Spiny Rhombic Dodecahedral CuPt Nanoframes with Enhanced

Jun 16, 2017 - It was surprising that galvanic replacement was dominant during the Pt coating and made most nanocrystals become hollow structures (Fig...
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Spiny Rhombic Dodecahedral CuPt Nanoframes with Enhanced Catalytic Performance Synthesized from Cu Nanocube Templates Lian-Ming Lyu,†,# Ya-Chuan Kao,†,# David A. Cullen,§ Brian T. Sneed,∥ Yu-Chun Chuang,‡ and Chun-Hong Kuo*,†,⊥ †

Institute of Chemistry, Academia Sinica, Taipei 11529, Taiwan Materials Science and Technology Division and ∥Center for Nanophase Materials Sciences, Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831-6496, United States ‡ National Synchrotron Radiation Research Center, Hsinchu 30076, Taiwan ⊥ Institute of Materials Science and Engineering, National Central University, Jhongli 32001, Taiwan §

S Supporting Information *

ABSTRACT: Platinum was coated on the surfaces of copper nanocubes to form Cu−CuPt core−alloy−frame nanocrystals with a rhombic dodecahedral (RD) shape. Co-reduction of Pt2+ ions and residual Cu+ ions in the supernatant of the Cu nanocube solution followed by the interdiffusion of Cu and Pt atoms over the core−shell interface allowed their formation. Growth in the ⟨100⟩ directions of the {100}-terminated Cu nanocubes resulted in the {110}-faceted rhombic dodecahedra. By the introduction of additional Pt precursor, the {100} vertices of the Cu−CuPt RD nanocrystals could be selectively extended to form spiny CuPt RD nanocrystals. After removing the Cu core template, both CuPt alloy RD and spiny CuPt alloy RD nanoframes (NFs) were obtained with Pt/Cu ratios of 26/ 74 and 41/59, respectively. Abundant surface defects render them highly active catalysts due to the open frame structure of both sets of NFs. The spiny RD NFs showed superior specific activity toward the oxygen reduction reaction, 1.3 and 3 times to those of the RD NFs and the commercial Pt/C catalysts, respectively. In 4-nitrophenol reduction, both NFs displayed better activity compared to commercial Pt NPs in the dark. Their activities were improved ∼1.3 times under irradiation of visible light, attributed to the effect of LSPR enhancement by the Cu-rich skeleton.

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INTRODUCTION Extraordinary efforts have been launched into the development of devices for energy conversion and storage, such as in proton exchange membrane fuel cells (PEMFCs) and metal−air (M− air) batteries.1−6 Unfortunately, reaching the goal of widespread adoption of these technologies is hindered by the high cost of catalytic and storage materials. Low abundance and high cost of Pt is of primary concern with regard to PEMFCs. Pursuing higher Pt mass activity with more efficient catalysts has received much interest recently to alleviate this problem. According to US DOE targets, the catalysts used in commercial PEMFCs should reach 0.44 A/mgPt for the oxygen reduction reaction (ORR) by 2020.7,8 Pt blended with earth-abundant elements has been investigated, i.e., PtNi alloys, which were found superior for catalyzing the ORR.9−12 Apart from alloying Pt with Ni, or other late transition metals, such as Co, Cu is also an excellent potential alternative as it is similar in abundance and price. Furthermore, nanoscale Cu possesses localized surface plasmon resonance (LSPR) that induces strong absorption of visible light.13−15 LSPR is an important phenomenon that can be applied to enhance optical sensing and catalytic performance. Accordingly, blending Pt with Cu is © 2017 American Chemical Society

promising to form nanocrystals with integrated functions of plasmonics and catalysis, comparable to, but more cost-effective than other plasmonic combinations such as AuPd.16,17 Numerous works of CuPt binary nanocatalysts have been demonstrated with regard to control of morphology and composition.18−25 However, upon surveying the recent development of Pt-based alloy nanostructures, those of open 3D frameworks with interior and exterior surfaces exposed for high surface area and molecular accessibility provide superior catalytic performance in both oxidation and reduction reactions.24−27 For example, Jia et al. fabricated excavated rhombic dodecahedral PtCu3 alloy nanocrystals which had ultrathin nanosheets of high-energy {110} facets, more active than octahedral ones constructed with {111} facets for the formic acid oxidation reaction (FOR).27 Ding and co-workers synthesized rhombic dodecahedral (RD) CuPt nanoframes with 3D-accessible surface areas for the methanol oxidation reaction (MOR). The nanoframes generated 2.77 and 3.9 times Received: April 14, 2017 Revised: June 7, 2017 Published: June 16, 2017 5681

DOI: 10.1021/acs.chemmater.7b01550 Chem. Mater. 2017, 29, 5681−5692

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Chemistry of Materials

color gradually turned from greenish blue to red, indicating the formation of Cu nanocubes. Synthesis of Cu−CuPt Core−Frame RD Nanocrystals. After getting Cu nanocubes, the platinum precursor was added during the stirring and heating to synthesize core−frame RD nanocrystals. The platinum precursor was prepared by dissolving 0.0076 g (0.0193 mmol) of Pt(acac)2 in a mixed solution of 1 mL of toluene and 1 mL of OAm in a glass vial at room temperature. The reaction temperature was held at 200 °C for 60 min followed by cooling down to room temperature naturally. To collect nanocrystals, 10 mL of n-hexane was added into the solution, and the products were collected by centrifuging at 6000 rpm for 3 min. The supernatant solution was discarded to remove the unreacted precursors. In the same way, the precipitation was washed by centrifuging in 10 mL of n-hexane at 6000 rpm for 3 min, which was repeated five times to remove the capping agent. Finally, the precipitate was redispersed in 1 mL of toluene for further use. Synthesis of Spiny Cu−CuPt Core−Frame RD Nanocrystals. The synthesis procedure is the same to that of making Cu−CuPt core−frame RD nanocrystals except that the Pt(acac)2 amount was increased to 0.0152 g (0.0386 mmol) and the reaction time was extended to 110 min. Nanocrystals obtained at different reaction times were monitored to understand the growth mechanism. Treatment for Surface Cleaning. After synthesis, both core− frame nanocrystals had to be treated to remove the capping agents on their surfaces toward further catalytic reactions. The cleaning solution was prepared by dissolving 10 mg (0.264 mmol) of NaBH4 in 10 mL of DI water in an ice bath. In typical, 1 mL of the toluene solution containing Cu−CuPt nanocrystals was diluted with 4 mL of ethanol (95%) followed by centrifuging at 4000 rpm for 3 min. Next, the supernatant was removed and the precipitate was diluted with 5 mL of NaBH4 solution and well dispersed. The solution was left undisturbed for 30 min at room temperature; then, the Cu−CuPt core−frame nanocrystals were collected by centrifuging at 4000 rpm for 3 min. Afterward, the precipitation was washed with 1 mL of a mixed solution of ethanol/water (volume ratio equal to 1) for three times to wash the remained chemicals off. Finally, the precipitate was collected and redispersed in 1 mL of ethanol. Preparation of CuPt Nanoframes. A total of 100 μL of ethanol solution containing NaBH4-treated Cu−CuPt core−frame nanocrystals was taken to mix with 900 μL of 0.5 M HNO3, which was sonicated for 30 min to etch the Cu cores off. The etched products were then collected by centrifuging at 7000 rpm for 5 min. The collected products were washed with 1 mL of a mixed solution of ethanol/water (volume ratio equal to 1) three times to remove the remaining chemicals. Lastly, the products were redispersed in 100 μL of ethanol. Electrochemical Measurement. To prepare the working electrodes for oxygen reduction reaction (ORR), both of sets of NFs were loaded onto carbon black (Vulcan XC-72, CABOT) as the supports with a metal amount of 46 wt % (determined by ICP-OES). Typically, 1.7 mg of the nanoframe catalysts and 2 mg of Vulcan XC-72 were dispersed together in 2 mL of isopropanol (IPA) under ultrasonication for 1 h. A 7 μL aliquot of the suspension was dropped on a glassy carbon rotating disk electrode (GC-RDE) with a geometric area of 0.07 cm2 and naturally dried under ambient conditions. Then, 1 μL of 0.1 wt % Nafion/IPA solution was dropped on the electrode and further dried for 5 min at room temperature. The commercial Pt/C catalyst (20 wt %) was used for comparison. Typically, 5 mg of the commercial Pt/C catalyst was dispersed in 1.25 mL of 0.1 wt % Nafion/IPA solution with ultrasonication for 1 h to prepare an ink. Then 5 μL of the ink was dropped on a GC-RDE and naturally dried at ambient conditions. For the ORR, the working electrodes were connected to a CHI 705E potentiostat (CH Instruments). Ag/AgCl was used as the reference electrode, and a Pt wire was used as the counter electrode. The measured potential EAg/AgCl was converted versus the reversible hydrogen electrode (RHE) scale according to the following equation using the known potential of Ag/AgCl (3 M KCl) of 0.21 V vs RHE.

higher specific activity and Pt mass activities over that of the commercial Pt/C catalysts.26 More recently, Luo et al. made concave CuPt octopod nanoframes.28,29 They achieved 7.2 times higher MOR and 20 times higher ORR enhancement in Pt mass activities over that of the commercial Pt/C catalysts. This performance is believed to arise from high-index crystal faces on the surfaces of the octopods. In this work, we utilize a two-step Cu-templated method to create Cu−CuPt core−frame RD and spiny RD nanocrystals. This unprecedented strategy allows us to study the growth mechanism by decoupling effects of galvanic replacement. Highly monodisperse Cu nanocubes were first synthesized followed by epitaxial deposition of Pt atoms to form the RD nanocrystals. These nanocrystals were Cu−CuPt core−frame structures in which the cores remained pure Cu while the frames CuPt alloy. By increasing the Pt precursor amount, further growth at the {100} vertices of the RD frames took place, which led to the formation of the vertex-extended spiny Cu−CuPt core−frame RD nanocrystals. After their creation, the cubic Cu template was etched from the corresponding core to create CuPt alloy nanoframes (NFs). The two types of NFs were determined to be CuPt alloys by ICP-OES. Abundant defects such as terracing and adatoms on the surfaces of the NFs were observed through aberration-corrected STEM imaging. Such surface defects primarily occurred due to Cu leaching which rendered the NFs as highly active catalysts in both electrochemical oxygen reduction (ORR) and 4-nitrophenol (4-NP) reduction reactions. In ORR, the spiny RD NFs showed the best specific activity. However, the Pt mass activity presented a lower value than that of the regular NFs due to the higher content of Pt in the frameworks. Nevertheless, both of the NFs performed better than those of the commercial Pt/C catalysts. In 4-NP reduction, the two kinds of NFs also displayed much better activities than that of the commercial Pt catalysts in the dark. Furthermore, their activities were both improved at least 1.3 times as much under irradiation of visible light, attributed to the effect of SPR enhancement by the nanophase Cu domains in the Cu-rich skeletons. These findings serve as a benchmark and provide a new route for the fabrication of cost-efficient and function-added (LSPR) nanocatalysts.

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EXPERIMENTAL SECTION

Chemicals. Copper(I) bromide (CuBr, 98%, Acros Organics), platinum(II) acetylacetonate (Pt(acac)2, ≥97%, Sigma-Aldrich), tri-noctylphospine oxide (TOPO, 99%, Alfa Aesar), oleylamine (OAm, 80−90%, Acros Organics), n-hexane (C6H14, 95%, Macron Fin Chemicals), toluene (C6H5−CH3, ≥99.7%, Sigma-Aldrich), sodium borohydride (NaBH4, ≥98%, Sigma-Aldrich), perchloric acid (HClO4, 70%, Sigma-Aldrich), 4-nitrophenol (C6H5O3, 99%, Alfa Aesar), nitric acid (HNO3, ≥65%, Sigma-Aldrich), Pt/C catalysts (20 wt % on carbon black, Alfa Aesar), Pt catalysts (200 nm nanopowder, 99.9%, Aldrich) and Nafion (5 wt %, Aldrich) were used without further purification. Ultrapure deionized water (18.2 MΩ cm−1) was used for all solution preparations. Synthesis of Cu Nanocubes. In a typical synthesis, 0.0615 g (0.428 mmol) of CuBr and 0.4143 g (1.07 mmol) of TOPO were dissolved into 5 mL of OAm in a three-necked flask connected with a cooling condenser for refluxing. A thermocouple in a glass tube was placed for temperature measurement and a Teflon tube for N2 purging. The solution initially underwent a gentle N2 flow at room temperature to remove O2 with vigorous stirring for 30 min. The solution was then heated to 80 °C at the rate of 20 °C/min and kept at this temperature for 15 min. Afterward, it was further heated up to 200 °C at the same rate with refluxing and kept for 60 min. The solution

E RHE = EAg/AgCl + 0.059pH + 0.21 5682

DOI: 10.1021/acs.chemmater.7b01550 Chem. Mater. 2017, 29, 5681−5692

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Chemistry of Materials

Figure 1. (a) SEM, (b) bright-field (BF) TEM, and (c) HAADF-STEM images of the Cu−CuPt core−frame RD nanocrystals. (d) STEM-EDS spectrum images of Cu and Pt core−frame RD nanocrystals along two different orientations. (e) ac-HAADF-STEM image of a single core−frame RD nanocrystal and its corresponding crystal model. The scale bar represents 10 nm. (f) Atom-resolved ac-HAADF-STEM image of a selected vertex (dashed-line area in (e)). (g) Synchrotron X-ray diffraction pattern of the Cu−CuPt core−frame RD nanocrystals. The inset pattern is an enlarged pattern of the 111 peak with deconvoluted curves. Additional surface cleaning was carried out by cyclic voltammetry (CV) scanning in 60 mL of 0.1 M HClO4 aqueous solution at room temperature between 0.0 and 1.27 V (vs RHE) at a sweep rate of 100 mV/s for 100 scans. After consecutive scans stabilized, proton adsorption/desorption CV curves were recorded in an Ar-saturated 60 mL of 0.1 M HClO4 solution at room temperature across the range of 0.0 and 1.27 V at 50 mV/s. The specific electrochemical active surface area (ECSA) of catalysts was calculated according to the charge associated with the adsorption of protons by integrating the area in the region from 0.02 to 0.35 V after double-layer calibration with a reference value of 210 μC/cm2 for the desorption of a monolayer of protons from the Pt surface. The ORR was performed with catalysts on a GC-RDE rotating at 1600 rpm under linear-sweeping voltammetry (LSV) from 0.1 to 1.1 V (vs RHE) in a 90 mL of O2saturated 0.1 M HClO4 solution at 10 mV/s. The accelerated deterioration test (ADT) was carried out under a sweep rate of 100 mV/s between 1.1 and 0.45 V for 1200 cycles. Apart from 1600 rpm, the measurements under disk rotation rates of 100, 225, 400, and 900 rpm were also carried out. The apparent number of electrons transferred (n) in the ORR reaction on the carbon catalysts was determined by the Koutecky−Levich equation given by

catalysts, 0.2 mL of 0.1 M NaBH4, and 0.02 mL of 0.01 M 4nitrophenol was prepared in a quartz cuvette with an inner edge length of 1 cm. The cuvette was placed in a water bath controlled at a constant temperature of 32 °C throughout the reaction. The reaction was carried out in the dark and under the irradiation of visible light using a Xe lamp providing a true incident light power of 120 mW to the sample at the irradiated distance of 25 cm. An HITACHI U-3310 UV−visible spectrometer was set to monitor the absorbance in the range of 250 to 550 nm where the λmax of 4-NP is at 400 nm and 4-AP at 305 nm. For comparison, commercial Pt nanopowder (Aldrich no.771937) with an average size of ∼200 nm was used as a reference. The calibration line of 4-NP was carried out (Figure S27) to obtain the extinction coefficient with the Beer’s law for calculating the specific conversion of 4-NP. Characterization. To prepare samples for SEM, 1 μL of concentrated sample solution was dropped on 2 × 2 mm2 silicon wafers which were slowly dried in ambient conditions. The TEM samples was prepared by dropping 1 μL of 100-fold-diluted sample solution onto the carbon-coated gold grids with slow drying at ambient conditions. Centrifuging was done using an Eppendorf Centrifuge 5804 and a Thermo Scientific Heraeus Pico 17. SEM images were recorded by a ZEISS ULTRA PLUS equipped with the OXFORD EDX detector, operated at the accelerating voltage of 10 keV. Lowand high-magnification TEM bright-field images were taken by a JEOL JEM-2100F microscope operating at 200 kV. High-angle annular darkfield (HAADF)-STEM imaging and energy dispersive X-ray spectroscopy (EDS) spectrum images were acquired using an FEI Talos F200X STEM operated at 200 kV, which takes advantage of a system of four orthogonal EDS detector units for high collection efficiency. Aberration-corrected (ac-) HAADF-STEM imaging was conducted on an aberration-corrected JEOL 2200FSTEM/STEM operated at 200 kV. UV−vis absorption spectra were measured on a HITACHI U3310 spectrophotometer. The X-ray diffraction experiments were performed at BL01C2 in the National Synchrotron Radiation Research Center (NSRRC). The diffraction data were collected using 18 keV Xrays (0.68888 Å in wavelength) and Mar345 image plate detector with Debye−Scherrer geometry. The patterns were converted by a GSAS-II

1/J = 1/JL + 1/JK = 1/(Bω1/2) + 1/JK

B = 0.201nFC0(D0)2/3 V −1/6 where J is the measured current density, JL is the diffusion-limited current density, JK is the kinetic current density, ω is the electrode rotation rate (rpm), F is the Faraday constant (96485.33 C mol−1), C0 is the bulk concentration of O2 (1.38 × 10−3 mol L−1 for 0.1 M HClO4 solution), D0 is the diffusion coefficient of O2 (1.9 × 10−5 cm2 s−1 for 0.1 M HClO4 solution), and V is the kinetic viscosity of the electrolyte (9.87 × 10−3 cm2 s−1 for 0.1 M HClO4 solution).30,31 Catalytic Reduction of 4-Nitrophenol. The reduction of 4nitrophenol (4-NP) to 4-aminophenol (4-AP) in an aqueous solution was performed to test the plasmon effect on the catalytic activity by monitoring the degradation of the 4-NP featured absorption peak at 400 nm. Typically, a solution containing 2.78 mL of DI water, 1 μg of 5683

DOI: 10.1021/acs.chemmater.7b01550 Chem. Mater. 2017, 29, 5681−5692

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Figure 2. (a) SEM and (b) BF TEM images of the CuPt RD NFs. (c−e) STEM-EDS spectrum images of Cu and Pt in a single NF with (f) HAADFSTEM image. The scale bar represents 10 nm. (g) Atom-resolved ac-HAADF-STEM image of a selected area (the dashed box in (f)). (g) Synchrotron X-ray diffraction pattern of the CuPt RD NFs compared with that of the Cu−CuPt core−frame RD nanocrystals. The inset pattern is the enlarged pattern of the 111 peak. program, and the angle calibration was performed according to LaB6 (SRM 660c) standard. Inductively coupled atomic emission spectroscopy (ICP-OES) analyses were performed on Varian 720-ES (Agilent Technologies).

along with OAm (Figure S2). Addition of TOPO also promotes the disproportionation as it brings down the required temperature.35 As a consequence, we synthesized highly monodispersed Cu nanocubes in an average size of 30.4 nm with a standard deviation of 10.5% (Figure S3a,b). The sizes of Cu nanocubes are tunable by changing the heating temperature. As shown in Figure S4, Cu nanocubes have the sizes of 30.97, 36.99, and 44.47 nm obtained in the syntheses at 220, 240, and 260 °C, respectively. It could be attributed to the accelerated rate of disproportionation which led to the size growth. In contrast, the disproportionation was not triggered if the heating temperature is lower than 200 °C (e.g., 180 °C), and thus no Cu nanocubes were obtained in an hour. For controlling the final-formed bimetallic nanocrystals in small size, the Cu nanocubes in 30.4 nm were always used as the templates for further experiments of Pt coating. Figure 1a is an SEM image showing well-defined rhombic dodecahedral (RD) nanocrystals obtained by adding Pt(acac)2 to the OAm solution of as-synthesized Cu nanocubes at 200 °C for an hour. The RD nanocrystals were verified as singlecrystalline structures by selected-area electron diffraction (SAED) (Figure S5a,b). The RD nanocrystals are uniform with an average edge length of 35.2 nm and a standard deviation of 8.5% (Figure S3c,d). Both bright-field (BF) TEM and HAADF-STEM images of the RD nanocrystals present clear differences in the contrast among their faces, edges, and vertices. This implies that the distribution of heavy Pt atoms were deposited preferentially at edges and vertices (Figure 1b,c). Figure 1d shows the STEM-EDS results of two RD nanocrystals viewed along [111] and [211] directions. The elemental maps of Cu and Pt reveal that the RD nanocrystals are core−frame structures. This supports the observations in TEM images and firmly validates that Pt indeed distributes heterogeneously. The crystal models of the two core−frame RD nanocrystals are shown in Figure S6. To carefully examine the crystal structures of the core−frame RD nanocrystals, diffraction of synchrotron radiation with an X-ray energy of 18 keV (λ = 0.68888 Å) was carried out. In Figure 1g, the resulting

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RESULTS AND DISCUSSION Pt Coating on Cu Nanocubes. The Cu nanocubes were synthesized by thermal reduction of CuBr in oleylamine (OAm) purged with N2 gas at a typical temperature of 200 °C, in which trioctylphosphine oxide (TOPO) was also added before heating. OAm is the chemical which functions as a solvent, a reducing and a capping agent which is widely used in the synthesis of metallic nanoparticles.32−34 The double bond (CC) in the chain of OAm acts as an electron donor capable of reducing metal ions. As a capping agent, OAm can confine particle size and modify the growth rates of certain crystal facets in the nanocrystals. As a solvent, it stabilizes particles without aggregation because of the repulsive force between the surfaces of nanocrystals. However, it has been demonstrated by Guo et al. that the shape-controlled formation of Cu nanocubes is mainly driven by the disproportionation of Cu+ ions.35 The disproportionation reaction, generating Cu0 atoms and Cu2+ ions due to the inter-Cu+ electron transfer at low temperature, allows the prompt nucleation of Cu seeds for further growth with the new-formed Cu0 atoms. Figure S1a and S1b show the color evolution of the reaction solutions from greenish blue to brownish red, suggesting the formation of Cu nanocubes from Cu+ ions. In Figure S1c, the blue supernatant of the as-prepared solution of Cu nanocubes after the first centrifuging reveals the existence of Cu2+ ions generated from the disproportionation. The transparent n-hexane supernatant in Figure S1d is evidence that the Cu nanocubes are not leached or dissolved again after the second centrifuging within n-hexane at ambient. Referring to the conclusion demonstrated by Wang and co-workers, TOPO, similar to TOP, has weak binding to the Cu surfaces.36 However, as shown in Figure S2, the Cu nanocrystals synthesized without TOPO formed as nonuniform truncated cubes or cuboctahedra, indicating TOPO had an auxiliary role in shaping the growing Cu nanocrystals into the cubic shape 5684

DOI: 10.1021/acs.chemmater.7b01550 Chem. Mater. 2017, 29, 5681−5692

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XRD peak shifts and ICP-OES analysis, confirm some Cu leaching of the frames and their surfaces, in addition to the cores. Preparation and Characterization of Spiny RD Nanocrystals. Highly monodispersed Cu−CuPt core−frame nanocrystals were typically obtained with RD shape when the synthesis was done at 200 °C for 1 h after addition of Pt(acac)2. To discover whether the RD shape came from the structural evolution of Cu nanocubes by heating or was caused by the addition of Pt(acac)2, a series of control experiments were carried out. Figure S9a is the result of Cu nanocubes obtained in the same condition except for a prolonged 2-h heating (no Pt added). As result, the Cu nanocubes remained in the cubic shape. This reveals that the presence of Pt atoms caused the difference in the particle shape. If the heating time was extended to 2 h in the synthesis after addition of Pt(acac)2 (3 h in total), the RD shape of the core−frame nanocrystals was retained, except for some growth at {100} vertices (indicated by the arrow in Figure S9b). It implies a potential in further growth at the vertices can be induced over long time if the Pt0 source is sufficient. Apart from this, the formation of some small particles also occurred (Figure S9b). Turning down the heating temperature to 170 °C for an hour after addition of Pt(acac)2 created the RD shape, but in a poor size distribution. We note here that no bulges at corners or small particles were observed (Figure S9c). However, heating at 230 °C gave a surprising result (Figure S9d). The core−frame RD nanocrystals formed showed growth at vertices in the ⟨100⟩ directions, while growth at the {111} vertices was not obvious (highlighted by arrows). It is important to note here that the RD shape contains these two different vertex types, where 4 and 3 rhombic facets meet, pointing in the ⟨100⟩ (6 total) and ⟨111⟩ (8 total) directions, respectively. The selective growth of {100} vertices is likely a result of the cube substrate exposing primarily {100} facets. In addition, small particles formed along with the spiny RD nanocrystals (in dashed circles). The mixture of spiny RD and small nanocrystals was treated with the nitric acid solution. As a consequence, the spiny RD NFs were obtained, but small particles remained (Figure S9e). These small particles were identified as Cu−Pt alloy nanocrystals by STEM-EDS, and thus this suggests they were able to resist corrosion like the CuPt NFs (Figure S10). To further confirm if the temperature was a factor to induce growth at the {100} vertices of the NFs, a control experiment was done by heating the reaction solution at 200 °C for an hour followed by 230 °C for another hour after addition of Pt(acac)2. This resulted in the spiny RD nanocrystals with further growth at {100} vertices but only a limited amount of the small particle alloy byproduct (Figure S9f). This suggests a competition in the consumption of Pt atoms to growth of spiny nanostructures and newly formed smaller alloyed nanoparticles. Following the information gained from control experiments, 200 °C was recognized as the optimum temperature for the steady growth of CuPt alloy crystal domains on the Cu nanocubes, leading to the core−frame RD nanocrystals. However, further growth of Pt atoms at the {100} vertices of the CuPt frames required prolonged heating at higher temperature. This suggests that the reduction rate of Pt precursor influences the formation of spiny RD nanocrystals. In order to increase the deposition without losing the uniformity of Cu−CuPt RD nanocrystals, the amount of Pt(acac)2 was doubled in the synthesis of Cu−CuPt core−frame RD nanocrystals. Figures 3a,b and S11 show the well-defined

pattern indicates that they are of f.c.c. crystal structure with 111, 200, 220, and 311 peaks at 18.92°, 21.90°, 31.18°, and 36.74°, respectively. All peaks are in-between those of f.c.c. Pt (17.28°, 19.98°, 28.42°, and 33.44°; 41525-ICSD) and f.c.c. Cu (19.00°, 21.98°, 31.26°, and 36.84°; 43493-ICSD). Notably, the peaks have weak shoulders at their low 2-theta sides. Looking into the shape of the (111) peak, a weak peak at 18.83° and a strong one at 18.94° can be separated by deconvolution (inset pattern). The weak peak denotes the alloy structures of the CuPt frames in an estimated Pt/Cu ratio of 13/87. The alloying CuPt crystal domain is observed in the aberration-corrected acHAADF-STEM image of a single core−frame RD nanocrystal viewed along the [100] direction (Figures 1e and S7). In this image, the CuPt frame is distinguishably thicker at the corners but thin on the edges. Figure 1f is the atom-resolved image from the margined area in Figure 1e. The alloying region shows that the {200} d-spacing is 1.88 Å (Pt 1.99 Å; Cu 1.80 Å) and that of the {220} planes is 1.31 Å (Pt 1.40 Å; Cu 1.28 Å). In contrast to the weak peak in the inset XRD pattern in Figure 1g, the strong peak is attributed to the (111) peak of the Cu cores, albeit a shift of 0.06° in the 2-theta value comparing with that of the pure Cu metal reference. The shift is possibly due to a little alteration in the unit cell constants of the Cu cores induced by the surface-growing CuPt frames and little-alloying {110} faces, in addition to inconsistent experimental parameters in diffraction between the core−frame sample and the quoted Cu standard. Structural Analysis of CuPt RD Nanoframes. Treating the Cu−CuPt core−frame RD nanocrystals with diluted nitric acid removes the Cu from the cores via the reaction Cu(s) + 4HNO3 → Cu2+ + 2NO32− + 2NO2(g) + 2H2O and results in nanoframes with a 3D-open structure. Figures 2a and S3e are the SEM images of the CuPt RD NFs. The NFs have similarly high uniformity of the Cu−CuPt core−frame nanocrystals in both size and shape. The histogram in Figure S3f shows the size distribution of the NFs, giving an average size of 29.6 nm in the edge with a standard deviation of 6.1%. The average size shrinks compared with that of the core−frame nanocrystals after acidic etching. Possibly some Cu atoms were also leached for frame restructuring and consequently resulted in size shrinking. This is possibly due to Cu leaching from the frames in addition to the cores. Figure 2b contains a BF TEM image of the NFs to show that the Cu core was removed from inside the frames. Figure 2c−e contains the elemental maps of Cu and Pt acquired by STEM-EDS analysis on a single RD NF that indicate an alloy composition of the NFs. The result was further confirmed by synchrotron X-ray diffraction. Figure 2h is the diffraction pattern showing the (111), (200), (220), and (311) peaks of the NFs at 18.3°, 21.1°, 30.0°, and 35.36°, respectively, revealing that they remained the f.c.c. crystal structure. However, all peaks broadened and shifted to lower 2-theta values after acid etching compared with those of the core− frame nanocrystals. The peak broadening is likely a result of the smaller crystal domains, and the peak shift is evidence of the removal of Cu atoms, leading to an increase in the relative content of Pt in the alloy NFs. Through measuring the average composition of the CuPt NFs by ICP-OES, a value of 26/74 was obtained for the Pt/Cu ratio, higher than that of the frames on the Cu−CuPt core−frame nanocrystals. Figure 2f (enlarged in Figure S8a) is the ac-HAADF-STEM image of a single CuPt NF. Looking into the atomic structure of this NF, abundant defects such as adatoms and vacancies are observable (yellow arrows in Figures 2g and S8b). These results, together with the 5685

DOI: 10.1021/acs.chemmater.7b01550 Chem. Mater. 2017, 29, 5681−5692

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Figure 3. (a) SEM and (b) BF TEM images of Cu−CuPt spiny RD nanocrystals and the crystal model. (c) HAADF-STEM image of a spiny RD nanocrystal with (d) atomic-resolution ac-HAADF-STEM image of a selected vertex (the boxed area in (c)). (e) BF TEM image of CuPt spiny RD NFs and the corresponding crystal model. (f) ac-HAADF-STEM image of a spiny RD NF. (g) Atom-resolved ac-HAADF-STEM image over the boxed area in (f)). (h) Synchrotron X-ray diffraction pattern of CuPt spiny RD NFs overlapping with that of RD NFs. Both scale bars in (c) and (f) indicate 10 nm.

resulting in poor growth control of CuPt nanostructures. However, the rapid reduction of Pt precursors efficiently prevented the Cu nanocubes from undergoing corrosion prematurely and led to the epitaxial growth of CuPt alloy frames with cores intact. Notably, the frame structures formed as alloys, instead of purely Pt. To elucidate this seemingly peculiar growth mechanism, UV−vis spectroscopy and ICPOES were carried out to investigate the variation of Cu and Pt ions from precursors. Figure S13a and S13b show the timedependent UV−vis spectra of Cu(I)Br and Cu(II)Br 2 precursors reduced without addition of Pt(acac)2. The solid and dashed gray curves represent the absorption of Cu+ and Cu2+ ions which peak at 655 and 620 nm, respectively. The peak of Cu+ ions has a blue shift from 655 to 630 nm after 3 h while that of Cu2+ ions remains at 620 nm over time. This difference indicates that a portion of the Cu+ ions were gradually oxidized owing to disproportionation (2Cu+ → Cu2+ + Cu0) and hence caused the blue shift. Figure S13e proves that TOPO was not a factor inducing oxidation of Cu+ ions but actually somewhat eased it. According to the ICP-OES data, a gradual drop in the Cu content of the Cu(I)Br supernatant over time indicates that a certain amount of Cu+ ions were slowly reduced by OAm and disproportionation (black columns in Figure S14a). However, there was a 6% Cu content drop in the Cu(II)Br2 supernatant because no disproportionation occurred and Cu2+ ions were unable to be reduced by OAm (Figure S14b). Comparing the results with those of the supernatants from the solutions of RD and spiny RD Cu−CuPt RD nanocrystals, the absorption peaks have a further blue shift for the addition of Pt(acac)2. As shown in Figure S14c, the absorption peak of the supernatant from the solution of RD nanocrystals is at 630 nm (here Pt(acac)2 was added after 1 h,

vertex-extended nanostructures synthesized at 200 C for 2 h after addition of Pt(acac)2. The corresponding crystal model is shown in Figure 3a. The HAADF-STEM image of a single spiny nanocrystal in Figure 3c shows brighter contrast located at the {100} vertices (pointed to by the arrow), denoting a higher Pt content. The atomic structure of one tip is shown in Figure 3d (boxed area in Figure 3c). In Figure S12, the crosssection of the area around the tip at this orientation suggests high index faceting (terraced {111} steps). Figure 3e presents a TEM image of the spiny CuPt NFs obtained by etching with the nitric acid, and Figure 3f is an ac-HAADF-STEM image of a single spiny NF viewed along the [111] direction. The atomresolved ac-HAADF-STEM image in the boxed area in the Figure 3f shows no segregated Pt nanocrystals but some crystal domains with high brightness which indicates higher Pt content (Figure 3g). The CuPt alloy structure is confirmed where measurement of d-spacings gives 1.34 and 1.91 Å for 220 and 200 crystal planes, respectively. Both are between those of pure Cu and Pt metal references. Synchrotron X-ray diffraction was also performed, by which a broadened pattern in the f.c.c. crystal structure was obtained (Figure 3h). It is worth noting that the peaks of the spiny RD CuPt NFs shift to lower 2-theta angle compared the counterpart CuPt NFs without extended vertices. The average composition of the spiny NFs was further analyzed by ICP-OES and was shown to have a Pt/Cu ratio of 41/59. These results agree with an increase of Pt content in the spiny NFs attributed to addition of Pt atoms, in particular, the region of {100} vertices. Formation Mechanism of the CuPt Alloy Nanoframes on the Cu Nanotemplates. In the coating of Pt layers on the surfaces of Cu nanocubes, it might be expected that the cores would be sacrificed through galvanic replacement37 of Cu, 5686

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Figure 4. (a−h) BF-TEM images display different intermediates over time in the syntheses of Cu−CuPt RD and spiny RD nanocrystals. For RD nanocrystals, the intermediates were collected at the (a) 5th, (b) 10th, (c) 20th, and (d) 30th minute while the spiny ones were collected at the (e) 10th, (f) 50th, (g) 70th, and (h) 90th minute. (i) Possible mechanism of surface evolution from the Cu nanocube to the Cu−CuPt spiny RD nanocrystal through the Cu−CuPt RD shape when the Pt source is sufficient.

Nevertheless, the remaining existence of the core−frame RD nanocrystals reveals two important clues. First, galvanic replacement was effectively slowed in the synthesis conditions by virtue of the protection of TOPO and the rapid reduction of Pt2+ ions by OAm. Moreover, because there were suspended Cu+ ions in the solution, galvanic replacement was further inhibited by disproportionation taking up Pt2+ ions. The estimation was further proved by carrying out two control experiments. The first experiment was to synthesize the Cu− CuPt RD nanocrystals without adding any TOPO in the whole synthetic process. It turned out that the etching on the 110 faces of Cu−CuPt core−frame nanocrystals was stronger and thus led to semihollow Cu−CuPt RD nanocrystals (Figure S16a,b). Since knowing TOPO prefers to weakly cap on Cu 100 faces but not 110 edges, we estimated for the suspended TOPO molecules the role to ease the etching of galvanic replacement on the Cu 110 faces through trapping Pt2+ ions (possibly by PO groups). Without TOPO, the excessive etching by free Pt2+ ions would take place. The second experiment was to synthesize the Cu−CuPt RD nanocrystals with washed Cu nanocubes in pure oleylamine without suspended TOPO and Cu+ ions. It was surprising that galvanic replacement was dominant during the Pt coating and made most nanocrystals become hollow structures (Figure S16c,d), which indicated that the suspended Cu+ ions were the suppressor to free Pt2+ ions in addition to TOPO. Notably,

total reaction time was 2 h), which is 5 nm shorter in wavelength to that from the Cu(I)Br solution after heating for 2 h. For the case of spiny RD nanocrystal solution (double amount of Pt(acac)2 was added after 1 h, total reaction time was 3 h), the blue shift is 8 nm. These further blue shifts are due to the electron transfer from a portion of Cu+ to Pt2+ ions toward the formation of Pt0 atoms. Hence, more Pt2+ ions caused a larger shift in absorption. In the ICP-OES results, low content of Pt2+ ions was left in the solutions of the RD and the spiny RD Cu−CuPt nanocrystals, owing to their rapid reduction by OAm (Figure S14c,d). However, no significant difference in the Cu contents of both samples (orange columns in Figure S14a) is observed if comparing with the black columns. This indicates that disproportionation and slow reduction of Cu+ ions by OAm were independent of the reduction of Pt2+ ions. On the basis of these results, the formation of CuPt alloy can be attributed to the coreduction of Cu+ and Pt2+ ions. For further validation, a control experiment was performed for washed Cu nanocubes (separated from the original supernatant and redispersed in fresh OAm with TOPO) used for the coating of Pt atoms in the same condition as in synthesizing the RD nanocrystals. As a consequence, the Cu−CuPt core−frame RD nanocrystals formed as the major product; however, there were hollow CuPt alloy spheres present (Figure S15). The hollow spheres likely come from galvanic replacement. 5687

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alloy distribution of Pt and Cu is shown in Figure S19d, where a slight enrichment of Pt at corners is evident. After the reaction reached 30 min, the growth of Cu−CuPt core−frame RD nanocrystals was nearly complete (Figures 4d and S20). At this time, a slight heterogeneity in the distribution of Pt signal is observed on the {100} vertices. In stage II, the intermediates were studied during the growth of spiny core−frame RD nanocrystals to interpret their formation mechanism. Figure 4e−h shows the intermediates captured at different reaction times of 10, 50, 70, and 90 min after addition of Pt(acac)2. At the 10th minute, intact RD nanocrystals were already observed, 20 min faster than that in the stage I. This is likely due to the doubling of [Pt(acac)2]. Since the 50th minute, epitaxial growth and/or coalescence of smaller nanoparticles at the {100} vertices (arrows in Figure 4f) was slow until the 90th minute when well-defined {100} vertices were observed (Figure 4g,h). Figure S21 records the tip-to-tip diameters of the spiny RD nanocrystals collected at different reaction times. The size evolves from 45 to 59.8 nm across 110 min. It can be concluded from these results that crystal growth along ⟨100⟩ directions dominated and thus results in the RD shape. One factor for this overall dominant growth of 100 faces arose from the codeposition of Cu0 and Pt0 atoms in the form of CuPt small nanoparticles. Although TOPO was recognized as being in the auxiliary role for shaping the Cu nanocubes with OAm owing to its preferential capping on the Cu {100} faces, the weak capping of TOPO did not strictly stop the growth of 100 crystal faces because of their dynamic adsorption− desorption activities under a high-temperature heating. For galvanic replacement, it was the factor hard excluded in spite of the suppression of Pt2+ ions by TOPO and Cu+ ions. Hence, galvanic replacement, in fact very slow and weak, was also contributing to the formation of the RD shape with a fast etching rate on the {110} edges of a Cu nanocube. Overall, the shape transformation from a Cu nanocube to a Cu−CuPt RD nanocrystal was achieved by the synergistic and balancing effects of weak galvanic replacement and codeposition of Cu0 and Pt0 atoms in the form of CuPt small nanoparticles. Figure 4i is the scheme interpreting the possible formation mechanism of Cu−CuPt core−frame RD nanocrystals across the two stages. From the beginning (left), Pt2+ and Cu+ ions are coreduced to nucleate CuPt alloy monomers and small particles which epitaxially adsorb on the Cu nanocubes and undergo reconstruction through migration on the Cu surfaces (blue arrows) and interdiffusion across the interface (light green and red arrows). The Pt0 atoms on the Cu surfaces come from either reduction by OAm and Cu+ ions (green dashed arrows) or galvanic replacement by taking electrons from Cu0 atoms on the Cu nanocube (red dashed arrows). The first intermediates are RC nanocrystals. The optimized temperature for the CuPt alloy coating should not be higher than 200 °C; otherwise, the growth of small CuPt alloy nanoparticles becomes dominant, and hence bigger CuPt nanoparticles are generated (Figure S9d). On the RC surfaces, reduction, deposition, and reconstruction ultimately led to the RD and spiny RD morphologies. The anisotropy of the shape across the two stages relies on the concentration of the Pt precursor; in particular the formation of terraced highly pointed vertices required sufficient Pt atoms. Regarding fast growth on the cube {100} crystal faces generating the rhombic dodecahedral {110} facets, it is not clear yet. With regard to the capping agent, it is possible that the distribution of TOPO on the square {100} facets is

the hollow structures were the RD cages rather than frame structures (Figure S16d). The critical difference tells that Pt0 atoms generated by galvanic replacement could deposit on all Cu crystal faces but had a relatively higher growth rate on Cu 100 faces (coupling with a faster etching rate by Pt2+ on Cu 110 faces) than others, thus leading to the shape evolution from Cu nanocubes to CuPt RD cages. Second, coreduction was not the only factor to achieve CuPt alloy crystal domains on the Cu cores. Interdiffusion, a phenomenon that generally occurs in binary metallic systems owing to the Kirkendall effect, is another driving force by which inhomogeneous crystals could restructure if enough energy is available to overcome the barrier of atomic motion.38,39 This process typically occurs in response to a concentration gradient of atoms. In this way, the Pt atoms of the deposited crystal domains and Cu atoms on the surfaces of Cu nanocubes diffused into each other. However, hollow CuPt nanostructures or the voids arisen from the Kirkendall effect were absent. It is due to the far lower quantity of Pt atoms (0.0193 mmol for RD and 0.0386 for spiny RD) than that of Cu ones (0.428 mmol) which lags the interdiffusion rate and hence limits the blending depth at 2 nm over the Cu surfaces. Apart from the driving forces for Cu−Pt alloying, the process forming RD and spiny RD core−frame nanostructures was also investigated. For this purpose, two stages of control experiments were carried out to realize the formation of RD (stage I) and spiny RD nanocrystals (stage II). Figure 4a−d shows the stage I growth process in which four intermediates were captured at different reaction times of 5, 10, 20, and 30 min after adding the corresponding amount of Pt(acac)2. First, polyhedra are obtained by the fifth minute (Figure 4a). The polyhedral nanocrystals were confirmed the rhombicubocthedral (RC) shape which are single-crystalline structures bounded by {100}, {110}, and {111} faces (Figure S17a−c). Because the cubic substrates lack {110} and {111} facets, the RC structure is reasonably arrived at as the result of relatively faster growth on the square 100 faces with slower growth on {111} corners and {110} edges of the Cu nanocubes. When the reaction time reached 10 min, the resulting products were quasi-RD nanocrystals showing anisotropic growth at the vertices and a number of smaller nanoparticles (Figure 4b). The irregular quasi-RD shape formed because of deposition at vertices and possibly through the coalescence of smaller nanoparticles (Figure S18a). The structures were determined to be singlecrystalline by selected-area electron diffraction (SAED) in spite of their somewhat irregular appearance (Figure S18b). Judging from this and the local FFT pattern, the region around the vertex showed an atomic arrangement consistent with that of the central core (Figure S18c). This suggests the growth is occurring epitaxially from the substrate, in addition to some alignment/rearrangement of crystal domains of small particles before and after their coalescence. Importantly, because there are a relatively large number of smaller nanoparticles which are observed to be consumed later in the reaction and because these and the resulting outgrowth have alloyed composition (Figure S18d), this might suggest the latter mechanism (particle attachment) to be dominant; however, since the structure already presented an RC structure, we cannot rule out the possibility of ripening at the expense of these smaller particles in continued epitaxial growth. When the reaction time reached 20 min, nanocrystals took on an RD shape. However, they expose rough surfaces, which are attributed to the incomplete surface reconstruction (Figures 4c and S19). The 5688

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Figure 5. (a) CV and (b) ORR polarization curves for CuPt RD and spiny RD NFs and commercial Pt/C catalysts, respectively. (c) Compared results of specific activities and (d) Pt mass activities for the three kinds of catalysts at 0.9 V vs RHE. (e) ln(A0/At)−time plot and (f) Pt mass specific conversions of 4-NP reduction with the three kinds of catalysts in the dark and under irradiation of visible light.

Catalytic Performance of CuPt RD and Spiny RD Nanoframes. Pt is the primary component in the majority of fuel cell catalysts; however, the metal’s rarity hinders the possibility of mass production of fuel cell vehicles. The as-made RD and spiny Rd CuPt NFs both formed 3D-open frameworks with Cu-rich compositions (>50%) and thus are advantageous to scaling up because of their high surface area and potential reduction in cost through Pt thrifting. Compared with the NFs, the Cu−CuPt core−frame RD nanocrystals were not desired catalysts for ORR because the oxidation potential of pure Cu (Cu0 → Cu2+ + 2e−, E0 = 0.342 V versus RHE) results in the dissolution of the cores during catalytic tests.24 Figure S22 shows the dissolution of the Cu cores during the CV scans in 0.1 M HClO4. In addition, the Cu−CuPt core−frame RD nanocrystals are not stable for long storage. In the stability test, both the regular and the spiny core−frame RD nanocrystals

disturbed after adding Pt(acac)2 , where TOPO might preferentially occupy and have easier access to exposed higher energy corners and edges of the cubes, initially promoting growth in the [100] direction, yet, this does not account for the further stabilization of {110} facets. The shape appears to be somewhat accessible for first row transition metal alloys with Pt, despite the {110} facets having the highest surface energy of the low index set for f.c.c. metals.10 The alloy composition may promote these facets due to Pt surface enrichment, where a Pt skin may help to stabilize this surface relative to others when the transition metal is present underneath. Lastly, we note that though the (110) surface may be thermodynamically the least stable, the surface area to volume ratio of the RD compared to the alternative low index faceted cubes and octahedra is the smallest for the same volume, which may also contribute to the presence of the morphology. 5689

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Chemistry of Materials were collected following the typical procedure and finally redispersed in 0.5 mL of toluene at ambient. After 1 week, most of the Cu cores were oxidized, whereas the frame structures remained (Figure S23a,b). If they were stored in water or ethanol, the process took less than 3 days. According to the XRD results, CuO had become the major composite arisen from the oxidation of the nonalloying Cu crystal domains (Figure S23c,d). Hence, all Cu−CuPt nanocrystals had their Cu cores removed right after as-prepared to form more stable CuPt NFs for further use. The CuPt NFs were also stored no longer than 3 days. To examine the potentials of these two NFs for fuel cell electrochemistry, the oxygen reduction reaction (ORR) was carried out on the catalysts on a rotating disc electrode (RDE) and compared with the performance of commercial Pt/C catalysts. Typically, the NFs were prepared as described in the experimental section and loaded on a glassy carbon RDE (GCRDE, geometric area: 7.068 mm2) with the Pt mass loading of 29.20 μg/cm2 for the regular NFs, 15.42 μg/cm2 for the V−S NFs, and 20.37 μg/cm2 for the commercial Pt/C catalysts. The cyclic voltammetry (CV) was performed in an Ar-saturated 0.1 M HClO4 aqueous solution at a sweep rate of 50 mV/s to measure the specific electrochemical active surface areas (ECSAHupd) of all catalysts (Figure 5a). The ECSAHupd was derived from the proton adsorption/desorption peak areas (0.02 to 0.35 V versus RHE) which are 53.87 m2/gPt for the RD NFs, 39.09 m2/gPt for the spiny RD NFs, and 16.99 m2/gPt for the commercial Pt/C. Figure 5b shows the ORR polarization curves normalized by the GC-RDE geometric area. The halfwave potentials were 0.87 V for the RD NFs, 0.86 V for the spiny RD NFs, and 0.78 V for the commercial Pt/C, showing the enhanced ORR activities of the NFs in both types. The specific and mass activities were normalized by the ECSAHupd and the total mass of the loaded Pt (Figure 5c,d). The spiny RD NFs gave a specific activity of 2.19 mA/cm2 at 0.9 V versus RHE, higher than those of regular NFs (1.70 mA/cm2) and Pt/ C (0.74 mA/cm2) tested under the same conditions. This indicates that the abundant terraces on the surfaces of the extended vertices benefit oxygen reduction. However, the specific activity of the spiny RD NFs decayed to 1.48 mA/cm2 (32.4% decrease) in the accelerated deterioration test (ADT) under a sweep rate of 100 mV/s in an O2 saturated 0.1 M HClO4 after 1200 cycles. Here the RD NFs (24.7%) retained activity longer, suggesting a more stable structure. The degradation of spiny RD NFs was characterized by examining the catalyst structures after ADT. Figure S24a show that a portion of the RD NFs after ADT aggregated and had more segregated Pt crystal domains in the skeletons (dashed circles). However, the RD frame structures were still had crystallinity (Figure S24b). In contrast, the spiny RD NFs after ADT had distortion or collapse in the skeletons and their vertices became rounded in their contours, apart from aggregation (Figure S24c,d). Nevertheless, both the NFs had better performance than that of the commercial Pt/C (48.3% decay in ECSA activity). Comparing the mass activities, the spiny RD NFs gave a lower value (0.86 A/mgPt) than that of the regular NFs (1.08 A/mgPt). This makes sense considering that the spiny RD NFs have higher Pt content in the structures. After ADT, the mass activity of the spiny RD NFs dropped to 0.58 A/mgPt and that of the RD NFs to 0.81 A/mgPt. Importantly, the two kinds of NFs showed at least 8 times higher mass activities than that of the commercial Pt/C in the initial sweep (0.13 A/mgPt) and after ADT (0.06 A/mgPt). Figure S25a−c shows the ORR

polarization curves of the three samples acquired at different rotating speeds of the GC-RDE. The corresponding Koutecky− Levich plots reveal that diffusion current linearly scales with ω1/2 and demonstrates a first-order ORR kinetics at the potentials of 0.8 V (Figure S25d). The electron-transfer number (n) was quantitatively calculated to be 3.9 for the RD NFs, 4.0 for the spiny RD NFs, and 3.6 for the commercial Pt/C catalysts by using the slopes of these plots, indicating the defected surfaces and the terraced vertexes of NFs benefit the 4-electron pathway in the ORR process. Since the NFs are Cu-rich structures, the dense Cu nanocrystal domains in their skeletons generate the localized surface plasmon resonance (LSPR) after absorbing visible light, which possibly enhances the reactivity on the surfaces of the NFs and therefore the catalytic performance. Figure S26a shows the UV−vis absorption spectra of Cu nanocubes which indicates an LSPR absorption peak at 646 nm. Compared with those of Cu−CuPt RD and spiny RD nanocrystals (dash curves in Figure S26b and S26c), the LSPR absorption peak of Cu shifted to 687 and 694 nm owing to nanocrystals that changed in size and morphology. After removing Cu cores with nitric acid, the strong LSPR peak of Cu was shifting to 550 nm for the CuPt RD NFs and 595 nm for the CuPt spiny RD NFs. Both peaks became broad and weak owing to the limited crystal volume and inhomogeneous distribution of Cu nanophases in the NFs. In addition, the difference in absorption wavelength between the NFs reveals a significant influence of the extended CuPt vertices to the dielectric surroundings. To confirm the phenomenon of the LSPR-enhancement, the reduction of 4nitrophenol was carried out at a constant temperature of 32 °C with the NFs and the commercial Pt nanoparticles (ca. 200 nm in size) as the catalysts. The 4-nitrophenol (4-NP) is a strong visible light absorber owing to an absorption peak at 400 nm, yet the reaction product 4-aminophenol (4-AP) has absorption at 305 nm. In the presence of excess NaBH4, the reaction is known to correspond with pseudo-first-order reaction kinetics and is dependent upon the concentration of 4-NP. Hence, the slope in the plot of ln(A0/At) versus reaction time, where A is the absorbance at 400 nm, represents the reaction rate k. Figure S27 collects the time-dependent UV−vis spectra of 4-NP degradation at 400 nm with no catalyst, two types of NFs, and commercial Pt NPs. The results were obtained in the dark and under irradiation of visible light generated by a Xe lamp providing a true incident light power of 120 mW to the sample at an irradiated distance of 25 cm. Figure S28a is the calibration line of UV−vis absorption at 400 nm versus the concentration of 4-NP and Figure S28b is the ln(A0/At) versus time plot of 4NP reduction with no catalyst, showing no significant enhancement on the reaction rate after applying visible-light irradiation. In Figure 5e, the results show that the reaction rates of 4-NP reduction correspond with the first order for all catalysts after the first 3 min. Nevertheless, their rates are different, which are 0.002 min−1 for the commercial Pt NPs, 0.175 min−1 for the RD CuPt NFs, and 0.197 min−1 for the spiny RD CuPt NFs. It is worth noting that the activity of the Pt NPs was not improved under the visible-light irradiation but elevated to 0.266 min−1 for the RD NFs and 0.261 min−1 for the spiny RD CuPt NFs. The improvement came from the effect of LSPR enhancement with the factors of 1.52 and 1.32, respectively. The two types of NFs seem to have close mass activities under light irradiation (the same weight of 1 μg). However, the curves of the specific conversion of 4-NP over time and Pt mass reveals that the RD NFs were the most cost5690

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efficient catalysts (Figure 5f). This again is a result of a comparatively high content of Pt in the spiny RD NFs not exposed on the surfaces for reaction.

CONCLUSIONS In conclusion, we report the synthesis of CuPt rhombic dodecahedral alloy nanoframes via a Cu nanocube templated method. We have further investigated the overgrowth mechanism on the Cu nanocubes, where introduction of Pt ions to the Cu nanocube solution triggered growth in the [100] direction, which led to the RD and spiny RD structures with extended {100} vertices. Consequently, the two kinds of open RD CuPt NFs with abundant surface defects showed improved catalytic performance in both ORR and 4-NP reduction reaction. The Cu-rich composition gave plasmon-enhancement to their catalytic activities in 4-NP reduction by excitation of the LSPR from visible light, at least 1.3 times higher than that in dark. Future work is aimed at extending this strategy to alternate cost-effective metal combinations for catalysis. ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.7b01550. Additional SEM and TEM images, SAED patterns, HAADF-STEM-EDS mapping images, UV−vis spectra, ICP-OES results, and catalytic performance of catalysts (PDF)

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AUTHOR INFORMATION

Corresponding Author

*(C-H.K.) E-mail: [email protected]. Phone: +8862-27898615. Fax: +886-2-27898611. ORCID

Brian T. Sneed: 0000-0002-5656-6180 Chun-Hong Kuo: 0000-0001-6633-8985 Author Contributions #

L.-M. L. and Y.-C. K. contributed equally to this work.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We are grateful for the technical support from NanoCore, the Core Facilities for Nanoscience and Nanotechnology at Academia Sinica in Taiwan. A portion of the electron microscopy was performed as part of a user project through Oak Ridge National Laboratory’s Center for Nanophase Materials Sciences, which is a U.S. Department of Energy (DOE) Office of Science User Facility, and using instrumentation provided by the U.S. DOE Office of Nuclear Energy, Fuel Cycle R&D Program, and the Nuclear Science User Facilities. We specially thank Ms. Mei-Ying Chung, the technician in the Institute of Chemistry at Academia Sinica in Taiwan, for carrying out SEM analyses and measurements of ICP-OES. This work is financially supported by the Ministry of Science and Technology, Taiwan (MOST 104-2113-M-001-007-MY2), Academia Sinica, Taiwan (Program of Nanotechnology 2393), and Executive Yuan, Taiwan (Taiwan’s Deep Decarbonization Pathways toward a Sustainable Society). 5691

DOI: 10.1021/acs.chemmater.7b01550 Chem. Mater. 2017, 29, 5681−5692

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DOI: 10.1021/acs.chemmater.7b01550 Chem. Mater. 2017, 29, 5681−5692