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Crystal Structure Induced Preferential Surface Alloying of Sb on Wurtzite/Zinc Blende GaAs Nanowires Martin Hjort, Peter Kratzer, Sebastian Lehmann, Sahil J. Patel, Kimberly A. Dick, Chris J. Palmstrom, Rainer Timm, and Anders Mikkelsen Nano Lett., Just Accepted Manuscript • Publication Date (Web): 24 May 2017 Downloaded from http://pubs.acs.org on May 30, 2017

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Crystal Structure Induced Preferential Surface Alloying of Sb on Wurtzite/Zinc Blende GaAs Nanowires Martin Hjort1, †, Peter Kratzer2, Sebastian Lehmann1, Sahil J. Patel3, Kimberly A. Dick1,4, Chris J. Palmstrøm3,5, Rainer Timm1,6, and Anders Mikkelsen1* 1

Department of Physics and NanoLund, Lund University, P.O. Box 118, 221 00 Lund, Sweden

2

Faculty of Physics and Center for Nanointegration (CENIDE), University Duisburg-Essen, Lotharstrasse 1, 470 48 Duisburg, Germany 3

Materials Department, University of California-Santa Barbara, Santa Barbara, California 93106, USA

4

Centre for Analysis and Synthesis, Lund University, P.O. Box 124, 221 00 Lund, Sweden

5

Department of Electrical and Computer Engineering, University of California-Santa Barbara, Santa Barbara, California 93106, USA 6 California Nanosystems Institute, University of California-Santa Barbara, Santa Barbara, California 93106, USA †

Present address: Department of Materials Science and Engineering, Stanford University. 476 Lomita Mall, Stanford CA 94305

*

e-mail: [email protected]

ABSTRACT: We study the surface diffusion and alloying of Sb into GaAs nanowires (NWs) with controlled axial stacking of wurtzite (Wz) and zinc blende (Zb) crystal phases. Using atomically resolved scanning tunneling microscopy, we find that Sb preferentially incorporates into the surface layer of the 110-terminated Zb segments rather than the 1120-terminated Wz segments. Density functional theory calculations verify the higher surface incorporation rate into the Zb phase and find that it is related to differences in the energy barrier of the Sb-for-As exchange reaction on the two surfaces. These findings demonstrate a simple processing-free route to compositional engineering at the monolayer level along NWs.

Keywords: STM, GaAs, Sb-for-As exchange, Surface alloy, wurtzite

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Nanowires (NWs) with controlled axial stacking of different crystal structures enable welldefined surface science experiments. Recently, a lot of attention has been devoted to not only grow NWs with pure wurtzite (Wz) or zinc blende (Zb) structure1, but also to switch between the two within single NWs2-5. Due to the differences in crystal stacking, such structures also experience variations in the exposed surface facets6, 7. It is common for the Zb phase to expose 110 facets (at the growth settings used here), whereas the Wz typically exposes either 1120 or 1010 facets (110 or 112 zinc blende equivalents, respectively). From a surface science viewpoint, the exposure of both Wz and Zb facets in the same nanocrystal presents an excellent opportunity to study preferential surface adsorption and diffusion on different semiconductor bulk crystal structures. Here we can isolate the effect of crystal structure from other material variations and since all phases are present in the same nanocrystal they can be studied under the same temperature and gas conditions. In addition, since these facets can be controllably generated they represent a clear opportunity for tailoring surface and interface structures both radially and axially down to the limit of single layers.

GaAs nanowire structures have been studied intensely and recently several interesting novel structures have been realized with applications in photonics and photovoltaics8-10. The combination of GaAs and GaSb presents several interesting opportunities due to the type-II band alignment11 and allows for the development of infrared photodetectors12, long wavelength lasers13-15, and charge based memories16. However, the growth of sharp GaAs/GaSb interfaces is challenged by strong anion segregation effects, such as Sb-for-As exchange reactions17-19, although some progress for more advanced heterostructures using crystal phase engineering has been achieved20. The incorporation of Sb atoms in epitaxially grown Zb GaAs via Sb-for-As exchange is a well-established procedure, also called “Sb soaking”. During that step, a GaAs 2 ACS Paragon Plus Environment

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growth surface is exposed to a flux of Sb atoms, without the presence of Ga atoms, at increased temperatures, typically for several seconds to minutes. This procedure was found to be selflimiting, where a maximum of about 1 to 2 monolayers of GaAs were replaced by GaSb, as confirmed by reflection high-energy electron diffraction21, X-ray diffraction11, or (in the case of Sb soaking of InAs) by X-ray photoemission spectroscopy18. This progress apart, new routes are needed for efficient and easy creation of GaAs/GaSb heterostructures with down to atomic precision, where the nanowire surfaces might represent an interesting route with its special crystal structures and morphology.

The scanning tunneling microscope (STM) has in recent years become applicable to III-V NW studies giving access to atomic level structural and electronic properties6, 7, 22. STM allows direct studies of the NW surfaces, which inherently (due to the large surface-to-volume ratio) are a necessity to control for tuning of electronic and optical properties and even predict the growth of entire NWs23. STM has also been used for studies of Sb overlayers on GaAs(110) surfaces

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as

well as cross-sectional STM (XSTM) studies of the group-V exchange mechanism between Sb and As atoms in GaAs19, 25 and in GaSb26, 27. However, both the surface and the cross-sectional STM studies on bulk crystals can only access Zb facets, since these are the thermodynamically stable facets for bulk GaAs. STM has also been used to study Sb containing III–V NWs22, 28, 29 where the Sb was incorporated during growth. No studies have been presented showing controlled Sb surface modification on NWs after growth. However, such controlled incorporation (followed by further radial overgrowth to protect the Sb containing layer from oxidation) would be of large interest as a mean to create well-defined, low-dimensional heterostructures.


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In this letter, we show atomic scale imaging of crystal structure induced preferential surface alloying of Sb in III–V NWs. Using STM, we probe the atomically sharp interface between axial segments of Wz and Zb structure on NWs cleaned in atomic hydrogen. By diffusion of Sb4 from an adjacent elemental source, we can explore surface diffusion of Sb on the GaAs NWs. The Sbfor-As exchange reaction occurs on all surface facets. However, we find a factor 4 higher incorporation rate into the Zb surfaces compared to the Wz counterparts. This higher incorporation is investigated by ab-initio DFT calculations and it is found to be due to a lower energy barrier for the site exchange of an As with an Sb atom on Zb GaAs (110) compared to Wz GaAs1120.

Heterostructured GaAs NWs with a Wz bottom and a Zb top segment with designed Wz inclusions were prepared by metal-organic vapor phase epitaxy (MOVPE) in an 3x2” Aixtron close coupled showerhead reactor (CCS) following the particle-assisted growth mode and the use of Au particles. The latter were deposited onto GaAs111 substrates by aerosol technique30 with a total areal density of 1 µm-2 and nominal diameters of 30, 50, and 70 nm. The NWs were grown at a set temperature of 550 °C with trimethylgallium (TMGa) and arsine (AsH3) as precursor materials at a total reactor flow of 8 slm, and a total reactor pressure of 100 mbar. The molar fractions were set to χTMGa = 1.9x10-5 and χAsH3 = 4.5x10-5 and 4.4x10-3 for Wz and Zb conditions, respectively2, 3. In order to remove surface oxides and allow proper substrate preconditioning, a 10 min annealing step was carried out prior to growth at about 630 °C in an AsH3/H2 atmosphere. After that step the temperature was reduced to the growth temperature, and after thermal stabilization the precursors were introduced to initiate growth. For detailed information about the growth of sharp crystal structure interfaces in III–V NW systems, see Ref. 2 and references therein. Scanning electron microscopy (SEM) characterization was carried out 4 ACS Paragon Plus Environment

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in a ZEISS Leo Gemini 1560 setup. For structural characterization, the NWs were placed on copper grids covered with a lacey carbon layer and investigated in a JEOL-3000F transmission electron microscope (TEM).

The NWs were transferred to an n-type epi-ready GaAs111 substrate by mechanical break off and loaded into ultra-high vacuum (UHV, with a base pressure of 1000 nm2). On the other hand, Figure 3(b) depicts a part of the Wz stem of the NWs, showing both a 1120-type facet, but also a small 1010-type facet. Both facets show surface incorporated Sb-atoms at similar area densities of about 0.16±0.07 Sbatoms/nm2 (measured in images covering >1000 nm2). This significant difference in Sb incorporation is notable and its origins will be discussed in detail below. By performing


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autocorrelation analysis, it is possible to study the placement of the Sb-atoms

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. In short,

autocorrelation of the positions of the Sb-atoms shows the available vectors between the Sbatoms. Through such an analysis, it is e.g. possible to determine whether there are certain positions that the Sb-atoms do not occupy such as neighbouring sites. For both the 110 and 1120 facets, we found random placement of the Sb-atoms on As-sites and that all vectors between Sb-atoms could be found, i.e. all sites – even neighbouring – can be occupied. This observation is compatible with the idea that the species that trigger the exchange are highly mobile single Sb atoms. While each impact of an Sb4 molecule could produce up to four Sb adatoms, no spatial correlation between the point of impact and the sites of the exchange reaction could be observed. The random placement of Sb atoms found here is different to what was recently observed on InAsSb NWs28. In that study, it was found that Sb atoms incorporated during MBE growth into Zb{110} surfaces would not occupy neighboring positions along the atomic chains, but rather give rise to short range ordering along certain crystallographic directions. We observe neighboring Sb-Sb atoms, see Figure 2 and 3, which further corroborates the idea that the Sb incorporation observed here is a stochastic process based on adatom exchange, and not step-flow growth (as in the other study). The lack of locally ordered phases is important when scaling down to smaller diameter NWs since local differences in stoichiometry might act as scattering centers for electron transport In order to better understand the thermodynamics and kinetics of the Sb dissociation, diffusion, and the Sb-for-As exchange reaction, we developed a five-step model based on DFT calculations. The model comprises a three-step creation of Sb adatoms (Sbads), the exchange reaction itself, and eventually the removal of the As atoms resulting from the exchange. The reaction enthalpies (at zero temperature and pressure) for the relevant steps are summarized in


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Table 1. Table 1: Calculated reaction enthalpies in eV. Negative numbers indicate endothermic reactions. Surface

GaAs1120

GaAs(110)

DFT functional

LDA

PBE

LDA

PBE

Sb4 → 2(Sb2)ads

-2.01

-2.60

-1.91

-2.44

(Sb2)ads → 2Sbads

0.42

-0.14

-0.42

-0.87

Sbads + Assurf → Sbsurf + Asads

-0.31

-0.24

-0.28

-0.11

Asads + Sbads→ AsSb

0.04

0.00

0.51

0.43

As a first step, we assume that the Sb4 molecule breaks apart into two adsorbed Sb2 species upon arrival on the surface. These species are anchored in the troughs of the relaxed surface and stand upright. This first step is endothermic by about 1.9 eV to 2.6 eV, which means that only a small fraction of the encounters of the Sb4 molecules with the surface will lead to the reaction. In a second step, the adsorbed Sb2 species break apart into two Sb adatoms (Sbads). Again, this step is mostly endothermic, but the energy required is clearly smaller than in the first step. Here the first significant difference between Zb GaAs(110) surface compared to Wz GaAs1120 is observed as the Sbads is bound more strongly on the Zb surface and as a result the energy needed to dissociate the (Sb2)ads species is lower. When using the PBE functional, the binding of the final state, the Sb adatom, to the surface is calculated to be somewhat weaker than in the LDA calculation. The LDA functional seems to overestimate the binding of adatoms to the surface. However, the energy trends using the two different functionals are similar when comparing the Zb and Wz surfaces. The consequence of the energy difference of this second step could be a difference in concentration of Sbads on the two facets, however, to evaluate this the third step needs to be taken into account. In the third step, the Sb adatoms diffuse on the surface. According to our DFT calculations, Sb 13 ACS Paragon Plus Environment

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diffusion takes place preferentially along troughs, with a diffusion barrier of about 0.2 eV on the Wz surface and 0.26 eV on the Zb surface. This is much lower than the Asads diffusion barrier. The As adatoms were found to diffuse on Wz on top of the Ga-As-Ga-As chains with a barrier of 0.64 eV47 along these chains. The relative ease of Sb diffusion ensures that the Sb adatoms, once created, are available with equal concentration everywhere on the surface, despite of a higher (Sb2)ads dissociation rate on Zb GaAs(110). Moreover, we may conclude that diffusing Sb adatoms will move too fast to be resolved in STM images. The fourth, most important step comprises the exchange of an Sb adatom, located in the trough, with an As atom being part of the original substrate surface. The energy barriers for this process, both on the Zb GaAs(110) and the Wz GaAs1120 surface, are shown in Figure 4. Also this reaction step is slightly endothermic as seen in Table 1. Since the Sb atom is bulkier than the As adatom, it does not fit as well as As into the bonding environment on the surface. In addition to reaction enthalpies, the reaction barriers are also relevant in determining the exchange reaction. The reaction barrier, determined with the PBE functional for the exchange process is 0.70 eV on the Zb GaAs(110) surface, and 0.97 eV on the Wz GaAs1120 surface. The LDA results (0.81 and 1.05 eV, respectively) confirm the trend found in PBE. The difference in barrier height on the two surfaces, even if Sbads is available in the same concentration everywhere, will result in a significant difference in incorporation rate. If pre-exponential factors are assumed to be the same, we estimate the ratio of the rates for the exchange reaction to be about 72 at T= 723K. This gives the right trend, but is much larger than the experimental ratio of about 4, pointing to some influence of other kinetic factors. One such factor that will indirectly influence this difference is the fifth step of the As adsorbate removal which will be discussed below. It is noteworthy that, despite similar structural motifs (Ga-As-Ga-As zig-zag chains of three-fold


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coordinated surface atoms), the transition state for the exchange reaction, and hence the barrier, are different on Zb and Wz surfaces. The zig-zag chains on either surface are running in different crystallographic directions. More importantly, the chains on the GaAs(110) surface have constant bond angles (110°), leaving room for the attack of the Sb atom, whereas the chains on the GaAs1120 surface have alternating in-plane bond angles of 120° at the Ga atoms, but only 98° at the As atoms. Finally, as a fifth step, the reaction mechanism must answer the question what happens to the exchanged As atoms. Since their diffusion barrier is rather high and they are then quite immobile, it is unlikely that they will find another Asads as reaction partner to form As2 and subsequently desorb. We find it more plausible that one of the fast diffusing Sbads adatom reacts with Asads to form a mixed dimer AsSb, which then desorbs into the gas phase. Desorption of those species would also be promoted by the UHV environment used in our experiments. As our DFT calculations show (see Table 1), this final step is thermoneutral on GaAs(110), and even exothermic on GaAs1120. Therefore, this final step is expected to occur with a much higher rate than the previous ones. This could explain why no As adatoms produced by the exchange reaction are detectable in the STM images.


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Figure 4. Calculated diffusion paths for an Sb adatom on Zb/Wz surfaces (a) View of the ZB GaAs(110) surface. A zig-zag chain of Ga atoms (green) and As atoms (blue) forms the top layer. Bottom left: The Sb adatom (brown) is located in its most stable adsorption site in the trench between zig-zag rows. Bottom right: The Sb atom inserts itself into the zig-zag row, pushing an As atom out. The barrier, measured from the Sb adatom state, is 0.70 eV in PBE (0.81 eV in LDA). The energy scale is taken relative to the ideal GaAs(110) surface + ½ Sb2 molecule. (b) View of the Wz GaAs1120 surface. A zig-zag chain of Ga atoms (green) and As atoms (blue) forms the top layer. Bottom left: The Sb adatom (brown) is located in its most stable adsorption site in the trench between zig-zag rows. Top right: The Sb atom inserts itself into the zig-zag row, pushing an As atom out. The barrier, measured from the Sb adatom state, is 0.97 eV in PBE (1.05 eV in LDA). The energy scale is taken relative to the ideal GaAs1120surface + ½ Sb2 molecule. Solid line is calculated using LDA, whereas the dashed line is calculated using PBE. The fast As removal also plays a role for the concentration of substituted Sb atoms. From the potential diagram in Figure 4 one can realize that the reverse exchange reaction of As-for-Sb might also happen, and in fact the barrier is smaller than for the Sb-for-As reaction. However, if the Asads is removed, this reaction becomes impossible (it needs an Asads) and we are now left with substituted Sb atoms. Still as the barrier for the As-for-Sb substitution is lowest on GaAs(110), the reaction is more likely to occur on this surface before the Asads atom is removed by a surface Sb atom: This would lower the number of substituted Sb atoms that we observe. As a result, the actual overweight of Sb incorporation into the GaAs(110) will be smaller than predicted by the pure Sb-for-As substitution barrier. We also consider the possibility of interstitial Sb atoms, which gives interesting insights into the different behavior of the two surfaces. It is possible to place the Sb atom at an interstitial site 16 ACS Paragon Plus Environment

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below the surface. In this case, different behavior is observed for Zb and Wz respectively: Below the Zb surface, the Sb interstitial is quite high in energy, but the interstitial cage is a local energy minimum. The Sb interstitial atom has an energy that is 2 to 2.5 eV higher than the energy of the Sb adatom, but it stays inside. In the Wz slab, the subsurface interstitial Sb atom is unstable. It always pushes out an As atom or a Ga atom and replaces it in the lattice. The substituted atom is pushed to the surface and becomes an As adatom or Ga adatom. The final state reached after that is rather stable, comparable to an Sb atom substituting a surface As atom in the top layer. It appears that it is easier to push an atom out of the wurtzite surface. This could be because of the geometrical structure, there is more space to move because the 1120 surface has channels along the surface normal. Theoretically it has been shown that there is a difference of As diffusion in Zb and Wz GaAs. The diffusion of an As atom within Wz ab planes is easier than diffusion along the c axis48. Moreover, it has been found that Sb interstitials also play a significant role in the differences between Ga and Sb diffusion in GaSb49. To summarize, while confirming the experimental observations, the theoretical calculations present a more comprehensive picture. From the calculations it appears that for Zb GaAs(110) the Sb incorporates into the top-most surface layer. Substitution in deeper layers through interstitials is not likely, see Table 2. For the Wz GaAs, Sb incorporation in the surface is less likely (but still much more probable than bulk substitution) while there is a small probability for substitution into subsurface layers through interstitial diffusion. Table 2: Formation energy (in eV) for a substitutional Sb atom on As lattice site as a function of depth below the surface. Values are calculated in LDA; the chemical potentials µAs and µSb are taken to be – ½ Ebind(As2) and – ½ Ebind(Sb2), respectively. bilayer#

GaAs(110) 17 ACS Paragon Plus Environment

GaAs1120

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0.15

0.14

2

0.38

0.31

3

0.35

0.30

4

0.39

0.34

While the crystal structure dependent Sb surface incorporation in itself is an interesting fundamental observation, it is worth considering what would happen during further overgrowth with e.g. GaAs - to protect the structure upon exposure to air. Since the As-for-Sb reaction is energetically favorable (the reverse reaction as described in Figure 4), and the barrier for this exchange is slightly lower than the one for Sb-for-As, some of the Sb might be removed from the initial surface layer again and then either be incorporated in higher lying layers, evaporate, or even float at the surface upon further overgrowth if As adsorbates become available. Sb is known for the latter behavior, acting as a surfactant during epitaxial growth, which means that Sb atoms tend to float at the growth surface and influence crystal growth, without being incorporated45, 50. However, as the exchange reactions are kinetically limited, these effects might be controlled by tuning the temperature and Ga to As ratio during overgrowth. Indeed, studies of overgrowth of GaSb quantum dots in GaAs19, 51-53 have shown that careful control of growth conditions and strain determine the preservation of the quantum dots or their transformation into ring-like structures due to significant Sb segregation. In the present case we have seen that the formation of a single layer of GaAsxSb1-x is possible on the nanowire surfaces with very different x depending on the crystal type. Overgrowing these layers would then form GaAsxSb1-x rings which would be axially tailored following the crystal structure, reaching from extended, through atomically thin, shell structures to ring-shaped single atomic chains, or even arrays of equally spaced rings along the nanowire axis. Such structures could be used to confine charge carriers in complex and well-defined geometries, creating novel electronic phases.


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In conclusion, we have shown that Sb incorporates differently in Zb and Wz GaAs NW surfaces. GaAs NWs were grown to exhibit Wz segments of varying lengths in a Zb matrix. Atomic scale images were obtained from the NWs showing unreconstructed 110-, 1120-, and 1010surfaces. After Sb-exposure, Zb segments showed an approximately 4 times higher incorporation rate of Sb. By modelling the Sb incorporation using DFT calculations, it was possible to conclude that the higher Sb incorporation at the Zb segments was related to a lower energy barrier of the Sb-for-As exchange reaction at that surface. We have developed a five-step model for the incorporation of Sb into the As sublattice, that can qualitatively explain our results. As the theoretical model predicts a larger difference in Sb incorporation this would indicate that additional kinetic factors such as the back reaction of As-for-Sb and As-Sb adsorbate recombination play a role. Still the experimental and theoretical data agree that the surface concentration of Sb can be strongly varied by diffusion across a surface with variable crystal structure and that this works even for segments of a few atomic layers. The findings indicate a potential way to modify the surface in a controllable way for the creation of atomically thin GaAs1-xSbx layers on GaAs NWs. This allows for processing-free opportunities in the engineering of advanced semiconductor heterostructures. While this is the first observation of such crystal dependent alloying on semiconductors, we believe that this type of surface modification is not limited to Sb alloying into GaAs, but should be relevant also for other elements and compounds. SUPPORTING INFORMATION AVAILABLE Additional images Figures S1–S3. ACKNOWLEDGEMENTS


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This work was supported by the Swedish Research Council (VR), the Swedish Foundation for Strategic Research (SSF), the Swedish energy agency, the Crafoord Foundation, the Knut and Alice Wallenberg Foundation (KAW), and the European Research Council under Grant Agreement No. 259141. M.H. gratefully acknowledges financial support from the Royal Physiographic Society in Lund as well as a postdoctoral fellowship from KAW. Experiments at the University of California, Santa Barbara, were supported in part by the Office of Naval Research (Grant number N00014-13-1-0660) and the MRSEC Program of the National Science Foundation through DMR-1121053. REFERENCES 1. Joyce, H. J.; Wong-Leung, J.; Gao, Q.; Tan, H. H.; Jagadish, C. Nano Letters 2010, 10, (3), 908-915. 2. Lehmann, S.; Wallentin, J.; Jacobsson, D.; Deppert, K.; Dick, K. A. Nano Letters 2013, 13, (9), 4099–4105. 3. Lehmann, S.; Jacobsson, D.; Deppert, K.; Dick, K. Nano Research 2012, 5, 470-476. 4. Caroff, P.; Dick, K. A.; Johansson, J.; Messing, M. E.; Deppert, K.; Samuelson, L. Nature Nanotechnology 2009, 4, (1), 50-55. 5. Algra, R. E.; Verheijen, M. A.; Borgstrom, M. T.; Feiner, L.-F.; Immink, G.; van Enckevort, W. J. P.; Vlieg, E.; Bakkers, E. P. A. M. Nature 2008, 456, (7220), 369-372. 6. Hjort, M.; Lehmann, S.; Knutsson, J.; Timm, R.; Jacobsson, D.; Lundgren, E.; Dick, K. A.; Mikkelsen, A. Nano Letters 2013, 13, (9), 4492–4498. 7. Hjort, M.; Lehmann, S.; Knutsson, J.; Zakharov, A. A.; Du, Y. A.; Sakong, S.; Timm, R.; Nylund, G.; Lundgren, E.; Kratzer, P.; Dick, K. A.; Mikkelsen, A. Acs Nano 2014, 8, (12), 12346-12355. 8. Heiss, M.; Fontana, Y.; Gustafsson, A.; Wüst, G.; Magen, C.; O’Regan, D. D.; Luo, J. W.; Ketterer, B.; Conesa-Boj, S.; Kuhlmann, A. V.; Houel, J.; Russo-Averchi, E.; Morante, J. R.; Cantoni, M.; Marzari, N.; Arbiol, J.; Zunger, A.; Warburton, R. J.; Fontcuberta i Morral, A. Nature Materials 2013, 12, 439-444. 9. Heurlin, M.; Magnusson, M. H.; Lindgren, D.; Ek, M.; Wallenberg, L. R.; Deppert, K.; Samuelson, L. Nature 2012, 492, (7427), 90-94. 10. Mariani, G.; Scofield, A. C.; Hung, C.-H.; Huffaker, D. L. Nat Commun 2013, 4, 1497. 11. Ledentsov, N. N.; Böhrer, J.; Beer, M.; Heinrichsdorff, F.; Grundmann, M.; Bimberg, D.; Ivanov, S. V.; Meltser, B. Y.; Shaposhnikov, S. V.; Yassievich, I. N.; Faleev, N. N.; Kop’ev, P. S.; Alferov, Z. I. Physical Review B 1995, 52, (19), 14058-14066. 12. Carrington, P. J.; Wagener, M. C.; Botha, J. R.; Sanchez, A. M.; Krier, A. Applied Physics Letters 2012, 101, (23), 231101. 13. Quochi, F.; Kilper, D. C.; Cunningham, J. E.; Dinu, M.; Shah, J. Photonics Technology Letters, IEEE 2001, 13, (9), 921-923. 20 ACS Paragon Plus Environment

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