GaAs(0 0 1) heterostructures studied by analytical transmission electron microscopy

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Journal of Physics and Chemistry of Solids 69 (2008) 335–342 www.elsevier.com/locate/jpcs

Composition fluctuations and clustering in (Ga,In)(N,As)/GaAs(0 0 1) heterostructures studied by analytical transmission electron microscopy X. Kong1, A. Trampert, K.H. Ploog Paul-Drude-Institut fu¨r Festko¨rperelektronik, Hausvogteiplatz 5–7, D-10117 Berlin, Germany

Abstract In this article, we summarize our studies on the spatial element distribution in (Ga,In)(N,As) quantum wells and epilayers grown on GaAs(0 0 1) substrates by molecular beam epitaxy. Nanometer-sized composition fluctuations are detected in (Ga,In)(N,As) layers with In and N concentration above 20% and 2%, respectively, by dark-field transmission electron microscopy and spatially resolved electron energy-loss spectroscopy. The fluctuations and clustering are inherently present in these quaternary alloys due to the phase separation tendency. Morphological instabilities, such as the surface roughening due to the elastic strain relief (i.e., a 2D- to 3D-growth mode transition), are succeeding processes. The origin of the fluctuations is discussed with respect to the selected growth conditions and the post-growth annealing procedure. r 2007 Elsevier Ltd. All rights reserved. PACS: 68.65.Fg; 68.55.Nq; 79.20Uv; 68.37.Lp Keywords: A. Semiconductors; A. Quantum wells; C. Electron microscopy; C. Electron energy-loss spectroscopy

1. Introduction (Ga,In)(N,As) dilute nitride heterostructures are considered as attractive systems to realize GaAs-based laser diodes operating in the 1.3–1.55 mm optical fiber window [1,2]. In these ‘‘dilute’’ nitride alloys, a small N fraction of only a few percent, which is added to the conventional III–V semiconductor (Ga,In)As, induces a strong band-gap reduction and thus paves the way for the fabrication of GaAs-based long-wavelength vertical cavity surface emitting lasers [3]. However, when incorporating more than about 2% N and 20% In as required for reaching the desired wavelength range, the structural quality of the (Ga,In)(N,As) quantum wells (QWs) and epilayers significantly deteriorates due to the large miscibility gap and the phase separation tendency of the Ga–N–As system [4,5]. Consequently, large composition fluctuations and strong interface roughnesses are observed, which finally lead to a low photoluminescence (PL) efficiency [6–8]. Corresponding author. Tel.: +49 30 20377 280; fax: +49 30 20377 515. 1

E-mail address: [email protected] (A. Trampert). On leave from University of Glasgow, UK.

0022-3697/$ - see front matter r 2007 Elsevier Ltd. All rights reserved. doi:10.1016/j.jpcs.2007.11.025

A detailed understanding of the origin of this intrinsic composition fluctuation is required in order to further clarify their impact on the structural and optical properties. In this paper, we summarize our results about inherent composition fluctuations in (Ga,In)(N,As) QWs and epilayers for the composition range of relevance, i.e., for [In] 420% and [N] 42%. Analytical transmission electron microscopy (TEM) including dark-field imaging and electron energy-loss spectroscopy (EELS) is applied to detect the nanometer-sized fluctuations of In and N, and to determine their origin during molecular beam epitaxy (MBE) growth. The contribution of local strain energy versus cohesive bond energy to the composition fluctuations is discussed with respect to the growth and postgrowth annealing processes, respectively. 2. Experimental details 2.1. Samples grown by MBE The (Ga,In)(N,As) samples were grown on GaAs (0 0 1) substrates in an MBE system equipped with conventional

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sources for the group III elements as well as the As element, and with an rf-plasma source for N. First, a GaAs buffer layer was grown at about 580 1C. The growth temperature was then decreased to 400–430 1C for growing the Ga1xInxNyAs1y QWs of about 10 nm thickness followed by a GaAs layer of about 3 nm. The remaining GaAs top barrier layer (65 nm) was grown at 580 1C again. More details about the growth can be found in Ref. [9]. In the case of the complete laser structure, the QWs and GaAs barriers were deposited under similar conditions, and the (Al,Ga)As cladding layers were grown at 700 1C including an in situ annealing process [10]. Besides the QW structures, thick (Ga,In)(N,As) epilayers were grown under conditions required to maintain the 2D growth mode. The thickness of the films varies from 0.7 to 2.1 mm. During the epilayer growth, the surface retains the 2D growth mode thus providing a planar growth front, as monitored by the in situ reflection high-energy electron diffraction (RHEED) pattern. The samples studied in this paper were additionally characterized by PL spectroscopy and X-ray diffraction. 2.2. TEM and EELS analysis The cross-sectional specimens used for the dark-field TEM and EELS analysis were prepared by the standard method of mechanical grinding and dimpling down to below 25 mm. The specimens were then thinned by an Ar ion beam with an energy of 3 keV under an incident angle of 31 at room temperature in a Gatan precision ion polishing system. The TEM investigation was carried out in a JEOL 3010 microscope with LaB6 cathode operating at 300 kV. All the dark-field images were directly recorded with a charge coupled device (CCD) camera with 1024  1024 pixels mounted on the TEM. Moreover, this microscope is equipped with a post-column Gatan Enfina parallel electron energy-loss spectrometer system. In the low-loss EEL spectra, the energy resolution is about 1.5 eV according to the full width at half maximum (FWHM) of the zero-loss peak. For the EELS measurements, the QWs were oriented parallel to the incident electron beam (‘‘endon’’) so that the spectra taken from the QWs were not affected by overlapping GaAs regions. In order to remove the effect of multiple inelastic scattering and the multiple plasmon excitations, the standard Fourier-log deconvolution technique was applied to obtain a single scattering distribution (SSD) [11]. 2.2.1. Dark-field TEM The ability of dark-field images to reveal interfaces in heterostructures and layers of different composition is well established [12]. The (0 0 2) and the (2 2 0) Bragg reflections are used to analyze the composition and the strain variations in III–V compound semiconductors, respectively. The dark-field micrographs taken under two-beam condition with the (0 0 2) reflection are sensitive to the chemical composition of crystals with zinc-blende structure

(structure factor imaging). The high contrast of such images appears because the structure amplitude F0 0 2 is equal to the difference between the atomic scattering factors fi for the group III and group V atoms. For Ga1xInxNyAs1y, F0 0 2is given by F 0 0 2 ¼ 4½xðf In  f Ga Þ þ yðf As  f N Þ þ ðf Ga  f As Þ.

(1)

In the kinematical approximation, the intensity of the (0 0 2) reflection producing in the dark-field image is then proportional to jF0 0 2j2. In the case of the (2 2 0) reflection, any intensity modulation is induced by a local distortion of the (2 2 0) lattice planes indicating shear strain perpendicular to the growth direction. The shear strain is caused by the elastic relaxation, which is, in the case of a planar growth front, only related to the composition variation. In the case of the TEM thin foil preparation, additionally introduced surface relaxation processes can affect the intrinsic strain distribution in the sample and thus modify the strain contrast in the micrograph. In quaternary alloys, like Ga1xInxNyAs1y, composition variations in (0 0 2) dark-field images can originate from fluctuations of both the group III and group V elements, i.e., x and y. Therefore, the unambiguous determination of the In and N content from one (0 0 2) dark-field image is not possible and thus, complementary techniques with similar spatial resolution and detecting the same sample area are necessary, such as the strain analysis based on high-resolution TEM [13]. Alternatively, the spatially resolved EELS method can be applied to investigate the local element distribution. 2.2.2. Low-loss EELS In order to quantitatively and simultaneously determine the N and In content, the EELS in the low-loss region is utilized. The incident electrons lose part of their energy due to the inelastic scattering, when they pass through the samples. Some of these energy losses provide a way to identify the elements in the sample and can be used to measure the element distribution on a nanometer scale [11]. Therefore, the EELS in the TEM offers the unique combination of a very high spatial resolution and the ability to provide information upon chemical composition and bonding of conventional III–V semiconductors [14–18]. Fig. 1 shows an original EEL spectrum and the corresponding SSD taken from (Ga,In)(N,As) in the lowloss region reflecting two main characteristic contributions: the plasmon excitation at about 16 eV and a broad peak superimposed on the rapidly falling tail of the plasmon excitation, which includes the transitions from the In 4d (above 18 eV) and the Ga 3d (above 20 eV) level to the conduction band [18]. Both features are utilized to quantitatively measure the N and the In content independently and simultaneously from one low-loss spectrum. The In content is analyzed from the In 4d transition intensity, which has been isolated from the influence of the plasmon peak and the overlapping Ga 3d transitions by a

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Single scattering distribution

Plasmon (~16 eV)

In 4d (~18 eV) Ga 3d (~20 eV)

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whereas the third sample C is a complete QW laser structure lasing at 1.4 mm [21]. The three QWs in the reference sample A contain a constant N (2.5%) but increasing In content (23%, 30%, 40%), whereas the QWs in the reference sample B consists of a fixed In (35%) and variable N content (1%, 2%, 3.5%). The weak contrast variation in the chemically sensitive (0 0 2) dark-field image of the reference samples shown in the Fig. 2(a), indicates a relatively small composition fluctuation along the QWs in the lateral direction. Fig. 2(b) shows a bright-field micrograph of the reference sample C with a (Ga,In)(N,As) QW in the center embedded between two GaAs barriers and (Al,Ga)As cladding layers. The nominal concentrations were set to xIn ¼ 35% and yN ¼ 2.5%. Along this QW, no contrast fluctuations are visible in the (0 0 2) darkfield micrograph (not shown) indicating the high structural quality, which is in agreement with the high PL efficiency and the lasing characteristics [22]. Low-loss EEL spectra

Fig. 1. The low-loss EEL spectrum and its corresponding single scattering distribution (inset) taken from (Ga,In)(N,As).

In Eq. (2), n is the number of valence electrons per unit volume and m* is the effective electron mass. From this equation it follows that Ep depends on the density of electrons, i.e., on the lattice parameter, and on the effective electron mass, which is related to the band structure. In the case of dilute nitrides, two behaviors are noticeably different from the conventional semiconductor compounds: (i) The incorporation of a few percent of N in GaAs and (Ga,In)As results in a huge reduction of the band gap (band-gap ‘‘bowing’’) [20]. This is in contrast to the linear and very small variation of the band gap with the composition of conventional III–V ternary alloys, like (Ga,In)As. (ii) The addition of N in GaAs reduces the band gap while simultaneously reducing the lattice constant. As a consequence, quaternary (Ga,In)(N,As) alloys can be grown lattice matched to GaAs. Therefore, we can conclude that in (Ga,In)(N,As) the stain effect due to lattice parameter mismatch should play a minor role in the variation of Ep than the band structure. This conclusion is additionally supported by our results from well-defined reference samples as described in the following. In order to verify the dependency of Ep on the band structure of Ga1xInxNyAs1y, the plasmon energies were experimentally determined for three reference samples: samples A and B contain three QWs with a systematic variations of the In and N concentrations, respectively,

50 nm

GaAs

400 nm

AlGaAs III

GaAs (Ga,In)(N,As)

II

GaAs

(Ga,In)(N,As) AlGaAs

g=(002) I GaAs

GaAs In content (%) 16.2 Plasmon energy Ep (eV)

complex method described in Ref. [19]. The small amount of N incorporating in these dilute nitride alloys can be measured quantitatively by the plasmon energy shift with respect to a reference GaAs signal. Based on the free electron theory, the energy of the plasmon excitation Ep is given by [11] sffiffiffiffiffiffiffiffiffiffi ne2 Ep ¼ _ . (2) 0 m

10

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GaAs

16.1

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laser structure of (b)

16.0

In

15.9 15.8 N 15.7 0

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Fig. 2. (a) Cross-sectional (0 0 2) dark-field TEM image of reference sample A: QWs with [N] ¼ 2.5%, but [In] ¼ 23% (I), [In] ¼ 30% (II) and [In] ¼ 40% (III). (b) The micrograph of the reference sample C of laser structure containing QW with [N] ¼ 2.5% and [In] ¼ 35%. (c) Experimental plasmon energies versus N and In content, obtained from reference sample A, sample B and from the laser structure shown in (b).

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were measured at about 10 positions along every QW of these three reference samples. The corresponding average plasmon energies and their standard variations are summarized in Fig. 2(c). These experimental results demonstrate that, within the composition range of interest, Ep is sensitively dependent on the N content but is almost independent of the In content. This conclusion is in agreement with the theoretical and experimental results [20], which demonstrate a much stronger bowing effect on the band structure for the incorporation of N compared to In atoms. Therefore, the N concentration can be quantitatively determined for our (Ga,In)(N,As) samples according to the shift of Ep relative to its position for the GaAs substrate. 3. Results and discussions 3.1. (Ga,In)(N,As) QWs on GaAs (0 0 1) substrate 3.1.1. The impact of composition fluctuation on the morphological instabilities and the optical properties In the following, we will focus on the (Ga,In)(N,As) QWs with a composition range of [In] 30% and [N] 3%, which generally contain an epitaxial strain of eE1.5%. Due to the large miscibility gap of the alloy, a low substrate temperature during MBE and a careful adjustment of the growth conditions are required to achieve high structural and optical properties of QWs. Fig. 3 presents (0 0 2) dark-field cross-sectional TEM

images of the microstructure along the /1 1 0S-projection and the corresponding PL spectra for a series of three QWs grown under basically identical conditions (regarding growth rate, III/V ratio, layer thickness), but at growth temperatures of 430, 420 and 410 1C, respectively. The In and N mole fractions are about 30% and 3%, respectively. Under the chemically sensitive g0 0 2 imaging condition, the TEM images reveal that an increase of the growth temperature leads to an enhancement of the lateral composition modulation as identified by means of the contrast variation along the QW. The strongest modulation with a period of 10–15 nm was found in the QW grown at 430 1C. Moreover, a stronger interface undulation is also detected with the increase of the growth temperature. At the same time, the PL line spreads out and the intensity drops dramatically with raising composition modulation. The highest PL efficiency (Fig. 3(f)) is only obtained in the case of perfectly 2D-grown QWs possessing the slightest possible composition variation, as demonstrated in Fig. 3(c). The small interface width and the smoothness of the interface in the lateral direction correlate well with the comparably narrow PL line width of 28 meV at room temperature. Further, it is remarkable that in this case the designed wavelength of 1.5 mm is achieved. We therefore conclude that with the increase of growth temperature, the composition fluctuation accumulates and, thus an interface undulation and a 2D-to-3D growth mode transition is observed, which finally plays a dominant role for the optical properties of (Ga,In)(N,As) QWs.

GaAs Tg = 430°C (Ga,In)(N,As) 20 nm 0.8

Tg = 420°C (Ga,In)(N,As) 20 nm

1.3 GaAs Tg = 410°C

1.2

1.4

1.6

FWHM = 31.05meV λ = 1.51 μm

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GaAs

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1.7

FWHM = 28.46meV λ = 1.56 μm

(Ga,In)(N,As) 20 nm 1.3

1.4

1.5

1.6

1.7

Wavelength (μm) Fig. 3. Cross-sectional (0 0 2) dark-field TEM images of (Ga,In)(N,As) quantum wells grown at 430 1C (a), 420 1C (b), and 410 1C (c) and their corresponding PL spectra (d)–(f).

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the QW grown at a lower temperature (410 1C), where uniform contrast and smooth interfaces are observed in the (0 0 2) dark-field image (cf. Fig. 5(a)), indicate a more homogeneous composition with less fluctuations of both elements within the experimental error bars (see Fig. 5(b)). The ideal element distribution in a quaternary alloy, i.e., the nearest-neighbor bond configuration, is determined by minimizing the alloy free energy, which includes local strain and cohesive bond energy terms. The local strain is produced by differences in bond lengths between the corresponding III–V atom combinations [23]. During the epitaxial growth, the physisorbed adatoms are able to relieve local strain toward the free surface. Therefore, the development of nearest neighbor configurations is more driven by maximizing the cohesive bond energy than minimizing the local strain. Taking into account the values for the cohesive energies of the various bond configurations (compare values in Table 1), the formation of Ga–N and In–As bonds should be favored inducing an inherent composition modulation near the growing surface. This inherent composition modulation is associated with a strong increase of the epitaxial strain energy. Areas with preferred Ga–N and In–As bond configurations cause a higher epitaxial strain to the GaAs substrate compared to

GaAs (Ga,In)(N,As)

g=(002)

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Tg=410°C

GaAs

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4

a

GaAs

g=(002)

(Ga,In)(N,As)

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In content (%)

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Fig. 5. Cross-sectional (0 0 2) dark-field TEM image of the as-grown (Ga,In)(N,As) QW grown at 410 1C (a), and the corresponding N and In distribution (b) from a scan of low-loss EEL spectra marked by circles.

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3.1.2. The origin of composition fluctuation In order to understand the influence of composition fluctuations on the structural and finally optical properties, it is required to clarify the origin of this intrinsic composition fluctuation, which is generated during epitaxial growth. For this purpose, the low-loss EELS method is used to determine the element distribution in as-grown (Ga,In)(N,As) QWs grown at 430 and 410 1C, respectively. In order to analyze the element distribution along the QWs, an electron beam with about 8 nm spot size is generated that is smaller than the periodicity of the measured contrast modulation in the (0 0 2) dark-field images. A series of low-loss EEL spectra is taken along the QW as indicated by the row of circles in Fig. 4(a). The associated plasmon energies indicate the inhomogeneous lateral distribution of N atoms. The maximal variation in Ep amounts to 0.4 eV, which corresponds to a wide fluctuation of the N content between 0% and 3.5%. The corresponding In content varies at these positions between 10 and 42%. The results of the spatially resolved EELS measurements of the local In and N distribution along the QW are summarized in Fig. 4(b). The trend of the curves reflecting the local elemental distributions shows a periodic oscillation similar to the contrast modulation in the darkfield TEM image. However, both curves follow an opposite trend, i.e., positions with high N content correspond to low In concentration and vice versa. Therefore, the composition fluctuation, which is detected by contrast modulation in the dark-field image of Fig. 4(a), is not only the result of fluctuations in the N but also in the In composition. This result indicates that there are regions along the QW with preferred formation of Ga–N and In–As bond configurations, respectively, although these configurations lead to a higher local strain. On the other hand, the EELS results for

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Table 1 The cohesive energies and the bond lengths of Ga–N, In–N, Ga–As and In–As bond configurations [24] Bond type

0 Position (arb. units) Fig. 4. Cross-sectional (0 0 2) dark-field TEM image of the as-grown (Ga,In)(N,As) QW grown at 430 1C (a), and the corresponding N and In distribution (b) from a scan of low-loss EEL spectra marked by circles.

Ga–N In–N Ga–As In–As

Cohesive energy (eV/bond)

Bond length (A˚)

1.94 2.15 2.45 2.61

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the case of Ga–As and In–N (see values for bond lengths given in Table 1). Therefore, the elastic strain energy is accumulated during growth until surface roughening starts that finally initiates the 2D-to-3D growth mode transition. Lowering the growth temperature results in a reduction of the adatom mobilities, which is responsible for a lower probability of forming the favorable bonds. The composition modulation is hence suppressed, in agreement with our observation of the QW grown at 410 1C (Figs. 3(c) and 5(a)). In this case, the composition fluctuation is smaller than the detection limit of the dark-field image and EELS method. Such a perfect QW structure is taken as a reference for studying the effects of annealing on the microstructure. For that purpose, a perfect QW structure with nominal 35% In and 3.5% N was heated to 700 1C for 40 min. As expected, the heat treatment generates a strong enhancement of the radiative efficiency with a blue-shift of the PL peak [6,7]. This behavior is displayed in Fig. 6 showing the PL spectra before and after annealing at room temperature. The structural characterization of the asgrown and annealed samples by a combination of (0 0 2) dark-field (inset in Fig. 6) and high-resolution TEM indicates that there is no change in the QW thickness, interface width and average composition revealing that there is no diffusion of In and/or N from the QW into the GaAs barriers. However, a local redistribution of In and N atoms is detected inside the well from the interface to the central region (for details of these measurements, see [6,7]). The driving force for this redistribution is related to the reduction of local strain (cf. bond length in Table 1) and can be interpreted as a change from the Ga4N configuration into a more stable Ga3In–N cluster (in agreement with

300 K as-grown g002

PL Intensity (arb. units)

x 3000 x100 (Ga,In)(N,As)

10 nm

GaAs annealed

g002

(Ga,In)(N,As) GaAs

1.4

1.5

1.6 1.7 1.8 Wavelength (µm)

10 nm

2.0

Fig. 6. PL spectra taken from an as-grown (dotted line) and annealed (solid line) (Ga,In)(N,As) QWs with [In]35% and [N]3.5%. Inset: the cross-sectional (0 0 2) dark-field TEM images of as-grown and annealed (Ga,In)(N,As) QW.

recently published Raman measurements [25]). During the annealing process, the development of the nearest neighbor configuration is thus more driven by the minimization of the local strain than by the maximum of the cohesive bond energy, in contrast to the situation at the growing epitaxial surface. It should be additionally mentioned that the respective diffusion mechanism during ex situ annealing depends on the initial structure of the as-grown sample. Large composition fluctuations (as shown in Fig. 3(a)) and the consequent interface roughness are favorable sources for QW intermixing as observed by Albrecht et al. [26]. These large fluctuations will definitely open a further diffusion path that is activated in addition to the just mentioned local redistribution process that we have detected in our perfect QW structure. 3.2. Composition fluctuation in bulk-like (Ga,In)(N,As) epilayers As mentioned in the previous section, the composition fluctuation is an inherent property of the (Ga,In)(N,As) alloy existing within the considered composition range. This fluctuation appears at the growth front during MBE without the need of surface roughness based on elastic strain relief. The 2D-to-3D growth mode transition is thus a secondary process initiated if sufficient strain is accumulated, that, of course, depends on the actual In and N content as well as the dominant growth parameter. Macroscopic epitaxial or local strain is definitely not the primary cause of the present composition fluctuation but the phase separation tendency due to the miscibility gap, which is realized by spinodal decomposition. In this context, we have also investigated thick (Ga,In) (N,As) epilayers, which were grown under conditions similar to the 2D QWs retaining a planar growth front. The nominal In and N concentrations are 20% and 2%, respectively, resulting in a lattice mismatch of around 1%. Part of this epitaxial strain is relieved by the formation of misfit dislocations without affecting the 2D surface morphology as verified by in situ RHEED and TEM measurements. Fig. 7(a) and (b) displays {1 1 0} crosssectional dark-field TEM images with g ¼ (2 2 0) taken from a 0.7 mm-thick (Ga,In)(N,As) epilayer before and after annealing (700 1C, 40 min), respectively. The specklelike contrast with elongated feature along the growth direction is caused by local distortions of the (2 2 0) lattice planes. Neglecting the influence of surface relaxation due to the thin TEM foil preparation, these distortions arise from composition modulations proceeding shear strain at the (2 2 0) planes. Because the contrast does not vanish in the annealed sample, but is even slightly enhanced (cf. Fig. 7(b)), the corresponding composition fluctuations are based on the spinodal decomposition process, as already discussed in relation to the QW structure. Unfortunately, the composition fluctuations cannot be appropriately detected under the chemical sensitive (0 0 2)

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(Ga,In)(N,As) g=(220) GaAs

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GaAs

200 nm

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(Ga,In)(N,As) g=(002) GaAs

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Fig. 7. Cross-sectional (2 2 0) and (0 0 2) dark-field TEM images of (Ga,In)(N,As) epilayer before (a) and (c) and after annealing (b) and (d).

In

N content (%)

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Fig. 8. The lateral distributions of N and In concentrations in the annealed (Ga,In)(N,As) epilayer.

between 0% and 2.5%. The absence of detectable In variations in the present case is explained by the relative small In concentration of 20% compared to the average 35% measured for the QW shown in Fig. 4. Therefore, from these measurements, two main conclusions can be drawn: First, the N clustering must be considered as the primary driving force for the composition fluctuation and In will support this process when its concentration exceeds the 20% boundary. And second, composition fluctuations can take place even in the absence of elastic surface relaxation, i.e., during 2D growth. 4. Summary

dark-field condition (Fig. 7(c) and (d) for as-grown and annealed sample). The calculation of the (0 0 2) reflection intensity in the kinematical approximation shows a parabolic behavior in dependence of the In and N concentration with a local minimum in the range between 20% and 25% In and 0% and 4% N. [13,27] This is just the composition range of the present (Ga,In)(N,As) alloy, and variations in the composition can thus not reliably be detected. In order to clarify the character and the amount of the composition fluctuation existing in the annealed epilayer, a lateral scan of low-loss EEL spectra was performed in the same way as we did for QWs. The results are summarized in Fig. 8 demonstrating that the In concentration shows a very week variation around 20%, whereas the N concentration reflects a large fluctuation

We have studied the microstructure and the local chemical composition of epitaxially grown (Ga,In)(N,As)/ GaAs(0 0 1) heterostructures containing In and N concentrations above 20% and 2%, respectively, with dark-field TEM and spatially resolved EELS. The main results are summarized as follows: (1) Nanometer-sized composition fluctuation, which are generally detected, are driven by maximizing the cohesive bond energy forming areas with preferred Ga–N and In–As bond configurations, respectively, in spite of the increasing local strain. Lowing the growth temperature reduces the level of fluctuation and results in a more homogeneous QW structure with smooth 2D interfaces.

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(2) Composition fluctuations are also present in the case of layers grown in the 2D growth mode. Surface roughening and the 2D-to-3D growth mode transition is initiated if sufficient strain is accumulated, which is generated by the composition fluctuation. (3) The effect of ex situ annealing on the chemical composition and interface character depends on the microstructure of the as-grown QWs and epilayers. In the case of structurally perfect QWs, the annealing step gives rise to a local N and In rearrangement within the QWs resulting in a more homogeneous element distribution by maintaining the abrupt interfaces. Strong composition fluctuations and interface roughness in the as-grown heterostructure lead to interdiffusion during annealing and thus to a broadening of the interface. The spinodal decomposition process is additionally initiated.

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