Metals 2014, 4(1), 8-19; doi:10.3390/met4010008

Article
Magnetism-Structure Correlations during the ε→τ Transformation in Rapidly-Solidified MnAl Nanostructured Alloys
Felix Jiménez-Villacorta 1, Joshua L. Marion 1, John T. Oldham 1, Maria. Daniil 2, Matthew A. Willard 2 and Laura H. Lewis 1,*
1
Department of Chemical Engineering, Northeastern University, 360 Huntington Ave., Boston, MA 02115, USA; E-Mails: fjv2003@gmail.com (F.J.-V.); marion.jo@husky.neu.edu (J.L.M.); jtoldham10@gmail.com (J.T.O.)
2
Department of Materials Science and Engineering, Case Western Reserve University, Cleveland, OH 44106-7204, USA; E-Mails: mxd421@case.edu (M.D.); maw169@case.edu (M.A.W.)
*
Author to whom correspondence should be addressed; E-Mail: lhlewis@neu.edu; Tel.: +1-617-373-3419.
Received: 13 December 2013; in revised form: 14 January 2014 / Accepted: 16 January 2014 /
Published: 21 January 2014

Abstract

: Magnetic and structural aspects of the annealing-induced transformation of rapidly-solidified Mn55Al45 ribbons from the as-quenched metastable antiferromagnetic (AF) ε-phase to the target ferromagnetic (FM) L10 τ-phase are investigated. The as-solidified material exhibits a majority hexagonal ε-MnAl phase revealing a large exchange bias shift below a magnetic blocking temperature TB~95 K (Hex~13 kOe at 10 K), ascribed to the presence of compositional fluctuations in this antiferromagnetic phase. Heat treatment at a relatively low annealing temperature Tanneal ≈ 568 K (295 °C) promotes the nucleation of the metastable L10 τ-MnAl phase at the expense of the parent ε-phase, donating an increasingly hard ferromagnetic character. The onset of the ε→τ transformation occurs at a temperature that is ~100 K lower than that reported in the literature, highlighting the benefits of applying rapid solidification for synthesis of the rapidly-solidified parent alloy.
Keywords:
nanostructured magnetic materials; Mn-based magnets; rapid-solidification

1. Introduction

Current market and supply pressures on rare earth elements—components of technologically important advanced permanent magnets—are driving investigation into magnetic materials that can replace rare-earth-based alloys with more abundant and less strategically important elements. A promising strategy for realization of new concepts in permanent magnetic materials is to revisit the “classic” hard magnetic alloys that were widely used before the advent of the rare-earth-based supermagnets [1,2] and, by employing advanced fabrication and analysis methods, gain insight into the fundamentals governing their structure-magnetic property behavior for improved performance. Among the list of materials targeted for this purpose include intermetallic ferromagnetic Mn-based alloys, such as MnAl [3,4] or MnBi [5,6], which were considered as permanent magnet candidates in the 1960s and 1970s [7,8,9]. In particular, the MnAl compound with the tetragonal L10 type crystal structure (τ phase) combines favorable aspects, such as low density, large corrosion resistance, and low cost [10,11], with good permanent magnetic performance. The MnAl magnetic properties include a Curie temperature of TC~525 K, a robust magnetocrystalline anisotropy, K1~106–107 J/m3 (K1~107–108 erg/cm3) that fosters appreciable room-temperature coercivity values, Hc~0.15 to 0.4 T (Hc~1.5 to 4.0 kOe), and a relatively large saturation magnetization of Ms~88–98 emu/g that has enabled experimentally-realized maximum energy products of (BH)max~16–56 kJ/m3 ((BH)max~2–7 MGOe) [12,13]. While some success has been recently achieved in the development of these materials [3,14,15] experience demonstrates that optimization of the microstructure and the magnetic features of new magnetic materials may take several years (or even decades) [4], making a strong case for renewed investigation of structure-property relationships in former generations of magnetic materials [16,17,18].

Non-equilibrium processing techniques, such as melt-spinning, splat-quenching, and mechanical milling, allow access to thermodynamically metastable phases and microstructures that can be beneficial, including nanoscaled features and enhanced magnetic properties [19,20]. Within the Mn-Al phase diagram, the ferromagnetic (FM) τ-MnAl phase forms via a two-step reaction process that originates from the parent hexagonal ε-MnAl phase through the intermediate B19-structured ε'-phase, by annealing at temperatures in the range 723–823 K [21,22]. These transformations are assumed to be compositionally invariant, controlled by nucleation and growth processes [11,23] that commence at the grain boundaries [24,25]. Hence, MnAl fabrication methods that provide a refined microstructure with small ε-crystallites are envisioned to yield enhanced surface-to-volume ratios that promote formation of the τ-phase, with lowered phase transformation energy barriers and lower values for the transformation temperatures. Recent reports on ε–phase MnAl thin films (5 nm thickness) [26,27] and on ε–phase MnAl ultrafine nanocrystallites (~8.2 nm) obtained by mechanical milling [28] reveal an ε→τ transformation temperature in the range ~623–673 K, confirming the hypothesis that the temperature required for τ-phase formation is dependent on details of the microstructure.

In this work, the annealing-induced evolution of magnetic and structural properties of nanostructured, rapidly-solidified Mn55Al45 are evaluated at early stages of the ε-MnAl to τ-MnAl transformation confirmed to begin at T~568 K (295 °C), 50–100 degrees lower than previously-reported transformation onset temperatures. The evolution of the magnetic parameters of these heat-treated MnAl ribbons is correlated with phase transformation and microstructural attributes. Clarification of the character of the ε→τ transformation is anticipated to lead to tailoring of existing pathways, or identification of new pathways, for stabilization of the L10-type τ-phase, enabling the application of advanced processing methods, and associated improved performance as a rare-earth-free permanent magnetic material.

2. Experimental Section

Arc-melted ingots of nominal composition Mn55Al45 made from initial charges of Mn granules (Alfa Aesar, 99.98% purity, metals basis) and Al foil (Alfa Aesar, 99.9% purity, metals basis) were used as precursors for melt-spinning under a He atmosphere at a pressure of 0.3 atm, with a tangential velocity of 64 m/s. The rapidly-quenched alloys produced by this process were shiny, brittle, approximately 50 µm thick, and 1 to 5 cm in length. The average composition of the as-solidified ribbons was determined to within about 2 at.% using scanning electron microscopy—energy-dispersive X-ray spectroscopy (SEM–EDX; Hitachi S4800, Dallas, TX, USA). Crystal structures were confirmed with X-ray diffraction (XRD; Philips PANalytical X’Pert-Pro, Westborough, MA, USA) using Cu-Kα radiation. The Kα2 contribution to the diffracted X-rays was removed by a software-based correction algorithm. Lattice parameters were calculated with a least-squares cell parameter program [29], and crystallite sizes were determined from the Scherrer equation [30]. The phase transformation character of the alloy was investigated using Differential Scanning Calorimetry (DSC; Netzsch STA 449, Burlington, MA, USA) under an Ar atmosphere using heating rates of 5 K/min.

Annealing procedures were performed on samples of the as-quenched Mn55Al45 alloy sealed in thin vitreous silica tubes under a vacuum pressure of 1 × 10−6 Torr to protect against oxidation. Two samples were placed at the same location in the center zone of a programmable tube furnace (MTI GSL1100X, Richmond, CA, USA) and simultaneously annealed in sequential 30-min intervals in the nominal temperature range 373 K ≤ Tanneal ≤ 618 K (the actual Tanneal was within 2 K of the target value). After each heat treatment interval, the magnetic properties of one of the vacuum-sealed sample were analyzed. Magnetic measurements were carried out using a superconducting quantum interference device magnetometer (SQUID; Quantum Design MPMS, San Diego, CA, USA). Zero-field-cooling–field-cooling (ZFC-FC) magnetization vs. temperature curves (M vs. T) were collected in the temperature range 10 K < T < 400 K, using a probe field of 0.1 T. Magnetization loops (M vs. H) were collected between 10 K and 300 K under ZFC and FC conditions (applied field during cooling: HFC = 5 T). All magnetization measurements were corrected for demagnetization effects; each specimen was approximated as a cylinder and the demagnetization factor (Nd) was determined from the plot of Nd vs. sample aspect ratio [31]. After observation of a significant variation in the magnetic features of the first sample, which occurred at Tanneal = 568 K (295 °C), the first sample was removed from the silica tube for XRD investigation. Thermal treatment and magnetic examination procedures continue with the second sample up to Tanneal = 618 K (345 °C); after this annealing step this second sample was removed from the silica tube and subjected to XRD examination to obtain phase and crystal structure information.

As discussed more thoroughly in Section 3, magnetic hysteresis loops obtained from the samples often indicated the presence of more than one magnetic phase. Contributions from different magnetic components were extracted from the hysteresis curves collected at 10 K and at 300 K, using an empirical equation that describes the magnetization response of a ferromagnetic (FM) sample comprised of n independently-switching magnetic phases [32]:

Metals 04 00008 i001
where M = total magnetization (emu/g), H = effective internal magnetic field (Oe), MS = total saturation magnetization of the multi-phase material (emu/g), Metals 04 00008 i002 = effective intrinsic coercivity of magnetic phase-(i) (Oe), Metals 04 00008 i003 squareness ratio (remanence vs. saturation; unitless), xi = relative fractional contribution of each constituent magnetic phase respect to the total magnetization ( Metals 04 00008 i004), and n = number of constituent magnetic phases contributing to M (unitless).

3. Results

Melt-spun Mn55Al45 ribbons analyzed by XRD (Figure 1) reveal a majority hexagonal (hcp) phase, identified as ε-MnAl [28,33]. As reported in detail elsewhere [34], each hcp ε-MnAl Bragg peak occurs in pairs, consistent with the presence of two phases with the ε-MnAl crystal structure, denoted as ε1 and ε2, with similar unit cell lattice parameters around a~2.70 Å and c~4.37 Å and identical c/a ratios ~ 1.62. The ε1 and ε2 phases possess unit cell volumes of V~82.9 ± 0.3 Å3 and V~82.1 ± 0.5 Å3, respectively, and average crystallite sizes (Dε) of approximately 30 and 60 nm, as calculated from the Scherrer formula (Figure 2). Therefore, the ε1 phase has a larger unit cell with a significantly smaller grain size than the ε2 phase. The remaining minority Bragg peaks in the XRD pattern are associated with a rhombohedral γ2-MnAl phase with lattice parameters a = 10.1 ± 0.2 Å, b = 7.63 ± 0.2 Å, c = 16.0 ± 0.3 Å, and unit cell volume V~1230 ± 40 Å3.

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Figure 1. XRD patterns from Mn55Al45 as-spun ribbons and those annealed at 568 K and 618 K.

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Figure 1. XRD patterns from Mn55Al45 as-spun ribbons and those annealed at 568 K and 618 K.
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Figure 2. Thermal evolution of crystallite size (D) of both hexagonal ε-MnAl regions—ε1 (filled circles) and ε2 (open diamonds)—in melt-spun Mn55Al45 ribbons. Dashed lines drawn between markers to guide eye.

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Figure 2. Thermal evolution of crystallite size (D) of both hexagonal ε-MnAl regions—ε1 (filled circles) and ε2 (open diamonds)—in melt-spun Mn55Al45 ribbons. Dashed lines drawn between markers to guide eye.
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The coherently-diffracting size of the hcp ε-phases in the Mn55Al45 ribbons varies only slightly with annealing to Tanneal = 568 K, and then reaches similar values (Dε~45 nm) upon heating to 618 K. This results is consistent with the ε1 phase growing at the expense of the ε2 phase with annealing. At the higher annealing temperature an additional phase identified as the metastable tetragonal τ-MnAl phase with L10-type structure is visible in the X-ray diffraction scans. The τ-phase possesses unit cell lattice parameters a = 2.77 ± 0.01, c = 3.56 ± 0.01 Å, unit cell volume V = 27.3 ± 0.1 Å3, in close agreement with previous literature [8,14,28], and approximate crystallite size Dτ~20 nm.

The magnetic response of the melt-spun MnAl alloys is also modified by annealing. The ZFC-FC M(T) curves of the as-quenched material (Figure 3) exhibit a cusp at Tpeak = 95 K, with a divergence in the magnetization for T < Tpeak that is associated with the Néel temperature of the AF ε-phase [35]. Both, the cusp and the magnetization divergence, vanish with annealing to Tanneal = 618 K, accompanied by a significant increase of the magnetization. Field-cooled hysteresis curves measured from the as-spun ribbons at T = 10 K (Figure 4a) display a prominent shift along the applied field axis of Hex~1.3 T (Hex~13 kOe). The field-axis shift, not present under ZFC conditions, decreases with increased measurement temperature T and vanishes for TTpeak. Further, it is noted that the hysteresis loops displayed in Figure 4a feature a small step at low fields in the ascending and descending branches of the curves, indicative of a lower-coercivity Hci~0.15 T (Hci~1.5 kOe) phase present in the sample in addition to the majority high-coercivity phase, Hci~1.9 T (Hci~19 kOe) phase. Subsequent annealing at moderate temperatures (see hysteresis loop for the Tanneal = 568 K sample in Figure 4a) yields a decrease of the hysteresis loop shift previously noted at low temperature, accompanied by an increased magnitude of the second phase magnetization at room temperature (Figure 4b). Further annealing suppresses the contribution corresponding to the shifted hysteresis loop, and only the second phase is observed at both T = 10 and 300 K, revealing an increased coercivity at both temperatures to a value of Hci~0.25 T (Hci~2.5 kOe).

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Figure 3. M vs. T curves collected under a probe field of 1 kOe from as-spun Mn55Al45 ribbons (filled circles) and those treated at 618 K (open diamonds).

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Figure 3. M vs. T curves collected under a probe field of 1 kOe from as-spun Mn55Al45 ribbons (filled circles) and those treated at 618 K (open diamonds).
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Figure 4. FC magnetization loops at (a) T = 10 K and (b) T = 300 K for as-spun Mn55Al45 ribbons (filled circles) and those annealed at Tanneal = 568 K (open diamonds) and Tanneal = 618 K (filled triangles), showing presence of two magnetic phases. Ribbons treated at 618 K display Ms~18 emu/g.

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Figure 4. FC magnetization loops at (a) T = 10 K and (b) T = 300 K for as-spun Mn55Al45 ribbons (filled circles) and those annealed at Tanneal = 568 K (open diamonds) and Tanneal = 618 K (filled triangles), showing presence of two magnetic phases. Ribbons treated at 618 K display Ms~18 emu/g.
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Finally, DSC data (presented in Figure 6c for discussion in Section 4) reveal a dominant exothermal peak centered at T = 654 K, associated with the annealing-induced ε→τ transformation, although a more complex trace is observed in the lower temperature region, where an inflection point at around T~550 K is detected. The onset observed by DSC agrees with the evolution observed by magnetic characterization, revealing a higher sensitivity for small presence of emerging phases compared to X-ray diffraction results.

4. Discussion

Analysis of the structural and magnetic data obtained from rapidly-solidified Mn55Al45 ribbons provides insight into the evolution of the ε-phase to τ-phase transformation in these nanostructured alloys. The as-quenched ribbons reveal metastable retention of the hexagonal (hcp) ε-MnAl phase, distributed into nanoscaled regions that are comparatively Mn-rich (ε1) and Mn-poor (ε2) upon rapid solidification from the melt [34]. Areas that are more concentrated in Mn contain a greater number of Mn-Mn nearest-neighbor pairs with AF character, while regions of lower Mn content contain more Mn-Mn next-nearest-neighbor pairs, with ferromagnetic character [36,37]. Strong interactions between these two regions lead to large exchange bias (Hex) values at T < Tpeak, analogous to the reported mictomagnetic character of AgMn or CuMn alloys in which compositional fluctuations and strong coupling between Mn-rich and Mn-poor areas yield the formation of exchange-biased loops [36,38]. Employing the assumptions of an atomic radius for Al of ~1.25 Å and ~1.40 Å for Mn [39], and utilizing the calculated unit cell volumes of the ε1 and ε2 phases in the alloy (Section 3) obtained by X-ray diffraction, it is determined that the Mn-rich phase contains ~64 ± 3 at.% Mn, whereas the Mn-poor phase contains ~60 ± 3 at.% Mn, within the Mn content range for the ε-phase proposed by Liu et al. [40]. These high Mn concentrations, as compared to the nominal Mn content, are consistent with the coexistence of the ε-phases with the lower-Mn-content γ2-phase of approximate composition 35–45 at.% Mn.

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Figure 5. Modeling of the hysteretic demagnetization curve obtained at 10 K from Mn55Al45 ribbons annealed at 568 K (open circles) into two magnetic phases (dashed lines), in accordance with Equation (1).

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Figure 5. Modeling of the hysteretic demagnetization curve obtained at 10 K from Mn55Al45 ribbons annealed at 568 K (open circles) into two magnetic phases (dashed lines), in accordance with Equation (1).
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The fraction of magnetic phases produced upon annealing was determined in the MnAl ribbon from hysteresis loop measurements obtained at 10 K, according to Equation (1) with typical results illustrated in Figure 5. Two magnetic phase contributions are noted: the first contribution is attributed to the exchange-biased magnetization response, predominantly driven by the hexagonal ε-phase (labeled as xε), and the second magnetic contribution is associated with the FM τ-MnAl phase (denoted as xτ). As the magnetic features are well defined at low temperature, data collected at 10 K were used to display the evolution of the coercivity Hci and the relative phase fraction of each phase at 10 K, Figure 6a,b, respectively. The decrease of the exchange bias field Hex and the coercivity Hci noted for Tanneal ≥ 473 K (Figure 6a), along with the simultaneous increase in the magnetization of the ε-MnAl phase, suggests a reduction of antiparallel spin arrangement between nearest neighbors of the hcp ε-phase regions with annealing attributed to Mn homogenization in the alloy. The increased coercivity of the τ-MnAl phase in the annealing temperature range Tanneal ≈ 523-568 K, reaching a stable value of Hci~0.25 T (Hci~2.5 kOe), could be attributed to higher anisotropy associated with improved chemical order of the tetragonal L10 phase, as well as improved microstructural features.

The contribution of the τ-phase (xτ) grows at the expense of the ε-phases (xε) at Tanneal ≥ 568 K, as confirmed by the T = 10 K data trend (Figure 6b). Further treatment at Tanneal = 618 K causes a precipitous rise in xτ from approximately 15% to 90% of the total magnetization. DSC data (Figure 6c) indicates the onset of phase transformation, by the inflection point at around T~530–570 K, right below the exothermal peak centered at T = 654 K signaling the massive ε→τ phase transformation, in agreement with the onset observed by magnetic examination. Assuming a room-temperature saturation magnetization value Ms~98 emu/g for the τ-MnAl phase [3,12], it is estimated that the τ-MnAl content increases from 0.08 wt. % to 0.28 wt. % upon annealing in the range Tanneal = 523–568 K. Subsequent heat treatment at Tanneal = 618 K causes a significant increase in the τ-phase contribution from 0.28 wt. % to at least 18.0 wt. %, confirmed by the qualitative increase of magnetization featured by the alloy (Ms~18 emu/g) and the XRD evaluation of its crystallographic constituents. Based on the magnetic measurements, it can be concluded that nucleation of the τ-MnAl phase seems to occur at Tanneal ≈ 568 K, around 100 K lower than previously-reported values for the onset of the ε→τ transformation in quenched MnAl alloys, and similar to those observed in nanostructured samples, as summarized in Table 1. The significantly-reduced τ-MnAl nucleation temperature confirmed in this work is thus attributed to the refined ε-MnAl nanocrystalline structure and a large amount of crystal defects featured by rapid solidification via melt-spinning to directly obtain the ε-MnAl parent phase, opening potential pathways to develop energy-efficient routes to fabricate MnAl permanent magnets.

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Figure 6. Thermal evolution of (a) Hci and (b) relative fraction of the hysteretic magnetization (xi) contributed by the ε-MnAl phase regions and τ-MnAl at 10 K, determined by Equation (1); dashed lines were drawn to guide eye; (c) DSC trace of as-spun Mn55Al45 ribbon showing τ-MnAl formation (Inset: first derivative of the DSC plot).

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Figure 6. Thermal evolution of (a) Hci and (b) relative fraction of the hysteretic magnetization (xi) contributed by the ε-MnAl phase regions and τ-MnAl at 10 K, determined by Equation (1); dashed lines were drawn to guide eye; (c) DSC trace of as-spun Mn55Al45 ribbon showing τ-MnAl formation (Inset: first derivative of the DSC plot).
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Table 1. Reported initial formation temperature of the τ-MnAl phase as a function of material processing method.

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Table 1. Reported initial formation temperature of the τ-MnAl phase as a function of material processing method.
Fabrication/processing methodτ-MnAl formation temperature (K)MethodReferences
Water-quenched Mn53Al46C2723 KAnnealed[23]
Oil-quenched Mn55Al45800 KAnnealed[13]
Nonequilibrium synthesis methods
Splat-quenching~700 KDSC (onset)[20]
723 KDSC (peak)[20]
Mechanical milling~600 KDSC (onset)[28]
665 KDSC (peak)[28]
623 KXRD[28]
Melt-spun Mn54A l44C2 (at 25 m/s)~725 KDSC (onset)[27]
779 KDSC (peak)[27]
748 KM(T) curves[27]
Melt-spun Mn55Al45 (at 40 m/s)~710 KDSC(onset)[19]
758 KDSC(onset)[19]
Melt-spun Mn55Al45 (at 64 m/s)~550 KDSC (onset)This work
654 KDSC (peak)This work
618 KXRDThis work
568 KM(H) curvesThis work

5. Conclusions

This work provides insight into the phase transformation of rapidly-solidified nanostructured Mn55Al45 alloys, from the metastable hexagonal ε-phase to the L10-ordered τ-phase, by means of detailed annealing procedures. The as-spun ribbons are confirmed to exhibit fluctuations in the local Mn content, causing phase separation into antiferromagnetic Mn-rich regions of the hexagonal ε-MnAl phase and Mn-poor regions that possess ferromagnetic character; exchange interactions between these regions result in large exchange bias shifts of approximately ~13 kOe at T = 10 K. Heat treatment at moderate temperatures (Tanneal ≥ 473 K) causes Mn diffusion and homogenization throughout the matrix, reducing the exchange bias effect. The ε→τ transformation revealed by magnetic measurements starts after annealing at Tanneal ≥ 568 K (295 °C), signifying the onset of the tetragonal FM τ-MnAl nucleation from the hcp ε-MnAl parent phase at about 100 K lower than indicated by previous reports. The fine grains of the ε-MnAl parent phase are thought to be responsible for this large reduction in transition temperature. Understanding the nature of the ε→τ transformation could enable further reduction of the transformation temperature, which may play an impactful role for improving fabrication and processing techniques and strategies, and engineering of L10-type τ-MnAl permanent magnets.

Acknowledgments

This research has been funded by the Office of Naval Research (ONR), under the auspices of grant no. N00014-10-1-0553 and by ARPA-E Award# DE-AR0000186 (J.T.O.).

Conflicts of Interest

The authors declare no conflict of interest.

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