Cyril Langlois
email: cyril.langlois@insa-lyon.fr

Ziyou L Li

Jinjiang Yuan

Damien Alloyeau

Jaysen Nelayah

Davide Bochicchio

Riccardo Ferrando

Christian Ricolleau

Jun Yuan

Transition from core-shell to Janus chemical configuration for bimetallic nanoparticles

come

Introduction

Bimetallic nanoparticles have been a subject of numerous studies due to the increased capabilities for tuning their properties. Compared to monometallic nanoparticles, they have additional degrees of freedom to adjust the magnetic, optical and catalytic properties, by playing with the composition and the chemical configuration. [1][2][3][4] Bimetallic nanoparticles can adopt either alloyed or segregated configurations. The alloyed configuration can be either ordered or disordered, accompanied by very different properties. For instance, ordered CoPt nanoparticles present a tetragonal structure which exhibits a high magnetocrystalline anisotropy, whereas disordered nanoparticles are face centered cubic and do not exhibit such anisotropy. [5][6][7] Another example, in the field of catalysis, is the use of Ru-Pt nanoparticles to remove carbon monoxide from hydrogen obtained from reformatted hydrocarbons. The article from Alayoglu et al. illustrates perfectly the distinct properties of particles with different chemical configuration. 8 Core-shell nanoparticles allow to initiate the selective oxidation reaction at lower temperatures compared to their alloyed counterpart, by combining surface strain effect and a modification of the electronic structure of the Pt shell.

Many different segregated chemical arrangements have been identified, including core-shell, onions-like and side-segregated. 9 The latter is referred as Janus particles because of the similarities between this structure, made of two juxtaposed metallic phases, and the two faces of the Roman god symbolizing 'transitions'. These segregated configurations bring new possibilities for obtaining multifunctional particles. 10 One unique object can easily present both magnetic and optical properties. For instance Co-Ag core-shell particles have magnetic properties due to the Co core, while as the same time exhibit a plasmon resonance associated to the Ag shell that can be tuned by varying the ratio between the shell thickness and the core diameter. 11 For the Janus configuration, the idea is no more to protect the core from external reactions. On the contrary, both metals are accessible for interactions with the environment, which makes possible the 50 design of a new class of catalysts. Such adjustments of the properties need a precise control of the structure of the particles and especially of their chemical configuration. This requires significant efforts for controlling the growth of the nanoparticles, either by chemical or physical routes. 55 It is of fundamental interest to study under which conditions a core-shell growth is favored compared to the Janus configuration. In the present study, we have chosen the Cu-Ag system as a model system to determine the possibilities to control the chemical configuration using a physical route. The synthesis is 60 made by pulsed laser deposition under ultra-high vacuum conditions, via a sequential deposition procedure. We show that this particular growth technique is perfectly adapted to obtain segregated bimetallic nanoparticles, with the temperature of the substrate and the absolute quantity of Ag in a particle as the main 65 parameters that drive the chemical configuration. To explain the transition from a core-shell configuration to a Janus configuration as a function of Ag quantity, we have conducted densityfunctional theory (DFT) calculations and atomistic molecular dynamics simulations to investigate the stability of this system.

2 | Journal Name, [year], [vol], 00-00
with a pulse duration of 25 ns and a repetition rate of 5 Hz. The energy of the pulses was adjusted separately for Ag and Cu depending on their respective ablation threshold. The pulse energy of the laser was 160 mJ for Ag and 210 mJ for Cu. The resulting deposition rates were typically 0.14 nm per minute for Cu and 0.5 nm per minute for Ag. Commercial TEM grids supporting a 10-15 nm thick amorphous carbon film were used as substrate in this study. The temperature of the substrate during the deposition could be adjusted in the range of 20°C to 800°C. The thickness deposited for a given element was measured by an in situ quartz crystal monitor that gives the nominal thickness in a continuous thin film approximation. For each sample, we carried out Cu deposition first, immediately followed by Ag deposition. A 10 nm-thick protective layer of amorphous carbon could be optionally deposited on the sample after cooling down to room temperature. Several Cu@Ag samples were synthesized for this study. The 2@1 samples were obtained by deposition of a nominal thickness of 2 nm of Cu followed by deposition of 1 nm of Ag, with the substrate temperature held at 270°C during all the process. With the same scheme, 1@2 samples were also prepared. The 2@1-100 samples were obtained with the same deposition procedure as for 2@1 samples, except the temperature of the substrate for Ag deposition that was decreased to 100°C. For each synthesis, a reference sample of pure Cu was grown under the same conditions, referred as 2@0 sample. When a protective layer of amorphous carbon was deposited, the samples were labeled with the 'ac-' prefix on their name, such as 'ac-2@0', 'ac-2@1', and 'ac-2@1-100' respectively.

Several TEM techniques have been employed in this study: conventional Bright Field imaging (BF), Energy Filtered Imaging (EFTEM), High Resolution imaging (HRTEM) and High Angle Annular Dark Field imaging in scanning mode (HAADF-STEM) on two different transmission electron microscopes operating at 200 kV: a JEOL JEM-2100F (University Paris 6, France) and a FEI Tecnai F20 STEM (University of Birmingham, UK).

Elemental maps have been obtained by the three-window technique in EFTEM, and the filtered images at different energies have been aligned using the script developed by Schaffer et al.. [START_REF] Reimer | Energy-Filtering Transmission Electron Microscopy[END_REF][START_REF] Schaffer | [END_REF] The EFTEM technique is particularly adapted to study the chemical configurations of bimetallic nanoparticles, as we shown in our previous work. 14

Results

Cu-Ag chemical configuration as a function of substrate temperature and deposited thickness

HAADF-STEM observations have been carried out on 2@0 and 2@1 samples. The microstructural evolution due to the Ag deposition is clearly evidenced in the two micrographs in Fig. 1, together with the size distributions. Size distribution characteristics and the particle density on the substrate are given in Table 1. For sample 2@0, the pure Cu nanoparticles show a round shape, characterized by a size distribution on the substrate centered around 10 nm. For sample 2@1, a bimodal distribution is observed with the mean peak positions at 3 and 12 nm. It is worth noting that, apart from the small particles around 3 nm, the particle density on the substrate is significantly lower than for pure Cu particles, decreasing from 4700 µm -2 to 3800 µm -2 . on the 2@1 sample. Figure 2(a) shows the zero-loss BF image of a representative area on sample 2@1. The silver chemical map of the same area is presented in Fig. 2(b). A majority of the nanoparticles presents a core-shell configuration, with a mean shell thickness of 1 nm  0.5 nm. From this EFTEM chemical 70 map, it is clear that the small particles of 3 nm are attributed to pure Ag nanoparticles nucleating on their own on the surface, a result confirmed by HRTEM observations. The capping of the Cu core can be evidenced by Moiré fringes on zone-axis particles and previous work on 2@1 samples using electron tomography 75 showed that the 3D shape of the particles was a truncated sphere with aspect ratio of 1.4. 16 Some particles are linked by a silver bridge, showing signs of sintering process at different stages. On 1.

The temperature of the substrate can have a strong influence on the chemical configuration of the sample. EFTEM experiments were carried out on 2@1-100 samples. A color chemical map of Cu and Ag on a representative area of the sample is presented in Fig. 3. It is clear that most of the particles are now multi-cores, 5 with Ag bridging several Cu cores inside one large ensemble. Though usually two or three Cu cores are present in one ensemble, up to six Cu cores per ensemble can also be observed, as pointed out by an arrow in Fig. 3. Small core-shell nanoparticles were also formed on 2@1-100 samples, with a 10 circular projection in the substrate plane that likely corresponds to a direct growth of the shell from Ag atoms diffusing on the substrate together with those directly landing on the particle. The large multi-core ensembles were also present in sample 2@1, but in very small proportion (see particles labeled 6 on Fig. 2(b)). To explore the influence of global composition on the chemical configuration of the bimetallic nanoparticles, 1@2 samples have been synthesized. Figure 4 presents the TEM results, with a representative view of the nanoparticles on (a), the size 25 distribution on (b), EFTEM chemical map on (c) and HRTEM on (d). The mean size is around 12 nm (Table 1), very similar to the 2@1 particles with identical total nominal thickness. On the contrary, it is clear from the chemical map of Fig. 4(c) that the chemical configuration is different from that of sample 2@1. A 30 majority of bimetallic particles exhibits a Cu core that is not at the center of the particle, and sometimes touches the outer rim of the particle to form a Janus-like configuration. This is a general tendency over the whole sample, confirmed by high-resolution imaging. A representative example is shown in Fig. 4(d).

35

Thermal stability of CuAg system

Thermal stability of the CuAg particles is addressed through annealing experiments under high vacuum of samples ac-2@1-100 and 2@1-100, with and without amorphous carbon capping 40 layer respectively. The temperature of the sample was raised to 550°C and held for 90 min, to allow for a complete evolution of the particle assembly toward equilibrium. Figure 5 shows the initial and final configurations of the ac-2@1-100 sample. The initial configuration is typical of as-grown sample 2@1-100. The 45 post-synthesis deposition of the amorphous carbon did not alter the particles characteristics. Pure Ag nanoparticles were present between the bimetallic particles, the mean diameter of the latter being around 30 nm. The only change after annealing seems to be the disappearance of the small pure Ag nanoparticles. Most grain 50 boundaries visible in the initial configuration were still there after annealing, indicating little changes within the particles. The same annealing experiment for the uncapped 2@1-100 sample resulted in more pronounced evolution. Figure 6 shows the particle assembly after annealing, in BF imaging (Fig. 6(a)) 55 and EFTEM imaging (Zero-loss (ZL) image in Fig. 6(b) and Ag elemental map in Fig. 6(c)). The size distribution is given in the inset of Fig. 6(a). The mean size of 32 nm and the particle density of 720 µm -2 demonstrate a very strong coalescence between the particles. This phenomenon completely altered the initial 60 structure to larger particles, with a typical Janus chemical configuration, that is, with Cu on one side of the particle and Ag on the other side. In most cases, Cu and Ag were separated by a unique planar interface crossing the particle. The planar interfaces could also present an angle, as for the particles #1 in Fig. 6(c). 65 Here, two distinct structural features are noted: (1) small coreshell nanoparticles were still present on the substrate, under a critical size of 20 nm, and (2) the Cu part of Janus nanoparticles is still surrounded by a thin Ag shell up to 3.5 nm thick.

Discussion

Growth mechanism of the 2@1 microstructure

From the results on samples 2@0 and 2@1, the mechanisms leading to the structural characteristics can be inferred. The 20 decrease of the particle density from N 1 = 4700 to N 2 = 3800 particles per squared micron due to silver deposition corresponds roughly to n = 900 coalescence events per squared micron during the deposition of silver. It means that, assuming only paircoalescence events of particles, almost 40 % of the existing Cu particles undergo sintering. Let us consider for simplicity sake that all Cu particles coalescing are initially identical, with the same size d 1 (equal to the mean size of the size distribution of 2@0 sample) and an hemi-spherical morphology (see the model on Fig. 7). The coalescence results theoretically in a double-peak size distribution for Cu cores with peaks positions at d 1 and 2 1/3 d 1 . The mean size d 2 of the CuAg particles for sample 2@1 should then be:

  3 1 1 1 2 2 3 1 1 22 2 21 12 s s N n d n d dt N d n t Nn               
with s t the Ag shell thickness measured after coalescence. The second term inside the brackets corresponds to the increase percentage compared to the initial size, that is 6%. This leads to a theoretical increase of the mean size to d 2 = 12.6 nm, which corresponds nicely to the effective mean size of the particles in the size distribution on sample 2@1 (see Table 1). On the size distribution, we of course do not see the two theoretical peaks because of the polydispersity of the size distributions. From this analysis, the most probable growth mechanism of the 2@1 coreshell nanoparticles is the formation of an Ag shell around each copper nanoparticles, with pair-coalescence for ~40 % of the particles to give larger core-shell nanoparticles. One possible driving force for the coalescence to happen is through diffusion of the metallic species on the substrate. To get insight on the first steps of the initial core-shell structure formation, the diffusion of silver and copper on the carbon substrate were lowered, by reducing the substrate temperature for the Ag deposition from 270°C to 100°C. Firstly, we observed that the size distribution and particle density of the Cu cores of the 2@1-100 samples were close to the ones of 2@0 sample. This means that for the 2@1-100 sample, the temperature of 100°C was not sufficiently high to drive the Cu particle diffusion to lead to their coalescence. Secondly, the Ag elemental map reveals that this temperature was not high enough to cause the atomic rearrangement of the incoming Ag atomic species around the Cu particles. This effect could be related to the high silver deposition rate of 0.5 nm per minute. Here, we can conclude that, at 100°C, the pair-coalescence of particles originates from Ag bridges formed directly during the deposition of silver.

It is worth noting that for other substrates, similar results will be obtained provided that the following general criterion is fulfilled: interactions between the metallic species considered and the substrate must be weak. This is the key point to obtain particles instead of a thin film in the first step of the growth process and to attract the second metallic species around the already formed particles in the second step of the growth. Different metals can exhibit different behavior when interacting with the same substrate. 17 However, the positive point is that an increase of substrate temperature during the growth can overcome metal/substrate interactions to a certain extent and leads to the same kind of nanoparticles. Zero-loss EFTEM image on a representative area of sample 2@1-100 sample and (c) superimposed Ag and Cu elemental map of sample 2@1-100 obtained by EFTEM using the three-window technique on the same area. Color code identical to Fig. 3. 12,13 60

Diffusion mechanisms in the bimetallic Cu-Ag nanoparticles

It is clear from the annealing results that a capping layer of amorphous carbon completely blocks any coalescence of the particles on the substrate, only allowing the small Ag particles to 65 dissolve themselves in the carbon matrix by Ostwald ripening mechanism. Moreover, no internal re-organization of Cu and Ag inside an initial ensemble was noticed. The latter observation is important because theoretically, the temperature was high enough to induce a chemical re-arrangement by volume diffusion 70 mechanisms of Cu and Ag in the particles. 18 If no re-arrangement is observed, it means that the presence of carbon at the surface blocks the mechanisms responsible for the atomic displacements inside a particle. We could deduce from this that it is mainly the diffusion on the surface of the particle that drives any evolution

6 | Journal Name, [year], [vol], 00-00
of the chemical configuration in the CuAg particles. Another reason could be also that the surface energy of the particles changes due to the carbon matrix, impeding the driving force toward internal re-arrangement with temperature. 19,20 On the contrary, the annealing has a strong effect on the uncapped particles that are free to undergo Ostwald ripening and to coalesce through facile diffusion on the substrate. 21 Without the carbon layer, the coalesced particles can change their chemical arrangement toward Janus configuration. These conclusions are in good agreement with previous simulations carried out on the 10 Cu-Ag system in which the authors show that the surface diffusion processes are energetically favored compared to diffusion inside the core. It seems also important to underline that oxidation hardly affects the nanoparticles. The convergence of several strong arguments 20 allows us to rule out the formation of any oxide in the particle, at least in a significant way: -all the EFTEM experiment on the oxygen threshold that we carried out show very low signal-to-noise ratio and do not evidence the presence of oxide. EELS spectra do not exhibit the typical fine structure of oxidised Cu neither. -all the HRTEM images acquired on these samples are free of any Cu or Ag oxide lattice fringes, and often show perfect epitaxy between the two metals. -in case of oxidation of copper, an inversion of the core-shell configuration would be expected, with Ag in the core and Cu oxide as a shell because the mean surface energy of Cu 2 O is less than the one of Ag. 23 -we never noticed any difference between carbon-capped and air-exposed 2@1 sample.

Competition between core-shell and Janus chemical configuration

It is interesting to note that, although the 2@1 and 1@2 samples 40 are very similar from the point of view of synthesis apart from their chemical composition, the EFTEM characterization reveals a different internal structure of CuAg particles. For the 1@2 sample, not only the average Ag shell thickness is larger as the Ag content increases, but also the core is off-centred and very 45 close to the rim of the particle. These results support the idea that the increase in shell thickness is responsible for the evolution of the chemical configuration toward Janus configuration. This hypothesis can be verified by the annealing experiments. In our annealing experiments, the amount of matter was kept constant 50 (equivalent to 2@1 samples) but coalescence events occurred, resulting in larger particles, hence larger core radius and a larger amount of silver available for re-arrangement around the Cu core. This suggests that both an increase of Ag concentration and a coalescence event have a consequence in common, that is an 55 increase of the amount of silver for a given CuAg particle. Therefore, the Janus configuration identified for the annealed samples may originate from a larger shell thickness that turns finally to the side-segregated configuration, as for the 1@2 samples. From the measurement on the annealed sample and the 60 1@2 sample, the critical thickness c t at which the transition occurs from the core-shell to the Janus configuration is estimated to be 3-4 nm. Because of the large size distribution of the particles in the samples, we always find both type of chemical configurations in a given sample. This result is confirmed by 65 measurements carried out on larger CuAg nanoparticles, synthesized by thermal evaporation in a previous work. 14 These particles correspond to a 6@3 system. Both Janus and perfect core-shell nanoparticles are present on the substrate, but for the perfect core-shell configuration, there are no shell with a 70 thickness greater than 4 nm, independently of the Cu core size.

To gain further insight of the driving force for the transition from the core-shell to Janus configuration, we carried out DFT calculations and molecular dynamics simulations (MD) using atomistic potentials derived from second-moment approximation 75 to the tight-binding model. DFT calculations have been made by using the Perdew-Burke-Ernzerhof exchange-correlation functional (PBE) and the Quantum Espresso code. 24 Details are given in the work from Bochiccio et al.. 25 We report atomistic calculations by means of two different parameter sets of the 80 Gupta potential. 26 Parametrization P1 has been fitted on DFT calculations of small Ag-Cu clusters, 25 whereas parametrization P2 was fitted on the experimental energetics of single impurities of the bulk host matrix of the other metal. 22 The energy differences between these positions in the cluster are small, so that zero-point energy contribution may be important. 28 However, when the concentration of impurities is finite, the energy differences between aggregates in central and subsurface positions greatly increases. For example, we have considered the truncated octahedron of 201 atoms with 13 Cu atoms. According to P2, the best subsurface position of the Cu aggregate is lower by about 1 eV than the central position of the aggregate. On the contrary, an Ag impurity in a Cu cluster would prefer lowcoordination sites, such as vertex sites on the surface. We have verified this for the truncated octahedron of 201 atoms, in which the vertex site is favoured by 0.9 eV compared to central and subsurface sites. All these results are consistent with the formation of off-center Cu aggregates in Ag-rich clusters. If we consider only the composition inside the particle, it is clear that a low concentration of Ag (i.e. thin Ag shell) corresponds to a core-shell configuration (as clearly shown by the global optimization results 25 ) whereas a large concentration of Ag (i.e. a thick Ag shell) should correspond to a Janus configuration.

In order to verify this hypothesis, we have performed globaloptimization searches of chemical ordering in Ag-rich clusters.

We have considered an fcc truncated octahedron of 586 atoms containing 90% Ag and 10% Cu. We have initially placed the Cu atoms randomly in the cluster and then we have performed extensive global optimization runs in which only exchange moves between atoms of different species were used. The lowest energy 50 configurations according to P1 and P2 are shown in Fig. 9. For both models, an off-center Cu core is obtained so that a quasi-Janus particle is formed. The chemical ordering of these nanoparticles clearly resembles that of the experimental cluster of Fig. 4(d).

is worth noting that this result has not been experimentally observed for nanoparticles, even though it is likely to be representative of a quite general behavior. In fact, atomistic simulations are predicting that the same kind of behavior is likely to be observed in AgCo and AgNi too. 30,31 In order to have a complete view of the problem, it is important to consider also the surface science literature on the CuAg system. A brief review shows that the growth modes of Ag on Cu or Cu on Ag are very different, they strongly depend on the orientation of the substrate, on the substrate temperature and on the amount of matter deposited. The growth of Ag on Cu (111) has been widely studied: although some discrepancy exists, a layer-bylayer growth mode is favoured. [32][33][34] Twinning has been observed, resulting in similar heteroepitaxy relationships as for the bulk. 35,36 When the Cu substrate is [110] oriented, the growth mode is reported to be layer-by-layer at low temperature, and Stranski-Krastanov at high temperature. 37 For the growth of Cu on Ag, the Stranski-Krastanov growth mode is the most frequently reported. 29,38 From these experimental data, it is clear that the situation is very complex in the case of a nanoparticle because

Conclusion

We have presented in this study the synthesis of bimetallic Cu-Ag nanoparticles via pulsed laser ablation, with various chemical compositions. We demonstrated that the deposition of silver on pre-existing copper cores (two-step growth procedure) induced a sintering between the particles, if the substrate temperature was 45 high enough to enable surface diffusion on the substrate. The thermal stability of these nanoparticles has also been studied by annealing experiments under high vacuum conditions. Cu and Ag adopted a segregated chemical configuration, according to the immiscibility of this alloy and the strong difference of their mean 50 surface energy. Two types of segregated configurations were identified: a core-shell configuration for shell thickness under 3 to 4 nm, and a Janus configuration when the amount of Ag inside a particle was in excess compared to this critical thickness. For the Janus configuration, we determined that the Cu part of a

Fig. 2 Fig. 1 (

 21 Fig.1(Color online) HAADF-STEM images of samples (a) 2@0 Cu and (b) 2@1 CuAg. The size distribution for each micrograph is also presented. Characteristics of the distribution are detailed in Table1.

Fig. 2 (

 2 Fig. 2 (Color online) (a) Zero-loss EFTEM image on a representative area of sample 2@1 and (b) Ag elemental map obtained by EnergyFiltered20

Fig. 3 (

 3 Fig. 3 (Color online) Superimposed Ag and Cu elemental map of sample 2@1-100 obtained by EFTEM using the three-window technique. The arrow designs a CuAg ensemble with six copper cores linked by silver.

Fig. 4 (

 4 Fig. 4 (Color online) TEM observations of sample 1@2. (a) Bright field TEM image on a representative area of the sample, (b) the size distribution of sample 1@2, (c) Ag elemental map obtained by EFTEM using the three-window technique and (d) HRTEM image of a Janus nanoparticle.

Fig. 5

 5 Fig. 5 Annealing experiment on ac-2@1-100 sample: bright field TEM image of (a) the initial configuration and (b) final configuration after 90 15

Fig. 6 (

 6 Fig. 6 (Color online) Annealing experiment on 2@1-100 sample. (a) Bright field TEM image of the final configuration of the bimetallic nanoparticles. The inset shows the size distribution of the particles. (b) 55

22

 22

Fig. 7 (

 7 Fig. 7 (Color online) Schematic model showing the result of coalescence of two identical hemispheric Cu particles, with conservation of the total volume of Cu and the hemispherical morphology.

Fig. 8 (

 8 Fig. 8 (Color online) Model for the truncated octahedral Ag78Cu1 cluster used in the DFT, P1 and P2 calculations. The external surface of the Ag78Cu1 cluster is represented in (a). The cluster configuration in (b) represent the Cu atom in the most favorable place according to DFT calculations.

20 different 30 Fig. 9 (

 20309 Fig. 9 (Color Online) Structure of a truncated octahedron of 586 35

  55particle remains covered by a thin Ag layer. The above observations are supported by DFT calculations and MD simulations, which suggest that the energetically favorite position for a Cu impurity in an Ag cluster is a sub-surface position. These sub-surface impurities can act as seeds for the nucleation 60 of off-center Cu clusters, and lead to the formation of quasi-Janus structures for Ag-rich compositions. The present work is a step forward to the understanding of the mechanisms leading to the formation of a particular chemical configuration for bimetallic particles. Further studies on the 65 influence of the chemical configuration on the optical properties of such systems are currently under investigation.

Table 1

 1

		Data relative to the particle assemblies of samples 2@0, 2@1 and
		1@2 obtained by TEM image thresholding. 15
	60
		Sample	2@0	2@1	1@2
		Data from particle distribution
		Particles mean diameter (nm)	9.9	12	11.8
		Peak FWHM (nm)	5.3	8.3	6.2
		Particle density
		(particles per µm²)	4700	3800 3700
		In order to investigate the chemical configurations that Cu and
		Ag adopt on the substrate after the complete deposition
	65	procedure, EFTEM imaging experiments have been carried out

  Bochicchio et al. 25 ). We have calculated the energetics of single 90 Cu impurities placed inside truncated octahedral Ag clusters (see Fig.8) in order to determine the most favorable position inside the cluster. We have considered three cluster sizes: 79, 201 and 586 atoms. For the smallest cluster, we performed both DFT and atomistic potential calculations. For 201 and 586 atoms, DFT calculations become very cumbersome, so that only atomistic calculations were performed. The results are presented in Table2, with the energies associated to Cu positions under different types of facets. For the Ag 78 Cu 1 the DFT calculations are in good agreement with the results of P2 potential, both showing that for the Cu impurity, the sub-surface positions, just one layer below the surface, are always more stable than central and surface positions. Potential P1 does not agree with DFT and P2 for this size, because it prefers the central positions. For Ag 200 Cu 1 and Ag 585 Cu 1 , P2 confirms the tendency to favor sites that are just one layer below the surface. Also P1 does not favor the central sites anymore when size increases. According to P1, the most favorable sites for Cu impurities are located two layers below the surface. Our previous work on the simulation of the Cu-Ag system has shown that, when Cu atoms are deposited on a Ag cluster (inversed deposition model), the Cu atoms agglomerates under the surface in order to form a pure Cu cluster.27

		These atomistic
	85	potentials have been thoroughly tested against DFT calculations,
		which showed that P1 agrees very well for what concerns the
		energetics of core-shell nanoparticles. The agreement of P2 is
		slightly worse, but the qualitative behaviour is the same for both
		P1 and P2 (see the supplementary online material of Ref.

Journal Name, [year], [vol], 00-00 | 7

K. Major, C. De and S. Obare, Plasmonics, 2009, 4, 61-78.

N. S. Sobal, M. Hilgendorff, H. Mahwald, M. Giersig, M. Spasova, T. Radetic and M. Farle, Nano Letters, 2002, 2, 621-624.

H. Zeng, S. Sun, J. Li, Z. L. Wang and J. P. Liu, Applied Physics Letters, 2004, 85, 792-794.

Acknoledgments

We are grateful to the Egide (France) and British Council (UK) for supporting this study through the Alliance project, as 70 well as the European COST Action MP0903. R.F. and D.B are grateful for support from CINECA supercomputing center for the project HP10CF1XI1. ZY.L. and J.Y. thank supports from UK EPSRC. We are also grateful to Region Ile-de-France for the support to the JEOL 2100F electron microscope installed at

Table 2 Results from DFT and atomistic calculations on Ag clusters containing a single Cu atom. For each size, the data report the energy difference from the configuration with Cu on the center of the cluster, which is taken as the zero of energy.


D Alloyeau
email: damien.alloyeau@onera.fr

C Langlois

C Ricolleau

Y Le Bouar

A Loiseau

TEM in situ experiment as a guideline for the synthesis of asgrown ordered CoPt nanoparticles

come

A TEM in situ experiment as a guideline for the synthesis of as-grown ordered CoPt nanoparticles

Introduction

The control of nanoparticle size and structural arrangement are essential challenges in nanofabrication.

These parameters are particularly important in the case of CoPt and FePt bi-metallic nanoparticles since they govern their magnetic properties. These alloys exhibit, in the bulk state, a phase transition between a tetragonal chemically ordered phase (L1 0 ) at low temperature and a disordered Face Centred Cubic (FCC) structure at high temperature. The L1 0 ordered phase, due to the alternative atomic stacking of pure Co or Fe and pure Pt layers, has large magnetocrystalline anisotropy. By means of this magnetic property, important research efforts are devoted to study L1 0 ordered nanoparticles which are expected to be the future information storage media for extremely high density recording (EHDR) [START_REF] Sellmyer | Nanostructured magnetic films for extremely high density recording[END_REF][START_REF] Yu | Nanostructure and magnetic properties of composite CoPt:C films for extremely high-density recording[END_REF].

Among the different physical techniques currently used, only few succeed in obtaining as-grown L1 0 structure for FePt and CoPt alloyed nanoparticles [START_REF] Miyazaki | Size effect on the ordering of L1 0 FePt nanoparticles[END_REF]. The disordered structure (A1 phase) is largely predominant for a growth temperature ranging from 20°C to 600°C. This is a non-equilibrium structure compared to the bulk state phase diagram (figure 1) [START_REF] Bouar | Origin of the complex wetting behavior in Co-Pt alloys[END_REF][START_REF] Leroux | Etude du diagramme de phase et des transformations ordre désordre du système Co-Pt par microscopie électronique et résistivité[END_REF]. The ordered structure is obtained usually by a post-synthesis annealing at temperature higher than 600°C. Because of this elevated annealing temperature (~T m /2, T m being the melting temperature of the bulk alloy), this procedure may induce coarsening or coalescence of the nanoparticles, which results in a particle size out of the range considered for EHDR [START_REF] Yu | Nanostructure and magnetic properties of composite CoPt:C films for extremely high-density recording[END_REF][START_REF] Castaldi | Coevaporation of CoPt nanoparticles[END_REF]. Indeed, considering an area density of nearly 1Tbit/in 2 for the next generation of hard disks, the bit dimension will reach a nanometric size of approximately 10 nm [START_REF] Hannour | D arrays of CoPt nanocluster assemblies[END_REF].

Avoiding the annealing step would be a clear technological breakthrough, and numerous research groups aim at the fabrication of ordered as-prepared nanoparticles by optimization of growth parameters. In the case of vapour phase deposition techniques, adjustments of key parameters like the deposition rates, the nature of the substrate [START_REF] Langlois | Growth and structural properties of CuAg and CoPt bimetallic nanoparticles[END_REF] and particularly the temperature of the substrate can improve dramatically the ordering of the as-grown nanoparticles, without having recourse to annealing.

In this paper we show how a finely tuned pulsed laser deposition (PLD) experiment on amorphous alumina results in ordered as-grown CoPt nanoparticles. In a first part, we show that we can reliably select any composition for the Co x Pt 1-x nanoparticles. Then, to better understand the role of the substrate temperature on the degree of order and on the morphology of CoPt nanoparticles embedded in alumina, we performed in situ heating experiments in a transmission electron microscope (in situ TEM). The temperature regimes observed during the in situ experiments are used to optimize the PLD parameters. Finally, the growth mechanisms during the synthesis and the key role of temperature in the ordering and coalescence mechanisms are discussed.

Experimental procedures

Sample preparation

The CoPt nanostructured thin films on amorphous alumina (a-Al 2 O 3 ) were produced by pulsed laser deposition in a high vacuum chamber. The pressure in the chamber is better than 10 -7 mbar. A typical target-substrate configuration is used to deposit separately the a-Al 2 O 3 and the metals by PLD using a KrF excimer laser at 248 nm with pulse duration of 25 ns at repetition rates in the range of 1 to 20 Hz.

The laser energy can be chosen in the range of 150 to 250 mJ depending on the ablation threshold of the selected target. The distance from targets to substrate was approximately 5 cm. Samples are obtained by alternated irradiation of the two metal targets to produce the plasma of each metal. The substrates used were commercial TEM grids on which we have deposited an amorphous carbon layer with a thickness of 10 nm. On the top of the amorphous carbon, a 3 nm layer of a-Al 2 O 3 was deposited using our PLD setup. Then, cobalt and platinum were alternatively deposited using pure Co and Pt targets. The deposition rate of each element is controlled by an in situ quartz crystal monitor, which indicates the thickness of deposited materials on the quartz surface. This thickness is then defined as the nominal thickness in a continuous film approximation. The deposition rate depends on the target nature and on the laser frequency. In this work, all the samples were synthesized with a laser frequency of 5 Hz. With this frequency, the ablation rates are respectively 3 nm, 0.15 nm and 0.02 nm per minute for a-Al 2 O 3 , cobalt and platinum. The quantity of Pt and Co deposited at each step depends on the targeted composition. In addition, to insure a good homogeneity of the film, the nominal thickness corresponding to two successive steps (one Co and one Pt) is set to 0.1 nm. During the deposition, the substrate is held at a fixed temperature T s , between room temperature and 725°C. At the end of synthesis, samples were covered by a 3 nm thick layer of a-Al 2 O 3 in order to protect alloys from air oxidation. The various tuneable experimental parameters in PLD technique allow a high control on vapour-phase growth with slow deposition rate.

In situ heating experiment

In situ TEM is a powerful method to observe order-disorder phenomena in metallic alloys. In situ annealing has been performed using the GATAN heating specimen holder model 652 which allows heating from room temperature to 1000°C. The thermal stability of the sample during the heating procedure is promoted by very efficient water cooling. The evolution of the sample morphology was followed in bright field mode and the structural transformation in diffraction mode performed on a JEOL 4000FX electron microscope operating at 400kV. The experimental protocol has been defined by taking into account the equilibrium regions of the phase diagram of the bulk CoPt material (figure 1). It was however slightly modified during the experiments as a function of the evolution of the observed phenomena. The heating procedure used to observe the ordering-disordering process on a selected area was as follows: the temperature was progressively increased from 25°C up to 600°C (in 50 minutes). Then, the temperature was increased by steps of 30°C every 10 minutes, until 700°C.

Then, the steps have been reduced to 15°C every 20 minutes, from 700°C to 850°C. After half an hour at 850°C we started the temperature decrease by steps of 30°C every 10 minutes until 570°C and finally, the sample was cooled down to room temperature in one hour.

First of all, heating procedure was performed on Co 45 Pt 55 bulk sample in order to check the temperature calibration of the heating holder. For this composition the CoPt phase diagram exhibits a narrow two phase stability region at 780°C (figure 1). During the experiment, this first order phase transition has been observed at 815°C indicating that the temperature given by the thermocouple of the heating holder may differ from that of the sample by a few tens of degrees. In addition, when using carbon film deposited on copper grids, the temperature is not perfectly homogeneous on the carbon film and slightly changes with the distance from the copper bar. Therefore, during the in situ experiments, we have always observed the same area of the sample.

Results

Control of the nanoparticles composition

The synthesis of nanoparticle films by the PLD method has been previously realized with a single target containing the Co and Pt elements at chosen concentration [START_REF] Trelenberg | Femtosecond pulsed laser ablation of metal alloy and semiconductor targets[END_REF]. However, because cobalt and platinum have widely differing vapour pressures, the stoichiometry of the ablated target is not reproduced in the collected nanoparticle film. Instead of using a single target, we show here that the alternate use of two pure metal targets allows a composition-controlled synthesis of homogeneous nanoparticle films. Moreover, this approach has the obvious advantage that we can easily select any composition. We have first synthesized seven samples to cover the whole composition range between pure Co and pure Pt. The nominal thickness of the samples was 2.5 nm and the substrate temperature was fixed at 550°C. The as-grown nanoparticles are not chemically ordered and their size is in a range from 10 to 20 nm. Energy Dispersive X-Ray spectroscopy (EDX) was performed on large areas (0.4 µm) of the samples. The lattice parameter of the FCC structure was determined by means of electron diffraction measurements calibrated with reference of the bulk Co 45 Pt 55 alloy. Because the substrate is amorphous, the nanoparticles are randomly oriented on the surface and the diffraction diagram displays a typical ring pattern characteristic of a powder-like diffraction. We found that the composition of the films is homogeneous on each TEM grid within the accuracy of EDX measurements (2% at.) and we showed that it varies linearly with the lattice parameter of the nanoparticles (figure 2). In addition, the values presented in figure 2 are close to the line joining the values of the bulk Co to that of the bulk Pt [START_REF] Villars | Person's Handbook ASM International[END_REF], i.e. close to the Végard law observed for the electrodeposited CoPt thin film [START_REF] Zana | Enhancing the perpendicular magnetic anisotropy of Co-Pt(P) films by epitaxial electrodeposition onto Cu(1 1 1) substrates[END_REF]. A behaviour close to that of the thin film is consistent with the rather large size of the nanoparticles (above 10 nm in diameter) for which the size effect on the lattice parameter is small [START_REF] Wasserman | On the determination of the surface stress of copper and platinum[END_REF].

Note that we can now use the curve presented in figure 2 to deduce the composition of the disordered nanoparticle films using the lattice parameters obtained by electron diffraction. More importantly, following the experimental protocol improved during the course of the above study, we are now able to synthesize homogeneous Co x Pt 1-x nanoparticle films of any composition, with a measured composition closer than 2% to the targeted value.

Order-disorder and coalescence phenomena

To study the phase transition from chemically disordered FCC phase to the L1 0 chemically ordered phase, the in situ heating protocol described in the experimental procedures section (part 2.2) was performed on CoPt nanoparticles assemblies obtained for a nominal thickness of 2 nm (figure 3a). This sample, with a Co 55 Pt 45 composition, was prepared with a substrate temperature of 500°C and has been covered by a 3 nm thick a-Al 2 O 3 layer. As-grown nanoparticles observed at room temperature have a coverage on the substrate close to 67% and the nanoparticle size range from 5 nm to 25 nm.

The coverage was determined by measuring, using TEM bright field images, the percentage of the a-Al 2 O 3 substrate covered by the nanoparticles. At this temperature, the tendency of the metal to wet the surface results in nanoparticles of irregular shapes elongated in the plane of the substrate, which makes it difficult to define their mean size from a statistical point of view (figure 3a). The mean distance between the particles is close to 4 nm. As shown by the diffraction pattern in the inset of figure 3a, the particles are in the disordered FCC phase.

Increasing the temperature up to 600°C in 50 minutes had no significant influence on the nanoparticles mean size, but the nanoparticle morphologies tend to be slightly more spherical (figure 3b). However, the superstructure 001 ring appears at 570°C and it is clearly visible on the electron diffraction of figure 3b at 600°C. Therefore, at temperatures above 570°C, atomic diffusion inside the nanoparticles is fast enough to drive the transformation to the ordered L1 0 structure. We can also note that this 001 ring has a diffuse intensity, which can be either due to short range ordering or due to the nucleation of small ordered domains on the surface or in the volume of the nanoparticles. During the next steps of the heating, the intensity of superstructure reflections increases rapidly and the diffuse ring is replaced by a sharp one. It is also expected that the FCCL1 0 phase transition changes the values of the lattice parameters. By comparing the interplanar spacing calculated from the position of the 001 and 110 superstructure rings, we observe that the initial cubic structure becomes tetragonal.

The c/a ratio of the tetragonal structure decreases to 0.97 (i.e. a = 3.782 Å and c = 3.669 Å), which is very close to the value in bulk alloys [START_REF] Leroux | Comparative investigation of structural and transport properties of L1 0 NiPt and CoPt phases: the role of magnetism[END_REF].

No visible evolution of the morphology takes place during the heating procedure from 600°C to 690°C at the rate of 30°C every 10 minutes. When reaching 710°C, the particles start to coalesce, resulting in larger and thicker particles. The size increase of the nanoparticles is also visible on the electron diffraction pattern where sharp diffraction spots appear along the diffraction rings. The sample morphologies observed after 20 minutes at 740°C (figure 3c) show a broadening of the size distribution of the particles from 5 nm to 80 nm. After 20 minutes at 750°C (figure 3d) the size distribution of the particles is bimodal. The first population has a mean particle size of 5 nm whereas the second one has a mean particle size in the range of 20 to 100 nm. Note also that the small particles have an almost spherical shape and that the larger ones have a platelet-like shape. This point has been established at the end of the in situ experiment by tilting the sample in order to get different projected images of the particles. Detailed inspections of figures 3c and 3d and also video recording bring some clues about the coalescence mechanisms. Two coalescence mechanisms are distinguished in the literature: (a) static coalescence when the nanoparticles are immobile and coalesce by mutual rearrangement of their surface and (b) dynamic coalescence when the nanoparticles are mobile and coalesce together upon contact [START_REF] Pai | Island Diffusion and Coarsening on Metal (100) Surfaces[END_REF][START_REF] Zinke-Allmang | Clustering on surfaces[END_REF]. The latter phenomenon was not observed during our experiments. This can be due to different reasons: it occurs too rapidly to be detected or it concerns too small particles or it is inhibited by the presence of the Al 2 O 3 film. The first phenomenon takes place by atomic rearrangement of two close enough nanoparticles, leading to the formation of atom inter-diffusion bridges [START_REF] Palasantzas | Coalescence aspects of cobalt nanoparticles during in situ high-temperature annealing[END_REF]. These bridges are well visible on the examples arrowed in figures 3c and 3d. This mechanism can be favoured by the growth of the nanoparticles supplied by the diffusion of Co and Pt atoms in the Al 2 O 3 matrix, through the well known Ostwald ripening phenomenon [START_REF] Lifshitz | The kinetics of precipitation from supersaturated solid solutions[END_REF]. This mechanism involves the growth of large particles at the expense of smaller ones, when the substrate is annealed at high temperature. This coalescence results in a drastic evolution of the coverage down to 52% and an increase of the mean inter-particles distance up to 8 nm (figure 3d). At 750°C the system reaches a morphological stability: the morphology of the film does not evolve anymore during the heating up to 850°C in 130 minutes, and also during half an hour step at 850°C as well as during the temperature decrease to room temperature in 150 minutes. This stability is due to the fact that both inter-particles distance and particle size become very large which inhibit their static and dynamic coalescences.

During the coalescence phenomenon, we observe at 750°C the disordering of the nanoparticles (see electron diffraction in the left inset of figure 3d). This temperature is close to the order-disorder transition temperature (785°C) expected from the bulk CoPt phase diagram for this alloy composition (figure 1), taking into account the temperature difference between the heating system and the sample (a few tens of degrees) and the accuracy of the composition measurements (±2%). This is consistent with the fact that, in this range of size (a few tens of nanometers), nanoparticles have the same thermodynamic behaviour than the bulk material. The temperature, at which the disordering appears, is expected to be lower for the smallest nanoparticles. However, a dedicated in situ experiment, using a smaller density of nanoparticles to avoid the coalescence mechanisms and a lower heating ramp would be necessary to reveal the phenomenon.

After heating up to 850°C, we have decreased the temperature to test the reversibility of the transition. At 715°C, few sharp diffraction spots appear on the position of the 001 ring, and at 700°C the superstructure rings are clearly visible (see electron diffraction in the right inset of figure 3d). The appearance of the first sharp diffraction spots is probably due to the ordering of the largest domains. Note that the temperature at which order is observed, 715°C, differs from the temperature at which the nanoparticles become disordered during the heating ramp (750°C). This difference may be partly explained by the delay between the temperature heating / cooling of the sample and of the sample holder or by a kinetic effect. However, at these temperatures, the atomic diffusion is much faster than the temperature ramp (30°C every 10 min). We have checked that this phenomenon is not observed for the bulk sample. It is therefore likely that this hysteresis can be either due to a surface effect influenced by the alumina matrix [START_REF] Takahashi | On low-temperature ordering of FePt films[END_REF] or due to the low thermal conductivity of the alumina insulating layer [START_REF] Langlois | Growth and structural properties of CuAg and CoPt bimetallic nanoparticles[END_REF]. EDX analysis shows that the composition of the particles remains unchanged during the in situ experiment. Further work is underway to clarify this point.

Discussion

To summarize the in situ TEM experiment, we have characterized the morphological and structural evolution of an initially disordered Co-Pt nanoparticle film. The initially disordered particles order at 570°C and coalesce at 710°C inducing a rapid increase of the particle size. We now comment on the temperature dependence of the order phenomena and of the coalescence processes observed in our experiment. First, the ordering of the initially disordered nanoparticles is a kinetic process, requiring atomic jumps both at the surface and in the bulk of the nanoparticles. So, the temperature at which order is observed is not a quantity of thermodynamical equilibrium, and depends on the annealing time. However, because diffusion is a thermally activated process, the diffusion coefficients evolve very strongly with temperature. Therefore, when increasing temperature, the transition between a process 'too slow to be observed in a reasonable time' and a rapid process is usually very fast. That value of the temperature of transition between these two regimes can be estimated by assuming that the mean diffusion length (γ) equals to the first nearest neighbour distance [START_REF] Rellinghaus | On the L1 0 ordering kinetics in Fe-Pt nanoparticles[END_REF]. The mean diffusion length is given by:

γ= t kT E exp D Dt a 0        
where D is the diffusion coefficient, t the time, D 0 the pre-exponential factor, E a the activation energy, k the Boltzmann constant and T the temperature in Kelvin. We have chosen the impurity diffusion of Pt (in Co) since, owing to its size it is the least diffusing constituent of the alloys. Using the values D 0 = 0.65×10 -4 m 2 s -1 and E a = 279.3 kJmol -1 [START_REF] Landolt | Numerical Data and Functional Relationships in Science and Technology Madelung O Berlin[END_REF], and t = 10 minutes as the characteristic annealing time in our experiment, the transition is found close to 550°C. This estimation is in a good agreement with our experiment where an annealing temperature above 570°C is necessary to really observe ordering of the particles. This conclusion is also consistent with the observation of the disordered structure for as-grown Co-Pt nanoparticles when the substrate temperature is well below 600°C [START_REF] Favre | Structural and magnetic properties of CoPt mixed clusters[END_REF][START_REF] Wang | Pt nanoparticles in C matrix[END_REF]. Similarly, the temperature at which the coalescence is observed has a kinetic origin. In addition, we want to stress that this temperature may also slightly differ between in a sample with a very high coverage of particles (as the one used in the in situ experiment) and in a sample containing well-isolated nanoparticles.

Our objective was to use the results of the in situ experiment to optimize the choice of the substrate temperature during the PLD synthesis. Before doing so, it is important to stress the limits of this approach. First of all, during a PLD synthesis, the nanoparticles are formed in a highly out-ofequilibrium condition, which differs from a simple annealing of the nanoparticles. Secondly, the nanoparticles observed in situ are covered with alumina, and this is obviously not the case during the PLD synthesis. We therefore expect that the temperature at which the coalescence is activated inside the PLD apparatus is smaller than the one suggested by the in situ experiment. The temperature range that can be used to obtain small as-grown ordered nanoparticles is thus smaller than the (570°C-710°C) range.

We have selected a substrate temperature of 650°C, roughly in the middle of the temperature range suggested by the in situ experiment. The other parameters of the PLD synthesis are unchanged (laser frequency of 5 Hz and nominal thickness 2 nm). After the synthesis, the as-grown nanoparticles are characterized using TEM. When compared to the sample obtained at 500°C, the particles have a more regular shape, often close to a sphere. EDX analysis, in a nanoprobe mode, has shown that they also have a Co 55 Pt 45 composition. Moreover, the coverage (51%) is lower than the one observed at 500°C indicating that the growth is more three-dimensional. The size distribution of the particles (right inset of figure 4a) is close to a Gaussian with a mean size of 7.6 nm and a standard deviation of 3.3 nm.

Most importantly, the as-grown nanoparticles are ordered in the L1 0 structure as it is shown on the electron diffraction pattern of the nanoparticles (left inset of figure 4a). The same synthesis protocol has been realized with a substrate temperature of 725°C. In this case, the histogram of the particle sizes (right inset of figure 4b) indicates a bimodal size distribution with a first size population from 1.5 to 3.5 nm and a second one centred on 10 nm. Moreover the coverage has been reduced to 42%. The largest particles are polycrystalline suggesting that the first particles formed during the synthesis have coalesced. In this scheme, new nucleation sites become available on the substrate between the coalesced nanoparticules [START_REF] Sondergard | Cheyssac P and Stella A Production of nanostructures by self-organization of liquid Volumer-Weber films[END_REF] making possible the growth of smaller CoPt nanoparticles in the late stage of the synthesis. The electron diffraction (left inset of figure 4b) shows the ordered state of these nanoparticles. Moreover, on both samples, the order is also clearly visible on HREM images of the nanoparticles. Figures 4c and4d

Conclusion

We have synthesized as-grown L1 0 ordered CoPt nanoparticles using a PLD experiment where the substrate temperature was selected on the basis of an in situ annealing experiment in a transmission electron microscope. Moreover, the size of the particles has been shown to depend on the temperature of the substrate. Thus, controlling the temperature and the amount of matter deposited on it makes possible the manufacturing of as-grown L1 0 ordered nanoparticles of definite size, density and composition. Even if the particles size distribution is still not narrow enough to be used in the magnetic applications for EHDR, the control of the structural and morphological parameters are very promising from the point of view of magnetism studies and applications. It becomes therefore possible to directly synthesize film of nanoparticles chemically ordered and magnetically isolated. Additional

  show nanoparticles of 5 nm oriented along the [001] direction and 8 nm oriented along the [112] direction, coming from the samples respectively heated at 650°C and 725°C during the synthesis. The power spectrum of each image clearly shows the superstructure reflections (encircled in insets of figures 4c & 4d) characteristic of the tetragonal L1 0 structure oriented along the [001] and [112] directions.

Figures and figure captions

  Figures and figure captions:

  Figures and figure captions:

Figure 1 .

 1 Figure 1. Equilibrium phase diagram of bulk CoPt.

Figure 2 .

 2 Figure 2. Rectangles: Atomic percent of Cobalt measured by EDX and FCC lattice parameter, determined on electron diffraction patterns, in Co x Pt 1-x nanoparticles samples. Line: Végard Law deduced from the values of the FCC lattice parameter of pure bulk Co and Pt metals.

Figure 3 .

 3 Figure 3. Evolution of a selected area as a function of increasing temperature observed in bright field and electron diffraction modes (diffraction pattern in the inset) (a) 25°C, the diffraction pattern has been indexed according to the FCC structure; (b) 600°C; (c) 740°C, the diffraction pattern has been indexed according to the L1 0 structure; (d) Morphology of the area at 750°C which has remained unchanged during the temperature decrease down to 700°C. Left inset: 750°C during the temperature increase, after the order-disorder phase transition. Right inset: 700°C during the temperature decrease, after the disorder-order phase transition. The arrows show atom inter-diffusion bridges between nanoparticles.

Figure 4 .

 4 Figure 4. CoPt nanoparticles grown at T s = 650°C (a & c), T s = 725°C (b & d). Electron microscopy images in Bright Field mode, corresponding diffraction patterns (left inset) of as-grown ordered CoPt nanoparticles and size distributions of the particles in the right inset (a & b). HRTEM images and power spectra of ordered nanoparticles, (c) projected along the [001] direction, (d) projected along the [112] direction. The ordered reflections of the L1 0 structure are encircled. Images acquired on a JEOL 2010F electron microscope.

work is currently under progress to quantify the degree of order as a function of the particles size, and to study their magnetic properties.