Direct gas-phase formation of complex core–shell and three-layer Mn–Bi nanoparticles

Abstract

We report the formation of complex core–shell and three-layer Mn–Bi nanoparticles in a single step inert-gas condensation process. These structures have been achieved by controlling the thermal environment of the nanoparticles. High resolution transmission electron microscopy, high-angle annular dark-field imaging in scanning transmission electron microscopy mode, and elemental mapping by energy dispersive spectroscopy have been used to determine the crystal structure and chemical composition of the nanoparticles. These particles exist in two forms: (1) a crystalline Bi core with an amorphous Mn-rich shell, and (2) a crystalline Bi annular shell between two amorphous layers with high Mn concentration. These particles show significant magnetic hysteresis possibly arising from the change in bond length between Mn atoms introduced by Bi atoms in the bonding environment of the Mn atoms.

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1 Introduction

The formation of novel nanostructures is critical to understand the complex nucleation and growth mechanisms at the nanoscale and to explore the possibility of using these engineered structures for potential applications. We use inert-gas condensation (IGC) to form A–B-type core–shell and A–B–A-type three-layer structure in the Mn–Bi nanosystem. In IGC we utilize experimental parameters such as inert-gas flow rate, sputtering power, and condensation temperature to control the nucleation and growth of nanoparticles. This is a single step process where the desired nanostructures are formed by manipulating the thermal environment of the nanoparticles. This is unlike the previous works^1–3 on multi layer heterostructured nanoparticles that utilized either successive reduction of chemical precursors to form a core first and a shell over the core or a nanoscale Kirkendall effect to form hollow-core heterogeneous nanostructures.^4,5 In fact, the only three-layer structure that has been reported in the Au–Pd system uses successive reduction of precursors to form the structures.⁶ Here, we have been able to control the diffusion of Bi atoms in Mn core layers in a direct singlestep process to form complex core–shell and three-layer Mn–Bi nanostructures.

The Mn–Bi bimetallic system has a wide miscibility gap.⁷ Apart from a hexagonal MnBi line compound at 1 : 1 stoichiometry and its high temperature variant, almost the entire phase diagram consists of phase mixtures due to the immiscibility between Mn and Bi. The hexagonal MnBi phase is highly anisotropic and has shown excellent hard magnetic properties.^8,9 Significant attention is being paid to study the hexagonal MnBi phase and its magnetic properties,^10–13 but there is little information about the phase-segregated regions on both sides of the MnBi hexagonal phase. This is surprising given the fact that there are reports of solid solubility extension^14,15 and stabilization of unique phases^16,17 at the nanoscale in other immiscible bimetallic systems. Recently, with the report of the formation of core–shell Fe–Bi particles,¹⁸ there is interest in understanding the complex structures and their growth mechanism in Bi–M (M = transition metal) systems. In this report, we show that in spite of the equilibrium thermodynamics principles, the complex core–shell and three-layer structures are formed at the nanoscale. Also, these particles show significant magnetic hysteresis behavior with a saturation magnetization of 60 emu cm⁻³.

2 Experimental methods: synthesis and characterizations

Mn–Bi nanoparticles were fabricated by inert-gas condensation.¹⁹ An alloy target with 50 at% each of Mn and Bi was used for sputtering at a base pressure of 10⁻⁷ Torr. A high power of 150 W was used for sputtering as it produces high-temperature annealing conditions and results in near-equilibrium structures.²⁰ A mixture of Ar and He gas was used for sputtering as well as a carrier gas. The nucleation chamber was water cooled to condense the sputtered atoms to form nanostructures. The condensed particles were then carried to the deposition chamber by the Ar/He mixture gas by differential pressure. In the deposition chamber, the particles were deposited directly on cleaned Si substrates for magnetic, X-ray measurements and on TEM grids for structural and chemical characterization.

The high-resolution transmission electron microscopy (HRTEM), high-angle annular dark field-scanning transmission electron microscopy (HAADF-STEM), and energy dispersive spectroscopy (EDS) measurements were performed using a FEI Tecnai Osiris® microscope. The microscope is fitted with four X-ray detectors for energy dispersive X-ray mapping in STEM mode. This results in high quality elemental maps that can be superimposed on the HAADF images. The false coloring and superposition of individual elemental maps were carried out in Bruker Espirit® software. The quantification of EDS spectra was performed using Cliff–Lorimer thin film approximation. The quantitative elemental maps were then used to compute the composition of the individual layers in the nanoparticle. The mapping of the scanned area enables us to derive ED X-ray spectrum from any desired part of the particle, provided sufficient signal has been detected from that part. Although there is a provision of measuring EDS along a line, but this mode is not drift-corrected. So, in the present work, to nullify effects of small sample drift we have collected all spectrum in area mode. The area scan mode has a specimen stabilizer as opposed to the line scan mode and maps produce results with less spatial error. Furthermore, from the quantified maps taken in area mode, results can be compared from different points along a line. This effectively reproduces a line scan. The quantification results presented in the main text are average over several small area scans within a particle. The image analysis for structural characterization was carried out using ImageJ®, Crystal Maker®, and Crystal Diffract® softwares.

For the magnetic measurements, the Mn–Bi nanoparticles were deposited directly on a Si substrate. The thickness of the cluster film, as measured by a quartz-crystal thickness monitor, was 30 nm. The as-deposited films were covered with a SiO₂ film using a second RF gun to prevent oxidation. Two types of magnetic measurements were performed using a Quantum Design Magnetic Property Measurement System (MPMS) superconducting quantum interference device (SQUID) magnetometer with a maximum field of 7 T. The first one was a measurement of magnetization (M) as a function of applied field (H) at 10 K and 300 K. The magnetic signal coming from the diamagnetic Si substrate and SiO₂ was subtracted from the sample signal by fitting a straight line to the high-field region and subtracting the linear portion from the measured signal. The saturation magnetizations were determined by plotting M vs. 1/H² in the high-field regions and extrapolating to 1/H². The measurement of magnetization (M) as a function of temperature (T) was done in field-cooled (FC) and zero-field-cooled (ZFC) modes in a second set of experiments. This gives us information about changes in magnetic phases and blocking temperature in case the particles are superparamagnetic.

The X-ray diffraction patterns were taken in a Rigaku® D/Max-B diffractometer operating at 40 kV and 40 mA. It uses a Co-kα radiation (λ = 0.1789190 nm) for diffraction in a Bragg–Brentano θ–2θ geometry. The obtained diffraction pattern was matched with a standard powder diffraction pattern of rhombohedral Bi obtained from the database of International Center for Diffraction data (ICDD®).

3 Results and discussions

The average size of the as-deposited Mn–Bi nanoparticles is 10.5 ± 2.8 nm (Fig. 1a). The crystal structure of the particles was determined by X-ray diffraction (XRD). The XRD pattern is indexed to a rhombohedral structure with a lattice parameter a = 0.453 nm and c = 1.18 nm (Fig. 1b). The lattice parameters are very close to those reported for pure rhombohedral Bi. There is also a small peak at 2θ = 65° which is coincident with a peak characteristic of bcc-Bi. The crystal structures of individual nanoparticles of three different sizes were also determined from the HRTEM images and are shown in Fig. 2a–c. All of them are single crystalline and the fast Fourier transform (FFT) of these high-resolution images are indexed to the same rhombohedral structure as determined from the XRD pattern (Fig. 2d–f). No traces of Mn in crystalline form can be found in them. A case of special interest is Fig. 2c, which shows that the particle with largest diameter (among the three particles) has its crystalline part spread in an annular shell between two amorphous layers. To avoid any confusion that may arise from microscope defocus, the image was taken in three defocus conditions near the Scherzer defocus (ESI Fig. 1†). The shape of the annular shell remains the same with a small amount of shift of the core amorphous region. It has been observed in the present study that the three-layer structures are less common than the double layer core–shell structures and can only be found in particles with a size >20 nm. The HRTEM images (Fig. 2a and b) of particles with sizes 8.8 nm and 12.5 nm show no signs of an amorphous core within the particles. Although the high-resolution images can clearly distinguish between a regular core–shell and a three-layer structure with an amorphous core (ESI Fig. 2†), it is hard to determine if these structures are indeed arising from different chemical compositions or due to microscope conditions (aberration, orientation of the particle with respect to zone axis, and defocus). The problem is further compounded by the fact that the Mn-rich regions are amorphous. To determine the chemical nature of individual layers, the nanoparticles were further investigated in HAADF-STEM mode, which is capable of showing atomic number contrast. Furthermore, X-ray mapping in STEM mode provides chemical information from within the nanoparticle, particularly if there is chemical heterogeneity.

Fig. 1 Size distribution and phase identification of Mn–Bi nanoparticles. The average size of the Mn–Bi nanoparticles measured by bright field TEM images is 10.5 ± 2.8 nm. The X-ray diffraction is indexed to a rhombohedral Bi phase with a = 0.453 nm and c = 1.18 nm. There is a small peak at 2θ = 65° that can be indexed to a bcc Bi phase.

Fig. 2 High-resolution TEM images of Mn–Bi nanoparticles. (a–c) The high-resolution images of Mn–Bi nanoparticles with sizes 8.9 nm, 12.5 nm, and 20.6 nm, respectively. (d–f) The fast Fourier transformations show particles belong to a crystalline Bi phase. The particle at (c) shows crystalline Bi phase exists in the annular shell only.

Fig. 3 shows HAADF-STEM images of three nanoparticles. Each of these nanoparticles shows different contrasts at the core and the shell, indicating that these two regions consist of predominantly two different kinds of atoms. Further, the particles at the top and at the bottom in Fig. 3 show a different contrast near the core (marked with arrow) similar to that of the shell (see ESI Section 2† for more detailed analysis of this image). The elemental mapping of particles using EDS decisively determines if the contrast is arising from different elements. Fig. 4 shows maps of individual elements and their superposition in two nanoparticles along with their HAADF-STEM images, respectively. Fig. 4a–d show that there are three layers in the nanoparticle. There is a Mn-rich core and shell, and between them there is a Bi-rich annular shell. Quantitative analysis of EDS spectrum (Fig. 5) from individual layers show that the compositions of Mn at the core, inner shell, and outer shell are 33 at%, 9 at%, and 67 at%, respectively. Fig. 4e–h show that the particle consists of predominantly Bi core and a Mn shell. The quantification (Fig. 6) results show that the Mn composition at the core is 11 at% and that at the shell is 68 at%. The EDS quantification results are based on 15 measurements and the maximum error of 6 at% was observed in case of Mn with an average value of 68 at%. For all other cases the error was below 2 at%. The marked change in intensities in Mn-K lines and Bi-L lines in the EDS spectra for different layers are visible in Fig. 5 and 6. The compositional analysis show that each layer (for both kinds of structures) consists of Mn and Bi atoms. The observation is similar to that found in the three-layer Au–Pd structures.⁶

Fig. 3 High-angle annular dark-field images of Mn–Bi nanoparticles. The images show the formation of core–shell structures in these particles with distinguishable contrast between the core and the shell. In the top and the bottom particles there are some regions inside the particles (marked with arrow) that have the same contrast as that of the outer shell of the nanoparticles. This indicates the formation of three-layer A–B–A-type structure.

Fig. 4 X-ray maps of individual nanoparticles. (a–d) The individual HAADF images with false-color maps of each elements show that the Mn atoms are present predominantly at the core and the surface. This is a three-layer structure. (e–h) The similar X-ray maps show that the particle core consists mainly of Bi atoms and the shell of Mn atoms. This is regular core–shell structure.

Fig. 5 The energy dispersive spectra of a three-layer nanoparticle. The spectra were taken from three different points along a line. The Mn-K lines and Bi-L lines have different intensities at three different layers. Compositions are presented in the text.

Fig. 6 The energy dispersive spectra of a core–shell nanoparticle. The spectra were taken from two different points along a line. The Mn-K lines and Bi-L lines have different intensities at the core and shell. Compositions are presented in the text.

The nucleation and growth of NPs in IGC, as mentioned earlier, are dependent on three main experimental parameters: inert gas-flow rate, sputtering energy, and condensation temperature. The use of a high power (150 Watt) increases the sputtering rate and generates a greater metal flux near the sputtering target. An increase in metal flux (vapor density of metallic atoms) reduces the cooling efficiency of the inert gas, so the nucleation rate is lower than the further growth of such clusters into nanoparticles.²¹ A higher gas-flow rate decreases the time the clusters spend in the growth region of high metal vapor density. It also shortens the growth distance by maximizing metal vapor density very close to the target.²¹ In the present experiment an optimum balance between power and gas flow rate was maintained so that a sufficient number of nanoparticles with a size around 10 nm were formed. We have used water cooling as opposed to liquid N₂ cooling in the condensation chamber so that atomic diffusion in the condensed particles is sustained to form crystalline structures.

The observed phenomenon of phase separation matches well with the equilibrium phase diagram. The Mn–Bi system has positive enthalpy of mixing (ΔH^mix)²² and the difference in surface free energies between Mn and Bi are also large.²³ At smaller sizes ΔH^mix decreases and as a consequence solid solubility increases. We are in a size range (≈10 nm) where ΔH^mix is not enough to overcome surface effects and core–shell particles are formed. Similar surface energy-driven core–shell formation has been reported in other bimetallic systems.²⁴ At larger particle sizes, the Mn–Bi system tends to follow equilibrium thermodynamics-i.e., phases form which are commonly observed in bulk systems. A closer look at the formation of low temperature MnBi phase (LT-MnBi) from the high temperature MnBi (HT-MnBi) show that the reaction is accompanied by Mn precipitation.⁷ In the present case, the Mn precipitates are further mixing with Bi to form Mn-rich alloy shell in the core–shell structure and first and third layers of the three-layer structure. The formation of a Bi-rich core and a Mn-rich shell and the formation of three-layered structure, however, are contrary to that expected of equilibrium thermodynamics. A core–shell structure with Mn as core would be more favorable as Bi would reside on the surface owing to its lower surface energy and larger atomic size compared to Mn.²³ Also, the distribution of Mn simultaneously at the core and the outermost shell in the A–B–A structure cannot be explained by surface energy difference. There must be an inverse migration of Mn or Bi atoms to form these kinds of structures. As these nanoparticles are formed in a nonequilibrium process, the phases with nonequilibrium solid solubility cannot be ruled out.

When encountered with the formation of a similar kind of complex structures in Ag–M (M = Pd, Cu, Ni) systems, ​Baletto et al.²⁵ used nonequilibrium molecular dynamics simulations to model the inert gas condensation process. The assumption of their model is that the core layer (of A atoms) forms initially and depending on the morphology of the core layer and temperature of the condensation chamber, the second kind of atoms (B atoms) are incorporated below the A layer. The morphology of the core layer is critical to the incorporation process as the simulation depicts that if the core is a truncated octahedron (TO), the most favorable site for B atoms (to nucleate and grow) are one layer below the surface, especially below the edges and vertices. An A–B–A-type three-layer structure is formed at the end of the growth process. On the other hand, if the core layer is an icosahedron (IH), the B atoms preferred to diffuse to the center of the core (as opposed to the subsurface) and grow. The core–shell structure formed here has a core of B atoms and a shell of A atoms, although the initial core was made of A atoms. Although in the present case the core layers are spherical, locally IH or TO structures can form resulting in these unusual structures.

These core–shell and three-layer nanoparticles show significant magnetic hysteresis behavior at room temperature and at 10 K (Fig. 7). The saturation magnetization of these Mn–Bi cluster films is approximately 60 emu cm⁻³. The M vs. T curves in field-cooled and zero-field-cooled conditions show no significant loss in magnetization up to 300 K. This indicates that the particles are ferromagnetic. However, the particles have very small (≈20 Oe at 10 K) coercivity both at 300 K and 10 K. Given the fact that most of these particles are single crystalline Bi with a Mn rich coating on them, the ferromagnetic behavior is unusual. Recently, however, Wei et al. have shown that antiferromagnetic MnAu nanoparticles show ferromagnetic behavior when coated with a thin layer of Mn.²⁶ Also, small Mn_NBi_M alloy clusters show ferromagnetic behavior when the ratio of M to N is close to 2.²⁷ This study indicates that, at this composition, the separation between Mn atoms due to the presence of Bi atoms in the lattice is optimal for ferromagnetic behavior. The reason for ferromagnetism in these clusters is derived from the change in local bonding environment of Mn atoms either through covalent bonding that changes the local magnetic moment of Mn or through a change in interatomic distance of Mn as described earlier. In our case there can be two reasons for a change in the local environment of Mn: (1) at the interface between crystalline Bi and amorphous Mn, in some atomic layers the bonding of Mn–Bi results in ferromagnetism, and (2) there is some amount of interdiffusion of Mn and Bi the opposite layers. The latter case (2) is supported by the quantification of X-ray maps, which shows that there is a significant amount of Mn in the core Bi-rich layer of the core–shell structure.

Fig. 7 Magnetic measurements of Mn–Bi nanoparticles. (a) The M vs. H curves at 300 K and 10 K show the hysteresis behavior of Mn–Bi particles. The saturation magnetization is approximately 60 emu cm⁻³ at both 10 K and 300 K. The inset shows second quadrant behavior. (b) The magnetization in field cooled and zero filled curves as a function of temperature follow the same path. There is no significant decrease in the magnetization.

4 Conclusions

In summary, ambient stable complex core–shell and three-layer Mn–Bi nanoparticles have been synthesized by inert-gas condensation. These layers were formed in-flight during the deposition process without any extra change of experimental conditions, such as the change in position of the sputtering guns so that one element is preferentially deposited over the other at different points in the chamber, forming core–shell structures. There exists a strong thermodynamic driving force for the formation of these phase-separated structures. The formation of three-layered structures, however, cannot be completely explained by equilibrium thermodynamics. These structures were formed in a nonequilibrium process, so several factors including local inhomogeneity can play a role in their formation. A nonequilibrium model that uses the preferential position of Bi nucleation sites in a Mn-rich core is used to understand how local inhomogeneity can possibly play a role in their formation. The nanoparticles are ferromagnetic. Their magnetism can be attributed to the change in local bonding environments between Mn atoms. To determine the exact reason for the formation of these complex layers and their magnetic properties a first-principle calculation is needed. Nonetheless, the direct one-step formation of such layered nanostructures is a significant result as this is the first experimental result showing complex three-layer nanomagnets.

Acknowledgements

This research was supported by the Army Research Office Grant no. WF911NF-10-2-0099. This research was performed in part in Central Facilities of the Nebraska Center for Materials and Nanoscience, which is supported by the Nebraska Research Initiative. Electron microscopy research was supported by NSF-MRI (DMR-0960110).

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Footnotes

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c6ra15799h

‡ Present address: Materials Science and Engineering, Rutgers University, Piscataway, NJ 08854, USA.

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