nanomaterials

Article Magnetic Phase Coexistence and Hard–Soft Exchange Coupling in FePt Nanocomposite

O. Crisan 1,*, I. Dan 2, P. Palade 1, A. D. Crisan 1, A. Leca 1 and A. Pantelica 3

1 National Institute for Materials Physics, 077125 Magurele, Romania; palade@infim.ro (P.P.); alina.crisan@infim.ro (A.D.C.); aurel.leca@infim.ro (A.L.) 2 R&D Consulting and Services S.R.L., 023761 Bucharest, Romania; [email protected] 3 Horia Hulubei National Institute for Physics and Nuclear Engineering, P.O. Box MG-6, 077125 Magurele, Romania; apantel@ifin.nipne.ro * Correspondence: ocrisan@infim.ro

 Received: 2 July 2020; Accepted: 11 August 2020; Published: 18 August 2020 

Abstract: With the aim of demonstrating phase coexistence of two magnetic phases in an intermediate annealing regime and obtaining highly coercive FePt nanocomposite magnets, two alloys of slightly off-equiatomic composition of a binary Fe-Pt system were prepared by dynamic rotation switching and ball milling. The alloys, with a composition Fe53Pt47 and Fe55Pt45, were subsequently annealed at 400 ◦C and 550 ◦C and structurally and magnetically characterized by means of X-ray diffraction, 57Fe Mössbauer spectrometry and Superconducting Quantum Interference Device (SQUID) magnetometry measurements. Gradual disorder–order phase transformation and temperature-dependent evolution of the phase structure were monitored using X-ray diffraction of synchrotron radiation. It was shown that for annealing temperatures as low as 400 ◦C, a predominant, highly ordered L10 phase is formed in both alloys, coexisting with a cubic L12 soft magnetic FePt phase. The coexistence of the two phases is evidenced through all the investigating techniques that we employed. SQUID magnetometry hysteresis loops of samples annealed at 400 ◦C exhibit inflection points that witness the coexistence of the soft and hard magnetic phases and high values of and are obtained. For the samples annealed at 500 ◦C, the hysteresis loops are continuous, without inflection points, witnessing complete exchange coupling of the hard and soft magnetic phases and further enhancement of the coercive field. Maximum energy products comparable with values of current permanent magnets are found for both samples for annealing temperatures as low as 500 ◦C. These findings demonstrate an interesting method to obtain rare earth-free permanent nanocomposite magnets with hard–soft exchange-coupled magnetic phases.

Keywords: L10 FePt; nanocomposite magnets; mechanical alloying; exchange coupling

1. Introduction During the last decade, a lot of research effort has been dedicated to the quest for new classes of permanent magnets, RE-free and capable of operating at high temperatures. Among the possible solutions, FePt stands as one of the most promising due to significant magnetocrystalline anisotropy, high coercivity and remanence of its ordered tetragonal L10 phase. In its nanoparticulate form, FePt has good catalytic properties, high corrosion resistance and is biocompatible. For these reasons, FePt nanoparticles have been intensively investigated for uses in biomedicine (magnetic hyperthermia [1], MRI imaging [2] and targeted drug delivery [3]) and in catalysis [4] as sensors [5] or biosensors [6]. Various chemical and physical synthesis methods, such as thermal decomposition using high-temperature boiling point solvents [7], polyol [8], microemulsions [9], microwave processing [10], sonochemistry [11], pulsed plasma in liquid [12], or inert gas condensation [13], have been used to obtain

Nanomaterials 2020, 10, 1618; doi:10.3390/nano10081618 www.mdpi.com/journal/nanomaterials Nanomaterials 2020, 10, 1618 2 of 12

FePt nanoparticles with well-defined shape and controllable sizes. Various other reports deal with the specific link between synthesis methods and magnetic performances in a large range of magnetic materials [14–26]. As a building block for various devices, FePt is also of interest. Self-assembled nanostructures realized using FePt nanoparticles onto block copolymers templates [27] are considered for high density magnetic recording in the near future using either the heat-assisted magnetic recording method (HAMR) [28] or as self-organized magnetic array (SOMA) bit-patterned media [29]. On the other hand, as bulk materials, FePt-based alloys are very promising as permanent magnets operating in extreme conditions when taking into account some of their attributes, such as corrosion resistance at high temperatures, increased when with rare earth (RE)-containing permanent 7 3 magnets (477 ◦C), high magnetocrystalline anisotropy (10 MJ/m ) and saturation (1.4 T) and considerable energy product. Of importance for a series of applications, FePt-based alloys also show increased hardness and Young modulus [30]. As chemical routes or vacuum deposition methods do not allow production of a high amount of material needed for certain applications, other techniques can be envisaged. Non-equilibrium metallurgical methods such as melt-spinning [31] or ball-milling [32] are proven to be prolific in obtaining sufficiently high amounts of homogeneous Fe-Pt alloys. Furthermore, there is much more versatility in metallurgical methods since they allow for slight changes in composition and doping with a small amount of other elements in order to promote the formation of the L10 phase directly from the synthesis without the need for additional annealing, or to promote the formation of additional magnetic phases, needed for instance to construct a nanocomposite exchange-spring FePt-based . Mechanical alloying via ball milling is one of these non-equilibrium techniques which allows, for instance, one to obtain industrially significant quantities of nanocomposite powders in a single run. Moreover, further processing of the FePt-based nanocomposites into geometries and shapes (discs; cylinders) required for applications can be conducted by using, e.g., spark plasma sintering [32,33]. Several studies on FePt nanocomposites made by ball milling have been reported. Surfactant-assisted ball milling of Fe-Pt using oleylamine and oleic acid was employed in [34], however, it proved to be somewhat ineffective in obtaining L10 FePt with high coercivity. Lyubina et al. [35] reported synthesis by ball milling of Fe and Pt powders at liquid nitrogen temperature in order to counteract Pt tendency to stick onto balls during milling. However, after annealing, they obtained a mixture of hard magnetic L10 and soft magnetic cubic FePt phases. Hu et al. [36] obtained L10 FePt from ball milling of a molecular complex precursor and the method was deemed suitable for diminishing the formation temperature of the L10 phase down to 450 ◦C. In the present work, we report on the coexistence of hard L10 and soft L12 FePt phases, we follow the evolution with the temperature of the structural phase transition with occurrence of the L10 phase at early stages of annealing by means of temperature-dependent X-ray diffraction of synchrotron radiation, and we show that the coexistence of the two phases, partially exchange-coupled at 400 ◦C, leads to a fully coupled exchange spring magnet at temperatures as low as 500 ◦C.

2. Experiment

Two FePt compounds with compositions close to the equiatomic ratio, Fe53Pt47 and Fe55Pt45, were synthesized using ball milling, starting from constituent powders Pt (99.9+%, 325 mesh, Alfa Aesar, Thermo Fisher GmbH, Kandel, Germany) and Fe (99.9%, iron powder spherical < 10 µm, Alfa Aesar). A Retsch PM400 (Retsch GmbH, Haan, Germany) planetary mill dotted with stainless steel vials of 125 mL was used for the synthesis process. Each mixture of constituent Fe and Pt powders weighed about 3.3 g. During milling, one 20 mm stainless steel ball and eight 10 mm stainless steel balls were loaded in the vials together with the composite powder, ensuring a ball-to-powder mass ratio of 20:1. Milling cycles, 10 min long, were separated by a 3 min stop to avoid the powder overheating and element segregation. The direction of rotation was reversed at each pause. A small amount of hexane (30 mL for each vial) was used to cover the powder and balls in the vials with the aim of diminishing the Pt adherence to the balls and loss of effective powder mass of the final composition. The total effective milling time was 8 h at a rotation frequency of 350 rpm. The powders were loaded to—and afterwards Nanomaterials 2020, 10, 1618 3 of 12

extracted from—vials under a protective argon atmosphere (<1 ppm O2, <1 ppm H2O), and operations were performed in an MBraun Labstar (MBraun Inertgas-Systeme GmbH, Garching, Germany) glove box. With the purpose of following the disorder–order transformation that conveys the formation of the hard magnetic tetragonal L10 phase, the as-milled powders were subsequently annealed under argon flow (99.9999% purity, 100 mL/ min) for 1.5 h at 400 ◦C and 500 ◦C for samples Fe53Pt47 and Fe55Pt45, respectively. The composition of the milled alloys was verified using proton induced X-ray emission spectrometry (PIXE). This analysis was carried out using a 3 MeV proton beam generated at the 3 MV Tandetron accelerator of National Institute for Physics and Nuclear Engneering, NIPNE. The crystalline structure of the composite powders was investigated by powder X-ray diffraction (XRD) using a Bruker D8 Advance (Bruker Corporation, Billerica, MA, USA) diffractometer with Cu Kα radiation in θ–2θ geometry. The evolution with the temperature of the disorder–order phase transformation was monitored with temperature-dependent X-ray diffraction of synchrotron radiation. The experiments were performed between 200 ◦C and 600 ◦C using monochromatic radiation (λ = 0.066 nm) at the X04 SA materials science beamline at the Swiss Light Source (Paul Scherrer Institute, Villigen, Switzerland). The heating during measurements was carefully performed and, upon reaching the set temperature value, an additional 2 s settle time was allowed for temperature stabilization at the desired value within 1 ◦C. The measurement protocol involves the collection of two diffractograms for 4 s exposure at two different detector positions followed by merging the two results into one pattern to remove the geometrical effects. The total exposure time of the sample at a given temperature was about 10 s. Magnetometry measurements were performed in order to characterize the magnetic properties of the samples. Hysteresis loops at 27 ◦C were acquired with a Quantum Design (Quantum Design Europe GmbH, Darmstadt, Germany) SQUID magnetometer running under RSO (reciprocate sample option) mode. The spin structure was investigated using 57Fe Mössbauer spectroscopy performed at room temperature with a linear spectrometer in a conventional set-up in transmission geometry. The gamma source we used was a 57Co source in the Rh matrix.

3. Results and Discussion The chemical composition of the resulted powder was verified using proton induced X-ray emission spectroscopy (PIXE). Proton-induced X-ray emission is an elemental analysis technique which uses a highly energetic beam of heavy charged particles (usually protons of kinetic energy between 1 and 4 MeV) to produce an element-specific X-ray emission from solid samples. Compared to other compositional analysis techniques, for instance, energy-dispersive X-ray spectroscopy, PIXE has 5 7 the advantage of low detection limits of 0.1–10 mg/kg (or 10− –10− ) of the concentration, especially for low Z elements. The composition of the as-milled alloys was found to be very close to nominal composition, within 0.5 at%. The alloys compositions that resulted from PIXE are Fe52.6Pt47.4 and Fe55.3Pt44.7. No oxygen or carbon contamination was found in the as-milled alloys, as proven by structural analysis. Both as-milled samples Fe53Pt47 and Fe55Pt45 were structurally investigated by XRD, and the obtained diffractograms are shown in Figure1. It can be seen that the di ffractograms for both samples are quite similar, as expected, and present several Bragg lines, typical for nanocrystalline materials. Full-profile Rietveld-type analysis was performed and we were able to index these peaks as belonging to the (111), (200), (220), (311) and (222) reflections of the face-centered-cubic A1 FePt phase. The quantitative analysis was conducted by full profile refinement of the patterns using a Rietveld-type fitting software named MAUD (Materials Analysis Using Diffraction) [37]. The fit provided the accurate peak positions and full width at half maximum (FWHM) of the peaks. From these values, we calculated the average grain size and the lattice parameters. The average grain size for the Fe53Pt47 was calculated to be 22 nm, while for Fe55Pt45 the calculated value was quite close: 27 nm. The calculated lattice parameters were also very close to each other: 3.803 Å for Fe53Pt47 and 3.804 Å for Fe55Pt45. Nanomaterials 2019, 9, x FOR PEER REVIEW 4 of 13

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Figure 1. XRD diffractograms for Fe53Pt47 and Fe55Pt45 as-milled powders.

The quantitative analysis was conducted by full profile refinement of the patterns using a Rietveld-type fitting software named MAUD (Materials Analysis Using Diffraction) [37]. The fit provided the accurate peak positions and full width at half maximum (FWHM) of the peaks. From these values, we calculated the average grain size and the lattice parameters. The average grain size for the Fe53Pt47 was calculated to be 22 nm, while for Fe55Pt45 the calculated value was quite close: 27 nm. The calculated lattice parameters were also very close to each other: 3.803 Å for Fe53Pt47 and 3.804 Å for Fe55Pt45. In Figure 2aFigure,b we comparatively 1. XRD diffractograms plotted for the Fe 53XPt-ray47 and diffractograms Fe55Pt45 as-milled of the powders. as-milled samples and Figure 1. XRD diffractograms for Fe53Pt47 and Fe55Pt45 as-milled powders. annealedIn Figure at 4002a,b °C weboth comparatively the Fe53Pt47 and plotted Fe55Pt the45 samples. X-ray di ffItractograms can be observed of the that as-milled many samplesmore Bragg and lines appear after annealing. The MAUD analysis and peak indexation allowed us to unambiguously annealedThe atquantitative 400 ◦C both analysis the Fe53 wasPt47 andconducted Fe55Pt45 bysamples. full profile It can refinement be observed of that the many patterns more using Bragg a assign these supplementary peaks to the face-centered tetragonal L10 FePt phase. Most significant are linesRietveld appear-type after fitting annealing. software The named MAUD MAUD analysis (Material and peaks Analysis indexation Using allowed Diffraction) us to unambiguously [37]. The fit the two (001) and (110) so-called main superlattice peaks whose occurrence is a widely accepted assignprovided these the supplementary accurate peak positions peaks to theand face-centeredfull width at h tetragonalalf maximum L10 FePt(FWHM) phase. of Mostthe peaks. significant From criterion of correct identification of the tetragonal phase. However, the occurrence of the L10 phase is arethese the values, two (001) we calculated and (110) so-calledthe average main grain superlattice size and the peaks lattice whose parameters. occurrence The is average a widely grain accepted size proven also by some other superlattice peaks and the tetragonal splitting of the convoluted Bragg criterionfor the Fe of53Pt correct47 was identificationcalculated to be of the22 nm tetragonal, while for phase. Fe55Pt However,45 the calculated the occurrence value was of quite the L1 close:0 phase 27 peaks. In the cubic A1 lattice, the (200), (220) and (311) Bragg lines are the second most intense peaks isnm. proven The calculated also by some lattice other parameters superlattice were peaks also very and theclose tetragonal to each other: splitting 3.803 of Å the for convoluted Fe53Pt47 and Bragg 3.804 after the main (111) peak. However, due to tetragonal distortion, after the disorder–order transition peaks.Å for Fe In55P thet45.cubic A1 lattice, the (200), (220) and (311) Bragg lines are the second most intense peaks and occurrence of the L10 phase, these Bragg lines split as follows: (200) into (200) and (002); (220) afterIn the Figure main (111)2a,b we peak. comparatively However, due plotted to tetragonal the X-ray distortion, diffractograms after theof the disorder–order as-milled samples transition and into (220) and (202); and (311) into (311) and (113). The more obvious the splitting, the more andannealed occurrence at 400 of °C the both L10 thephase, Fe53Pt these47 and Bragg Fe55 linesPt45 samples. split as follows: It can be (200) observed into (200) that and many (002); more (220) Bragg into significant the phase transformation and the more abundant the formed L10 phase. For better (220)lines appear and (202); after and annealing. (311) into The (311) MAUD and (113). analysis The and more peak obvious indexation the splitting, allowed the us more to unambiguously significant the comparison, all the Bragg lines that are unambiguously attributed to the tetragonal L10 phase are phaseassign transformationthese supplementary and the peaks more to abundantthe face-centered the formed tetragonal L10 phase. L10 FePt For phase. better Most comparison, significant all theare denoted in italics. Braggthe two lines (001) that and are (110) unambiguously so-called main attributed superlattice to the tetragonalpeaks whose L10 occurrencephase are denoted is a widely in italics. accepted criterion of correct identification of the tetragonal phase. However, the occurrence of the L10 phase is proven also by some other superlattice peaks and the tetragonal splitting of the convoluted Bragg peaks. In the cubic A1 lattice, the (200), (220) and (311) Bragg lines are the second most intense peaks after the main (111) peak. However, due to tetragonal distortion, after the disorder–order transition and occurrence of the L10 phase, these Bragg lines split as follows: (200) into (200) and (002); (220) into (220) and (202); and (311) into (311) and (113). The more obvious the splitting, the more significant the phase transformation and the more abundant the formed L10 phase. For better comparison, all the Bragg lines that are unambiguously attributed to the tetragonal L10 phase are denoted in italics.

Figure 2. XRD diffractograms for as cast samples of (a) Fe53Pt47 and (b) Fe55Pt45 powders annealed at 400 ◦C.

Quantitative full-profile analysis was performed on the annealed samples, with the aim of quantizing the amount of the L10 phase formation in the two investigated samples. The fitting results are shown in Figure3 for the Fe 53Pt47sample annealed at 400 ◦C, where both experimental (black dots) and the fitting (blue line) are plotted. The fitting model included two crystal structures: the face-centered-cubic fcc FePt and the tetragonal L10 FePt. The fitting shows good agreement with the experiment, and all major Bragg lines are accounted for. The results of the annealed samples suggest that the powders consist of a nanoscale size-mixed structure of fcc and L10 FePt phases with probably

Nanomaterials 2019, 9, x FOR PEER REVIEW 5 of 13

Figure 2. XRD diffractograms for as cast samples of (a) Fe53Pt47 and (b) Fe55Pt45 powders annealed at 400 °C.

Quantitative full-profile analysis was performed on the annealed samples, with the aim of quantizing the amount of the L10 phase formation in the two investigated samples. The fitting results are shown in Figure 3 for the Fe53Pt47sample annealed at 400 °C, where both experimental (black dots) and the fitting (blue line) are plotted. The fitting model included two crystal structures: the face- centered-cubic fcc FePt and the tetragonal L10 FePt. The fitting shows good agreement with the Nanomaterialsexperiment,2020 and, 10 all, 1618 major Bragg lines are accounted for. The results of the annealed samples suggest5 of 12 that the powders consist of a nanoscale size-mixed structure of fcc and L10 FePt phases with probably well-formed L10 grains separated by less ordered regions of cubic symmetry. Similar results were well-formed L1 grains separated by less ordered regions of cubic symmetry. Similar results were also also obtained for0 the Fe55Pt45sample annealed at 400 °C. The fitting quantitative results are given in obtainedTable 1. for the Fe55Pt45sample annealed at 400 ◦C. The fitting quantitative results are given in Table1.

Figure 3. MAUD full-profile fitting of the diffractogram of sample Fe53Pt47 annealed at 400 ◦C. Figure 3. MAUD full-profile fitting of the diffractogram of sample Fe53Pt47 annealed at 400 °C. Table 1. Lattice parameters and average grain sizes obtained from full-profile MAUD analysis. Table 1. Lattice parameters and average grain sizes obtained from full-profile MAUD analysis. Lattice Parameters Average Grain Sample Phase Amount (%) Sample Phase Lattice Parameters(Å) (Å) Average SizeGrain (nm) Size (nm) Amount (%) Fe53Pt47—as-milled fcc FePt a = 3.8037 ± 0.0002 D=22 ± 3 100 Fe53Pt47—as-milled fcc FePt a = 3.8037 0.0002 D=22 3 100 fcc FePt a = 3.7972 ± 0.0002± D=25 ±± 4 27.2 fcc FePt a = 3.7972 0.0002 D=25 4 27.2 Fe53Pt47—anneal. 400 °C L10 FePt a = 3.8383 ± 0.0001± D=38 ±± 7 72.8 Fe53Pt47—anneal. 400 ◦C L1 FePt a = 3.8383 0.0001 D=38 7 72.8 0 c = 3.7228 ± 0.0016± ± c = 3.7228 0.0016 Fe55Pt45—as-milled fcc FePt a = 3.8045 ± 0.0004± D=27 ± 2 100 Fe Pt —as-milled fcc FePt a = 3.8045 0.0004 D=27 2 100 55 45 fcc FePt a = 3.7925 ± 0.0003± D=32 ±± 3 21.3 Fe55Pt45—anneal. 400 °C L10 FePtfcc FePta = 3.8249a = 3.7925 ± 0.00040.0003 DD=49=32 ± 5 3 21.378.7 ± ± Fe55Pt45—anneal. 400 ◦C L1 FePt a = 3.8249 0.0004 D=49 5 78.7 0 c = 3.7116 ± 0.0013± ± c = 3.7116 0.0013 ± As can be seen, the lattice parameter does not change a lot with the small changes in composition in theAs as can-milled be seen, samples, the lattice and the parameter average doesgrain not size, change calculated a lot with using the the small integral changes breadth in composition method, is inbetween the as-milled 20–30 samples,nm for both and the samples. average The grain annealed size, calculated samples, using which the consist integral of breadth both cubic method, and istetragonal between phases 20–30, nmshow for a slightly both samples. lower lattice The annealedparameter samples, for the cubic which phase consist than ofthe both as-milled cubic case, and tetragonalwhile the tetragonal phases, show phase a slightly lattice parameters lower lattice (a parameterand c) are forvery the similar cubic for phase the thantwo annealed the as-milled samples. case, whileThis indicates the tetragonal that the phase micros latticetructure parameters of the two (a andsamplesc) are is very very similar much alike, for the with two very annealed little influence samples. Thisof the indicates compositional that the changes microstructure that differentiate of the two samplesthe two as is- verymilled much powders. alike, withThe veryL10 phase little influenceis highly ofpredominant the compositional in both changesannealed that samples differentiate (between the 72% two and as-milled 78%). powders. The L10 phase is highly predominantIn order in to both monitor annealed the samples phase structure(between 72% and and the 78%). evolution of the disorder–order phase transformation,In order to a monitortemperature the-dependent phase structure synchrotron and the X-ray evolution diffraction of thestudy disorder–order was undertaken. phase X- transformation,ray diffractograms a temperature-dependent on the as-milled Fe55 synchrotronPt45 sample X-ray were di recordedffraction study after heating was undertaken. the sample X-ray at ditempffractogramseratures ranging on the as-milled from 200 Fe °C55 Ptto45 400sample °C in were 50°C recorded steps and after from heating 400 °C the to sample 600 °C at in temperatures 25 °C steps. rangingFor the first from three 200 recorded◦C to 400 diffractograms,◦C in 50◦C steps at 200 and °C, from 250 400°C and◦C to300 600 °C,◦ onlyC in the 25 ◦peaksC steps. of the For cub theic firstFePt threephase recorded are observed. diffractograms, Starting with at 350 200 °C,◦C, the 250 superla◦C andttice 300 peaks◦C, only of the the L1 peaks0 phase of start the cubicto appear. FePt phaseFor this are temperature, observed. Startingthe L10 phase with 350 abundance◦C, the superlatticewas estimated peaks from of full the- L1profile0 phase fitting start to to be appear. about For10%. this Upon temperature, heating at the 400 L1 °C,0 phase this value abundance increases was even estimated more, fromreaching full-profile about 24%. fitting Starting to be about with 10%. 450 Upon heating at 400 ◦C, this value increases even more, reaching about 24%. Starting with 450 ◦C, the L10 phase becomes predominant, its abundance being larger than that of the parent cubic phase (59:41). The phase transformation towards complete formation of the L10 phase is steady as we increase the temperature of the XRD study, and this transformation is complete at about 575 ◦C. The relative abundance of the cubic and tetragonal phases in the as-milled Fe55Pt45 sample vs. the temperature of the XRD experiment is presented in the Figure4. The tetragonal L1 0 structure remains single-phased in the alloy and is thermally stable up to 600 ◦C. Of particular interest for the stability of the structure and phase evolution with temperature is also the monitoring of the lattice parameters of the tetragonal phase Nanomaterials 2019, 9, x FOR PEER REVIEW 6 of 13

°C, the L10 phase becomes predominant, its abundance being larger than that of the parent cubic phase (59:41). The phase transformation towards complete formation of the L10 phase is steady as we increase the temperature of the XRD study, and this transformation is complete at about 575 °C. The relative abundance of the cubic and tetragonal phases in the as-milled Fe55Pt45 sample vs. the temperature of the XRD experiment is presented in the Figure 4. The tetragonal L10 structure remains

Nanomaterialssingle-phased2020 ,in10 ,the 1618 alloy and is thermally stable up to 600 °C. Of particular interest for the stability6 of 12 of the structure and phase evolution with temperature is also the monitoring of the lattice parameters of the tetragonal phase as a function of temperature. The evolution of the tetragonal lattice parameters as(a aand function c) as revealed of temperature. from the The fitting evolution of XRD of data the tetragonalwith the temperature lattice parameters of up to (a 600and °Cc) asis shown revealed in fromFigure the 5. fitting Both lattice of XRD parameters, data with thea and temperature c, increase ofslowly up to at 600 the◦C same is shown pace inbetween Figure 3505. Both °C and lattice 475 parameters,°C. From 475a and °C toc, increase600 °C, the slowly temperature at the same at which pace between the L10 350phase◦C becomes and 475 ◦ predominantC. From 475 in◦C the to 600sample,◦C, the a parameter temperature start at to which increase the more L10 phaseabruptly becomes while the predominant increase in inc parameter the sample, becomea parameter slower. startAs a to consequence, increase more the abruptly tetragonality while the ratio increase c/a inexhibitsc parameter an abrupt become decrease slower. As from aconsequence, 475°C. This themeasurement tetragonality indicates ratio c/ athatexhibits there anis a abrupt predominant decrease in from-plane 475 expansion◦C. This measurementof the tetragonal indicates lattice from that there475 °C. is aOne predominant can observe in-plane that, between expansion 475 of °C the and tetragonal 600 °C, the lattice total from in-plane 475 ◦ expansionC. One can along observe the that,a = b betweenaxes (Δa 475≈ 0.015C and Å) 600is 5 timesC, the larger total in-plane than the expansion total expansion along the alonga = btheaxes c axis (∆a (Δ0.015c ≈ 0.003 Å) is Å). 5 times As a ◦ ◦ ≈ largerconsequence, than the we total note expansion that the alongthermal the expansionc axis (∆c of the0.003 tetragonally Å). As a consequence, distorted unit we cell note is thatstrongly the ≈ thermalanisotropic, expansion occurring of the essentially tetragonally in-plane. distorted This unit phenomenon cell is strongly has also anisotropic, been observed occurring in ternary essentially L10 in-plane.phases in This Fe-Mn phenomenon-Pt [38] and has is in also agreement been observed with thermal in ternary expansion L10 phases behavior in Fe-Mn-Pt of Fe-Pt [38-Ag] and-B alloys is in agreement[39]. The corresponding with thermal expansion tetragonality behavior ratio of(c/ Fe-Pt-Ag-Ba) is shown alloysalso in [39 Figure]. The 5 corresponding. It can be seen tetragonality that there is ratioan optimal (c/a) is showntetragonality also in ratio Figure at5 about. It can 475 be °C seen where that therethe value is an reaches optimal a tetragonality maximum of ratio about at about0.9735. 475Beyond◦C where this temperature, the value reaches the tetragonality a maximum ofratio about decreases 0.9735. to Beyond about this0.97 temperature, at 600 °C, due the to tetragonality the in-plane ratiothermal decreases expansion to about of th 0.97e unit at cell. 600 ◦ However,C, due to thethe in-planevariation thermal in the tetragonality expansion of parameter the unit cell. only However, reaches the3.5% variation over the in thewhole tetragonality temperature parameter range, only and reachesa value 3.5% of 0.97 over tetragonality the whole temperature ratio stands range, for a and high a valuedegree of of 0.97 tetragonal tetragonality ordering ratio of stands the L1 for0 phases a high with degree important of tetragonal consequences ordering on of the overall L10 phases magnetic with importantbehavior of consequences the alloy. The on c the/a ratio overall defines magnetic the degree behavior of tetragonal of the alloy. distortion The c/a ratio of the defines atomic the structure, degree ofthus, tetragonal it is a valuable distortion parameter of the atomic of interest, structure, as it thus, is correlated it is a valuable to the parameter degree of magnetocrystallin of interest, as it ise correlatedanisotropy to (MCA) the degree of the of L magnetocrystalline10 phase [40]. anisotropy (MCA) of the L10 phase [40].

Figure 4. The relative ratio of the cubic and tetragonal phase abundance as a function of temperature, Figure 4. The relative ratio of the cubic and tetragonal phase abundance as a function of temperature, as obtained from the full-profile fitting of the synchrotron X-ray diffractograms. as obtained from the full-profile fitting of the synchrotron X-ray diffractograms. Another parameter of interest for defining the high degree of ordering of the tetragonal phase is represented by the crystallographically coherent (or coherence) length, assimilated to the average grain size of the tetragonal phase. To estimate the coherence length, we employed the integral breadth method [41]. The method calculates the volume-averaged coherence length Ghkl from two parameters issued from the full-profile fitting: (a) the linewidth FWHM of superlattice L10 Bragg peaks belonging to (hkl) = (001), (110), (111) and (200) planes and (b) the Bragg peak asymmetry, as revealed by the mixing parameter (or the proportion of Lorentz-type function in the overall pseudo-Voigt profile function) of each Bragg line. The coherence length was calculated then by averaging the Ghkl values of all superlattice reflections for the samples annealed at temperatures from 450 ◦C up to 600 ◦C, i.e., for temperatures where the tetragonal phase is predominant, as seen in Figure4. It was obtained that the coherence length increases from 63 nm for the sample annealed during the synchrotron XRD study at 450 ◦C to about 75 nm for the sample annealed during the synchrotron XRD study at 600 ◦C. Nanomaterials 2019, 9, x FOR PEER REVIEW 7 of 13 Nanomaterials 2020, 10, 1618 7 of 12

Figure 5. Tetragonal lattice parameters (a and c) and the c/a parameter as a function of temperature Figureobtained 5. Tetragonal from the full-profilelattice parameters fitting of (a the and synchrotron c) and the X-rayc/a parameter diffractograms. as a function of temperature obtained from the full-profile fitting of the synchrotron X-ray diffractograms. Additional information regarding structural features of the annealed FePt samples was furnished byAnother means of parameter Mössbauer of spectroscopy.interest for defining The Fe the55Pt high45powder degree sampleof ordering annealed of the attetragonal 400 ◦C for phase 1 h wasis 57 representedmeasured usingby the crystalFe Mössbauerlographically spectroscopy. coherent The(or spectrumcoherence) was length, recorded assimilated at 27 ◦C to in the transmission average geometry in a velocity range of 12/+12 mm/s, and it is shown in Figure6. The transmission pattern grain size of the tetragonal phase. To− estimate the coherence length, we employed the integral breadth methodreveals [41 the]. The shape method of a convoluted calculates the magnetic volume sextet,-averaged typical coherence for ferromagnetic length Ghkl fro materials.m two parameters Taking into issuedaccount from the the large full- absorptionprofile fitting: factor (a) forthe X-rays, linewidth exhibited FWHM by of heavy superlattice metals L1 such0 Bragg as Pt, peaks and also belonging the large to Pt(hkl) content = (001), in the (110), sample, (111) the and magnitude (200) planes of the and measured (b) the absorptionBragg peak (Mössbauer asymmetry, eff asect) revealed was surprisingly by the mixingof a su parameterfficient standard (or the toproportion accurately of fit Lorentz the spectrum.-type function In agreement in the withoverall the pseudo structural-Voigt findings, profile we function)assumed of that each there Bragg are line. two unequivalentThe coherence Fe length ions in was the chemical calculated environment, then by averaging one with the cubic Ghkl symmetry, values of attributableall superlattice to the reflections residual cubicfor the FePt samples phase, annealed and the otherat temperatures one with tetragonal from 450 symmetry,°C up to 600 attributable °C, i.e., forto temperatures the ordered tetragonal where the FePttetragonal phase. phase Satisfactory is predominant, fitting with as two seen individual in Figure sextets4. It was provided obtained us tha witht thethe coherence hyperfine length parameters increases of the from two 63 sublattices: nm for the isomer sample shift annealed(IS) = 0.29(2) during mm the/s (measured synchrotron relative XRD to studymetallic at 450 Fe), °C quadrupole to about 75 splittingnm for the (QS) sample= 0.23(3) annealed mm/s during and hyperfine the synchrotron field (HF) XRD= 27.8(1) study T atfor 600 the °C. first componentAdditional (I) and information IS = 0.21(4) regarding mm/s, QS structural= 0.01(2) mm features/s and HF of = the30.2(1) annealed T for the FePt second samples component was furnished(II). Taking by means into account of Mössbauer the large spectroscopy. value of the quadrupole The Fe55Pt45 splitting,powder sample which isannealed typical for at 400 a tetragonally °C for 1 h distorted was measur cubiced Fe using environment, 57Fe Mössbauer as well spectroscopy. as the value of The the spectrum hyperfine was field, recorded the first at component 27 °C in is transmissionattributed to geometry the L10 phase. in a velocity The near-zero range quadrupole of −12/+12 mm/s, splitting and of the it is second shown component in Figure is 6. typical The transmissionfor a cubic Fepattern environment reveals the and, shape taking of also a convoluted into account magnetic the larger sextet, hyperfine typical field for ferr thanomagnetic in the first materials.component, Taking we into can concludeaccount the that large the absorption second component factor for is X attributed-rays, exhibited to the by fcc heavy FePt phase. metals In such what asfollows, Pt, and we also will the show large that Pt the content coexistence in the of cubicsample, and the tetragonal magnitude phases of obtain the measured more fundament absorption from (Mössbauerthe magnetic effect) measurements. was surprisingly of a sufficient standard to accurately fit the spectrum. In agreementThe initialwith the magnetization structural findings, and hysteresis we assumed loop of that the samplethere are Fe 55twoPt45 unequivalentannealed at 400Fe ions◦C (Figure in the7 ) chemicalwere recorded environment, using aone Quantum with cubic Design symmetry, SQUID attributable magnetometer to the with residual the applied cubic field FePt parallel phase, and to the thesample other plane. one w Theith magnetic tetragonal field symmetry, was applied attributable up to 50 kOe to the and orderedmeasurements tetragonal were FePttaken phase. at 27 ◦ C. SatisfactoryInitial magnetization fitting with showstwo individual quite slow sextets approach provided to saturation, us with the and hyperfine an inflection parameters point of is the observed two sublattices:on the magnetization isomer shift (IS) curve = 0.29(2) at about mm/s 6 kOe. (measured The curve relative continues to metallic its slow Fe), approachquadrupole to saturation;splitting (QS)however, = 0.23(3) this mm/s is not and reached hyperfine up to field the maximum(HF) = 27.8(1) applied T for field the offirst 50 component kOe. The hysteresis (I) and IS loop = 0.21(4 shows) mm/s,considerably QS = 0.01(2) high mm/s remanent and HF magnetization = 30.2(1) T for of aboutthe second 65% ofcomponent the maximum (II). Taking magnetization into account value the and, largemore value interestingly, of the quadrupole shows a considerably splitting, which high value is typical of coercivity for a tetragonally (Hc = 6.67 kOe). distorted Such cubic values Fe are environment,consistent with as well the predominanceas the value of ofthe a hardhyperfine magnetic field, phase the first in the component sample and is attributed witness the to formation the L10 phase.of the The L10 nearphase-zero at quite quadrupole low temperatures splitting of of thephase second transformation component (400 is ◦ typicalC). It is for to be a cubic noted Fe that environmentusually in FePt and, alloys, taking temperatures also into account higher the than larger 550 hyperfine◦C are needed field inthan order in the to promotefirst component, the formation we canof conclude the L10 phase, that the and second this value component is even is higher attributed in the to case the fcc of FePtFePt nanoparticles.phase. In what Tofollows, our knowledge, we will

Nanomaterials 2020, 10, 1618 8 of 12

incipient formation of the L10 phase at temperatures lower than 400 ◦C (around 300 ◦C) has only been reported by Wang and Barmak [42] in ultrathin FePt films co-sputtered on hot substrates. Two more inflectionNanomaterials points, 2019, signaled9, x FOR PEER by arrowsREVIEW on the graph, are observed on the hysteresis loop. These inflection8 of 13 points indicate the presence of a second magnetic phase, with significantly lower coercivity but high magnetization.show that the coexistence This second of magnetic cubic and phase tetragonal is clearly phases softer obtain than L1 more0 and fundament all clues point from to the the magnetic fact that itmeasurements. is indeed the fcc FePt phase that was identified both in XRD and the Mössbauer spectrum.

Nanomaterials 2019, 9, x FOR PEER REVIEW 9 of 13 57 Figure 6. Fe Mössbauer spectrum of the Fe55Pt45 sample annealed at 400 ◦C. Figure 6. 57Fe Mössbauer spectrum of the Fe55Pt45 sample annealed at 400 °C.

The initial magnetization and hysteresis loop of the sample Fe55Pt45 annealed at 400 °C (Figure 7) were recorded using a Quantum Design SQUID magnetometer with the applied field parallel to the sample plane. The magnetic field was applied up to 50 kOe and measurements were taken at 27 °C. Initial magnetization shows quite slow approach to saturation, and an inflection point is observed on the magnetization curve at about 6 kOe. The curve continues its slow approach to saturation; however, this is not reached up to the maximum applied field of 50 kOe. The hysteresis loop shows considerably high remanent magnetization of about 65% of the maximum magnetization value and, more interestingly, shows a considerably high value of coercivity (Hc = 6.67 kOe). Such values are consistent with the predominance of a hard magnetic phase in the sample and witness the formation of the L10 phase at quite low temperatures of phase transformation (400 °C). It is to be noted that usually in FePt alloys, temperatures higher than 550 °C are needed in order to promote the formation of the L1Figure0 phase, 7. Initial and magnetizationthis value is even and hysteresis higher in loop the ofcase the of sample FePt Fenanoparticles.55Pt45 annealed To at our 400 ◦knowledge,C. Figure 7. Initial magnetization and hysteresis loop of the sample Fe55Pt45 annealed at 400 °C. incipient formation of the L10 phase at temperatures lower than 400 °C (around 300 °C) has only been reportedWhile by the Wang coexistence and Barmak of the two[42]magnetic in ultrathin phases FePt was films confirmed co-sputtered by complementary on hot substrates. investigations Two more While the coexistence of the two magnetic phases was confirmed by complementary (XRD,inflection MS, SQUID)points, signaled in the Fe by55Pt arrows45 sample on annealed the graph, at are 400 observed◦C, these ontwo the phases loop. are These not investigations (XRD, MS, SQUID) in the Fe55Pt45 sample annealed at 400 °C, these two magnetic exchangeinflection coupled, points indicate as proven the presence by the existence of a second of magnetic the inflection phase, points with insignificantly the initial lower magnetization coercivity phases are not exchange coupled, as proven by the existence of the inflection points in the initial andbut thehigh hysteresis magnetization. loop. In This contrast, second the magnetic inflection phase points is clearly in the hysteresis softer than loops L10 and allow all in clues some point cases to magnetization and the hysteresis loop. In contrast, the inflection points in the hysteresis loops allow determinationthe fact that ofit ismagnetic indeed parameters the fcc FePt of phase each individual that was magneticidentified sublattice. both in XRD Their and occurrence the Mössbauer means in some cases determination of magnetic parameters of each individual magnetic sublattice. Their thatspectrum. the diff erent magnetic components—in our case the tetragonal (hard) and cubic (soft) magnetic phases—areoccurrence notmeans completely that the exchange different coupled, magnetic and components annealing— atin a higher our case temperature the tetragonal might ( behard needed) and forcubic achieving (soft) magnetic full exchange phases coupling—are of not the completely two phases. exchange For this purpose, coupled we, and measured annealing hysteresis at a higher loops temperature might be needed for achieving full exchange coupling of the two phases. For this of the Fe55Pt45 and Fe53Pt47 samples annealed at 500 ◦C. The two loops, recorded in same conditions purpose, we measured hysteresis loops of the Fe55Pt45 and Fe53Pt47 samples annealed at 500 °C. The as for the sample annealed at 400 ◦C (Quantum Design SQUID with a magnetic field of up to 50 kOe two loops, recorded in same conditions as for the sample annealed at 400 °C (Quantum Design applied parallel to the sample plane at 27 ◦C), are shown in Figure8. Contrary to the case of the SQUID with a magnetic field of up to 50 kOe applied parallel to the sample plane at 27 °C), are shown sample annealed at 400 ◦C, here, the hysteresis loops of both samples show no inflection points and the in Figure 8. Contrary to the case of the sample annealed at 400 °C, here, the hysteresis loops of both samples show no inflection points and the allure of the hysteresis loop resembles what a single magnetic phase would look like. As in Figure 7, the curves exhibit a considerably slow approach to saturation with no inflection points. The saturation has not been reached for neither samples up to the maximum applied field of 50 kOe. The hysteresis loops of both Fe55Pt45 and Fe53Pt47 samples annealed at 500 °C show an even larger value of coercivity than the Fe55Pt45 sample annealed at 400 °C. While the values of coercive field are not very different between the two samples, there is a large difference in the maximum magnetization values (assimilated to Ms) obtained at 50 kOe. The magnetic parameters recorded for the Fe55Pt45 sample annealed at 500 °C were Hc = 7.05 kOe and Ms = 119 emu/g. The same parameters recorded for Fe53Pt47 sample annealed at 500 °C were Hc = 8.21 kOe and Ms = 97 emu/g. The diminished magnetization at 50 kOe for the Fe53Pt47 sample annealed at 500°C can be explained by the smaller amount of the ordered tetragonal L10 phase, which was already formed at 400 °C (see Table 1). After annealing at 400 °C, there was a 6% difference between the abundances of the already-formed L10 phase in the two samples; thus, it is quite probable that this trend was preserved also after annealing at 500 °C. Between the two alloys with a slightly different composition, a delay in the incipiency of the tetragonal phase for the sample with the lower Fe content was observed. Synchrotron XRD studies proved that the amount of the L10 phase formed in Fe53Pt47 sample at 500 °C annealing is about 5% lower than the amount of L10 already formed in the same annealing conditions for the Fe55Pt45 sample.

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allure of the hysteresis loop resembles what a single magnetic phase would look like. As in Figure7, the curves exhibit a considerably slow approach to saturation with no inflection points. The saturation has not been reached for neither samples up to the maximum applied field of 50 kOe. The hysteresis loops of both Fe55Pt45 and Fe53Pt47 samples annealed at 500 ◦C show an even larger value of coercivity than the Fe55Pt45 sample annealed at 400 ◦C. While the values of coercive field are not very different between the two samples, there is a large difference in the maximum magnetization values (assimilated to Ms) obtained at 50 kOe. The magnetic parameters recorded for the Fe55Pt45 sample annealed at 500 ◦C were Hc = 7.05 kOe and Ms = 119 emu/g. The same parameters recorded for Fe53Pt47 sample annealed at 500 ◦C were Hc = 8.21 kOe and Ms = 97 emu/g. The diminished magnetization at 50 kOe for the Fe53Pt47 sample annealed at 500◦C can be explained by the smaller amount of the ordered tetragonal L10 phase, which was already formed at 400 ◦C (see Table1). After annealing at 400 ◦C, there was a 6% difference between the abundances of the already-formed L10 phase in the two samples; thus, it is quite probable that this trend was preserved also after annealing at 500 ◦C. Between the two alloys with a slightly different composition, a delay in the incipiency of the tetragonal phase for the sample with the lower Fe content was observed. Synchrotron XRD studies proved that the amount of Nanomaterials 2019, 9, x FOR PEER REVIEW 10 of 13 the L10 phase formed in Fe53Pt47 sample at 500 ◦C annealing is about 5% lower than the amount of L10 already formed in the same annealing conditions for the Fe55Pt45 sample.

Figure 8. The 27 ◦C hysteresis loops of the Fe55Pt45 and Fe53Pt47 samples annealed at 500 ◦C.

Figure 8. The 27 °C hysteresis loops of the Fe55Pt45 and Fe53Pt47 samples annealed at 500 °C. In order to quantify the potential magnetic performance of these alloys, we also calculated the maximumIn order energyto quantify product the (BH)potentialmax, a magnetic parameter performance which measures of these the amount alloys, ofwe magnetic also calculated energy usablethe maximumin these magnetsenergy product in various (BH) applications.max, a parameter Being which essentially measures proportional the amount to the of area magnetic encompassed energy by usablethe hysteresis in these loops, magnets (BH) inmax variouswas calculated applications. from the Being experimental essentially magnetization proportional vs. to applied the area field encompassedvalues and wasby the found hysteresis to be 19.8 loops, MGOe (BH) formax the was Fe 55calculatedPt45 sample from annealed the experimental at 500 ◦C and magnetization 18.4 MGOe for vs. theapplied Fe53Pt field47 sample values annealed and was at found 500 ◦C. to Thesebe 19.8 values MGOe are for situated the Fe towards55Pt45 sample the higher annealed limits at reached500 °C in andthe 18.4 literature MGOe forfor FePtthe Fe films—for53Pt47 sample instance, annealed Shen at et 500 al. in°C. [43 These] reported values values are situated of about towards 20 MGOe the for higher300 nmlimits thick reached FePt thinin the films literature annealed for atFePt 500 films◦C. —for instance, Shen et al. in [43] reported values of aboutThese 20 MGOe findings for 300 lead nm to thick the creation FePt thin of films a potentially annealed exchange-coupled at 500 °C. nanocomposite magnet whereThese both findings soft and lead hard to magneticthe creation phases of a coexistpotentially and yield exchange a large-coupled enough nanocomposite maximum energy magnet product wherethat both is comparable soft and hard with magnetic existing phase permanents coexist magnets. and yield a large enough maximum energy product that is comparable with existing permanent magnets. 4. Conclusions 4. ConclusionsTwo alloys of a composition slightly off-stoichiometric from the equiatomic FePt system were synthesized by ball milling in a dynamic, rotation-switched regime of a planetary ball-milling device Two alloys of a composition slightly off-stoichiometric from the equiatomic FePt system were from high purity elemental powders dispersed in liquid carrier (hexane) in order to avoid Pt adherence synthesized by ball milling in a dynamic, rotation-switched regime of a planetary ball-milling device to the balls. While the as-milled state of samples Fe Pt and Fe Pt show a face-centered cubic from high purity elemental powders dispersed in liquid53 47 carrier (hexane)55 45 in order to avoid Pt adherence to the balls. While the as-milled state of samples Fe53Pt47 and Fe55Pt45 show a face-centered cubic A1 crystalline symmetry, the annealing at 400 °C produced multiple additional Bragg lines in the X-ray diffractograms. The occurrence of additional so-called superlattice peaks as well as the splitting observed for some of the Bragg peaks of the initial cubic symmetry due to tetragonal distortion showed that the L10 phase FePt is formed after annealing at 400 °C, a temperature lower than what is usually needed to form the L10 phase in conventional FePt alloys and nanoparticles. Quantitative full-profile analysis of the XRD diffractograms allowed us to prove a high amount of the L10 phase (72% and 78%, respectively, for Fe53Pt47 and Fe55Pt45), while 57Fe Mössbauer spectroscopy brought further proof for the two-phase behavior. Two magnetic sublattices were identified from the fitting of the Mössbauer spectrum with significantly different quadrupole splitting and hyperfine field values. Magnetometry measurements provided further evidence of the coexistence of the hard magnetic tetragonal L10 and soft magnetic cubic FePt phases, with strong predominance of the hard magnetic phase. This was proven by the existence of the inflection points in the hysteresis loops, which is a sign of the interplay between two magnetic phases, one hard magnetic and another one soft magnetic. For the Fe53Pt47 and Fe55Pt45 samples annealed at 500 °C, the two magnetic phases became fully exchange coupled as witnessed by the lack of inflection points in the hysteresis loops. As a consequence, large coercivity (up to 8.2 kOe) and significant remanent magnetization were obtained, leading to a maximum energy product of up to 19.8 MGOe, which is comparable with those

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A1 crystalline symmetry, the annealing at 400 ◦C produced multiple additional Bragg lines in the X-ray diffractograms. The occurrence of additional so-called superlattice peaks as well as the splitting observed for some of the Bragg peaks of the initial cubic symmetry due to tetragonal distortion showed that the L10 phase FePt is formed after annealing at 400 ◦C, a temperature lower than what is usually needed to form the L10 phase in conventional FePt alloys and nanoparticles. Quantitative full-profile analysis of the XRD diffractograms allowed us to prove a high amount of the L10 phase (72% and 78%, 57 respectively, for Fe53Pt47 and Fe55Pt45), while Fe Mössbauer spectroscopy brought further proof for the two-phase behavior. Two magnetic sublattices were identified from the fitting of the Mössbauer spectrum with significantly different quadrupole splitting and hyperfine field values. Magnetometry measurements provided further evidence of the coexistence of the hard magnetic tetragonal L10 and soft magnetic cubic FePt phases, with strong predominance of the hard magnetic phase. This was proven by the existence of the inflection points in the hysteresis loops, which is a sign of the interplay between two magnetic phases, one hard magnetic and another one soft magnetic. For the Fe53Pt47 and Fe55Pt45 samples annealed at 500 ◦C, the two magnetic phases became fully exchange coupled as witnessed by the lack of inflection points in the hysteresis loops. As a consequence, large coercivity (up to 8.2 kOe) and significant remanent magnetization were obtained, leading to a maximum energy product of up to 19.8 MGOe, which is comparable with those of current permanent magnets. These results provide an interesting method to obtain an exchange spring nanocomposite magnet, which is promising in view of use of such materials as a future class of rare earth-free permanent magnets.

Author Contributions: Conceptualization, O.C.; methodology, A.D.C.; validation, I.D.; formal analysis, A.L.; investigation, P.P. and A.P. All authors have read and agreed to the published version of the manuscript.

Funding: Project PCCDI 47 contract 597/2018; POC Projects: “Materiale multifunct, ionale inteligente pentru aplicat, ii de înaltă tehnologie (MATI2IT)” (code MySMIS 105726), and: “Materiale Avansate Speciale pe Baza de Bor si de Pamanturi Rare—REBMAT” POC 28/2016, both funded by the Romanian Ministry of European Funds; Core Program project PN21N funded by the Romanian Ministry of Education and Research. Acknowledgments: Technical support from the X04 SA materials science beamline at the Swiss Light Source (Paul Scherrer Institute, Villigen, Switzerland) is gratefully acknowledged. Conflicts of Interest: The authors declare no conflict of interest.

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