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L10 phase formation in ternary FePdNi alloys
Journal of Alloys and Compounds 648 (2015) 845e852
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L10 phase formation in ternary FePdNi alloys A.M. Montes-Arango a, N.C. Bordeaux b, J. Liu c, K. Barmak c, L.H. Lewis a, b, * a
Department of Mechanical and Industrial Engineering, Northeastern University, Boston, MA 02115, USA Department of Chemical Engineering, Northeastern University, Boston, MA 02115, USA c Department of Applied Physics and Applied Mathematics, Columbia University, New York, NY 10027, USA b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 11 May 2015 Received in revised form 1 July 2015 Accepted 2 July 2015 Available online 7 July 2015
Metallurgical routes to highly metastable phases are required to access new materials with new functionalities. To this end, the stability of the tetragonal chemically ordered L10 phase in the ternary FeePd eNi system is quantified to provide enabling information concerning synthesis of L10-type FeNi, a highly attractive yet highly elusive advanced permanent magnet candidate. Fe50Pd50xNix (x ¼ 0e7 at%) samples were arc-melted and annealed at 773 K (500 C) for 100 h to induce formation of the chemically ordered L10 phase. Coupled calorimetry, structural and magnetic investigations allow determination of an isothermal section of the ternary FeePdeNi phase diagram featuring a single phase L10 region near the FePd boundary for x < 6 at%. It is demonstrated that increased Ni content in Fe50Pd50xNix alloys systematically decreases the order-disorder transition temperature, resulting in a lower thermodynamic driving force for the ordering phase transformation. The Fe50Pd50xNix L10 / fcc disordering transformation is determined to occur via a two-step process, with compositionally-dependent enthalpies and transition temperatures. These results highlight the need to investigate ternary alloys with higher Ni content to determine the stability range of the L10 phase near the FeNi boundary, thereby facilitating kinetic access to the important L10 FeNi ferromagnetic phase. © 2015 Elsevier B.V. All rights reserved.
Keywords: Chemically-ordered L10 structure FePdNi ternary alloy FePd alloy FeNi alloy Order to disorder phase transformations
1. Introduction and motivation Ferrous compounds exhibiting the tetragonal L10 chemicallyordered structure (space group P4/mmm) are attractive for advanced magnetic material applications due to their high uniaxial magnetocrystalline anisotropy (typically Ku ~ 107 erg/cc [106 J/m3]) [1,2] and high saturation magnetization (MS ~ 1100e1140 emu/cc [1100e1140 kA/m]) [1], which combine to yield room-temperature energy densities comparable to those of Nd2Fe14B magnets [3]. Composed of crystallographic planes stacked along the tetragonal c-axis which alternate in atomic species, the L10 structure type is a superstructure which is generally restricted to the vicinity of the equiatomic composition. The high uniaxial magnetic anisotropy in L10-type compounds is strongly related to the reduced symmetry of the crystal structure, which provides a uniaxial easy direction of magnetization along the tetragonal c-axis. Typically, chemicallyordered L10 phases form by nucleation and growth from a parent low-anisotropy face-centered cubic (fcc) phase at temperatures
* Corresponding author. Department of Chemical Engineering, Northeastern University, Boston, MA 02115, USA. E-mail address: [email protected] (L.H. Lewis). http://dx.doi.org/10.1016/j.jallcom.2015.07.019 0925-8388/© 2015 Elsevier B.V. All rights reserved.
below a critical chemical order-disorder temperature (TOD) [4,5]. While the L10 phase is found in the binary FePt and FePd systems that have attracted considerable attention for thin film magnetic recording media applications, their high cost limits their use for bulk applications. Consequently, technological interests are shifting to other transition-metal-based alloys in the L10 family, including the nominally equiatomic FeNi compound, which is known to also adopt the L10 structure [6e10]. The ready availability of the constituent elements, their comparatively low cost and promising magnetic properties [8,9] make L10-structured FeNi an exceptionally interesting material for bulk advanced permanent magnets, with relevance for energy harvesting and conversion applications. Formation of the highly attractive L10 FeNi phase in laboratory time frames is extremely challenging due to exceptionally slow lattice diffusion below the equilibrium chemical order-disorder temperature [7,11], TOD ¼ 593 K (320 C) [12]. Therefore, cooling rates of the order of 0.1e100 K/106 years are necessary for natural formation of L10 FeNi (also known as terataenite), reported to occur in selected stony, stony-iron, and iron meteorites [3,6,10e13] where it is typically found either as sub-50-mm-sized grains in contact with other phases or as sub-20-mm-wide rims surrounding fcc FeNi (taenite) grains [10]. Laboratory synthesis of very small amounts of
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the L10 phase in FeNi has been reported using electron and neutron irradiation techniques [7e9], and recently ultrathin films deposited using molecular beam epitaxy have displayed characteristics consistent with the tetragonal form of the compound [14e16]. However, bulk synthesis of L10 FeNi by conventional metallurgical techniques at technologically-relevant time scales is unlikely, and therefore other metallurgical routes to facilitate kinetic access to this phase are required. One possibility is the use of ternary alloying additions to the FeNi system: given the low TOD reported in binary FeNi and the accompanying low lattice diffusivity, it is anticipated that ternary alloying additions to increase TOD will be required either to amplify the driving force for nucleation of the L10 phase [17] or to permit annealing at higher temperatures, fostering increased lattice diffusivities. A useful starting point for the selection of ternary additions to FeNi considers thermodynamicallystable L10 phases containing Fe or Ni (FeX or NiX alloys); evaluation of the effects of substitution of element X by either Ni (for an FeX alloy) or Fe (for a NiX alloy) may then be used to determine the amount of alloying element addition to FeNi required for laboratory formation of the L10 phase. This current work presents structural, magnetic and calorimetric results concerning L10 phase formation in the alloy family FePdNi to inform synthesis routes to achieve L10-type FeNi. Additions of Ni are substituted into the nominally-equiatomic bulk FePd system that is known to readily transform to the L10 structure through a chemical disorder-order transformation. The stability of the L10 phase in the ternary FePdNi system is analyzed, to provide improved understanding of the development of chemical ordering, magnetocrystalline anisotropy and phase formation in the FeNi system. The binary FePd compound exhibits an equilibrium chemical order-disorder transformation temperature in the vicinity of 933e1063 K (660e790 C) for the compositional range 50e60 at % Pd [18]. The magnetic character and chemical ordering behavior in the FePd system have been widely studied in nanoparticles [19,20], in thin films [21e24] and in bulk form [25e32]. Studies of the effects of ternary alloying additions on the FePd chemical orderdisorder transformation are limited to only a few investigations [33e38]. In particular, the effects of Ni additions in bulk equiatomic L10 FePd have been reported only by Horiuchi et al. [38] who also performed phenomenological calculations of phase equilibria in the FePdNi ternary system at temperatures between 823 K (550 C) and 923 K (650 C). Based on these calculations, they concluded that Ni substitution for Pd is energetically more favorable than for Fe in the L10-compound. The present work quantifies L10 phase stability in Fe50Pd50-xNix (x 7 at%) alloys, allowing determination of the 773 K isothermal section of the ternary FePdNi phase diagram near the FePd boundary. While it is demonstrated that the order-disorder transition temperature decreases with increased Ni, resulting in a lower thermodynamic force for an ordering phase transformation, future work requires study of ternary alloys with higher Ni content to determine the stability range of the L10 phase near the FeNi boundary, thereby facilitating kinetic access to the important L10 FeNi ferromagnetic phase.
12 mm. EDS spectra were obtained at a take-off angle of 35.71 at three different locations of selected arc-melted samples that were polished with 1.0 and 0.5 mm alumina powder slurry. Alloy compositions were determined from the spectra after application of the atomic number, absorption and fluorescence correction factors using the ZAF scheme [39]. No standards were employed in the alloy composition analysis; however, as shall be seen, the measured alloy compositions were close to the nominal (i.e., intended) values, and thus the measured compositions are considered reliable. To induce L10 chemical ordering from the metastably-retained fcc phase, samples were encapsulated in evacuated (0.13 103 Pa) silica tubes and annealed for 100 h at 773 ± 5 K (500 ± 5 C). The annealing conditions were selected according to the time-temperature-transformation (TTT) diagram reported by prior studies on bulk FePd [40] that indicate the maximum chemical-ordering rate is in the vicinity of 773 K, with full transformation into the L10 phase guaranteed for annealing times of 100 h. Additionally, the temperature 773 K is intermediate between the order-disorder temperatures of the binary FePd and FeNi systems. It is anticipated that this temperature will allow growth of the L10 phase at laboratory time scales in the ternary system, provided that TOD for the ternary alloy is sufficiently high to permit phase nucleation. The lattice diffusivities of the FeNi, NiPd and FePd binary alloys at 773 K are in the range 1015-1020 cm2/s [7,41e43]. The crystal structure of the samples was examined by X-ray diffraction (XRD) using a PANalytical q-2q diffractometer employing CuKa-radiation. The resultant Bragg peaks from XRD data were fit with a pseudo-Voigt function and a least-square fitting procedure was used to estimate the lattice parameters of the phase(s) [44]. Line broadening analysis using a variation of the WilliamsonHall approach [45,46] was employed to estimate the crystallite size D and strain ε within the lattice; after correction for instrumental broadening, the full-width at half-maximum b of Bragg reflections found at an angle 2q was used to generate a plot of
b* ¼
1 d* þ 2s ; D Ehkl
(1)
where b* ¼ b cos q=l and d* ¼ 2 sin q=l, with l the radiation wavelength. The stress s is related to the strain through Hook's law s ¼ εE. The Young's moduli Ehkl perpendicular to particular (hkl) planes were used to account for elastic anisotropy. These moduli were calculated from the reported elastic constants of fcc-Fe50Pd50 and L10-Fe50Pd50 [29], and were assumed to be equal for all studied Fe50Pd50-xNix compositions. As a chemically-ordered atomic arrangement, the L10 phase may be characterized by the long-range order parameter LRO, which describes the degree of chemical order ranging from zero (absence of order) to unity (perfect order). The LRO may be estimated from Xray diffraction data as
LRO ¼
sffiffiffiffiffiffiffiffiffiffi IS $I *F IF $I *S
(2)
2. Experimental Alloys with the nominal starting composition Fe50Pd50-xNix (at %), x ¼ 0, 3, 5 and 7, were synthesized from high-purity elemental granules by arc-melting in an Ar atmosphere, and repeating the process at least two times for homogeneity. Disks (~3 mm in diameter) were sliced out of the arc-melted charges for further processing and analyses. Composition and chemical homogeneity were verified by means of an energy-dispersive X-ray spectrometer (EDS) coupled to a field emission scanning electron microscope (SEM, Hitachi S4800) operating at 20 kV and a working distance of
where IS and IF are the experimental integrated intensities of the superstructure and fundamental Bragg reflections, while IS* and IF* are the theoretical values calculated for a sample with LRO ¼ 1 [47]. The LRO calculations conducted in this current study employed the measured (110) superlattice and (220) fundamental Bragg peak intensities, corrected for effects of finite sample size as the samples tested were smaller than the incident X-ray illuminated area. The correction factor as a function of the X-ray incident angle was determined by comparison of the experimental integrated intensities of the Bragg reflections of two polycrystalline randomly
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oriented Si powder disk references, with diameters ~3 mm (equal to diameter of the FePdNi samples) and ~22 mm (larger than X-ray illuminated area) [48]. Finally, in circumstances of multiphase coexistence, the volume fraction fo of a given phase (either fL10 for the L10 phase or ffcc for the fcc phase), was determined from XRD data using a relation of the Bragg peak intensities of fcc and L10 reflections as described by Cebollada et al. [49]. The magnetic character of the samples was investigated as a function of magnetic field using vibrating sample magnetometry (VSM, Quantum Design Versalab). In-plane hysteresis loops of the disk-shaped samples were obtained at room temperature with a maximum applied field of 30 kOe (2387 kA/m) and demagnetization corrections for the specific sample geometries were applied to the data [50]. For samples in which the maximum applied field was not sufficient to achieve saturation, MS was estimated graphically from the plot of M vs. 1/H, as described by the law of approach to saturation for polycrystalline materials [51]. Structural and magnetic phase transformations in the temperature range 300 K < T < 1050 K were studied using calibrated differential scanning calorimetry (DSC, Netzsch STA 449 F3) in an argon atmosphere under a heating rate of 20 K/min. The reversibility of any transformation observed during the first DSC heating scan was confirmed by cooling at a rate of 20 K/min and performing a second heating run on the same sample. First-order phase transitions were identified as sharp exothermic or endothermic peaks, with onset temperatures Tonset that define the beginning of the transition and the area under the peak representing the enthalpy of transformation DH. For the chemical order-disorder transformation, Tonset was taken as an indication of the order-disorder temperature TOD. Second-order phase transitions such as the Curie temperature TC were observed as a clear shift in background DSC signal rather than as a sharp peak. Estimated errors for reported temperatures are ±3 K and for reported enthalpies ±0.2 kJ/mol. In specific cases the reported enthalpy values are determined as the sum of enthalpies from all observed endothermic events. 3. Results The nominal and measured alloy compositions are given in Table 1. The measured compositions are the mean of three separate measurements, with standard deviations that are in the range of 0.1e1.0 at% and thus evidence compositional homogeneity of the samples. The table also shows that the measured alloy compositions are close to the nominal values, with the largest difference seen for the FePd sample as 2.5 at%. For simplicity, from this point onward the alloys will be referred to by their determined Ni content as x ¼ 0, 3.5, 5.2 and 7.5. As confirmed by XRD, all as-arc-melted samples consist of the fcc chemically-disordered phase, with no secondary phases observed. As an example, Fig. 1 presents X-ray diffractograms for the two extreme compositions (x ¼ 0 and 7.5) in the range 32 < 2q < 52 , where the fcc (111) and (200) Bragg peak reflections are observed. Upon annealing, the presence of the L10 phase is evidenced by the appearance of (001), (110), (201), (112), (221), (310), (203), and (312) superlattice Bragg peaks associated with chemical ordering, as well as by the splitting of the (200),
847
Fig. 1. XRD diffractograms obtained using Cu Ka radiation for as-arc-melted and annealed x ¼ 0 and 7.5 samples.
(220) and (311) fundamental Bragg reflections signaling the symmetry reduction from cubic to tetragonal. Fig. 1 displays XRD results for the annealed x ¼ 0 and 7.5 samples, where the (110)L10 superlattice peak characterizing the L10 phase is observed around 2q z 33 , with an accompanying (200) peak splitting in the interval 46 < 2q < 50 . While samples of Ni content x 3.5 have transformed into a nominal single-phase L10 structure upon annealing, annealed samples with x 5.2 are multiphase. Fig. 2 presents the X-ray diffractogram for these samples in the range 46 < 2q < 50 , illustrating tetragonal L10 (200) and (002) Bragg reflections along with additional peaks marked by an asterisk “*”. The additional peaks observed in the XRD data obtained from the x ¼ 5.2 annealed sample (Fig. 2a) correspond to an fcc-based phase with a lattice parameter that is the same, within experimental error, as that of the fcc phase in the x ¼ 5.2 as-arc-melted sample. Thus, this fcc phase has been tentatively attributed to a remnant fcc parent phase. Further, XRD data from the x ¼ 7.5 annealed sample (Fig. 2b) indicate the presence of two additional fcc-based phases, one with a lattice parameter that within experimental error corresponds to that of the x ¼ 7.5 as-arc-melted fcc phase, and the other with a unit cell volume that is 1.3% smaller than that of the x ¼ 7.5 as-arcmelted fcc phase. These phases have thus been tentatively attributed to a remnant fcc parent phase and a “new” fcc phase, respectively. The L10 phase fractions fL10 in the x ¼ 5.2 and 7.5 annealed samples are estimated as 95 vol% and 20 vol%, respectively, while the fcc phase fractions ffcc in the x ¼ 7.5 annealed sample are estimated as 20 vol% (parent phase) and 60 vol% (new phase). Fig. 3 illustrates trends in the structural parameters of the fcc asarc-melted and L10 annealed FePdNi alloys as a function of nickel content. The unit cell volumes decrease slightly with increased Ni
Table 1 Alloy nominal and measured compositions. Nom.
Fe Pd Ni
50 50
Measured
Nom.
Avg.
Std. Dev.
47.5 52.5
1 1
50 47 3
Measured
Nom.
Avg.
Std. Dev.
49.3 47.2 3.5
0.5 0.4 0.3
50 45 5
Measured
Nom.
Avg.
Std. Dev.
47.5 47.3 5.2
0.8 0.7 0.2
50 43 7
Measured Avg.
Std. Dev.
48.2 44.3 7.5
0.4 0.4 0.1
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Fig. 2. Multi-peak structures in the vicinity of the (200) diffraction peak in annealed x 5.2 samples. Peaks labeled with “*” are attributed to fcc phases.
Fig. 3. Variation with Ni content of unit cell volume, c/a ratio, LRO parameter, crystallite size (D) and strain (ε) of the fcc as-arc-melted and L10 annealed phases (lines drawn to guide the eye). Data points for LRO in x 5.2 annealed samples and crystallite size in the x ¼ 7.5 annealed sample have been omitted considering the large error associated with the calculated values.
concentration, outside of experimental error. Lattice parameter(s) obtained for the binary FePd composition in its fcc (a ¼ 3.82 Å) and L10 (a ¼ 3.862 Å, c ¼ 3.697 Å) states are consistent with values reported in the literature [18,32]. Ni additions to the L10 phase
cause an initial increase and then decrease in the c/a ratio at x ¼ 5.2, while the L10 LRO parameter decreases with 3.5 at% Ni additions. Data points for the LRO parameter of x 5.2 at% samples have been omitted from Fig. 3 on account of their large associated error, attributed to the de-convolution process of the individual L10 and fcc reflections in the multi-peak observed. The calculated crystallite sizes (D) for as-arc-melted fcc samples decrease from ~43 nm to ~ 20 nm with increased Ni concentration, Fig. 3. Upon annealing and formation of the L10 phase, an increase in the crystallite size is observed for x 5.2 samples. For the x ¼ 7.5 annealed sample, the crystallite size has a large associated error and thus this data point has been omitted from Fig. 3. The calculated lattice strain values (ε) are moderate, in the range 0.10% < ε < 0.25%. For a given composition, the strain manifest in the cubic fcc phase is the same, within experimental error, as that characterizing the tetragonal L10 phase. The room-temperature magnetization of all as-arc-melted fcc alloys increases rapidly with increased applied field, reaching saturation at anisotropy fields around HA z 3000 Oe (238.7 kA/m; see Fig. 4a). Measured MS values for these alloys (MS ¼ 1066e1110 emu/cc (1066e1110 kA/m)) are in good agreement with those reported in the literature for bulk Fe50Pd50 [1]. While nickel additions slightly increase MS for the fcc as-arc-melted phase (Fig. 4b), they have a negligible effect in the coercivity and remanence, with values HCi < 25 Oe (2 kA/m) and MR < 60 emu/cc (60 kA/m). In contrast, the magnetization of the annealed x 5.2 samples slowly increases with applied field without reaching saturation at 30 kOe (2387 kA/m), indicative of increased magnetocrystalline anisotropy. Anisotropy fields and saturation magnetization for these L10 alloys are larger than those of their fcc counterparts, with HA > 27,000 Oe (2148.6 kA/m) and extrapolated saturation values (shown in Fig. 4b) between MS ¼ 1160e1180 emu/cc (1160e1180 kA/m). The coercivity and magnetic remanence values for annealed L10-type samples (x 5.2) are in the range 94e160 Oe (7.5e12.7 kA/m) and 72e157 emu/cc (72e157 kA/m). For the x ¼ 7.5 annealed sample, a 2% increase in MS relative to that of the fcc as-arc-melted state is noted, while no significant change in HA, HCi or MR was detected despite noted crystallographic differences. All as-arc-melted fcc samples exhibit only a single calorimetric event of second-order phase transformation character that is associated with the Curie temperature TC. As an example, Fig. 5 presents DSC scans for the two extreme compositions (x ¼ 0 and 7.5, arc-melted). On the other hand, annealed L10-type samples (x 5.2) show multiple thermal features, an example of which can be seen in Fig. 5 for the x ¼ 0 annealed alloy; the first heating scan reveals the L10 Curie transition followed by a broad irreversible
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849
Fig. 4. a) Room temperature field-dependent magnetization of binary FePd as-arc-melted (fcc) and annealed (L10) sample. b) Evolution of MS with Ni content for as-arc-melted and annealed samples. Note that the annealed x ¼ 7.5 sample is plotted separately from the L10 trend, as it has a multiphase character with a low L10 volume fraction. The error bars for the x ¼ 7.5 annealed sample are smaller than the symbol.
Fig. 5. DSC scans of x ¼ 0 (green) and 7.5 (pink) samples in their as-arc-melted state (top) and annealed state (bottom). For the annealed sample, consecutive heating scans are labeled “1st scan” and “2nd scan”. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
endothermic event preceding a first-order endothermic transition associated with the L10/fcc order-disorder phase transformation to produce a “recovered” fcc phase. A subsequent heating scan (2nd scan in Fig. 5) reveals only a single thermal event associated with the Curie transition of the recovered fcc phase, where TC(recovered fcc) > TC(as-arc-melted fcc) for x > 0 alloys. For the x ¼ 7.5 annealed sample, only a Curie transition is detected in calorimetry (Fig. 5), despite the fact that the XRD data evidenced approximately 20 vol% of the L10 phase in this sample; no measurable enthalpy associated with the chemical order-disorder transformation phase was obtained. Addition of Ni to the fcc and L10 lattices increases the Curie temperature, Fig. 6, with a Curie temperature difference between as-arc-melted and recovered fcc phases that increases with increased Ni content. A similar trend is observed between the saturation magnetization of as-arc-melted and recovered fcc phases (not shown), with a maximum increase of 2.1% exhibited by the recovered fcc x ¼ 5.2 sample. Increased Ni content produces a decrease in the onset temperatures and the associated enthalpies of the L10/fcc phase transition, Fig. 6. For reference, Fig. 6 includes an
Fig. 6. Evolution of the Curie temperature of L10 and fcc phases and of the onset temperature and enthalpy of the order-disorder transformation with Ni content. An extrapolated TOD value calculated for the x ¼ 7.5 alloy is plotted for reference.
extrapolated value of TOD (x ¼ 7.5) obtained from a polynomial fit to the experimental TOD values as a function of Ni content. 4. Discussion Structural differences between the as-arc-melted and annealed FePdNi alloys are reflected in their magnetic behavior. Attainment
850
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of the L10 phase in x 5.2 annealed samples is accompanied by an increase in both intrinsic (i.e., MS, TC) and extrinsic (i.e., HCi, MR, HA) magnetic properties. Annealing of the x ¼ 7.5 sample contributes to a moderate increase in MS and TC, while no detectable changes in HCi, MR, or HA are observed. This moderate increase in MS and TC, that deviates largely from the trends observed for the L10 phase in other samples, is attributed to the multi-phase character of the x ¼ 7.5 annealed sample, with a low L10-phase fraction (20 vol%). In analogy, noted differences in TC and MS between the as-arc-melted and DSC-recovered fcc phases allude to subtle structural differences between them; it is proposed that retention of a small degree of short-range order occurs after the material undergoes the L10/fcc phase transition upon heating in the DSC. The detection of a very broad endothermic feature preceding the main chemical order-disorder peak in FePdNi alloys is consistent with a two-stage disordering phase transformation (L10/fcc). While broad endothermic signals that precede a chemical orderdisorder transformation have been reported for thermoresistometry [51] and calorimetry [52] experiments conducted on chemically-ordered CuAu samples, their origin remains unexplained. In this current study, the broad endothermic peak is associated with atomic short-range motion to produce a readjustment of the LRO parameter below the temperature at which the nucleation and growth of the chemically-disordered fcc phase may be measured. In this context, the total enthalpy of transformation DHL10/fcc is the sum of both endothermic events. The phases detected in all annealed (T ¼ 773 K) samples are in qualitative agreement with those presented in the calculated (T ¼ 823 K) FePdNi ternary diagram of Horiuchi et al. [38]. Horiuchi's phase diagram indicates a single L10 phase centered around the Fe50Pd50 composition next to a two-phase “disorder” þ L10 region, Fig. 7a. The solubility limit of Ni in a single L10 phase at a fixed Fe content (~50 at%) is provided by Horiuchi as z 5 at%. Consistent with this scenario, in the present study a single L10 phase was detected in annealed x 3.5 samples. For the x ¼ 5.2 annealed alloy, the transformation did not go to completion, with a small volume fraction (~5 vol%) of the parent fcc phase remaining. It is anticipated that this alloy should transform into a single L10 phase if sufficient time is allowed to achieve equilibrium. For the x ¼ 7.5 annealed alloy, the observed phases are also consistent with Horiuchi et al.'s proposed diagram, for which a two-phase mixture is expected to exist in equilibrium; these two phases correspond to the observed “new” fcc and L10 phases, while the presence of a remnant fcc parent phase indicates that the transformation was not
complete. This apparently good agreement between the current results and those of Horiuchi et al. should be considered with caution as (i) the calculations of Horiuchi et al. [38] predict TOD ¼ 800 K for binary L10 FeNi, significantly higher that the experimentally determined value of 320 K [7,52,53] and (ii) the L10 FePd boundary in Horiuchi et al.'s calculated ternary phase diagram does not agree with the published FePd binary phase diagram [18]. Nonetheless, Horiuchi et al.'s proposed L10 phase boundary near the 50 at% Fe composition has been used to estimate the Ni concentrations of the phases present in the x ¼ 7.5 annealed alloy. By invoking assumptions that mass fraction is equivalent to volume percent and that the L10 phase and the parent disordered fcc alloy possess the same Fe concentration (48.2 at%), the Ni concentrations of the phases present in the x ¼ 7.5 annealed alloy are estimated as x ¼ 6 at% (L10 phase) and as x ¼ 8 at% (new fcc phase), (Fig. 7b). This information allows determination of a very narrow fcc þ L10 twophase region in the composition range 6 at% < Ni < 8 at% for 48.2 at% Fe at 773 K. Using the Ni content limits determined here for the fcc þ L10 two-phase region, along with phase equilibria information at 773 K obtained from the binary FeNi [54], NiPd [55], and FePd [18] phase diagrams, an updated isothermal section of the ternary FePdNi phase diagram is proposed in Fig. 8. It is noted that the phase boundaries presented for single-phase regions other than L10 are arbitrary and are included for illustration purposes only. Subjecting the Ni ¼ 5.2 and Ni ¼ 7.5 alloys to longer annealing times should prove useful in the determination of the L10 and the fcc þ L10 phase boundaries with greater certainty. Assignment of a single-phase L12 region extending across the full composition range is made based on the complete mutual solubility of Ni and Pd in the fcc phase (g) in the NiePd binary system. In the proposed isothermal section, an fcc phase is shown to be more stable than an L10 phase at 773 K for Fe50Pd50-xNix alloys with higher Ni contents than studied here. However, only experimental results from FePd alloys made with higher Ni contents can confirm whether a stable or metastable L10 region can be found near the Fe50Ni50 composition at this temperature. 5. Conclusions This study confirms that Ni additions to FePd significantly impact the stability of the L10 phase in the FePdNi ternary system, reflecting important changes in the structural characteristics and magnetic properties of these alloys to inform future synthesis of the
Fig. 7. a) FePdNi ternary phase diagram at 823 K proposed by Horiuchi et al. [38], along with the alloy compositions of the present study. (b) Estimated compositions of the multiple phases present in the x ¼ 7.5 annealed sample, namely the parent fcc phase (blue) and the product L10 (red) þ new fcc (green) phases. Horiuchi et al.'s L10 boundary is left for reference as a dashed line as it was used to estimate the compositions. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
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[2] [3]
[4]
[5]
[6] [7] [8]
[9]
[10] [11] Fig. 8. 773 K isothermal section proposed for the updated FePdNi ternary phase diagram (a represents the body centered cubic crystal structure while g represents an fcc phase). L10 and fcc þ L10 phase boundaries have been selected to match experimental observations, while phase boundaries displayed for other regions are arbitrary.
alternative permanent magnet candidate tetrataenite. Comparison of structural data and phase determinations obtained here with those reported by Horiuchi et al.'s calculated ternary FePdNi phase diagram [38] reveal that samples with nickel content x 3.5 at% achieve equilibrium after 100 h of annealing, while longer annealing times are required to form the L10 phase in compositions of higher Ni content. The sluggish kinetics observed for x 5.2 at% alloys result from a decreased fcc/L10 phase transformation driving force that is attributed to a reduced order-disorder temperature upon Ni addition. Furthermore, a narrow fcc þ L10 twophase region is proposed for the range 6 at% x 8 at% in Fe48.2Pd51.8xNix alloys. Study of the reverse phase transformation (L10/fcc) that occurs upon heating the annealed alloys of this study reveals a two-stage disordering process that initiates with short-range atomic motion to readjust the LRO parameter prior to the first-order chemical disordering transition. Additionally, retention of a certain degree of short range order was observed after undergoing the L10/fcc transition, which is amplified with increased Ni content. This confirms that Ni additions have a considerable effect on the kinetics of the order-disorder phase transformation. The results of this work indicate that L10 formation in the ternary Fe50Pd50-xNix is highly sensitive to slight compositional variations. It is therefore necessary to investigate ternary alloys with higher Ni contents than studied here to confirm if an L10 phase can be found near the FeNi boundary, thus allowing kinetic access to the highly attractive ferromagnetic L10 FeNi phase, relevant for the next-generation permanent magnets.
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[26]
Acknowledgments [27]
This study was supported by the National Science Foundation eCMMI Divisione under grants No. 1129433 and No. 1259736, and by Northeastern University.
[28]
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
L10 phase formation in ternary FePdNi alloys A.M. Montes-Arango a, N.C. Bordeaux b, J. Liu c, K. Barmak c, L.H. Lewis a, b, * a
Department of Mechanical and Industrial Engineering, Northeastern University, Boston, MA 02115, USA Department of Chemical Engineering, Northeastern University, Boston, MA 02115, USA c Department of Applied Physics and Applied Mathematics, Columbia University, New York, NY 10027, USA b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 11 May 2015 Received in revised form 1 July 2015 Accepted 2 July 2015 Available online 7 July 2015
Metallurgical routes to highly metastable phases are required to access new materials with new functionalities. To this end, the stability of the tetragonal chemically ordered L10 phase in the ternary FeePd eNi system is quantified to provide enabling information concerning synthesis of L10-type FeNi, a highly attractive yet highly elusive advanced permanent magnet candidate. Fe50Pd50xNix (x ¼ 0e7 at%) samples were arc-melted and annealed at 773 K (500 C) for 100 h to induce formation of the chemically ordered L10 phase. Coupled calorimetry, structural and magnetic investigations allow determination of an isothermal section of the ternary FeePdeNi phase diagram featuring a single phase L10 region near the FePd boundary for x < 6 at%. It is demonstrated that increased Ni content in Fe50Pd50xNix alloys systematically decreases the order-disorder transition temperature, resulting in a lower thermodynamic driving force for the ordering phase transformation. The Fe50Pd50xNix L10 / fcc disordering transformation is determined to occur via a two-step process, with compositionally-dependent enthalpies and transition temperatures. These results highlight the need to investigate ternary alloys with higher Ni content to determine the stability range of the L10 phase near the FeNi boundary, thereby facilitating kinetic access to the important L10 FeNi ferromagnetic phase. © 2015 Elsevier B.V. All rights reserved.
Keywords: Chemically-ordered L10 structure FePdNi ternary alloy FePd alloy FeNi alloy Order to disorder phase transformations
1. Introduction and motivation Ferrous compounds exhibiting the tetragonal L10 chemicallyordered structure (space group P4/mmm) are attractive for advanced magnetic material applications due to their high uniaxial magnetocrystalline anisotropy (typically Ku ~ 107 erg/cc [106 J/m3]) [1,2] and high saturation magnetization (MS ~ 1100e1140 emu/cc [1100e1140 kA/m]) [1], which combine to yield room-temperature energy densities comparable to those of Nd2Fe14B magnets [3]. Composed of crystallographic planes stacked along the tetragonal c-axis which alternate in atomic species, the L10 structure type is a superstructure which is generally restricted to the vicinity of the equiatomic composition. The high uniaxial magnetic anisotropy in L10-type compounds is strongly related to the reduced symmetry of the crystal structure, which provides a uniaxial easy direction of magnetization along the tetragonal c-axis. Typically, chemicallyordered L10 phases form by nucleation and growth from a parent low-anisotropy face-centered cubic (fcc) phase at temperatures
* Corresponding author. Department of Chemical Engineering, Northeastern University, Boston, MA 02115, USA. E-mail address: [email protected] (L.H. Lewis). http://dx.doi.org/10.1016/j.jallcom.2015.07.019 0925-8388/© 2015 Elsevier B.V. All rights reserved.
below a critical chemical order-disorder temperature (TOD) [4,5]. While the L10 phase is found in the binary FePt and FePd systems that have attracted considerable attention for thin film magnetic recording media applications, their high cost limits their use for bulk applications. Consequently, technological interests are shifting to other transition-metal-based alloys in the L10 family, including the nominally equiatomic FeNi compound, which is known to also adopt the L10 structure [6e10]. The ready availability of the constituent elements, their comparatively low cost and promising magnetic properties [8,9] make L10-structured FeNi an exceptionally interesting material for bulk advanced permanent magnets, with relevance for energy harvesting and conversion applications. Formation of the highly attractive L10 FeNi phase in laboratory time frames is extremely challenging due to exceptionally slow lattice diffusion below the equilibrium chemical order-disorder temperature [7,11], TOD ¼ 593 K (320 C) [12]. Therefore, cooling rates of the order of 0.1e100 K/106 years are necessary for natural formation of L10 FeNi (also known as terataenite), reported to occur in selected stony, stony-iron, and iron meteorites [3,6,10e13] where it is typically found either as sub-50-mm-sized grains in contact with other phases or as sub-20-mm-wide rims surrounding fcc FeNi (taenite) grains [10]. Laboratory synthesis of very small amounts of
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the L10 phase in FeNi has been reported using electron and neutron irradiation techniques [7e9], and recently ultrathin films deposited using molecular beam epitaxy have displayed characteristics consistent with the tetragonal form of the compound [14e16]. However, bulk synthesis of L10 FeNi by conventional metallurgical techniques at technologically-relevant time scales is unlikely, and therefore other metallurgical routes to facilitate kinetic access to this phase are required. One possibility is the use of ternary alloying additions to the FeNi system: given the low TOD reported in binary FeNi and the accompanying low lattice diffusivity, it is anticipated that ternary alloying additions to increase TOD will be required either to amplify the driving force for nucleation of the L10 phase [17] or to permit annealing at higher temperatures, fostering increased lattice diffusivities. A useful starting point for the selection of ternary additions to FeNi considers thermodynamicallystable L10 phases containing Fe or Ni (FeX or NiX alloys); evaluation of the effects of substitution of element X by either Ni (for an FeX alloy) or Fe (for a NiX alloy) may then be used to determine the amount of alloying element addition to FeNi required for laboratory formation of the L10 phase. This current work presents structural, magnetic and calorimetric results concerning L10 phase formation in the alloy family FePdNi to inform synthesis routes to achieve L10-type FeNi. Additions of Ni are substituted into the nominally-equiatomic bulk FePd system that is known to readily transform to the L10 structure through a chemical disorder-order transformation. The stability of the L10 phase in the ternary FePdNi system is analyzed, to provide improved understanding of the development of chemical ordering, magnetocrystalline anisotropy and phase formation in the FeNi system. The binary FePd compound exhibits an equilibrium chemical order-disorder transformation temperature in the vicinity of 933e1063 K (660e790 C) for the compositional range 50e60 at % Pd [18]. The magnetic character and chemical ordering behavior in the FePd system have been widely studied in nanoparticles [19,20], in thin films [21e24] and in bulk form [25e32]. Studies of the effects of ternary alloying additions on the FePd chemical orderdisorder transformation are limited to only a few investigations [33e38]. In particular, the effects of Ni additions in bulk equiatomic L10 FePd have been reported only by Horiuchi et al. [38] who also performed phenomenological calculations of phase equilibria in the FePdNi ternary system at temperatures between 823 K (550 C) and 923 K (650 C). Based on these calculations, they concluded that Ni substitution for Pd is energetically more favorable than for Fe in the L10-compound. The present work quantifies L10 phase stability in Fe50Pd50-xNix (x 7 at%) alloys, allowing determination of the 773 K isothermal section of the ternary FePdNi phase diagram near the FePd boundary. While it is demonstrated that the order-disorder transition temperature decreases with increased Ni, resulting in a lower thermodynamic force for an ordering phase transformation, future work requires study of ternary alloys with higher Ni content to determine the stability range of the L10 phase near the FeNi boundary, thereby facilitating kinetic access to the important L10 FeNi ferromagnetic phase.
12 mm. EDS spectra were obtained at a take-off angle of 35.71 at three different locations of selected arc-melted samples that were polished with 1.0 and 0.5 mm alumina powder slurry. Alloy compositions were determined from the spectra after application of the atomic number, absorption and fluorescence correction factors using the ZAF scheme [39]. No standards were employed in the alloy composition analysis; however, as shall be seen, the measured alloy compositions were close to the nominal (i.e., intended) values, and thus the measured compositions are considered reliable. To induce L10 chemical ordering from the metastably-retained fcc phase, samples were encapsulated in evacuated (0.13 103 Pa) silica tubes and annealed for 100 h at 773 ± 5 K (500 ± 5 C). The annealing conditions were selected according to the time-temperature-transformation (TTT) diagram reported by prior studies on bulk FePd [40] that indicate the maximum chemical-ordering rate is in the vicinity of 773 K, with full transformation into the L10 phase guaranteed for annealing times of 100 h. Additionally, the temperature 773 K is intermediate between the order-disorder temperatures of the binary FePd and FeNi systems. It is anticipated that this temperature will allow growth of the L10 phase at laboratory time scales in the ternary system, provided that TOD for the ternary alloy is sufficiently high to permit phase nucleation. The lattice diffusivities of the FeNi, NiPd and FePd binary alloys at 773 K are in the range 1015-1020 cm2/s [7,41e43]. The crystal structure of the samples was examined by X-ray diffraction (XRD) using a PANalytical q-2q diffractometer employing CuKa-radiation. The resultant Bragg peaks from XRD data were fit with a pseudo-Voigt function and a least-square fitting procedure was used to estimate the lattice parameters of the phase(s) [44]. Line broadening analysis using a variation of the WilliamsonHall approach [45,46] was employed to estimate the crystallite size D and strain ε within the lattice; after correction for instrumental broadening, the full-width at half-maximum b of Bragg reflections found at an angle 2q was used to generate a plot of
b* ¼
1 d* þ 2s ; D Ehkl
(1)
where b* ¼ b cos q=l and d* ¼ 2 sin q=l, with l the radiation wavelength. The stress s is related to the strain through Hook's law s ¼ εE. The Young's moduli Ehkl perpendicular to particular (hkl) planes were used to account for elastic anisotropy. These moduli were calculated from the reported elastic constants of fcc-Fe50Pd50 and L10-Fe50Pd50 [29], and were assumed to be equal for all studied Fe50Pd50-xNix compositions. As a chemically-ordered atomic arrangement, the L10 phase may be characterized by the long-range order parameter LRO, which describes the degree of chemical order ranging from zero (absence of order) to unity (perfect order). The LRO may be estimated from Xray diffraction data as
LRO ¼
sffiffiffiffiffiffiffiffiffiffi IS $I *F IF $I *S
(2)
2. Experimental Alloys with the nominal starting composition Fe50Pd50-xNix (at %), x ¼ 0, 3, 5 and 7, were synthesized from high-purity elemental granules by arc-melting in an Ar atmosphere, and repeating the process at least two times for homogeneity. Disks (~3 mm in diameter) were sliced out of the arc-melted charges for further processing and analyses. Composition and chemical homogeneity were verified by means of an energy-dispersive X-ray spectrometer (EDS) coupled to a field emission scanning electron microscope (SEM, Hitachi S4800) operating at 20 kV and a working distance of
where IS and IF are the experimental integrated intensities of the superstructure and fundamental Bragg reflections, while IS* and IF* are the theoretical values calculated for a sample with LRO ¼ 1 [47]. The LRO calculations conducted in this current study employed the measured (110) superlattice and (220) fundamental Bragg peak intensities, corrected for effects of finite sample size as the samples tested were smaller than the incident X-ray illuminated area. The correction factor as a function of the X-ray incident angle was determined by comparison of the experimental integrated intensities of the Bragg reflections of two polycrystalline randomly
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oriented Si powder disk references, with diameters ~3 mm (equal to diameter of the FePdNi samples) and ~22 mm (larger than X-ray illuminated area) [48]. Finally, in circumstances of multiphase coexistence, the volume fraction fo of a given phase (either fL10 for the L10 phase or ffcc for the fcc phase), was determined from XRD data using a relation of the Bragg peak intensities of fcc and L10 reflections as described by Cebollada et al. [49]. The magnetic character of the samples was investigated as a function of magnetic field using vibrating sample magnetometry (VSM, Quantum Design Versalab). In-plane hysteresis loops of the disk-shaped samples were obtained at room temperature with a maximum applied field of 30 kOe (2387 kA/m) and demagnetization corrections for the specific sample geometries were applied to the data [50]. For samples in which the maximum applied field was not sufficient to achieve saturation, MS was estimated graphically from the plot of M vs. 1/H, as described by the law of approach to saturation for polycrystalline materials [51]. Structural and magnetic phase transformations in the temperature range 300 K < T < 1050 K were studied using calibrated differential scanning calorimetry (DSC, Netzsch STA 449 F3) in an argon atmosphere under a heating rate of 20 K/min. The reversibility of any transformation observed during the first DSC heating scan was confirmed by cooling at a rate of 20 K/min and performing a second heating run on the same sample. First-order phase transitions were identified as sharp exothermic or endothermic peaks, with onset temperatures Tonset that define the beginning of the transition and the area under the peak representing the enthalpy of transformation DH. For the chemical order-disorder transformation, Tonset was taken as an indication of the order-disorder temperature TOD. Second-order phase transitions such as the Curie temperature TC were observed as a clear shift in background DSC signal rather than as a sharp peak. Estimated errors for reported temperatures are ±3 K and for reported enthalpies ±0.2 kJ/mol. In specific cases the reported enthalpy values are determined as the sum of enthalpies from all observed endothermic events. 3. Results The nominal and measured alloy compositions are given in Table 1. The measured compositions are the mean of three separate measurements, with standard deviations that are in the range of 0.1e1.0 at% and thus evidence compositional homogeneity of the samples. The table also shows that the measured alloy compositions are close to the nominal values, with the largest difference seen for the FePd sample as 2.5 at%. For simplicity, from this point onward the alloys will be referred to by their determined Ni content as x ¼ 0, 3.5, 5.2 and 7.5. As confirmed by XRD, all as-arc-melted samples consist of the fcc chemically-disordered phase, with no secondary phases observed. As an example, Fig. 1 presents X-ray diffractograms for the two extreme compositions (x ¼ 0 and 7.5) in the range 32 < 2q < 52 , where the fcc (111) and (200) Bragg peak reflections are observed. Upon annealing, the presence of the L10 phase is evidenced by the appearance of (001), (110), (201), (112), (221), (310), (203), and (312) superlattice Bragg peaks associated with chemical ordering, as well as by the splitting of the (200),
847
Fig. 1. XRD diffractograms obtained using Cu Ka radiation for as-arc-melted and annealed x ¼ 0 and 7.5 samples.
(220) and (311) fundamental Bragg reflections signaling the symmetry reduction from cubic to tetragonal. Fig. 1 displays XRD results for the annealed x ¼ 0 and 7.5 samples, where the (110)L10 superlattice peak characterizing the L10 phase is observed around 2q z 33 , with an accompanying (200) peak splitting in the interval 46 < 2q < 50 . While samples of Ni content x 3.5 have transformed into a nominal single-phase L10 structure upon annealing, annealed samples with x 5.2 are multiphase. Fig. 2 presents the X-ray diffractogram for these samples in the range 46 < 2q < 50 , illustrating tetragonal L10 (200) and (002) Bragg reflections along with additional peaks marked by an asterisk “*”. The additional peaks observed in the XRD data obtained from the x ¼ 5.2 annealed sample (Fig. 2a) correspond to an fcc-based phase with a lattice parameter that is the same, within experimental error, as that of the fcc phase in the x ¼ 5.2 as-arc-melted sample. Thus, this fcc phase has been tentatively attributed to a remnant fcc parent phase. Further, XRD data from the x ¼ 7.5 annealed sample (Fig. 2b) indicate the presence of two additional fcc-based phases, one with a lattice parameter that within experimental error corresponds to that of the x ¼ 7.5 as-arc-melted fcc phase, and the other with a unit cell volume that is 1.3% smaller than that of the x ¼ 7.5 as-arcmelted fcc phase. These phases have thus been tentatively attributed to a remnant fcc parent phase and a “new” fcc phase, respectively. The L10 phase fractions fL10 in the x ¼ 5.2 and 7.5 annealed samples are estimated as 95 vol% and 20 vol%, respectively, while the fcc phase fractions ffcc in the x ¼ 7.5 annealed sample are estimated as 20 vol% (parent phase) and 60 vol% (new phase). Fig. 3 illustrates trends in the structural parameters of the fcc asarc-melted and L10 annealed FePdNi alloys as a function of nickel content. The unit cell volumes decrease slightly with increased Ni
Table 1 Alloy nominal and measured compositions. Nom.
Fe Pd Ni
50 50
Measured
Nom.
Avg.
Std. Dev.
47.5 52.5
1 1
50 47 3
Measured
Nom.
Avg.
Std. Dev.
49.3 47.2 3.5
0.5 0.4 0.3
50 45 5
Measured
Nom.
Avg.
Std. Dev.
47.5 47.3 5.2
0.8 0.7 0.2
50 43 7
Measured Avg.
Std. Dev.
48.2 44.3 7.5
0.4 0.4 0.1
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Fig. 2. Multi-peak structures in the vicinity of the (200) diffraction peak in annealed x 5.2 samples. Peaks labeled with “*” are attributed to fcc phases.
Fig. 3. Variation with Ni content of unit cell volume, c/a ratio, LRO parameter, crystallite size (D) and strain (ε) of the fcc as-arc-melted and L10 annealed phases (lines drawn to guide the eye). Data points for LRO in x 5.2 annealed samples and crystallite size in the x ¼ 7.5 annealed sample have been omitted considering the large error associated with the calculated values.
concentration, outside of experimental error. Lattice parameter(s) obtained for the binary FePd composition in its fcc (a ¼ 3.82 Å) and L10 (a ¼ 3.862 Å, c ¼ 3.697 Å) states are consistent with values reported in the literature [18,32]. Ni additions to the L10 phase
cause an initial increase and then decrease in the c/a ratio at x ¼ 5.2, while the L10 LRO parameter decreases with 3.5 at% Ni additions. Data points for the LRO parameter of x 5.2 at% samples have been omitted from Fig. 3 on account of their large associated error, attributed to the de-convolution process of the individual L10 and fcc reflections in the multi-peak observed. The calculated crystallite sizes (D) for as-arc-melted fcc samples decrease from ~43 nm to ~ 20 nm with increased Ni concentration, Fig. 3. Upon annealing and formation of the L10 phase, an increase in the crystallite size is observed for x 5.2 samples. For the x ¼ 7.5 annealed sample, the crystallite size has a large associated error and thus this data point has been omitted from Fig. 3. The calculated lattice strain values (ε) are moderate, in the range 0.10% < ε < 0.25%. For a given composition, the strain manifest in the cubic fcc phase is the same, within experimental error, as that characterizing the tetragonal L10 phase. The room-temperature magnetization of all as-arc-melted fcc alloys increases rapidly with increased applied field, reaching saturation at anisotropy fields around HA z 3000 Oe (238.7 kA/m; see Fig. 4a). Measured MS values for these alloys (MS ¼ 1066e1110 emu/cc (1066e1110 kA/m)) are in good agreement with those reported in the literature for bulk Fe50Pd50 [1]. While nickel additions slightly increase MS for the fcc as-arc-melted phase (Fig. 4b), they have a negligible effect in the coercivity and remanence, with values HCi < 25 Oe (2 kA/m) and MR < 60 emu/cc (60 kA/m). In contrast, the magnetization of the annealed x 5.2 samples slowly increases with applied field without reaching saturation at 30 kOe (2387 kA/m), indicative of increased magnetocrystalline anisotropy. Anisotropy fields and saturation magnetization for these L10 alloys are larger than those of their fcc counterparts, with HA > 27,000 Oe (2148.6 kA/m) and extrapolated saturation values (shown in Fig. 4b) between MS ¼ 1160e1180 emu/cc (1160e1180 kA/m). The coercivity and magnetic remanence values for annealed L10-type samples (x 5.2) are in the range 94e160 Oe (7.5e12.7 kA/m) and 72e157 emu/cc (72e157 kA/m). For the x ¼ 7.5 annealed sample, a 2% increase in MS relative to that of the fcc as-arc-melted state is noted, while no significant change in HA, HCi or MR was detected despite noted crystallographic differences. All as-arc-melted fcc samples exhibit only a single calorimetric event of second-order phase transformation character that is associated with the Curie temperature TC. As an example, Fig. 5 presents DSC scans for the two extreme compositions (x ¼ 0 and 7.5, arc-melted). On the other hand, annealed L10-type samples (x 5.2) show multiple thermal features, an example of which can be seen in Fig. 5 for the x ¼ 0 annealed alloy; the first heating scan reveals the L10 Curie transition followed by a broad irreversible
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Fig. 4. a) Room temperature field-dependent magnetization of binary FePd as-arc-melted (fcc) and annealed (L10) sample. b) Evolution of MS with Ni content for as-arc-melted and annealed samples. Note that the annealed x ¼ 7.5 sample is plotted separately from the L10 trend, as it has a multiphase character with a low L10 volume fraction. The error bars for the x ¼ 7.5 annealed sample are smaller than the symbol.
Fig. 5. DSC scans of x ¼ 0 (green) and 7.5 (pink) samples in their as-arc-melted state (top) and annealed state (bottom). For the annealed sample, consecutive heating scans are labeled “1st scan” and “2nd scan”. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
endothermic event preceding a first-order endothermic transition associated with the L10/fcc order-disorder phase transformation to produce a “recovered” fcc phase. A subsequent heating scan (2nd scan in Fig. 5) reveals only a single thermal event associated with the Curie transition of the recovered fcc phase, where TC(recovered fcc) > TC(as-arc-melted fcc) for x > 0 alloys. For the x ¼ 7.5 annealed sample, only a Curie transition is detected in calorimetry (Fig. 5), despite the fact that the XRD data evidenced approximately 20 vol% of the L10 phase in this sample; no measurable enthalpy associated with the chemical order-disorder transformation phase was obtained. Addition of Ni to the fcc and L10 lattices increases the Curie temperature, Fig. 6, with a Curie temperature difference between as-arc-melted and recovered fcc phases that increases with increased Ni content. A similar trend is observed between the saturation magnetization of as-arc-melted and recovered fcc phases (not shown), with a maximum increase of 2.1% exhibited by the recovered fcc x ¼ 5.2 sample. Increased Ni content produces a decrease in the onset temperatures and the associated enthalpies of the L10/fcc phase transition, Fig. 6. For reference, Fig. 6 includes an
Fig. 6. Evolution of the Curie temperature of L10 and fcc phases and of the onset temperature and enthalpy of the order-disorder transformation with Ni content. An extrapolated TOD value calculated for the x ¼ 7.5 alloy is plotted for reference.
extrapolated value of TOD (x ¼ 7.5) obtained from a polynomial fit to the experimental TOD values as a function of Ni content. 4. Discussion Structural differences between the as-arc-melted and annealed FePdNi alloys are reflected in their magnetic behavior. Attainment
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of the L10 phase in x 5.2 annealed samples is accompanied by an increase in both intrinsic (i.e., MS, TC) and extrinsic (i.e., HCi, MR, HA) magnetic properties. Annealing of the x ¼ 7.5 sample contributes to a moderate increase in MS and TC, while no detectable changes in HCi, MR, or HA are observed. This moderate increase in MS and TC, that deviates largely from the trends observed for the L10 phase in other samples, is attributed to the multi-phase character of the x ¼ 7.5 annealed sample, with a low L10-phase fraction (20 vol%). In analogy, noted differences in TC and MS between the as-arc-melted and DSC-recovered fcc phases allude to subtle structural differences between them; it is proposed that retention of a small degree of short-range order occurs after the material undergoes the L10/fcc phase transition upon heating in the DSC. The detection of a very broad endothermic feature preceding the main chemical order-disorder peak in FePdNi alloys is consistent with a two-stage disordering phase transformation (L10/fcc). While broad endothermic signals that precede a chemical orderdisorder transformation have been reported for thermoresistometry [51] and calorimetry [52] experiments conducted on chemically-ordered CuAu samples, their origin remains unexplained. In this current study, the broad endothermic peak is associated with atomic short-range motion to produce a readjustment of the LRO parameter below the temperature at which the nucleation and growth of the chemically-disordered fcc phase may be measured. In this context, the total enthalpy of transformation DHL10/fcc is the sum of both endothermic events. The phases detected in all annealed (T ¼ 773 K) samples are in qualitative agreement with those presented in the calculated (T ¼ 823 K) FePdNi ternary diagram of Horiuchi et al. [38]. Horiuchi's phase diagram indicates a single L10 phase centered around the Fe50Pd50 composition next to a two-phase “disorder” þ L10 region, Fig. 7a. The solubility limit of Ni in a single L10 phase at a fixed Fe content (~50 at%) is provided by Horiuchi as z 5 at%. Consistent with this scenario, in the present study a single L10 phase was detected in annealed x 3.5 samples. For the x ¼ 5.2 annealed alloy, the transformation did not go to completion, with a small volume fraction (~5 vol%) of the parent fcc phase remaining. It is anticipated that this alloy should transform into a single L10 phase if sufficient time is allowed to achieve equilibrium. For the x ¼ 7.5 annealed alloy, the observed phases are also consistent with Horiuchi et al.'s proposed diagram, for which a two-phase mixture is expected to exist in equilibrium; these two phases correspond to the observed “new” fcc and L10 phases, while the presence of a remnant fcc parent phase indicates that the transformation was not
complete. This apparently good agreement between the current results and those of Horiuchi et al. should be considered with caution as (i) the calculations of Horiuchi et al. [38] predict TOD ¼ 800 K for binary L10 FeNi, significantly higher that the experimentally determined value of 320 K [7,52,53] and (ii) the L10 FePd boundary in Horiuchi et al.'s calculated ternary phase diagram does not agree with the published FePd binary phase diagram [18]. Nonetheless, Horiuchi et al.'s proposed L10 phase boundary near the 50 at% Fe composition has been used to estimate the Ni concentrations of the phases present in the x ¼ 7.5 annealed alloy. By invoking assumptions that mass fraction is equivalent to volume percent and that the L10 phase and the parent disordered fcc alloy possess the same Fe concentration (48.2 at%), the Ni concentrations of the phases present in the x ¼ 7.5 annealed alloy are estimated as x ¼ 6 at% (L10 phase) and as x ¼ 8 at% (new fcc phase), (Fig. 7b). This information allows determination of a very narrow fcc þ L10 twophase region in the composition range 6 at% < Ni < 8 at% for 48.2 at% Fe at 773 K. Using the Ni content limits determined here for the fcc þ L10 two-phase region, along with phase equilibria information at 773 K obtained from the binary FeNi [54], NiPd [55], and FePd [18] phase diagrams, an updated isothermal section of the ternary FePdNi phase diagram is proposed in Fig. 8. It is noted that the phase boundaries presented for single-phase regions other than L10 are arbitrary and are included for illustration purposes only. Subjecting the Ni ¼ 5.2 and Ni ¼ 7.5 alloys to longer annealing times should prove useful in the determination of the L10 and the fcc þ L10 phase boundaries with greater certainty. Assignment of a single-phase L12 region extending across the full composition range is made based on the complete mutual solubility of Ni and Pd in the fcc phase (g) in the NiePd binary system. In the proposed isothermal section, an fcc phase is shown to be more stable than an L10 phase at 773 K for Fe50Pd50-xNix alloys with higher Ni contents than studied here. However, only experimental results from FePd alloys made with higher Ni contents can confirm whether a stable or metastable L10 region can be found near the Fe50Ni50 composition at this temperature. 5. Conclusions This study confirms that Ni additions to FePd significantly impact the stability of the L10 phase in the FePdNi ternary system, reflecting important changes in the structural characteristics and magnetic properties of these alloys to inform future synthesis of the
Fig. 7. a) FePdNi ternary phase diagram at 823 K proposed by Horiuchi et al. [38], along with the alloy compositions of the present study. (b) Estimated compositions of the multiple phases present in the x ¼ 7.5 annealed sample, namely the parent fcc phase (blue) and the product L10 (red) þ new fcc (green) phases. Horiuchi et al.'s L10 boundary is left for reference as a dashed line as it was used to estimate the compositions. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
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[2] [3]
[4]
[5]
[6] [7] [8]
[9]
[10] [11] Fig. 8. 773 K isothermal section proposed for the updated FePdNi ternary phase diagram (a represents the body centered cubic crystal structure while g represents an fcc phase). L10 and fcc þ L10 phase boundaries have been selected to match experimental observations, while phase boundaries displayed for other regions are arbitrary.
alternative permanent magnet candidate tetrataenite. Comparison of structural data and phase determinations obtained here with those reported by Horiuchi et al.'s calculated ternary FePdNi phase diagram [38] reveal that samples with nickel content x 3.5 at% achieve equilibrium after 100 h of annealing, while longer annealing times are required to form the L10 phase in compositions of higher Ni content. The sluggish kinetics observed for x 5.2 at% alloys result from a decreased fcc/L10 phase transformation driving force that is attributed to a reduced order-disorder temperature upon Ni addition. Furthermore, a narrow fcc þ L10 twophase region is proposed for the range 6 at% x 8 at% in Fe48.2Pd51.8xNix alloys. Study of the reverse phase transformation (L10/fcc) that occurs upon heating the annealed alloys of this study reveals a two-stage disordering process that initiates with short-range atomic motion to readjust the LRO parameter prior to the first-order chemical disordering transition. Additionally, retention of a certain degree of short range order was observed after undergoing the L10/fcc transition, which is amplified with increased Ni content. This confirms that Ni additions have a considerable effect on the kinetics of the order-disorder phase transformation. The results of this work indicate that L10 formation in the ternary Fe50Pd50-xNix is highly sensitive to slight compositional variations. It is therefore necessary to investigate ternary alloys with higher Ni contents than studied here to confirm if an L10 phase can be found near the FeNi boundary, thus allowing kinetic access to the highly attractive ferromagnetic L10 FeNi phase, relevant for the next-generation permanent magnets.
[12] [13]
[14]
[15]
[16]
[17]
[18] [19]
[20] [21]
[22]
[23]
[24]
[25]
[26]
Acknowledgments [27]
This study was supported by the National Science Foundation eCMMI Divisione under grants No. 1129433 and No. 1259736, and by Northeastern University.
[28]
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