How do atoms rearrange in a solid-state reaction to form new structures, especially in catalytic systems with dramatic property difference? Understanding atomic dynamics is key to achieving controlled synthesis of catalysts and other materials. Using advanced electron microscopy, we directly monitor the rearrangement of atoms during a reaction between nanocrystalline Pt and SnO 2 and observe consecutive formation of Pt 3 Sn and PtSn. This real-time imaging provides invaluable understanding of the alloying mechanisms in intermetallic nanoscale systems, which enables precise phase selection and thereby control of the catalytic properties, as we demonstrate for the semi-hydrogenation of acetylene. In pursuit of fully controlled synthesis of intermetallic nanoparticles, our results have made an important step forward in the structure manipulation to serve society with more efficient and affordable energy and chemistry.

Intermetallic nanoparticles (iNPs) have yielded enormous successes in catalytic applications by the formation of ordered phases. However, atomic-level understanding of the alloying mechanism, which plays a pivotal role in controlling intermetallic phases and tailoring their catalytic properties, is still elusive. In this study, we discovered a consecutive formation of ordered Pt 3 Sn and PtSn phases during the growth of Pt-Sn iNP inside a well-defined nano-reactor at elevated temperature by using in situ scanning transmission electron microscopy. We found that the surface-mediated diffusion of Sn controls overall dynamics of the reaction, while the unique coherent interfacial structure is determinative for the PtSn transformation. We then further controlled the phase selection of Pt-Sn iNPs and demonstrated their distinguishable catalytic behaviors. Our findings not only provide detailed experimental evidence on the alloying mechanism in intermetallic nanoscale systems but also pave the way for mechanistic control of synthesis and catalytic properties of iNPs.

This study aimed to unravel these competing effects by directly monitoring the formation of ordered phases in iNPs at the atomic level. This was achieved with a silica-confined nano-reactor, in which single reactant pairs of Pt/SnOparticles were encapsulated by mesoporous silica shells (Pt/SnO@mSiO). This elaborate structure was synthesized by polymerizing the silica layer around the surface of nanoparticles according to a sol-gel process (see Experimental Procedures for details). The nano-reactor can isolate the reactant pairs up to 750°C without sintering and aggregation,providing a stable platform by which we can study chemical diffusion and phase transformations in the reactant pairs at elevated temperature. High-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM), which can simultaneously provide atomic spatial resolution and chemical sensitivity (Z contrast),was employed to observe the movement of Pt and Sn atoms during the reaction. In addition, the formation of ordered Pt-Sn iNPs requires the reduction of SnO, which can be facilitated by the electron beam (as the reducing agent). Therefore, we were able to precisely control and monitor the transformation process using a sub-angstrom electron probe. We discovered that the reaction pathway between Pt and SnOfollows a consecutive formation of ordered PtSn and PtSn. A unique coherent interface,and, was observed during the PtSn→PtSn transformation. Combining in situ atomic-scale observations and theoretical calculations, we explored the effects of the surface diffusion and the interfacial structure. Furthermore, we demonstrated that the phase formation and morphology of Pt-Sn iNPs could be predicted and controlled according to our findings. These iNPs showed dramatically altered catalytic properties in the semi-hydrogenation of acetylene, emphasizing the importance of understanding the alloying mechanism of iNPs.

Intermetallic structures tend to minimize their free energy by forming atomic configurations commensurate with strong chemical bonds. As a result, different elements can alternate in specific sites and form ordered phases. Such ordering transformations are of great interest because these structures can lead to distinguishable and often controllable electronic,magnetic,and chemicalproperties, often predictable from thermodynamics and kinetics considerations.For bulk materials, phase diagrams that describe the alloying behaviors of metals have been established in many systems, providing key guidance for the alloy design and development.However, at the nanoscale, the limited dimensions of nanoparticles raise a quantum-size effect that can strongly modify the thermodynamics and kinetic stability of many intermetallic compounds.Indeed, different alloying behaviors are evident between bulk and iNPs. For example, in the bulk Pt-Sn system, obtaining the ordered PtSn and PtSn requires melting of pure Pt and Sn in the appropriate ratios and then high-temperature treatments such as zone refining.In contrast, these phases can be easily formed in the iNPs: PtSn is obtained, e.g., by annealing the Pt-Sn nanoparticles at 300°C to 600°C; PtSn forms at temperatures even below 300°C.The increased surface area, which facilitates mass transportation via surface diffusion, is believed to reduce the energy barrier of the ordering transformations in iNPs.On the other hand, when ordering occurs, the introduction of phase boundaries increases the system energy, opposing the formation of ordered phases.These two competing effects are generally active in nanoscale systems and can be influenced by the size, shape, and interfacial energies of iNPs. Such complicated pathways deviate iNPs from the well-established bulk phase diagrams, making it more difficult to design and fabricate multifunctional iNPs at the nanoscale.

Intermetallic compounds have drawn great attention for catalytic applications because of their dramatically improved activity, selectivity, and stability.Design and synthesis of intermetallic nanoparticles (iNPs), bearing a high surface-to-volume ratio, are thus highly desired for an increased number of active sites. Being a solid-state mixture of two or more metal components, iNPs possess not only modified crystal structures to accommodate different elements but also altered surface geometric and electronic structures that could directly correlate with catalytic activity and selectivity,given that the adsorption of reactive species is highly dependent on the structure of the active sites.

To account for the situation in the nanoparticles, we extracted a 5 nm sphere-shaped particle from the defect-free region in Figure 5 A. A section view of the resulting atomic arrangement is presented in Figure 5 C, in which all atoms are colored by the same cohesive energy scale used in Figures 5 A and 5B. Note that the atoms on the surface of the particle had slightly higher cohesive energy than the internal ones, indicating that the elastic strain induced by the lattice misfit was released to the particle surface. However, the misfit cannot be eliminated entirely at the interface, as indicated by the dashed lines in Figure 5 C, where the Pt frames were obviously distorted. During the transformation from PtSn to PtSn, the Pt atoms were substituted by Sn in the PtSn phase. The excess energy induced by the substitution process can be assessed by combining the atomic stress field and the elasticity theory.Because Sn has a larger atomic radius than Pt, the substitution of Pt by Sn induces a dilatational strain at the original Pt site, which can be represented by a dilatational volume, V(Åper Sn atom), with a positive value. The excess energy (binding energy) for the substitution can be expressed as E= −V, whereis the local stress field on the Pt site.Apparently, a tensile stress field that gives a positive value ofwill result in a negative binding energy and vice versa. It can be seen in the calculated stress tensors ( Figures 5 D–5F) that the Pt atoms near the interface-affected region of PtSn were always subjected to relatively high tensile stress, indicating that these atoms are energetically preferable for the substitution by Sn. The result reveals that the unique interfacial structure of PtSn/PtSn induces tensile stress on the Pt atoms at the interface and serves as the driving force for the PtSn→PtSn transformation.

To investigate the mechanism of the formation of ordered phases in the Pt/SnO@mSiO, we calculated the interfacial energy of PtSn/PtSn by using molecular dynamics (MD) simulations. Two types of interfaces were investigated: (1)and, which matches the STEM observation and is hereafter referred to as the type 1 interface; and (2)and, the type 2 interface, which obeys the conventional interfacial relation between face-centered cubic and hexagonal close-packed crystals.Note that the type 2 interface gave a lattice misfit of 17.6%, 7.3 times larger than that of the type 1 interface. The atomic configurations of type 1 and type 2 interfaces are shown in Figures 5 A and 5B , respectively, and the atoms are colored by their cohesive energy (enthalpy per atom). The simulated model for the type 1 interface ( Figure 5 A) yielded an interfacial structure that was exactly the same as the in situ STEM results, except for some defected regions caused by the lattice misfit (marked by the dashed box). The interfacial energy for this configuration was as low as 0.459 J/m. The Pt atoms at the interface were all in their low-energy state, and the strongest cohesive energy was calculated to be −6.53 eV. For the type 2 interface ( Figure 5 B), the calculated interfacial energy was 2.47 J/m, 5.4 times larger than that of type 1. In addition, the distribution of the cohesive energy of atoms appeared to be random. The strongest cohesive energy of Pt atoms at the interface was −5.91 eV, consistent with its high interfacial energy and large lattice misfit. The result indicates that the type 1 interface is energetically more favorable than the type 2 interface.

The final stage of the process involved the coalescence of PtSn grains. This process followed the classical sintering theory, in which the larger grains that have lower surface energy per unit volume tend to increase their size by consuming the smaller and higher surface energy grains.As shown in Figures 4 A–4D , the growth of the lower left and lower right grains toward the upper ones was first seen from 406.5 to 423 s, and the growth directions were indicated by arrows. Grain boundaries between them migrated and finally converged at the surface. This process was very fast—finished within only 16.5 s—resulting in a two-grain particle. Then, these two grains merged and transformed into a single grain (see Figure 4 E). Note that from 423 s, the notch at the lower part of the particle started to be filled, along with the dissolution of a smaller particle in the vicinity ( Figure 4 D). This phenomenon is the typical Ostwald ripening that is often observed in the coalescence of nanoparticles.A HAADF-STEM image of the sample ( Figure 4 F) was captured after it was cooled to room temperature. FFT in the inset reveals that the particle contained a single PtSn grain, which slightly rotated from the original [001] orientation during cooling. EDS mapping on the particle after the reaction ( Figure S3 ) shows that Pt and Sn were uniformly distributed throughout the particle, indicating that the starting sample, Pt/SnO, was successfully converted to single crystalline PtSn.

In addition, we noticed that a different PtSn grain formed in the top part of the particle at 148.5 s ( Figure 3 H). The lattice fringes were measured as ∼2.15 Å, matching those of PtSn (102). The formation of this PtSn grain was also accompanied by the consumption of the PtSn phase. No orientation preference of this nucleation relative to the Pt particle was identified. This PtSn grain was hardly grown afterward. A similar phenomenon was observed at 283.5 s in the bottom part of the particle, as shown in Figure 3 N. The particle at the time was composed of four distinct PtSn grains and one unreacted Pt/PtSn grain. This tentative equilibrium was broken as annealing continued. At 381 s, the four PtSn grains started to grow toward the center of the particle, and the PtSn fringes disappeared ( Figure 3 O). Finally, the reaction ended up with four PtSn grains, between which boundaries were built up ( Figure 3 P).

Using the aforementioned Fourier filters, we were able to track the evolution of PtSn. To do that, we superimposed the IFFT images from the PtSn (001) superlattice spots on the HAADF-STEM images with yellow color, as shown in Figure 3 . The FFTs and the original IFFT images are presented in Figure S2 . The PtSn phase mainly existed near the surface of the particle in the early stages ( Figure 3 A) and propagated gradually to the center of the particle as a result of Sn diffusion. The PtSn nucleus started to grow into the PtSn domain, leading to the migration of theinterface, seen at 69 s in Figure 3 B. In contrast to the continuous growth of PtSn, the growth of PtSn was discontinuous, which included a reshuffle of atoms within a nanometer region. For example, at 79.5 s, the lattice fringes of PtSn in front of the PtSn nucleus suddenly disappeared, leaving a disordered region ( Figure 3 C, as marked by the dashed line). This region showed a darker contrast relative to the Pt/PtSn domains, suggesting a high concentration of Sn in the area. Pt columns with the hexagonal periodicity evolved from this Sn-rich region, as observed in the following frame ( Figure 3 D). This atomic-level rearrangement completed at 82.5 s, resulting in an ordered PtSn region on the right side of the particle ( Figure 3 E). When the PtSn phase was depleted in front of the PtSn/PtSn interface, the progression of the PtSn domain arrested, possibly as a result of insufficient Sn supply. As can be seen in Figures 3 E–3H, the PtSn and PtSn domains in the sample were nearly static, except for a slight growth of the other PtSn nucleus near the left edge (see Figure 3 G). Assuming the reaction was limited by the supply of Sn, we increased the beam current from 30 to 50 pA at 192 s. Confirming our hypothesis that the electron beam is critical to the reduction of SnO, the formation of PtSn commenced ( Figure 3 I). At 195 s, the region in front of the interface was dominated by PtSn ( Figure 3 J). Similar to the previous growth sequence of PtSn out of PtSn + Sn ( Figures 3 C–3E), the disappearance of PtSn lattice was again observed at 207 s ( Figure 3 K), followed by the appearance of a disordered Sn-rich area (encircled by the dashed line). By 244.5 s, PtSn lattice emerged from the Sn-rich area ( Figure 3 L), resulting in the growth of the PtSn domain and the migration of the interface at 246 s ( Figure 3 M).

Nucleation of the PtSn phase on the right corner of the Pt particle was observed after ∼18 s, as highlighted by the yellow box in Figure 2 D. The distance between the bright atomic columns with 6-fold symmetry, expected for PtSn, was measured as ∼4.1 Å, matching the Pt-Pt spacing of hexagonal PtSn along [001] (see the atomic model in Figure 2 K and a simulated HAADF image along [001] in Figure 2 M). The PtSn nucleus and the parent PtSn phase follow a specific orientation relationship:and. This configuration can be ascribed to the similar d-spacings of PtSn (002) (2.002 Å) and PtSn (110) (2.051 Å), giving a lattice misfit of only 2.4%. Another PtSn nucleus formed on the opposite corner was seen at 27 s, as shown in Figure 2 G. Note that the nucleus formed with the same orientation preference described previously.

The reaction between Pt and SnOwas successfully triggered by the electron beam at elevated temperature, yielding the formation of ordered PtSn and PtSn phases. Figure 2 shows a few selected frames at the early stage of the reaction. A shell with relatively dark contrast started to form around the Pt particle right after being exposed to the electron beam, as shown in Figures 2 A, 2D, and 2G. The darker contrast of the shell compared with the Pt core indicates that the shell was rich in Sn. We attribute the appearance of this Sn-rich shell to the SnOreduction initiated by the electron beam at 350°C. Interestingly, this shell was seen to rapidly “flow” on the particle surface throughout the reaction (see Video S1 ), indicative of the fast diffusivity of Sn via surface diffusion. As a result, an ordered PtSn phase formed at 3 s, as evident by the appearance of the superlattice spots in the FFT ( Figure 2 B). By applying masks in the Fourier space (marked by the yellow circles), we calculated a filtered real-space image containing only the information from PtSn by the inverse fast Fourier transform (IFFT) ( Figure 2 C). The IFFT image confirmed that the PtSn phase formed mainly at the surface of Pt. Similar Fourier filters were applied to the images at 18 s ( Figure 2 D) and 27 s ( Figure 2 G). The IFFT images in Figures 2 F and 2I confirmed that PtSn became more prominent as the reaction proceeded.

The Pt/SnO@mSiOnano-reactor featured the well-encapsulated Pt/SnOreactant pairs in the mesoporous silica shells. Figure 1 A shows a typical HAADF-STEM image of the sample. The Pt particles in bright contrast had an average diameter of 4.0 ± 1.4 nm, derived from a histogram of size distribution shown in the inset. The SnOparticles in light-gray contrast could be found either side by side or as shells in a core-shell structure with Pt nanoparticles. The silica shells that surround the particles rendered in dark-gray contrast. Elemental mapping on a single capsule by energy-dispersive X-ray spectroscopy (EDS) further confirmed the well-encapsulated Pt/SnO, as shown in Figure S1 . A reactant pair with a ∼5 nm Pt particle (top) and a ∼10 nm SnOparticle (bottom) was selected for in situ observation, as presented in Figure 1 B. Fast Fourier transform (FFT) from the Pt region ( Figure 1 C) shows the d-spacings matching those of the face-centered cubic Pt. The gray region shows two discernible α-SnOdomains with different crystal orientations, one of which is near [11], as confirmed by the FFT ( Figure 1 D).

Discussion

2 in the Pt/SnO 2 @mSiO 2 nano-reactor encompasses consecutive transformations: the Pt 3 Sn phase was first formed at the surface of the Pt particle once the SnO 2 was reduced; further reducing SnO 2 resulted in the transformation of Pt 3 Sn to PtSn. At the beginning of the reaction, the Pt surface facilitated diffusion of free Sn, as evident by the “flowing shell” on Pt (see 3 Sn, which was seen right after free Sn was generated ( 3 Sn suggests that this process is mainly controlled by Sn diffusion. PtSn nucleation was first seen on the particle corner (∗. 32 Porter D.A.

Easterling K.E. Phase Transformation in Metals and Alloys. ∗ = 1/2V∗ ⋅ ΔG V , where ΔG V is the difference of the free energies per unit volume of PtSn and Pt 3 Sn. More importantly, with a unique coherent interface, ( 110 ) Pt 3 Sn ∥ ( 001 ) PtSn and 1 1 ¯ 0 Pt 3 Sn ∥ 210 PtSn , the interfacial energy was minimized. Therefore, the energy barrier was further reduced, given that V∗ ∝ γ3, where γ is the interfacial energy. 32 Porter D.A.

Easterling K.E. Phase Transformation in Metals and Alloys. 3 Sn phase is required. The driving force for this uphill diffusion was the tensile stress (induced by the interface) on Pt sites. The Sn atoms in front of the interface were pulled toward the PtSn side, leading to the substitution of Pt by Sn and the formation of PtSn. When the region in front of the interface was depleted of Sn, e.g., after 82.5 s of the reaction (see 2 and hence secure the Sn diffusion. As a result, we were able to resume the formation of Pt 3 Sn, as well as its transformation to PtSn. Therefore, Sn diffusion plays a crucial role in the formation of ordered phases in the Pt-Sn system. The above in situ observation demonstrated that the reaction between Pt and SnOin the Pt/SnO@mSiOnano-reactor encompasses consecutive transformations: the PtSn phase was first formed at the surface of the Pt particle once the SnOwas reduced; further reducing SnOresulted in the transformation of PtSn to PtSn. At the beginning of the reaction, the Pt surface facilitated diffusion of free Sn, as evident by the “flowing shell” on Pt (see Video S1 ). This surface-mediated Sn diffusion was responsible for the fast formation of PtSn, which was seen right after free Sn was generated ( Figure 2 ). The continuous growth of PtSn suggests that this process is mainly controlled by Sn diffusion. PtSn nucleation was first seen on the particle corner ( Figure 2 D), where the nucleus could minimize its critical volume, VThe energy barrier was hence lowered, given by ΔG= 1/2V⋅ ΔG, where ΔGis the difference of the free energies per unit volume of PtSn and PtSn. More importantly, with a unique coherent interface,and, the interfacial energy was minimized. Therefore, the energy barrier was further reduced, given that V∝ γ, where γ is the interfacial energy.In order for the PtSn nucleus to grow and the interface to migrate, a net flux of Sn atoms from the PtSn phase is required. The driving force for this uphill diffusion was the tensile stress (induced by the interface) on Pt sites. The Sn atoms in front of the interface were pulled toward the PtSn side, leading to the substitution of Pt by Sn and the formation of PtSn. When the region in front of the interface was depleted of Sn, e.g., after 82.5 s of the reaction (see Figures 3 E–3H), the transformation halted because of insufficient Sn supply. Then, the beam current was increased to speed up the reduction of SnOand hence secure the Sn diffusion. As a result, we were able to resume the formation of PtSn, as well as its transformation to PtSn. Therefore, Sn diffusion plays a crucial role in the formation of ordered phases in the Pt-Sn system.

3 Sn formed at 3 min ( 3 Sn was noticed at 18 min in the ∼15 nm particle ( 43 Jiang Q.

Zhang S.H.

Li J.C. Grain size-dependent diffusion activation energy in nanomaterials. Moreover, particle size also alters the reaction dynamics. We conducted similar observation on ∼10 and ∼15 nm Pt particles, and the selected sequential HAADF-STEM images are presented in Figures S4 and S5 , respectively. It can be seen that reaction speed slowed down as the Pt size increased, e.g., PtSn formed at 3 min ( Figure S4 B) and PtSn nucleation occurred at 5 min ( Figure S4 C) in the ∼10 nm particle, whereas the formation of PtSn was noticed at 18 min in the ∼15 nm particle ( Figure S5 B), and its transformation to PtSn was observed at 23 min ( Figure S5 C). Nevertheless, the reaction followed the same pathway as we discussed in the ∼5 nm particle, suggesting that the particle size mainly affects Sn diffusion because the bigger particle has a lower diffusivity.In addition, stacking faults were observed in the ∼10 nm particle ( Figure S4 A), which seemed to facilitate Sn diffusion ( Figure S4 D). As a result, the PtSn front propagated preferentially along the defective region, which was slightly rich in Sn. However, the stacking faults eventually annihilated as the reaction proceeded ( Figure S4 E). Because of the short lifetime of this structural imperfection, its effect on the reaction dynamics is likely to be subordinate.

2 @mSiO 2 at 300°C in flowing hydrogen for varied time. 2 @mSiO 2 and the samples reduced for 5 min, 30 min, and 4 h, showing gradual transformation of Pt→Pt 3 Sn→PtSn. The crystallite size was estimated to be 4–5 nm from the XRD results ( 3 Sn-enriched surface were obtained in the sample reduced for 5 min given that Sn diffusion was limited ( 3 Sn/PtSn bicrystal structure with the coherent interface was reproduced by allowing Sn diffusion for 30 min ( 2 to PtSn phase was reached in the sample reduced for 4 h ( Figure 6 Controlled Phase Formation of Pt-Sn iNPs and Their Catalytic Performance for Acetylene Semi-hydrogenation Show full caption (A–D) Representative HAADF-STEM images of Pt/SnO 2 @mSiO 2 reduced in flowing hydrogen at 300°C for 5 min (A), 30 min (C), and 4 h (D). (B) IFFT image from (A) generated by the superlattice spots of Pt 3 Sn (circled in the inset). (E) Concentration of ethane measured at the reactor outlet versus acetylene conversion for the synthesized Pt-Sn iNP catalysts. 0% ethane corresponds to the highest selectivity of acetylene hydrogenation. The reactant gas mixture contains 13.35 mL/min He, 1.5 mL/min H 2 , 15 mL/min C 2 H 4 , and 0.15 mL/min C 2 H 2 at 1 bar. The above results also provide possibilities for controlling the formation of ordered phases in iNPs by accommodating diffusion of the second element, e.g., at different reaction conditions. As a proof of concept, we demonstrate that controlled phase formation and morphology of Pt-Sn iNPs can be achieved by thermal reduction of Pt/SnO@mSiOat 300°C in flowing hydrogen for varied time. Figure S6 presents X-ray diffraction (XRD) profiles of the starting Pt/SnO@mSiOand the samples reduced for 5 min, 30 min, and 4 h, showing gradual transformation of Pt→PtSn→PtSn. The crystallite size was estimated to be 4–5 nm from the XRD results ( Table S1 ), in agreement with the HAADF-STEM images in Figures 6 A–6D , S7 , and S8 . Particles with PtSn-enriched surface were obtained in the sample reduced for 5 min given that Sn diffusion was limited ( Figures 6 A, 6B, and S7 ). The PtSn/PtSn bicrystal structure with the coherent interface was reproduced by allowing Sn diffusion for 30 min ( Figures 6 C and S8 ). Full conversion of Pt/SnOto PtSn phase was reached in the sample reduced for 4 h ( Figure 6 D).