. Here, one can see two disconnected domains. From the TEM image, one can clearly see the different crystal orientations between the two domains. Their atomic diffraction patterns, as denoted by F and G, are shown in the lower portion of Fig. 7. Here, one can see quite different crystal orientations, differing by 6 degrees in the c-plane. With such a large orientation difference, the coalescence should be quite difficult. In a thicker overgrowth sample, in which the large-area coalescence is implemented, the overgrowth on one of the aforementioned domains may be terminated at a certain height such that other domains of more similar crystal orientations can be coalesced [35,36].
Fig. 7 A plan-view TEM image of sample I showing two domains in proximity of different crystalline orientations, as indicated by the diffraction patterns of regions F and G at the bottom of the figure.
4. Discussions
In Fig. 1(b), one can see the two-step coalescence process during the overgrowth stage. First, several nanocolumns of a couple hundred nm in individual size merge into a thicker column, as demonstrated in the middle layer of Fig. 1(b). Those nanocolumns can be quite close and have almost the same crystal orientations for easier coalescence. Figure 8 shows a plan-view SEM image of a nanocolumn sample grown by MBE under the same conditions as those for samples I and II. In this image, one can see the non-uniform nanocolumn distribution. Closely packed nanocolumns form a group of several units in one dimension. It is speculated that those nanocolumns in the same group have similar crystal orientations such that they can be merged in the early stage of overgrowth through the growth of small bridging domains, as shown in Fig. 4. After the formation of thicker columns in the middle layer of Fig. 1(b), they are developed into the horn-shape structures for further coalescence.
Fig. 8 A plan-view SEM image of a nanocolumn sample showing the grouped distribution of nanocolumns.
Because the nanocolumns stem from the naturally formed AlN nucleation layer, in an extended area, the c-plane orientations of different nanocolumns can vary. For connecting those columns, some of them may need to twist or tilt the orientations during overgrowth for merging into a continuous layer. When two columns touch each other in the lateral dimension, they can be coalesced through the formation of stacking faults for compensating the atomic misalignments. However, the high stacking fault density around the domain boundaries form the major defect source even though the threading dislocation density in such a coalesced overgrown nanocolumn sample can be tremendously reduced. If the nanocolumns can be grown from a GaN template based on certain patterned growth techniques [37–39], their crystal orientations can be more uniform for easier coalescence overgrowth. In addition to the use of coalescence overgrowth of GaN nanocolumn method, other relevant approaches by using nanoheteroepitaxy of GaN on nanopatterned sapphire substrates had been reported for achieving dislocation density reduction in the materials [40–42].
5. Conclusion
In summary, we have demonstrated the SEM images to show the two-stage MOCVD coalescence overgrowth of GaN nanocolumns. First, a group of nearby nanocolumns were merged into a thicker column. One of the possible merging processes is the growth of a bridging domain as demonstrated by the plan-view TEM images. The thicker columns were then developed into horn-shaped structures for the second-stage coalescence. Because different columns might have different crystal orientations, stacking faults could be formed for implementing the connection between two domains. Such stacking faults around the boundaries of merged domains represent the major defect centers after the threading dislocation density is reduced based on the nanocolumn overgrowth technique.
Acknowledgment
This research was supported by the National Science Council, The Republic of China, under the Grants NSC 100-2221-E-194-043, 101-2221-E-194-049, 102-2221-E-194-045, and 102-2622-E-194-004-CC3.
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