be observed that the HRTEM image of the real boundary shows one bright layer of spots framed by a darker layer on each side. This feature is not reproduced by any of the sim-ulated images, indicating that some other structure than antibixbyite could be present at the IDB.
In x Ga 1 − x N islands
This chapter starts with a literature survey about InGaN islands and QDs. In the following, the term island refers to any three dimensional structure, whereas the term QD is used if QD like luminescence was observed or - in the case of referring to the results of other authors - if these authors used the term QD for their own results.
The experimental results concerning the InGaN islands are subdivided into three dif-ferent sections. As Ga or In droplets can form on the surface during epitaxial growth of the specimen, it is important to know if there are metallic droplets on the surface and how they can be distinguished from InGaN islands. Therefore, the first experimental section addresses the recognition of droplets. The second section describes results from InGaN island samples grown by MBE. Finally, the results obtained for MOVPE grown InGaN samples are presented in the third experimental section.
For the uncapped MBE samples and some of the uncapped MOVPE grown samples, InxGa1−xN island with distinct size and geometry are observed on the surface. In the following, these samples and capped samples where the same InGaN growth conditions were employed will be referred to as “geometric InxGa1−xN island samples”. Changing the growth approach for MOVPE samples, a series was obtained where some samples exhibit QD like luminescence centres as observed by PL. This sample series will be referred to as “InxGa1−xN sample series with QD like PL”, as no geometric islands or In rich regions could be identified up to now using TEM.
5.1 Literature survey
The literature survey is meant to give an overview over the topic of InGaN islands. It cannot give a complete description of the whole literature available on InGaN islands, QDs, and related topics. The main purpose is to describe different aspects which are interesting or important in view of this chapter.
71
There are several approaches to realise three dimensional growth of islands and QDs [44]. In the following, the different approaches for island and QD growth of InGaN will be presented.
5.1.1 Patterning
One approach for realisation of QDs is patterning. This can be done by structuring an InGaN QW e. g. by use of a focused ion beam (FIB).
On the other hand, the surface on which the InGaN QDs are intended to grow can be patterned before deposition of InGaN. Kapolnek et al. [106, 107] used a SiO2 mask and lithography to induce position selective growth on underlying (0001) GaN. Further growth of GaN by MOVPE resulted in hexagonal GaN pyramids with {1101} facets. For larger mask openings or equivalently shorter GaN growth time, a {0001} top facet is present.
Haffouz et al. [108] showed that the pyramidal form is maintained for Mg doped GaN.
In case of n-type doping with Si, a change to columnar growth with {1100} side facets is obtained for higher flow rates of the Si precursor. This is explained by a strong increase of growth rate in h0001i. After the patterning and the site selective growth of the GaN pyramids, InGaN can then be deposited [107, 109, 110]. Tachibana et al. [111, 112] (see also [113, 114]) used this technique to grow nominal In0.2Ga0.8N layers on top of thin In0.02Ga0.98N. The cap layer also contained 2% In. The lateral size of the resulting QDs is supposed to be 30 nm at most and is given by the radius of curvature of the tips of the pyramids. The FWHM of emitted luminescence measured by PL was reported to be as low as 0.65 meV [115]. P´erez-Sol´orzano et al. [116, 117] explained the formation of QDs at the tip of the pyramids by the relaxation of InGaN. The incorporation of In at the tip produces less strain compared to incorporation at the side facets. Therefore, they expected a higher In concentration at the tips. Nevertheless, also the temperature of the substrate could be lower at the exposed GaN tips, which have a height varying between 0.4 µm and 4 µm.
At lower temperatures, the decomposition of InGaN and reevaporation of In is reduced.
Therefore, a reduced temperature at the tips of the pyramids can also be considered to lead to an increase of In concentration at these sites.
Another growth approach for QDs using structuring is to damage the subjacent GaN by e. g. creating small holes in the GaN with a FIB. On further growth, the stress on the surface above the damaged region is changed, which can lead to positions for preferential incorporation of In. This stress driven growth is observed for other material systems, where multiple layers of islands exhibit a vertical correlation due to the strain field of buried islands [118].
The advantage of the selective growth of InGaN islands is their high homogeneity in size and in In concentration of different islands. These parameters strongly influence the luminescence and hence the performance and reproducibility of intended devices. For a
single QD the luminescence isδ-like comparable to an atom. For an ensemble of QDs with a broad distribution of sizes and In concentrations, theδ-like emission lines would overlap, resulting in one broad luminescence peak. These distributions can be controlled quite well by selective growth. Nevertheless, the density of islands is limited by the structuring technique used, and is much lower than what can be reached with other growth techniques for QDs. If high densities of islands are required, as is the case for QD lasers, other growth techniques for QDs need to be considered. In addition, the structuring process can induce severe damage of the surface and thus can be detrimental for luminescence properties of intended devices.
5.1.2 Antisurfactants
InGaN islands can also grow self assembled. There are different approaches to achieve self assembled growth. One is the self assembled growth by use of an antisurfactant, which can alter the sticking coefficients of impinging species, change the surface diffusion or the surface energy. Several reports exist where Si was used as antisurfactant [119–126]. Keller et al. [119, 120] observed spiral InGaN islands at dislocations, grown on a GaN/sapphire substrate by MOVPE. These authors argue that the Si pretreatment causes passivation of surface sites. During the subsequent growth of InGaN, the impinging species propably diffuse to the more “reactive” dislocation cores, where spiral growth can be observed in case of dislocations having a Burgers vector with screw component. As dislocations can lead to non radiative recombination, the location of the QDs at dislocations can be detrimental for the optical properties of devices. Hirayama et al. [121] grew InGaN QDs of 10 nm diameter and 5 nm height on Al0.12Ga0.88N using Si as antisurfactant. Despite the effects of Si on evoking three dimensional growth, the choice of a subjacent AlGaN layer with larger a lattice parameter compared to GaN increases the strain for the InGaN and can be considered to facilitate the strain induced Stranski-Krastanow (SK) growth mode for InGaN islands. N¨orenberg et al. [122] grew InN on Si pretreated GaN on a Si(111) surface and also succeeded to grow three dimensional InN. Nevertheless, the growth of InN islands directly on Si(111) was not successfull. Cheng et al. [124, 125] analysed the effect of Si doping of barriers and InGaN QWs on formation of QDs. While they observed only weak fluctuations of In for undoped InGaN QWs, the well doped and especially the barrier doped QWs are reported to contain QDs of (3. . .5) nm diameter. This also points out the role of Si not only as dopant but also as antisurfactant. If n-type doping is not desired, In can be used as antisurfactant. This was studied by Zhang et al. [123], who pretreated an In0.01Ga0.99N surface with In before growth of InGaN with higher In content. With this approach they obtained In rich QDs. The capping of the about 4 nm wide and 1.5 nm high QDs was done with 4 nm In0.01Ga0.99N and 5 nm GaN. As explanation for the QD growth they propose that during In pretreatment, In−N bonds are formed at the sample surface.
This is claimed to reduce the sticking coefficient of Ga [127] and thus the Ga content of the
growing material. Furthermore, if positions of higher In concentration are formed, the local a lattice parameter can be larger. The further incorporation of In is thus more favourable at these positions, as the strain energy is reduced in comparison to In incorporation directly on GaN or low In content InGaN. This is also called composition pulling effect or substrate pulling effect [128]. The In antisurfactant induced QD growth was not observed by Han et al. [126]. On the contrary, they observed for their growth parameters a step flow growth mode after In pretreatment and argue that the composition pulling effect is responsible for the formation of the QDs observed by Zhang et al. [123], rather than the effects of In on the surface kinetics. Han et al. [126] analysed also the effect of Ga pretreatment prior to InGaN growth. In comparison to growth directly on GaN, the observed QDs show the tendency to agglomerate and coalescence. The most homogeneous size distribution of QDs was obtained by passivating the GaN surface for 24 h in air [126]. After a cleaning procedure, a GaN seeding layer was grown by MOVPE at only 550◦C, resulting in GaN QDs [129]. On this seeding layer, InGaN QDs were grown at Tg = 850◦C. The authors attribute the successful formation of the QDs to a reduction of the surface diffusion length, which is proportional to e−E/2kB T. Here, E is the activation energy for surface diffusion, T the temperature, and kB the Boltzmann constant. In case of a small surface diffusion length, the adsorbate is limited to a smaller area, which can induce three dimensional growth. With the described growth method, the surface diffusion length is decreased by decreasing the growth temperature T and additionally by increasing the activation energy in this case by the passivation process. Furthermore, surface contaminations which are not removed during the cleaning of the passivated surface may act as nucleation centres for three dimensional growth [130]. For growth of multiple InGaN QD stacks using this method [130–132], the QD density decreases, the average QD height increases, and the homogeneity of the QD sizes worsens for increasing stack number. Further drawbacks of this method are the complicated growth procedure especially for multiple stacks, which requires an air passivation for each single InGaN QD layer, and the quite uncontrolled incorporation of foreign atoms in the heterostructure.
5.1.3 Stranski-Krastanow growth mode
Another approach for the realisation of self assembled QDs discussed here is the strain induced Stranski-Krastanow (SK) growth, which was already schematically described in section 2.3. Many groups report the successful realisation and analysis of InGaN islands grown via the SK growth mode [133–148]. Generally, the island density decreases with increasing growth temperature (see e. g. [149]). A detailed analysis of growth modes of InGaN on GaN grown by MOVPE was published by Oliver et al. [148]. They report the transition from two dimensional to three dimensional SK growth mode for increasing pressure in the growth chamber. By analysis of the composition it was additionally observed that the In concentration was decreased for increasing pressure. This means, that even
though the strain in the InGaN is reduced, the transition to three dimensional growth takes place, i. e. the strain in the film is not the only driving force for island growth, and other aspects such as growth kinetics have to be considered. Feuillet et al. [133] report the growth of InN islands on GaN by MBE and state that these islands were formed by the SK growth mode. The InN islands, which are grown at 420◦C, are completely relaxed due to formation of misfit dislocations. The InN grew completely unstrained right from the start of growth, i. e. the InN was already relaxed before islands were formed, and a transition to three dimensional growth was observed after deposition of approximate 2 ML, as was determined from the RHEED pattern. Ng et al. [142] analysed the growth of InN on GaN by MBE with RHEED. For temperatures exceeding 420◦C and simultaneous In rich growth conditions, two dimensional growth was observed, else three dimensional SK growth. The evaluation of the a lattice parameter of the InN revealed that relaxation in (0001) starts immediately. For SK growth, islands occur after growth of about 6 ML, when approximately 80% of the strain is relieved. They suggest that the InN is first partly relaxed by formation of misfit dislocations (with a critical thickness for misfit dislocation generation less than 2 ML [142] for InN on GaN according to the criterion proposed by Matthews and Blakeslee [49]). Then, the remaining strain is released by island formation.
The islands are supposed to contain dislocations and to be free of strain. Grandjean and Massies [134] analysed MBE grown InxGa1−xN of concentrations up to approximately 0.5.
They observed strain induced island formation for x exceeding 0.1 [134] or 0.12 [135].
The critical wetting layer (WL) thickness for the transition from two to three dimensional growth was determined by RHEED and decreases from 7 ML for x = 0.12 to 2.5 ML for x= 0.47. Below an In concentration of 0.1, no island formation was observed but surface roughening occurred for thicknesses exceeding (9. . .13) ML. In comparison, Adelmann et al. [136,138] and Daudin et al. [137,139] report a SK growth if the In concentration exceeds 0.18. For In0.3Ga0.7N QDs analysed with PL they observed a FWHM of the luminescence peak of 150 meV, which they attribute to size or concentration fluctuations. Taliercio et al. [141] grew InGaN QDs by MBE. They explicitly stated that even though they assume the In concentration to be constant in their QDs and WL, an In enrichment in the QDs could be possible due to “strain induced migration”.
An advantage of the SK growth mode is that the densities of InGaN islands are usually high (&1010 cm−2). Even though the distribution of islands in one InGaN stack may not be homogeneous, multiple stacks may be vertically correlated due to strain fields of buried islands [118]. This can lead to very homogenous island distributions, as was for example observed in case of GaN QDs in AlN [150–152].
5.1.4 Growth interruptions
Ji et al. [153–156] and Hung et al. [157] used a growth interruption technique to realise InGaN QDs using MOVPE and also grew light emitting diode (LED) structures. They
grew nominally In0.3Ga0.7N at 730◦C. After deposition of 4.5 ML they stopped for 12 s before depositing another 4.5 ML InGaN. The QD density of the samples with growth interruption (2·1010 cm−2) is increased by about one order of magnitute compared to an InGaN structure without growth interruption. The diameters and average height of the uncapped QDs was measured with atomic force microscopy (AFM) and amounted to (20. . .35) nm and 4 nm, respectively. HRTEM analysis of QDs capped with GaN showed pyramidal QDs of usually 3 nm height and 10 nm diameter. While the authors [154]
argue that ripening of the uncapped QDs may be responsible for the larger QD sizes of the uncapped InGaN surface, it seems equally likely that the InGaN QDs start to dissolve during the growth of the GaN cap layer.
No satisfactory argument is given in the cited publications concerning the changes in growth mode induced by the growth interruption. Only Hung et al. [157] state that their growth interruption should lead to SK growth rather than phase separation, arguing that their supply of NH3 during growth interruption prevented the decomposition of InGaN.
It again strongly depends on the growth conditions, if islands are obtained or a smooth film. E. g. Kim et al. [147] and Kwon et al. [158–161] used a growth interruption technique to grow smooth QWs with high In concentration up to approximate 0.7 by MOVPE.
5.1.5 Droplet epitaxy
Johnson et al. [162] proposed a model that In is cumulated at In droplets, leading to a reduction of the In concentration in the vicinity, i. e. to a lateral inhomogeneity of the In concentration x. For a high amount of supplied In, the In concentration of the growing InGaN film could even be reduced to an equilibrium value with increasing InGaN film thickness. This was demonstrated by B¨ottcher et al. [163] and Kim et al. [164]. In case of GaN growth on InGaN, Shimizu et al. [165] observed that In droplets present on the surface after InGaN growth act as source of In. Thus ternary InGaN is grown instead of GaN at the positions of the In droplets. In comparison to the reduced In concentration in the vicinity of In droplets [162–164], Selke et al. [166] argue that the In droplets may also act as In source for the further growth of InGaN and thus increase the In incorporation in the adjacent material. In any case, the presence of droplets induces an inhomogeneity of In concentration. This was used by Oliver et al. [167] and Rice et al. [168], who grew self assembled InGaN QDs by “droplet epitaxy” using MOVPE.
A thin InGaN film was deposited and subsequently annealed in a N2 atmosphere before growth of a GaN cap layer. Analysis of annealed but uncapped structures with AFM revealed small In droplets with a density of 1010cm−2 and pits in the surface. The authors propose that spinodal decomposition of the InGaN film takes place during the annealing procedure, and that the In droplets are formed due to this decomposition. The In droplets may then give rise to formation of InGaN QDs during growth of the cap layer. The
authors analysed the capped samples with PL and demonstrated the QD nature of their luminescence centres [145, 167, 168]. They showed e. g. that µ-PL of a masked QD sample at 4.2 K exhibited sharp emission lines of FWHM of 0.27 meV, typical for QDs. The PL peak intensity attributed to the QDs decreased with increasing temperature. Nevertheless, the decrease for comparable QWs was much more pronounced.
5.1.6 Fluctuation of QW width
Fluctuations of the width of QWs can also act as localisation centres for charge carriers.
There are some reports, where luminescence centres were observed (e. g. by PL) and QW width fluctuations were supposed or could not be excluded as possible origins [169, 170].
Soltani Vala et al. [171] calculated that already interface fluctuations of 1 ML may be sufficient to give rise to QD like properties. The interface fluctuations cause fluctuations of the local potential comparable to the changes in the potential induced by the presence of QDs. These local potential fluctuations are localisation centres for charge carriers. Usually the localisation energy is lower for fluctuations of the QW width compared to real QDs.
This implies that the localisations induced by fluctuations of the QW width will not be as effective for increasing temperature, as the charge carriers can be excited into the barrier already at lower temperature.
5.1.7 Phase separation
Considering the phase diagrams of Ho and Stringfellow [15] and Karpov [17] (p. 10) a miscibility gap exists for relaxed InGaN as well as InGaN constrained to the a lattice parameter of GaN. Also other authors calculated phase diagrams and report a miscibility gap for InGaN [172–174]1. Nevertheless, there are many reports of InGaN grown by MBE or MOVPE with In concentrations in the range of the miscibility gap (see e. g. [176] and references therein). For example Adelmann et al. [127] found that InxGa1−xN films can be grown by MBE in the whole concentration range from x = 0 to x = 1 for growth temperatures below 565◦C. Above this temperature, InN decomposes for MBE growth, i. e. in UHV. The growth of InxGa1−xN in the entire composition region between x = 0 and x = 1 is possible as growth by MBE as well as MOVPE does not take place under equilibrium conditions, as was also pointed out by Gerthsen et al. [177]. Whether phase separation occurs or not depends on the actual growth conditions. Several publications of the group of Gerthsen [177–183] concern the analyses of InGaN QWs grown on either GaN or InxGa1−xN of x = 0.02 [177] by MOVPE and/or MBE. They observed fluctuations of the In concentration on two different length scales for both growth techniques, even for average In concentration as low as 0.065 [178]. In rich regions up to approximately
1Teles et al. [175] calculated the phase diagram of sphalerite InGaN for different strain states.
pure InN of diameter less than 5 nm are observed as well as weaker fluctuations on a length scale of some 10 nm to 100 nm. Even though this is in qualitative agreement with phase separation [177, 178], two further explanations are proposed. These are random alloy fluctuations and the formation of nanometer sized pits. Surface pits with diameters of (1. . .2) nm were observed by Chen et al. [184,185]. The extensions of these surface pits, which supposedly are built to relieve strain, did not increase during growth. In and in close vicinity of the pits the In concentration was found to be increased. According to Gerthsen et al. [178] these nanometer sized pits could be filled with In due to surface diffusion, which would lead to the observed In rich regions. These nanometer sized pits should not be intermixed with V-shaped pits, which are often observed in samples containing InGaN layers [186–189]. The V-shaped pits have {1101} facets, are connected to dislocations or IDBs, and usually start at the InGaN area. Based on DFT calculations of In on GaN {1101} facets, Northrup et al. [190, 191] argued that the formation of V-shaped pits is enhanced by strong segregation of In on{1101} facets.
Rao et al. [192] reported phase separation in MOVPE grown (220. . .660) nm thick InGaN films only if the average concentration exceeds 0.12. As this is not in the miscibility gap of the strained InGaN as calculated by Karpov [17], they argued that the phase diagram by Ho and Stringfellow [15] has to be applied, as the thick InGaN films are supposed to relax within the the first few monolayers of the deposited InGaN.
Nistor et al. [193] observed InGaN QDs with sizes of few nanometers and density of about 1010 cm−2, formed by phase separation in an approximately 300 nm thick InGaN layer of average In concentration of 0.22.
The good performance of commercially available InGaN QWs in diodes is also at-tributed to strong localisation of excitons in QD like structures induced by phase separa-tion [169, 194–196].
Another possibility to change the microstructure of an InGaN film is given by applying an annealing procedure [197–207]. The effects of thermal annealing were found to depend e. g. on the width of InGaN QWs [203] and on the actual temperature used [201]. If phase separation is not present in the as-grown InGaN layer, it can be induced or enhanced by a post growth temperature treatment. For example Gerthsen et al. [181] annealed an InGaN QW sample for 1 h at 980◦C under nitrogen atmosphere. They observed a slight decrease of the average In concentration and increase of the average QW width.
The pronounced In fluctuations on the short length scale were not affected by the heat treatment. In addition, larger clusters of high In concentration were observed, which often contained stacking faults. The authors stated that by choosing appropriate parameters for the annealing procedure, the size of the In fluctuations can be adjusted, which would offer the possibility to tailor QDs.
On the other hand, Chuo et al. [199,200] observed a microstructure similar to embedded QDs in as-grown InGaN multiple QWs. These gradually diminished upon annealing for
10 minutes under nitrogen atmosphere when the temperature was increased from 800◦C to 900◦C. They attribute this to an enhanced interdiffusion which results in an improved structural quality.
Additional to the phase separation, the existence of ordered InxGa1−xN, mostlyx= 0.5, was discussed by Shimotomai and Yoshikawa [176] and also by Northrup et al. [190] and experimentally observed by Doppalapudi et al. [208] and Ruterana et al. [209, 210].
5.1.8 Segregation
Segregation can be described by the phenomenological model of Muraki et al. [211]. For this model it is assumed that for a monolayer of deposited material, a fraction R of in the present case In atoms segregates to the surface, forming the “floating layer”, and a fraction 1−R of In atoms is incorporated in the growing monolayer. The concentration x(n) of monolayer n of e. g. an InxGa1−xN film embedded in GaN is then given by
x(n) =
0 for n < N0
x0(1−Rn−N0) for N0 ≤ n ≤ N1
x0(1−RN1−N0)Rn−N1 for n > N1
. (5.1)
x0 is the asymptotic In concentration, which can be obtained for thick films. R is called segregation efficiency or segregation probability and is found to increase with increasing temperature e. g. for InGaAs embedded in GaAs [212, 213]. The nominal thickness of the Inx0Ga1−x0N film is N =N1−N0, where N0 represents the onset of nominal Inx0Ga1−x0N and N1 is correlated with the cessation of the In supply. The In atoms incorporated for n > N1 stem from the floating layer.
For growth of InGaAs on GaAs, a close connection between segregation and island formation was found. InxGa1−xAs island formation on GaAs is observed as soon as the amount of In in the floating layer exceeds a critical value. This enables to predict the WL layer thickness for arbitrary growth conditions if the segregation efficiency is known.
Nevertheless, the exact correlation between island formation and segregation is not fully understood until now. Still, a similar connection between both effects may also be present for the InGaN system. Therefore, some aspects of In surface segregation in InGaN are addressed in the following.
O’Steen et al. [214] analysed the incorporation of In for MBE growth in dependence of growth temperature and V/III flux ratio. As explanation for the observed decrease of In concentration x in the InxGa1−xN films with either increase of growth temperature or decrease of V/III flux ratio, surface segregation and desorption of In was proposed. Later, Sharma et al. [215] analysed the influence of a growth interruption of 0 s to 60 s before GaN cap layer deposition on InGaN QWs. They observed QDs inside the QWs in case of
0 s interruption, whereas for all other times homogeneous InGaN wells were obtained. The QDs are supposed to be related to the In floating layer, with the possibility of In enrichment in the nanometer sized pits reported by Chen et al. [184,185]. The homogeneous QWs were attributed to fast desorption of the In floating layer during the growth interruption.
Chen et al. [185, 216–218] observed strong In segregation for MBE growth of InGaN on (0001) and (0001) surfaces for metal rich growth conditions. The segregation of In is mainly attributed to weaker In−N bonds in comparison to Ga−N bonds. On top of the GaN, a metal bilayer is observed in case of (0001) surfaces, which is in agreement with DFT calculations of Northrup et al. [7]. In addition, the surface exhibits nanometer sized pits (see section 5.1.7). The pits are believed to relieve strain induced by In incorporation in the lower metal layer of the bilayer.
Waltereit et al. [219,220] and Brandt et al. [221] report strong In segregation for growth by MBE in case of metal rich growth conditions. Additionally, In segregation is found to be reduced or suppressed in case of N rich conditions. Potin et al. [183] and Gerthsen et al. [180, 181] report In surface segregation for MBE and MOVPE growth. The segregation efficiency was found to be reduced for MOVPE growth at 720◦C (R = 0.69) in comparison to MBE growth at 530◦C (R= 0.85±0.05), even though the growth temperature is higher.
Both findings are in qualitative agreement with the reports of Schowalter et al. [212, 213], who analysed In segregation for InGaAs capped by GaAs. They observed a reduction of the segregation efficiency for increasing IIIV ratio of MBE grown heterostructures at constant growth temperature. In addition, they report lower segregation efficiencies for MOVPE in comparison for MBE grown specimens.
Moon et al. [222] analysed InGaN QD like regions, which were positioned inside thicker In0.12Ga0.88N layers of 30 nm and 100 nm thickness, grown by MOVPE. For increasing InGaN film thickness, the surface was found to roughen. The formation of the QD like regions is attributed to In segregation to the surface and preferential bonding at ridges, which are embedded in the film during further growth. For growth of multiple quantum well (MQW) InGaN structures, Moon et al. [223] observed two different kinds of In rich regions. One is identified as self assembled QD like luminescence centres, the other is positioned at the upper interface of the QWs and is attributed to segregated In atoms.
For cross section analysis of QWs in TEM, an asymmetric In concentration profile as described by equation (5.1) could also be observed if the electron beam propagates through interface roughnesses occurring during InGaN QW growth, as proposed by Smeeton et al. [224]. If the electron beam “averages” over a sample thickness which is too thick in comparison to the interface roughnesses, a distinction between interface roughnesses and In segregation may not be possible anymore. To eliminate this possibility, investigations should be performed at adequate, thin TEM specimen positions only.
5.1.9 Other effects concerning island formation
From inspection of literature one can find the common trend that QDs are grown at lower temperatures than QWs. There are two main reasons for the reduced temperature (see e. g. [147]). First, the dissociation and reevaporation of In is decreased for lower temperature and if desired, more In can be incorporated in the growing material, which can lead to SK growth. Second, the surface diffusion length is reduced, thus three dimensional growth is facilitated.
Another point to consider to evoke three dimensional growth is the surface of the growing film, as was already briefly described above. Growth under Ga rich conditions is known to commonly lead to smooth films. Smith et al. [225] report metal like (0001) and (0001) surfaces, and report the presence of a Ga bilayer on the surface2. This bilayer could act as surfactant [148] and suppress the island formation. For growth of InGaN on (0001), this bilayer consists mainly of In atoms [185, 216–218].
The effect of surface diffusion of Ga and N on Ga terminated (0001) and (0001) surfaces and Ga on N terminated (0001) and (0001) surfaces was analysed by Zywietz et al. [226].
They obtained energy barriers for surface diffusion E employing DFT. For equilibrium surfaces, which are Ga terminated, the energy barrier for Ga is much lower than for N ((0001): E(Ga) = 0.4 eV, E(N) = 1.4 eV, (0001): E(Ga) = 0.2 eV,E(N) = 0.9 eV). For N terminated surfaces, as can be present in case of N rich grow conditions, the energy barrier for surface diffusion of Ga is much larger ((0001): E(Ga) = 1.8 eV, (0001): E(Ga) = 1.0 eV). This indicates that especially for N rich growth conditions, three dimensional growth can occur, as the diffusion length is reduced. For a smooth surface, the authors [226]
thus recommend slightly Ga-rich growth3. Nevertheless, the Ga bilayer present at the surface was not taken into account for the calculations. This was done applying DFT by Dai et al. [227], who also calculated the energetically favourable positions of Ga atoms in the bilayer. The first adlayer of Ga was found directly on top of the Ga atoms which terminate the (0001) GaN surface. For the second adlayer, the Ga atoms can be adsorbed at two different positions, which are energetically almost equivalent. Analysing the diffusion barrier for a N atom on this surface, they observed a pronounced decrease of energy if the Ga adlayer is allowed to relax in the presence of N. On basis of this finding the authors state that Ga atoms in the adlayer are not bonded tightly, but form metallic Ga−Ga bonds. The obtained energy barrier for surface diffusion for N on this Ga bilayer is E(N) = 0.65 eV, which is much lower than the value reported by Zywietz et al. [226]. Still, the general conclusions of Zywietz et al. [226] and Dai et al. [227] are the same, and Dai et al. [227]
also recommend Ga rich growth conditions for smooth surfaces, as the energy barrier for surface diffusion of N is much lower on the Ga bilayer than on an ideal Ga terminated GaN
2For comparison it shall be mentioned that according to Zywietz et al. [226] the metal like surface for III-V semiconductors is usually not observed; for example GaAs does not exhibit a metal like surface.
3If the growth is very Ga rich, Ga droplets can form on the surface, leading to a perturbed surface.
surface. Even though no comparable results are reported in literature for N diffusion on InGaN, a similar explanation seems likely for InGaN surfaces. Thus, if three dimensional growth of InGaN is desired, N rich growth conditions are usually chosen. Nevertheless, these may induce a higher stacking fault density for the following reason. In bulk material, N is positioned on top of Ga (see figure 2.1, p. 6). For an isolated N atom on (0001) surfaces, this position is energetically unfavourable in comparison to the fcc location, which is above the centroid of the triangle made up by three adjacent Ga atoms in (0001). For a change from the fcc position to the bulk position an energy barrier is present. If the energy barrier is not overcome, stacking faults may be generated. Taking the Ga bilayer into account, the epitaxial position is energetically favourable for the N atom [227]. But also in this case stacking faults are expected to be generated easily as the fcc position requires only an increase of energy of 0.14 eV.
Another aspect related to island formation was published by Tersoff [228], who con-sidered the free energy of island formation and reported that strained layers should be metastable against island and pit formation [229]. According to Tersoff [228] it is expected that the component producing larger misfit segregates to the islands, which is In in case of InGaN. This segregation is stated to evoke first a lower nucleation barrier for islands and second that islands nucleate and grow at enriched In concentration x. For large misfits or low growth temperatures, islands are supposed to form or directly nucleate with misfit dislocations, presumably leading to a strong enrichment of the segregating atom species in the island.