Cathodoluminescence spectral mapping of III-nitride structures

Transcription

1 phys. stat. sol. (a) 201, No. 4, (2004) / DOI /pssa Cathodoluminescence spectral mapping of III-nitride structures R. W. Martin *, 1, P. R. Edwards 1, K. P. O Donnell 1, M. D. Dawson 2, C.-W. Jeon 2, C. Liu 2, G. R. Rice 2, and I. M. Watson 2 1 Physics Department, Strathclyde University, 107 Rottenrow, Glasgow, G4 0NG, UK 2 Institute of Photonics, Strathclyde University, 106 Rottenrow, Glasgow, G4 0NW, UK Received 2 October 2003, revised 26 February 2004, accepted 27 February 2004 Published online 10 March 2004 PACS Hk, Fd, De, Hc The application of cathodoluminescence spectral mapping to the characterisation of a range III-nitride semiconductor structures is described. Details are presented of the instrumentation developed to carry out such measurements using an electron probe micro-analyser. The spatial resolution of the luminescence data is ~100 nm. The technique is enhanced by the ability to simultaneously perform X-ray microanalysis and electron imaging. Results are presented from epitaxially laterally overgrown GaN and InGaN/GaN structures using both single-layer SiO 2 and multilayer SiO 2 /ZrO 2 masks. Effects of strain and microcavity formation are resolved. Application of the technique to InGaN epilayers shows spatially-dependent shifts in the peak wavelength of the luminescence spectrum which correlate directly with microscopic variations in the indium content. Regions emitting at lower energy and with decreased intensity are shown to have higher InN contents, mirroring equivalent macroscopic observations. Finally the spectral mapping technique is used to analyse the luminescence from micron-scale selectively grown III-N pyramids, indicating possible formation of quantum dots at the sharp tips. 1 Introduction Characterisation tools that enable the measurement of light emission from materials with high resolution in both spatial and spectral domains are highly desirable. One method for achieving this is the use of cathodoluminescence (CL) spectra collected using a scanning electron microscope (SEM). The spectral resolution is provided by using one of several light-collection arrangements to guide light, emitted in response to electron beam excitation, from the microscope to a spectrograph or similar device. The interaction volume of the electrons with the material determines the spatial resolution. By using low electron beam energies (~ 2 kev) and field emission sources a spatial resolution of 20 nm has been achieved, demonstrated by imaging GaAs/AlGaAs quantum wells in cross-section [1]. Using higher electron beam energies and standard tungsten sources the spatial resolution is larger, but 100 nm is still feasible with energies of approximately 5 kev. If the light emission is restricted to a thin layer within the material an improved resolution may be possible, depending on the excitation pathway. The power of CL spectroscopy can be further enhanced by CL spectral mapping in which use of the beam-scanning facility within an SEM, or the stage scanning within an electron probe micro-analyser (EPMA), allows a CL spectrum to be collected at each point within a 2D map. The resulting three-dimensional CL dataset, or hyperspectral image, can be analysed using a variety of visualisation or numerical tools to extract dependencies or variations not discernable using individual spectra or panchromatic imaging. III-Nitride semiconductor materials are now well established as highly efficient and commercially important light emitters, lying at the heart of violet/blue laser diodes and UV/blue/green/white LEDs. Investigations of the cathodoluminescence from such materials have provided important information on * Corresponding author: r.w.martin@strath.ac.uk, Phone: , Fax:

2 666 R. W. Martin et al.: Cathodoluminescence spectral mapping of III-nitride structures the material properties [2 5]. The machines used to collect CL data are generally fitted with a variety of detectors for complementary analysis techniques, such as electron imaging and X-ray microanalysis. Acquiring such information at the same time, and from the same region, as CL further enhances the usefulness of the technique. For example, a previous study by the authors [3] employed combined CL and wavelength-dispersive X-ray (WDX) measurements with a static electron beam to investigate InGaN alloys. This allowed the dependence of the luminescence peak energy on the InN fraction to be determined with greater accuracy than previously possible. This work was extended by using a scanning beam to simultaneously map both the complete emission spectrum and the elemental composition of InGaN layers with sub-µm spatial resolution [4]. 2 Experimental details 2.1 Instrumentation The measurements were carried out using a Cameca SX100 electron probe micro-analyser (EPMA). The EPMA incorporates an automated stage with 100 nm step-size to allow images to be acquired by scanning either the sample or the beam. A built-in optical microscope, coaxial and confocal with the electron beam, allows optical monitoring of the region excited by the electron-beam. We have modified the EP- MA by the addition of an optical spectrometer, equipped with a cooled silicon CCD detector array, into the light path of the optical microscope [3-5]. This enables the fast acquisition of a room temperature CL spectrum at each point in the raster scan. The resultant CL spectral map is then treated computationally to extract two-dimensional images representing various aspects of the CL data, such as spectrallyintegrated intensity, peak wavelength, peak width or chromaticity. The EPMA is designed to allow quantitative elemental analysis and composition mapping. In our case this is provided by three WDX spectrometers, whose high spectral resolution and peak-background ratios result in excellent composition detection limits (<0.05 atomic % demonstrated for rare-earth ions in GaN [5].) This facility allows the CL hyperspectral image to be compared directly with WDX element al composition maps. The sub-µm spatial resolution of the WDX technique is well-matched to that of CL measurements. The CL acquisition times are determined by the brightness of the sample (monitored at room temperature) and the desired total measurement time. Good quality CL spectral maps can be acquired from high brightness In- GaN using acquisition times as low as 25 ms, whilst the stability of the EPMA also allows times of several seconds to be used for emitters of lower intensity. 2.2 Beam conditions and interaction volume The EPMA is currently configured with a tungsten electron gun. In addition to considering the spot-size produced on the sample using this source the interaction volume of the electrons within sample must be taken into account when determining the beam conditions. The interaction volume for a particular material is estimated using Monte Carlo electron trajectory simulation [6]. Fig. 1 shows the calculated energy deposition of a 6 kev electron beam for a 250 nm layer of InGaN on GaN. This plot indicates that the sampled region is ~ 150 nm in diameter and ~ µm 3 in volume. Such a beam energy is generally sufficient to produce CL spectra and WDX data of suitable quality and, in the case shown, confines the excitation to the top layer. In the case of quantum well samples the excited values covers a range of layers within the sample and slightly different factors determine the optimum beam energy, but values of order 5 kev are often suitable. The electron current is selected to produce suitable CL intensity and/or WDX count rates. It is typically na when compostion data is required but can be sub-na when only measuring CL. A beam regulator actively controls the current, allowing maps to be acquired over long periods. 2.3 Structures studied The results presented are all from structures grown in an Aixtron 200 series metalorganic vapour phase epitaxy (MOVPE) reactor at the University of Strathclyde. Specific sample details are provided below.

3 phys. stat. sol. (a) 201, No. 4 (2004) / Depth (nm) % 10% 5% InGaN GaN Lateral distance (nm) Fig. 1 Monte-Carlo simulation (10 6 electrons) of the energy deposition of a 10 nm wide, 6 kev electron beam within a 250 nm In 0.2 Ga 0.8 N on GaN layer. The contours indicate the normalised rate of energy deposition. 3 Results 3.1 Laterally overgrown GaN The use of lateral epitaxial overgrowth of III-nitrides to generate huge reductions in the defect density has been well documented. Such material has been characterised by a wide range of techniques, including CL mapping [7, 8]. Here we present the use of plan-view CL spectral mapping to map the seed regions between the mask stripes and the laterally grown wing regions. Conventional lateral overgrowth employs single-layer striped SiO 2 masks, with a thickness of ~100 nm and wing and seed widths of ~ 5-10 µm. We have used such masks to fabricate and characterise laterally overgrown GaN and In- GaN/GaN structures [9]. The cross-sectional image of a structure with ~4 µm of overgrown GaN is shown in Fig. 2. The CL spectral mapping technique has been used to analyse the grown surface of such material. Fig. 3 shows a 67 µm wide region, covering several stripes of material, imaged using a 5 kev electron beam. The upper panel shows the integrated CL intensity. The CL clearly shows the expected improvements associated with lateral overgrowth. The material grown above the mask shows brighter luminescence whilst the seed regions emit weaker CL and have a mottled appearance, related to the larger dislo- Void-free overgrown GaN SiO 2 mask stripe 5 microns Fig. 2 Cross-sectional secondary electron image of laterally overgrown GaN above a single-layer silica mask.

4 668 R. W. Martin et al.: Cathodoluminescence spectral mapping of III-nitride structures counts counts nm nm Fig. 3 (online colour at: CL spectral mapping of a 67 µm wide region of laterally overgrown GaN in plan-view. The upper panel shows the integrated luminescence intensity and the lower panel shows the weighted mean CL wavelength. cation density in these areas. A dark band is visible down the centre of the overgrown regions, associated with reduced luminescence at the coalescence of the two laterally growing faces. The lower plot shows the weighted mean wavelength plotted over the same area. The scale range is less than 1 nm but the image shows these very small shifts in CL peak to be clearly related to the overgrowth. The shifts are due to changes in strain within the GaN. The red shift observed in the wing regions is caused by relaxation of the compressive strain conventionally found in GaN-on-sapphire layers as a result of differential contraction during cool-down from the growth temperature. The compressive strain increases again at the coalescence boundary and this region shows shorter wavelength CL than the rest of overgrown area. 3.2 Lateral overgrowth above dielectric mirror elements In previous reports we have discussed the advantages of using both upper and lower all-oxide DBR mirrors for the fabrication of III-N microcavities and the possibility of achieving this using various forms of lateral epitaxial overgrowth [10, 11]. One approach is to replace the single-layer SiO 2 mask used in conventional epitaxial lateral overgrowth with a multi-layer mirror stack. A 10½-period SiO 2 /ZrO 2 DBR mirror has been formed into a series of ~9 µm wide stripes on a trench patterned 2-inch diameter GaNon-sapphire pre-layer. Lateral overgrowth of GaN was then performed to bury this DBR, with a series of InGaN/GaN quantum wells included near the top of ~6 µm of overgrowth. A second SiO 2 /ZrO 2 DBR mirror was deposited on the planarised top surface. The structure was then studied using room temperature CL spectral mapping. Fig.4 shows data from a µm region imaged with a 20kV electron beam. The CL intensity map shows the bright regions above the dielectric mask. In this case the luminescence is enhanced both by the improved material quality due to lateral overgrowth and by the provision of a highly reflecting DBR mirror buried within the structure. Individual spectra are shown in the right panel of Fig. 4 and clearly show the effect of the buried mirror. The InGaN emission in the spectra of the wing region is strongly modulated by fringes related to the cavity formed by the two DBR mirrors. In the seed regions the fringes are much weaker due to the much lower reflectance at the GaN-sapphire interface. Analysis of the fringe spacing shows the cavity length to be increased for the seed regions compared to the wings, for whom the DBR (positioned higher than the sapphire) forms the lower reflecting element. Although all-oxide DBR mirrors have been successfully buried by high-quality overgrown III-N material the processing steps associated with mask preparation are highly complex and in need of simplification.

5 phys. stat. sol. (a) 201, No. 4 (2004) / CL counts (a.u.) Wing Seed Wavelength (nm) Fig. 4 (Left) CL intensity map from a 50 x 50 µm region of InGaN/GaN quantum wells laterally overgrown above a SiO 2 /ZrO 2 DBR mirror mask. (Right) Individual spectra extracted from the CL spectral image showing typical spectra in the wing and seed regions. 3.3 InGaN epilayers The CL spectral mapping technique has previously been used to investigate spatial variations in luminescence from InGaN layers (e.g. [12]. Here this is enhanced by the use of simultaneous data on elemental composition to examine the relationship between optical emission and InN fraction in a range of InGaN epitaxial layers [3, 4]. Results from one structure are presented in this paper. The sample consists of a ~1 µm GaN buffer on a sapphire substrate, followed by ~180 nm of InGaN. Rutherford backscattering spectrometry and WDX measurements [3] showed the area-average indium cation fraction to be 12%. The data was collected using an electron beam energy of 7 kev, calculated to be optimal for this thickness of InGaN, and a beam current of 40 na to provide a sufficient WDX count rate for indium. Fig. 5 shows data from a 20 x 20 µm region obtained by scanning the sample stage. The elemental maps of indium and gallium show many similar contrast features which, by comparison with the backscattered electron image, can be identified as resulting from the surface roughness. This false contrast has been minimised by dividing the two WDX images to yield a map (Fig. 5a) showing variations in the cation ratio which has been quantified using analysis of the relative peak and background WDX counts for sample and standards. The InN fraction ranges between 0.11 and Figs. 5(b-c) show two 2-D aspects A B x = 0.13 A B A B 428 nm x = nm (a) In:Ga ratio (b) CL counts c) Peak CL wavelength All to same scale: 5 µm Fig. 5 (online colour at: Composition and room temperature cathodoluminescence images of an In x Ga 1 x N epilayer, acquired simultaneously using a step stage scan. Examples of indium-rich (A) and indium-poor (B) regions are highlighted.

6 670 R. W. Martin et al.: Cathodoluminescence spectral mapping of III-nitride structures Normalised CL intensity K E 40 mev Indium rich region (A) Indium poor region (B) Wavelength (nm) Fig. 6 Room temperature spectra extracted from the CL dataset, corresponding to the two labelled points on Fig. 3. of the 3-D CL dataset; respectively the integrated intensity and peak wavelength of the main, higher energy, luminescence band. Whilst the total CL emission intensity will be modified by the presence of surface roughness, correlation is still observable between this intensity and the indium content. The emission wavelength, less sensitive to surface morphology, is seen to exhibit a strong dependence on the composition. These images show that less intense and longer wavelength light is emitted from regions with locally higher indium content, in clear agreement with macroscopic measurements. Fig. 5 also shows the positions of the two points (labelled A and B) representing indium-rich and indium-poor areas respectively. The spectra from these points are shown in Fig. 6. The peak wavelengths of the two spectra are separated by 5.5 nm, corresponding to an energy difference at this wavelength of 40 mev. Using the linear relationship between peak emission energy and composition [3], such a change in energy corresponds to a difference in InN fraction of ~0.01. This is consistent with the difference in x seen between the two points in Fig. 5a. 3.4 Selectively grown III-N structures As described above the lateral overgrowth of GaN-based materials using patterned silica masks has enabled significant advances in material quality, through dramatic reductions in defect density. Termination of overgrowth above an aperture patterned mask prior to the coalescence of features formed in the earliest growth stages allows the design of novel structures, including arrays of hexagonal pyramids or prisms. This is selective epitaxy, upon which lateral overgrowth depends. Incorporation of InGaN quantum wells in the final stages of pyramid growth can lead to the formation of quantum dot arrays [13]. Selective growth also has important applications in the fabrication of field-emitter arrays and III-N microcavity devices [14, 15]. Microring and microdisk devices have previously been fabricated by lithography and etching [16] but the selective growth approach potentially produces more efficient cavities due to smoother facet walls. The CL system described above was used to produce spectral maps from arrays of selectively grown InGaN/GaN quantum wells. A similar CL study of GaN pyramids, without the quantum wells, has previously been reported [17]. Our structures were grown on GaN-on-sapphire seed layers covered with 100 nm thick silica mask layers, which were patterned into arrays of holes by lithography and wet etching. Overgrowth of GaN results in arrays of sharp-tipped pyramids as shown by the secondary electron images in Fig. 7. The six facets of each pyramid have formed naturally due to the symmetry of the GaN and are extremely smooth. The pyramids include templated InGaN/GaN quantum wells, grown using similar steps to those used to produce conventional planar InGaN quantum wells. The parameters are those which would be expected to give 440 nm emission when grown conventionally.

7 phys. stat. sol. (a) 201, No. 4 (2004) / µm Fig. 7 (online colour at: Left: Plan-view secondary electron image of the III-N pyramids taken with a 5 kev beam. Right (colour): map of CL peak position across one pyramid (violet = 410 nm, red = 460 nm and above) Data from a room temperature CL hyperspectral image of one of the pyramid structures are shown in Fig. 8, representing a µm area mapped using a 5 kev, 500 pa electron beam. Individual CL spectra show two bands related to the InGaN emission. A single blue luminescence peak, centred at 440 nm and corresponding to a conventional quantum well, is collected from the pyramid sidewalls whilst an intense green emission is localised at the peak, which shows a reduced blue emission. The images of the CL intensity within the wavelength bands nm and nm are shown in Fig. 8 along with representative spectra. A linescan of the nm CL intensity across the apex of the pyramid (Fig. 9) shows that the bright luminescence to be restricted to a region of less than 300 nm. This suggests Artikel I facet apex Wavelength (nm) Fig µm images of the CL intensity within the wavelength bands nm (left) and nm (right) for the 5-quantum well structure. Representative CL spectra are shown. CL counts (a.u.) Distance (µm) Fig. 9 CL intensity linescan across the apex of a pyramid.

8 672 R. W. Martin et al.: Cathodoluminescence spectral mapping of III-nitride structures the formation of quantum dots at the tip of the pyramid [13], although we can not rule out contributions due to light guiding effects, defects and/or variations in InN concentration. Confirmation of the presence of quantum dots has been provided by low temperature micro PL [18]. The observation of quantum dots emitting at lower energies than the parent quantum wells may seem surprising. However the critical dimension of the dots will be larger than the width of the wells, leading to the possibility of lower confinement, and in this case there will also be a stronger red-shift due to the intense in-built electric fields within strained InGaN structures. Similar behaviour has been reported for the analogous situation of quantum wires formed in V-grooved substrates [19]. 4 Summary The application of room temperature cathodoluminescence spectral mapping to the characterisation of µm scale III-N semiconductor structures has been described. For laterally overgrown III-N structures, using both single layer SiO 2 and multi-layer SiO 2 /ZrO 2 DBR masks, the data clearly shows the improvement in material due to the lateral growth and the effect of cavity formation. The spectral resolution is such that important information is extracted from CL peak shifts of less than 1 nm. Results from selectively grown III-N pyramids demonstrate a spatial resolution of ~100 nm and provide possible evidence for quantum dot formation at the pyramid apices. The power of the CL spectral mapping is enhanced by combination with simultaneously acquired WDX data as exemplified by a study of the relationship of composition and CL on a sub-micron scale in an InGaN epilayer. Acknowledgements We are grateful to Mr. H. M. H. Chong and Prof. R. M. De La Rue for preparing the masks for selective growth and to the UK Engineering and Physical Sciences Research Council and the Research and Development Fund of the University of Strathclyde for financial support. References [1] C. E. Norman, Inst. Phys. Conf. Ser. 169, 557 (2001). [2] J. Christen, M. Grundmann, and D. Bimberg, J. Vac. Sci. Technol. B 9, 2358 (1991). [3] R. W. Martin, P. R. Edwards, K. P. O Donnell, E. G. Mackay, and I. M. Watson, phys. stat. sol. (a) 192, 117 (2002). [4] P. R. Edwards, R. W. Martin, K. P. O Donnell, and I. M. Watson, phys. stat. sol. (c) 0, 2474 (2003). [5] R. W. Martin, S. Dalmasso, K. P. O Donnell, Y. Nakanishi, A. Wakahara, A. Yoshida & the RENiBEl Network, Mat. Res. Soc. Symp. Proc. 743, 411 (2003). [6] P. Hovington, D. Drouin, and R. Gauvin, Scanning 19, 1 (1997); CASINO software. [7] J. Christen and T. Riemann, phys. stat. sol. (b) 228, 419 (2001). [8] E. Feltin, B. Beaumont, P. Vennegues, M. Vaille, P. Gibart, T. Riemann, J. Christen, L. Dobos, and B. Pecz, J. Appl. Phys. 93, 182 (2003). [9] I. M. Watson, C. Liu, K.-S. Kim, H.-S. Kim, C. J. Deatcher, J. M. Girkin, M. D. Dawson, P. R. Edwards, C. Trager-Cowan, and R. W. Martin, phys. stat. sol. (a) 188, 743 (2001). [10] R. W. Martin, P. R. Edwards, R. Pecharroman-Gallego, C. Trager-Cowan, T. Kim, H.-S. Kim, K.-S. Kim, I. M. Watson, M. D. Dawson, T. F. Krauss, J. H. Marsh, and R. M. De La Rue, phys. stat. sol. (a) 183, 145 (2001). [11] R. W. Martin, P. R. Edwards, H.-S. Kim, K.-S. Kim, T. Kim, I. M. Watson, M. D. Dawson, Y. Cho, T. Sands, and N. W. Cheung, Appl. Phys. Lett. 79, 3029 (2001). [12] F. Bertram, S. Srinivasan, R. Liu, L. Geng, F. A. Ponce, T. Riemann, J. Christen, S. Tanaka, H. Omiya and Y. Nakagawa. Mater. Sci. Eng. B 93, 19 (2002). [13] K. Tachibana, T. Someya, S. Ishida, and Y. Arakawa, Appl. Phys. Lett. 76, 3213 (2000). [14] H. X. Jiang, J. Y. Lin, K. C. Zeng, and W. Yang, Appl. Phys. Lett. 75, 763 (1999). [15] R. E. Pritchard et al., J. Appl. Phys. 90, 475 (2001). [16] H. W. Choi, C. W. Jeon, M. D. Dawson, P. R. Edwards, R. W. Martin, and S. Tripathy, J. Appl. Phys. 93, 5978 (2003). [17] F. Bertram, J. Christen, M. Schmidt, K. Hiramatsu, S. Kitamura, and N. Sawaki, Physica E 2, 552 (1998). [18] R. A. Taylor, J. H. Rice, J. W. Robinson, and J. H. Na, unpublished work. [19] R. Roshan, N. I. Cade, A. C. Maciel, J. F. Ryan, A. Schwarz, T. Schapers, and H. Luth, Physica E 13, 174 (2002).