Solution Growth of Screw Dislocation Driven α-GaOOH Nanorod

Jul 31, 2017 - Using the α-GaOOH/α-Ga2O3/ZnGa2O4 wide bandgap transparent conductor materials as a demonstration, we synthesize vertical arrays of h...
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Solution Growth of Screw Dislocation Driven α‑GaOOH Nanorod Arrays and Their Conversion to Porous ZnGa2O4 Nanotubes Hanfeng Liang,†,‡ Fei Meng,† Brandon K. Lamb,† Qi Ding,† Linsen Li,† Zhoucheng Wang,‡ and Song Jin*,† †

Department of Chemistry, University of WisconsinMadison, 1101 University Avenue, Madison, Wisconsin 53706, United States College of Chemistry and Chemical Engineering, Xiamen University, Xiamen 361005, China

‡

S Supporting Information *

ABSTRACT: The solution synthesis of ternary metal oxides is difficult due to the competing hydrolysis of metal ions. There are reports of hydro-/solvothermal growth of nanoparticles, but one-dimensional (1D) nanoarrays are less common. Here, we report an alternative and general strategy to circumvent this challenge by converting the 1D binary metal oxide/hydroxide nanostructures initially grown driven by screw dislocations into ternary oxides. Using the α-GaOOH/ α-Ga2O3/ZnGa2O4 wide bandgap transparent conductor materials as a demonstration, we synthesize vertical arrays of high aspect ratio α-GaOOH nanorods (NRs) on conducting substrates with controllable length for the first time using a continuous flow reactor and confirm their growth mechanism to be dislocation-driven. Then the α-GaOOH NR arrays can be converted into porous α-Ga2O3 NR arrays, which can be further converted via a solution method into porous ZnGa2O4 nanotube (NT) arrays due to the Kirkendall effect. This work presents a new and general strategy to prepare 1D nanostructure arrays of various binary and ternary oxides at low cost and large scale, and such facile solution growth and the unique structure of porous ZnGa2O4 NT arrays will facilitate their practical applications.

O

The screw dislocation driven (SDD) growth is a powerful and versatile technique for growing 1D (as well as 2D) nanomaterials26−32 that are suitable for solution synthesis. In general, dislocation-driven growth requires two ingredients, the presence of dislocation sources and a suitable low precursor supersaturation.27 By rational synthetic design, various materials with 1D morphologies have been successfully prepared following the SDD growth mechanism.29,33−44 Because SDD growth does not rely on catalysts or surfactants and can be easily conducted not only in vapor phase but also in aqueous solutions,27 1D nanomaterials can be synthesized in aqueous solutions at mild temperature instead of high temperature vapor phase growth, which can help to lower the cost and increase the throughput. Gallium oxyhydroxide (α-GaOOH) and gallium oxide (Ga2O3) are wide bandgap (∼4.8 eV) semiconductors that exhibit high transparency in a wide wavelength range and a low refractive index, and they are attractive for power electronics and can serve as efficient surface passivation layers.45−48 Ga2O3 has also been shown to be the best n-type semiconductor to pair with p-type Cu2O semiconductor in a heterojunction to lead to efficient photovoltaic solar cells.49,50 As potential contact materials or semiconductors, vertically aligned NR

ne-dimensional (1D) nanostructures, such as nanorods (NRs), nanowires (NWs), and nanotubes (NTs) have exhibited new fundamental physical properties and demonstrated promising applications in nanoelectronics, nanophotonics, solar energy conversion, thermoelectrics, catalysis, energy storage, and chemical and biological sensing.1−7 The many synthetic methods of 1D nanostructures can be generally categorized into two groups: vapor phase (e.g., vapor−liquid− solid, or VLS) method and solution phase (e.g., hydro/ solvothermal) methods.7−16 The vapor phase method is the more extensively studied approach and has been successfully used to synthesize a wide range of 1D nanostructures with excellent control; however, it requires high temperature, is usually more energy intensive and costly, and therefore is more suitable for electronics, photonics, or other smaller scale device applications. In contrast, low temperature aqueous solution growth could be a facile and economic approach to the synthesis of 1D nanostructures on a large scale while maintaining sufficient crystal quality7,13 and therefore is much more suitable for large scale renewable energy applications. For example, solution growth of TiO2 and ZnO NR arrays on conducting substrates enabled their applications in dyesensitized solar cells, quantum dot solar cells, and solar water splitting with improved performance.17−23 Large-area NW arrays of metal oxides and their heterostructures on conducting substrates have also been very useful for electrochemical energy storage applications.24,25 © 2017 American Chemical Society

Received: May 10, 2017 Revised: July 28, 2017 Published: July 31, 2017 7278

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Figure 1. Schematic illustration of the growth of Ga-based oxyhydroxide, oxide, and ternary spinel oxide nanostructures on fluorine-doped tin oxide (FTO) substrate. (A) Solution growth of α-GaOOH NR arrays through screw dislocation-driven growth (SDD) in a continuous flow reactor (CFR). (B) Porous α-Ga2O3 NR arrays acquired by thermal treatment of α-GaOOH NRs at 450 °C for 4 h. (C) ZnGa2O4 microcubes converted from αGaOOH NRs at 200 °C in solution. (D) Porous ZnGa2O4 NT arrays converted from porous α-Ga2O3 NRs at 200 °C in solution. The colors are only for illustration purposes and are not the actual colors of the materials.

Figure 2. Structural characterization of the α-GaOOH NR arrays grown by a static reaction and a continuous flow reaction (flow rate: 1 mL min−1) with an initial Ga3+ concentration of 5 mM at 95 °C for 24 h. (A) Top view and cross-sectional view (inset) SEM images of the static grown αGaOOH NR arrays. (B) PXRD pattern and top view (C and D) and cross-sectional view (E) SEM images of the α-GaOOH NR arrays grown on FTO substrate using a CFR. The asterisks (*) in B mark the diffraction peaks from the FTO substrate.

commonly used substrates, such as FTO coated glass, cannot tolerate temperatures higher than 600 °C. Therefore, lowtemperature solution growth of ZnGa2O4 (and other MGa2O4, such as CdGa2O4 and NiGa2O4) 1D arrays is highly desirable for their potential application as TCO materials but has not been achieved yet. However, nanostructures of complex ternary metal oxides are difficult to synthesize directly under typical hydrothermal conditions due to the competing hydrolysis of metal ions.58,59 There are reports of hydro-/solvothermal growth of ternary oxide nanoparticles,59−61 but 1D NRs or NWs,62−64 especially aligned 1D arrays, are less common. An alternative strategy to circumvent the challenges associated with directly growing 1D nanostructures of more complex ternary oxides is to convert the precursor 1D binary oxide/hydroxide nanostructures into ternary oxides. Solution conversion offers a versatile route to convert nanostructures of one composition into another and could be equally effective as vapor conversion conducted at high temperatures, especially when the diffusion of the pertinent species is fast under low temperature solution conditions, as shown in the example of conversion of βNi(OH)2 to NiSe2 nanosheets.65 Here using ZnGa2O4 1D arrays as a demonstration, we take advantage of the relatively facile cation diffusion in nanoscale spinel oxide structures to prepare more complex ternary oxide compounds that are difficult to synthesize directly in solution. Specifically, by reacting precursor α-GaOOH NR or α-Ga2O3 NR arrays with Zn2+ ions in hydrothermal reactions, microcube thin films

arrays on conductive substrates could lead to more efficient devices,22,23 because such NR arrays can serve as the conducting scaffolds to contact to semiconductors or part of bulk heterojunctions. Here we first demonstrate the controllable solution synthesis of α-GaOOH NR arrays driven by screw dislocations for the first time. Vertical arrays of long (up to 3.5 μm) and high aspect ratio α-GaOOH NRs were grown on conducting fluorine-doped tin oxide (FTO) substrates (Figure 1A) using a home-built continuous flow reactor (CFR) under mild conditions. The as-obtained α-GaOOH NR arrays can be further thermally converted into single-crystalline but porous α-Ga2O3 NR arrays (Figure 1B) in solution. Furthermore, many gallium-based ternary spinel oxides (MGa2O4), such as ZnGa2O4 and CdGa2O4, are also wide bandgap semiconductors and promising p-type transparent conductive oxide (TCO) materials due to their excellent optical properties and high chemical and thermal stability.51,52 These spinel oxide materials have also been used as photocatalysts53 and luminescent probes.54 However, a lack of cost-effective and scalable growth of MGa2O4 1D arrays has limited their practical applications. Disordered ensembles of ZnGa2O4 NWs were previously synthesized by high-temperature (∼1000 °C) vapor reactions via the VLS growth using Au nanoparticle catalysts.55−57 Such vapor growth suffers from the high cost and difficulty in compositional control. Another seemingly inconspicuous but practically important drawback is that such high temperatures greatly limit the choice of substrates, because 7279

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Figure 3. TEM characterization of the as-grown α-GaOOH NRs. (A and B) TEM images showing the twist contour bands. (C) A collection of TEM images showing the nanotube morphology. (D and E) TEM images showing the dislocation contrast at the center of the NR and the emerging hollow core. (F) A collection of TEM images showing the dark-contrast longitudinal lines along the NRs.

the nucleation of α-GaOOH. To further control the supersaturation, we used urea instead of the NaOH used previously66 as the source of OH− ions. Urea can gradually release OH− through hydrolysis at temperatures above 90 °C (eq 1), which provides a steady supply of OH− ions during the hydrolysis reaction (eq 2), thus minimizing the fluctuation in precursor supersaturation. After the flow reaction, the product grown on FTO substrate was collected and first characterized by PXRD (Figure 2B). Except for the peaks originating from the FTO substrate, the diffraction peaks of the product can be assigned to the orthorhombic α-GaOOH phase (JCPDS No. 06-0180, space group Pbnm, a = 4.58 Å, b = 9.8 Å, c = 2.97 Å). A representative SEM image of the α-GaOOH NR arrays (Figure 2C) grown in a CFR at 95 °C for 24 h shows the surface of FTO substrate is uniformly and fully covered with vertically aligned α-GaOOH NRs. The small lattice mismatch (3.4%) of α-GaOOH with SnO2 (a = 4.74 Å) could facilitate the heterogeneous nucleation of α-GaOOH on the FTO substrate (note that the α-GaOOH cannot grow in nice NR arrays on a glass substrate, see Figure S2). The magnified SEM image (Figure 2D) reveals that the as-grown α-GaOOH NRs have quadrangular cross sections with rhombus top facets, the typical growth behavior for orthorhombic crystal structure. These NRs vary in width from 200 to 500 nm but have similar length. After a 24 h CFR growth, the length of the NRs was ∼2 μm as determined from the side-view SEM image (Figure 2E). These results suggest that the α-GaOOH NRs can be efficiently grown on FTO substrate via a simple, low-temperature solution reaction, and the CFR could enable a better control over the NR growth. We then carried out bright-/dark-field TEM characterization on the α-GaOOH NRs to confirm the existence of axial screw dislocations and sometimes hollow NTs. Since α-GaOOH is unstable under high-energy electron beam exposure, we used a low accelerating voltage (120 kV), a small condenser aperture, and a large beam spot size to minimize the damage on these NRs during imaging. However, a slightly longer exposure to electron beam still damaged the NRs, making the TEM experiments quite difficult. TEM images of individual NRs (Figure 3A,B) clearly reveal twist contour bands, which are

(Figure 1C) or porous NT arrays (Figure 1D) of ZnGa2O4 were obtained, respectively. Study of the conversion mechanisms revealed that the porous ZnGa2O4 NT arrays were likely formed due to the nanoscale Kirkendall effect, but the microcubes were formed through a dissolution−crystallization process. Our work provides a general strategy for the rational solution synthesis of dislocation-driven metal oxide and hydroxide NR arrays and their further conversion to 1D nanostructures of more complex ternary metal oxides, which could facilitate their applications.

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RESULTS AND DISCUSSION α-GaOOH NR Arrays. The chemical reaction that leads to the formation of GaOOH NRs is essentially a well-controlled hydrolysis of Ga3+ ions by reacting Ga(NO3)3·xH2O with urea at 95 °C in aqueous solution: (NH 2)2 CO + 3H 2O → 2NH4 + + 2OH− + CO2

(1)

Ga 3 + + 3OH− → GaOOH + H 2O

(2)

Note that there have been attempts at growing GaOOH NR arrays hydrothermally using Ga(NO3)3·xH2O precursor.66 We first carried out a similar static solution reaction (see details in the Materials and Methods section), which led to the growth of α-GaOOH NRs on FTO substrates but with a rather short length (∼490 nm, Figure 2A), similar to the products in the previous report.66 Such static reactions suffer from a significant drop of precursor supersaturation, leading to convoluted growth kinetics and uncontrolled products. Here we sought to promote the SDD growth and, thus, achieve better control of the NR growth. Because the SDD growth is favored in low supersaturation conditions, we conducted the growth of αGaOOH NRs using a home-built CFR,29 which helps to deliver and maintain a constant low supersaturation by continuously feeding the precursor solution into the reaction system.27 We also intentionally used a low concentration of Ga(NO3)3·xH2O (5 mM) to achieve a low supersaturation. Note that even lower concentrations, e.g., 2.5 mM, only yielded sparse particles (Figure S1, Supporting Information). This means that the concentration of 5 mM is sufficiently low but could still enable 7280

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Figure 4. Structural characterization of the converted porous α-Ga2O3 NR arrays. (A) PXRD pattern, (B) SEM, (C) TEM, and (D) high-resolution TEM (HRTEM) images of the converted porous α-Ga2O3 NRs. The inset of C shows the indexed selected area electron diffraction (SAED) pattern taken from the region marked by a rectangle in C. The asterisks (*) in A mark the diffraction peaks from the FTO substrate.

Figure 5. Structural characterization of the porous ZnGa2O4 NT arrays converted from porous α-Ga2O3 NRs. (A) PXRD pattern, (B, C) SEM and (D, E) TEM images, and (F) SAED pattern of the as-converted porous ZnGa2O4 NTs. (G, H) HRTEM images of the nanoparticles that make up the ZnGa2O4 NTs. (I) HADDF-STEM image of an individual NT and the corresponding EDS mapping showing the distribution of O, Zn, and Ga elements. The asterisks (*) in A mark the diffraction peaks from the FTO substrate.

above strongly indicate that the growth of α-GaOOH NRs is driven by dislocations. Therefore, the continuous feeding of precursor that is achieved by a CFR would effectively maintain the constant low supersaturation and thus promote the SDD growth. The CFR not only helps to deliver and maintain a constant low supersaturation but also allows “real-time” adjustment of the experimental conditions since it is a relatively open system. We can conveniently vary the precursor concentration, the flow rate, and the reaction time, thus controlling the length of the αGaOOH NR product. We first examined the influence of initial Ga3+ concentration on the morphology of the α-GaOOH NRs. As we increased the Ga3+ concentration from 5 (Figure 2D) to 7.5 mM (Figure S3A), the length of NRs increased from 2 to 3.5 μm. However, further increase in Ga3+ concentration (e.g., 10 and 15 mM) led to slightly shorter NRs but denser arrays (Figure S3B,C). This may be due to the fact that, at high starting concentrations, the nucleation density is higher. This means there are many more sources to deplete the precursor more quickly during the growth, resulting in shorter NRs. We

generally due to the large stress and strain created in the NRs by screw dislocations.26,27,29,41,42 Besides solid NRs, we also observed many hollow NTs (Figure 3C), which is a signature of the dislocations-driven grown materials. The strain energy associated with screw dislocations is related to the magnitude of the Burgers vector of the screw dislocation. With increasing Burgers vector magnitude, the strain energy within the crystal eventually exceeds the surface energy required for creating a new inner surface and therefore causes the dislocation core to hollow out.26,27,29,41,42 Indeed, small voids near the dislocation core can be observed (Figure 3D,E), indicating the beginning of NT formation. Note that we did not observe these void or hollow features in the static solution grown α-GaOOH sample, further confirming the SDD growth of the CFR grown αGaOOH NRs. Furthermore, we also observed dark contrast lines associated with dislocations26,27,29,41,42 going through the center of NRs (Figure 3F). Although we were not able to perform detailed structural analysis on these screw dislocations to assign their Burgers vectors due to the instability of αGaOOH under electron beam, all of the features presented 7281

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Figure 6. Schematic illustration of the conversion process from α-Ga2O3 NRs to ZnGa2O4 NTs due to the nanoscale Kirkendall effect together with the SEM images of the corresponding examples observed. (A) α-Ga2O3 NRs, (B) NTs with voids at interface, (C) NTs with skeletal bridges, (D) NTs with solid particles at interface, and (E) completely hollow NTs.

further investigated the impact of flow rate on the morphology of the product. Starting with 7.5 mM Ga3+ precursor solution, the α-GaOOH NRs can grow up to 1 μm in length at a relatively low flow rate of 0.1 mL min−1 (Figure S3D). At higher flow rates, the morphology of the as-grown α-GaOOH NRs does not show any significant changes, but the length of the NRs becomes much longer, up to 1.9 μm at 0.5 mL min−1 (Figure S3E) and 3.5 μm at 1.0 mL min−1 (Figure S3A), respectively. This is not surprising because a higher flow rate means more precursor is involved in the growth (note that the nucleation density does not change as it is determined by the initial precursor concentration). Meanwhile, at high flow rates, the mass transport is faster, which means a greater degree of collisions in classical chemical kinetics. Further, since the growth of α-GaOOH NRs is driven by dislocations, it is expected that the NR length can also be easily tuned by varying the reaction time when the precursor concentration is constant. Indeed, as we increased the reaction from 4 (Figure S3F) to 24 h (Figure S3A), the length of the NRs increased from 1.2 to 3.5 μm. These results clearly demonstrate the advantages and efficacy of the CFR in controlling the length of NRs. Conversion of α-GaOOH to Porous α-Ga2O3 NR Arrays. The α-GaOOH NRs can be readily converted into porous α-Ga2O3 NRs by annealing in air at 450 °C for 4 h. PXRD (Figure 4A) confirms that the annealed NRs are hexagonal α-Ga2O3 (JCPDS No. 06-0503, space group R3̅c, a = 4.979 Å, c = 13.429 Å) with high purity. No additional peaks from α-GaOOH are observed. The converted α-Ga2O3 NRs preserve the alignment and morphology of the α-GaOOH precursor NRs very well (Figure 4B). Depending on the αGaOOH NR precursors, the width of the α-Ga2O3 NRs ranges from 200 to 500 nm. TEM image (Figure 4C) reveals that many pores were introduced to the α-Ga2O3 NRs during the conversion, which is commonly seen for other metal hydroxide precursor/annealed oxide systems.41,67 This is due to the volume contraction associated with the loss of water and the transformation from low density α-GaOOH (5.23 g cm−3) to the denser α-Ga2O3 (6.44 g cm−3). The selected area electron diffraction (SAED) pattern (Figure 4C, inset) also confirms the phase, and the lattice spacing measured from the highresolution TEM (HRTEM) image (Figure 4D) matches well with that of α-Ga2O3 (012) planes. Conversion of α-Ga2O3 Nanorods to Porous ZnGa2O4 Nanotube Arrays. We then further converted the porous αGa2O3 NR arrays to porous ZnGa2O4 NT arrays in a solution process. Because of the small width of the NRs, the typical challenges with slow ionic diffusion in oxides are much less of an issue here. For a typical reaction, the porous α-Ga2O3 NRs

on FTO substrate were immersed into an aqueous solution containing Zn(CH3COO)2 and hydrothermally heated at 200 °C for 24 h (see details in the Materials and Methods section). The PXRD pattern (Figure 5A) clearly shows the product to be the cubic ZnGa2O4 phase (JCPDS No. 38-1240, space group Fd3̅m, a = 8.3349 Å). The low-magnification SEM image (Figure 5B) shows that the overall 1D structure and morphological homogeneity of the precursor α-Ga2O3 NRs are preserved, but interestingly, the NRs hollow out and form NTs. No peeling off from the FTO substrate was observed for ZnGa2O4 NTs after the conversion. Close examination (Figure 5C) reveals that there are three typical types of hollow nanostructures, i.e., NTs with voids at the core−shell interface (as indicated by arrow 1), NTs with voids and solid “skeletal bridges” (as indicated by arrow 2), and completely hollow NTs (as indicated by arrow 3). More SEM images showing these structural characteristics can be found in Figure S4 in the Supporting Information. We also found that the inner diameters of these NTs were generally smaller than the width of the precursor α-Ga2O3 NRs. A similar phenomenon was observed in Kirkendall effect-driven hollow Co3S4 nanoparticles, wherein the size of the final void was noticeably smaller than the particle size of the Co precursor.68 These observations suggest that the ZnGa2O4 NTs were likely formed following the Kirkendall effect.69,70 TEM (Figure 5D,E) further confirms the porous hollow structure and reveals that these NTs were composed of crystalline nanoparticles with a typical grain size of ∼30 nm (Figure 5G). The corresponding SAED pattern (Figure 5F) can be indexed as the cubic ZnGa2O4. The orderly arranged spots suggest that the as-converted sample is quasi-single-crystalline, i.e., all nanoparticles within a NT have the same crystallographic orientation. The lattice-resolved HRTEM (Figure 5H) shows lattice spacing of 0.251 nm, which can be assigned to the (311) planes of cubic ZnGa2O4, further confirming the converted product is ZnGa2O4. The EDS elemental mapping clearly reveals that the Zn and Ga are distributed homogeneously in an individual ZnGa2O4 NT. These results suggest that the low-temperature solution conversion method is effective in producing ZnGa2O4 NT arrays that are otherwise difficult to synthesize directly. It also avoids the use of vapor conversion or solid-state conversion reaction (see an example of CuMn2O471) at high temperatures, which facilitates large-scale preparation with low cost and high throughput. The nanoscale Kirkendall effect has been widely utilized for the preparation of hollow nanostructures69,70,72,73 but has been less commonly reported for complex compounds (e.g., ternary metal oxides), especially through solution conversion methods. 7282

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Figure 7. Structural characterization of the porous ZnGa2O4 microcubes converted from α-GaOOH NRs. (A) PXRD pattern, (B, C) SEM images, (D) TEM image, (E) SAED pattern, and (F) HRTEM image of the as-converted porous ZnGa2O4 microcubes. (G) HADDF-STEM image of an individual microcube and the corresponding EDS mapping showing the distribution of O, Zn, and Ga elements. The asterisks (*) in A mark the diffraction peaks from the FTO substrate.

also proceed via direct dissolution through the pores or gaps and eventually lead to completely hollow NTs (Figure 6E). It is worth mentioning that unlike spherical nanocrystals (in which voids grow isotropically), 1D nanostructures also allow the material transport along the longitudinal axis, which may result in uneven growth of the voids, therefore leading to the formation of segmented tubes (partly hollow and partly solid).57 Conversion of α-GaOOH Nanorods to ZnGa2O4 Microcubes. It has been reported that ZnGa2O4 nanoparticles could also be formed by reacting α-GaOOH particles with Zn2+ ions under hydrothermal conditions.53,76 Thus, we further attempted to use α-GaOOH NR arrays as the precursor to directly convert them to ZnGa2O4. The PXRD (Figure 7A) confirms that cubic ZnGa2O4 phase was formed after the conversion, and the sharp diffraction peaks indicate the good crystallinity of the product. Interestingly, the product did not inherit the NR morphology of the α-GaOOH precursor; instead they are microcubes with a typical size of 1 μm (Figure 7B,C). The TEM image (Figure 7D) and the corresponding SAED pattern (Figure 7E) further confirm the structure and indicate that the resulting ZnGa2O4 microcubes are single crystalline. Figure 7F shows the HRTEM image taken from the area marked by a rectangle in Figure 7D. The lattice spacing of ∼4.81 Å corresponds to the d spacing of (111) lattice planes of cubic ZnGa2O4. The EDS mapping (Figure 7G) reveals the homogeneous distribution of Zn and Ga elements throughout an individual ZnGa2O4 microcube. Clearly, the direct solution conversion of α-GaOOH NRs is also effective in producing ZnGa2O4, but it is interesting that a completely different morphology is formed, in contrast to the NT product when porous α-Ga2O3 NRs were used as the precursor. To understand why totally different morphologies were observed from conversion using different precursors, we further examined the conversion mechanisms. Even though the same hydrothermal conditions were employed for these two reactions, it seems that direct conversion and Kirkendall effect did not occur when α-GaOOH NRs were used as the precursor. Therefore, the observed disparity in morphology of the converted ZnGa2O4 products has to be due to the different structures and properties of the α-GaOOH and α-Ga2O3

The formation of Kirkendall voids is due to the nonreciprocal mutual ionic diffusion: the unequal flow of core and shell materials generates a net inward vacancy diffusion, which leads to the formation of voids and eventually hollow structures.69,70,72,73 For the conversion of α-Ga2O3 to ZnGa2O4 here, the α-Ga2O3 NRs serve not only as the Ga source but also as a sacrificial template. Putative chemical reactions involved during the conversion could be described as the following: α‐Ga 2O3 + Zn 2 + + 2OH− → ZnGa 2O4 + H 2O

(3)

At the initial stage, the porous structure of the α-Ga2O3 NRs enables facile diffusion and provides sufficient surface sites so that the Zn2+ ions could effectively adsorb on the surface of αGa2O3 NRs. It is known that ZnGa2O4 is thermodynamically more stable than a physical mixture of ZnO and β-Ga2O3 (note that β-Ga2O3 is the most stable among the five polymorphs of Ga2O3).74 Therefore, it is possible that the adsorbed Zn2+ ions react with α-Ga2O3 to form a ZnGa2O4 layer at the solid−liquid interface (eq 3). The PXRD pattern of the product after a conversion reaction for a shorter amount of time (1 h) clearly shows a mixture of both the α-Ga2O3 and the ZnGa2O4 phases (Figure S5). Such an initially formed ZnGa2O4 layer then prevents the direct chemical reaction of α-Ga2O3 with Zn2+, and thus further reaction depends on the diffusion of Ga3+ and Zn2+ ions through the solid interface. The Ga3+ ions diffuse outward faster than the Zn2+ ions diffuse into the α-Ga2O3 core (Figure 6A), thus creating a net outward movement of ions and simultaneously resulting in a flow of vacancies to the vicinity of the interface and eventually the formation of small voids (Figure 6B). When these voids contact the inner surface of the ZnGa2O4 layer, the α-Ga2O3 NR core would continue hollowing out as the Ga atoms diffuse outward along the skeletal bridges (Figure 6C). Such solid bridges connecting the core and the shell could facilitate the transport of the core materials outward and were usually found during the intermediate state of Kirkendall effect-induced hollow nanostructures.73,75 The hollow interior enlarges as the outward diffusion of the core material continues (Figure 6D). Because of the porous structure of α-Ga2O3 NRs, the adsorption of Zn2+ is inhomogeneous, and thus pores or gaps may form during the growth of the initial ZnGa2O4 layer. The material exchange can 7283

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Chemistry of Materials precursors. The α-GaOOH is amphoteric and could dissolve in alkaline solutions to form soluble [Ga(OH)4]− (for reference, see the Pourbaix diagram of the Ga−O−H2O system),77 which would subsequently react with Zn2+ ions to form ZnGa2O4 nanocubes. Thus, the reaction between α-GaOOH and Zn(CH3COOH)2 could proceed as follows:

precursor would go through a dissolution−recrystallization process and form ZnGa2O4 microcubes. Due to the generality of dislocation-driven nanomaterials growth, this facile route could be extended to synthesize 1D (NRs, NWs, and NTs) arrays of a wide variety of metal oxide/ hydroxides, for instance, α-FeOOH and Co(OH)2 NWs.38,66 This facile approach to grow vertical NR arrays of α-GaOOH (and converted Ga2O3) and other metal oxides/hydroxides can enable a myriad of applications of these materials, just as the vertical NR arrays of ZnO17,20 and TiO218 have. Furthermore, these 1D nanomaterials of metal oxide/hydroxides could serve as precursors and be readily converted into ternary or more complex metal oxides, such as spinel oxides (as demonstrated herein) and perovskites, as well as metal chalcogenides,65 nitrides,79,80 and phosphides, through such simple solution conversion with the preservation of the nanomorphology. The facile solution method can enable the low-cost and large-scale production of the 1D nanostructures of these more complex materials and promote their utility in practical large-scale applications in electrocatalysis, solar energy conversion, or electrochemical storage.

α‐GaOOH + H 2O + OH− → [Ga(OH)4 ]− (soluble) (4)

2[Ga(OH)4 ]− + Zn 2 + → ZnGa 2O4 + 4H 2O

(5)

Due to this dissolution−recrystallization process, the morphology of the ZnGa2O4 products does not depend on the morphology of the α-GaOOH precursor. The cubic morphology of the product reflects the intrinsic crystal habits of the cubic spinel ZnGa2O4 structure. We noted that ZnGa2O4 nanocubes could also be obtained by hydrothermally reacting α-GaOOH of other morphologies, such as nanoplates or microrods, with Zn(CH3COOH)2, as reported previously.53 In contrast, although α-Ga2O3 is also slightly amphoteric, it is much more stable than α-GaOOH and dissolves very slowly in base; thus, it could serve as both a sacrificial template and a gallium source for the conversion into ZnGa2O4 as discussed earlier. In fact, the GaO6 octahedra in ZnGa2O4 have bonding configurations comparable to those of GaO6 octahedra in αGa2O3 (Figure S6). Their average Ga−O bond lengths were also found to be very close (1.9980 Å for ZnGa2O4 and 1.9989 Å for α-Ga2O3).78 Moreover, the interplanar spacing of some lattice planes of ZnGa2O4 is very close to that of α-Ga2O3.78 For example, the d(111̅ ) spacing of 4.81 Å is almost equal to 2 × d(1̅20) spacing (4.98 Å) of α-Ga2O3. Such structural similarity facilitates the conversion into ZnGa2O4 on the surface of α-Ga2O3 and further allows the preservation of the 1D geometry of α-Ga2O3 precursor during conversion. The porous structure of α-Ga2O3 can alleviate the volume change during the phase transformation and thus also helps the smooth conversion. These results suggest that the solution conversion can serve as a powerful route to synthesize nanomaterials of ternary compounds with specific morphology that otherwise could be difficult to synthesize directly, which provides the opportunity to evaluate their properties and performance and enable their large-scale applications.

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MATERIALS AND METHODS

Reagents. Gallium(III) nitrate hydrate [Ga(NO3)3·xH2O, 99.9%, CAS no. 69365-72-6], urea [CO(NH2)2, CAS no. 57-13-6], and zinc acetate [Zn(CH3COO)2, 99.99%, CAS no. 557-34-6] were purchased from Sigma-Aldrich. Ethanol (200 proof, CAS no. 64-17-5) was purchased from Decon Laboratories, Inc. All chemical reagents were used as received. Continuous Flow Synthesis of α-GaOOH NR Arrays. A homebuilt continuous flow cell reactor (CFR), whose setup has been previously detailed,29 was used to synthesize α-GaOOH NR arrays. For the preparation of growth substrate, a piece of FTO-coated glass substrate (size: 1 × 2 cm) was mounted onto an 11 mm magnetic stage using a double-sided tape. Then the substrate was mounted upside down near the exit port of a clean jacketed CFR column. Subsequently, the CFR column was connected to the heating water supply and return lines connected to a circulating water bath set to 95 °C, as well as the precursor feed and waste lines. To prepare a feed precursor solution for a typical reaction, 2.56 g of Ga(NO3)3·xH2O (from 10 mmol with x = 0 to 9.35 mmol with x = 1) and 1.2 g (20 mmol) of urea were dissolved in 2.0 L of nanopure water (Thermo Scientific, Banstead Nanopure, 18.2 MΩ cm) to result in a clear solution with a nominal 5 mM Ga3+ concentration [assuming x = 0 in Ga(NO3)3·xH2O]. After the CFR column was warmed to 95 °C, the precursor solution was pumped through the CFR at a rate of 1.0 mL min−1 for 24 h. Once completed, the heating and flow were stopped and the substrate was removed from the CFR column, rinsed with water and ethanol, and dried in a flow of N2 gas. Static Solution Growth of α-GaOOH NR Arrays. The static hydrothermal reaction employed the same concentration and temperature conditions as those in a CFR reaction. In a typical reaction, 0.192 g of Ga(NO3)3·xH2O and 0.09 g of urea were dissolved in 100 mL of water to result in a solution with a nominal 5.0 mM Ga3+ concentration, which was heated in a sealed glass bottle at 95 °C for 24 h. The FTO substrate was immersed in the solution facing down just below the air−water interface. At the end of the reaction, the substrate was rinsed with water and ethanol and dried with a stream of N2 gas. Conversion of α-GaOOH NRs to Porous α-Ga2O3 NR Arrays. The α-GaOOH NR arrays were annealed in air at 450 °C for 4 h with a heating rate of 1 °C min−1 to yield porous α-Ga2O3 NR arrays. Conversion of α-Ga2O3 NRs to Porous ZnGa2O4 NT Arrays. In a typical synthesis, 46 mg (0.25 mmol) of Zn(CH3COO)2 was dissolved in 15 mL of water and then transferred into a Teflon-lined stainless steel autoclave with a capacity of 20 mL. The as-obtained αGa2O3 NR arrays on FTO (1 × 2 cm) were used as the precursor as well as the growth substrate. After the sealed autoclave was heated at

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CONCLUSIONS We demonstrate a simple and general strategy to grow dislocation-driven 1D NR arrays of metal oxyhydroxides on FTO substrates and further convert them into porous metal oxide and ternary metal oxide arrays using the α-GaOOH/αGa2O3/ZnGa2O4 system as examples. Vertical arrays of αGaOOH NRs were synthesized in solutions through the controlled hydrolysis of Ga3+ ions using a CFR. Compared to static reactions, the continuous flow reactions enable the growth of long and high aspect ratio α-GaOOH NRs up to 3.5 μm with controlled length. Structural characterization confirms that the growth of α-GaOOH NRs follows the screw dislocation-driven growth mechanism. The α-GaOOH NR arrays can be converted to porous α-Ga2O3 NR arrays by thermal annealing. Furthermore, ZnGa2O4 NT arrays were formed through the solution conversion of α-Ga2O3 NRs due to the nanoscale Kirkendall effect. However, when α-GaOOH NRs were used as the precursor, because α-GaOOH is more soluble than α-Ga2O3 in alkaline solutions, the α-GaOOH NR 7284

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Chemistry of Materials 200 °C for 24 h, the substrate was removed, rinsed with water and ethanol, and dried in a flow of N2 gas. Conversion of α-GaOOH to ZnGa2O4 Microcube Thin Films. The procedure was the same as that for the conversion of α-Ga2O3 NR arrays to porous ZnGa2O4 NT arrays described above except that αGaOOH NR arrays on FTO substrates were used as the precursor. Structural Characterization. Powder X-ray diffraction (PXRD) measurements were performed on a Bruker D8 Advance powder X-ray diffractometer using a Cu Kα radiation. For scanning electron microscopy (SEM), samples were imaged using a LEO Supra 55 VP field emission scanning electron microscope. For transmission electron microscopy (TEM) observation, powder that was scraped from FTO substrates was dispersed into ethanol and sonicated for 5 min. Then a few drops of suspension were casted onto lacey carbon supported TEM grids. TEM imaging was performed on a FEI Tecnai T12 transmission electron microscope operated at an accelerating voltage of 120 kV. High-resolution TEM (HRTEM) was carried out on a FEI Titan transmission electron microscope operated at an accelerating voltage of 200 kV. Energy-dispersive X-ray spectroscopy (EDS) elemental mapping was performed on the same instrument in a scanning TEM (STEM) mode.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.7b01930. Additional figures and data (PDF)

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AUTHOR INFORMATION

Corresponding Author

*(S.J.) E-mail: [email protected]. ORCID

Hanfeng Liang: 0000-0002-1778-3975 Brandon K. Lamb: 0000-0002-3141-4439 Song Jin: 0000-0001-8693-7010 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This research is supported by NSF Grant DMR-1508558. S.J. also thanks the UW-Madison H.I. Romnes Faculty Fellowship for support. H.L. thanks the China Scholarship Council for support. Z.W. thanks the National Natural Science Foundation of China (Grant 51372212) for support.

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