Origin of Spontaneous CoreâShell AlGaAs Nanowires Grown by

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Origin of spontaneous core-shell AlGaAs nanowires grown by molecular beam epitaxy V. G. Dubrovskii, I. V. Shtrom, R. R. Reznik, Yu. B. Samsonenko, A. I. Khrebtov, I. P. Soshnikov, S. Rouvimov, N. Akopian, T. Kasama, and G. E Cirlin Cryst. Growth Des., Just Accepted Manuscript • DOI: 10.1021/acs.cgd.6b01412 • Publication Date (Web): 12 Oct 2016 Downloaded from http://pubs.acs.org on October 17, 2016

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Crystal Growth & Design

Origin of spontaneous core-shell AlGaAs nanowires grown by molecular beam epitaxy V. G. Dubrovskii1,2,3,*, I. V. Shtrom1,2, R. R. Reznik1,4, Yu. B. Samsonenko1,5, A. I. Khrebtov1, I. P. Soshnikov1,2,5, S. Rouvimov6, N. Akopyan7, T. Kasama8, G. E. Cirlin1,3,5 1 2

St. Petersburg Academic University, Khlopina 8/3, 194021 St. Petersburg, Russia

Ioffe Physical Technical Institute of the Russian Academy of Sciences, Politekhnicheskaya 26, 194021 St. Petersburg, Russia 3 4

ITMO University, Kronverkskiy pr. 49, 197101 St. Petersburg, Russia

Peter the Great St. Petersburg Polytechnic University, Polytechnicheskaya 29, 195251, StPetersburg, Russia 5

Institute for Analytical Instrumentation RAS, Rizhsky 26, 190103, St-Petersburg, Russia 6

7

University of Notre Dame, Notre Dame, IN 46556 USA

Department of Photonics Engineering, Technical University of Denmark, 2800 Kongens Lyngby, Denmark

8

Center for Electron Nanoscopy, Technical University of Denmark, 2800 Kongens Lyngby, Denmark

*corresponding author, E-mail: [email protected]

ABSTRACT: Based on the high-angle annular dark-field scanning transmission electron microscopy and energy dispersive X-ray spectroscopy studies, we unravel the origin of spontaneous core-shell AlGaAs nanowires grown by gold-assisted molecular beam epitaxy. Our AlGaAs nanowires have a cylindrical core and a tapered shell. The composition of the shell is close to nominal, while the aluminum content in the core is systematically smaller than nominal. After switching off the group III fluxes, the aluminum content in the droplet and in the topmost part of the nanowire rapidly tends to zero, while gallium remains there at a high percentage. We present a quantitative model to explain these findings. Lower aluminum composition in the core is attributed to its lower surface diffusivity, with the aluminum collection length of 250 nm against 780 nm for gallium at the substrate temperature 510 oC and under the nominal aluminum content of 0.2. These values decrease to 8 nm and 160 nm when the nominal aluminum content is raised to 0.6. On the other hand, aluminum leaves the droplet at least 100 times faster than gallium, with a typical bonding rate with arsenic of the order of 1000 nm/s. 1 ACS Paragon Plus Environment


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INTRODUCTION

Ternary III-V nanowires (NWs) and NW heterostructures are promising building blocks for scalable bottom-up photonic devices which can be integrated with silicon electronic platform1,2. In particular, core-shell NWs based on AlGaAs material are of great interest for high-speed lasers3,4, single photon sources5,6 and THz emitters7. Core-shell GaAs/AlGaAs NW structures have been fabricated by either the vapor-liquid-solid (VLS)3,4,8 or selective area epitaxy9 techniques where the GaAs cores and AlGaAs shells are deliberately separated by changing the deposition conditions (for example, temperature) and vapor fluxes. Alternatively, AlGaAs heterostructures can be obtained via the gold-assisted VLS growth with simultaneous deposition of gallium and aluminum. In this case, self-separation of aluminum composition in the NW core and shell occurs spontaneously during growth without changing the vapor environment. For example, spontaneous AlxGa1-xAs core-shell NWs with different compositions x in the cores and shells were grown in Ref. [10] by molecular beam epitaxy (MBE) and in Ref. [11] by metalorganic chemical vapor deposition. Self-separation of ternary III-V materials into core-shell structures with different composition is not specific for AlGaAs and has also been reported for InGaAs12,13, GaAsP14, InAsSb15 and GaAsSb16 NWs, where the core and shell compositions are highly dependent on the growth conditions and the substrate used. Clearly, the growth procedure for spontaneous core-shell III-V NWs is much simpler than the deliberate separation of the core and the shell. On the other hand, the mechanisms that govern such self-separation and determine the inhomogeneous NW composition remain unclear to this end. Furthermore, previously published results are often qualitatively different – for example, in Ref. [10] the aluminum content in the core was found to be lower than in the shell, while in Ref. [11] this tendency was inverted. Stationary composition of gold-assisted ternary III-V NWs versus different growth parameters and droplet size has recently been thoroughly studied in the case of InGaAs (see Ref. [17] and references therein), where there is no spontaneous formation of distinct core-shell structures. In Ref. [18], aluminum was found to 2 ACS Paragon Plus Environment

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crystallize much faster than gallium at the liquid-solid interface of self-catalyzed axial AlGaAs NW heterostructures. Overall, understanding stationary and non-stationary compositional evolution in ternary AlGaAs NWs is highly desired for the controlled fabrication of photonic NW structures with pre-defined properties. Consequently, in this work we closely examine the compositions of AlGaAs NWs obtained by gold-assisted MBE on Si(111) and present a quantitative model that unravels the origin of their inhomogeneous spatial composition along and across the growth axis. Fitting the data by a quantitative model allows us to deduce some important kinetic parameters such as the collection lengths of aluminum and gallium and the bonding rates of Al-As and Ga-As pairs.

EXPERIMENTAL DETAILS The samples were grown by MBE using Riber 21 setup equipped with a separate vacuum

chamber to deposit gold. After HF treatment in water solution (1:10), Si(111) substrates were loaded into the vacuum chamber and outgassed at 850 oC before the gold deposition at 550 oC (with 0.1 nm effective thickness for the nominal aluminum composition z = 0.2 and 0.3 and 0.2 nm for other z ). The substrates were kept at 550 oC for 1 min after the gold deposition to improve the droplet size homogeneity and subsequently cooled down to room temperature. After that, the substrates were transferred to the main growth chamber with no vacuum brake. For all NW growth runs, the substrate temperature was set at 510 oC, the equivalent AlGaAs growth rate was fixed at 0.3 nm/s and the total V/III flux ratio was 3.

Five different nominal Al

compositions were used, z = 0.2, 0.3, 0.4, 0.5 and 0.6, as calibrated using reflection high energy diffraction (RHEED) oscillation technique on standard GaAs(001) substrates. The appearance of pronounced wurtzite-type RHEED patterns was documented after ~ 20 s of the growth in most cases. The total growth time for all samples was 20 min.

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Figure 1. SEM images of AlGaAs NWs grown by Au-assisted MBE with 0.3 (a) and 0.6 (b) nominal aluminum content.

Typical scanning electron microscopy (SEM) images of AlGaAs NWs grown with z = 0.3 and 0.6 are shown in Fig. 1. The maximum NW length is drastically decreased form ~ 4000 nm at z = 0.2 to only about 1000 nm at z = 0.6. This should be due to a higher aluminum content which suppresses surface diffusivity of metal adatoms on the NW sidewalls19,20. Consequently, the increase of nominal aluminum composition yields more tapered NWs toward higher z . At the same time, the diameter of the NWs tops remains constant (15 - 25 nm) in all cases. Quasi-two dimensional AlGaAs layer is also grown on Si(111) substrate. Structural properties of AlGaAs NWs were studied by transmission electron microscopy (TEM). The NWs were found to have rather small density of structural defects, such as twins or stacking faults, and exhibit wurtzite crystal structure, consistent with our RHEED observations. For the low nominal aluminum content from 0.2 to 0.3, we did not detect any noticeable number of defects along ten NWs studied. As the aluminum content increases, the defect density increases too, reaching 10 µm-1 at z =0.6. More defects are located in the bottom part of the NWs. To identify the compositions of selected NWs, high-angle annular dark-field imaging (HAADF) and energy dispersive X-ray spectroscopy (EDX) were performed in the Scanning TEM mode using an FEI Titan 80-300ST TEMs operated at 120 keV. 4 ACS Paragon Plus Environment

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RESULTS AND DISCUSSION In Fig. 2, we show a typical TEM image of the AlGaAs NW grown at z = 0.2, along with

the height-dependent compositional profile obtained by EDX (not accounting for the relative composition of the core and the shell). Figure 3 (a) shows a typical STEM-EDX lateral crosssection of another NW grown with much higher z = 0.6, unambiguously revealing the spontaneous core-shell structure with a higher aluminum composition in the shell. We note that the central part of the NW (core) shown in the Fig. 3 (a) has higher Ga content in comparison to the given numbers because the elements profiles count both core and shell parts. The core-shell structure is also clearly seen in the TEM image of Fig. 3 (b). Figure 4 shows the average core (which were taken just under the droplet) and shell compositions (taken just from thick shell part at 500 nm apart of the droplet) of the NWs grown with different z from 0.2 to 0.6. These data reveal three specific features that were repeatable for different NWs: (i) Average aluminum content in the core is always lower than in vapor ( x < z ), while the shell composition is close to z (see Fig. 4). (ii) There is practically no aluminum left in the droplet after growth, while the gallium content remains significant (~ 0.38 in the case shown in Fig. 2); (iii) The aluminum content in the topmost NW part (~ 50 nm below the droplet in Fig. 2) is also close to zero. Over only 30 nm height, it changes rapidly to a value which is close to its composition in the core ( x ≅ 0.11 in Figs. 2 and 3). Far enough from the droplet, the average aluminum content is close to that of the shell because the shell is much thicker than the core. Gallium-rich Au-Ga alloys after growth (with the gallium concentrations as high as 0.50.7) were reported earlier in the case of Au-assisted growth of binary GaAs NWs21,22. The observed decrease of aluminum concentration toward the NW top and the absence of aluminum in the droplet are consistent with the findings of Ref. [10]. Based on the data of Ref. [18], we may assume that aluminum leaves the droplet much faster than gallium upon the termination of both group III materials inputs under the arsenic flux, which explains its absence in the droplet 5 ACS Paragon Plus Environment


and in a short NW section below the droplet. On the other hand, aluminum should be less diffusive on the NW sidewalls19,20 and hence enter the droplet slower than gallium, incorporating more into the vapor-solid shell than by the VLS process into the core.

Figure 2. EDX profiles of the average gold, gallium and aluminum concentrations in the liquid and solid phases versus the distance from the droplet, with a typical TEM image of an AlGaAs NW having nominal

z = 0.2 shown in the insert.

As

Concentration

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Al Ga Position (a)

(b)

Figure 3. (a) STEM-EDX lateral x-section of a NW grown with z = 0.6. The core has higher Ga while the shell has higher Al content. Measured values for Al and Ga in a central part of the wire composed from the contributions from core, upper and lower shell parts. (b) TEM image revealing the spontaneous core-shell structure of the NW where the aluminum-rich shell appears brighter. 6 ACS Paragon Plus Environment

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Figure 4. Average aluminum contents in the cores and shells of AlGaAs NWs versus the aluminum content in vapor (symbols). The data for the NW cores are fitted by Eq. (5) with γ = 2.6. The aluminum composition in the shells is close to the nominal one for most data points (straight line x = z ).

To quantify these data, we use the model of Ref. [23] for a non-stationary quaternary Au Al-Ga-As liquid alloy, illustrated in Fig. 5. The atomic concentrations of aluminum, gallium and arsenic within the regular VLS growth concept are described by the kinetic equations dci α = (Vi − U i ci − K i c As ci ) , i = Al , Ga ; dt R dc As α = [V As − U As c As − c As ( K Al c Al + K Ga cGa )] , dt R

(1)

Here, Vk and U k for k = Al, Ga and As are the kinetic coefficients (measured in nm/s) that determine the material influxes and outgoing fluxes, respectively, for each growth species. For gallium and aluminum, the Vi terms include surface diffusion of adatoms and the U i ci are the reverse diffusion fluxes from the droplet onto the NW sidewalls (desorption of both group III species is negligible at 510 oC). The VAs and U As c As stand for the adsorption and desorption rates of arsenic and do not contain any diffusion-induced contributions24,25. The Ki c As ci for i = Al and Ga are the crystallization rates of Al-As and Ga-As pairs at the growth interface, with K i as the



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corresponding bonding rates. The α = 3Ω L /[ f ( β )Ω S ] is the geometrical factor, with Ω L as the elementary volume per atom in the liquid phase, Ω S = 0.0452 nm3 as the volume of a GaAs pair in the solid phase and f ( β ) = (1 − cos β )(2 + cos β ) /(1 + cos β ) sin β as the geometrical function relating the volume of spherical cap to the radius of its base through the contact angle of the droplet β . The axial grow rate of the NW core in the VLS process is given by G = c As ( K Al c Al + K Ga cGa ) .

(2)

The aluminum composition is defined as

x=

K Al c Al . K Al c Al + KGa cGa

(3)

Figure 5. VLS growth of AlxGa1-xAs core. Aluminum and gallium enter the droplet at the rates

Vi = vi [sin−2 β + (2 sinθ / π )(λi / R)] for i = Al, Ga due to the direct fluxes vi and the surface diffusion of adatoms with the collection lengths λi . They leave the droplet only by crystallization at the liquidsolid interface with the bonding rates K Al c As c Al and K Ga c As cGa . The arsenic influx equals V As and its desorption rate is U As c As . The group III fluxes are directed at the angle θ to the substrate normal. Contact angle of the droplet equals β . We want to find the solid composition of the core x as a function of the vapor composition z in stationary regime and after the group III fluxes are switched off.


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Our first goal is to find the stationary aluminum composition in the core. We note that, whenever the reverse diffusion fluxes are neglected ( U Al = U Ga = 0 ), the stationary solutions for the aluminum and gallium concentrations become ci( s ) = Vi /( Ki c As ) and hence the stationary composition equals26

x=

VAl VAl + VGa

(4)

regardless of the arsenic concentration and the K i values. This result is anticipated, because without any outgoing fluxes the solid composition should be entirely determined by the material transports into the droplet (it does not mean, however, that the liquid and solid compositions are identical, as suggested in Ref. [17]). In our MBE system, the aluminum and gallium beams have the same inclination angle θ = 30 deg with respect to the substrate normal. For β ≥ π / 2 + α (Ref. [27]), the simplest expression for Vi writes19 Vi = vi [sin−2 β + (2 sinθ / π )(λi / R)] , with vi as the vapor fluxes and λi as the effective collection lengths of adatoms on the NW sidewalls to the droplet for i = Al, Ga. The nominal composition is determined by the vapor fluxes according to

z = v Al /(v Al + vGa ) . Using these expressions in Eq. (4), we get x=

πR + 2λGa sin θ sin 2 β z ; γ = . (1 − γ ) z + γ πR + 2λ Al sin θ sin 2 β

(5)

This result has the same form as in Ref. [17] but is re-arranged for the MBE growth. The line in Fig. 4 show a good fit to the measured core compositions by Eq. (5) with the z − independent γ = 2.6. In Ref. [17], the indium composition of gold-catalyzed VLS InxGa1-xAs NWs was found higher than in vapor due to a larger collection length of indium with respect to gallium. The situation is reversed for gold-catalyzed AlxGa1-xAs NWs where the aluminum composition in the VLS cores is lower than in vapor because of a smaller collection length of aluminum with respect to gallium.



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Assuming a typical contact angle of the droplet β = 120 deg, at θ = 30 deg and for the average NW core radius R = 20 nm, Eq. (5) at γ = 2.6 yields a linear relationship

λGa = 2.6λ Al + 135 nm for any z , meaning that the λGa / λ Al ratio is larger than 2.6 and the λGa is larger than 135 nm for all the data points. From Eq. (1) for arsenic, its stationary concentration ( s) equals c As = VAs / U As when VAs >> G , and does not depend on the group III concentrations.

Under such arsenic-rich conditions, which can be enhanced by re-emission of the arsenic species24, the axial NW growth rate equals G = VGa + VAl , corresponding to the standard group III limited

regime19,20,25.

This

expression

can

be

re-arranged

as

G = v[sin−2 β + (2 sinθ / π )( zλAl + (1 − z)λGa )] for our Vi , with the total flux of the group III atoms v = v Al + vGa of 0.36 nm/s corresponding to a fixed equivalent AlGaAs growth rate of 0.3 nm/s at

θ = 30 deg. The maximum NW growth rate G of 4.17 nm/s is observed for z = 0.2 . This value is estimated as the length of the longest NWs (5000 nm) divided by the total growth time of 20 min. The maximum NW length decreases to 1000 nm for z = 0.6 , yielding G = 0.833 in this case. Using the above linear relationship between the two collection lengths, we deduce the values λ Al = 250 nm and λGa =780 nm at z = 0.2 , decreasing to λ Al = 8 nm and λGa = 155 nm at z = 0.6 . Therefore, both collection lengths strongly decrease with increasing the aluminum content in vapor, consistent with a much shorter NW lengths and stronger tapering of aluminum-rich shells toward larger z . At z = 0.6 , the aluminum diffusivity on the NW sidewalls is almost completely suppressed due to a higher density of surface steps where the adatoms incorporate into a tapered shell. The values for λGa of the order of several hundred nanometers are consistent with previously reported19,20,28-30. The inequality λGa / λ Al > 2.6 agrees with the result of Ref. [19], obtained from the analysis of axial growth rates of AlGaAs NWs. We now analyze a very rapid decrease of the aluminum content in the topmost part of the NW after the group III fluxes are switched off under the arsenic flux. Form Eqs. (1) for


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( s) = VAs / U As , we may aluminum and gallium at Vi = 0 and under the above assumption of c As

assume that the arsenic concentration stays constant also

in the kinetic stage. Then the

aluminum composition of the NW given by Eq. (3) decreases with time as VAl e − AK Al t V Al x= ≅ e − AK Al t , − AK Al t − AK Ga t VAl e + VGa e VAl + VGa

A=

α VAs R U As

(6)

where the approximate expression applies when K Al >> K Ga . At z = 0.2 , the 30 nm height over which the aluminum concentration decreases to zero is grown in approximately 7.2 s after the ( s) = VAs / U As = 0.03 flux termination. Assuming a typical value of the arsenic concentration c As

(Ref. [31]), β = 120 deg and Ω L = 0.02 nm3, we obtain A = 3.8 × 10-4 nm-1. This requires K Al ≅ 1100 nm/s to ensure that x tends to zero after 7.2 s upon the growth stop. The extremely high value of the aluminum bonding rate with arsenic is two or three orders of magnitude larger than the NW growth rate (~ 1-4 nm/s). In contrast, gallium remains at approximately 40% even in the droplet, meaning that the gallium bonding rate with arsenic should be at least two orders of magnitude lower than for aluminum ( KGa ~ 10 nm/s or less). Much higher K Al with respect to KGa agrees well with the results of Ref. [18], where the estimate K Al ≅ 200 × KGa was obtained for the analysis of interfacial abruptness in self-catalyzed AlGaAs axial heterostructures. Although the catalysts are different, we believe that a high crystallization rate of AlAs should pertain regardless of the liquid composition, i.e. for both Al-Ga and Au-Al-Ga droplets provided that arsenic is available for bonding in sufficient amounts. The huge K Al obtained here does not contradict the earlier assumption VAs >> G , equivalent to VAs >> VAl + VGa in our case. This condition does not contain K i and simply requires that the atomic influx of arsenic into the droplet (including re-emitted species) is much larger than the total influx of the group III atoms. To conclude, we presented TEM and EDX data and a supporting model that explains spontaneous formation of core-shell AlGaAs NWs entirely on the grounds of different material transport into the droplet for aluminum and gallium. Lower diffusivity of aluminum on the NW 11 ACS Paragon Plus Environment


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sidewalls leads to its lower content in the VLS core and at the same time the vapor-solid overgrowth of tapered shells. When both group III fluxes are switched off in arsenic environment, aluminum leaves the droplet almost instantaneously while gallium remains there at a high percentage. Huge difference between the bonding rates of aluminum and gallium with arsenic (at least two orders of magnitude) should be very advantageous for obtaining sharp NW heterointerfaces by switching the metal fluxes. The model describes quantitatively the compositional trends in the stationary VLS process as well as emptying the droplet at the end of growth, with a minimum number of parameters. That is why we were able to estimate some unknown kinetic constants, in particular the collection lengths and bonding rates for aluminum and gallium. Overall, these results can be useful for the compositional control and bandgap engineering in different ternary III-V NWs and NW heterostructures.

ACKNOWLEDGMENTS The NW samples were grown under the support of the Russian Science Foundation

(Project No. 14-12-00393). V. G. D. thanks the Russian Ministry of Education and Science for the financial support under the Contract No. 14.613.21.0055 (project ID: RFMEFI61316X0055).

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2016, 16, 1917. (19) Cirlin, G. E.; Dubrovskii, V. G.; Sibirev, N. V.; Soshnikov, I. P.; Samsonenko, Yu. B.; Tonkikh, A. A.; Ustinov, V. M. Semiconductors 2005, 39, 557. (20) Dubrovskii, V. G.; Hervieu, Yu. Yu. J. Cryst. Growth 2014, 401, 431. (21) Harmand, J. C.; Patriarche, G.; Péré-Laperne, N.; Mérat-Combes, M. N.; Travers, L.; Glas, F. Appl. Phys. Lett. 2005, 87, 203101. (22) Soda, M.; Rudolph, A.; Schuh, D.; Zweck, J.; Bougeard, D.; Reiger, E. Phys. Rev. B 2012, 85, 245450. (23) Dubrovskii, V. G.; Sibirev, N. V. Cryst. Growth Des. 2016, 16, 2019. (24) Ramdani, M. R.; Harmand, J. C.; Glas, F.; Patriarche, G.; Travers, L. Cryst. Growth Des.

2013, 13, 91. (25) Dubrovskii, V. G. Appl. Phys. Lett. 2014, 104, 053110. (26) Dubrovskii, V. G. Cryst. Growth Des. 2015, 15, 5738. (27) Glas, F. Phys. Stat. Solidi B 2010, 247, 254. (28) Plante, M. C, LaPierre, R. R. J. Appl. Phys. 2009, 105, 114304. (29) Harmand, J. C.; Glas, F.; Patriarche, G. Phys. Rev. B 2010, 81, 235436. (30) Dubrovskii, V. G.; Xu, T.; Díaz Álvarez, A.; Larrieu, G.; Plissard, S. R.; Caroff, P.; Glas, F.; Grandidier, B. Nano Lett. 2015, 15, 5580. (31) Glas, F. J. Appl. Phys. 2010, 108, 073506.


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For Table of Contents Use Only Origin of spontaneous core-shell AlGaAs nanowires grown by molecular beam epitaxy V. G. Dubrovskii1, I. V. Shtrom, R. R. Reznik, Yu. B. Samsonenko, A. I. Khrebtov, I. P. Soshnikov, S. Rouvimov, N. Akopyan, T. Kasama, G. E. Cirlin

Synopsis: We unravel the origin of spontaneous core-shell AlGaAs nanowires grown by gold-assisted MBE. Our data show that the composition of tapered shells is close to nominal, while the aluminum content in cylindrical cores is systematically smaller. After the growth stop, the aluminum concentration in the droplet and in the topmost part of the wire rapidly tends to zero, while gallium remains there at a high percentage. According to the model, smaller aluminum content in the core is due to its lower surface diffusivity. On the other hand, aluminum leaves the droplet at least 100 times faster than gallium, with a typical bonding rate with arsenic of 1000 nm/s.



Al in nanowire x

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0.6 0.5 0.4

Shell

0.3

Core

0.2 0.1 0.0

0.2

0.3

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Al in vapor z

ACS Paragon Plus Environment

0.6

50 nm

Origin of Spontaneous CoreâShell AlGaAs Nanowires Grown by

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Origin of Spontaneous CoreâShell AlGaAs Nanowires Grown by