ARTICLES PUBLISHED ONLINE: 2 NOVEMBER 2014 | DOI: 10.1038/NMAT4115

Heterogeneous nucleation and shape transformation of multicomponent metallic nanostructures Soon Gu Kwon1, Galyna Krylova1†, Patrick J. Phillips2, Robert F. Klie2, Soma Chattopadhyay3,4, Tomohiro Shibata3,4, Emilio E. Bunel5, Yuzi Liu1, Vitali B. Prakapenka6, Byeongdu Lee7* and Elena V. Shevchenko1* To be able to control the functions of engineered multicomponent nanomaterials, a detailed understanding of heterogeneous nucleation at the nanoscale is essential. Here, by using in situ synchrotron X-ray scattering, we show that in the heterogeneous nucleation and growth of Au on Pt or Pt-alloy seeds the heteroepitaxial growth of the Au shell exerts high stress (∼2 GPa) on the seed by forming a core/shell structure in the early stage of the reaction. The development of lattice strain and subsequent strain relaxation, which we show using atomic-resolution transmission electron microscopy to occur through the slip of {111} layers, induces morphological changes from a core/shell to a dumbbell structure, and governs the nucleation and growth kinetics. We also propose a thermodynamic model for the nucleation and growth of dumbbell metallic heteronanostructures.

C

ombining multiple components within individual nanoparticles (NPs) is a simple way to control chemical and physical properties at the nanoscale to obtain efficient catalysts and advanced energy-conversion and storage systems1–9 . A broad range of heterostructures at the nanoscale that combine different constituents, such as metals, metal oxides and semiconductors, can be obtained through wet chemistry approaches as a result of the nucleation and growth of the overgrowth phase on the preformed seed NPs (refs 10–13). The seed-mediated growth method allows the synthesis of heteronanostructures in diverse shapes, including core/shells14–17 , nanodumbbells13,18–22 , nanorods23,24 and tetrapods25,26 . The synthesis of multicomponent heterostructured NPs is associated with the formation of heteroepitaxial structures, which makes the synthetic reaction more complicated compared with singlecomponent NP synthesis. As the seed and overgrowth phases have different crystal structures, heterostructured NPs can have lattice mismatch and strain27 . When the lattice mismatch is relatively small, the morphology of the NPs can be simply explained by the preferential growth of the overgrowth phase on the seed21–26,28 . However, as the lattice mismatch becomes larger, it starts to affect the morphology of the heterostructured NPs (ref. 29). Moreover, the lattice strain resulting from the mismatch changes the physicochemical properties of heterostructured NPs. Recently it was shown that, owing to lattice strain, the formation of a heteroepitaxial core/shell structure of II–VI semiconductors with a large lattice misfit can markedly change the bandgap structure30 . The growing importance of heterostructured NPs demands a detailed understanding of how the heteroepitaxial structure appears and evolves during synthesis;

this cannot be achieved by ex situ studies of static structures alone. Further progress in the synthesis of nanomaterials with precisely tailored compositions and functions for targeted applications requires a fundamental understanding of the processes involved in the synthesis of heteronanostructures. Here we report an in situ study of the nucleation and growth kinetics, and of the temporal changes in the crystal structure, of metal dumbbell NPs by using synchrotron small- and wide-angle X-ray scattering (SAXS/WAXS) techniques, which allow for the observation of the transient structural and volumetric changes of the NPs. We found that in the early stage of the reaction an intermediate core/shell heterostructure is formed. The lattice of the core/shell is deformed by a huge stress (∼2 GPa), which we attribute to the formation of a pseudomorphic Au shell. The strain energy originating from this deformation substantially increases the free energy of the core/shell NP and thus its chemical potential, hence affecting the kinetics of nucleation and growth. We found that the transition from a core/shell to a dumbbell structure occurs through strain relaxation of the pseudomorphic Au shell and results in the nucleation of a strain-free Au domain. Atomic-resolution scanning transmission electron microscopy (STEM) analysis revealed that dislocation takes place by the slip of {111} layers at the seed/Au interface, which relieves the strain. This mechanism explains the anisotropic growth of Au in a certain crystallographic direction of the seed.

In situ synchrotron X-ray scattering study In the seed-mediated growth of multicomponent NPs, the nucleation of the overgrowth phase is accompanied by only minimal

1 Nanoscience

and Technology Division, Argonne National Laboratory, 9700 South Cass Avenue, Argonne, Illinois 60439, USA, 2 Department of Physics, University of Illinois at Chicago, Chicago, Illinois 60607, USA, 3 CSRRI-IIT, MRCAT, Argonne National Laboratory, 9700 South Cass Avenue, Argonne, Illinois 60439, USA, 4 Physics Department, Advanced Materials Group, Illinois Institute of Technology, Chicago, Illinois 60616, USA, 5 Chemical Science and Engineering Division, Argonne National Laboratory, 9700 South Cass Avenue, Argonne, Illinois 60439, USA, 6 Center for Advanced Radiation Sources, University of Chicago, Argonne, Illinois 60439, USA, 7 Advanced Photon Source, Argonne National Laboratory, 9700 South Cass Avenue, Argonne, Illinois 60439, USA. †Present address: Center for Sustainable Energy at Notre Dame, University of Notre Dame, Indiana 46556, USA. *e-mail: [email protected]; [email protected] NATURE MATERIALS | ADVANCE ONLINE PUBLICATION | www.nature.com/naturematerials

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NATURE MATERIALS DOI: 10.1038/NMAT4115

a

c

b

20 nm

20 nm

CoPt3

e

d

20 nm

CoPt3/Au (f = 5.3%)

20 nm

Pt

FePt

f

Pt/Au (f = 3.7%)

20 nm

FePt/Au (f = 6.5%)

20 nm

Figure 1 | TEM images of the seeds and nanodumbbells. a–c, Seed NPs with a mean size of 6.2 nm (CoPt3 ), 4.8 nm (Pt) and 6.2 nm (FePt). d–f, Nanodumbbells synthesized by using the seeds shown in a–c. Lattice misfit f of the seed and Au is shown in brackets.

volume changes of the NPs, which imposes challenges in spatial and compositional resolution. To solve this problem, we carried out real-time SAXS and WAXS measurements simultaneously. This allowed us to detect the temporal evolution of the size and the crystal structure of the NPs at sub-ångström resolution. We studied the synthesis of Pt alloy/Au and Pt/Au nanodumbbells (Fig. 1) because the morphology and composition of Pt-based heteronanostructures are of particular interest in a broad range of catalytic applications31 . The in situ SAXS/WAXS study of nucleation and growth of CoPt3 /Au nanodumbbells was performed under ‘realistic’ conditions using the same reaction volume, stirring and heating procedures as those typically used to synthesize NPs (Supplementary Fig. 1). The reaction was initiated by the injection of a solution of 6.2 nm CoPt3 seeds into a hot solution of gold precursor at 95 ◦ C (refs 18–20). Small aliquots (330 µl) were pumped out of the reaction mixture (10 ml) at different reaction times and returned back to the solution after acquisition of SAXS and WAXS data through a quartz window. Only partial conversion of the seeds to dumbbells was observed under the conditions chosen for the in situ studies (Fig. 2a). The conditions to achieve 100% conversion are discussed below. The formation of the dumbbells is accompanied by: an increase in the intensities of the SAXS curves at q < 0.05 Å−1 (Fig. 2b); and the development of Au diffraction peaks in the WAXS patterns (Fig. 2c). We modelled the SAXS data assuming a mixture of dumbbells and seeds without a Au domain (‘un-nucleated seeds’). By fitting SAXS data (Supplementary Equation 1) we estimated the number and the size distribution of the seeds and the dumbbells with very high reliability (χ 2 > 0.99; Supplementary Fig. 2)32,33 . The fitting results showed that the total number concentration of NPs is kept constant at 0.20 µmol l−1 throughout the reaction time, indicating that the reaction is under ideal heterogeneous-nucleation conditions and no homogeneous nucleation of Au NPs takes place. As a result, we can estimate the heterogeneous-nucleation probability for the seeds by counting the fraction of dumbbells, fdb , defined as34,35 fdb = 2

Ndb Nseed + Ndb

where Nseed and Ndb are the numbers of un-nucleated seeds and dumbbells per unit volume, respectively (Fig. 2d). The growth rate of the NPs was calculated from the electron densities of the NPs and the solvent, ρP and ρsolv , respectively, and from the invariant Q, which is directly determined by the scattering intensity function I (q) (refs 32,33): 1 Q= 2 2π

Z∞

I (q)q2 dq

0

For a colloidal solution, Q/(ρp − ρsolv )2 is equal to the total volume fraction of the particles (Supplementary Equation 2). Therefore, the growth of the NPs can be directly monitored by plotting Q/(ρp − ρsolv )2 versus the reaction time (Fig. 2e). The temporal changes in the size of the seeds and the dumbbells from SAXS data are shown in Fig. 2f. The nucleation and growth kinetics obtained from SAXS data revealed that the formation of the dumbbells can be described by three distinct periods (Fig. 2d–f): pre-nucleation, nucleation and growth. The pre-nucleation period (0–20 min) is characterized by no dumbbell formation (fdb < 0.005). During the nucleation period, the Au domain nucleates on the seeds, as is evident from the steady increase of the dumbbell fraction fdb up to 52%. In the growth period (140–240 min), no further nucleation of dumbbells takes place, and the dumbbells nucleated grow from 97 Å up to 120 Å. The pre-nucleation period revealed a number of unexpected trends. The growth rate of the NPs (dQ/dt in Fig. 2e) is the highest in the pre-nucleation period even though the nucleation of the dumbbells does not take place yet. Nevertheless, in the prenucleation period the radius of the seeds increased by 3.2 Å (Fig. 2f). Moreover, the WAXS data demonstrated an increase in the intensity of the CoPt3 (111) peak by 13%, possibly indicating the growth of an epitaxial shell on the surface of the seeds (Fig. 2g). Also, the CoPt3 (111) peak position seemed to be shifted to lower q values, suggesting lattice expansion of the seeds (Fig. 2h). Similar changes were also observed for the CoPt3 (200) peak, indicating isotropic growth and expansion of the seeds (Supplementary Fig. 3). These

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ARTICLES

NATURE MATERIALS DOI: 10.1038/NMAT4115 a

d

Pre-nucleation

Nucleation

Growth

0.6 fdb

CoPt3 Au

10–2 240 120 30 10 0

min

CoPt3

10–1

I (10–5 cm–1 Å–1)

20

Au (111)

100

Dumbbell 80

Au domain

Seed

CoPt3 (111)

CoPt3 (200)

0

CoPt3(111) intensity 2.7

2.4 2.80

q (Å–1)

30 10

h

Au (200)

240 120 min

I(q) (10−4 cm−1)

15

5

2

g

q (Å–1)

10

dQ/dt

60

10–2

c

Invariant Q

120

10–1

10–5

6 4

f

Size (Å)

I(q) (cm–1)

100

10–4

8

Q (10−5/Δ ρ 2)

101

10–3

0.2 0.0

e

50 nm

b

Dumbbell fraction

0.4

240 min

CoPt3(111) position

2.78

0 2.4

2.6

2.8

3.0

3.2

3.4

0

40

80

q (Å–1)

120

160

200

240

Time (min)

Figure 2 | In situ SAXS/WAXS data from the synthesis reaction of CoPt3 /Au dumbbells. a, TEM images of CoPt3 seeds, and the NPs obtained after 4 h of reaction. b,c, In situ SAXS and WAXS patterns from the reaction solution. The time interval for each measurement is set as 30 s for the first 20 min, then 2.5 min until 1 h, and 5 min afterwards. d–f, Dumbbell fraction, fdb , invariant Q normalized by (1ρ)2 (1ρ = ρAu − ρsolv ) and its time derivative dQ/dt, and size distribution of the seeds and dumbbells as a function of time. All parameters are drawn from the fitting calculation of SAXS except for the Au-domain size, which is calculated using the Debye–Scherrer equation from the Au(111) peak in WAXS. g,h, Intensity and position of the CoPt3 (111) peak in WAXS as a function of time. The plots show time averages for every five data points. The error bars are representative standard deviations of the time-series data.

trends come to a sudden stop when the nucleation period starts at 20 min (Fig. 2).

Strain-induced shape transformation of core/shell The unusual behaviours during the pre-nucleation period suggest the formation of a CoPt3 /Au core/shell structure before the nucleation of the Au domain; this is analogous to pseudomorphic two-dimensional (2D) layer growth on a flat substrate in the Stranski–Krastanov model27,36,37 (Fig. 3a). According to this model, the overgrowth phase forms a 2D layer on the substrate if its thickness t is smaller than the critical thickness tc . Under such conditions, the overgrowth phase is strained so that its crystal lattice is coherent with the lattice of the substrate (pseudomorphism), which allows minimization of the lattice mismatch and the interface energy12,27,37 . However, as the layer thickness increases above tc , increasing strain energy in the overgrowth layer leads to the recovery of the original crystal structure by dislocation. As a result, a transition from the pseudomorphic 2D layer to a 3D island of the overgrowth phase takes place. At this transition, the interface energy increases owing to the distortion of the lattice at the interface.

Figure 3b illustrates the morphological evolution of the seed/Au NPs through the pre-nucleation, nucleation and growth periods on the basis of the Stranski–Krastanov model. The thickness of the Au shell in the pre-nucleation period is estimated to be 3.2 Å (Fig. 2f); that is, below the tc = 3.9 Å for Au on CoPt3 (ref. 37). The low interface energy between the pseudomorphic Au shell and the seed facilitates the fast deposition of Au atoms at the surface of CoPt3 seeds, which explains the high growth rate of NPs in the pre-nucleation period. As the growing Au shell adopts the same lattice structure as the CoPt3 seeds, the intensity of the CoPt3 diffraction peak increases. The strained lattice of the Au shell exerts mechanical stress on the seed, leading to its lattice expansion. As the lattice constant of Au (a = 4.079 Å) is larger than that of CoPt3 (a = 3.854 Å), the shell lattice is two-dimensionally compressed by the stress σshell (Fig. 3c, left). The restoring force of the lattice-contracted shell exerts the stress σseed on the seed in the radial direction. As a result, the lattice of the seed NPs expands until the restoring forces from the shell and the seed are at equilibrium30 (Supplementary Equations 2 and 3). As the shell thickness increases, the position of the CoPt3 diffraction peak

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NATURE MATERIALS DOI: 10.1038/NMAT4115

a 2D layer

3D island tc

b

θ c > 0°

θ c = 0°

Pre-nucleation

Nucleation

Growth

c

σinter t

σseed

σseed

2r

σshell

Seed

Pseudomorphic Au

Au

d

Au

Pt

Co

3 nm

Figure 3 | Core/shell to dumbbell transition of the seed/Au NP during synthesis. a, Lattice structure of the overgrowth phase in the 2D and the 3D mode of the Stranski–Krastanov model. b, Morphological evolution of the seed/Au heterostructure during the pre-nucleation, nucleation and growth periods. The contact angle θc between the seed and the Au is indicated for the core/shell and the dumbbell. c, Schematic illustration of the stresses, σshell , σseed and σinter , within core/shell and dumbbell NPs. d, STEM-EDX mapping of a CoPt3 /Au dumbbell. Elemental mapping was performed by acquiring EDX spectrum images including the Au M series (red), Pt M series (green) and Co L series (blue) X-rays. The rightmost one is the layered image of those mappings. The electron probe size is approximately 100 pm, and the acquisition time was 287 s.

is shifted further owing to lattice expansion. The stress on the seed calculated from the lattice expansion is 2.4 GPa at the end of the pre-nucleation period16 (Supplementary Equation 3). The lattice expansion of the seed NPs is also observed during the growth period (Fig. 3c, right). This is most likely a result of the extensional stress σint on the seed caused by the surface tension of Au and thus the reduction of the negative curvature of the Au domain at the seed/Au interface38 . Extended X-ray absorption fine structure (EXAFS) data show neither Au–Pt nor Au–Co alloy phases (Supplementary Fig. 4 and Table 1). Note that alloying of CoPt3 , FePt and Pt with Au is very unlikely at the reaction temperature of 95 ◦ C because of their high enthalpies of mixing39–41 . Thus, we can exclude the possibility of seed lattice expansion by alloying. The presence of a Au shell on the seeds is further evidenced by energy-dispersive X-ray spectroscopy (EDX)-STEM mapping (Fig. 3d). Note that the thickness of the Au shell estimated from 4

the SAXS data (3.2 Å) is comparable to the spatial resolution of this image. Catalytic activity tests confirm the presence of a continuous Au layer on the surface of CoPt3 seeds. The surface of CoPt3 NPs catalyses the hydrogenation reaction of unsaturated hydrocarbon molecules with 100% conversion42 . On the contrary, using the CoPt3 /Au NPs as a catalyst, the conversion yield dropped down to 0%, confirming that the whole surface of the CoPt3 seeds is covered with a catalytically inactive Au layer (Supplementary Fig. 5). Together, the results from EDX, catalytic activity tests, and WAXS data (isotropic growth and strain of the seeds in the pre-nucleation period) indicate that a continuous pseudomorphic Au shell is formed on the CoPt3 seeds. The contribution of lattice strain to the reaction kinetics is considered by calculating the strain energy and the chemical potential of the NPs. Assuming isotropic elasticity and a very thin shell (t  r), the strain energy Us of a core/shell NP with a spherical

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NATURE MATERIALS DOI: 10.1038/NMAT4115 a

µ

the 3D mode in the Stranski–Krastanov model. In this transition, the interface energy between the seed and the Au is increased by the distortion of the lattice (Fig. 3a, right). Thus, the morphological change from core/shell to dumbbell takes place to reduce the contact area between the seed and the Au domain by changing the contact angle from θc = 0 to θc > 0 (Fig. 3b).

µ soln,1

Nucleation

µc µsoln,2

Chemical potential and nucleation kinetics

nc

b

nAu

0.6

[AuCl]o = 10.5 mM

Dumbbell fraction

0.5 0.4 0.3

3.5 mM

0.2 0.1

1.2 mM

0.0 0

40

80

120

160

200

240

Time (min)

Figure 4 | Heterogeneous nucleation kinetics of the dumbbells. a, Chemical potential of the core/shell (blue) and dumbbell (green) NPs as functions of the amount of Au deposited on the NPs, nAu . See the main text for details. b, Temporal change in dumbbell fraction fdb during synthesis using solutions containing AuCl of various concentrations. The plot for [AuCl] = 10.5 mM is identical to that in Fig. 2d.

seed of radius r and a pseudomorphic shell of thickness t (Fig. 3c, left) is written as: Us = Us,seed + Us,shell Eseed Eshell = · (εseed )2 · 2π r 3 + · (εshell )2 · 4π r 2 · t (1 − 2νseed ) (1 − νshell )

(1)

where r is the radius of the seed NP, and Ei , νi and εi are Young’s modulus, Poisson’s ratio and lattice strain, respectively (see Supplementary Equations 5–7 for the derivation of equation (1)). According to SAXS/WAXS data, at the end of the pre-nucleation period the lattice of the CoPt3 seed is expanded by an εseed of 0.38% and the Au shell lattice is contracted by an εshell of 4.2% (Fig. 2g and Supplementary Fig. 3). Applying the bulk moduli of CoPt3 and Au, the strain energy Us for a CoP3 /Au core/shell structure with a seed diameter of 62 Å and a shell thickness of 3.2 Å is calculated to be 70 eV. This contribution of the strain energy to the free energy of the NP leads to an increase in the chemical potential of the core/shell NP, µNP , of 31 meV (≈Us /nAu , Supplementary Equation 8). As the chemical potential of Au in the solution (µsoln ) is approximately 36 meV higher than that of the seeds (Supplementary Equation 9 and Fig. 6), the increase of µNP by 31 meV can substantially suppress the deposition of Au from the solution onto the core/shell. This explains the rapid decrease in the growth rate of the NPs during the pre-nucleation period (Fig. 2e). When the thickness of the Au shell exceeds the critical value (tc = 3.9 Å), nucleation of the dumbbells may occur forming a strainfree Au domain; this is analogous to the transition from the 2D to

The effects of the chemical potential and the morphological changes on the nucleation and growth kinetics of the dumbbells can be explained by a simple thermodynamic model, illustrated in Fig. 4a. In classical heterogeneous nucleation theory, the contribution of heteroepitaxy to the free energy of the particles is not taken into account35,43 (Supplementary Equation 10 and Fig. 7). However, in our study, the strain energy Us can substantially increase the free energy of the core/shell NP and its chemical potential. In the pre-nucleation period, the chemical potential of the NP, µNP , increases sharply with the amount of Au deposited on the seed, nAu , owing to the strain energy Us (Supplementary Equation 8). For a dumbbell with a strain-free Au domain, the value of µNP is mainly determined by the size of the Au domain following the Gibbs–Thompson relation (1µ = 2γ Vm /r, where γ , Vm and 1/r are the surface free energy, molar volume, and the surface curvature of the NP, respectively)44 . In other words, dumbbells with a larger Au domain have a lower chemical potential. Thus, µNP of the dumbbell (green plot) decreases with nAu ; however, µNP for the core/shell increases (blue plot). As a result these two plots intersect at µc = µNP (nc ). In the pre-nucleation period, the core/shell structure is formed from the seed. The growth of the shell is driven by the difference between µNP and µsoln . As µsoln = kT ln S, where S is the supersaturation, the probability for the core/shell to pass through the peak point at µc and nucleate the Au domain is critically dependent on the supersaturation level. If the supersaturation is lowered from µsoln,1 to µsoln,2 , the height of µc relative to µsoln,2 acts as the energy barrier to block the nucleation of the dumbbells. In control experiments, the dumbbell fraction fdb decreases from 52 to 3% as the concentration of the Au precursor is lowered by 9 times (Fig. 4b). On the other hand, when the high supersaturation level was maintained by adding the Au precursor continuously throughout the reaction time using a syringe pump, the value of fdb was increased up to ∼100% (Supplementary Fig. 8)18–20 . Our model predicts that if the core/shell is free from lattice strain it may grow without the nucleation of dumbbells. The synthesis using Pd seed NPs supported our model. Unlike CoPt3 , Pt and FePt, Pd is known to be perfectly miscible with Au, and intermixing at their interface effectively alleviates misfit strain39–41 . Indeed, Pd/Au core/shell NPs were synthesized under the same reaction conditions as the CoPt3 /Au dumbbell. Note that the lattice misfit of Pd/Au is 4.6%, which is between the values of 5.3% for CoPt3 /Au and 3.7% for Pt/Au (Supplementary Fig. 9). Unrestricted Au shell growth on Pd seeds is in agreement with the Frank–van der Merwe model developed for thin films12,27 .

Dislocation mechanism in the heteronanostructure To study the heteroepitaxial structure of the dumbbells we conducted atomic-resolution STEM analysis (Fig. 5). In the fast Fourier transform (FFT) image in Fig. 5b, the diffraction spots from the seed (CoPt3 ) and Au domains are well resolved for {220} and the higher index planes (see Supplementary Fig. 10 for the larger size FFT image). We visualize CoPt3 and Au domains in the dumbbell by applying inverse FFT for CoPt3 and Au {220} spots, respectively45,46 (Fig. 5c). Notably, the boundary between these domains is roughly ¯ and (111) ¯ parallel to (111) planes. The distortion of the lattice is shown in Fig. 5d,e. It is clearly seen that the {111} lines are bent from the centre of the dumbbell to either the left-hand (Fig. 5d) or the right-hand side (Fig. 5e) with respect to the seed lattice (yellow lines)

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NATURE MATERIALS DOI: 10.1038/NMAT4115

a

b

b

c − (002) (111) − (111) Au(220)

− ) (111

− 1) (11

− (220)

CoPt3(220) 4 nm

e

d

(001

(001 )

)

−11) (1

− ) (111

f

g 5

1

10 Au 7 ML

Intensity (a.u.)

7 6

6 ML 5 ML 4 ML 3 ML 2 ML

− 1) (11

5 4

1 ML 5

1

3 2

0

1

5

10

10 Seed 15

20

25

Position (Å)

Figure 5 | STEM analysis on the seed/Au interface of the dumbbell. a, Atomic-resolution annular bright-field STEM image of a CoPt3 /Au dumbbell. The spatial resolution of this probe is approximately 73 pm. b, FFT image of a (white) overlapped with those from the seed (blue) and Au (red) domains. The corresponding domain areas are indicated in a. c, The seed (blue) and Au (red) domains in the dumbbell obtained by inverse FFT of CoPt3 and Au(220) spots in b. The area where two domains are superimposed is coloured in green. d,e, Magnified images of the area indicated in the insets. As a guide to the eyes (¯111) and (1¯11) lattice planes are indicated with white lines in d and e, respectively. Yellow straight lines are extensions of the seed lattice. The direction of the displacement of the lattice planes (white) with respect to the seed lattice (yellow) is indicated with arrows. f,g, Magnified image showing 7 adjacent monolayers over the (1¯11) interface (f) and intensity profiles along the dashed lines in the image (g).

as they pass through the interface. The displacement directions (arrows) of the bent lattice lines are parallel to the {111} and (001) planes of the seed. Displacement in atomic positions at the interface was further confirmed by measuring the interatomic distances in 7 adjacent monolayers (ML) over the interface (Fig. 5f,g). The length of 10 atoms is 24.2 Å for 1, 2, and 3 ML, and 25.8 Å for 6 and 6

7 ML (Fig. 5g). These values are close to 23.6 Å for CoPt3 and 25.0 Å for√Au, respectively, calculated from their lattice constants as 10×a 6/4 (ref. 47). The transition from CoPt3 to the Au lattice structure occurs in two monolayers (4 and 5 ML). This is in agreement with the thickness of the Au shell (∼3.2 Å) estimated from SAXS data, which is about 2 ML thick. Note that the dumbbell

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NATURE MATERIALS DOI: 10.1038/NMAT4115 a

(001)

(11− 1)

−111) (

See See eed e (CoPt (C Co CoPt3)

b

2.2 3

Å

d=

2. 3



Gold G Go old

c

d=

Seed Se eed (C (CoPt (CoP Co oPt3)

d 〈001〉 (φ = 0°)

φ PAu 〈001〉

PAu 〈111〉

φ

〈111〉 ( φ = 54.7°)

〈001〉

002 111 111 110 110

〈111〉

5 nm

e 〈001〉

f

φ

φ 〈111〉

Frequency (a.u.)

〈111〉

Frequency (a.u.)

φ Frequency (a.u.)

g 〈001〉

〈001〉

Pt/Au

CoPt3/Au

〈111〉

FePt/Au

Figure 6 | Lattice structure at the seed/Au interface, and distribution of Au domain positions for CoPt3 /Au, Pt/Au and FePt/Au dumbbells. a,b, 2D projection of lattice structures of core/shell (a) and dumbbell (b) along the h110i zone axis. The ratio of CoPt3 and Au{111} d-spacings is the same as the actual value (dCoPt3 :dAu = 100:105). Yellow straight lines are extensions of the seed lattice, the same as in Fig. 5. c, TEM image of a Pt/Au dumbbell as an example of measuring the position angle of the Au domain (φ = 0.6◦ ) with respect to the crystallographic axis. d, Schematic for the definition of position vector PAu and its angle φ. PAu is from the centre of mass of the seed to that of the Au domain. e–g, Histograms in polar coordinates showing the distribution of Au domain position angles for CoPt3 /Au, Pt/Au and FePt/Au dumbbells. The interval is 5◦ .

is free from the crystal defect even though it has the distorted lattice structure (lattice line bending in Fig. 5). In heteroepitaxial thin films dislocation is accompanied by distortion of the lattice as well as the formation of the line defect (misfit dislocation). However, 6.2 nm CoPt3 seeds are smaller that the misfit dislocation period Pd of 7 nm for CoPt3 /Au. (Supplementary Equation 11)37 . As a result, the dislocation in the dumbbells gives rise only to distortions but no line defect. The heteroepitaxial lattice structure at the interface revealed by STEM analysis is depicted in Fig. 6a,b. Unlike the flat thin film, the pseudomorphic shell at the surface of the spherical seed NP can select the lattice plane for dislocation to minimize the frictional force between the atomic planes. The {111} layers of face-centred cubic metals have the smallest interlayer friction (Peierls stress) among

the other lattice planes (Supplementary Equation 12)47 . Figure 6a,b illustrates the dislocation mechanism of the NP. On the expansion of the pseudomorphic Au lattice, a pair of adjacent Au {111} layers, ¯ and (111), ¯ for example, (111) slip in opposite directions, leading to the formation of a strain-free Au nucleus centred on the (001) plane of the seed. This mechanism resembles a threading dislocation in a heteroepitaxial thin film in the sense that the (001) layer expansion occurs by the slip of the (111) layer36 . The growth directions of the Au domain were investigated by measuring the position of the Au domain (centre of mass) with respect to the crystallographic direction of the seeds (Fig. 6c,d). Most of the Au domains are in the h001i direction of the seed lattice, with only a few examples showing the h111i orientation (Fig. 6e–g). These results confirm that the strain relaxation mechanism based on friction-force minimization

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NATURE MATERIALS DOI: 10.1038/NMAT4115

(Fig. 6b) is energetically most favoured during the nucleation of the Au domain (see Supplementary Figs 11–14 for TEM analysis of various dumbbells). Our study reveals that in the synthesis of multicomponent NPs the evolution of the heteroepitaxial structure induces mechanical deformation of the seeds and the overgrowth phase, which in turn affects the heterogeneous nucleation and growth reaction kinetics. In the synthesis of Pt alloy/Au and Pt/Au dumbbells, intermediate core/shell heterostructures with a coherent lattice structure are formed first at the expense of a lattice stress of 2.4 GPa on the seed. The relief of the lattice strain in the shell takes place through the slip of {111} Au atomic layers at the seed/Au interface. This leads to breaking of the lattice coherency and to the nucleation of a strain-free Au domain at the (001) plane, similar to the 2D-to-3D transition in the Stranski–Krastanov model. We have shown that mechanical stress affects the kinetics of the chemical reaction by changing the chemical potential of the heterostructured NPs. If the supersaturation of the Au precursor is lower than the chemical potential of the core/shell, the strain energy acts as an energy barrier, blocking the nucleation of the dumbbells and limiting their yield. The knowledge gained in this study should allow the development of a time-resolved synthesis mechanism that enables the control of structure–property functions in engineered multicomponent nanostructures.

imaging. Catalytic hydrogenation reactions using the seeds and the nanodumbbells as catalysts were performed and analysed following the same procedure described previously49 except for the use of citral (Aldrich, 95%) instead of 4-octyne. A Perkin Elmer Optima 3300DV inductively coupled plasma optical emission spectrometer was used to measure the concentration of the NPs.

Methods

References

Synthesis of seed NPs and nanodumbbells. CoPt3 , Pt, FePt and Pd seed NPs were synthesized following previously reported methods19,20,48 . For the synthesis of nanodumbbells, a Au solution was prepared by dissolving 24 mg AuCl (Aldrich, 99.9%), 153 mg didodecyldimethylammonium bromide (Fluka, 98%) and 270 mg 1-hexadecylamine (90%, Aldrich) in 10 ml toluene. The seed solution contained ∼3 mg of the seed NPs in 1 ml toluene. After injecting the seed solution into the Au solution at 90 ◦ C, the solution was heated at the same temperature for 4 h under a nitrogen atmosphere, and then excess acetone was added to separate the NPs.

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Synchrotron X-ray measurement. In situ SAXS/WAXS measurement was carried out at beamline 12-ID-B of the Advanced Photon Source (APS), Argonne National Laboratory. The wavelength λ of the X-ray beam was 1.03 Å (12 keV), the beam size was 0.6 × 0.05 mm2 , and the flux was ∼ 1012 s−1 . The beam exposure time was set to 2 s for each measurement. The sample-to-detector distance was 2,215.54 mm for SAXS and 455.26 mm (18.0◦ tilted towards sample) for WAXS. Detectors used for SAXS and WAXS were the Pilatus 2M and Pilatus 300k, respectively. The signal from the detectors was azimuthally averaged. The q(= 4π sin θ/λ) and absolute intensity calibrations are performed using silver behenate and a solution containing CoPt3 NPs of known concentration as standards, respectively. A Hamilton Microlab 600 syringe pump was used to inject the seed solution and sampling aliquots. For each measurement 330 µl of the sample aliquot was drawn from and pushed back to the solution at a rate of 40 µl s−1 through 18-gauge polymer tubing. In the middle of the tubing, a ∼2-cm-long quartz capillary with a diameter of 2.0 mm and a wall thickness of 0.01 mm was placed as a window for the X-ray beam. The syringe pump, beam exposure and data collection were controlled by custom-built codes run on a Linux-operated computer. Additional X-ray diffraction measurements were carried out at GSECARS Sector 13-ID-D (photon energy of 37 keV, λ = 0.3344 Å) of the APS. EXAFS measurements for Co K, Pt L3 and Au L3 edges of CoPt3 /Au dumbbells were carried out at MR-CAT Sector 10-ID-B of the APS. The solutions containing the dumbbells in toluene were loaded in cylindrical plastic cuvettes for measurements. Experiments were performed in transmission mode for the Pt L3 edge (11,564 eV) and in fluorescence mode for the Au L3 (11,919 eV) and Co K (7,709 eV) edges using ion chambers with Stern–Heald geometry. Characterization. Conventional TEM analysis of the NPs was performed with a JEM-2100F operated at 200 kV. STEM and EDX analysis were performed using an aberration-corrected JEOL JEM-200CF equipped with an Oxford X-MaxN 100TLE silicon drift X-ray detector and operated at 200 kV. The microscope is equipped with a cold field-emission source, which yields an energy resolution of 0.35 eV; it allows for 73 pm spatial resolution at 200 kV with the probe spherical-aberration corrector. To achieve a high-brightness probe for EDX mapping, we used a convergence semi-angle of 28 mrad at 200 kV primary energy, a 110–440 mrad collection semi-angle for high-angle annular dark-field imaging, and a 14–28 mrad collection semi-angle for annular bright field 8

Data analysis. In situ SAXS data were analysed by custom-built software that is run on Matlab after standard data correction including background correction. Before the seed injection, the scattering signal from the Au solution in the quartz capillary was measured first and used as the background. In situ WAXS data were background-corrected in the same way as SAXS data, and then noise-filtered by the moving-average method. The area and the position of each diffraction peak were calculated by Gaussian fitting (Supplementary Fig. 15). χ -square values (χ 2 ) of fitting were 0.94 for {111} peaks and 0.86 for {200} on average. The validity of applying the WAXS data analysis method is justified by simulating the coherent and incoherent diffraction from their atomic models (Supplementary Fig. 16). Analysis of TEM and STEM images was carried out using a Gatan DigitalMicrograph. Statistics on the size and the growth direction of the nanodumbbells were estimated from TEM images of 800k and 1,000k magnification. The number of nanodumbbells measured was 50 for CoPt3 /Au, 75 for Pt/Au, and 70 for FePt/Au. EDX data were analysed using Oxford AZtecTEM microanalysis software. The EDX maps shown in Fig. 3d are represented by pixels binned ×2, but are otherwise unprocessed. The apparent discontinuity of the Au shell in the core/shell structure is an artefact of the binning and the image contrast chosen in Fig. 3d.

Received 14 January 2014; accepted 19 September 2014; published online 2 November 2014

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Acknowledgements Use of the Center for Nanoscale Materials and Advanced Photon Source was supported by the US Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC02-06CH11357. MRCAT is funded by MRCAT host institutions. The authors thank C. Marshall for fruitful discussion and help in tests of catalytic activity. S.C. and T.S. would like to thank V. Zyryanov for help with experiments and C. Segre for beam-time allocation. P.J.P. and R.F.K. acknowledge support from the National Science Foundation (DMR-0959470) for the acquisition of the UIC JEOL JEMARM200CF. Support from the UIC Research Resources Center is also acknowledged. The work at GeoSoilEnviroCARS was supported by the National Science Foundation—Earth Sciences (EAR-0622171) and Department of Energy—Geosciences (DE-FG02-94ER14466).

Author contributions S.G.K. designed and implemented experiments, and analysed the data except for SAXS and EXAFS. G.K. carried out NP synthesis and sample preparations. P.J.P. and R.F.K. performed STEM and EDX measurements. S.C. and T.S. measured EXAFS and analysed the data. E.E.B. supervised catalytic reactions and analysis. Y.L. supervised TEM analysis. V.B.P. performed synchrotron X-ray diffraction measurements. B.L. supervised and performed SAXS/WAXS measurements and analysed SAXS data. E.V.S. supervised the project and wrote the manuscript with B.L. and S.G.K.

Additional information Supplementary information is available in the online version of the paper. Reprints and permissions information is available online at www.nature.com/reprints. Correspondence and requests for materials should be addressed to B.L. or E.V.S.

Competing financial interests The authors declare no competing financial interests.

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Heterogeneous nucleation and shape transformation of multicomponent metallic nanostructures.

To be able to control the functions of engineered multicomponent nanomaterials, a detailed understanding of heterogeneous nucleation at the nanoscale ...
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