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Article

Glass-Forming Ability and Magnetic Properties of Al82Fe16Ce2 and Al82Fe14Mn2Ce2 Alloys Prepared by Mechanical Alloying

by
Nguyen Hong Hai
,
Nguyen Hoang Viet
and
Nguyen Thi Hoang Oanh
*
School of Materials Science and Engineering, Hanoi University of Science and Technology, Hanoi 100000, Vietnam
*
Author to whom correspondence should be addressed.
Appl. Sci. 2024, 14(1), 152; https://doi.org/10.3390/app14010152
Submission received: 21 November 2023 / Revised: 15 December 2023 / Accepted: 21 December 2023 / Published: 23 December 2023
(This article belongs to the Section Materials Science and Engineering)
Browse Figures
Versions Notes

Abstract

:
Al82Fe16Ce2 and Al82Fe14Mn2Ce2 amorphous alloys were successfully synthesized by the mechanical alloying technique. The microstructural evolution of the milled powders was thoroughly investigated employing X-ray diffraction (XRD) and scanning electron microscopy (SEM). Additionally, their magnetic properties were quantitatively evaluated by a vibrating sample magnetometer (VSM). A full amorphous structure was obtained for both alloys after milling for 40 h. During the initial milling stage, extending from 5 to 20 h, an fcc solid solution phase was formed, coexisting with the residual Al phase. The partial substitution of 2 atomic percent (at.%) Mn for Fe in Al82Fe16Ce2 did not affect the alloy’s glass-forming ability. The amorphous Al82Fe16Ce2 and Al82Fe14Mn2Ce2 powders exhibited a nearly spherical shape, with diameters ranging from 1 to 3 µm and to 10 µm, respectively. Additionally, both the Al82Fe16Ce2 and Al82Fe14Mn2Ce2 alloys demonstrated characteristics of hard magnetism.

1. Introduction

Bulk metallic glass materials have attracted significant scientific and technological interests as both structural and functional materials [1,2,3,4,5,6]. This interest is largely due to their outstanding mechanical properties and excellent corrosion resistance, which are attributed to the absence of long-range order and defects such as point defects, dislocations, and stacking faults [2,7]. The tensile strength of Al-based amorphous alloys is often 2–3 times higher than that of conventional Al crystalline alloys [2,4,6,7,8,9]. Al-based amorphous alloys with high Al contents (>80 at.%) are promising candidates for structural applications because of their high specific strength at room temperature, low density, and excellent corrosion resistance [1,10,11]. There are two kinds of Al-based amorphous alloys such as aluminum–early transition metal–late transition metal (Al-ETM-LTM) and aluminum–late transition metal–rare earth (Al-LTM-RE), where LTM (Fe, Co, Ni, Cu), ETM (Ti, V, Cr, Mn), and RE (Y, Ce, La) are early transition metals, late transition metals, and rare earth elements, respectively [10]. The Al-LTM-RE alloys have a higher glass-forming ability (GFA) in comparison with Al-ETM-LTM alloys [12,13]. Additionally, most of the amorphous aluminum–transition metal–rare earth (Al-RE-TM) alloys are paramagnetic. However, certain Al-based glasses exhibit diverse magnetic behaviors, including diamagnetism, frustration, spin-glass behavior, and superparamagnetism. These unique properties make these alloys suitable for industrial applications, such as protective coatings, and for use in the electronics industry [14]. Mechanical alloying (MA) and rapid quenching (RQ) techniques are widely recognized as the predominant methods for preparing amorphous alloys, as evidenced by various studies [2,3,4,9,15,16,17,18]. For instance, J. Zhang et al. successfully produced a completely amorphous structure of Al86Ce10Fe4 using the melt-spinning RQ technique. This process involved rapid quenching at speeds around 40 m/s in a high-purity helium environment, resulting in ribbons with a thickness of 40 μm and a width of 4 mm [19]. However, achieving amorphization typically requires high cooling rates, ranging from 105 to 106 K/s, to prevent the crystalline-phase formation. Amorphous materials are generally categorized into two groups: conventional forms like thin ribbons [9,20,21] and bulk amorphous materials, which are more than 100 µm thick and can vary in shape, including rods, tubes, plates, etc. [22]. In comparison, the powder metallurgy technique, especially MA, provides certain advantages over RQ methods. MA can be conducted at room temperature and is particularly effective for producing amorphous alloys in powder form. These powders are conducive to being easily consolidated into bulk samples. Additionally, high-energy ball-milling, a key aspect of MA, has been employed for creating a range of materials, including nanocomposite materials, amorphous alloys, quasicrystalline alloys, and high-entropy alloys. This is achieved by milling elemental or compound powders in an inert atmosphere, as demonstrated in various studies [18,23,24,25,26]. In the milling process, under the impact of colliding balls, a mixture of powders undergoes plastic deformation, cold-welding, and fracture, repeatedly leading to a reduction in the particle size and obtaining a desirable structural phase [15,16,27,28,29,30]. Powdered aluminum-based amorphous alloys are utilized in a variety of applications that demand high strength, durability, wear resistance, and light weight. These alloys are integral in the production of diverse industrial and commercial products, including robot parts, tooling, molds, fishing reels, bicycle gears, and wheelchair components [31,32]. Among the Al-based amorphous alloys, Al-Fe alloys attract technological interests due to their high specific strength and excellent corrosion resistance at elevated temperatures under oxidizing, carburizing, and sulfurizing atmospheres media [33]. There have been several reports regarding synthesized Al-Fe amorphous alloys prepared by the MA technique [34,35,36,37,38,39,40,41,42,43,44]. Viet et al. reported that a full amorphous structure can be obtained for Al84Fe16 alloys prepared by MA after milling for 100 h [39]. A small amount of rare earth elements (Y) could enhance the GFA and thermal stability of Al-Fe amorphous alloys. By a partial substitution of Y of 2 at.% for Al in Al84Fe16, Viet et al. showed that a complete amorphization can be obtained for Al82Fe16Y2 after MA for 100 h [39]. It has been found that a minor substitution of Y (large radius of 182 pm and large negative heat of mixing among constituent elements) improved the thermal stability of Al84Fe16 alloy. The onset temperature of Al82Fe16Y2 is 655 K, which is higher than that of Al84Fe16 of about 626 K.
In this study, we attempted to improve the GFA of mechanically alloyed Al-Fe-based amorphous alloys by incorporating the rare earth element cerium (Ce) and the transition metal manganese (Mn). We investigated the phase transformation, GFA, and magnetic properties of Al82Fe16Ce2 and Al82Fe14Mn2Ce2 alloys. This involved comparing factors influencing the amorphization process in Al-Fe-TM(RE) systems to clarify the GFA of these amorphous alloys prepared using the mechanical alloying method.

2. Materials and Methods

Elemental Al, Fe, Ce, and Mn powders (purity ≥ 99.9%) were mixed to the nominal compositions of Al82Fe16Ce2 and Al82Fe14Mn2Ce2 in atomic percent. The metallic powders were then encapsulated in hardened steel vials. For the mechanical alloying process, a tungsten carbide (WC) ball with a powder-to-ball weight ratio of 10 g to 200 g was utilized. This ratio was selected to optimize the milling efficiency and ensure uniformity in the powder mixture. To prevent the sticking of the powders to the milling tools, 50 mL of n-hexane was employed as a process control agent. Prior to milling, the vials were evacuated and then filled with protective argon (Ar) gas at a pressure of 3 × 105 Pa to create an inert atmosphere, thus preventing oxidation and contamination during the milling process. The mechanical alloying was conducted using a high-energy planetary ball mill (AGO-II type), which operated at a rotational speed of 350 rpm. The milling vials were equipped with a water-cooling system to effectively manage any temperature increases that could potentially affect the material properties. The duration of milling was varied, ranging from 5 to 40 h, to investigate the evolution of the alloys’ microstructure over time. To maintain purity and minimize gaseous contamination, each milling run was executed continuously for its predetermined duration without any interruptions. The principle and function of a high-energy planetary ball mill in the preparation of amorphous alloy powders from elemental powders is illustrated in Figure 1. Post milling, the morphology of the milled powders was extensively characterized using scanning electron microscopy (SEM) employing a HITACHI TM4000 PLUS instrument (Hitachi High-Tech Corporation, Tokyo, Japan). Phase analysis of the as-milled powders was conducted using X-ray diffraction (XRD) with a SIEMENS D5000 diffractometer (Siemens, Berlin, Germany). This analysis utilized Cu Kα radiation (λ = 1.5405 Å), covering a 2θ range from 20 to 80 degrees, with a step size of 0.02° and a scanning speed of 1° per minute. The XRD data were analyzed and normalized using the MDI Jade 6.5 software, which accesses the ICDD’s PDF-2 database, allowing for accurate phase identification. Lastly, the magnetic properties of the as-milled powders were assessed using a vibrating sample magnetometer (VSM, MicroSense EV9, Lowell, MA, USA) capable of applying a maximum field strength of 1200 kA/m. This measurement provided crucial insights into the magnetic behavior of the alloys, including their saturation magnetization and coercivity, which are essential for understanding their potential applications in various technological fields.

3. Results and Discussion

Figure 2 shows the XRD patterns of Al82Fe16Ce2 alloy powders milled for 5, 10, 20, and 40 h, respectively. After 5 h of milling, an fcc solid solution is observed alongside the remaining Al phase. With further milling to 10 h, the intensity of the diffraction peaks decreases and broadens. A broad halo peak can be recognized in the 2Θ range between 40° and 50°, indicating the formation of an amorphous phase after milling for 20 h. The Al is consumed, and only broad diffraction peaks of an fcc solid solution are detected in powders milled after 30 h. No distinct in XRD peaks of any crystalline phases are observed in the powder milled for 40 h, suggesting that the sample has a fully amorphous structure.
In this work, the amorphization of Al82Fe16Ce2 is completed after 40 h of milling, which is as fast as the Al82Fe16Ti2 alloy (achieving a full amorphous structure after 40 h of MA), as mentioned by Oanh et al. in [43], and faster than the Al82Fe16Ni2 (full amorphous structure after 50 h of MA). However, the Al82Fe16Cu2 alloy does not attain complete amorphization even after 60 h of milling. Based on the XRD results for these alloys, the GFA is ranked as follows: Al82Fe16Cu2 < Al82Fe16Ni2 < Al82Fe16Ti2, Al82Fe16Ce2. The three criteria proposed by Inoue [11], which relate to the common features of metallic glasses (MGs) with a strong GFA, are as follows: (i) the alloy system consists of at least three elements, (ii) there is a significant negative mixing enthalpy among the elements, and (iii) the atomic size ratios among the main constituent elements are greater than 12%. All four alloys include three elements, thereby meeting the first criterion. For the second criterion, all element pairs exhibit a negative mixing enthalpy (Al-Ti, Al-Ni, Al-Cu, Al-Ce, Fe-Al, Fe-Ti, Fe-Ni) except for the Fe-Cu (ΔHmix = +13 kJ/mole) and Fe-Ce pairs (ΔHmix = +3 kJ/mole), as shown in Table 1. Among all element pairs, Al-Ce has the highest negative enthalpy of mixing −38 kJ/mole. However, the Al82Fe16Ti2 alloy shows a negative mixing enthalpy whereas the Al82Fe16Ce2 alloy Fe-Ce pair has a positive enthalpy of mixing. Considering the last criterion of Inoue, the size ratios among the main constituent elements, the atomic size differences of Ti, Ni, Cu, and Ce with respect to Al are 2.7%, 12.5%, 10%, and 21.8%, respectively. It can be observed that the atomic size differences of the transition metals (Ti, Ni, Cu) with solvent atoms (Al) are within the limiting boundary (15%) of the Hume-Rothery rule [45] except Ce with Al (21.8%). The transition metals (Ti, Ni, Cu) have a similar radius to Al, which leads to the formation of a substitutional solid solution with Al, whereas the presence of Ce leads to the formation of an fcc solid solution after milling Al82Fe16Ce2 for 5 h (Figure 2). There are two main factors contributing to the GFA of Al82Fe16Ce2 and Al82Fe16Ti2 alloys, which are the atomic size difference and the negative enthalpy of mixing. Additionally, the GFA of Al82Fe16Ni2 and Al82Fe16Cu2 alloys is lower than that of Al82Fe16Ti2, which may be attributed to their less negative enthalpy of mixing or to a mismatch compared to Al82Fe16Ce2. As Zhang et al. [46] reported, the GFA of the Al-Ni-RE alloy prepared by the rapid quenching technique exhibited a strong dependence on the size of the RE atom. As the atomic radius of the RE element increases, the GFA of the amorphous alloys also increases.
Furthermore, we consider the amorphization of Al82Fe16Ce2 and Al82Fe16Y2 [39] alloys, which contain rare elements and were prepared by mechanical alloying using the same planetary ball mill. Interestingly, both alloys satisfy all three of Inoue’s criteria. Each alloy consists of three elements, with the mixing enthalpies for element pairs Al-Y and Al-Ce being −38 kJ/mole, and for Fe-Ce and Fe-Y being +3 and −1 kJ/mole, respectively. Therefore, both alloys exhibit nearly similar mixing enthalpies between atomic pairs. Al exhibits atomic size mismatches of 21.8% with Ce and 21.4% with Y, while Fe displays mismatches of 32.2% and 31.8% with Ce and Y, respectively. However, the Al82Fe16Ce2 and Al82Fe16Y2 alloys achieved a full amorphous structure after 40 and 100 h of milling, respectively. Both alloys have similar physical properties, including atomic radius (Y: 180 pm, Ce: 181 pm), boiling point (Y: 3345 °C, Ce: 3426 °C), and Pauling electronegativity (Y: 1.22, Ce: 1.12). The electrical structural perspectives offer a deeper comprehension of the stability of metallic glasses. There is a significant difference in the electron configurations of Ce and Y. Cerium, with its 58 electrons, exhibits a more complicated electron configuration compared to yttrium, which contains just 39 electrons. This increased complexity in Ce’s electronic structure could enhance chemical disorder during the milling process, potentially contributing to a higher glass-forming ability (GFA) [48]. In addition, raw Ce powders have a particle size of 248 μm (−60 mesh), which is smaller than that of Y powders of 372.5 μm. The amorphization process is initiated by the MA, which deforms the powder particles through a series of cold welding and fragmentation cycles, ultimately resulting in a fully amorphous structure. The rapid size reduction process leads to a faster amorphous phase completion. In this context, Y has a bigger initial particle size than that of Ce, so it takes a longer time for the milling to produce amorphous alloys. Additionally, the inherent metallic ductility of Y tends to favor the cold-welding process, dominating over fragmentation during milling. This effect is clearly demonstrated in the detailed analysis presented in our previous article [39]. The reduction of particle size and cold-welding lead to the extension of the amorphization of Al-Fe-Y. On the contrary, because of the presence of Ce in the Al-Fe-Ce alloy, the fragmentation process dominates, so the particle size of the mixture powders decreases rapidly, causing the amorphization process to take place after only 40 h (Figure 3). As seen in SEM micrographs of the MAed Al82Fe16Ce2 powders (Figure 3) milled at different times, most particles are flattened due to the collision between powders, balls, and jar. After milling for 5 h, the powder particles assume a layered structure, and the particle size of the milled powders is about 10–15 μm. The powder particles of Al82Fe16Ce2 decrease gradually with increasing milling time from 10 to 40 h. This means the fracture process is dominant during the process of the mechanical alloying of the Al82Fe16Ce2 alloy powders. The particle size decreased to 1–3 µm, and some particle agglomerations can be seen in Figure 3g,h). The SEM images of the amorphous Al84Fe16 (without Ce) and Al82Fe16Y2 (substitute 2 at.% of Y for Al) alloys’ MA for 100 h [39] show that the particle powders of the amorphous alloy are about 10 μm, which is much higher than that of the Al82Fe16Ce2 alloy powders. This means that the addition or substitution of Ce for Al in the Al84Fe16 alloys or Y in Al82Fe16Y2, respectively, results in decreasing powder particles. In addition, Gan Luo et al. reported that Ce addition to aluminum alloys contributed to grain refinement [49]. This is achieved by lowering the recalescence temperature and the growth temperature of the Al-Fe eutectic structure, improving the morphology and distribution of the Fe-containing phase. Such enhancements also aid in the consolidation of amorphous powder, facilitating the formation of bulk materials.
In the Al82Fe14Mn2Ce2 alloy, amorphization occurs similarly to that observed in Al82Fe16Ce2 alloy powders. An fcc solid solution phase, along with an Al phase, appears after milling for 5 h. A halo peak, indicating the presence of an amorphous phase, is observed after 20 h of milling (see Figure 4). The amorphization process of the Al82Fe14Mn2Ce2 alloy powders is completed after 40 h.
The amorphization of Al82Fe14Mn2Ce2 alloy by MA is as fast as Al82Fe16Ce2 alloy. A fully amorphous structure in the Al82Fe14Mn2Ce2 alloy is achieved after 40 h of milling. Substituting 2 atomic percent (at.%) of Mn for Fe in the Al82Fe16Ce2 alloy does not significantly affect the amorphization process. Similar observations regarding the GFA of melt-spun Al-La-Ni-Fe alloys were reported by Bassu et al. [50]. As indicated in Table 1, the mixing enthalpy of the Mn-Fe pair is 0 kJ/mole, with Mn exhibiting atomic size mismatches of 5.5%, 13%, and 26% with Al, Fe, and Ce, respectively. Fe contributes slightly higher atomic radii mismatches compared to Mn. Meanwhile, the Al-Mn pair shows a greater negative mixing enthalpy (−19 kJ/mole) than the Al-Fe pair (−11 kJ/mole). Consequently, the Al82Fe16Ce2 alloy exhibits a GFA similar to that of the Al82Fe14Mn2Ce2 alloy. The substitution of Mn for Fe in Al-Fe-Ce alloy results in an unenhanced amorphization process. The addition of different chemical elements in these multicomponent alloys results in a dense ‘random-packed’ structure, which significantly restricts atomic mobility, a crucial factor for crystallization. This, in turn, may enhance the thermal stability of the alloys [51,52].
The particle sizes of Al82Fe14Mn2Ce2 powders are slightly larger than those of Al82Fe16Ce2 after milling for 5, 10, and 20 h (Figure 5). This can be explained by the enhanced softness of the powder material with Mn addition, making cold-welding processes during ball-collisions more favorable. The particle shapes of all three alloys appear rather spherical than flake-like, indicating that impact events during ball collisions dominated the sliding events. After milling for 40 h, the particle size of Al82Fe14Mn2Ce2 powders was in the range of 2 ÷ 10 µm due to agglomeration phenomena (Figure 5).
Figure 6a,b presents the magnetization curves (M-H) for Al82Fe16Ce2 and Al82Fe14Mn2Ce2 alloys after milling. All the synthesized alloys show similar hysteresis loops, indicating their ferromagnetic nature [53,54]. Table 2 summarizes the properties obtained from the M–H curves (saturation magnetization–Ms and coercive force–Hc). The maximum saturation Ms decreases from 29.5 and 24.69 emu/g to 1.53 and 1.41 emu/g for Al82Fe16Ce2 and Al82Fe12Mn2Ce2 alloys milled from 5 h to 40 h, respectively. The saturation magnetization values of the two amorphous alloys are similar to that of Al82Fe14Ti2Y2 amorphous alloys (Ms = 1.5 emu/g) and slightly greater than that of Al82Fe14Ni2Y2 amorphous alloys (Ms = 0.8 emu/g). The saturation magnetization depends on the phases present (crystalline structure) in the alloys and their chemical composition. By comparing Figure 2, Figure 3, Figure 4, Figure 5 and Figure 6, it is realized that by increasing the milling time from 5 to 40 h for the Al82Fe16Ce2 and Al82Fe14Mn2Ce2 alloys, the Ms decreases due to the content of the ferromagnetic phases–bcc-Fe decreasing, the content of paramagnetic phase–fcc solid solution increasing, and, at the final milling state, a transformation of solid solution to amorphous phase occurring [43]. The tendencies of the coercive force Hc of both alloys are similar. In the early milling state, there are large numbers of defects and microstrain introduced to the mixture of powder due to the strong collision between powders and milling tools. Therefore, a high Hc value of 366.45 and 398.99 Oe is obtained for the Al82Fe16Ce2 and Al82Fe14Mn2Ce2 alloys after 5 h of milling, respectively. For longer milling times of 10 h, line broadening of the peaks in Figure 2 and Figure 4 indicates the reduction in the crystallite size and lattice strain introduced by milling, so Hc decreases very quickly to 118 and 81.02 Oe for the Al82Fe16Ce2 and Al82Fe14Mn2Ce2 alloys, respectively. The Hc increases again with increasing milling time to 20 h when amorphous phases start forming in the alloys. In their amorphized state, the Al82Fe16Ce2 and Al82Fe14Mn2Ce2 alloys show coercivity (Hc) values of 156.69 Oe and 158.45 Oe, respectively. These values are lower than those exhibited by the Al82Fe14Ni2Y2 and Al82Fe14Ti2Y2 alloys, which are 186.5 Oe and 198.9 Oe. According to the standard classification of ferromagnetic materials, those with an Hc value less than 1000 A/m (approximately 12.57 Oe) are considered to be ‘soft magnetic’. Materials within the 1000–10,000 A/m range (approximately 12.57–125.7 Oe) are classified as ‘semi-hard’, and those exceeding 10,000 A/m (more than 125.7 Oe) are categorized as ‘hard’. Based on this classification, all these alloys are categorized as hard magnetic materials [53,54,55,56,57].

4. Conclusions

This study aimed to investigate the formation of amorphous structures in Al82Fe16Ce2 and Al82Fe14Mn2Ce2 alloy powders via mechanical alloying. The primary goal was to determine if a fully amorphous phase could be achieved and to understand the underlying mechanisms. Our results confirm that a fully amorphous structure was successfully obtained in both alloys after 40 h of milling. This demonstrates the effectiveness of mechanical alloying in producing amorphous structures in these specific alloys.
Initially, after 5 h of milling, both alloys exhibited an fcc solid solution formation alongside residual Al phases. The mixing enthalpies of Al-Ce (−38 kJ/mol) and Al-Fe (−11 kJ/mol), and the presence of Fe-Ce pairs with a positive mixing enthalpy (3 kJ/mol), did not hinder the amorphization process. This finding suggests that the appearance of Ce is a crucial factor in the amorphization mechanism, in conjunction with the three empirical rules of Inoue for bulk amorphous alloy fabrication. Additionally, factors like electron configuration and particle size of the raw powders were also found to influence the amorphization of Al-Fe-TM and Al-Fe-RE amorphous alloys.
Furthermore, our study observed significant changes in magnetic properties during the milling process. The saturation magnetization (Ms) values decreased dramatically from 29.5 and 24.69 emu/g to 1.53 and 1.41 emu/g for Al82Fe16Ce2 and Al82Fe12Mn2Ce2 alloys, respectively. The Ms values decreased according to the transformation of the ferromagnetic phase content: starting with the formation of fcc solid solution and subsequently transforming into an amorphous phase. The coercivity (Hc) of both alloys also displayed intriguing variations throughout the milling process. At first, Hc decreased as milling time increased from 5 to 10 h, attributable to the reduction in the crystallite size of the alloys. However, upon extending the milling time to 20 h, Hc values began to rise again due to the deformation of the powders. After 40 h of milling, the Al82Fe16Ce2 and Al82Fe14Mn2Ce2 alloys underwent a transformation from a solid solution to an amorphous state. This transformation resulted in their coercivity (Hc) values reaching their lowest values, which were 156.69 Oe and 158.45 Oe for the Al82Fe16Ce2 and Al82Fe14Mn2Ce2 amorphous alloys, respectively.
The amorphous powder of the Al82Fe16Ce2 alloy exhibited a uniform particle size range of 1 ÷ 3 µm with a nearly spherical particle shape, suggesting potential for consolidation into bulk samples.

Author Contributions

Conceptualization, N.H.H. and N.T.H.O.; methodology, N.H.H., N.H.V. and N.T.H.O.; software, N.H.V.; validation, N.H.H., N.H.V. and N.T.H.O.; formal analysis, N.H.V.; investigation, N.H.H. and N.H.V.; resources, N.H.H. and N.T.H.O.; data curation, N.H.V. and N.H.H.; writing—original draft preparation, N.H.H. and N.T.H.O.; writing—review and editing, N.H.V. and N.T.H.O.; visualization, N.H.V.; funding acquisition, N.H.H. All authors have read and agreed to the published version of the manuscript.

Funding

This work was financially supported by Project B2022-BKA-16, Ministry of Education and Training of Vietnam, which the authors sincerely appreciate.

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

All data generated or analyzed during this study are included in this published article.

Conflicts of Interest

The authors declare no conflict of interest.

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Applsci 14 00152 g001
Figure 1. Schematic diagram of mechanical alloying: (a) metallic powders, (b) view of ball milling vial, (c) alloying process during ball milling, and (d) final state of the amorphous alloy.
Figure 1. Schematic diagram of mechanical alloying: (a) metallic powders, (b) view of ball milling vial, (c) alloying process during ball milling, and (d) final state of the amorphous alloy.
Applsci 14 00152 g001
Applsci 14 00152 g002
Figure 2. XRD patterns of Al82Fe16Ce2 powders milled for different milling times.
Figure 2. XRD patterns of Al82Fe16Ce2 powders milled for different milling times.
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Applsci 14 00152 g003
Figure 3. SEM images of Al82Fe16Ce2 powders after milling for (a,b) 5 h, (c,d) 10 h, (e,f) 20 h, and (g,h) 40 h.
Figure 3. SEM images of Al82Fe16Ce2 powders after milling for (a,b) 5 h, (c,d) 10 h, (e,f) 20 h, and (g,h) 40 h.
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Figure 4. XRD patterns of Al82Fe14Mn2Ce2 powders milled for different milling times.
Figure 4. XRD patterns of Al82Fe14Mn2Ce2 powders milled for different milling times.
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Figure 5. SEM images of Al82Fe14Mn2Ce2 powders after milling for (a,b) 5 h, (c,d) 10 h, (e,f) 20 h, and (g,h) 40 h.
Figure 5. SEM images of Al82Fe14Mn2Ce2 powders after milling for (a,b) 5 h, (c,d) 10 h, (e,f) 20 h, and (g,h) 40 h.
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Figure 6. M–H curves of (a) Al82Fe16Ce2 and (b) Al82Fe14Mn2Ce2 milled for different times.
Figure 6. M–H curves of (a) Al82Fe16Ce2 and (b) Al82Fe14Mn2Ce2 milled for different times.
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Table 1. Atomic radii mismatch (in %) and enthalpies of mixing (in kJ/mole) for Al, Fe, Ni, Ti, and Y binary systems [47].
Table 1. Atomic radii mismatch (in %) and enthalpies of mixing (in kJ/mole) for Al, Fe, Ni, Ti, and Y binary systems [47].
AlFeTiNiCuMnYCe
Al-13 [%]2.7 [%]12.5 [%]10 [%]5.5 [%]21.4 [%]21.8 [%]
Fe−11 [kJ/mole]-15.6 [%]0.8 [%]3.1 [%]8.1 [%]31.8 [%]32.2 [%]
Ti−30 [kJ/mole]−17 [kJ/mole]-14.9 [%]12.9 [%]8.1 [%]19.2 [%]19.6 [%]
Ni−22 [kJ/mole]−2 [kJ/mole]−35 [kJ/mole]-2.3 [%]7.4 [%]31.3 [%]31.6 [%]
Cu−1
[kJ/mole]		+13
[kJ/mole]		−8
[kJ/mole]		−8
[kJ/mole]		-+5.1 [%]29.6 [%]30 [%]
Mn−19 [kJ/mole]0
[kJ/mole]		−8 [kJ/mole]−8 [kJ/mole]+4
[kJ/mole]		-25.8 [%]26 [%]
Y−38 [kJ/mole]−1 [kJ/mole]+15
[kJ/mole]		−31 [kJ/mole]−22
[kJ/mole]		−1 [kJ/mole]-0.5 [%]
Ce−38 [kJ/mole]+3 [kJ/mole]+18 [kJ/mole]−28 [kJ/mole]−21
[kJ/mole]		+1 [kJ/mole]0
[kJ/mole]		-
Table 2. Ms and Hc values obtained from VSM analyses for Al82Fe16Ce2, Al82Fe14Mn2Ce2 alloys (in this work) and Al82Fe14Ni2Y2, Al82Fe14Ti2Y2 alloys in [36] milled for different times.
Table 2. Ms and Hc values obtained from VSM analyses for Al82Fe16Ce2, Al82Fe14Mn2Ce2 alloys (in this work) and Al82Fe14Ni2Y2, Al82Fe14Ti2Y2 alloys in [36] milled for different times.
Al82Fe16Ce25 h10 h20 h40 h
Hc, Oe366.45118.49207.20156.69
Ms, emu/g29.515.703.841.53
Al82Fe14Mn2Ce25 h10 h20 h40 h
Hc, Oe398.9981.02218.36158.45
Ms, emu/g24.6910.174.621.41
Al82Fe14Ni2Y25 h10 h20 h40 h60 h
Hc, Oe-340.8-282.1186.5
Ms, emu/g-13.0-2.70.8
Al82Fe14Ti2Y25 h10 h20 h40 h60 h100 h
Hc, Oe-256.4-299.9281.0198.9
Ms, emu/g-14-2.81.21.5
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Hai, N.H.; Viet, N.H.; Oanh, N.T.H. Glass-Forming Ability and Magnetic Properties of Al82Fe16Ce2 and Al82Fe14Mn2Ce2 Alloys Prepared by Mechanical Alloying. Appl. Sci. 2024, 14, 152. https://doi.org/10.3390/app14010152

AMA Style

Hai NH, Viet NH, Oanh NTH. Glass-Forming Ability and Magnetic Properties of Al82Fe16Ce2 and Al82Fe14Mn2Ce2 Alloys Prepared by Mechanical Alloying. Applied Sciences. 2024; 14(1):152. https://doi.org/10.3390/app14010152

Chicago/Turabian Style

Hai, Nguyen Hong, Nguyen Hoang Viet, and Nguyen Thi Hoang Oanh. 2024. "Glass-Forming Ability and Magnetic Properties of Al82Fe16Ce2 and Al82Fe14Mn2Ce2 Alloys Prepared by Mechanical Alloying" Applied Sciences 14, no. 1: 152. https://doi.org/10.3390/app14010152

APA Style

Hai, N. H., Viet, N. H., & Oanh, N. T. H. (2024). Glass-Forming Ability and Magnetic Properties of Al82Fe16Ce2 and Al82Fe14Mn2Ce2 Alloys Prepared by Mechanical Alloying. Applied Sciences, 14(1), 152. https://doi.org/10.3390/app14010152

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