GaAsSe Ternary Alloy Nanowires for Enhanced Photoconductivity

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C: Physical Processes in Nanomaterials and Nanostructures

GaAsSe Ternary Alloy Nanowires for Enhanced Photoconductivity Kidong Park, Jinha Lee, Doyeon Kim, Jaemin Seo, Jung Ah Lee, In Hye Kwak, Ik Seon Kwon, Jae-Pyoung Ahn, and Jeunghee Park J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.8b11652 • Publication Date (Web): 23 Jan 2019 Downloaded from http://pubs.acs.org on January 29, 2019

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The Journal of Physical Chemistry

GaAsSe Ternary Alloy Nanowires for Enhanced Photoconductivity Kidong Park,† Jinha Lee,† Doyeon Kim,† Jaemin Seo,† Jung Ah Lee,† In Hye Kwak,† Ik Seon Kwon,† Jae Pyoung Ahn,‡ and Jeunghee Park*,† † Department

‡

of Materials Chemistry, Korea University, Sejong 339-700, Republic of Korea

Advanced Analysis Center, Korea Institute of Science and Technology, Seoul 136-791, Republic of Korea

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ABSTRACT Alloyed semiconductor nanowires are of great interest for next-generation integrated optoelectronic nanodevices owing to their tunable band gap and emission wavelength. In this study, we synthesized the GaAsSe ternary alloy nanowires (NWs) with various compositions between GaAs and Ga2Se3 using chemical vapor transport method. The band gap was continuously tuned in the range of 1.5–2.1 eV because of the completely miscible solid solution at all compositions. The alloy NWs (including Ga2Se3) consisted of cubic phase with the [011] growth direction, in contrast with the GaAs NWs grown along the [111] direction. In particular, the GaAs1xSex

(x = 0.3) alloy NWs were grown from Ga-rich Au nanoparticles such as cubic phase AuGa2

and had a defect-free single-crystalline nature. X-ray photoelectron spectroscopy analysis reveals much less surface oxide layers for x = 0.3, suggesting that Se incorporation at this composition effectively diminishes the surface defects. We fabricated photodetectors using the individual NW, showing that the photocurrent decreases with increasing Se composition. The alloy composition significantly diminished the dark current and thus greatly enhanced the photosensitivity for x = 0.3.


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1. INTRODUCTION One-Dimensional (1D) semiconductor nanowires (NWs) have attracted significant attention because they can be versatile building blocks for next-generation flexible and nanoscale electronics and optoelectronics.1 For almost two decades, their unique optical and electrical properties arising from the high aspect ratio have been extensively investigated using bottom-up approaches. As a typical III-V semiconductor, gallium arsenide (GaAs) NWs are among the most promising candidates for energy harvesting devices such as photodetectors and photovoltaic cells, because of their high carrier mobility as well as direct band gap.2-18 Compared to the original components, alloys offer the opportunity to tune the band gap and emission wavelength by adjusting the composition. A great amount of work has been performed in ternary alloys such as InGaAs,19-23 AlGaAs,15,17,24,25 GaPAs,26-28 and GaAsSb, using III or V compositions.29,30 The synthesis of alloy phase NWs further enables various complex architectures such as axial and radial heterostructures, and therefore extends the application range into nanolasers and light emitting diodes (LEDs). Since the controlled growth of alloy NWs remains difficult due to the complex thermodynamic and kinetic parameters involved in the growth process, the alloys with VI composition such as chalcogen (S, Se, Te) has been rarely reported yet. Gallium sesquiselenide (Ga2Se3) has a defective zinc blende (ZB)-like structure, in which one third of the cationic sites are vacant.31-33 There are several polymorphs that differ in the arrangement of the atoms and vacancies and the degree of order. The low-temperature α phase Ga2Se3 (cubic phase, sphalerite) contains vacancies randomly distributed at the cationic sites. Annealing first produces the monoclinic (β) and then the orthorhombic phases having zigzagordered and straightly aligned vacancies along the 110 direction, respectively. The optical band gap of α-Ga2Se3 was measured to be 2.0-2.1 eV.34-36 In the case of β-Ga2Se3, a broad range over 3 ACS Paragon Plus Environment


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2.0-2.85 eV was reported for the optical band gap by a number of research groups.37-39 Firstprinciples calculations predicted a nondispersing valence band maximum along the vacancy line direction, with a direct band gap of 2.56 and 2.0 eV for the monoclinic and orthorhombic phases, respectively.39 As an important application, chalcogenide passivation of GaAs has been actively developed to eliminate undesirable defect sites on the surface or interface.40-44 Recently, Se passivation of GaAs NWs was reported to suppress the surface recombination of charge carriers and in turn significantly improves the device performances.43,44 Inspired by these results, the GaAsSe ternary alloy NWs would be an attractive system to advance the optoelectronic devices. In the present work, we successfully synthesized the GaAsSe alloy NWs using a chemical vapor transport method. The composition of the cubic phase alloys, from GaAs to α-Ga2Se3, was varied over the whole range by changing the ratio of the precursors (GaAs and β-Ga2Se3 powders); i.e., GaAs1-xSex (x  0.5) and (GaAs)1-x(Ga2Se3)x (x > 0.5). As a result, the band gap (Eg) was tuned between 1.5 and 2.1 eV, confirming that the cubic phase α-Ga2Se3 NWs has a band gap of 2.1 eV. To our best knowledge, it is the first report of achieving completely miscible phase at all ternary compositions. Se incorporation into the GaAs NWs changed the growth direction from [111] to [011], accompanied by an enriched Ga composition at the Au nanoparticle tip. We further fabricated prototype photodetectors using the individual NW, demonstrating that incorporation of Se at x = 0.3 results in better performance than the GaAs NWs, promising a wide range of applications. 2. EXPERIMENTAL SECTION GaAs (99.999%, Alfa Aesar) and Ga2Se3 (99.99%, Alfa Aesar) powders were placed inside a quartz tube reactor. A silicon substrate on which a 3 nm thick Au film was deposited was 4 ACS Paragon Plus Environment

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positioned 10 cm away from the powder source. Argon gas was continuously supplied at a rate of 200 sccm during the synthesis. Hydrogen (H2) gas was co-flowed at a rate of 10 sccm. The temperature of the powder sources was set to 950 °C. The substrate was approximately maintained at 750°C for 20-60 min to synthesize the NWs. The composition was controlled by changing the ratio of the GaAs and Ga2Se3 powder. Characterization and fabrication of NW devices are described in Supporting Information, Experimental Details. 3. RESULTS AND DISCUSSION The GaAs, Ga2Se3, and alloy NWs were synthesized by evaporating GaAs, β-Ga2Se3, and their powder mixture, respectively. SEM images and energy-dispersive X-ray spectroscopy (EDX) data of each sample are shown in Supporting Information, Figure S1. The obtained X-ray diffraction (XRD) patterns are shown in Supporting Information, Figure S2. The XRD peaks of end compositions (x = 0 and 1, namely GaAs and Ga2Se3) matched respectively those of ZB phase GaAs (JCPDS Card No. 80-0016; F43m, a = 5.654 Å) and cubic phase α-Ga2Se3 (JCPDS Card No. 05-0724; F43m, a = 5.429 Å). As the Se content increases, the peaks of the GaAs phase blueshift towards those of the α-Ga2Se3 phase. The lattice constant of α-Ga2Se3 is 4.0% smaller than that of GaAs, and the effective radius of Se anion is smaller than that of As anion (184 pm vs. 210 pm). The composition (x) calculated using the XRD peak position was compared with the EDX data (see Supporting Information, Figure S3), and those of the alloy phase were determined as GaAs1-xSex for x  0.5 and (GaAs)1-x(Ga2Se3)x for x > 0.5. The peak widths of x > 0.5 are larger than that of x  0.5, which is ascribed to the defects of NWs that will be discussed later.



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Figure 1. (a) Scanning electron microscopy (SEM) images of GaAs1-xSex (x = 0.3) NWs homogeneously grown on a substrate. High-resolution transmission electron microscopy (HRTEM) and corresponding fast Fourier-transform (FFT) images of (b) GaAs, (c) GaAs1-xSex (x _

= 0.3), and (d) Ga2Se3 NWs (at the [011] zone axis marked by (ii)), showing the growth direction; [111], [011], and [011], respectively. Selected-area electron diffraction (SAED) pattern for the tip (marked by (i)) are matched to (b) hexagonal phase Au7Ga2 (zone axis = [0001]), (c) cubic phase _

AuGa2 (zone axis = [011]), and (d) orthorhombic phase Au2Ga (zone axis = [131]). The SEM image in Figure 1a shows that the GaAs1-xSex (x = 0.3) NWs grew densely on the substrates. TEM images revealed that the average diameter is 150 nm for x  0.5, and 80 nm for x _

> 0.5. HRTEM and FFT images of individual NW at the zone axis of [011] are shown in Figures 1b-1d. The GaAs NWs are consisted of the single-crystalline ZB phase with the [111] growth 6 ACS Paragon Plus Environment

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direction (Figure 1b). The d spacing between the neighboring (111) planes is 3.26 Å, which is consistent with that of bulk GaAs (d111 = 3.2644 Å with a = 5.654 Å). The GaAs NWs have straight and smooth surfaces. The round nanoparticle at the tip is attached to the (111) facet of GaAs NWs. The SAED pattern for the tip part revealed hexagonal phase Au7Ga2 (JCPDS No. 24-0424, P62m, a = 7.721 Å and c = 8.751Å). The GaAs1-xSex (x = 0.3) NWs exhibit single-crystalline nature and the growth direction of [011] (Figure 1c). The d spacing of their (011) planes is d011 = 3.9 Å, which is consistent with that linearly interpolated from d011 = 3.998 Å for GaAs and 3.838 Å for α-Ge2Se3. There exists a polygonal shaped nanoparticle at the tip, and its SAED pattern corresponds to the cubic phase AuGa2 (JCPDS No. 03-0969, Fm3m, a = 6.073 Å). The crystallographic directions of AuGa2 are all matched with those of GaAs. The rhombohedral phase -GaSe layers (JCPDS Card No. 81-1971; R3m, a = 3.730 Å and c = 23.860 Å) sheathed the AuGa2 nanocrystals, whose TEM images will be discussed later. As shown in Figure 1d, the cubic phase Ga2Se3 NWs are grown along the [011] direction. The d spacing of (011) planes is 3.8 Å, which is consistent with that of cubic phase -Ga2Se3 (d011 = 3.838 Å). The Ga2Se3 NW contain the 2-5 nm periodic stacking faults along the [111] direction, resulting in the lines in the FFT image. The spherical tip is tilted to the catalyst-NW interface along the [111] direction of the Ga2Se3 NW. The SAED pattern of the tip revealed orthorhombic phase Au2Ga (JCPDS No. 29-0619, a = 18.02 Å, b = 3.199 Å and c = 6.999 Å).



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Figure 2. HAADF STEM images (scale bar = 100 nm) and EDX elemental mappings (using the Au L-, Ga K-, As K-, and Se K-shell peaks) of (a) GaAs, (b) GaAs1-xSex (x = 0.3), and (c) Ga2Se3 NWs, and corresponding EDX line-scanned profiles along the axial direction of NWs and the tip. Figure 2a displays the high-angle annular dark-field scanning (HAADF STEM) image of GaAs NW, and its EDX elemental mapping/line profile (counts vs. distance) along the axial direction using the Au L-, Ga K-, and As K-shell peaks. The atomic ratio of Ga:As is 1:1 at the NW part, and that of Au:Ga is 7:2 at the tip part corresponding to the composition of Au7Ga2. The x = 0.3 alloy NWs shows a polygonal shaped nanoparticle at the tip (Figure 2b). The EDX mapping and 8 ACS Paragon Plus Environment

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line profile along the axial direction reveal the atomic ratio of Ga:As:Se = 1:0.7:0.3 for the NW, and Au:Ga = 1:2 for the center part of the tip to confirm the AuGa2 phase. The thin layer sheathing the AuGa2 nanocrystal mainly consisted of Ga and Se. The Se composition at the tip is 15 at% against the Au. For Ga2Se3 NWs, the NW and tip parts have the respective atomic ratio of Ga:Se = 2:3 and Au:Ga = 2:1 that corresponds to a composition of Au2Ga (Figure 2c). The Se dissolved into the Au2Ga nanoparticles is 20 at%. The GaAs1-xSex (x = 0.2) NWs possess a straight NW morphology and the [011] growth direction (see Supporting Information, Figure S4). The round nanoparticle tip consists of orthorhombic phase AuGa. The GaAs1-xSex (x = 0.5) NWs have the polygonal shaped (Se-doped) AuGa nanoparticles sheathed with rhombohedral phase -GaSe layers (see Supporting Information, Figure S5). The (GaAs)1-x(Ga2Se3)x (x = 0.7) NWs have the As-containing -Ga2Se3 nanoparticle at the tip. Those NWs have a high degree of crystalline disorder, similar to the -Ga2Se3 NWs (see Supporting Information, Figure S6). The EDX data of the tip shows that the Se composition increases with increasing x, reaching 20 at% against the Au. The As composition at the Au-Ga alloy tip is negligible for x  0.5 and 10 at% at the -Ga2Se3 phase tip for x > 0.5 alloy. We summarized our results in Table 1, including their (photo)electrical properties which will be described later. Now we discuss the growth mechanism of GaAs NWs. The round shape of Au7Ga2 nanoparticles indicates that the growth follows the typical vapor-liquid-solid (VLS) mechanism, since the melting point of Au7Ga2 (~500 °C) is lower than the growth temperature (750 °C).45,46 The H2 flow facilitates the production of precursor vapors via the reaction 2GaAs(s) + 3H2(g)  2Ga(g) + 2AsH3 and promotes the vapor transport. Then, the Ga dissolves into the Au nanoparticles by forming Au-Ga eutectic, while the concentration of As in the nanoparticle is negligibly small due to low solubility. Then, the saturated Ga precipitate with As vapors at the 9 ACS Paragon Plus Environment


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triple phase region (i.e., the interface between the NW solid, the Au alloy liquid droplet, and the gas phase), to produce the NWs. The catalyst-NW interface consists of (111) planes with the lowest surface energy, indicating that precipitation on the (111) surface leads to the largest decrease in the Gibbs free energy.47,48 Table 1. Summary of structure and properties of cubic phase GaAsSe NWs. Composition Formular a

x

[As]:[Se]

Growth Direction

Composition of Tip

Carrier Type

Mobility (cm2 V-1 s-1)b

Ip c

0.10 1.7 A 0 1:0 [111] Au7Ga2 p -0.2 8:2 [011] p -AuGa GaAs1-xSex 0.3 7:3 [011] p Au2Ga 7.9(0.8)10-5 1.4 A 0.5 1:1 [011] -AuGa -0.5 nA 5.3(0.5)10-3 0.7 1:7 [011] Ga2Se3 n 0.1 nA (GaAs)1-x 0.8 1:12 [011] Ga2Se3 n --(Ga2Se3)x 1 0:1 [011] n Au2Ga 1.2(0.1)10-2 0.05 nA a Determined by the TEM and EDX. b See the data in Supporting Information, Figure S10. c Photocurrents of NW electrode at a bias voltage of 2 V, under the irradiation of 365 nm. d Photosensitivity (I /I p dark), where Idark is the current under dark condition.

Ip/Idark d 4 -1.4103 103 102 -2.5102

Ohta et al. reported that the monoclinic phase -Ga2Se3 NWs are grown along the [011] direction and laterally on As-terminated Si substrates, which could be ascribed to the ordered Ga vacancies.49 However, the present -Ga2Se3 phase has random Ga vacancies, so the vacancies may not be important in driving the [011] growth. It is noteworthy that the Au2Ga tip of Ga2Se3 NWs is richer in Ga than the Au7Ga2 tip of GaAs NWs, and it also contains the 20 at% Se. This observation suggests that the Ga-Se vapor is dissolved in the Au nanoparticle owing to the reaction of Ga2Se3(s) + H2(g)  2Ga-Se(g) + H2Se and form the eutectic (melting point = 339.3 °C for Au2Ga). We presume that as the Se dissolves in the Au nanoparticles, the growth rate would be higher compared to the case of GaAs NWs. Then, the faster growth rate results in the [110] growth 10 ACS Paragon Plus Environment

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direction that is kinetically favorable.4 On the other hand, the [111] growth direction of GaAs NWs is thermodynamically favorable and has a slower growth rate due to the un-dissolved As. In one of the important models for the Au-assisted VLS growth of NWs, the diameter of NWs can be a crucial factor in controlling the growth direction.47,48 In the case of Si NWs, for example, at smaller diameters the [011] growth direction becomes favorable. Since the Ga2Se3 NWs have about half of the GaAs NWs in diameter, the [011] growth direction would also be related to their smaller diameter. The smaller diameter induces the faster growth rates, consistently with our model. The [011] growth direction benefits from the side walls comprising of four low-energy {111} facets, leading to the largest decrease in Gibbs free energy during growth.47,48 The AuGa2 tip for x = 0.3 alloy NWs has a higher Ga concentration than the Au7Ga2 tip for GaAs or AuGa tip for x = 0.2. Therefore, as x increases, the more Ga dissolved into the tip, probably due to the more dissolution of Ga-Se. The polygonal shape of AuGa2 nanoparticle suggests its solid state during NW growth instead of in the molten form. As the Ga composition increases, the melting point increases from 461.3 °C for AuGa, and 491.3 °C for AuGa2.46 However, the melting point of AuGa2 is still lower than the growth temperature of NWs. One possibility is that the Se dissolution may elevate the melting point, which is referred from the melting points of ~960 °C for 20% Se-Au alloy and ~920 °C for 10% Se-Ga alloy. If the catalytic nanoparticle is sold during growth, the growth mechanism would follow a vapor-solid-solid (VSS) process, in which the NWs are epitaxially grown on the Au-Ga-Se alloy. The lattice matching at the interface would direct the structure as well as the growth direction of the NWs. In this regard, the crystal orientation of the tip/NW interface becomes favored along the [011] direction of cubic phase NWs.



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Another possibility is that the Au2Ga nanoparticles are in molten phase and solidified with the polygonal shape after termination because of the GaAs sheathing layers that formed by the precipitation of the dissolved Ga and Se in the nanoparticles. The growth rate would be higher than that of GaAs NWs, resulting in the kinetically favorable [110] growth direction via the VLS process. We expected that the Ga concentration could be actually higher than that of AuGa2 during growth. The melting point of GaSe (938 °C) is high enough to allow it to exist in the solid during the termination. In fact, the d006 value of -GaSe and d011 value of GaAs are perfectly matched, which plays a critical role in initiating the gowth of GaSe layers. At x = 0.2, the less dissolution of Ga-Se produces the AuGa, and the GaAs NWs are grown along the [011] direction, following the VLS mechanism. At x = 0.5, the enhanced dissolution of Ga-Se produces thicker GaSe layers, and the AuGa composition results from the more severe depletion of Ga. As x further increases to 0.7, the Ga2Se3 nanoparticles are grown after growth termination, and cover fully the Au alloy nanoparticles. The above discussion could explain how the different growth directions of NWs are related to the phase and morphology of the tip. Further studies are required to understand the growth mechanism. The electronic structures of the GaAs1-xSex (x = 0, 0.3, and 0.5) and (GaAs)1-x(Ga2Se3)x (x = 0.7 and 1) NW samples were analysed using synchrotron X-ray photoelectron spectroscopy (XPS) where the photon energy is 600 eV. The survey scan is shown in Supporting Information, Figure S7. Figure 3a shows the fine-scanned Ga 3d peaks whose 3d3/2 and 3d5/2 (separated by 0.44 eV) are unresolved.50 The peaks were resolved into two components (Ga1 and Ga2), which are blue shifted from the Ga metal (Ga0) at 18.6 eV, using a Voigt function: Ga1 (red) at 18.4-19.4 eV and Ga2 (green) at 19.8-20.3 eV. The Ga1 and Ga2 were assigned to the Ga-As and/or Ga-Se and GaO, respectively. The fraction of Ga-O band was determined as 75%, 45%, 60%, 55%, and 55%, 12 ACS Paragon Plus Environment

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respectively, for x = 0, 0.3, 0.5, 0.7, and 1. The lowest value at x = 0.3 indicates the minimized surface oxide layers at this composition. The O 2s band appears at 24 eV for all sample except x = 0.3, confirming our analysis.

Figure 3. Fine-scanned XPS spectrum: (a) Ga 3d, (b) As 3d, and (c) Se 3d peaks for GaAs1-xSex (x = 0, 0.3, and 0.5) and (GaAs)1-x(Ga2Se3)x (x = 0.7 and 1) NW samples. The data (open circles) are fitted by a Voigt function, and the sum of the resolved bands is represented by a black line. The positions of the peaks corresponding to neutral species (3d5/2 of Ga0, As0, and Se0) are marked by dotted lines to highlight the corresponding shift. Figure 3b shows the As 3d5/2 and 3d3/2 peaks (separated by about 0.69 eV). The peak at 40.5 eV is redshifted from that of neutral As (As0 at 41.6 eV). This band is resolved into two components (As1 and As2); each consisted of 3d5/2 and 3d3/2 peaks. The As1 (red) at 40.2 eV is attributed to As in the fourfold coordinated GaAs. The As2 (green) at 41.5 eV is shifted by 0.7 eV towards higher binding energy, which originates from the dimeric As on the surface.50 The blue-shifted 13 ACS Paragon Plus Environment


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band (As3, blue) at 43.5-43.7 eV is assigned to As-O at surface. For x = 0.3, the fraction of As2 and As3 is lower than that of GaAs, indicating the dimeric surface defects as well as the surface oxide layers are much reduced. At x = 0.5 and 0.7, the As3 becomes a dominant component, owing to the increased surface oxide layers. Figure 3c shows the XPS Se 3d peaks for x = 0.3, 0.5, 0.7, and 1. Two components (each consisted of 3d3/2 and 3d3/2 peaks separated by about 0.86 eV) are apart by about 1 eV; Se1 (red) at 52.8-53.9 eV and Se2 (green) at 53.7-55 eV, redshifted from that of neutral Se 3d5/2 (Se0 at 55.6 eV), are attributed to the Ga-Se subsurface and surface, respectively.50,51 The energy difference of two components coincides quite well with the reported value (0.92 eV) and the ratio Se1/Se2 is 0.3-0.5.51 The Se1 has a narrowest width at x = 0.3, indicating the highest degree of crystallinity. We conclude that (i) the GaAs NWs have the surface defects originated from the oxide layers and the dimeric As forms, (ii) at x = 0.3 the surface oxide layers becomes minimized and the defects are also much reduced, and (iii) the x = 0.5-1 NWs have significant surface oxide layers. The UV-visible absorption spectra in the diffuse reflectance mode were measured for as-grown GaAs1-xSex (x = 0, 0.3, and 0.5) and (GaAs)1-x(Ga2Se3)x (x = 0.7 and 1) NW samples. As shown in Figure 4a, the absorption band is blueshifted as x increases: its onset appears at 1.4 eV (880 nm) for GaAs and 2.0 eV (620 nm) for Ga2Se3, increasing almost linearly with increasing x. Based on the Kubelka-Munk (K-M) transformation, the plot of [F()h]2 vs. photon energy, where F() is the diffuse reflectance spectrum, yielded the direct band gap (Eg), as shown in the inset. We measured the Raman spectrum for all NW samples, providing another evidence for the composition tuning (Supporting Information, Figure S8).


x=0 x=1

1.5

1.5

2.0

2.0

(b) 2.2 2.0

Eg (eV)

=0 = 0.3 = 0.5 = 0.7 =1

PL Intensity (arb. units)

x x x x x

2

(a)

(F(R) h )

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Absorption (arb. units)

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1.8 1.6 1.4

0.0 0.2 0.4 0.6 0.8 1.0

2.5

Energy (eV)

Composition (x)

Figure 4. (a) UV-visible diffuse reflectance spectrum (in absorption) and K-M plots of GaAs1-xSex (x = 0, 0.3, and 0.5) and (GaAs)1-x(Ga2Se3)x (x = 0.7 and 1) NW samples. (b) Dependence of the bandgap (Eg) on the Se content (x), determined by the K-M plots. Photoluminescence (PL) spectrum of GaAs and Ga2Se3 is plotted in (a), which are measured at 8 K using the excitation energy of 2.41 eV (514 nm) and 3.81 eV (325 nm), respectively. Figure 4b shows the Eg value (obtained by K-M plot) for all NW samples synthesized in this work with different compositions (x). The Eg increase almost linearly with increasing x, but has an uncertainty of at least 5% that comes from the extrapolation process of the linear region. The Eg value of GaAs NWs (1.45 eV) is consistent with the band gap of ZB phase GaAs (bulk). We first observed that the cubic phase α-Ga2Se3 NWs have a band gap of 2.07 eV, which is consistent with that of the bulk phase (2.0–2.1 eV).34-36 The PL spectrum was measured by delivering continuouswave excitation from a 514 nm Ar ion laser to as-grown GaAs NW samples at 8 K. The PL peak position, corresponding to the optical band gap (Eg), is 1.48 eV, which is consistent with the value obtained by UV-visible spectrum. For Ga2Se3 NW, the PL spectrum was measured using a 325 nm 15 ACS Paragon Plus Environment


He–Cd laser, showing the broad emission in the range of 1.4-2 eV, which would be originated from the defect sites. The PL spectrum of alloy NWs shows a nearly same feature as that of Ga2Se3 NW.

Figure 5. (a) Current–voltage (I–V) characteristics of GaAs1-xSex (x = 0, 0.3, and 0.5) and (GaAs)1x(Ga2Se3)x

(x = 0.7 and 1) NW under dark conditions and 365 nm (60 mW cm-2) irradiation. The

SEM image shows the NW aligned between Ti/Au bottom electrodes with a 2.2 μm gap and Pt top electrodes with a gap of 14 m (inset). Inset shows the photocurrent of x = 0 and 0.3 vs. light intensity (mW cm-2). (b) I–t curves at a bias voltage of 2 V under chopped irradiation. Next, photodetectors were fabricated using individual NW by employing the dielectrophoresis and focused ion beam techniques. Figure 5a shows the I-V curves measured for the GaAs1-xSex (x = 0, 0.3, and 0.5) or (GaAs)1-x(Ga2Se3)x (x = 0.7 and 1) NW under dark and under light irradiation of 365 nm (3.4 eV with 60 mW cm-2 using LED). The SEM image (inset) shows that the NW is aligned between the electrodes. The I-V curves under light irradiation are almost linear within the measured range (2 V). From x = 0.5 to 1, the current change upon light irradiation decreases 16 ACS Paragon Plus Environment

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significantly by three orders of magnitude. Figure 5b displays the I-t curves measured by applying a bias voltage of 2 V, collected in real time for five on/off cycles. The photocurrent (Ip), defined as Ilight-Idark and corresponds to the current increase upon light irradiation, decreases with increasing x. This decrease is probably due to the reduced absorption coefficient with increasing x. The photocurrent at 2 V is 1.7 and 1.4 A for x = 0 and 0.3, respectively, showing excellent stability over a long time. At x = 0.5, 0.7, and 1, the photocurrent is in nA scale (0.5 nA, 0.1 nA, and 0.05 nA, respectively). Compared to GaAs NWs, the dark current (Idark) of the alloy NWs is much reduced; 0.4 A, 1 nA, 0.5 pA, 1 pA, and 0.2 pA at 2 V, respectively, for x = 0, 0.3, 0.5, 0.7, and 1. The photosensitivity (Ip/Idark) is 1400 for x = 0.3, while it is only 4 for GaAs. The photosensitivity is 1000, 100, and 250 for x = 0.5, 0.7, and 1, respectively. For the GaAs and x = 0.3 NWs, the photocurrents increase linearly with the light intensity (0–60 mW cm-2), as shown in the inset of Figure 5a (see the data in Supporting Information, Figure S9). Spectral responsivity (R), defined as the photocurrent generated when light of unit intensity shines on the effective area of NW, can be expressed as R = Ip/PA, where P is the incident light intensity (= 60 W cm-2) and A is the effective area of NW. GaAs and x = 0.3 exhibited R = 5.5  103 and 4.5 103 A W-1, respectively. Another figure-of-merit of a photodetector is its specific detectivity (D*). When the noise from dark current is small, it can be defined as D* = R (A/2eIdark)1/2 The respective values of GaAs and x = 0.3 are 3.5  1011 and 1.8  1013 Jones (i.e., cm · Hz1/2 W-1). The D* value of x = 0.3 is two-times larger than the value (9.04  1012) of photodetectors recently fabricated using sulfur (S)-passivated GaAs NWs reported by Chen et al.43 A further comparison with the previous photodetector parameters of GaAs (including GaAsSb and InAs) NWs has been performed and shown in Table S1 (Supporting Information). 17 ACS Paragon Plus Environment


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The higher photodetector performance for x = 0.3 than that of GaAs NWs is due to the much lower dark current. The Se incorporation could act like surface passivation, which is consistent with the recent work by Chen et al. that showed S passivation of GaAs NWs decreases the dark current of single-NW photodetector.43 The decrease of dark current was also observed for the Spassivated InGaAsSb/GaSb p-i-n-photodiode.41 Since the substitution of Se (x = 0.3) for As effectively protects the NWs from oxidation and diminishes the defects (as proven by XPS data), it would eliminate the unsaturated dangling bonds of GaAs that cause the dark current.43,44 As the Se content increases further, the passivation effect reduces due to the increase of surface defect sites, which is supported by the increase in oxide layers. We fabricated field effect transistor (FET) using single NW, as shown in Supporting Information, Figure S10. The measurement of gate effect revealed that the GaAs and Ga2Se3 NWs are p-type and n-type semiconductors with the carrier mobilities of 0.102 and 0.012 cm2 V-1 s-1, respectively. Increasing x converts the NW from p-type to n-type. Calculation demonstrated that the As antisite defects is highly stable in the p-type GaAs NWs, which is supported by our XPS analysis of dimeric As peak.52 The carrier mobility/concentration and thus the gate effects are reduced significantly for the alloy NWs. The Ho group reported that the native surface oxide shell layer of GaAs NWs would induce the acceptor-like interface trapping defects, which reduce carrier concentrations in the NWs and lead to p-type conduction.44 The Se passivators diminish the space charge depletion effect and makes the NWs less of p-type.53 For alloy NWs, the carrier mobility and concentration both decrease, which is similar with their passivation effects. The elimination of the charge recombination or trapping at surface defect sites is responsible for the reduced dark current. Therefore, the Se-incorporated alloy NWs hold great promise for applications in optoelectronic devices. 18 ACS Paragon Plus Environment

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4. CONCLUSIONS We synthesized GaAsSe ternary alloy NWs on Au-deposited Si substrates by a chemical vapor transport method utilizing the thermal evaporation of GaAs/Ga2Se3 powders. The composition was determined as GaAs1-xSex for x  0.5 and (GaAs)1-x(Ga2Se3)x for x > 0.5. The formation of a completely miscible solid solution in the cubic phase over the whole range resulted in the continuous tuning of the direct band gap (1.5–2.1 eV). For the first time, we synthesized the cubic phase -Ga2Se3 NWs, which have a band gap of 2.1 eV. The GaAs NWs were grown in the [111] direction, while the Ga2Se3 NWs in the [011] direction, following a typical VLS growth mechanism using the catalytic Au nanoparticles. The alloy NWs exhibit the [011] growth direction, which is attributed to the VSS or VLS growth using Ga-rich Au nanoparticles. In particular, the x = 0.3 NWs were grown from the AuGa2 catalytic nanoparticles. XPS data reveal that the surface oxide layers are minimized at x = 0.3, and the defects are also much reduced. Photodetectors using individual NW show that the photocurrent decreases with increasing Se composition. The Se incorporation significantly reduces the dark current, and thus the photosensitivity becomes maximized at x = 0.3. The gate effect of FET devices diminished at x = 0.3, indicating the elimination of the charge recombination or trapping at surface defect sites. These results suggest a new strategy for enhancing the performance of photodetectors. ACKNOWLEDGEMENTS This study was supported by 2014R1A6A1030732, 2017K000494, and 2018R1A2B6003624, funded by the Ministry of Science and ICT. The HVEM measurements were supported by the KBSI under the R&D program (D38700). The experiments at the PLS were partially supported by MOST and POSTECH. 19 ACS Paragon Plus Environment


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AUTHOR INFORMATION Corresponding Authors *E-mail: [email protected]. ORCID Kidong Park (0000-0003-1273-2076) Jinha Lee (0000-0002-2836-4186) Doyeon Kim (0000-0002-0762-2529) Jaemin Seo (0000-0002-1499-622X) Jung Ah Lee (0000-0002-8418-3902) In Hye Kwak (0000-0003-4697-0566) Ik Seon Kwon (0000-0003-0611-2276) Jae-Pyoung Ahn (0000-0003-2657-7425) Jeunghee Park (0000-0002-6913-5569) ASSOCIATED CONTENT Supporting Information. Experimental details and Figures S1-S10. This material is available free of charge via the Internet at http://pubs.acs.org. SEM, XRD, TEM, EDX, XPS, Raman, power-dependent photocurrents, and I-V data of FET devices. REFERENCES (1) Wang, J. -L.; Hassan, M.; Liu, J. -W.; Yu, S. -H. Nanowire Assemblies for Flexible Electronic Devices: Recent Advances and Perspectives. Adv. Mater. 2018, 1803430. (2) Gudiksen, M. S.; Lauhon, L. J.; Wang, J.; Smith, D. C.; Lieber, C. M. Growth of Nanowire Superlattice Structures for Nanoscale Photonics and Electronics. Nature 2002, 415, 617-620. (3) Persson, A. I.; Larsson, M. W.; Stenström, S.; Ohlsson, B. J.; Samuelson, L.; Wallenberg, L. R. Solid-Phase Diffusion Mechanism for GaAs Nanowire Growth. Nat. Mater. 2004, 3, 677681. (4) Shtrikman, H.; Popovitz-Biro, R.; Kretinin, A.; Heiblum, M. Stacking-Faults-Free Zinc Blende GaAs Nanowires. Nano Lett. 2009, 9, 215-219. (5) Joyce, H. J.; Gao, Q.; Tan, H. H.; Jagadish, C.; Kim, Y.; Fickenscher, M. A.; Perera, S.; Hoang, T. B.; Smith, L. M.; Jackson, H. E. et al. Unexpected Benefits of Rapid Growth Rate for 20 ACS Paragon Plus Environment

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GaAsSe Ternary Alloy Nanowires for Enhanced Photoconductivity

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