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GaN superlattices by NH3-based molecular beam epitaxy
Journal of Crystal Growth 346 (2012) 50–55
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Journal of Crystal Growth journal homepage: www.elsevier.com/locate/jcrysgro
NH3-rich growth of InGaN and InGaN/GaN superlattices by NH3-based molecular beam epitaxy J.R. Lang n, J.S. Speck Materials Department, University of California, Santa Barbara, CA 93106-5050, USA
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a b s t r a c t
Article history: Received 11 February 2012 Accepted 25 February 2012 Communicated by H. Asahi Available online 7 March 2012
N-rich growth by NH3-based molecular beam epitaxy was investigated for intermediate-temperature GaN and InGaN on c-plane GaN templates. The dependences of growth mode and surface morphology on group-V overpressure, In/Ga ratio, and temperature were explored with atomic force microscopy and high resolution x-ray diffraction. Extension to an ‘‘ultra-NH3-rich’’ regime of very high NH3-flows showed a decreased growth rate and increased In-content for InGaN alloys for constant group III source fluxes. Rapid modulation of NH3 overpressure, growth rate, and substrate temperature has enabled the growth of high quality, many-period InGaN/GaN superlattices, while suppressing morphological instabilities and subsequent stress relaxation. Published by Elsevier B.V.
Keywords: A1. Crystal morphology A3. Molecular beam epitaxy A3. Superlattices B1. Nitrides B2. Semiconducting indium compounds
1. Introduction NH3 MBE offers an avenue to combine the beneficial features of both metal organic chemical vapor deposition (MOCVD) and traditional plasma-assisted MBE (PAMBE) for III-nitride epitaxy. NH3 MBE can combine the high potential growth rates and N-rich growth environments of MOCVD with the low carbon environment of PAMBE. The reduced hydrogen incorporation and realization of p-type behavior in as-grown GaN:Mg [1] may create new opportunities for vertical device structures. While MBE has found success in majority carrier electron devices, such as high electron mobility transistors (HEMTs) [2–4], the pursuit of vertical optoelectronic minority carrier devices has been more difficult. There have been recent reports of increasingly long-wavelength emitters by PAMBE, likely owing much to the availability of much higher quality substrates [5–8]. The reports of visible-spectrum and vertical devices from NH3 MBE have recently shown improvement as well, with violet laser diodes [9], very low leakage and high ideality p–n diodes [10], and low leakage high IQE solar cells [11]. These results demonstrate that the benefits of an N-rich or NHX-terminated surface and high vacuum can be effectively realized in an NH3 MBE to produce high quality vertical devices. The great challenge of group-III nitride MBE is the group-V source, as it is available only in gaseous form and is naturally inert, owing to the very high N2 bond strength of approximately 9.7 eV [12]. The approach of PAMBE utilizes a RF-plasma cracking cell to produce ‘‘active’’ nitrogen, but this method also likely produces significantly higher energy species than a thermal source, and these may
n
Corresponding author. E-mail address: [email protected] (J.R. Lang).
0022-0248/$ - see front matter Published by Elsevier B.V. doi:10.1016/j.jcrysgro.2012.02.036
contribute to defect formation [13]. Additionally, the optimal growth morphology for PAMBE was found to occur for metal-rich growth with a metal adlayer or bilayer [14]. This metal layer behaves as an auto-surfactant, but the metal-rich conditions may also result in metal-rich threading dislocation cores that give rise to high vertical leakage currents [15–17]. Similar to MOCVD, NH3 MBE uses a nitrogen precursor with significantly reduced bond strength to make available active nitrogen. Unlike PAMBE, which can incorporate active nitrogen species into the metal adlayer and behave analogous to a vapor–liquid–solid growth, optimal conditions for NH3 MBE have been realized for N-rich conditions [18,19]. The accumulation of metal on the surface may inhibit growth by blocking the adsorption of NH3, as demonstrated by a decrease in growth rate with increasing Ga flux beyond stoichiometric conditions, while the accumulated Ga droplets lead to surface pitting. The chemical pathway for NH3 decomposition on the GaN surface is unclear. Additionally, the steady-state nature of the GaN surface is also poorly understood. It seems likely that the physisorbed or chemisorbed NHx (x¼1–3) species are catalytically thermolysed on the GaN surface—an endothermic process that appears to decrease in efficiency with decreasing temperature [20]. These issues of active nitrogen supply in MBE motivate the work in this study: increasing supply of active nitrogen for low temperature growth by increasing gas flow from a conventional NH3-showerhead injector to increase solubility limits of InN in GaN for a given growth temperature.
2. Experimental All samples were grown on 3.5 mm thick conductive MOCVDgrown GaN:Si templates on sapphire by Lumilog with nominal Sidoping concentration of 2 1018 cm 3 and threading dislocation
J.R. Lang, J.S. Speck / Journal of Crystal Growth 346 (2012) 50–55
density of 5 108 cm 2, or in-house MOCVD-grown templates with similar structural and morphological quality with nominally the same doping level. These single-sided polished wafers were backside coated with 0.5 mm of Ti metal using standard electronbeam evaporation to provide an infrared absorbing layer for heating in the MBE chamber. The growth experiments were carried out on a Veeco Gen930 MBE system with dual-filament thermal effusion cells for In and Ga sources. High purity NH3 was provided through a mass flow controller (MFC) with gas flows between 0.2 and 1.0 SLM (standard liters per minute), leading to beam equivalent pressures (BEPs) of 0.25–1.12 mTorr. The groupIII sources were operated with BEPs between 2.5 10 8 and 2 10 7 Torr, leading to beam flux ion gage measured V/III ratios of 103–104, which are similar to molar NH3/metalorganic ratios in nitride MOCVD [21]. The growth temperatures were measured by both a thermocouple embedded in the heater and a line-of-sight IR pyrometer that was calibrated to the melting point of Al. We report the substrate surface temperature determined by pyrometry. Film characterization was performed using high resolution xray diffraction (HRXRD) on a Philips Panalytical MRD PRO diffractometer, atomic force microscopy (AFM) using a Digital Instruments D3000 AFM, and in-situ reflection high energy electron diffraction (RHEED). Film thickness and composition were determined by HRXRD measurements with film peak ¨ Pendellosung fringe spacing and peak separation, respectively. For these thin layers, fully-strained films were assumed with elastic stiffness coefficients as calculated by Wright [22].
3. Results The results are presented here in a similar order to that in which they were observed, with an overarching goal of obtaining device-quality InGaN thin films. The primary effects of large NH3 flux on InGaN alloy composition and growth rate are presented first. Growth rate effects are then compared with calculated gasphase scattering processes. Following this understanding of growth rate, this variable is eliminated to further examine and verify InGaN composition trends with NH3 flux and substrate temperature. Finally, morphological trends are examined with RHEED and AFM, and the acquired information is employed in the growth of high-quality multi-quantum-well and superlattice structures. 3.1. Dependence of InGaN composition and growth-rate on NH3 flux A series of growths were performed at fixed In and Ga fluxes ( 4.5 and 5 10 8 Torr BEP, respectively) and a substrate growth temperature of 615 1C. The NH3 flows as provided by the MFC, were varied from 0.2 to 1.0 SLM, corresponding to BEPs of 0.25–1.12 mTorr. All InGaN films for these data were grown with an indium flux below the indium-accumulation threshold. The composition and growth rate trends for the InGaN films are shown in Fig. 1. The InN fraction (XIn) was clearly shown to increase from 0.06 to 0.105 with increasing NH3 supply, a qualitatively similar behavior to the results obtained by increasing nitrogen supply during PAMBE growth of InGaN layers [8].
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Fig. 1. Composition and growth rate trends for InGaN films grown at 615 1C as a function of NH3 beam equivalent pressure (BEP).
¨ x-ray Pendellosung fringes for thickness measurements. An identical trend of decreasing growth rate with increased NH3 (a growth rate reduction of approximately 50% for the highest NH3 flux) was observed for the GaN films. Loss of incident groupIII species through gas-phase scattering processes was suspected as the cause for the decrease in growth rate. Fig. 2(A) shows the dependence of normalized growth rates for the different films on the NH3 BEP, co-plotted with a family of curves describing scattering-attenuated Ga fluxes for different beam interaction distances. The relevant reactor geometry is shown and the beam interaction distance (x) is defined in the schematic in Fig. 2(B). This simple approximation of beam attenuation through scattering can be described by a statistical scattering process for a population of Ga atoms traveling through a distance, x, with mean free path, l, as follows [23]: N ¼ fraction unscattered ¼ ex=l : N0
ð1Þ
Based on the source geometry as described in Fig. 2(B), values of x (5, 7, and 9 cm) were selected to approximate the 7.5 cm source– substrate distance of the NH3-injector. The values of l appropriate for the gas densities during growth were calculated by considering the Ga and NH3 beams to interact as a mixture of two isotropic gases having a Maxwellian velocity distribution. A full treatment of the problem has been presented by Kennard [23]. We consider a low density of Ga relative to NH3, such that self-scattering in the Ga beam is negligible. The mean free path of gas A in gas B can be defined as the ratio of average velocity (vA ) and scattering rate per particle of A (yAB):
lAB ¼
vA
yAB
,
ð2Þ
with the scattering rate given as the product of the number density of B (nB), the mutual cross section of A and B (SAB), and the relative velocity (vr):
3.2. Dependence of growth rate on NH3 flux: gas scattering
yAB ¼ nB SAB vr :
As the indium incorporation increased with NH3 flux, however, the growth rates were observed to decrease by approximately 50% across the range of NH3 fluxes investigated. GaN films were also grown at 830 1C over the same range of NH3 fluxes to provide a control for growth rate trends. The GaN films were grown with an internal AlN marker such that the top GaN layer produced
The process described by (3) counts the collisions per unit time occurring in a ‘‘volume’’ of area SAB and length vr. By assuming the Maxwellian velocity distribution and integrating over all velocities, the scattering rate can be reduced to
yAB ¼ nB SAB ðvA 2 þ vB 2 Þ1=2 :
ð3Þ
ð4Þ
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J.R. Lang, J.S. Speck / Journal of Crystal Growth 346 (2012) 50–55
Fig. 2. (A) Normalized film growth rates for GaN with two differently-located but otherwise equivalent Ga effusion sources (Ga1 and Ga2) and InGaN compared to scattering attenuated Ga fluxes (calculated). (B) Reactor geometry schematic. X is the ‘‘beam interaction distance’’ used for calculating scattering attenuation.
By combining (2) with (4) and reducing, one may come to a convenient form for the mean free path as follows: !1=2 v 2 lAB 1 ¼ nB SAB 1þ B 2 : ð5Þ vA If one assumes the hard sphere model, the mutual cross section becomes simply related to the average particle diameter (d) of A and B: 2
SAB ¼ pdAB ¼
p 4
ðdA þdB Þ2 :
ð6Þ
Because average particle velocity can be related to temperature (T) and particle mass (m) as 8kB T , ð7Þ v¼ pm with kB being Boltzmann’s constant, a final useful form can be found by substituting (6) and (7) in (5): p T NH3 mGa 1=2 nNH3 ðdGa þ dNH3 Þ2 1 þ lGa 1 ¼ : ð8Þ 4 T Ga mNH3 The density of NH3 was calculated assuming the ideal relation between pressure (p), temperature and particle density [24]: nNH3 ¼
pNH3 : kB T NH3
Fig. 3. Composition vs. temperature of 30 nm thick InGaN films at two NH3 pressures—all at constant growth rate. Group-III fluxes were adjusted upwards to compensate for the drop in growth rate with higher NH3, while In/Ga ratio was kept constant. Lines are linear least-square fits to the data shown.
ð9Þ
Pressures used were as measured by the beam flux ionization gage. Temperature for Ga was assumed to be the source temperature (1200 K), as the effused atoms escape through an aperture in the crucible from the gas in equilibrium with the melt. A full treatment of the behavior of the NH3 beam (whereby a finite amount of expansion cooling may occur) is beyond the scope of this work, but a first approximation of the temperature is taken as the injector temperature, which remained at room temperature (300 K). The particle diameters of Ga and NH3 species were taken as 270 pm and 380 pm, respectively, and masses as 17 amu (NH3) and 69.73 amu (Ga) [25–27]. 3.3. Scattering-corrected constant growth rate: composition vs. NH3 flux and temperature To account for any effect of the difference in growth rate when comparing InGaN films grown at different NH3 fluxes, an additional series of samples was grown at constant growth rate, each to a thickness of 30 nm for three temperatures and two NH3 fluxes. To
compensate for the reduced growth rate at higher NH3 flux, the In and Ga source fluxes were adjusted while the In/Ga ratio was fixed at 9/10, as in the previous experiment. The results, given in Fig. 3, show a marked increase in In-content for all temperatures with the increase in NH3 flux. These results indicate that decreasing the temperature and increasing the V/III ratio are the most effective means of increasing the InN solubility in GaN with NH3 MBE. 3.4. Growth mode and surface morphology To probe the surface morphology of these films, in-situ RHEED and ex-situ AFM were used. The observation of time-dependent RHEED specular spot intensity can provide insight to the growth mode in real time. NH3-rich growth offers a unique opportunity to observe surface processes absent on a metal-rich surface [28,29]. In this work, the observation of RHEED intensity behavior during growth demonstrated a transition from step-flow to layer-bylayer growth modes as the substrate temperature was cooled below 700 1C. Above this approximate transition temperature, intensity transients upon growth initiation are singular and
J.R. Lang, J.S. Speck / Journal of Crystal Growth 346 (2012) 50–55
dampen continuously to a constant intensity, a behavior indicative of step-flow growth. In contrast, RHEED specular spot intensity for growth initiated below the transition temperature exhibits many (410) periods of oscillation within a dampening envelope, indicating layer-by-layer growth. Following the standard procedure, the substrate was kept under an NH3 flux of 0.25 mTorr (BEP) during heating to prevent thermal decomposition. Growth was initiated by opening the group-III shutters under this NH3 flux. Fig. 4 shows the evolution of the normalized specular spot intensity after the onset of growth for an InGaN layer at 620 1C. The symmetric dampening of the amplitude (both maxima and minima decaying to the same steady state) is indicative of a competition between step flow and island nucleation at this temperature and flux rate [30]. To confirm these expectations of growth mode, ex-situ AFM was performed on bulk-layer InGaN samples grown at three temperatures and two NH3 fluxes, as shown in Fig. 5. The trend with decreasing growth temperature is the expected increase in island density as adatom surface mobility is decreased. As a consequence of this high island nucleation rate, growth rates were reduced to less than 0.1 mm/h to prevent significant surface roughening. Because of the low growth temperatures necessary, bulk
53
layers of indium atomic fraction greater than 0.11 and thickness greater than 60 nm still remained quite challenging. While growth temperature and V/III ratio was sufficient to prevent metal accumulation on the surface, the morphology did not seem NH3flux dependent. For growth temperatures below 600 1C, at the lowest NH3 flux used in this study (0.25 mTorr) and In and Ga BEPs of 4.5 and 5.0 10 8, respectively, a qualitative change in morphology occurred with a significant increase in film roughness, correlated with growth in an indium-rich regime. These samples showed In surface accumulation that was verified in-situ by a decrease in RHEED intensity with a streaky pattern and ex-situ by In-metallic peaks in x-ray measurements. (In-metal XRD peaks vanished after etching the samples in HCl.) The indium-rich films also showed negligible indium incorporation from x-ray measurements, indicating a likely surface segregation effect [31]. These results suggest that maximum indium incorporation for a given set of conditions would be found at the onset of indium accumulation. This is the expected result as the maximum achievable indium chemical potential for the growing film is that of indium metal in local equilibrium with the growth front: mIn_liquid ¼ mIn_InGaN. 3.5. Growth of capped, temperature-modulated MQWs and SLs As demonstrated in the previous section, the high island nucleation rate for sustained growth at the temperatures necessary to incorporate significant indium into InGaN films makes the growth of thick layers challenging. However, in many device applications of InGaN, thick layers are unnecessary and the multiquantum-well active region design allows for a solution via temperature modulated growth. The additional surface stability afforded by a large NH3 overpressure enabled a cyclic MQW growth procedure as follows: (1) InGaN quantum well was initiated under high NH3 flux at the desired InGaN growth temperature and then (2) the In source was shuttered, leading to the immediate growth of a GaN ‘‘cap layer’’ to stabilize the surface, and finally (3) the substrate was heated to ideal GaN growth temperature for growth of a barrier layer. It has been found that a 2 nm thick cap layer is sufficient to stabilize the QW for heating under NH3 to a barrier growth temperature 200 1C hotter than the InGaN growth temperature. RHEED intensity oscillations, as shown in Fig. 6, allow the absolute calibration of
Fig. 4. RHEED specular spot intensity oscillations for the onset of InGaN growth at 620 1C, indicating a layer-by-layer growth mode.
Fig. 5. Atomic force microscopy (AFM) scans of surface morphology vs. temperature at two NH3 pressures, with constant group III fluxes. All sizes are 5 5 mm2. Two lower right panes (boxed) were found to be in the ‘‘indium accumulation’’ regime.
Fig. 6. RHEED specular spot intensity of InGaN QW growth with target thickness of 2.5 nm. Approximately 10 MLs of deposition are indicated by the oscillations prior to the sudden increase in intensity that occurred with the growth of the lowtemperature GaN cap layer.
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Fig. 7. HRXRD o 2y scan of a 28 period InGaN/GaN superlattice, with an AFM micrograph (inset) of the GaN-terminated superlattice surface morphology.
precise quantum well thicknesses. The RHEED specular spot average intensity and oscillation amplitude exhibited a sudden increase at the onset of cap layer growth. By heating across the transition between layer-by-layer and step-flow growth modes, a highly-stepped and smooth surface can be recovered during barrier growth, and the time-average morphology can remain stable. (The quantum wells may also benefit from some annealing effects.) The stability of the surface allows for growth of nearly arbitrary active region thicknesses without significant morphological degradation. Fig. 7 shows an x-ray diffraction measurement and AFM micrograph of a 28 period InGaN (2.5 nm)/GaN (4 nm) superlattice terminated at a 4 nm GaN layer. Multiple higherorder peaks indicate good interfacial quality throughout the structure. The sample exhibited strong room-temperature photoluminescence at a wavelength centered at 440 nm, indicating efficient radiative recombination from the InGaN layers. The surface morphology shows the recovery of step-flow growth following 70 nm of total InGaN thickness grown well below the critical temperature for step flow. Thinner 3-well MQW LED devices grown using the same methods have exhibited electroluminescence ranging from 400 to 500 nm in wavelength. This modulated growth method demonstrates that thick InGaN/GaN active regions are possible and may improve the quality of thinner active region devices in NH3 MBE.
4. Discussion It is known from bulk measurements of InGaN decomposition [32] and NH3-MBE growth of GaN at high temperature [19,33] that increased NH3 overpressure will suppress the thermal decomposition of nitride films. Thus, it is expected that the relatively weakly bound InN component of an InGaN film would be stabilized by the increased supply of NH3 as well. The temperature and N-flux dependent results in [8] are interpreted through the activation energy of nitrogen loss from InN and correlated with the bond dissociation energy of InN. While these PAMBE results are also consistent with observations of InGaN growth by NH3 MBE, behavior in the In-accumulation regime seems to differ, by exhibiting little surface segregation effect. Work in PAMBE of InGaN has found indium incorporation for a given temperature and N-flux to saturate beyond the accumulation threshold In-flux and not subsequently
decrease [34]. In contrast, In-droplets on the surface of an InGaN film in NH3 MBE seem to act as indium sinks, reducing the indium incorporation into the film and blocking the adsorption of NH3. These effects result in rough morphology and negligible indium content in the films. MOCVD studies [21] have found a dependence of the indium accumulation threshold on NH3, while previous studies of NH3 MBE [31] did not. This work does find an NH3dependent indium accumulation threshold because a larger fraction of the incident indium is incorporated for higher V/III ratio, though it is not necessarily in disagreement with the previous studies. To the authors0 best knowledge, this is the first report of NH3 MBE operating at such high NH3 fluxes, and this enables access to a broader growth parameter space than was previously available. The maximum achieved V/III ratio in this work was limited by maximum installed MFC flow-rate capacity and minimum group-III limited practical growth rates. While a further optimized NH3injector and source geometry could enhance V/III ratios, the gas phase scattering process described herein provides the practical upper limit on V/III ratio in NH3 MBE. A further concern is the potentially increased sensitivity to NH3 source contamination, as the ratio of effective partial pressures of NH3-borne contaminants to group-III sources increases with V/III ratio. Thus, a growth system and device structure dependent optimum point for high V/III ratio NH3 MBE is required for maximizing indium incorporation and surface quality for InGaN layers. The surface stabilizing effect of the high NH3-overpressure and thin GaN ‘‘capping layers’’ on the InGaN surface has become an important technique in MOCVD growth of high quality MQW active regions [35–37], but has not, until this work, been implemented in MBE growth. While this technique has improved InGaN/GaN MQW morphology, such as the reduction in V-defect depth and increased optical powers as measured by PL and EL, the benefits are qualitatively different when compared to NH3 MBE. Because of the lower growth temperature in MBE required to incorporate indium, there is a transition to an unstable island nucleation/layer-by-layer growth mode that is not generally present at MOCVD growth temperatures. For thin layers, such as quantum wells, the alloy composition of the InGaN layers may be significantly extended as compared to bulk layers without transition to 3-dimensional growth, as evidenced by the sample shown in Fig. 7. The limit for device quality bulk layers, explored through the fabrication of solar cell devices, seems to appear near an indium fraction XIn 0.11. In contrast, quantum-well based solar cells have exhibited favorable properties compared to bulk-like devices, such as a significant increase in open-circuit voltage of up to 0.5 V for indium fraction held constant and the ability to extend the indium fraction much further before significant performance breakdown. The ability to grow functioning device structures with XIn 40.2 has been demonstrated by significant RT EL from LED structures extending to 495 nm in output wavelength. Thus, through the use of temperature modulated growth with high NH3-overpressure wells and low temperature capping layers, useful InGaN compositions may be extended for both solar cell and LED devices by NH3 MBE.
5. Conclusion In conclusion, this work focused on the effects of expanding NH3 MBE to the upper limits of NH3 flux available in a conventional MBE system on the growth of InGaN and GaN films at intermediate growth temperature. It was observed that upon increasing the NH3 BEP from 0.25 to 1.1 mTorr during the growth of InGaN films, a growth rate decrease of 50% was observed, while In-content increased by nearly 50% for constant group-III source fluxes and growth temperature. The growth rate decrease was
J.R. Lang, J.S. Speck / Journal of Crystal Growth 346 (2012) 50–55
found to be independent of alloy content and growth temperature by verifying with growth of pure GaN films. The growth rate data was consistent with estimations of gas phase scattering effects from the high NH3 pressure near the growth surface. All InGaN films were determined to grow in a layer-by-layer mode through the observation of RHEED oscillations and ex-situ AFM measurements. Indium incorporation was found to be maximized through the use of decreased growth temperature and increased NH3 flux, while maintaining growth parameters below the indium surfaceaccumulation threshold. Through the use of these observed growth conditions and temperature modulated growth to mitigate progressive roughening of the InGaN, high quality InGaN/ GaN superlattices and MQWs were grown. This set of growth techniques, when combined with the other advantages of MBE, has the potential to enable new device structures for high performance and scientific measurement. References [1] A. Dussaigne, B. Damilano, J. Brault, J. Massies, E. Feltin, N. Grandjean, Journal of Applied Physics 103 (2008) 013110. [2] Christiane Poblenz, Andrea L. Corrion, Felix Recht, Chang Soo Suh, Rongming Chu, Likun Shen, James S. Speck, Umesh K. Mishra, IEEE Electron Device Letters 28 (11) (2007) 945–947. [3] Rongming Chu, Christiane Poblenz, Man Hoi Wong, Sansaptak Dasgupta, Siddharth Rajan, Yi Pei, Felix Recht, Likun Shen, James S. Speck, Umesh K. Mishra, Applied Physics Express 1 (2008) 061101. [4] Uttam Singisetti, Man Hoi Wong, Sansaptak Dasgupta, James S. Speck, Umesh K. Mishra, Applied Physics Express 4 (2011) 024103. [5] C. Skierbiszewski, P. Wis´niewski, M. Siekacz, P. Perlin, A. Feduniewicz-Zmuda, G. Nowak, I. Grzegory, M. Leszczyn´ski, S. Porowski, Applied Physics Letters 88 (2006) 221108. [6] C. Skierbiszewski, M. Siekacz, P. Perlin, A. Feduniewicz-Zmuda, G. Cywinski, I. Grzegory, M. Leszczynski, Z.R. Wasilewski, S. Porowski., Journal of Crystal Growth 305 (2007) 346–354. [7] C. Skierbiszewski, Z.R. Wasilewski, I. Grzegory, S. Porowski, Journal of Crystal Growth 311 (2009) 1632–1639. [8] M. Siekacz, M. Sawicka, H. Turski, G. Cywinski, A. Khachapuridze, P. Perlin, T. Suski, M. Bockowski, J. Smalc-Koziorowska, M. Krysko, R. Kudrawiec, M. Syperek, J. Misiewicz, Z. Wasilewski, S. Porowski, C. Skierbiszewski., Journal of Applied Physics 110 (2011) 063110. [9] J. Heffernan, M. Kauer, K. Johnson, C. Zellweger, S.E. Hooper, V. Bousquet, Physica Status Solidi (a) 202 (5) (2005) 868–874. [10] C.A. Hurni, O. Bierwagon, J.R. Lang, B.M. McSkimming, C.S. Gallinat, E.C. Young, D.A. Browne, U.K. Mishra, J.S. Speck, Applied Physics Letters 97 (2010) 222113. [11] J.R. Lang, C.J. Neufeld, C.A. Hurni, S.C. Cruz, E. Matioli, U.K. Mishra, J.S. Speck, Applied Physics Letters 98 (2011) 131115.
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Journal of Crystal Growth journal homepage: www.elsevier.com/locate/jcrysgro
NH3-rich growth of InGaN and InGaN/GaN superlattices by NH3-based molecular beam epitaxy J.R. Lang n, J.S. Speck Materials Department, University of California, Santa Barbara, CA 93106-5050, USA
a r t i c l e i n f o
a b s t r a c t
Article history: Received 11 February 2012 Accepted 25 February 2012 Communicated by H. Asahi Available online 7 March 2012
N-rich growth by NH3-based molecular beam epitaxy was investigated for intermediate-temperature GaN and InGaN on c-plane GaN templates. The dependences of growth mode and surface morphology on group-V overpressure, In/Ga ratio, and temperature were explored with atomic force microscopy and high resolution x-ray diffraction. Extension to an ‘‘ultra-NH3-rich’’ regime of very high NH3-flows showed a decreased growth rate and increased In-content for InGaN alloys for constant group III source fluxes. Rapid modulation of NH3 overpressure, growth rate, and substrate temperature has enabled the growth of high quality, many-period InGaN/GaN superlattices, while suppressing morphological instabilities and subsequent stress relaxation. Published by Elsevier B.V.
Keywords: A1. Crystal morphology A3. Molecular beam epitaxy A3. Superlattices B1. Nitrides B2. Semiconducting indium compounds
1. Introduction NH3 MBE offers an avenue to combine the beneficial features of both metal organic chemical vapor deposition (MOCVD) and traditional plasma-assisted MBE (PAMBE) for III-nitride epitaxy. NH3 MBE can combine the high potential growth rates and N-rich growth environments of MOCVD with the low carbon environment of PAMBE. The reduced hydrogen incorporation and realization of p-type behavior in as-grown GaN:Mg [1] may create new opportunities for vertical device structures. While MBE has found success in majority carrier electron devices, such as high electron mobility transistors (HEMTs) [2–4], the pursuit of vertical optoelectronic minority carrier devices has been more difficult. There have been recent reports of increasingly long-wavelength emitters by PAMBE, likely owing much to the availability of much higher quality substrates [5–8]. The reports of visible-spectrum and vertical devices from NH3 MBE have recently shown improvement as well, with violet laser diodes [9], very low leakage and high ideality p–n diodes [10], and low leakage high IQE solar cells [11]. These results demonstrate that the benefits of an N-rich or NHX-terminated surface and high vacuum can be effectively realized in an NH3 MBE to produce high quality vertical devices. The great challenge of group-III nitride MBE is the group-V source, as it is available only in gaseous form and is naturally inert, owing to the very high N2 bond strength of approximately 9.7 eV [12]. The approach of PAMBE utilizes a RF-plasma cracking cell to produce ‘‘active’’ nitrogen, but this method also likely produces significantly higher energy species than a thermal source, and these may
n
Corresponding author. E-mail address: [email protected] (J.R. Lang).
0022-0248/$ - see front matter Published by Elsevier B.V. doi:10.1016/j.jcrysgro.2012.02.036
contribute to defect formation [13]. Additionally, the optimal growth morphology for PAMBE was found to occur for metal-rich growth with a metal adlayer or bilayer [14]. This metal layer behaves as an auto-surfactant, but the metal-rich conditions may also result in metal-rich threading dislocation cores that give rise to high vertical leakage currents [15–17]. Similar to MOCVD, NH3 MBE uses a nitrogen precursor with significantly reduced bond strength to make available active nitrogen. Unlike PAMBE, which can incorporate active nitrogen species into the metal adlayer and behave analogous to a vapor–liquid–solid growth, optimal conditions for NH3 MBE have been realized for N-rich conditions [18,19]. The accumulation of metal on the surface may inhibit growth by blocking the adsorption of NH3, as demonstrated by a decrease in growth rate with increasing Ga flux beyond stoichiometric conditions, while the accumulated Ga droplets lead to surface pitting. The chemical pathway for NH3 decomposition on the GaN surface is unclear. Additionally, the steady-state nature of the GaN surface is also poorly understood. It seems likely that the physisorbed or chemisorbed NHx (x¼1–3) species are catalytically thermolysed on the GaN surface—an endothermic process that appears to decrease in efficiency with decreasing temperature [20]. These issues of active nitrogen supply in MBE motivate the work in this study: increasing supply of active nitrogen for low temperature growth by increasing gas flow from a conventional NH3-showerhead injector to increase solubility limits of InN in GaN for a given growth temperature.
2. Experimental All samples were grown on 3.5 mm thick conductive MOCVDgrown GaN:Si templates on sapphire by Lumilog with nominal Sidoping concentration of 2 1018 cm 3 and threading dislocation
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density of 5 108 cm 2, or in-house MOCVD-grown templates with similar structural and morphological quality with nominally the same doping level. These single-sided polished wafers were backside coated with 0.5 mm of Ti metal using standard electronbeam evaporation to provide an infrared absorbing layer for heating in the MBE chamber. The growth experiments were carried out on a Veeco Gen930 MBE system with dual-filament thermal effusion cells for In and Ga sources. High purity NH3 was provided through a mass flow controller (MFC) with gas flows between 0.2 and 1.0 SLM (standard liters per minute), leading to beam equivalent pressures (BEPs) of 0.25–1.12 mTorr. The groupIII sources were operated with BEPs between 2.5 10 8 and 2 10 7 Torr, leading to beam flux ion gage measured V/III ratios of 103–104, which are similar to molar NH3/metalorganic ratios in nitride MOCVD [21]. The growth temperatures were measured by both a thermocouple embedded in the heater and a line-of-sight IR pyrometer that was calibrated to the melting point of Al. We report the substrate surface temperature determined by pyrometry. Film characterization was performed using high resolution xray diffraction (HRXRD) on a Philips Panalytical MRD PRO diffractometer, atomic force microscopy (AFM) using a Digital Instruments D3000 AFM, and in-situ reflection high energy electron diffraction (RHEED). Film thickness and composition were determined by HRXRD measurements with film peak ¨ Pendellosung fringe spacing and peak separation, respectively. For these thin layers, fully-strained films were assumed with elastic stiffness coefficients as calculated by Wright [22].
3. Results The results are presented here in a similar order to that in which they were observed, with an overarching goal of obtaining device-quality InGaN thin films. The primary effects of large NH3 flux on InGaN alloy composition and growth rate are presented first. Growth rate effects are then compared with calculated gasphase scattering processes. Following this understanding of growth rate, this variable is eliminated to further examine and verify InGaN composition trends with NH3 flux and substrate temperature. Finally, morphological trends are examined with RHEED and AFM, and the acquired information is employed in the growth of high-quality multi-quantum-well and superlattice structures. 3.1. Dependence of InGaN composition and growth-rate on NH3 flux A series of growths were performed at fixed In and Ga fluxes ( 4.5 and 5 10 8 Torr BEP, respectively) and a substrate growth temperature of 615 1C. The NH3 flows as provided by the MFC, were varied from 0.2 to 1.0 SLM, corresponding to BEPs of 0.25–1.12 mTorr. All InGaN films for these data were grown with an indium flux below the indium-accumulation threshold. The composition and growth rate trends for the InGaN films are shown in Fig. 1. The InN fraction (XIn) was clearly shown to increase from 0.06 to 0.105 with increasing NH3 supply, a qualitatively similar behavior to the results obtained by increasing nitrogen supply during PAMBE growth of InGaN layers [8].
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Fig. 1. Composition and growth rate trends for InGaN films grown at 615 1C as a function of NH3 beam equivalent pressure (BEP).
¨ x-ray Pendellosung fringes for thickness measurements. An identical trend of decreasing growth rate with increased NH3 (a growth rate reduction of approximately 50% for the highest NH3 flux) was observed for the GaN films. Loss of incident groupIII species through gas-phase scattering processes was suspected as the cause for the decrease in growth rate. Fig. 2(A) shows the dependence of normalized growth rates for the different films on the NH3 BEP, co-plotted with a family of curves describing scattering-attenuated Ga fluxes for different beam interaction distances. The relevant reactor geometry is shown and the beam interaction distance (x) is defined in the schematic in Fig. 2(B). This simple approximation of beam attenuation through scattering can be described by a statistical scattering process for a population of Ga atoms traveling through a distance, x, with mean free path, l, as follows [23]: N ¼ fraction unscattered ¼ ex=l : N0
ð1Þ
Based on the source geometry as described in Fig. 2(B), values of x (5, 7, and 9 cm) were selected to approximate the 7.5 cm source– substrate distance of the NH3-injector. The values of l appropriate for the gas densities during growth were calculated by considering the Ga and NH3 beams to interact as a mixture of two isotropic gases having a Maxwellian velocity distribution. A full treatment of the problem has been presented by Kennard [23]. We consider a low density of Ga relative to NH3, such that self-scattering in the Ga beam is negligible. The mean free path of gas A in gas B can be defined as the ratio of average velocity (vA ) and scattering rate per particle of A (yAB):
lAB ¼
vA
yAB
,
ð2Þ
with the scattering rate given as the product of the number density of B (nB), the mutual cross section of A and B (SAB), and the relative velocity (vr):
3.2. Dependence of growth rate on NH3 flux: gas scattering
yAB ¼ nB SAB vr :
As the indium incorporation increased with NH3 flux, however, the growth rates were observed to decrease by approximately 50% across the range of NH3 fluxes investigated. GaN films were also grown at 830 1C over the same range of NH3 fluxes to provide a control for growth rate trends. The GaN films were grown with an internal AlN marker such that the top GaN layer produced
The process described by (3) counts the collisions per unit time occurring in a ‘‘volume’’ of area SAB and length vr. By assuming the Maxwellian velocity distribution and integrating over all velocities, the scattering rate can be reduced to
yAB ¼ nB SAB ðvA 2 þ vB 2 Þ1=2 :
ð3Þ
ð4Þ
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Fig. 2. (A) Normalized film growth rates for GaN with two differently-located but otherwise equivalent Ga effusion sources (Ga1 and Ga2) and InGaN compared to scattering attenuated Ga fluxes (calculated). (B) Reactor geometry schematic. X is the ‘‘beam interaction distance’’ used for calculating scattering attenuation.
By combining (2) with (4) and reducing, one may come to a convenient form for the mean free path as follows: !1=2 v 2 lAB 1 ¼ nB SAB 1þ B 2 : ð5Þ vA If one assumes the hard sphere model, the mutual cross section becomes simply related to the average particle diameter (d) of A and B: 2
SAB ¼ pdAB ¼
p 4
ðdA þdB Þ2 :
ð6Þ
Because average particle velocity can be related to temperature (T) and particle mass (m) as 8kB T , ð7Þ v¼ pm with kB being Boltzmann’s constant, a final useful form can be found by substituting (6) and (7) in (5): p T NH3 mGa 1=2 nNH3 ðdGa þ dNH3 Þ2 1 þ lGa 1 ¼ : ð8Þ 4 T Ga mNH3 The density of NH3 was calculated assuming the ideal relation between pressure (p), temperature and particle density [24]: nNH3 ¼
pNH3 : kB T NH3
Fig. 3. Composition vs. temperature of 30 nm thick InGaN films at two NH3 pressures—all at constant growth rate. Group-III fluxes were adjusted upwards to compensate for the drop in growth rate with higher NH3, while In/Ga ratio was kept constant. Lines are linear least-square fits to the data shown.
ð9Þ
Pressures used were as measured by the beam flux ionization gage. Temperature for Ga was assumed to be the source temperature (1200 K), as the effused atoms escape through an aperture in the crucible from the gas in equilibrium with the melt. A full treatment of the behavior of the NH3 beam (whereby a finite amount of expansion cooling may occur) is beyond the scope of this work, but a first approximation of the temperature is taken as the injector temperature, which remained at room temperature (300 K). The particle diameters of Ga and NH3 species were taken as 270 pm and 380 pm, respectively, and masses as 17 amu (NH3) and 69.73 amu (Ga) [25–27]. 3.3. Scattering-corrected constant growth rate: composition vs. NH3 flux and temperature To account for any effect of the difference in growth rate when comparing InGaN films grown at different NH3 fluxes, an additional series of samples was grown at constant growth rate, each to a thickness of 30 nm for three temperatures and two NH3 fluxes. To
compensate for the reduced growth rate at higher NH3 flux, the In and Ga source fluxes were adjusted while the In/Ga ratio was fixed at 9/10, as in the previous experiment. The results, given in Fig. 3, show a marked increase in In-content for all temperatures with the increase in NH3 flux. These results indicate that decreasing the temperature and increasing the V/III ratio are the most effective means of increasing the InN solubility in GaN with NH3 MBE. 3.4. Growth mode and surface morphology To probe the surface morphology of these films, in-situ RHEED and ex-situ AFM were used. The observation of time-dependent RHEED specular spot intensity can provide insight to the growth mode in real time. NH3-rich growth offers a unique opportunity to observe surface processes absent on a metal-rich surface [28,29]. In this work, the observation of RHEED intensity behavior during growth demonstrated a transition from step-flow to layer-bylayer growth modes as the substrate temperature was cooled below 700 1C. Above this approximate transition temperature, intensity transients upon growth initiation are singular and
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dampen continuously to a constant intensity, a behavior indicative of step-flow growth. In contrast, RHEED specular spot intensity for growth initiated below the transition temperature exhibits many (410) periods of oscillation within a dampening envelope, indicating layer-by-layer growth. Following the standard procedure, the substrate was kept under an NH3 flux of 0.25 mTorr (BEP) during heating to prevent thermal decomposition. Growth was initiated by opening the group-III shutters under this NH3 flux. Fig. 4 shows the evolution of the normalized specular spot intensity after the onset of growth for an InGaN layer at 620 1C. The symmetric dampening of the amplitude (both maxima and minima decaying to the same steady state) is indicative of a competition between step flow and island nucleation at this temperature and flux rate [30]. To confirm these expectations of growth mode, ex-situ AFM was performed on bulk-layer InGaN samples grown at three temperatures and two NH3 fluxes, as shown in Fig. 5. The trend with decreasing growth temperature is the expected increase in island density as adatom surface mobility is decreased. As a consequence of this high island nucleation rate, growth rates were reduced to less than 0.1 mm/h to prevent significant surface roughening. Because of the low growth temperatures necessary, bulk
53
layers of indium atomic fraction greater than 0.11 and thickness greater than 60 nm still remained quite challenging. While growth temperature and V/III ratio was sufficient to prevent metal accumulation on the surface, the morphology did not seem NH3flux dependent. For growth temperatures below 600 1C, at the lowest NH3 flux used in this study (0.25 mTorr) and In and Ga BEPs of 4.5 and 5.0 10 8, respectively, a qualitative change in morphology occurred with a significant increase in film roughness, correlated with growth in an indium-rich regime. These samples showed In surface accumulation that was verified in-situ by a decrease in RHEED intensity with a streaky pattern and ex-situ by In-metallic peaks in x-ray measurements. (In-metal XRD peaks vanished after etching the samples in HCl.) The indium-rich films also showed negligible indium incorporation from x-ray measurements, indicating a likely surface segregation effect [31]. These results suggest that maximum indium incorporation for a given set of conditions would be found at the onset of indium accumulation. This is the expected result as the maximum achievable indium chemical potential for the growing film is that of indium metal in local equilibrium with the growth front: mIn_liquid ¼ mIn_InGaN. 3.5. Growth of capped, temperature-modulated MQWs and SLs As demonstrated in the previous section, the high island nucleation rate for sustained growth at the temperatures necessary to incorporate significant indium into InGaN films makes the growth of thick layers challenging. However, in many device applications of InGaN, thick layers are unnecessary and the multiquantum-well active region design allows for a solution via temperature modulated growth. The additional surface stability afforded by a large NH3 overpressure enabled a cyclic MQW growth procedure as follows: (1) InGaN quantum well was initiated under high NH3 flux at the desired InGaN growth temperature and then (2) the In source was shuttered, leading to the immediate growth of a GaN ‘‘cap layer’’ to stabilize the surface, and finally (3) the substrate was heated to ideal GaN growth temperature for growth of a barrier layer. It has been found that a 2 nm thick cap layer is sufficient to stabilize the QW for heating under NH3 to a barrier growth temperature 200 1C hotter than the InGaN growth temperature. RHEED intensity oscillations, as shown in Fig. 6, allow the absolute calibration of
Fig. 4. RHEED specular spot intensity oscillations for the onset of InGaN growth at 620 1C, indicating a layer-by-layer growth mode.
Fig. 5. Atomic force microscopy (AFM) scans of surface morphology vs. temperature at two NH3 pressures, with constant group III fluxes. All sizes are 5 5 mm2. Two lower right panes (boxed) were found to be in the ‘‘indium accumulation’’ regime.
Fig. 6. RHEED specular spot intensity of InGaN QW growth with target thickness of 2.5 nm. Approximately 10 MLs of deposition are indicated by the oscillations prior to the sudden increase in intensity that occurred with the growth of the lowtemperature GaN cap layer.
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Fig. 7. HRXRD o 2y scan of a 28 period InGaN/GaN superlattice, with an AFM micrograph (inset) of the GaN-terminated superlattice surface morphology.
precise quantum well thicknesses. The RHEED specular spot average intensity and oscillation amplitude exhibited a sudden increase at the onset of cap layer growth. By heating across the transition between layer-by-layer and step-flow growth modes, a highly-stepped and smooth surface can be recovered during barrier growth, and the time-average morphology can remain stable. (The quantum wells may also benefit from some annealing effects.) The stability of the surface allows for growth of nearly arbitrary active region thicknesses without significant morphological degradation. Fig. 7 shows an x-ray diffraction measurement and AFM micrograph of a 28 period InGaN (2.5 nm)/GaN (4 nm) superlattice terminated at a 4 nm GaN layer. Multiple higherorder peaks indicate good interfacial quality throughout the structure. The sample exhibited strong room-temperature photoluminescence at a wavelength centered at 440 nm, indicating efficient radiative recombination from the InGaN layers. The surface morphology shows the recovery of step-flow growth following 70 nm of total InGaN thickness grown well below the critical temperature for step flow. Thinner 3-well MQW LED devices grown using the same methods have exhibited electroluminescence ranging from 400 to 500 nm in wavelength. This modulated growth method demonstrates that thick InGaN/GaN active regions are possible and may improve the quality of thinner active region devices in NH3 MBE.
4. Discussion It is known from bulk measurements of InGaN decomposition [32] and NH3-MBE growth of GaN at high temperature [19,33] that increased NH3 overpressure will suppress the thermal decomposition of nitride films. Thus, it is expected that the relatively weakly bound InN component of an InGaN film would be stabilized by the increased supply of NH3 as well. The temperature and N-flux dependent results in [8] are interpreted through the activation energy of nitrogen loss from InN and correlated with the bond dissociation energy of InN. While these PAMBE results are also consistent with observations of InGaN growth by NH3 MBE, behavior in the In-accumulation regime seems to differ, by exhibiting little surface segregation effect. Work in PAMBE of InGaN has found indium incorporation for a given temperature and N-flux to saturate beyond the accumulation threshold In-flux and not subsequently
decrease [34]. In contrast, In-droplets on the surface of an InGaN film in NH3 MBE seem to act as indium sinks, reducing the indium incorporation into the film and blocking the adsorption of NH3. These effects result in rough morphology and negligible indium content in the films. MOCVD studies [21] have found a dependence of the indium accumulation threshold on NH3, while previous studies of NH3 MBE [31] did not. This work does find an NH3dependent indium accumulation threshold because a larger fraction of the incident indium is incorporated for higher V/III ratio, though it is not necessarily in disagreement with the previous studies. To the authors0 best knowledge, this is the first report of NH3 MBE operating at such high NH3 fluxes, and this enables access to a broader growth parameter space than was previously available. The maximum achieved V/III ratio in this work was limited by maximum installed MFC flow-rate capacity and minimum group-III limited practical growth rates. While a further optimized NH3injector and source geometry could enhance V/III ratios, the gas phase scattering process described herein provides the practical upper limit on V/III ratio in NH3 MBE. A further concern is the potentially increased sensitivity to NH3 source contamination, as the ratio of effective partial pressures of NH3-borne contaminants to group-III sources increases with V/III ratio. Thus, a growth system and device structure dependent optimum point for high V/III ratio NH3 MBE is required for maximizing indium incorporation and surface quality for InGaN layers. The surface stabilizing effect of the high NH3-overpressure and thin GaN ‘‘capping layers’’ on the InGaN surface has become an important technique in MOCVD growth of high quality MQW active regions [35–37], but has not, until this work, been implemented in MBE growth. While this technique has improved InGaN/GaN MQW morphology, such as the reduction in V-defect depth and increased optical powers as measured by PL and EL, the benefits are qualitatively different when compared to NH3 MBE. Because of the lower growth temperature in MBE required to incorporate indium, there is a transition to an unstable island nucleation/layer-by-layer growth mode that is not generally present at MOCVD growth temperatures. For thin layers, such as quantum wells, the alloy composition of the InGaN layers may be significantly extended as compared to bulk layers without transition to 3-dimensional growth, as evidenced by the sample shown in Fig. 7. The limit for device quality bulk layers, explored through the fabrication of solar cell devices, seems to appear near an indium fraction XIn 0.11. In contrast, quantum-well based solar cells have exhibited favorable properties compared to bulk-like devices, such as a significant increase in open-circuit voltage of up to 0.5 V for indium fraction held constant and the ability to extend the indium fraction much further before significant performance breakdown. The ability to grow functioning device structures with XIn 40.2 has been demonstrated by significant RT EL from LED structures extending to 495 nm in output wavelength. Thus, through the use of temperature modulated growth with high NH3-overpressure wells and low temperature capping layers, useful InGaN compositions may be extended for both solar cell and LED devices by NH3 MBE.
5. Conclusion In conclusion, this work focused on the effects of expanding NH3 MBE to the upper limits of NH3 flux available in a conventional MBE system on the growth of InGaN and GaN films at intermediate growth temperature. It was observed that upon increasing the NH3 BEP from 0.25 to 1.1 mTorr during the growth of InGaN films, a growth rate decrease of 50% was observed, while In-content increased by nearly 50% for constant group-III source fluxes and growth temperature. The growth rate decrease was
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found to be independent of alloy content and growth temperature by verifying with growth of pure GaN films. The growth rate data was consistent with estimations of gas phase scattering effects from the high NH3 pressure near the growth surface. All InGaN films were determined to grow in a layer-by-layer mode through the observation of RHEED oscillations and ex-situ AFM measurements. Indium incorporation was found to be maximized through the use of decreased growth temperature and increased NH3 flux, while maintaining growth parameters below the indium surfaceaccumulation threshold. Through the use of these observed growth conditions and temperature modulated growth to mitigate progressive roughening of the InGaN, high quality InGaN/ GaN superlattices and MQWs were grown. This set of growth techniques, when combined with the other advantages of MBE, has the potential to enable new device structures for high performance and scientific measurement. References [1] A. Dussaigne, B. Damilano, J. Brault, J. Massies, E. Feltin, N. Grandjean, Journal of Applied Physics 103 (2008) 013110. [2] Christiane Poblenz, Andrea L. Corrion, Felix Recht, Chang Soo Suh, Rongming Chu, Likun Shen, James S. Speck, Umesh K. Mishra, IEEE Electron Device Letters 28 (11) (2007) 945–947. [3] Rongming Chu, Christiane Poblenz, Man Hoi Wong, Sansaptak Dasgupta, Siddharth Rajan, Yi Pei, Felix Recht, Likun Shen, James S. Speck, Umesh K. Mishra, Applied Physics Express 1 (2008) 061101. [4] Uttam Singisetti, Man Hoi Wong, Sansaptak Dasgupta, James S. Speck, Umesh K. Mishra, Applied Physics Express 4 (2011) 024103. [5] C. Skierbiszewski, P. Wis´niewski, M. Siekacz, P. Perlin, A. Feduniewicz-Zmuda, G. Nowak, I. Grzegory, M. Leszczyn´ski, S. Porowski, Applied Physics Letters 88 (2006) 221108. [6] C. Skierbiszewski, M. Siekacz, P. Perlin, A. Feduniewicz-Zmuda, G. Cywinski, I. Grzegory, M. Leszczynski, Z.R. Wasilewski, S. Porowski., Journal of Crystal Growth 305 (2007) 346–354. [7] C. Skierbiszewski, Z.R. Wasilewski, I. Grzegory, S. Porowski, Journal of Crystal Growth 311 (2009) 1632–1639. [8] M. Siekacz, M. Sawicka, H. Turski, G. Cywinski, A. Khachapuridze, P. Perlin, T. Suski, M. Bockowski, J. Smalc-Koziorowska, M. Krysko, R. Kudrawiec, M. Syperek, J. Misiewicz, Z. Wasilewski, S. Porowski, C. Skierbiszewski., Journal of Applied Physics 110 (2011) 063110. [9] J. Heffernan, M. Kauer, K. Johnson, C. Zellweger, S.E. Hooper, V. Bousquet, Physica Status Solidi (a) 202 (5) (2005) 868–874. [10] C.A. Hurni, O. Bierwagon, J.R. Lang, B.M. McSkimming, C.S. Gallinat, E.C. Young, D.A. Browne, U.K. Mishra, J.S. Speck, Applied Physics Letters 97 (2010) 222113. [11] J.R. Lang, C.J. Neufeld, C.A. Hurni, S.C. Cruz, E. Matioli, U.K. Mishra, J.S. Speck, Applied Physics Letters 98 (2011) 131115.
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