Structure and magnetic properties of melt-spun Mn-Ga-Cu-Al ribbons

In this work, we investigated structure and magnetic properties of Mn65Ga20Al15−xCux (x = 0, 5, 10 and 15) alloy ribbons prepared by melt-spinning method combined with annealing. The annealing temperature was varied from 250 °C to 350 °C, and the annealing time was changed from 5 h to 20 h. Concentration of Cu and annealing process significantly influence on the formation of the desired phases in the alloy ribbons. The D022-type Mn3Ga crystalline phase with the hexagonal structure, which characterizes hard magnetic property of Mn-Ga based alloys, is enhanced after an appropriate annealing process. The change of grain size after annealing also contributes to the increased coercivity of the alloy ribbons. The highest coercivity of 12.9 kOe and saturation magnetization of 18.7 emu g−1 are achieved on the alloy ribbons with Cu concentration of 10%. The simultaneous enhancement of these magnetic parameters has an important significance for application possibility of the Mn-Ga based alloys.


Introduction
With good magnetic properties, rare earth magnets (Nd-Fe-B and Sm-Co) are widely used in practice [1][2][3]. However, the issue of the exploitation and utilization of rare earth elements is a global concern because of the scarcity, the monopoly of the nations, the environmental pollution... [4][5][6]. Therefore, the technological research to avoid using rare earth in the hard magnetic materials is essential. Up to now, a number of rare earthfree hard magnetic materials such as Co-Zr, Co-Hf, Fe-Pt, Fe-Ni, Mn-Bi, Mn-Al... have been found and they are under development to meet application requirements in practice [7][8][9][10][11][12][13][14][15][16][17][18][19][20]. Among the developing rare earth-free hard magnetic materials, Mn-based alloys (Mn-Bi, Mn-Ga and Mn-Al) have shown a great potential because of their low cost and good magnetic properties. Each Mn-based hard magnetic alloy has its advantage properties. Mn-Bi alloys have the possibility for very high coercivity due to large magnetocrystalline anisotropy (K u = 1.6 MJ m −3 ) of the low temperature phase MnBi (LTP-MnBi). Besides, their theoretical maximum energy product (BH) max reaches 16.2 MGOe [18]. Notably, the coercivity of the Mn-Bi alloys increases with raising temperature leading to possibility for applications in high temperature conditions [16,17]. However, the formation of the LTP-MnBi is very difficult, requiring a long annealing time [21,22]. With magnetocrystalline anisotropy K u = 1.7 MJ m −3 , saturation magnetization M s = 70 emu g −1 and Curie temperature T C = 653 K of τ-MnAl phase, Mn-Al alloys have potential to create permanent magnets with their theoretical maximum energy product of 19 MGOe, but it is also very difficult to create a hard magnetic phase for these alloys [19,23,24]. For Mn-Ga alloys, their magnetic properties are better than those of the Mn-Al alloys. Besides, the natural abundance of Ga is more than that of Bi [18]. Like the Mn-Bi and Mn-Al alloys, the phase formation in the Mn-Ga system is rather complicated due to the existence of many intermetallic phases such as MnGa 2 , Mn 3 Ga, Mn 5 Ga 7 , Mn 2 Ga 5 , Mn 8 Ga 5 K [25]. Among them, the D0 19 -Mn 3 Ga (hexagonal), D0 22 -Mn 3 Ga (tetragonal) and Cu 3 Au-type (cubic) phases are mainly formed [26]. The D0 19 antiferromagnetic phase is easily formed in bulk alloys [27]. Any further distribution of this work must maintain attribution to the author(s) and the title of the work, journal citation and DOI.
Meanwhile, the D0 22 ferrimagnetic phase is obtained by annealing at temperature in range of 350°C-450°C [26,28]. The D0 22 phase has low saturation magnetization (M s ∼ 43 emu g −1 ) due to the antiparallel arrangement of magnetic moments in ferrimagnetic ordering. However, with high T C (∼762 K), large K u (∼2 MJ m −3 ) and relative high maximum energy product (BH) max (∼28 MGOe) [18,29] of this phase, Mn-Ga alloys can be widely used in magnetoelectronic devices [30][31][32][33][34]. The creation of this desired phases strongly depends on composition and fabrication conditions of the alloy. Up to now, various forms of Mn-Ga alloys including bulk, ribbon, film and nanoparticle have been fabricated by respective methods of arc-melting, rapidly quenching, sputtering and high energy ball milling [35][36][37][38][39][40][41][42][43][44][45]. With bulk samples prepared by arc-melting method, the desired magnetic phases are usually formed with long annealing time, up to several days [28,35,36]. Meanwhile, for ribbons prepared by rapidly quenching method, the annealing time is only a few hours [26,37,38]. The fabrication of the material with a large amount and easy control of the desired microstructure is the advantage of rapidly quenching method. The possibility to create hard magnetic phases with different structures in alloy ribbons is favorable for studying the relationship between structure and magnetic properties of materials. Early works indicated that the magnetic properties of Mn-Ga alloys are significantly enhanced by adding some elements (Al, CuK) and applying appropriate technology conditions. According to report of T. Saito et al [39], by partial replacement of Al for Ga in the Mn 65 Ga 35−x Al x (x = 0-25) alloy ribbons annealed at 700°C for 1 h, the coercivity increases from 5 kOe for un-added alloy to 9 kOe for 20% Al-added alloy. The addition of Al reduces the particle size, leading to an improvement of the coercivity of this alloy. Meanwhile, Cu is believed to widen the region of the D0 22 phase [42]. Therefore, the coercivity is significantly enhanced to 14.8 kOe for the Mn 65 Ga 25 Cu 10 ribbon annealed at 300°C for 10 h. In our previous study [45], the coercivity of Mn 65 Ga 25-x Al 10+x (x = 0, 5 and 10) alloy ribbons was improved to 10.5 kOe when concentration of Al is 15% (x = 5). In this work, the influence of Cu concentration and the annealing process on the structure and magnetic properties of the Mn 65 Ga 20 Al 15-x Cu x (x = 0, 5, 10 and 15) alloy ribbons was investigated.

Experiment
The pre-alloys with composition of Mn 65 Ga 20 Al 15−x Cu x (x = 0, 5, 10 and 15) were prepared from pure metals of Mn, Ga, Cu and Al in a water-cooled copper crucible of arc-melting furnace. Among the components of the alloy only Mn is evaporated. After each time of arc-melting in the same preparation conditions, Mn is lost of about 3 wt% and each sample was overturned and arc-melted five times to ensure homogeneity. Therefore, 15 wt% of Mn was added to the alloy for compensation of its evaporation. The obtained ingots were placed in a cylindrical quartz tube with an orifice at the bottom with a diameter of 0.5 mm to create alloy ribbons on a melt-spinning system using a single copper roller operating at a tangential velocity of 20 m s −1 . With these conditions, the obtained alloy ribbons have width of ∼ 3 mm and thickness of ∼ 20 μm. The ribbons were then annealed with various conditions to obtain the desired structure. The annealing temperature, T a , was changed from 250°C to 350°C, and the annealing time, t a , was varied from 5 to 20 h. All the arc-melting, melt-spinning and annealing processes were performed under Ar atmosphere to avoid oxygenation. x-ray diffraction (XRD), scanning electron microscopy (SEM) and energy-dispersive x-ray spectroscopy (EDX) methods were used to analyze crystalline structure, morphology and component content of the alloy ribbons at room temperature, respectively. Magnetic properties of the alloy ribbons were investigated on physical property measurement system (PPMS) and vibrating sample magnetometer (VSM).

Results and discussion
In order to check concentration of elements in the alloy ribbons after fabrication, the samples with x = 0 and x = 10 were selected for the EDX measurement (figure 1). We can see that, the concentration of elements has a deviation but not much from the nominal composition. This deviation probably is due to the nature of this analysis (semi-quantitative analysis). With adding Cu, the Al concentration decreases, meanwhile the Mn and Ga concentrations are almost unchanged.
The SEM images of the cross-section of the Mn 65 Ga 20 Al 15-x Cu x (x = 0, 5, 10 and 15) ribbons with the copper wheel contact-surface (wheel-side) at the bottom are shown in figure 2. We can see that, the shape, size and orientation of the crystalline grains strongly depend on the concentration of Cu. For the sample without addition of Cu (x = 0), the crystalline grains have columnar shape and their crystallization direction is tendly perpendicular to the ribbon surface. The size of the crystalline grains increases from the wheel-side to the freeside. The columnar-shaped grains have an average diameter and length in the range of about 2 μm − 4 μm and 10 μm − 20 μm, respectively. As partial replacement of Cu for Al, the directional crystallization of grains decreases with increasing the Cu concentration. For the sample with low Cu concentration of 5% (x = 5), both the uniformity and orientation of the grains are worse than those of the sample with x = 0. The grain size is nonhomogeneous for the sample with Cu concentration of 10% (x = 10). This sample contains both the small (1-3 μm) and large (5-8 μm) grains. On the other hand, the directional crystallization of the grains is not observed as in the samples with x = 0 and x = 5. With the highest Cu concentration of 15% (x = 15), it is difficult to observe the grain boundaries in the sample. The above analysis is consistent with the mechanical properties of the obtained alloy ribbons after fabrication. Without addition of Cu, the alloy ribbons are brittle and easily break. The ductility of the ribbons increases when the concentration of Cu increases. Figure 3 shows the XRD patterns of the as-quenched Mn 65 Ga 20 Al 15-x Cu x (x = 0, 5, 10 and 15) alloy ribbons. Through data of standard cards, diffraction peaks are identified for the crystalline structure types of    versus external magnetic field and do not have saturation state like conventional ferromagnetic materials. However, the samples still manifest coercivity in the range of 0.1 kOe to 0.6 kOe and magnetization at magnetic field of 50 kOe from 0.5 emu g −1 to 1.3 emu g −1 , depending on Cu concentration ( figure 4). This is in good agreement with the phase analysis from the XRD patterns in figure 3. The D0 19 -Mn 3 Ga antiferromagnetic phase is dominant, but the D0 22 -Mn 3 Ga ferrimagnetic phase hardly appears in all the as-quenched alloy ribbons. The magnetization at magnetic filed of 50 kOe decreases when the concentration of Cu increases from 0% to 15%. Figure 5(a) shows hysteresis loops at room temperature of the Mn 65 Ga 20 Al 15-x Cu x (x = 0, 5, 10 and 15) ribbons annealed at 275°C for 15 h. One can see that the annealing process strongly influences on the magnetic properties of the alloy ribbons. Both the coercivity and the saturation magnetization (for brevity, henceforth the magnetization at the magnetic field of 50 kOe is approximately called as the saturation magnetization M s ) are significantly enhanced. The variation trend of these two magnetic parameters on the Cu concentration after annealing is similar ( figure 5(b)). For the sample with Cu concentration of 10%, both the H c and M s reach the maximal value of 12.9 kOe and 6.8 emu g −1 , respectively. With Cu concentration lower or higher than 10%, both the H c and the M s of the samples decrease. The reduced virgin magnetization curves measured at room temperature of the Mn 65 Ga 20 Al 15-x Cu x (x = 0, 5, 10 and 15) ribbons annealed at 275°C for 15 h are presented in figure 5(c). The shape of these curves shows that the coercivity mechanism in the Mn-Ga based alloy includes both the reverse nucleation and domain wall pinning [46]. However, the reverse nucleation mechanism in the sample with x = 0 is dominant. Meanwhile, the domain wall pinning mechanism prevails in the sample with Cu concentration of 10%. Thus, the substitution of Cu for Al changed the coercivity mechanism of the alloy ribbons.
For further investigation of the influence of temperature and annealing time on structure and magnetic properties of the alloy, the composition of Mn 65 Ga 20 Al 5 Cu 10 (x = 10) is selected. It can be seen that, the coercivity and saturation magnetization significantly change after annealing at temperature of 250°C-325°C for 15 h (figure 6(a)) or annealing at 275°C for 5 h −20 h ( figure 7(a)). Figure 6(b) presents the dependence of H c and M s on various annealing temperatures as annealing time t a is kept for 15 h. When the annealing temperature T a is increased from 250°C to 275°C, the coercivity and saturation magnetization of the sample increase from 11 kOe to 12.9 kOe and from 1.6 emu g −1 to 6.8 emu g −1 , respectively. Continuing to increase the annealing temperature to 300°C, the coercivity decreases to 9 kOe, meanwhile the saturation magnetization increases and reaches its maximum value of 18.7 emu g −1 . When annealing at 325°C, H c and M s decrease to 6.8 kOe and 14.7 emu g −1 , respectively. The obtained results show that the formation of ferrimagnetic and antiferromagnetic phases in the samples depends not only on Cu concentration but also on the annealing temperature. The sample annealed at 650°C has lower coercivity but higher saturation magnetization in comparision with those of one annealed at 275°C. This probably is due to the formation of a ferromagnetic or ferrimagnetic phase in the sample with higher saturation magnetization but lower coercivity. On the other hand, if grain size is too small or large, then coercivity also decreases. To simultaneously increase these magnetic parameters, it is necessary to reduce the formation of magnetic phases with low H c and M s . Besides, the grain size of the hard magnetic phases should be controlled to be optimal for increasing the coercivity.
Continuing to keep annealing temperature at 275°C and change different annealing time for 5 h, 10 h, 15 h and 20 h, the dependence of the coercivity and saturation magnetization on annealing time is obtained ( figure 7(b)). We can realize that, the saturation magnetization rapidly increases from 2.7 emu g −1 to 6.8 emu g −1 when the annealing time increases from 5 h to 15 h. Meanwhile, the coercivity hardly changes and keeps the value of about 12.9 kOe. Further increasing the annealing time to 20 h, both the coercivity and saturation magnetization decrease. Thus, the annealing regime with temperature of 275°C − 300°C and time of 10 h − 15 h, is effective for enhancing hard magnetic parameters of the alloy ribbons. The obtained results are in good agreement with those reported by T. Saito et al for Mn 65 Ga 35-x Cu x alloy ribbons [40]. They showed that, the coercivity strongly depends on Cu concentration from 0 to 20% and annealing temperature from 200°C to 600°C. Meanwhile, these magnetic parameters almost un-change when the annealing time is prolonged from 2 h to 20 h. The annealing at 300°C for 10 h resulted in the formation of D0 22 -Mn 3 Ga and Cu 3 Mn 2 Ga phases. The presence of the D0 22 phase is responsible for the high coercivity (exceeding 20 kOe) of the annealed Mn 65 Ga 20 Cu 15 alloy ribbons. However, the obtained M s value is rather low, only 3 emu g −1 . In other studies [40,45], the substitution of Al for Ga in the Mn-Ga based alloy ribbons formed a phase with new cubic structure with space group P2 1 3 or P4 2  addition to the D0 22 ferrimagnetic phase. The coercivity in the Mn 65 Ga 15 Al 20 ribbon annealed at 700°C for 1 h is increased from 5 kOe to 9.1 kOe by contribution of the D0 22 phase. Thus, we can conclude that, each composition needs different optimal technological conditions.

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The reduced virgin magnetization curves measured at room temperature of the Mn 65 Ga 20 Al 5 Cu 10 ribbon with various annealing temperatures and time are shown in figures 6(c) and 7(c). It can be seen that, the coercivity mechanism of the alloy ribbons annealed at different temperatures is more distinct in comparision with those annealed for various time. This indicates that the annealing temperature influences on the magnetic properties of the sample more stronger than the annealing time. Normally, the formation of crystalline phases depends on annealing temperature, meanwhile annealing time needs for the growth and stability of the phases. When the annealing temperature and time increase, the reverse nucleation mechanism decreases and is replaced by the domain wall pinning one. Figure 8(a) shows the reduced thermomagnetization curves M(T) of Mn 65 Ga 20 Al 15-x Cu x alloy ribbons (x = 0, 5, 10 and 15) annealed at 275°C for 15 h with an applied magnetic field of 10 kOe. The shape of the reduced thermomagnetization curve as well as the phase transition temperature T C strongly depends on the Cu concentration. All the samples have a phase transition temperature, except that the sample with Cu concentration of 5% manifests two magnetic phase transitions. When the concentration of Cu is 10%, the Curie temperature is determined to be about 730 K, which is quite similar to the first phase transition temperature of the sample with 5% of Cu. This temperature is believed to be of the ferrimagnetic phase D0 22 -Mn 3 Ga and in good agreement with previous studies [31,38,39]. However, this value is still lower than that calculated by J. Kubler [29]. As presented above, the coercivity of this sample is also significantly enhanced, reaching 12.9 kOe. The second magnetic phase transition of the sample with 5% Cu is above 900 K, similar to that of the ribbons with Cu concentrations of 0 % and 15%. The obtained high coercivity for the samples with Cu concentrations of 5 % and 10% can be due to the existence of a hard magnetic phase with T C of about 730 K. Figure 8(a) indicates the temperature dependence of reduced magnetization in magnetic field of 10 kOe of the Mn 65 Ga 20 Al 5 Cu 10 ribbon annealed at various temperature for 15 h. At the annealing temperature of 250°C, T C of the sample reaches about 920 K. When the annealing temperature varies from 275°C to 325°C, T C changes slightly around temperature of about 730 K. This means that the D0 22 -Mn 3 Ga phase is formed in the sample annealed in this annealing temperature range. The sample behaves single magnetic phase with all the annealing temperatures. According to the obtained result in figure 6, the magnetic phase with higher T C has low saturation magnetization.
The XRD patterns of the Mn 65 Ga 20 Al 15-x Cu x (x = 0, 5, 10 and 15) ribbons after annealing at 275°C for 15 h were showed in figure 9. A clear change of structure in the samples after the annealing process is observed. Two crystalline phases of Mn 3 Ga with hexagonal (D0 19 ) and tetragonal (D0 22 ) structures are formed ( figure 9(a)). However, the D0 19 antiferromagnetic phase still dominates in the alloy ribbons with Cu concentrations of 0% and 5%. This indicated that the fraction of the D0 22 ferrimagnetic phase in these ribbons is low. Intensity of the diffraction peaks characterizing for the D0 19 phase decreases when the concentration of Cu increases to 10%. This is demonstrated by the fraction of the phases by Rietveld analysis for the typical samples with x = 5 and x = 10. For the ribbon with Cu concentrations of 5%, the fraction of phases D0 22 and D0 19 is 25% and 75%, respectively. When the Cu concentration increases to 10%, the fraction of D0 22 increases (68%) with a decrease of D0 19 (31%) and a small amount of AlCu 3 appeared (1%). It is difficult to observe the appearance of D0 19 phase when the content of Cu is 15%. Instead of it, the phases of D0 22 and AlCu 3 are formed with this concentration of Cu ( figure 9(b)). Combined with the results obtained on the M(T) curve ( figure 8(a)), the sample with x = 10 has only a phase transition temperature corresponding to the ferrimagnetic phase D0 22 -Mn 3 Ga, we realize that the fraction of D0 22 phase of the sample with Cu concentration of 10% Cu is highest. Thus, H c of this annealed ribbon strongly increases compared to that of the as-quenched one. The magnetic properties and fraction of the phases of the Mn 65 Ga 20 Al 15-x Cu x (x = 0, 5, 10 and 15) ribbons annealed at different annealing regimes are determined and listed in table 1. Figure 10 shows the SEM images of the cross-section of the Mn 65 Ga 20 Al 15-x Cu x (x = 0, 5, 10 and 15) ribbons annealed at 275°C for 15 h. We can see that the change of microstructure of the samples is depended on concentrations of Cu. For the sample with x = 0, structural change is unclear. The separation of small grains from large ones is observed in alloy ribbons with Cu concentration of 5%. These small grains appear more at the wheel-side than at the free-side and their size tends to increase towards the free-side of the ribbons. For the sample with x = 10, the separation of large grains into small ones is quite homogeneous throughout the ribbon and their size is relatively uniform. As for the ribbon with the largest Cu concentration (15%), the separation into small grains only occurs around the large ones. The separation of large grains into small ones can be due to the structural change of the crystalline phases during the annealing process. The formation of new crystalline phases change the magnetic properties of the alloy. For example, the strong enhancement of the coercivity of the

Conclusion
The structure and magnetic properties of the Mn 65 Ga 20 Al 15-x Cu x (x = 0, 5, 10 and 15) alloy ribbons were investigated. The appropriate Cu concentration and annealing process create the hard magnetic crystalline phases and optimize microstructure for coercivity of the alloy ribbons. The D0 22 -Mn 3 Ga crystalline phase plays an important role for the hard magnetic properties of the alloys. The hard magnetic phases are favourably formed when Cu concentration is in range from 5% to 10% combined with an annealing regime at 275°C − 300°C for 10 h − 15 h. The highest coercivity of 12.9 kOe and saturation magnetization of 18.7 emu g −1 are respectively obtained with different annealing temperatues of 275°C and 300°C but with the same Cu concentration of 10% and annealing time of 15 h. Controlling technological conditions to simultaneously improve magnetic parameters for application possibility of Mn-Ga based rare earth-free hard magnetic materials is significant.