Article pubs.acs.org/JACS
Cite This: J. Am. Chem. Soc. 2018, 140, 11680−11685
In Situ Kinetic and Thermodynamic Growth Control of Au−Pd Core− Shell Nanoparticles Shu Fen Tan,†,‡ Geeta Bisht,†,‡ Utkarsh Anand,†,‡ Michel Bosman,§,∥ Xin Ee Yong,‡ and Utkur Mirsaidov*,†,‡,§,⊥,# †
Department of Physics, National University of Singapore, Singapore 117551 Centre for BioImaging Sciences, Department of Biological Sciences, National University of Singapore, Singapore 117557 § Department of Materials Science and Engineering, National University of Singapore, Singapore 117575 ∥ Institute of Materials Research and Engineering, Agency for Science, Technology, and Research (A*STAR), Singapore 138634 ⊥ Centre for Advanced 2D Materials and Graphene Research Centre, National University of Singapore, Singapore 117546 # NUSNNI-NanoCore, National University of Singapore, Singapore 117411
J. Am. Chem. Soc. 2018.140:11680-11685. Downloaded from pubs.acs.org by UNIV OF NEW ENGLAND on 09/25/18. For personal use only.
‡
S Supporting Information *
ABSTRACT: One-pot wet-chemical synthesis is a simple way to obtain nanoparticles (NPs) with a well-defined shape and composition. However, achieving good control over NP synthesis would require a comprehensive understanding of the mechanisms of NP formation, something that is challenging to obtain experimentally. Here, we study the formation of gold (Au) core−palladium (Pd) shell NPs under kinetically and thermodynamically controlled reaction conditions using in situ liquid cell transmission electron microscopy (TEM). By controlling the reaction temperature, we demonstrate that it is possible to tune the shape of Au nanorods to Au−Pd arrowheaded structures or to cuboidal core−shell NPs. Our in situ studies show that the reaction temperature can switch the Pd shell growth between the kinetically and thermodynamically dominant regimes. The mechanistic insights reported here reveal how the reaction temperature affects the packing of the capping agents and how the facet selection of depositing shell atoms drives the shell formation under different kinetic conditions, which is useful for synthesizing NPs with greater design flexibility in shape and elemental composition for various technological applications.
■
INTRODUCTION Bimetallic nanoparticles (NPs) consisting of two different types of metals are important for many technological applications1 because of their enhanced catalytic performance,2 tunable optical properties,3 and improved thermal stability4 in contrast to their monometallic counterparts. From simple bimetallic systems, different-shaped structures of alloyed or intermetallic NPs, such as dimeric, dendritic, and core−shell bimetallic NPs, can be derived.1,5,6 Common chemical routes for synthesizing well-defined bimetallic NPs include co-reduction,7 thermal decomposition8 of two metal precursors at the same time, seed-mediated method,9 and galvanic replacement10 where the last two approaches utilize monometallic NPs as templates in a solution. Bimetallic NPs can be formed from the monometallic NP templates either by growing a different metal on them or by etching and replacing the template NPs by another metallic component. Despite tremendous recent advances in the synthesis of bimetallic NPs with different shapes, sizes, and compositions,1,11,12 the detailed mechanisms through which the NPs form in solution is still not clear. To understand the pathways of the bimetallic NP formation, a direct real-time © 2018 American Chemical Society
nanoscale probing of the NP evolution is needed. Current characterization methods are based on “quench-and-look” approach where the reaction is stopped at different stages and imaged with transmission electron microscopy (TEM),13 or indirect spectroscopic techniques compatible with in situ measurements, such as UV−vis14 and small-angle X-ray scattering,15 are used. These approaches have provided detailed insights into the transformation of NPs in a precursor solution. However, these methods lack either temporal or spatial resolution. Therefore, many important questions regarding the mechanisms of bimetallic NP formation are still not answered. For example, what is the role of the reaction temperature in controlling the nucleation and growth of the shell on the core NPs in a one-pot solution? What are the key factors that determine the facet selection for the deposition of shell atoms? In addition, we lack atomistic insights into the mechanisms of the shell atom deposition at core−shell interfaces.16 Received: May 23, 2018 Published: August 12, 2018 11680
DOI: 10.1021/jacs.8b05217 J. Am. Chem. Soc. 2018, 140, 11680−11685
Article
Journal of the American Chemical Society
Figure 1. Two different Au−Pd nanostructures synthesized from Au NRs. (A) Schematic showing the formation of two different Au−Pd nanostructures: arrow-headed NR (top) and cuboidal nanobar (bottom) that grow from Au NRs at T = 23 °C and T = 80 °C, respectively. (B) TEM images of Au NRs used as seeds. TEM images of core−shell Au−Pd (C) arrow-headed NRs and (D) nanobars synthesized at T = 23 °C and T = 80 °C, respectively. HAADF-STEM images and corresponding EDX chemical maps of a Au−Pd (E) arrow-headed NR and (F) nanobar, respectively. Nanoseedz Ltd., Hong Kong); we measured the NRs to be 35−50 nm in length and 10−15 nm in diameter. Ascorbic acid (cat. no. A596025G, Sigma-Aldrich Co., St. Louis, MO, USA), cetyltrimethylammonium bromide (CTAB) (cat. no. 52370-500G, Sigma-Aldrich Co., St Louis, MO, USA), and sodium tetrachloropalladate (II) (cat. no. 205818-1G, Sigma-Aldrich Co., St Louis, MO, USA) were used as received without further purification. We transferred 100 μL of the Au NR solution from the stock solution into a 1.5 mL centrifuge tube. Next, 100 μL of 100 mM aqueous CTAB solution and 60 μL of 50 mM aqueous ascorbic acid solution were added into the Au NR suspension. All aqueous solutions were prepared using deionized water with a resistivity of 18.2 MΩ·cm. Experimental Procedures. For in situ liquid cell TEM imaging experiments, ∼500 nL of the mixture was drop-casted onto the bottom membrane of our custom microfabricated liquid cell.26 Each liquid cell consisted of two chips with ultrathin (∼20 nm) electron translucent SiNx windows separated by ∼200 nm thick spacer. The dimensions of the SiNx windows of the liquid cells are 30 μm × 200 μm. For heating experiments, we replaced one of the chips with a chip that had a microfabricated heating element directly on top of the 50 nm-thick SiNx window (Hummingbird Scientific, Lacey, WA, USA). In these chips, a molybdenum (Mo) thin-film heating element was embedded within the SiNx that allows us to tune the reaction temperature during in situ TEM experiments. Before assembling the liquid cell and drop-casting the mixed solution of NRs, CTAB, and ascorbic acid on the bottom membrane, we cleaned the surfaces of both SiNx membrane windows with the oxygen plasma to render them hydrophilic. Each liquid cell was assembled and loaded into a liquid flow holder (Hummingbird Scientific, Lacey, WA, USA). After checking and confirming that there are no leaks in the flow cell, the holder was inserted into a JEOL 2010FEG TEM (JEOL Ltd. Akishima, Tokyo, Japan) operated at 200 kV for in situ imaging where the incident electron flux ranged from 10 to 20 e/(Å2·s). The image series were acquired at a rate of 10 frames per second using a OneView CMOS camera (Gatan, Inc., Pleasanton, CA, USA). The heating chip is calibrated using the software and parameters provided by Hummingbird Scientific. The accuracy for the temperature during the in situ heating TEM experiments was about ±10 °C. The temperature was calculated based on the resistivity of the heating element. We introduced the Pd precursor solution (10 mM Na2PdCl4) into the liquid cell via a flow tube (diameter of 200 μm and length of 50 cm) connected to a syringe pump at a flow rate of 10 μL/min. It takes a few minutes for the precursor solution to reach the window area of
Owing to recent developments in in situ liquid cell TEM imaging techniques,17−20 now it is possible to follow the entire synthesis process of individual NPs from a precursor solution. This direct real-time imaging of NP evolution is a powerful approach in revealing some of the underlying mechanism of the core−shell NP growth.21,22 For example, these studies showed that the shell growth is sensitive to the chemical environment23 and interfacial lattice strain between the core and shell.24,25 In these studies, the electron beam used for imaging also served as a reducing agent for metallic ions instead of commonly used chemical reduction methods. Xia et al.16 and our own earlier studies23 showed that the final morphology of core−shell NPs depends on the deposition and surface diffusion rates of shell atoms. The competition between the surface diffusion and the deposition rate, which is sensitive to temperature, should determine whether the growth is under thermodynamic or kinetic control. In order to resolve the detailed mechanisms of different growth modes, it is essential to perform in situ TEM experiments at different temperatures under chemically relevant conditions. Here, using in situ liquid cell dynamic TEM imaging at two different temperatures, we show that the growth of Au−Pd core−shell NPs switches from kinetically to thermodynamically controlled growth when the temperature rises, and we describe the dynamics of these growth modes. Distinct from the earlier in situ TEM studies of bimetallic NPs that were conducted at room temperature23 and with the electron beam as a reducing agent,22,24,25 our approach enables us to probe the overgrowth reactions that closely mimic the conditions used in typical large-scale synthesis (elevated temperature, titration, etc.). We chose Au nanorods (NRs) as seed NPs because they possess well-defined facets, which makes them ideal candidates for studying the facet selection that occurs during different stages of the overgrowth process; Au NR tips mainly consist of {111} facets, while the side facets are {110} and {100}.
■
EXPERIMENTAL DETAILS
Sample Preparation. We used a solution of cetyltrimethylammonium bromide (CTAB) stabilized Au NRs (cat. no. NR-10-750-50, 11681
DOI: 10.1021/jacs.8b05217 J. Am. Chem. Soc. 2018, 140, 11680−11685
Article
Journal of the American Chemical Society the liquid cell. The videos were recorded when we first observed visible changes to the Au NRs (t = 0 s). The high-angle annular dark-field scanning TEM (HAADF-STEM) imaging and chemical analysis were performed with an FEI Titan TEM and a JEOL ARM operated at 200 kV. EDX mapping was performed with a ∼0.5 nm diameter probe and with 300 ms acquisition time per pixel.
■
RESULTS AND DISCUSSION The schematic in Figure 1A shows the formation of Au−Pd nanostructures when Au NRs (Figure 1B) react with the Pd precursor solution at two reaction temperatures. At room temperature (T = 23 °C), the Pd is deposited at both tips of the Au NRs (Figure 1C) forming arrow-headed Au−Pd NRs. The atomic percentages of Au and Pd for the NR shown in Figure 1E are 79% and 21%, respectively. At an elevated reaction temperature (T = 80 °C), Pd is deposited as cuboidal shells encapsulating the Au NRs, resulting in Au−Pd nanobar structures (Figure 1D). The atomic percentages of Au and Pd for the nanobar shown in Figure 1F are 5% and 95% (also see SI Section 1 for more EDX maps). The scanning TEM (STEM) image of the Au−Pd nanobar reveals that there are irregular dark voids in the Pd shell surrounding the Au core (Figure 1F). These voids are not Moiré patterns arising from the superposition of the Au (aAu = 4.08 Å) and Pd (aPd = 3.89 Å) crystal lattices, as this would lead to alternating dark and bright bands,27,28 which we do not observe in our images. Instead, the voids formed at the Au−Pd interface are most likely due to the strain generated by the lattice mismatch at the interface.29 Depending on the shell thickness, this interfacial strain can be released through unit cell rearrangements,30 defect generation,31 diffusion of metals across the strained interfaces,32 and formation of interfacial voids33 similar to the STEM images shown in Figures 1F and S1C−D. Nonetheless, the detailed strain−release mechanism in bimetallic core−shell NPs still needs further investigation. To understand the formation of the Au−Pd core−shell structures, we investigated their growth mechanisms at different reaction temperatures using in situ liquid cell TEM. Figure 2B shows a typical structure that forms at T = 23 °C when the Pd precursor solution is introduced into a flow cell containing the drop-casted Au NRs on the bottom membrane (Supporting Video S1). The image recorded at t = 40.0 s and the subsequent images (Figure 2B) show that the Au NR is being transformed into a Au−Pd core−shell arrow-headed NR. During this process, a few small spherical Pd NPs also formed in the solution (Figure 2B, t = 0.0 s). Some of these Pd NPs contributed to the growth of the arrow heads by initially attaching to the tip of the NR and later to the NR corners between the {100} and {110} facets (Figure 2B, t = 40.0 s), and others simply pass by the NR without coming into a contact (Supporting Video S1: t = 0.0−27.0 s). Since the amount of the visible NPs attaching to the NR accounts only for a small fraction of the growth (∼3%), we concluded that the growth by monomer (Pd atoms) addition was the dominant growth mechanism. The schematic in Figure 2A shows our proposed model describing the formation of arrow-headed NRs. First, the Pd2+ ions are reduced to Pd0 atoms by the ascorbic acid (chemical reducing agent). Then, Pd atoms and small Pd NPs, which nucleate in the solution, diffuse and deposit preferentially onto the Au NR tips. Note that the reduction of Pd2+ ion on the Au NR surface is unlikely because of the electrostatic repulsion
Figure 2. Room-temperature growth of bimetallic arrow-headed NRs. (A) Schematic showing the overgrowth process of Pd on a Au NR. (B) Time series of in situ TEM images showing the formation of arrow-headed Au−Pd core−shell NR at T = 23 °C. The small Pd NPs that form in the solution and attach to the NR are highlighted by red dashed circles (Supporting Video S1). (C) Time-domain contour plots showing the formation of the arrow-headed Au−Pd NR shown in (A). (D) The length (L) and width (W) of the same arrow-headed NR as a function of time.
between the positively charged CTAB-capped Au NR (ζ potential of ∼40 mV)34 and positively charged Pd2+ ions. The preferential attachment of Pd to the NR tips is due to low capping density of CTAB surfactant at the tips associated with the weak CTAB affinity to {111} facets at the NR tips in comparison to {110} and {100} side facets.35 Furthermore, contour plots (Figure 2C) depicting the evolution of the NR shape reveal that the growth rate along the NR length is faster than along the width (Figure 2D). A similar growth mechanism is also observed in Supporting Video S2 under the same experimental conditions. The average density of in situ synthesized Au−Pd arrow-headed NRs in the field of view is ∼9 NRs per μm2 (SI Section 8). Moreover, the final morphology of our in situ synthesized arrow-headed NRs is consistent with the ex situ benchtop synthesis (Figure 1C) and with the results of Tsuji et al., who showed that this arrowheaded NRs forms only at low Pd/Au molar ratios.36 The surface diffusion of adatoms is commonly expressed through a relation:37 D = D0e−Ed / kBT
(1)
where D is the diffusion coefficient, D0 ≈ 1.2 cm /s, is the diffusion pre-exponential factor,38 Ed ≈ 1.32 eV, is the potential energy barrier to diffusion of Pd on Au,38 kB = 1.38 × 10−23 J/K is the Boltzmann constant, and T is the absolute 2
11682
DOI: 10.1021/jacs.8b05217 J. Am. Chem. Soc. 2018, 140, 11680−11685
Article
Journal of the American Chemical Society
ripening.17 As the small Pd NPs absorb onto the Au NR, we observe a thin shell formation (Figure 3B; t = 3.8 s) as the NR continues to grow into a cuboidal NP at t = 20.7 s. Here, the amount of the attaching Pd NPs alone cannot account for the observed growth. The dissolution and adsorption of small NPs suggest that the Au−Pd core−shell nanobar can form through multiple pathways, that is, Ostwald ripening, particle coalescence, and monomer addition occurring simultaneously. Schematics in Figure 3A illustrate the overgrowth process at T = 80 °C. The TEM image series (Figure 3B) taken at T = 80 °C and corresponding contour plots (Figure 3C) show that the shell grows predominantly at the sides of the NR and slower at the NR tips. The initial fast growth of Pd shell slowed down as the Au−Pd core−shell nanobar formed at t = 20.7 s (Figure 3B). Additional in situ TEM video and images of Au−Pd nanobar formation are presented in Supporting Video S4 and Figure S7. The average density of in situ synthesized Au−Pd nanobars is ∼15 NPs per μm2 (SI Section 8). In addition to the much faster adatom diffusion at T = 80 °C, it is also expected that the CTAB bilayers capping the Au NRs are likely to be desorbed. Lower CTAB packing density along the side facets of the Au NRs at T = 80 °C39,40 makes the Pd deposition more homogeneous than at T = 23 °C. Hence, we propose the following mechanism for the formation of Au−Pd nanobars: (i) At T = 80 °C, the Au NR is loosely capped with CTAB, facilitating the epitaxial growth of Pd in the ⟨100⟩ direction; (ii) Pd atoms (Pd2+ chemically reduced by ascorbic acid) diffuse in the solution and absorb onto the Au NR, preferentially to the Au NR tips and may be to a lesser extent to the side facets;39,40 (iii) subsequently, the adatoms undergo fast surface diffusion from the tip to {110} and {100} side facets and (iv) form a cuboidal Pd shell. Note that at T = 23 °C the surface diffusion rates are low and the CTAB packing on the NR side facets is dense; therefore, the Pd is confined to the NR tips. The diffraction pattern shown in Figure S4 and the Fast Fourier transform (FFT) shown in Figure S5 suggest that these Au−Pd nanobars are single-crystalline and the cuboidal shells (Figure S8) terminated by {100} faces are epitaxially grown on the Au NRs. The {111} Pd facets grow fastest because the growth of {100} and {110} side facets is hindered by the binding of CTAB molecules (CTAB has strong binding affinity {100} and {110} side facets).35 Moreover, the disappearance of {110} facets in the final morphology is associated with the substantially higher surface energy of {110} facets compared to {100} facets.41 Thus, there is a strong driving force for the nanobar formation as a result of overall surface energy minimization and CTAB facet blocking. We performed control experiments in which we flowed the Pd precursor into the Au NR suspension in the absence of the reducing agent, ascorbic acid. In these experiments, Au NRs did not evolve into arrow-headed or cuboidal Au−Pd core− shell nanostructures (Figure S6). In addition, for low electron fluxes used in our experiments, we did not observe detectable differences in the final morphology of Pd shells (Figure S7), and the end-products of these in situ experiments were consistent with the ex situ results (Figure 1D and Figure S3L). These results suggest that for our in situ observations, the electron beam does not significantly affect the growth mechanism. It also shows that ascorbic acid not only reduces the Pd2+ ions but also assists in guiding the Pd overgrowth
temperature. From eq 1, the adatom diffusion coefficient is 4 orders of magnitude bigger for T = 80 °C than for T = 23 °C (SI Section 5). To investigate whether this drastic increase in the diffusion rate of Pd adatoms plays a role in determining the final morphology of core−shell nanostructures, adjusting the reaction temperature is a straightforward approach to manipulate the diffusion rate and its effect on the overall bimetallic NP morphology.16 To test the effect of the reaction temperature and subsequent change in the diffusion of the deposited Pd adatoms, we performed in situ TEM experiments at T = 80 ± 10 °C using liquid cells with heating capability. Again, we observed the formation of small Pd NPs surrounding the Au NR (Figure 3B; t = 0.0 s, Supporting Video S3). Some of these NPs attach and adsorb onto the Au NR surface,23 while a few of the other NPs simply pass by or dissolve back into the solution presumably contributing to the growth via Ostwald
Figure 3. High-temperature growth of bimetallic nanobars. (A) Schematic illustration showing the overgrowth process of Pd on a Au NR at T = 80 °C. (B) Time series of in situ TEM images of the Au− Pd core−shell nanobar formation at T = 80 °C. The Au NR and small Pd NPs that attach to the NR are highlighted by green dashed rounded rectangle and red dashed circles, respectively (Also see Supporting Video S3 and Figure S7). (C) Time-domain contour plots of the Au−Pd nanobar perimeter at T = 80 °C. (D) The length (L) and width (W) of the Au−Pd nanobar as a function of time. 11683
DOI: 10.1021/jacs.8b05217 J. Am. Chem. Soc. 2018, 140, 11680−11685
Article
Journal of the American Chemical Society
from pre-existing seeds. Future work could address the possibility of examining the specific roles played by thermodynamic and kinetic parameters in a one-pot synthesis. More generally, insights into the formation and time evolution of NPs are valuable for the development of efficient processes to synthesize nanomaterials with desired properties.
process (SI Section 6), which is consistent with our previous findings.42 The final morphology of core−shell NPs strongly depends on the ratio between the deposition and surface diffusion rates.16 For example, we have previously shown that a seed can develop into a cube shape in an overgrowth reaction under thermodynamic control where the rate of surface diffusion is greater than the atomic deposition rate.23 If the atomic deposition rate is faster than the surface diffusion rate, the reaction is considered to be under kinetic control. Then, this shell growth mode will result in pointy or branched structures due to the preferential atomic attachment on certain facets. In the case of our Au−Pd core−shell nanostructure, the newly reduced atoms in a solution (Figure 4) deposit at the NR tips and either mostly remain there or migrate to the side facets of the NR via surface diffusion.
■
ASSOCIATED CONTENT
S Supporting Information *
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/jacs.8b05217. Materials characterization, control experiments, and additional discussions (PDF) Video S1: Formation of Au−Pd arrow-headed NRs when the Pd precursor solution was introduced to Au NRs in the presence of ascorbic acid at T = 23 °C (AVI) Video S2: Formation of Au−Pd arrow-headed NRs when the Pd precursor solution was introduced to Au NRs in the presence of ascorbic acid at T = 23 °C (AVI) Video S3: Formation of Au−Pd nanobars when the Pd precursor solution was introduced to Au NRs in the presence of ascorbic acid at T = 80 °C (AVI) Video S4: Formation of Au−Pd nanobars when the Pd precursor solution was introduced to Au NRs in the presence of ascorbic acid at T = 80 °C (AVI)
■
AUTHOR INFORMATION
Corresponding Author
*
[email protected] ORCID Figure 4. Schematic showing the shape evolution of a NR seed under kinetic (left) and thermodynamic (right) control. This model is adapted from Xia et al.16
Utkur Mirsaidov: 0000-0001-8673-466X Notes
The authors declare no competing financial interest.
■
At room temperature, the surface diffusion rate is slow, and the growth mode is accelerated along the ⟨111⟩ and ⟨110⟩ directions, promoting the formation of the arrow-headed NRs (Figures 1C,E and 2) as a kinetically favored product. At higher temperatures, the surface diffusion rate is fast, and most of the Pd adatoms diffuse from the NR tips toward the side facets, leading to the formation of cuboidal shell. This is reflected in the faster growth rate along the NR sides than tips (Figure 3C). This formation of the cuboidal shell is also favored by the thermodynamics (the shell morphology that has a global minimum in the Gibbs free energy of the reaction) in the presence of sufficient capping agents (see Figure S2).
ACKNOWLEDGMENTS This work was supported by the Singapore Ministry of Education Academic Research Fund Tier 2 (MOE2016-T2-2009) and the Singapore National Research Foundation’s Competitive Research Program funding (NRF-CRP16-201505). We thank Prof. S. J. Pennycook for allowing us to use the aberration-corrected STEM facility at NUS-MSE.
■
REFERENCES
(1) Gilroy, K. D.; Ruditskiy, A.; Peng, H.-C.; Qin, D.; Xia, Y. Chem. Rev. 2016, 116, 10414−10472. (2) Chen, C.; Kang, Y.; Huo, Z.; Zhu, Z.; Huang, W.; Xin, H. L.; Snyder, J. D.; Li, D.; Herron, J. A.; Mavrikakis, M.; Chi, M.; More, K. L.; Li, Y.; Markovic, N. M.; Somorjai, G. A.; Yang, P.; Stamenkovic, V. R. Science 2014, 343, 1339−1343. (3) Mayer, M.; Scarabelli, L.; March, K.; Altantzis, T.; Tebbe, M.; Kociak, M.; Bals, S.; García de Abajo, F. J.; Fery, A.; Liz-Marzán, L. M. Nano Lett. 2015, 15, 5427−5437. (4) Zhu, X.; Zhuo, X.; Li, Q.; Yang, Z.; Wang, J. Adv. Funct. Mater. 2016, 26, 341−352. (5) Lim, B.; Kobayashi, H.; Yu, T.; Wang, J.; Kim, M. J.; Li, Z.-Y.; Rycenga, M.; Xia, Y. J. Am. Chem. Soc. 2010, 132, 2506−2507. (6) Zhu, C.; Zeng, J.; Tao, J.; Johnson, M. C.; Schmidt-Krey, I.; Blubaugh, L.; Zhu, Y.; Gu, Z.; Xia, Y. J. Am. Chem. Soc. 2012, 134, 15822−15831. (7) Yin, A.-X.; Min, X.-Q.; Zhang, Y.-W.; Yan, C.-H. J. Am. Chem. Soc. 2011, 133, 3816−3819. (8) Chen, M.; Liu, J. P.; Sun, S. J. Am. Chem. Soc. 2004, 126, 8394− 8395.
■
CONCLUSION Our real-time observations reveal how the reaction temperature affects the packing of the capping agents and how the facet selection of the depositing shell atoms drives the shell formation under different kinetic conditions. This type of information is difficult to obtain from any ex situ studies.16,27,28,36,41 Our study provides experimental evidence that the morphology of bimetallic core−shell NPs is determined by the ratio between the deposition and surface diffusion rates, which depends on the reaction temperature. These results suggest that it is possible to engineer core−shell NPs of different shapes by simply tuning the reaction temperature. Currently, the mechanistic descriptions of thermodynamically or kinetically controlled growth are mostly based on the studies of seed-mediated growth whereby core−shell NPs grow 11684
DOI: 10.1021/jacs.8b05217 J. Am. Chem. Soc. 2018, 140, 11680−11685
Article
Journal of the American Chemical Society (9) Habas, S. E.; Lee, H.; Radmilovic, V.; Somorjai, G. A.; Yang, P. Nat. Mater. 2007, 6, 692−697. (10) Sasaki, K.; Naohara, H.; Cai, Y.; Choi, Y. M.; Liu, P.; Vukmirovic, M. B.; Wang, J. X.; Adzic, R. R. Angew. Chem., Int. Ed. 2010, 49, 8602−8607. (11) DeSantis, C. J.; Weiner, R. G.; Radmilovic, A.; Bower, M. M.; Skrabalak, S. E. J. Phys. Chem. Lett. 2013, 4, 3072−3082. (12) Wang, D.; Li, Y. Adv. Mater. 2011, 23, 1044−1060. (13) Langille, M. R.; Zhang, J.; Personick, M. L.; Li, S.; Mirkin, C. A. Science 2012, 337, 954−957. (14) Pong, B.-K.; Elim, H. I.; Chong, J.-X.; Ji, W.; Trout, B. L.; Lee, J.-Y. J. Phys. Chem. C 2007, 111, 6281−6287. (15) Harada, M.; Katagiri, E. Langmuir 2010, 26, 17896−17905. (16) Xia, Y.; Xia, X.; Peng, H.-C. J. Am. Chem. Soc. 2015, 137, 7947− 7966. (17) Zheng, H.; Smith, R. K.; Jun, Y.-w.; Kisielowski, C.; Dahmen, U.; Alivisatos, A. P. Science 2009, 324, 1309−1312. (18) Williamson, M. J.; Tromp, R. M.; Vereecken, P. M.; Hull, R.; Ross, F. M. Nat. Mater. 2003, 2, 532−536. (19) de Jonge, N.; Ross, F. M. Nat. Nanotechnol. 2011, 6, 695−704. (20) Loh, N. D.; Sen, S.; Bosman, M.; Tan, S. F.; Zhong, J.; Nijhuis, C. A.; Král, P.; Matsudaira, P.; Mirsaidov, U. Nat. Chem. 2017, 9, 77− 82. (21) Wu, J.; Gao, W.; Wen, J.; Miller, D. J.; Lu, P.; Zuo, J.-M.; Yang, H. Nano Lett. 2015, 15, 2711−2715. (22) Sutter, E. A.; Sutter, P. W. J. Am. Chem. Soc. 2014, 136, 16865− 16870. (23) Tan, S. F.; Chee, S. W.; Lin, G.; Bosman, M.; Lin, M.; Mirsaidov, U.; Nijhuis, C. A. J. Am. Chem. Soc. 2016, 138, 5190− 5193. (24) Niu, K.-Y.; Liu, M.; Persson, K. A.; Han, Y.; Zheng, H. ACS Nano 2016, 10, 6235−6240. (25) Liang, W.-I.; Zhang, X.; Zan, Y.; Pan, M.; Czarnik, C.; Bustillo, K.; Xu, J.; Chu, Y.-H.; Zheng, H. J. Am. Chem. Soc. 2015, 137, 14850− 14853. (26) Liu, Q.; Leong, F. Y.; Aabdin, Z.; Anand, U.; Si Bui Quang, T.; Mirsaidov, U. ACS Nano 2015, 9, 9020−9026. (27) Jing, H.; Wang, H. CrystEngComm 2014, 16, 9469−9477. (28) Annan, W.; Qing, P.; Yadong, L. Chem. Mater. 2011, 23, 3217− 3222. (29) Sneed, B. T.; Young, A. P.; Tsung, C.-K. Nanoscale 2015, 7, 12248−12265. (30) Wang, L.; Liu, P.; Guan, P.; Yang, M.; Sun, J.; Cheng, Y.; Hirata, A.; Zhang, Z.; Ma, E.; Chen, M.; Han, X. Nat. Commun. 2013, 4, 2413. (31) Bhattarai, N.; Casillas, G.; Ponce, A.; Jose-Yacaman, M. Surf. Sci. 2013, 609, 161−166. (32) Kwon, S. G.; Krylova, G.; Phillips, P. J.; Klie, R. F.; Chattopadhyay, S.; Shibata, T.; Bunel, E. E.; Liu, Y.; Prakapenka, V. B.; Lee, B.; Shevchenko, E. V. Nat. Mater. 2015, 14, 215. (33) Gutkin, M. Y.; Smirnov, A. M. Acta Mater. 2015, 88, 91−101. (34) Tan, S. F.; Raj, S.; Bisht, G.; Annadata, H.; Nijhuis, C. A.; Král, P.; Mirsaidov, U. Adv. Mater. 2018, 30, 1707077. (35) Chen, H.; Shao, L.; Li, Q.; Wang, J. Chem. Soc. Rev. 2013, 42, 2679−2724. (36) Tsuji, M.; Ikedo, K.; Uto, K.; Matsunaga, M.; Yoshida, Y.; Takemura, K.; Niidome, Y. CrystEngComm 2013, 15, 6553−6563. (37) Gomer, R. Rep. Prog. Phys. 1990, 53, 917. (38) Bukaluk, A. Appl. Surf. Sci. 2001, 175−176, 790−796. (39) Becker, R.; Liedberg, B.; Käll, P.-O. J. Colloid Interface Sci. 2010, 343, 25−30. (40) Gou, L.; Murphy, C. J. Chem. Mater. 2005, 17, 3668−3672. (41) Xiang, Y.; Wu, X.; Liu, D.; Jiang, X.; Chu, W.; Li, Z.; Ma, Y.; Zhou, W.; Xie, S. Nano Lett. 2006, 6, 2290−2294. (42) Tan, S. F.; Chee, S. W.; Lin, G.; Mirsaidov, U. Acc. Chem. Res. 2017, 50, 1303−1312.
11685
DOI: 10.1021/jacs.8b05217 J. Am. Chem. Soc. 2018, 140, 11680−11685
