Surface Passivation and Supersaturation: Strategies for

Nov 22, 2017 - Further, understanding how anisotropic nanocrystals form is of interest as such insight may define mechanisms for site-specific growth,...
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Surface Passivation and Supersaturation: Strategies for Regioselective Deposition in Seeded Syntheses Alexander N. Chen,† Mattea M. Scanlan,§ and Sara E. Skrabalak*,† †

Department of Chemistry, Indiana University, 800 E. Kirkwood Ave., Bloomington, Indiana 47405, United States Department of Chemistry, Currens Hall 214, Western Illinois University, 1 University Circle, Macomb, Illinois 61455, United States

§

S Supporting Information *

ABSTRACT: Crystal growth theory predicts that heterogeneous nucleation will occur preferentially at defect sites, such as the vertices rather than the faces of shape-controlled seeds. Platonic metal solids are generally assumed to have vertices with nearly identical chemical potentials, and also nearly identical faces, leading to the useful generality that heterogeneous nucleation preserves the symmetry of the original seeds in the final product. Herein, we test the limits of this generality in the extreme of low supersaturation, in an effort to expand the methods available for inducing anisotropic overgrowth. We formulate a strategy for favoring localized deposition that differentiates between both different vertices and different edges or faces, i.e., regioselective deposition. Deposition followed a simple kinetic model for nucleation rate, depending on wetting, supersaturation, and temperature. We demonstrate our ability to independently study the effects of varying supersaturation and surface passivation. Regioselective heterogeneous nucleation was achieved at low supersaturation by a kinetic preference for high-energy defect-rich sites over lower-energy sites. This outcome was also achieved by using capping agents to passivate facet sites where deposition was not desired. Collectively, the results presented herein provide a model for breaking the symmetry of seeded growth and for achieving regioselective deposition. KEYWORDS: anisotropic nanoparticles, kinetic control, site-selective deposition, heterogeneous nucleation, gold, palladium

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reports demonstrate that heterogeneous nucleation occurs preferentially, though not exclusively, at defect sites, such as edges and vertices on a nanocrystal.10−12 Highly strained sites can also induce directional growth.13−15 Considering the differences in surface energy between identical (zero) and nonidentical (nonzero) pairs of elements, deposition after nucleation should favor sites with the deposited material rather than the seed material. When the rate of deposition is faster than the rate of adatom diffusion on the seed surface, nonconformal growth results, e.g., with branches growing at the vertices of a seed.10 Supersaturation and surface differentiation can then determine the mode of anisotropic growth. Excess surface capping has previously shown the ability to prevent deposition on capped surfaces.6,16,17 Low supersaturation has promoted anisotropic growth, for example, by growing on only one or a few faces of a nanocube.18,19 Combining the two effects should then help to promote highly anisotropic growth at only select sites on a nanocrystal seed. In contrast, at high supersaturation, there are enough monomers

anomaterials show great promise in areas from energy conversion1,2 to medicine;3,4 however, before realizing the promise of their form-function relationships, synthetic routes to nanomaterials with defined size, shape, and architecture are needed. Bottom-up syntheses in pursuit of well-defined nanoscale and macroscale properties must contend with size, shape, and architecture-related polydispersity at the nanoscale. Thus, mechanistic insights into nanoparticle growth are necessary to exert maximal control over nanocrystal growth and organization. Further, understanding how anisotropic nanocrystals form is of interest as such insight may define mechanisms for sitespecific growth, which presents a challenge in solution-phase syntheses.5−7 The diffuse layer surrounding a growing isotropic nanoparticle is likely to contain approximately the same local concentrations of incoming monomers in every direction. Symmetry breaking by growing on, for example, only one facet out of a set of symmetrically related facets is not evident. And when anisotropic growth is achieved, a detailed mechanism that explains how to encourage a particular mode of anisotropic growth is often absent.8,9 The literature provides some qualitatively supported mechanisms for promoting anisotropic growth. Previous © XXXX American Chemical Society

Received: October 4, 2017 Accepted: November 13, 2017

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Figure 1. SEM images depicting behavior of heterogeneous nucleation on Pd cubes as the amounts of metal precursors and capping agents are varied, with (A) 400 μmol CTAB, (B) 200 μmol CTAB and 200 μmol NaBr, (C) 200 μmol CTAB and 400 μmol NaBr, (D) 200 μmol CTAB and 600 μmol NaBr, and (1) 10 μL, (2) 25 μL, (3) 50 μL, and (4) 100 μL each of 10 μM H2PdCl4 and 1 mM HAuCl4.

volume ν, and supersaturation ratio S, where γ and ν are assumed to take their bulk values:21

to nucleate and then grow at each of the defect-rich sites, preserving the symmetry of the seed when defects are reliably located at symmetry-related vertices of a shape-controlled seed. A variety of seed shapes and experimental conditions demonstrated the latter principle in seed-mediated co-reduction (SMCR), a synthetic technique where two metal precursors are simultaneously reduced to deposit metal on seeds.11 Here, we test the limits of SMCR as a model system in order to produce guidelines for regioselective deposition on crystalline seeds, similar to the known guidelines for selecting reaction sites on small molecules. As with small molecules, regioselective reactions clearly prefer one reaction site, such as one vertex out of a set of symmetrically related vertices, but not to the exclusion of other, chemically similar reaction sites. In general, the design of syntheses promoting anisotropic nucleation requires fine kinetic control. Crystal growth theory provides a well-developed model for the rate of nucleation dN

ΔG hom =

(2)

As the energy barrier for heterogeneous nucleation is equal to that for homogeneous nucleation multiplied by the wetting factor f(θ), we obtain the energy barrier for heterogeneous nucleation in eq 3:22

ΔG het = ΔG homf (θ )

(3)

From eqs 2 and 3, the theoretical influence of supersaturation on heterogeneous nucleation is evident: as S is greater than 1 for a supersaturated solution, increasing the supersaturation will always decrease the energy barrier for nucleation, thus increasing the nucleation rate. The wetting factor largely captures two variables at the seed surface. First, wetting implies removal of adsorbates such as capping agents as a prerequisite to forming a new metal−metal interface. Strongly bound capping agents should impose an additional energy barrier for their removal prior to the formation of a new interface, effectively increasing the wetting factor and so decreasing the heterogeneous nucleation rate. Second, regions of the seed that possess high defect density should readily form bonds between particularly undercoordinated surface atoms and the depositing material, effectively decreasing the wetting factor. And while capping agents that interact with the nucleating material may

dt

of isotropic structures,20,21 which approximates nucleation for other structures. From an energy barrier for nucleation ΔGN, the temperature T, the Boltzmann constant kB, and a preexponential factor A, the nucleation rate is given as ⎛ ΔG N ⎞ dN = A exp⎜ − ⎟ dt ⎝ kBT ⎠

16πγ 3ν 2 3kB 2T 2(ln S)2

(1)

The energy barrier for homogeneous nucleation is defined in eq 2, for a surface energy γ of the nucleating material, atomic B

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Figure 2. (A, G, M, S) Enlarged SEM images, (B, H, N, T) TEM images, and (C−F, I−L, O−R, U−X) STEM-EDS elemental maps of samples (A−F) A2, (G−L) C2, (M−R) A4, and (S−X) C4 from Figure 1. Yellow, Au; Red, Pd.

such a characteristic indicates Ratedeposition < Ratediffusion).10,11,23,24 Thus, relatively fast deposition is achieved by increasing defect richness and supersaturation, and reducing the passivation of targeted surface sites on the seed. In contrast, relatively slow deposition is achieved by reducing defect richness and supersaturation, while increasing the passivation of targeted surface sites.

also decrease the energy barrier for homogeneous nucleation by decreasing γ of the nucleating material, eq 3 shows that the energy barrier for heterogeneous nucleation will decrease by the same factor. So, we anticipate that the heterogeneous nucleation rate will be small on strongly passivated surfaces and high on weakly passivated, defect-rich surfaces. We observed morphological patterns associated with nucleation rate after tuning supersaturation, capping and, to a lesser degree, temperature. Overall, the final morphologies suggested that heterogeneous nucleation rate was the crucial kinetic factor in directing regioselective deposition. Supersaturation increased with increasing amounts of metal precursors. Wetting decreased on select surfaces by using seeds terminating with low-energy facets, and by adding surface-selective capping agents. Temperature was directly controlled by conducting the syntheses in an oil bath. The rates of nucleation and subsequent deposition could be determined qualitatively by comparison to the rate of adatom diffusion, the latter of which was assumed to remain constant under experimental conditions of constant temperature. The deposition rate was considered relatively fast when the deposited material was highly localized, and relatively slow when the deposited material was less localized on the seed surface (as

RESULTS AND DISCUSSION Applying the kinetic model for nucleation rates of isotropic nanostructures revealed competition between different surface sites in nanocrystal growth. Varying metal precursor concentrations gave different levels of supersaturation. Natural differences in defect density on nanocrystal seeds and the use of surface-selective capping agents made clear the effect of changing the seed’s surface chemistry. Deposition on Seeds with Crystallographically Equivalent Vertices: Cubes and Octahedra. In order to deconvolute the components of the rate expression, HAuCl4 and H2PdCl4 were co-reduced at 25 °C using L-ascorbic acid in the presence of Pd nanocubes capped with cetyltrimethylammonium bromide (CTAB) (Figure S1). Full experimental C

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shows that the Pd−Au surfaces bind NaBr more strongly and possibly more selectively than they bind CTAB, presumably due to the difference in electronic states of adsorbed bromide between CTAB and NaBr. And since wetting on the Pd surface involves removal of adsorbates, the more strongly bound bromide should impose a higher energy cost for desorption and so inhibit wetting. Enlarged SEM images from Figure 1, transmission electron microscopy (TEM) images, and elemental mapping by scanning transmission electron microscopy coupled with energy dispersive X-ray spectroscopy (STEM-EDS) of samples A2, C2, A4 and C4 are presented in Figure 2A−F, G−L, M−R, and S− X, respectively. Comparison of columns A and C, and of rows 2 and 4, make clear the significant differences in product caused by tuning supersaturation and capping in the reaction solution. Sample A4 shows a core−shell architecture, where a Au-rich phase is conformally deposited around the Pd seeds. Decreasing the supersaturation yields sample A2, which possesses highly diffuse Au-rich domains, but not enough deposited material to form a complete shell. Instead, the domains are elongated along edges and faces of the seed, but still unambiguously visible in TEM and STEM images (Figure 2B−C). Enlarged SEM images show little atomic number contrast (Figure 2A, M). Similar products are observed for sample A1 (Figure S6). Increasing the capping relative to sample A4 causes a progression from conformal to island growth, shown in both SEM and TEM for sample C4 (Figure 2S−T). Islands are generally also localized at the vertices, as evidenced by the lack of Au signal on the cube face in the STEM-EDS map (Figure 2U−X); however, some islands also grow on cube edges. Decreasing the supersaturation relative to sample C4 yields sample C2, whose elemental map displays smaller localized Au domains at the vertices (Figure 2I−L). In addition, the Au phases appear more localized at low supersaturation than they do at high supersaturation, without changing the amount of capping agent. From the elongated shape of the deposited phase in samples C2 and C4, growth appears to occur both outward into the reaction solution and along the edge of the Pd seeds (Figure 2G−H, S−T). The effects of supersaturation and surface passivation are intertwined during nucleation and growth. Sample C2 also shows more clearly than sample C4 that regioselectivity with respect to individual vertices is obtained at low supersaturation: not every vertex has had a successful nucleation event (Figure 2G−L). The Oh symmetry of the initial cube is broken, and the deposited phase is directed selectively toward only some of the seed’s vertices by a combination of decreased supersaturation and increased surface passivation. Asymmetric nucleation with respect to the vertices implies a difference in chemical potential between theoretically identical vertices, which was thought to arise in part from differences in defect richness between physically nonequivalent vertices of an atomically imperfect cube (Figure S7). Figure 3 shows structural models summarizing the trends observed for variations of surface passivation and supersaturation in Figures 1 and 2. Poor {100} capping gives diffuse deposited phases, which may form islands in the absence of sufficient material to form a full shell. More capping inhibits growth along faces and, to a lesser extent, edges. This situation results in larger and more defined domains growing at the vertices. At high supersaturation, the effect of capping is somewhat attenuated, as inhibiting wetting on a face decreases the rate of nucleation on that face while increasing S increases

details are in the Supporting Information. Figures 1 and 2 show the results obtained by varying the concentrations of the metal precursors (held at a 100-to-1 Au-to-Pd mole ratio) as well as the amount of an additional capping agent, NaBr. As the amount of the metal precursors increases from 10 to 250 nmol HAuCl4 (rows 1−4), the deposited phases grew increasingly large. Increasing amounts of NaBr from 0 to 600 μmol (columns A−D) yielded more localized Au-rich domains, which are evident in scanning electron microscopy (SEM) images in columns C and D by atomic number contrast, and which are preferentially situated at the vertices of the Pd cubes, i.e., where contact with {100} surfaces is minimized. Previous work shows that bromide preferentially adsorbs to the {100} faces of Pd nanocrystals, producing and stabilizing cubic morphologies.16,18,25,26 Thus, bromide-containing species raised the energy barrier for nucleation on {100} facets relative to other surfaces that were available for nucleation events during deposition on Pd nanocubes. At lower amounts of NaBr, the Au domains are less distinct, until they are nearly indistinguishable in the absence of NaBr (column A). The general morphological trends agree with the implications of the nucleation rate model, where the energy barrier for nucleation at a given surface site is predicted to increase both with stabilizing the seed’s surface by capping and with decreasing supersaturation. Thus, nucleation rate can become vanishingly small when, at low supersaturation, capping agents further disfavor wetting at select sites. The reaction solutions used to prepare the products in columns A and B contained identical molar amounts of bromide ions, but the origin of the bromide is different between these samples (CTAB only vs CTAB and NaBr at a 1-to-1 mole ratio). Interestingly, the product particles have different morphologies. Particles in column B show more localized Aurich domains, suggesting that better surface passivation is achieved in the presence of NaBr than there is in the presence of CTAB. It follows that the counterion must influence adsorption to the Pd surface. Indeed, this difference is evident from analysis of surface bromide species in the washed products by X-ray photoelectron spectroscopy (XPS). The Br 3d binding energies of reference CTAB and NaBr were found to be 67.0 and 69.0 eV, respectively, allowing for easy differentiation of the two (Figures S2, S3). The lower binding energy for CTAB suggests that the soft bromide may interact more strongly with the soft ammonium group than with the hard sodium cation. Sample A2 was fitted using a single doublet with a 3d5/2 binding energy of 67.0 eV. Sample B2 was fitted using two doublets, with 3d5/2 binding energies of 67.3 and 68.3 eV, with the higher binding energy doublet accounting for 62.7% of the peak area. The 67.0 and 67.3 eV binding energies are very similar to that of unbound CTAB; therefore, they were attributed to adsorbed CTAB. The higher binding energy doublet must then originate from NaBr, and its greater intensity indicates that Na+associated bromide is disproportionately present compared to CTA+-associated bromide on the washed surface. Also, the larger shift from 69.0 to 68.3 eV suggests that bromide from NaBr interacts more strongly with the surface (Figures S4, S5). Na 1s and N 1s peaks in the survey scans of washed samples further support differentiation of CTAB and NaBr (Figure S4), again showing that the strongly bound counterions continue to play a significant role at the particle surfaces. Strong interactions between the particle surface and the adsorbed bromide, especially for NaBr, account for differences in binding energy compared to the bulk references. Overall, XPS analysis D

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edges and with very few growing from faces. Elemental mapping shows that effectively all of the deposited Au is concentrated in these spheres (Figure 4D−F). In all of Figure 4, the symmetry of the original seed is lost. The relative lack of growth along a Pd−Au interface was attributed to the low surface energy of stable Pd{111} surfaces compared to Pd{100} surfaces. In addition, the high stability of Pd{111} faces is expected to favor surface diffusion less than the less stable Pd{100} faces of cubes, since diffusing atoms require attractive interactions with undercoordinated surface atoms, and such interactions are stronger with more undercoordinated atoms.27 Studies performed with only HAuCl4 or only H2PdCl4 as metal precursors produced similar products with HAuCl4 as the only metal precursor, and apparently unmarked octahedra with H2PdCl4 as the only metal precursor (Figure S9). This is to say that nucleation occurs preferentially at defect sites, with poor wetting for Au on Pd, but with sufficiently favorable wetting and subsequent diffusion of Pd on Pd to not produce a distinguishable deposited domain. In this case, the main benefit of SMCR over reduction of a single metal is the increased stability of the Au−Pd phase over pure Au systems, which has been noted in previous work.10 For the system with cubic seeds, the elongated Au-rich phases in Figure 2 suggest that diffusion of Au on Pd{100} is fast enough to compete with deposition, allowing for better apparent wetting. The core−shell structure in sample A4 presumably combines competitive diffusion with less passivated {100} surfaces. Increasing the temperature from 25 °C to 30 °C and 40 °C during the syntheses of samples B2 and D2 further supports the assertion that the competition between diffusion and deposition contributes to morphology development by producing results similar to sample A2, despite the nanocubes having highly passivated faces (Figure S10). Nevertheless, it should be noted that in eq 2, increasing T decreases the energy barrier for nucleation, so the effects of changing temperature are entangled. Still, the comparison between cubes and octahedra gives a useful idea of how the rates of diffusion compare to the rates of deposition at different surface sites, adding relevant kinetic parameters to the model. Clarifying the combined effects of supersaturation and capping on octahedra required a synthesis where significant deposition on the faces of the seed could be expected. According to eq 2, high supersaturation can help in achieving kinetically unfavorable deposition on stable faces; therefore, the amount of surface passivation was varied at very high supersaturation. When using {100}-terminated cubes, increasing the supersaturation relative to columns A and B of Figure 1 conforms to literature results.11,23 Eight sharp branches cover the whole cube with only CTAB as capping agent, but open square faces separate eight blunt and fused branches when both CTAB and NaBr are used (Figure S11). The formation of hopper-like nanocrystals was previously attributed to slowed precursor reduction due to complexation of metal ions with bromide and to the presence of high-index facets at the edge of the triangular face;23,28,29 however, this explanation may be expanded upon consideration of Figure 1. Figure 1 suggests that a second kinetic parameter associated with bromide is contributing: selective passivation of the square {100} facets on a cube selectively slows down face deposition and allows edge and vertex growth, producing only square concavities. When using {111}-terminated octahedra, increasing the supersaturation yields 24-branched structures with Oh symmetry (4 branches per seed vertex), as also reported from the literature.11

Figure 3. Schematic representation of morphological trends resulting from varying {100} capping and supersaturation.

it, and so yielding more comparable rates of nucleation along the seed’s faces relative to its vertices. In order to test the generality of conclusions drawn from using {100}-terminated cubic seeds, {111}-terminated octahedral seeds (Figure S8) were used in SMCR experiments analogous to those from column A in Figure 1. The resulting products were morphologically closer to column C than to column A, despite being synthesized in the presence of the selectively {100}-capping agent CTAB and no selectively {111}-capping agent. As shown in Figure 4A, the initial octahedral seeds are still visible and are attached to elongated domains of the deposited material. As supersaturation decreases in Figure 4B−C, the deposited material forms smaller, highly discrete spherical domains, mostly growing from vertices and

Figure 4. SEM images of product obtained when adding (A) 1 mL, (B) 100 μL, and (C) 50 μL of 10 mM HAuCl4 and 100 μM H2PdCl4 solutions to octahedral Pd seeds. (D) STEM image of a typical particle, and (E−G) elemental map by STEM-EDS analysis. Yellow, Au; Red, Pd. E

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ACS Nano Further increasing {111} face capping by addition of citric acid30,31 reduced the pH from 2.8 to 2.5 and produced more polydisperse particles. Addition of citric acid with the pH change neutralized by NaOH yielded visibly unmarked faces and overgrowth concentrated at the vertices (Figure S12). Overgrowth was also much less directional, taking the form of polyhedral domains rather than sharp branches. The change in overgrowth shape was ascribed to the passivation of newly formed {111} facets after nucleation, inhibiting further branch growth in the ⟨111⟩ direction. While very high supersaturation does support the dual roles of supersaturation and capping in directing growth away from relatively stable facets, it poses the additional problem that there is enough deposited material to produce significantly more surfaces for capping agents to passivate. Deposition on Seeds with Crystallographically Inequivalent Vertices: Right Bipyramids. Pd right bipyramidal seeds exemplify another method for using surface energies to guide regioselective nucleation. Produced as a major impurity from an adapted literature synthesis11 for cubes (Figure S13), Pd right bipyramids consist of six exposed {100} facets divided into two sets of three by an equatorial (111) twin plane.32 The two axial vertices terminate as {111} surfaces, as with cubes; however, the three equatorial vertices lie on the twin plane, and are different from the vertices of a cube. Thus, right bipyramids are expected to exhibit the same surface chemistry as cubes, except that they also possess a well-defined planar defect. SMCR performed at high supersaturation has previously yielded branch growth on each of the axial and equatorial vertices indiscriminately, giving rise to pentapods with D3h symmetry just like the seeds.11 SMCR performed at low supersaturation yielded products that typically showed bare axial vertices in TEM images (Figure 5A,B). Electron diffraction on a particle oriented near its C3 axis reveals that the C3 axis is also the ⟨111⟩ zone axis (Figure 5B, inset). A typical STEMEDS map shows a Au-rich phase surrounding the twin plane and not progressing to the single-crystalline axial vertices (Figure 5C−F). Recall that cubes under the same conditions produced diffuse Au-rich domains with little localization. Elemental mapping of more particles shows that Au overgrowth almost never reaches the visible axial vertex (Figure S14). Line scan analysis shows that the Au signal near the C3 axis is generally close to the Sm signal used to represent the background (Figure 5G). Weak, isolated Au signals near the C3 axis indicate that any Au found on the faces and axial vertices of the right bipyramid is not enough to form a complete shell and comprises a small minority of deposited material, as illustrated in Figure 5H. Notably, when the Au domain did not cover the entire perimeter of the equatorial plane, it still did not expand toward the center (Figure S14). These results align with a previous study concerning deposition of Pd on Pd right bipyramids,32 with the added benefit of Au−Pd contrast. Again, low supersaturation provides greater regioselectivity than high supersaturation. A discussion of varying nucleation rates on the surface of a single nanoparticle must be framed qualitatively in terms of defect density. Heterogeneous nucleation has been shown to be faster on defect-dense regions such as vertices than on defectpoor regions such as faces. In the case of right bipyramids, the two axial vertices have shown themselves to be significantly lower in energy than the twin plane defined by both the three equatorial vertices and the connecting edges. Wetting of equatorial edges and vertices over axial vertices suggests that

Figure 5. TEM images of (A) right bipyramids with overgrowth, (B) single particle and (inset) corresponding electron diffraction. (C) STEM image, (D−F) elemental map by EDS analysis and (G) line scan starting from the top of the yellow line in panel C, with Sm as background reference. (H) Structural models depicting right bipyramids before and after overgrowth.

surface sites located on planar defects have much higher energies than those on crystalline vertices. When comparing to Figure 2, the twin plane of a right bipyramid is far easier to wet than the crystalline faces, edges, and vertices of either right bipyramids or cubes, the latter of which required strong capping from NaBr to achieve similar regioselectivity. In other words, the rate of deposition at the twin plane is far greater than the rate of diffusion down the {100} faces.

CONCLUSIONS In order to exert fine control over the growth kinetics of heterogeneous nucleation, crystal growth theory predicts that manipulation of supersaturation, surface passivation, and temperature are key to achieving unusual morphologies. Using SMCR as a model system, tuning metal precursor concentration, surface passivation, and defect richness exerts kinetic control over nucleation events and over the final morphology, to the extent that nucleation differentiates between theoretically identical vertices. High supersaturation produces high rates of nucleation on more surfaces, so the majority of the results presented apply to low supersaturation. The results presented here theoretically also apply to homogeneous nucleation, as the variables noted above are all expected to play the same roles, aside from the wetting factor. However, achieving suitable supersaturation conditions may be more challenging given that the symmetry reduction occurs at the extreme of low supersaturation in our work and homogeneously nucleated particles often start out isotropic in contrast to the seeds used in these studies. From the theoretical F

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(2) Lai, J.; Luque, R.; Xu, G. Recent Advances in the Synthesis and Electrocatalytic Applications of Platinum-Based Bimetallic Alloy Nanostructures. ChemCatChem 2015, 7, 3206−3228. (3) He, C.; Liu, D.; Lin, W. Nanomedicine Applications of Hybrid Nanomaterials Built from Metal-Ligand Coordination Bonds: Nanoscale Metal-Organic Frameworks and Nanoscale Coordination Polymers. Chem. Rev. 2015, 115, 11079−11108. (4) Liao, L.; Liu, J.; Dreaden, E. C.; Morton, S. W.; Shopsowitz, K. E.; Hammond, P. T.; Johnson, J. A. A Convergent Synthetic Platform for Single-Nanoparticle Combination Cancer Therapy: Ratiometric Loading and Controlled Release of Cisplatin, Doxorubicin, and Camptothecin. J. Am. Chem. Soc. 2014, 136, 5896−5899. (5) Burrows, N. D.; Vartanian, A. M.; Abadeer, N. S.; Grzincic, E. M.; Jacob, L. M.; Lin, W.; Li, J.; Dennison, J. M.; Hinman, J. G.; Murphy, C. J. Anisotropic Nanoparticles and Anisotropic Surface Chemistry. J. Phys. Chem. Lett. 2016, 7, 632−641. (6) Cathcart, N.; Kitaev, V. Symmetry Breaking by Surface Blocking: Synthesis of Bimorphic Silver Nanoparticles, Nanoscale Fishes and Apples. Sci. Rep. 2016, 6, 32561. (7) Read, C. G.; Gordon, T. R.; Hodges, J. M.; Schaak, R. E. Colloidal Hybrid Nanoparticle Insertion Reaction for Transforming Heterodimers into Heterotrimers. J. Am. Chem. Soc. 2015, 137, 12514−12517. (8) Fantechi, E.; Roca, A. G.; Sepúlveda, B.; Torruella, P.; Estradé, S.; Peiró, F.; Coy, E.; Jurga, S.; Bastús, N. G.; Nogués, J.; Puntes, V. Seeded Growth Synthesis of Au−Fe3O4 Heterostructured Nanocrystals: Rational Design and Mechanistic Insights. Chem. Mater. 2017, 29, 4022−4035. (9) Burrows, N. D.; Harvey, S.; Idesis, F. A.; Murphy, C. J. Understanding the Seed-Mediated Growth of Gold Nanorods through a Fractional Factorial Design of Experiments. Langmuir 2017, 33, 1891−1907. (10) Weiner, R. G.; DeSantis, C. J.; Cardoso, M. B. T.; Skrabalak, S. E. Diffusion and Seed Shape: Intertwined Parameters in the Synthesis of Branched Metal Nanostructures. ACS Nano 2014, 8, 8625−8635. (11) DeSantis, C. J.; Skrabalak, S. E. Core Values: Elucidating the Role of Seed Structure in the Synthesis of Symmetrically Branched Nanocrystals. J. Am. Chem. Soc. 2013, 135, 10−13. (12) Weiner, R. G.; Kunz, M. R.; Skrabalak, S. E. Seeding a New Kind of Garden: Synthesis of Architecturally Defined Multimetallic Nanostructures by Seed-Mediated Co-Reduction. Acc. Chem. Res. 2015, 48, 2688−2695. (13) Luo, M.; Huang, H.; Choi, S. I.; Zhang, C.; da Silva, R. R.; Peng, H. C.; Li, Z. Y.; Liu, J.; He, Z.; Xia, Y. Facile Synthesis of Ag Nanorods with No Plasmon Resonance Peak in the Visible Region by Using Pd Decahedra of 16 nm in Size as Seeds. ACS Nano 2015, 9, 10523− 10532. (14) Gilroy, K. D.; Peng, H. C.; Yang, X.; Ruditskiy, A.; Xia, Y. Symmetry Breaking During Nanocrystal Growth. Chem. Commun. 2017, 53, 4530−4541. (15) Kunz, M. R.; McClain, S. M.; Chen, D. P.; Koczkur, K. M.; Weiner, R. G.; Skrabalak, S. E. Seed-mediated Co-reduction in a Large Lattice Mismatch System: Synthesis of Pd-Cu Nanostructures. Nanoscale 2017, 9, 7570−7576. (16) Peng, H.-C.; Li, Z.; Aldahondo, G.; Huang, H.; Xia, Y. SeedMediated Synthesis of Pd Nanocrystals: The Effect of Surface Capping on the Heterogeneous Nucleation and Growth. J. Phys. Chem. C 2016, 120, 11754−11761. (17) Lee, K. W.; An, H.; Haam, S.; Baik, H.; Lee, K. Regiospecific Growth of Au on a Concave PtZn Nanocube Forming an Au−PtZn Surface Mosaic Nanocube and an Au−PtZn Octapod. CrystEngComm 2015, 17, 6838−6842. (18) Zeng, J.; Zhu, C.; Tao, J.; Jin, M.; Zhang, H.; Li, Z. Y.; Zhu, Y.; Xia, Y. Controlling the Nucleation and Growth of Silver on Palladium Nanocubes by Manipulating the Reaction Kinetics. Angew. Chem., Int. Ed. 2012, 51, 2354−2358. (19) Peng, H.-C.; Park, J.; Zhang, L.; Xia, Y. Toward a Quantitative Understanding of Symmetry Reduction Involved in the Seed-Mediated Growth of Pd Nanocrystals. J. Am. Chem. Soc. 2015, 137, 6643−6652.

variables involved in nucleation rate, the work presented here leads to three general strategies for achieving regioselective heterogeneous nucleation: (1) Decreasing the supersaturation raises the energy barrier for nucleation at all available nucleation sites, forcing a kinetic choice between energetically different nucleation sites. This generality is exemplified by selective vertex deposition at low supersaturation. (2) Increasing the passivation of select surfaces inhibits wetting where it is not desired, thus increasing the energy barrier for nucleation. This generality is shown in the contrast between delocalized face, edge, and vertex growth on seeds with many high-energy surfaces, such as poorly capped Pd{100} and localized vertex growth on seeds with mostly low-energy surfaces, such as Pd{111} or well-capped Pd{100}. (3) Defect-rich regions engineered into seeds provide easily wetted, favorable nucleation sites. Edges and vertices of single-crystalline seeds have exhibited this behavior previously; the planar defects in right bipyramids are also preferential over vertices without a planar defect.

METHODS See Supporting Information for chemicals, characterization methods, and detailed synthetic procedures. Pd cubes33 and octahedra11 were synthesized according to the literature. Pd right bipyramids were synthesized according to a modified literature procedure.11 Seedmediated co-reduction was carried out according to the literature,11 using varying volumes of 10−100 μM H2PdCl4 and 1−10 mM HAuCl4·3H2O, and adding in 0.2 M NaBr, 0.2 M citric acid and 1 M NaOH as necessary.

ASSOCIATED CONTENT S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsnano.7b07041. Detailed experimental procedures and additional materials characterization (PDF)

AUTHOR INFORMATION Corresponding Author

*[email protected]. ORCID

Sara E. Skrabalak: 0000-0002-1873-100X Notes

The authors declare no competing financial interest.

ACKNOWLEDGMENTS This work was supported by NSF Award CHE-1602476. M. Scanlan was supported with a Research Experience for Undergraduates through NSF Award CHE-1460720. The authors would like to thank the IU Nanoscale Characterization Facility for access to instrumentation, Dr. David Morgan for assistance with TEM, and Dr. Yaroslav Losovyj for assistance with XPS. REFERENCES (1) Wang, H.; Zhang, L.; Chen, Z.; Hu, J.; Li, S.; Wang, Z.; Liu, J.; Wang, X. Semiconductor Heterojunction Photocatalysts: Design, Construction, and Photocatalytic Performances. Chem. Soc. Rev. 2014, 43, 5234−5244. G

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DOI: 10.1021/acsnano.7b07041 ACS Nano XXXX, XXX, XXX−XXX