GaN-based optoelectronics on silicon substrates

Materials Science and Engineering B93 (2002) 77
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GaN-based optoelectronics on silicon substrates Alois Krost *, Armin Dadgar Institut fu¨r Experimentelle Physik, Otto-von-Guericke Universita¨t Magdeburg, PO Box 4120, 39016 Magdeburg, Germany

Abstract Cracking of GaN on Si usually occurs due to the large thermal mismatch of GaN and Si when layer thicknesses exceed approximately 1 mm in metalorganic chemical vapor deposition (MOCVD) preventing the realization of device-quality material. The thermal stress can be reduced significantly by a combination of different concepts such as the insertion of low-temperature AlN interlayers, introducing multiple AlGaN/GaN interlayers, and growing on prepatterned substrates. The growth of crack-free GaNbased light emitting diodes (LEDs) on silicon on patterned Si(111) with areas of 100 mm /100 mm is reported # 2002 Elsevier Science B.V. All rights reserved. Keywords: GaN; Si; Light emitting diodes; Metalorganic chemical vapor deposition

1. Introduction For the last 15 years silicon as a substrate has attracted much attention for the epitaxial growth of III
/V compounds like GaAs and InP because of its low price and its availability in large diameters up to 12 inches now. However, in spite of huge efforts, no real breakthrough has been obtained because of the high mobility of dislocations in these materials leading to a rapid degradation of all devices fabricated so far. In contrast, GaN-based devices are known to operate very well without aging effects with dislocation densities as high as 1010 cm 2. Thus, the integration of Si- and GaN-based devices on the same chip becomes feasible as well as a silicon based optoelectronics technology, e.g. with the potential for small, high resolution, full color displays.

2. Silicon substrate From the point of view of economics Si offers a low price as compared to sapphire and SiC, high crystalline perfection, availabiltity of large size substrates, all types of conductivity, and high thermal conductivity (1.5 W cm 1). In most cases the Si(111) plane is chosen because * Corresponding author. E-mail address: [email protected] (A. Krost).

of its trigonal symmetry favouring epitaxial growth of the GaN(0001) plane. The large difference in the lattice parameters of GaN (aGaN /0.31892 nm) and Si (aSi(111) /0.38403 nm) yields a lattice parameter mismatch (f//16.9%) resulting in a high dislocation density of
/1010 cm 2 which is comparable to Ga on sapphire. The most severe problem is the large thermal dilatation mismatch between GaN and Si. The in-plane thermal expansion coefficient GaN is 5.59 /106 K 1 [1] as compared to 3.77 /10 6 K 1 [2] of Si (another value, often cited is 2.59 /10 6 K 1), which leads to a large tensile stress during cooling from the growth temperature to room temperature often resulting in cracked layers preventing device applications. The tensile stress causes a concave bending of the film/ substrate system. The stress values can be easily determined via the curvature of the sample which is proportional to the stress value [3]. The curvature leads to a strong broadening of the Bragg peaks in X-ray diffraction of the GaN thin film and the Si substrate. The radius of curvature (R ) can be measured, e.g. by measuring the width of the Si(111) Bragg peak using different apertures. Under typical MOCVD growth conditions, the stress amounts to 0.9 GPa mm 1 GaN [4]. There are three equivalent primary crack directions along [1120], [1210], and [2110]. The resulting cleavage ¯ ¯ (1010), and ¯ planes are¯ (1100), (0110), respectively. The ¯ ¯ epitaxial relationship is (0001)GaN¯ parallel (111)Si with [110]Si parallel [1120]GaN, and [112]Si parallel ¯ [1¯ 100]GaN. Since the ¯primary clavage planes of Si are ¯

0921-5107/02/$ - see front matter # 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 1 0 7 ( 0 2 ) 0 0 0 4 3 - 0

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of type {111} with Ž110 cleavage directions, GaN and Si always have a common cleavage direction, e.g. [1120]GaN and [110]Si. These relations are summarized ¯ ¯ in Fig. 1.

3. Cracking problem There are several possibilities to overcome the cracking problem. The critical thickness at which cracking occurs depends on the growth temperature. In MBEgrown samples the critical thickness is higher than in MOVPE grown because of the lower growth temperature. Indeed, crack-free GaN/AlN/Si(111) samples up to 3 mm in thickness have been reported [5]. The GaN layers grown at 790 8C exibit a residual tensile strain which depends on its thickness. From the energy position of the A free exciton, tensile stress values for 1 and 3 mm thick undoped GaN films of 0.3 and 0.7 GPa are deduced. This is in contrast to the findings of Nikishin et al. [6] who deduced from the shift of the E2 phonon in Raman scattering experiments an in-plane biaxial stress of
/0.16 GPa in GaN/AlN/Si(111) layers grown by gas source molecular beam epitaxy at 1000 K, independent of their thickness. Possibly, the critical thickness at which cracking occurs not only depends on the growth temperature but also on the quality of the buffer layer. According to the same authors the formation of amorphous SiNx favors cracking. It could be avoided by a proper nucleation process starting by exposing the Si surface to an Al flux and subsequent exposure to ammonia with the Al flux off. Repeating this process five to ten times suppresses the formation of SiNx . We found two different types of cracks: one type with open grooves with facets indicating that they occurred during growth (Fig. 2a) and the second type without grooves which most probably were generated during the

Fig. 1. Epitaxial relationship GaN(0001) on Si(111).

Fig. 2. Two types of cracks for GaN on Si: open groove with facets occurs during growth (a) and closed crack during cooling down (b); cross-section of open crack (c).

cooling process (Fig. 2b). Cracks which occur during growth are detrimental for the layer as well as for the substrate. As can be seen in Fig. 2c) the crack deeply (some mm) extends into the Si substrate. At high temperatures there is a strong chemical reaction between Ga and Si, which leads to deterioration of the substrate and the growing layer (Fig. 3). From an EDX analysis in

Fig. 3. Meltback etching causes a deteriotation of GaN layer and Si substrate; Nomarski image (top) and cross section TEM (bottom).

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TEM cross-section Ga-rich, Si rich and SiNx regions could be identified.

4. Buffer layers 4.1. Direct growth of GaN on silicon Direct epitaxial growth of GaN on Silicon suffers from a strong tendency towards the formation of amorphous Si3N4 when the Silicon substrate is exposed to ammonia [7]. Ishikawa et al. [8] studied the thermal stability of GaN on Si (111) substrates grown by MOVPE. They observed a strong degradation of low temperature (LT) grown (530 8C) GaN on Si upon thermal annealing at 1050 8C in nitrogen and ammonia ambients. From cross sectional SEM images a rough interface with hollows and gallium droplets are found. It was concluded that LT-grown GaN decomposes and a strong reaction between Si and elemental Ga is responsible for the degration. Therefore, LT-GaN is not suitable as buffer layer because the melt-back etching of Si by Ga. Chen et al. [9] deposited amorphous GaN buffers layers on Si (111) at 300 8C by MOVPE. The amorphous GaN layers crystallize upon annealing between 500 and 600 8C. In HRTEM an amorphous interfacial layer is detected which is considered to be Si3N4 which has been formed during the annealing process of the GaN buffer layer. In spite of the amorphous Si3N4 interlayer GaN grown on top of the amorphous buffer at 900 8C has a better structural and optical quality than GaN grown under the same conditions on LTbuffer GaN layers deposited at 550 8C. The formation of isolated islands of amorphous Six Ny preventing direct epitaxial growth was also observed when growing GaN on silicon in MBE by Ploog et al. [10]. Also Hashimoto et al. [11] reported on the formation of silicon nitride in the initial stages of the growth of GaN on Si. They have shown that a ten-monolayer thick Ga layer deposited at 500 8C in MOVPE suppresses the formation of silicon nitride considerably. In addition to this direct reaction of nitrogen with Si to form silicon nitride, Si was found to migrate to the surface at growth temperatures of 950 and 1010 8C during MOVPE growth as evidenced by photoelectron microscopy and X-ray absorption spectroscopy leading to the formation of SiNx on the top surface and to a severe surface roughening. In conclusion, GaN is not well suited for direct epitaxial growth on silicon substrate.

Kobayashi et al. [12] used oxidized AlAs as an intermediate layer for the growth of GaN on Si. From cathodoluminescene measurements, an oxygen contamination of the GaN layers seems to be possible. Strittmatter et al. [13] introduced successfully a AlAs buffer in a two-step growth process. Firstly, a lowtemperature AlAs nucleation layer is grown at 450
/ 500 8C followed by 50
/100 nm AlAs at 750 8C. Subsequently, the AlAs is nitridated at 970 8C whereby an As
/N exchange takes place resulting in an epitaxial AlN layer which served for the growth of high quality GaN on top of it with X-ray rocking curve width B/600 arcs for 1.2 mm GaN. From cross sectional TEM (Fig. 4) no formation of silicon nitride was observed at the interface and from secondary ion mass spectroscopy the Si incorporation was determined to be below the detection limit of 1/1016 cm 3. Using such buffer layers on n-type Si(111) a simple LED structure with a single InGaN/GaN quantum well was realized with (Al/ Au) n-contacts on the back side and 10 nm Pt p-contacts on the front side of the sample. Despite the formation of cracks and pits, which act as potential short circuits paths, and a high series resistance (12.5 V at 20 mA), the sample emitted at 439 nm, the luminescence being well visible at daylight [14]. Interestingly, the electroluminescence of a reference structure grown on sapphire substrate occurred at 429 nm. The redshift in the LED on Si was attributed to the high stress which persisted in spite of the cracks. 4.3. AlN buffer layer It is well known since thirty years, that AlN can be grown as single crystalline films on Si [15,16]. Amano et al. [17] were the first who used AlN as an intermediate buffer layer between GaN and Si. Meanwhile, AlN is used widely as a buffer layer for the growth of GaN on Si [18
/30] including ourselves, [31]. There is common sense that AlN of highest structural quality can be grown by MOVPE only at very high temperatures of 1100 8C and above [27,32,33]. How-

4.2. AlAs buffer layer The nitridation problems can be avoided using an ammonia-free atmosphere at the first growth step.

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Fig. 4. Cross-section TEM of GaN on nitridated AlAs/Si(111).

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ever, when the AlN buffer layer growth temperature is too high (1200 8C), the GaN layers on top of it showed an inclined c-axis and also pits where no coalescence occurred. In addition, inversion domain boundaries are observed in these layers and high-resolution TEM images reveal the presence of a non-continuous amorphous SiN layer. The effect is attributed to the beginning deteroriation of the Si substrate at 1200 8C. Surprisingly, GaN films on top of optimised AlN buffer layers (1150 8C) show a much broader rocking curve (1020 arcsec) as compared to GaN layers with AlN buffer layers grown at a lower temperature (1060 8C, 656 arcs). As-grown GaN layers are all highly resistive showing that high-quality AlN serves as an efficient diffusion barrier for Si. AlGaN/GaN heterostructures grown on top of AlN/Si(111) show a dimensional electron gas behaviour with mobilities of 725
/813 and 2200 cm2 V 1 s 1 at 300 and 77 K, respectively [32]. These values are comparable to those obtained by Schremer et al. [34] on AlGaN/GaN heterostructures grown by flow modulation MOVPE on 650 nm GaN/30 nm AlN/Si(111) with m300 K /600
/920 cm2 V1 s 1 and m77 K /1580 cm2 V 1 s 1. It should be noted that a 1.5
/3.7 nm amorphous interlayer believed to be SiNx was revealed by HRTEM on such structures which on one hand supports efficient relaxation of the large lattice mismatch f300 K //23.4% at the interface between AlN and Si(111) [35] and on the other hand causes a grain like GaN/AlN layer structure where the c -axes of the grain are tilted relative to the Si[111] direction. Even higher mobilities m300 K /1620 cm2 V 1 s 1, m77 K
/7000 cm2 V 1 s 1, m20 K /7500 cm2 V 1 s 1 are reported on AlGaN/GaN heterostructures grown by MBE on thick (up to 3 mm) GaN/50 nm AlN/Si(111) layers [5]. The high mobility prevails even at elevated temperatures m450 K /660 cm2 V1 s 1. This was achieved by optimising the quality of the AlN buffer layer using RHEED to rapidly obtain perfect layer by layer growth at a very low growth rate of 0.1 mm/h, a nucleation temperature of 650 8C followed by a ramping to 900 8C. The thick GaN layers were grown at 790 8C. The main parameter for the high mobility values at 300 K is suggested to be an extremely low residual free carrier background (Fig. 5). 4.4. Low-temperature AlN interlayers Amano et al. have [36] demonstrated that the insertion of low-temperature (LT) AlN interlayers between high-temperature (HT) GaN layers reduces the stress, improves the layer quality and allows for the growth of thick, strained AlGaN layers on GaN/sapphire. As has been shown by Dadgar et al. [4] these benefits of the insertion of LT
/AlN interlayers are also valid for the growth of GaN on Si. By introducing thin LT
/AlN interlayers the crack density was practically reduced

Fig. 5. Cross-section TEM of GaN/Si(111) with LT
/AlN interlayers. GaN on patterned Si(111). The ridges are oriented along Si [112]. ¯

from 240 mm 2 to zero in a 1.3 mm thick sample. A strong improvement was seen in the surface sensitive (2024) Bragg reflection in grazing incidence which ¯ showed a decrease in the full width at half maximum of the rocking curve from 270 to 65 arcs for a sample without and with LT
/AlN interlayers, respectively. It should be noted that the mechanisms leading to the stress reduction and structural improvements with LT
/ AlN interlayers are not clarified yet.

5. Compliant substrates Another possibility to reduce the cracking problem is the growth on compliant substrates. To release epilayer strain, compliant substrates have been proposed for lattice mismatched epitaxy [37]. The idea is that the strain in a lattice mismatched epilayer can be reduced via partial accomodation of the total strain in a thin compliant overlay. Silicon-on-insulator (SOI) substrates are fabricated commercially by SIMOX technology (separation by implanted oxygen). Hereby, a thin silicon surface layer (
/50 nm), which is separated by a thin (
/80 nm) SiO2 buried layer from the thick Si(001) or (111) substrate, acts as the compliant layer. Steckl et al. [38] have converted Si(111) SOI structures to SiC by carbonization of the thin Si layer. MOVPE growth of GaN on such substrates resulted in high-quality material (corrected FWHM of (0002) rocking curve of
/360 arcs) comparable to GaN grown on 6H
/SiC substrates. Using such Si (100) substrates Cao et al. [39,40] obtained GaN layers with a full width at half maximum of the GaN (0002) Bragg reflection of 366 arcs by MOVPE. Best results are reached when the Si surface was nitridated at 1040 8C prior to the GaN buffer growth

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at 500 8C and bulk growth temperature at 940 8C. Scanning electron microscopy revealed a rather rough surface showing hexagonal hillocks probably due to the growth on the amorphous silicon nitride interfacial layer caused by the nitridation process. Unfortunately, it was not stated whether the layers exhibit cracks or not. A similar attempt to reduce the stress was tried by Koh et al. [41] with N -implanted Si(111) substrates.

6. Wafer bonding An alternative approach to stress releasing is wafer bonding. The applicability of wafer bonding of GaN with Si was firstly demonstrated by Wong et al. [42]. A selective laser lift-off technique, in conjunction with a Pd
/In wafer bonding process was utilized to integrate GaN thin films, originally grown on sapphire substrates, with Si substrates. Recently, GaN was integrated with Si by the wafer bonding method using AuGe as an adhesion material [43]. GaN was first grown on AlAs/ GaAs(100) and either the backside of GaAs was bonded to Si or the GaN frontside with subsequent removal of the GaAs has been demonstrated. It should be stated, however, that the wafer bonding method is rather complicated and expensive as compared to direct growth methods.

7. LEDs on silicon The first GaN-based light emitting diode on Si substrate was reported by Guha et al. [24]. The double heterostructure was grown by MBE on n-type Si(111) with an intermediate 8 nm AlN buffer. The structure consisted of n-Alx GaN/6 nm GaN (Si-doped or undoped) as an active layer/p-Alx GaN/15 nm p-GaN layers with 0.05 B/x B/0.09. Ni/Au thin (14 nm) transparent served as p-type contacts and electron injection was carried out from the backside through the Si substrate. The diodes start light emitting at 4.5
/6.5 V with reverse leakage currents from 10 to 130 mA at /10 V. At 12 V the forward currents varied from 14 to 65 mA. These rather high values as compared to MOVPE grown devices on sapphire or SiC were attributed to a low p-type doping and nonoptimal p contacts. A device with a Si doped thin GaN layer showed a near band edge electroluminescence at 360 nm with a full width at half maximum of 17 nm, and a broad long wavelength tail that extended out into the visible spectral range, while a heterostructure with an undoped GaN layer showed a broad emission band centered at 420 nm most pobably due to deep radiative levels in the gap. The same authors reported on multicolored light emitters on silicon substrates using similar violet MBE-grown GaN LEDs as described above with somewhat higher Al-

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content (x/0.15). In conjunction with organic dye based color converters orange at
/600 nm and green
/yellow at
/530 nm electroluminescence on the same Si wafer is obtained. The output power was not given but the ‘visible part of the electroluminescence was bright enough to be clearly observed by the eye under normal room illumination’. It should be noted that the layers showed cracks. Tran et al. [44] reported the growth of InGaN/GaN multiple quantum well (MQW) blue LEDs on Si(111) grown by MOVPE. The structure consisted of a 20 nm AlN buffer deposited at 750 8C, 4 mm n-doped GaN, an undoped ten period MQW (2 nm In0.22Ga0.78N/9 nm GaN), a 40 nm p-doped Al0.1Ga0.9N layer and 0.3 mm pGaN cap layer. The structure showed blue electroluminescence at 465 nm. Light emitting started at 4 V, the reverse leakage current was 60 mA at /10 V. An optical power was not given. This structure showed also cracks. Yang et al. [45] fabricated an InGaN/GaN MQW LED by a combined MBE/MOVPE growth procedure in selective areas defined by openings in a SiO2 mask. The density of cracks was comparable to similar structures on flat SiC substrates. For the LED a forward turn-on voltage of 3.2 V was measured. The forward differential resistance was a factor of four higher than in comparable LEDs on sapphire substrate. At room temperature the device emitted at 465 nm. An MBE-grown ultraviolet electro-luminescence GaN/AlGaN single hetero-junction LED on Si(111) was also reported by Sa´nchez-Garcia et al. [46]. Room temperature electroluminescence centered at 365 nm with a FWHM of 8 nm was obtained. The turn-on voltage was 5 V, the structure suffered from a reverse leakage current of 200 mA at /5 V. The optical ultraviolet output power was estimated to be 1.5 mW at 35 mA. Adachi et al. [47] reported on MOCVD growth of an InGaN/GaN 3QW LED using Al/N/AlGaN intermediate layers. The intermediate layers resulted in a high resistance when using backside contacts. In addition, cracks are observed.

8. Reduced area growth Another way towards stress releasing could be the growth on reduced areas. Recently, Strittmatter et al. [48] used a maskless epitaxial overgrowth technique for the growth of high quality GaN layers on structured Si(111). The substrates were structured with parallel grooves along the [110] or [112] directions. GaN grows ¯ ¯ bottom of the grooves with two growth fronts on the and on top of the ridges. The material on the ridges can extend with its side wings over the grooves without contacting the substrate, similar to the pendeo-epitaxy process. The process is completely mask-free and does

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not require any growth interruption. Depending on the stripe orientation different type of facets with fast or slow lateral growth rate develop. When the grooves are aligned along Si[110] or Si [112] the fast or slowly ¯ growing (1122) facets or (1101)¯ facets are observed, ¯ respectively. In the former ¯ case, the layer can be planarized more easily, however, there are also cracks. Vice versa, planarization is more difficult to obtain in the latter case, with the crack density drastically reduced. When the ridges were oriented parallel [110], ¯ the material in the wing region was tilted by
/0.58 with respect to that on the ridges, as measured by X-ray diffraction. Remarkably, there is nearly no tilt of the wings when the stripes are oriented parallel Si[112] as shown in Fig. 6. As with other ELOG techniques¯ this process seems to be very promising to reduce the dislocation density as visualized by the surface morphology and the photoluminescence spectrum.

9. Growth on patterned substrates The idea behind is to reduce the dislocation density, e.g. by lateral epitaxial overgrowth and to releasing the stress via the free side facets. GaN/Si growth on reduced area, patterned substrates have been reported by several groups: Bidnyk et al. [49] reported on laser action in GaN pyramids grown on (111) silicon by selective lateral overgrowth. Kobayashi et al. [50] used the planar epitaxial lateral overgrowth technique for the overgrowth of GaN over AlOx . The AlOx was prepared by growing AlAs and GaAs on Si(111). After patterning GaAs stripes AlAs was converted cmpletely to AlOx by a wet oxidizinge steps. During this process the GaAs was partially converted to Ga2O3 stripes which served as nucleation templates for the GaN growth. Tanaka, Kwaguchi et al grew selectively GaN pyramids by MOVPE on Si(111) with a patterned dot structure of a SiO2 mask and on SiO2 stripe-patterned Si using an intermediate AlGaN layer [51
/53]. Seon et al. [54] grew high quality GaN/Si(111) on large window fields ran-

Fig. 6. X-ray v -scans arounnd GaN(0002) parallel and perpendicular to the Si ridges.

ging from 0.2 to 10 mm by gas-source MBE. Marchand et al. [29] grew on SiO2-patterned AlN/Si(111) substrates using the lateral epitaxial overgrowth technique. The 5 mm wide stripes were oriented in the [1100] direction, the width of the SiO2 mask region was ¯35 mm. Cracking occurred only when the stripes came into contact and formed a coalesced film. With 60 nm AlN buffer layer thickness the stripes were bound by the (0001) facet on top, by vertical {1120} facets on the edges, and by ¯ inclined {1122} sidewalls. The distance between cracks ¯ on a given stripe was in excess of 300 mm. A threading dislocation density B/106 cm 2 is reported. Zamir et al. [55] grew GaN on Si(111) substrates patterned into square mesas of different sizes by reactive ion etching. The trenches were 0.5
/0.6 mm deep and 2
/4 mm wide. It was found that crack-free GaN with a thickness of 0.7 mm could be grown on square islands smaller than 14 mm in size. We used square-patterned Si(111) directed along Si[110] and Si[112] with SixNy masks prepared by ¯ ¯ standard photolithography and wet chemical etching. Openings of 100/100 mm2 could be overgrown with 3.6 mm thick GaN without any cracks in the GaN (Fig. 7). The GaN side walls consisted of (1122) and (1101) ¯ ¯

Fig. 7. Crack-free GaN-based LED structure on masked-patterned Si(111) 100 mm  100 mm regions. There are only cracks in the masked area between the openings.

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facets as indicated in Fig. 7 (bottom). The two types of facets showed a completely different emission spectrum as examinated by spatially resolved cathodoluminescence (CL) measurements (Fig. 8). The CL spectrum of the (1101)-type facet is dominated by a strong pairband ¯ emission indicating a strong impurity incorporation, whereas CL of the (1122)-type facet shows a sharp ¯ donator related bound exciton emission indicating low impurity incorporation and high crystal quality. These findings can be understood in terms of the different growth velocities of the facets in MOVPE growth conditions: From cross-sectional electron micrographs the (1122)-type facet is a neutral one and favoured as compared to the nitogen terminated. As in the case of planar structures, 15 pairs of Al0.18Ga0.82N/GaN layers were inserted which resulted in a strong reduction of the tensile stress from 1.16 to 0.35 GPa otherwise observed by X-ray diffraction. In such structures cracks are only observed in the 10 mm wide Six Ny masked regions separating the GaN squares. On top of such material threefold InGaN/GaN multiquantum-well light emitting diodes were grown. Bright blue room-temperature electroluminescence centered at 422 nm with a light output of 100 mW at 20 mA was observed. In summary, low-cost, crack-free GaN light emitting diodes on Si substrates on device-relevant areas were demonstrated with an otical output power of 100 mW. Thus, an integration of III
/V optoelectronic devices with Si becomes feasible.

Fig. 8. Cathodoluminescence spectra of different type of facets of GaN on patterned Is.

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Acknowledgements This work was done within a fruitful collaboration with Ju¨rgen Christen and his group at OvG University, Magdeburg. Financial support by the Deutsche Forschungsgemeinschaft under contract no. KR 1239/10-1 is gratefully acknowledged.

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