Communication pubs.acs.org/cm
Observation of Surface Atoms during Platinum Nanocrystal Growth by Monomer Attachment Myoungho Jeong,†,‡,§ Jong Min Yuk,†,‡,§,# and Jeong Yong Lee*,†,‡ †
Department of Materials Science and Engineering, Korea Advanced Institute of Science and Technology (KAIST), Daejeon 305-701, Korea ‡ Center for Nanomaterials and Chemical Reactions, Institute for Basic Science (IBS), Daejeon 305-701, Korea S Supporting Information *
T
he electronic, optical, and catalytic properties of nanocrystals are strongly influenced by their size, shape, and morphology.1−3 Such structural properties of nanocrystals are mainly controlled or decided during the nanocrystal growth process and structural relaxation steps. Nanocrystal growth mechanisms have been classically considered as monomer attachment from solution or unstable small particles.4−7 On the other hand, nonclassical growth mechanisms such as aggregation,8−10 coalescence,11−13 and oriented attachment8,14 have recently been revealed to play a significant role in nanocrystal growth. Such nanocrystal growth mechanisms generally depend on overall growth rates.15 Under the conditions of a low growth rate, nanocrystals grow under a reaction-limited growth regime and have eventually faceted morphology. When the growth rate is high, however, polyhedral or near spherical nanocrystals are formed under a diffusion-limited growth regime. Once nanocrystal coalescence occurs prior to structural relaxation by surface diffusion or grain boundary migration,16 in addition, nanocrystals grow with anisotropic shapes such as rods or wires.17 Recent advances in in situ electron microscopy (EM) techniques have allowed direct observation of colloidal nanocrystal growth and have elucidated nonclassical growth mechanisms.11,18−20 Although atomic observations of nonclassical growth of nanocrystals have been reported, direct observation of surface atoms during monomer attachment growth has been hard to achieve because single atom attachments occur too rapidly to be detected on the nanocrystal surfaces.21 Recently, the atomic-resolution imaging of colloidal nanocrystal growth was demonstrated using in situ graphene liquid cell EM (GLC-EM) due to the low background noise of the GLC.12 In this communication, using in situ GLC-EM, we monitored time-serial trajectories of discrete atoms on Pt nanocrystal surfaces during monomer attachment growth and structural relaxation and demonstrated facet formation. To track surface atoms with a second temporal resolution during monomer attachment growth, we controlled three experimental conditions. First, we fabricated GLC encapsulating dilute Pt precursors (8 mg/mL), which increase the diffusion length (l) of monomers arriving at the nanocrystal surfaces and slow the atomic diffusion time (t) following Fick’s law (t ∼ l2). Figure 1a,b shows respectively a bright-field TEM image and the corresponding schematic illustration of GLC encapsulating dilute Pt growth solution between two laminated graphene layers. Two intact graphene sheets are bonded via van © XXXX American Chemical Society
Figure 1. (a) Bright-field TEM image of the intact GLC and (b) the corresponding schematic illustration. (c) Atomic force microscopy (AFM) image of (a). (d) Dark-field TEM image of (a) after e-beam irradiation for 180 s. Scale bars in (a) and (d) represent 50 nm.
der Waals attraction where the Pt growth solution has been removed and the remaining solution is entrapped and protected against evaporation of solution in the TEM with high vacuum.12,22,23 Second, we observed the edges of the liquid pockets with liquid thickness less than ∼10 nm (Figure 1c). The low quantity of Pt precursors produces sparse nanocrystal nucleation (∼0.01/nm2 in Figure 1d and Figure S1 in the Supporting Information) and prevents frequent nanocrystal coalescence after constant electron (e) beam irradiation with beam current of ∼4 × 108 A/cm2, which induces reduction of Pt precursors by the solvated free electrons.15 Third, we observed nearly fixed nanocrystals with diameters (d) smaller than 5 nm (Figure S2 in the Supporting Information) in the final growth stage because the nanocrystals become more stationary with growth due to the increase of the diffusion coefficient (D) following the Einstein−Stoke’s equation (D ∼ 2/d). In addition, e-beam irradiation for a long time at a fixed area results in solution being pushed out of view and thinning Received: January 27, 2015 Revised: April 22, 2015
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DOI: 10.1021/acs.chemmater.5b00334 Chem. Mater. XXXX, XXX, XXX−XXX
Communication
Chemistry of Materials
{144}, and {433}) coexist until 5 s after observation. The high indexed facets are gradually reduced and disappeared in 26 s, and only the {111} and {100} facets remain, which indicates that high-energy surfaces grow faster than low-energy surfaces and vanish as the nanocrystal grows, leading to nanocrystals with low energy facets.25 The {111} planes remarkably cover the full-grown nanocrystal facets. The {111} planes of the facecentered cubic crystal are the most stable planes with the lowest surface energy, and, thus, monomer attachments occur with the slowest rate on these planes for small nanocrystals in which monomers can freely adhere to surfaces, neglecting ligand effects.26 The {110} facets vanish because the {110} facets are intermediate surfaces between the {111} and {100} surfaces and have higher surface energy than either the {111} or {100} facets.21 We tracked the surface atoms of the nanocrystal during monomer attachment growth. Figure 3 shows sequential AR-
liquid thickness, which also slowed nanocrystal motion due to the substrate−nanoparticle interactions.24 The nanocrystal radius (r = (π/A)1/2, A is an area of the nanocrystal) continuously increases as t1/20 ∼ t1/7 except for occasional jumps owing to coalescence (Figure S3 in the Supporting Information), implying that monomer attachment is the dominant growth process for these nanocrystals. The increment of r is smaller than the t1/2 or t1/3 scaling predicted by a reaction- or diffusion-controlled growth, respectively, in the classical growth model. The smaller scaling of r may result from suppressed diffusion rate of monomers in our experimental conditions.15 We monitored facet development of Pt nanocrystals. Figure 2a shows the sequential atomic resolution TEM (AR-TEM)
Figure 3. Sequential AR-TEM images show the attachment of atoms on (a) {100} and (b, c) {111} facets during Pt nanocrystal growth. The blue and red dots indicate the newly growing Pt atoms with or without stacking fault (SF), respectively. Yellow dots indicate Pt atoms in pregrown planes. Scale bars represent 1 nm.
TEM images of the atomic attachment on {100} and {111} surfaces (Figure S5 in the Supporting Information). The discrete atoms (indicated by red dots in Figure 3) adhere to either the {100} or the {111} surface and nucleated (Figure 3a,b). The atoms that arrived later adhere beside the nucleated adatoms, and a new surface layer is eventually formed. Such an atomic attachment process leads the ledge-by-ledge growth of the {100} and {111} facets. The ledges on the {111} surfaces grow with atomic stacking sequence of either ABCA, normal stacking sequence, or ABC|B, stacking faulted sequence (Figure 3b,c). The stacking fault on the surface should be eliminated because the defects, such as stacking faults, grain boundaries, or dislocations, in a small homogeneous nanocrystal are energetically unstable.27 When the atoms that arrived on the ABC|B stacked ledges adhere on the C sites and two stacking faulted ledges (indicated by blue dots in Figure 3c) are built, the stacking fault is annealed by movement of the stacking faulted layers along the Burgers vector of 1/6 ⟨12̅1⟩. Once nanocrystal growth by attachments of monomers stops with imperfect ledges, the adatoms diffuse to stabilize the facets during structural relaxation after growth. Figure 4 shows sequential AR-TEM images of surface reshaping in a full-grown Pt nanocrystal (Figure S6 in the Supporting Information). The green and blue regions respectively indicate emerging and disappearing atoms with time that occurred by surface diffusion. The adatoms diffuse on the {100} or {111} surfaces to fill imperfect ledges, and, thus, unstable step edges are disappeared. The surface diffusion of adatoms can be assisted by the e-beam irradiation.16 In summary, we have observed surface atoms during Pt nanocrystal growth by attachments of monomers using in situ
Figure 2. (a) Sequential AR-TEM images of Pt nanocrystal growth by monomer attachment. Scale bar is 2 nm. Nanocrystal surface coverage of (b) {011}, {411}, {233}, {522}, {144}, and {433} facets and (c) {111} and {100} facets as a function of e-beam irradiation time during growth.
images during monomer attachment growth and structural relaxation (Figure S4 in the Supporting Information). The nanocrystals at the initial growth stage have a nearly spherical shape with ∼3.3 nm diameter. High indexed facets grow rapidly and vanish with time and, thus, the number of facets decreases. After e-beam irradiation for 32 s, the final nanocrystal with ∼3.8 nm diameter have a hexagonal shape with {100} and {111} facets. Nanocrystal facet development as a function of time is plotted in Figure 2b,c. The low indexed facets ({111}, {100}, and {011}) and the high indexed facets ({411}, {233}, {522}, B
DOI: 10.1021/acs.chemmater.5b00334 Chem. Mater. XXXX, XXX, XXX−XXX
Communication
Chemistry of Materials
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Figure 4. Sequential AR-TEM images of a Pt nanocrystal during structural relaxation. The green and blue regions indicate emerging and disappearing atoms with time, respectively. Scale bar is 2 nm.
GLC-EM. The final facets of the Pt nanocrystals consist of only {100} and {111} planes due to their low growth rates. These facets are formed through the ledge-by-ledge growth. During ledge growth, stacking faults can be formed but are annealed with nanocrystal growth. After growth, imperfect {100} or {111} ledges are stabilized during structural relaxation. Our findings will be likely of benefit to understand the monomer attachment growth mechanism of the nanocrystals.
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ASSOCIATED CONTENT
* Supporting Information S
Methods, Figures S1−S6, and Movies S1 and S2. The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.5b00334.
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AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. Present Address #
Present address for J.M.Y.: Department of Physics, University of California at Berkeley, California 94720, United States Author Contributions §
(M.J. and J.M.Y.) These authors contributed equally.
Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS This work was supported by IBS-R004-G3. REFERENCES
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DOI: 10.1021/acs.chemmater.5b00334 Chem. Mater. XXXX, XXX, XXX−XXX
