Microstructural and magnetic characterization of hydrogen-decrepitated CeFeB melt-spun ribbons fabricated at various wheel surface speeds

Microstructural and magnetic characterization of hydrogen-decrepitated CeFeB melt-spun ribbons fabricated at various wheel surface speeds | Research Square window.SnipcartSettings = { analytics: { enabled: false } }; (function() { var accessVector = localStorage.getItem('access_vector') || ''; window.dataLayer = window.dataLayer || []; if (accessVector) { window.dataLayer.push({ user: { profile: { profileInfo: { snid: accessVector } } } }); } })(); (function(w,d,s,l,i){w[l]=w[l]||[];w[l].push({'gtm.start':new Date().getTime(),event:'gtm.js'});var f=d.getElementsByTagName(s)[0],j=d.createElement(s),dl=l!='dataLayer'?'&l='+l:'';j.async=true;j.src='https://www.googletagmanager.com/gtm.js?id='+i+dl;f.parentNode.insertBefore(j,f);})(window,document,'script','dataLayer','GTM-K279D39R'); Browse Preprints In Review Journals COVID-19 Preprints AJE Video Bytes Research Tools Research Promotion AJE Professional Editing AJE Rubriq About Preprint Platform In Review Editorial Policies Our Team Advisory Board Help Center Sign In Submit a Preprint Cite Share Download PDF Research Article Microstructural and magnetic characterization of hydrogen-decrepitated CeFeB melt-spun ribbons fabricated at various wheel surface speeds Yasin Yılmaz, Muhammed Fatih Kilicaslan This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-8354310/v1 This work is licensed under a CC BY 4.0 License Status: Published Journal Publication published 24 Mar, 2026 Read the published version in Journal of Materials Science: Materials in Electronics → Version 1 posted You are reading this latest preprint version Abstract The demand for heavy rare earth-free permanent magnetic alloys has been increasing due to the limited reserves and high costs of critical rare earths such as Nd, Dy, and Pr. In this context, CeFeB alloys are considered a promising alternative because of the abundance and low cost of element Ce. In this study, CeFeB alloys (Ce: 35 wt%, Fe: 64 wt%, B: 1 wt%) were produced by the melt-spinning method at different wheel surface speeds of 3, 5, 15, and 25 m/s. Subsequently, the ribbons were subjected to a hydrogen decrepitation (HD) process to investigate the effects of cooling rate and hydrogen-induced cracking on their microstructural and magnetic properties. According to the X-ray diffraction results, all specimens exhibited a partially crystalline structure containing Ce₂Fe₁₄B as the main hard-magnetic phase, along with minor α-Fe and CeFe₂ phases. Increasing the wheel surface speed led to finer grains and more amorphization in the HD specimens. After the HD process, better microstructure and more homogeneous crack distribution were observed, in the specimen produced at the wheel surface speed of 5 m/s. The magnetic measurements revealed that coercivity (Hc), remanence (Mr), and maximum energy product ((BH)max) were strongly dependent on the cooling rate and the HD process. The highest Hc and (BH)max values were obtained as 4129.9 Oe and 13.62 MGOe, respectively, in the HD5 sample. However, formation of more amorphous regions, and oxidation at the higher cooling rates weakened the exchange-interaction mechanism, leading to a decrease in the Hc and (BH)max values. CeFeB alloy melt spinning hydrogen decrepitation magnetic properties Figures Figure 1 Figure 2 Figure 3 Figure 4 Figure 5 1. Introduction Rare earth-based NdFeB alloys have become a very important research topic in both academia and industry since their discovery in the 1980s [ 1 – 3 ]. They are the strongest magnetic alloys at room temperature, with superior coercivity (H c ) and maximum energy product ((BH) max ) [ 4 , 5 ]. Therefore, they play a key role in the advancement of next-generation clean energy systems such as wind turbines, hybrid electric vehicles (HEVs), and other applications. Their widespread use is also increasing day by day [ 6 – 10 ]. However, production of NdFeB alloys is highly dependent on heavy rare earth elements (HREEs) such as Dy, Sm, and Tb [ 11 – 15 ]. This leads to the rapid depletion of the resources containing these elements [ 3 , 16 – 18 ]. Nowadays, permanent magnetic alloys that not based on the HREEs have become increasingly important. The design and production of these alloys aim to overcome the limitations caused by these elements [ 3 , 19 , 20 ]. Among the rare-earth permanent magnetic alloys, CeFeB offers a promising low-cost alternative compared to their NdFeB counterpart. This is thanks to the lower cost and higher availability of the element Ce compared to the element Nd [ 3 , 20 – 25 ]. Nevertheless, these alloys suffer form notable disadvantages. The CeFeB alloys exhibits lower Curie temperature (T C ), magneto-crystallin anisotrophy (H A ), and saturation magnetisation (M s ), comapred to those of the NdFeB alloys. These drawbacks limits their applications [ 20 , 22 , 26 ]. Despite these drawbacks, the CeFeB alloys have potentials for the applications to fill the gap between ferrites and the NdFeB alloys [ 20 ]. Over the past few years, the CeFeB alloys have become the focus of extensive studies. All of these studies aim to enhance the magnetic properties such as coercivity (H c ), remanance (M r ), and maximum energy product ((BH) max ) [ 17 , 20 , 27 , 28 ]. The choice of manufacturing method plays a crucial role in obtaining the CeFeB magnetic alloys with improved magnetic properties [ 10 , 29 , 30 ]. Of these, melt spinning enables to manufacture the alloys with superior magnetic properties (H c , (BH) max , B r ) due to extremely high cooling rates of 10⁶–10⁹ K/s that form amorphous and/or nanocrystalline structures [ 20 , 31 – 34 ]. Nanocrystalline CeFeB alloys have been produced successfully by the melt spinning in previous studies [ 20 , 31 – 36 ]. Recently, hydrogen decrepitation (HD) process has also been utilized to break down the rare-earth-based magnetic alloys through their reaction with the hydrogen gas. This leads to mechanisms of hydride formation, volume expansion, and a highly brittle microstructure, respectively. This intrinsic brittleness induced by the hydrogenation is considered one of the key advantages of the HD technique. Because it facilitates the powder manufacturing, while maintaining the microstructural properties of the alloys [ 37 – 41 ]. The HD enables the manufacturing the fine powders with desirable particle sizes, and enhanced magnetic properties such as the H c and (BH) max . Therefore, it can offer significant advantages in preparing precursor powders for the obtainind the bulk magnetic alloys [ 38 ]. The HD process was first introduced in 1978 by Harris et al. for SmCo₅ alloys [ 38 ], and was later applied to Sm₂(Co,Fe,Cu,Zr)₁₇ alloys with rather high efficiency [ 42 , 43 ]. Following the discovery of the NdFeB alloys in 1983, the HD process was extended to such alloys, enabling the fabrication of the magnetic alloys with superior coercivity and maximum energy product values [ 38 , 44 ]. Comparative studies further revealed that the NdFeB alloys manufactured by HD exhibit the higher H c and a more favorable microstructure, compared to those obtained by conventional mechanical milling [ 45 ]. To the best of our knowledge, no systematic investigations have yet been reported on the hydrogen decrepitation (HD) process to the melt-spun CeFeB alloys. Another gap in the literature is the lack of studies investigating how different wheel surface speeds during the melt-spinning process influence the performance of the HD process in terms of magnetic and microstructural properties. The cooling rate used in the melt spinning method determines the amount of a-Fe soft magnetic phase in the microstructure. During the solidification at the relatively low cooling rates, this phase becomes more prominent [ 25 , 46 ]. These soft α-Fe phases prevent the hard magnetic 2:14:1 phases from magnetically aligning, resulting in deterioration of hard magnetic properties [ 47 ]. To prevent the formation of these α-Fe phases, it is very important to use higher cooling rates during the rapid solidification [ 46 , 48 ]. This highlights an important research topic regarding the systematic investigation of the effect of different wheel surface speeds on the HD process for the melt-spun CeFeB ribbons. Therefore, in the present study, we aim to fill this gap by systematically investigating the influence of the hydrogen decrepitation (HD) process on the microstructural evolution and magnetic performance of melt-spun CeFeB ribbons fabricated at different wheel surface speeds. As a result, the highest coercivity (H c ) and maximum energy product ((BH) max ) were obtained as 4129.9 Oe and 13.62 MGOe, respectively, in the specimen fabricated at the wheel surface speed of 5 m/s. These results indicate that an optimum cooling rate promotes uniform crack distribution and effective grain refinement during the HD process, leading to enhanced hard magnetic properties. Through a comprehensive correlation between the processing parameters, microstructural features, and magnetic properties, this work provides new insights into the optimization of CeFeB-based magnetic materials as cost-effective alternatives to their NdFeB counterparts. 2. Experimental CeFeB master alloys with a nominal composition of Ce 35 Fe 64 B 1 (wt.%) were prepared by vacuum induction melting (VIM). High-purity Ce, Fe, and Fe–B (18 wt.% B) raw materials were used to ensure compositional accuracy and homogeneity. The alloys were remelted and subsequently processed by the melt spinning at the wheel surface speeds of 3, 5, 15, and 25 m/s. The process was carried out under a vacuum of 10⁻⁴ bar, and a partial argon atmosphere of 470 mbar to minimize oxidation. Each of the melt-spun ribbons was subjected to hydrogen decrepitation (HD) in order to obtain fine and brittle powders suitable for subsequent milling stage. For each experiment, approximately 0.5 g of ribbon specimens were placed into hydrogen-resistant stainless steel reactors. Based on the previous studies in the literature, it can be said that the HD process is performed at an average hydrogen gas pressure of around 2 bar and at room temperature for 1 hour [ 38 , 39 , 45 , 49 ]. Therefore, the reactors were then purged and saturated with ultra-high purity hydrogen gas (99.999%) at a pressure of 4 bar, and room temperature (25°C) for 1 h. During the HD process, the rare-earth-rich phases in the CeFeB alloys readily absorbed hydrogen, leading to the formation of hydride phases. As a consequence, the melt-spun ribbon samples became mechanically brittle. After the HD process, a dehydrogenation step was carried out to remove the hydride phases formed in the powders, under a vacuum of 10⁻³ mbar at 300°C for 2 h. Then, the HD samples were milled into fine powders with jet-milling methods. A schematic representation of the hydrogen decrepitation (HD) apparatus employed in this study is presented in Fig. 1 . Phase identification was performed by X-ray diffraction (XRD, Bruker D8 Advance, Cu-Kα radiation, 20°–90°, 2°/min) analysis. Magnetic properties including M s , M r , H c , and (BH) max were determined from hysteresis loops using vibrating specimen magnetometer (VSM) (PPMS DynaCool-9, ± 4 T) analysis. The surface morphology of the samples was examined using a FEI Quanta FEG 250 scanning electron microscope (SEM) equipped with an energy-dispersive X-ray spectroscopy (EDS) detector. Before the microscopic examinations, the ribbon specimens were ground, mechanically polished, and then chemically etched in an HCl–HNO₃ solution containing 25 vol.% HNO₃ to expose the grain boundaries and reveal the microstructural features. 3. Results and Discussion Figure 2 presents the XRD patterns of the CeFeB-based ribbon specimens obtained by the melt-spinning method at the wheel surface speeds of 3, 5, 15, and 25, m/s followed by the HD process. As can be seen, all of the specimens contain paramagnetic CeFe 2 , soft magnetic a-Fe and hard magnetic Ce 2 Fe 14 B phases. Here, the increase in the wheel surface speed corresponding to the higher cooling rates led to broadening and reduction in the intensity of the diffraction peaks. This trend indicates the microstructure, which crystallinity gradually decreases and amorphous structure predominates. The increase in the cooling rate effectively limits long-range atomic diffusion and suppresses crystallization. Therefore, it inhibits the nucleation and subsequent growth of the crystalline phases [ 50 , 51 ]. However, at lower cooling rates, sufficient time for the atomic diffusion facilitates the formation of the nanocrystalline regions with relatively coarser grain sizes [ 52 , 53 ]. From the perspective of the HD process, hydrogen absorption into the lattice induces a local volumetric expansion, generating intergranular and/or transgranular stresses that promote the formation of microcracks. These microcracks play a crucial role in facilitating the refinement of the grain size [ 39 , 40 ]. Apparently, the formation of such finer crystallites became more pronounced in the specimens obtained at the higher wheel surface speeds, where the initial grain size was already smaller due to the higher cooling rate. Therefore, the HD process further facilitates the microstructural refinement in these specimens, leading to a higher proportion of ultrafine grains [ 38 ]. However, the specimens obtained at the lower wheel surface speeds, characterized by relatively coarser grains, underwent less pronounced hydrogen-induced fragmentation. This resulted in a less reduction in the grain size following the HD process. It can be said that the combined effects of the rapid solidification and HD in favor of the evolution of finer, predominantly amorphous microstructure characterized by suppressed crystallization and refined crystallite size. This is consistent with the findings reported in a similar studies about the strip-casting and HD, before [ 54 ]. These nanocrystalline structures are recognized for their positive influence on the hard magnetic behavior of the rare earth permanent magnetic materials [ 55 – 57 ]. The mean crystallite sizes (D) of the specimens were determined using the Debye–Scherrer equation, as given in Eq. 1 [ 1 , 56 ]. As presented in Table I, the crystallite sizes corresponding to the hard, soft, and paramagnetic phases in the HD specimens generally exhibited a reduction trend with increasing wheel surface speed. $$\:D=\frac{K.}{\beta\:.cos}\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\:\left(1\right)$$ In this equation, K represents the Scherrer constant (approximately 0.94 for cubic crystal systems), β denotes the full width at half maximum (FWHM) of the diffraction peak, λ is the X-ray wavelength (1.5406 Å for Cu-Kα radiation), and θ corresponds to the Bragg angle associated with the specific diffraction peak. The crystallite sizes corresponding to each identified phase are summarized in Table I. Based on the Debye–Scherrer eqaluation, the HD CeFeB alloys exhibit crystallite sizes in the range of 24.46–37.71 nm, 16.15–28.37 nm, and 25.40–58.06 nm for the Ce₂Fe₁₄B, α-Fe, and for CeFe₂ phases, respectively. Table I Crystallite size distribution of the soft magnetic (α-Fe), hard magnetic (Ce₂Fe₁₄B), and paramagnetic (CeFe₂) phases identified in the HDed CeFeB specimens. Specimen Calculated Crystallite Size (nm) Ce 2 Fe 14 B a-Fe CeFe 2 HD3 37.71 28.37 58.06 HD5 24.46 23.73 25.40 HD15 33.27 25.24 46.43 HD25 30.71 19.15 35.08 Figure 3 a) and 3b) show the magnetic hysteresis curves along with the corresponding second-quadrant regions of the CeFeB HDed specimens. It can be seen that the experimental data obtained reveal a clear dependence of the magnetic properties on the wheel surface speed and the HD process. Here, the HD3 specimen exhibits the highest saturation magnetization (M s ) of 131.49 emu/g, although the lowest coercivity (H c ) of 2113.82 Oe and maximum energy product ((BH) max ) of 8.10 MGOe. This indicates a dominance of the soft magnetic α-Fe and paramagnetic CeFe₂ phases, which increase the M s value and decrease the H c value. The M s value of the permanent magnetic alloys is closely related to the amount of soft magnetic phases [ 56 , 58 ]. As seen, the XRD pattern of the HD3 specimen indicates more a-Fe and CeFe 2 phases that those of the other specimens. Therefore, the M s value of the HD3 specimen was the highest. Unlike the α-Fe and CeFe 2 phases, the amount of the Ce 2 Fe 14 B phase is quite low. Furthermore, according to the Table I, the average grain size of this phase is larger than (i.e., 37.71 nm) that of the other specimens. In the rare earth magnetic alloys, the content and grain size of the hard magnetic phase affect their permanent magnetic properties (i.e. the H c , M r , and (BH) max ) [ 59 ]. As the grain size in the microstructure coarsens, the number of grain boundaries per unit area decreases. This leads to weaken the pinning effect of the domain walls, which is a significant mechanism for the high coercivity [ 60 ]. Therefore, these factors cause the HD3 specimen to exhibit the weakest permanent magnetic behavior. When the wheel surface speed increased to 5m/s, significant enhancements in the H c and (BH) max values were provided, while the M s value decreased. Here, the H c and (BH) max values ​​increased from 2113.82 Oe to 4129.91 Oe and from 8.10 MGOe to 13.62 MGOe, respectively. However, the M s value decrease from 131.49 emu/g to 118.88 emu/g. This is closely related to the decrease in the grain size of the hard magnetic Ce 2 Fe 14 B phase to 24.46 nm. The obtained coercivity (H c ) and maximum energy product ((BH) max ) values in this study are comparable to those reported in the literature for similar rare-earth magnetic alloys [ 3 , 20 , 27 , 61 , 62 ]. This indicates the effectiveness of the used processing parameters. The decreased grain size provides more grain boundaries for the pinning mechanism of the domain walls [ 63 ]. The XRD pattern of the HD5 specimen reveals to enhance the crystallization of the Ce₂Fe₁₄B hard magnetic phase. This indicates a more ordered and refined microstructure. Small amounts of α-Fe and CeFe₂ phases are still observed. Their average crystallite sizes are relatively small, measured as 22.87 nm and 37.65 nm, respectively. The fine dispersion of these soft magnetic phases minimize their detrimental effect on the overall permanent magnetic properties. As a result, the microstructure becomes more favorable for the domain wall pinning, thus enhancing the H c and (BH) max values. The exchange-interaction mechanism between the hard and soft magnetic phases also affects the H c and (BH) max values [ 56 , 59 ]. The presence of finer-sized hard and soft magnetic phases in the microstructure provides a better exchange-interaction mechanism between them [ 1 , 55 ]. Accordingly, the HD5 specimen has a microstructure exhibiting the optimal exchange-interaction mechanism among the existing specimens. This allows for the optimum H c and (BH) max values ​​to be achieved. The specimens obtained at the higher wheel surface speeds (i.e., HD15 and HD25) showed a fluctuation in the H c and (BH) max values in the ranges of 2568.45–3361.72 Oe, and 8.15–10.94 MGOe, respectively. This can be explained by the more formation of the amorphous phases, i.e., deviation of the microstructure from the optimum state, due to the increased wheel speed [ 64 , 65 ]. This is consistent with the XRD patterns of the corresponding specimens. Because they show the decrease in the crystallinity and increase in the amorphism due to the increased cooling rates, indicating a decrease in the peak intensities of each phase [ 66 ]. Under the influence of these factors, the exchange interaction mechanism between the hard and soft magnetic phases begins to weaken, thus, the H c and (BH) max properties deteriorate. The HD process induces the expansion the crystal lattice and microcrack formation via the hydrogen absorption. This process further refines the grain size [ 37 , 38 , 42 , 43 ]. According to the XRD analysis, in the lower wheel surface speeds, this refinement was at the low level. When the wheel surface speed increased to 5 m/s, the HD process further refined the grain size. Thus, it can be said that the efficiency of the HD process increases due to the increasing wheel surface speed. The evolution of the microstructure and the related changes in the magnetic properties observed in this study are in strong alignment with the findings of Yang et al. Here, the increase in the wheel surface speed and the application of the HD process facilitated the grain refinement and provided the enhanced coercivity. However, the increase in remanance, and hence maximum energy product values ​​did not occured [ 54 ]. The magnetic parameters including H c , M r , M s , and (BH) max , of the HD CeFeB specimens were listed in Table II to facilitate a detailed comparison of their magnetic characteristics. Additionally, Fig. 4 shows the efficiency of the HD process with respect to the magnetic properties, depending on the change in the wheel surface speed. Table II Magnetic properties of the various HDed CeFeB specimens. Specimen Magnetic Properties Coercivity (H c ), Oe Remanance (M r ), emu/g Maximum energy product ((BH) max ), MGOe Saturation magnetization (M s ), emu/g HD3 2113.82 99.37 8.10 131.49 HD5 4129.91 80.78 13.62 118.88 HD15 2568.45 74.84 8.15 113.19 HD25 3361.72 79.06 10.94 124.75 In Fig. 5 , SEM-EDS analysis of the HD specimens are shown. As observed, the variation in wheel surface speed, and the application of the HD process markedly influenced the microstructural evolution of the specimens. Because, increasing the wheel surface speed from 3 to 25 m/s significantly affected the HD behavior due to changes in the initial grain size and crack morphology. As shown in Fig. 5 (a), the HD3 specimen presents a coarse and heterogeneous morphology with irregular shaped particles distributed unevenly across the surface. The surface of the particles exhibits distinct cracks and fracture features formed during the HD process. As known, these microcracks are mainly intergranular in nature. Namely, they propagates along the grain boundaries due to the volume expansion associated with the hydrogen absorption and subsequent formation of brittle hydride phases [ 38 , 40 , 41 , 67 , 68 ]. During the dehydrogenation, these hydrides decompose, but the resulting microcracks remain and create a porous and fragmented structure [ 69 ]. The crack density is not homogeneous throughout the surface; instead, certain regions display more intense cracking formation. The coarse grain morphology, combined with the presence of extensive crack formation weakens the exchange-interaction between the neighboring magnetic grains [ 70 ]. The EDS spectrum shows that the HD3 specimen mainly consists of Fe (70.33 wt.%) and Ce (28.36 wt.%), together with small traces of oxygen (1.10 wt.%) and chlorine (0.51 wt.%). This composition reveals the existence of Fe-rich (α-Fe and CeFe 2 phases) regions, which aligns well with the XRD results. This indicates also insufficient proportion of the hard magnetic Ce₂Fe₁₄B phase for the hard magnetic behavior. The SEM micrograph of the HD5 specimen (Fig. 5 (b)) reveals a significantly refined and homogeneous morphology compared to that of the HD3 sample. The particle surfaces are composed of finer grain structure with a more compact appearance. This indicates the formation of smaller crystallites during the melt spinning at the higher wheel surface speed, followed the HD process. The microcracks is are finer and more uniformly distributed than that in HD3 specimen. In contrast to the wide and irregular cracks observed in the HD3 specimen, those in the HD5 specimen appear narrower. This suggests that the hydrogen absorption–desorption stresses during the HD process are more uniform due to the finer microstructure. The EDS spectrum confirms that the HD5 specimen consists mainly of Fe (64.48 wt.%) and Ce (32.67 wt.%), with oxygen (1.76 wt.%) and chlorine (1.09 wt.%) traces. The higher Ce content suggests an increased fraction of the hard magnetic Ce₂Fe₁₄B phase, reduced amount of α-Fe phase. The balanced Ce/Fe ratio also implies that the alloy is more uniform and showed minimum local compositional segregation at the existing conditions. The fine-grained microstructure, uniformly distributed microcracks and optimized Ce/Fe ratio confirm that the HD5 sample is the most suitable in microstructural and magnetic aspects. As shown in Fig. 5 (c), the SEM image of the HD15 specimen exhibits a highly refined nanocrystalline morphology, consisting of uniformly distributed equiaxed grains. Compared with the HD5 sample, the grain size distribution is narrower and the surface texture appears smoother in certain regions. This suggests more amorphous region formation due to the higher wheel surface speed of 15 m/s. As can be seen, the fracture morphology exhibits finer microcracks and their distributions are more uniform than those in the HD3 and HD5 specimens. According to the EDS analysis, the HD15 specimen contains Fe (70.59 wt.%), Ce (26.08 wt.%), and oxygen (3.33 wt.%). This oxygen enrichment indicates a stronger tendency for surface oxidation after the HD, likely promoted by the higher surface area of the nanocrystalline and amorphous regions. The increase in the amorphous regions and oxidation can weaken the exchange-interaction mechanism, leading to a decrease in the H c and (BH) max values [ 71 ]. At the highest wheel surface speed of 25 m/s, the microstructure is predominantly amorphous. It presents a smooth surface with dispersed nanograins and very fine and shallow cracks. This can be explained by the fact that the relatively high wheel surface speed of 25 m/s significantly suppresses the crystallization. This also confirms that rapid solidification limits the long-range atomic diffusion. This corresponds well with the XRD results showing a decrease in the diffraction intensity of crystalline phases. Here, the microcracks they are very fine and discontinuous, forming shallow grooves rather than deep intergranular cracks. This implies that the material fractured in a more ductile-like manner at the nanoscale, which is typical of amorphous metallic systems [ 72 ]. Although the nanograins can provide the coercivity enhancement, the excessive amorphous content prevented the formation of continuous Ce₂Fe₁₄B networks essential for the high coercivity in this specimen. The EDS analysis revealed a composition of Fe (68.54 wt.%), Ce (30.59 wt.%), and minor oxygen (0.87 wt.%). The relatively low oxygen content compared with the HD5 specimen indicates reduced surface oxidation, likely due to the lower defect density in the amorphous matrix. These microstructural and compositional results align with the resulting magnetic properties of the specimens. 4. Conclusion In this study, the variation in the microstructural and magnetic properties of the hydrogen-decrepitated (HD) CeFeB magnetic alloys was systematically investigated as a function of wheel surface speed during the melt spinning process. The results demonstrated that the CeFeB alloys, which are considered cost-effective alternatives to their NdFeB counterparts, can be successfully obtained with tunable microstructural and magnetic features. The XRD analyses confirmed that all specimens consisted mainly of the hard magnetic Ce₂Fe₁₄B phase, along with minor α-Fe and CeFe₂ phases. With the increase in the wheel surface speed, the crystallite size of the Ce₂Fe₁₄B and α-Fe phases decreased, while the fraction of amorphous and regions gradually increased. Among all specimens, the HD5 sample exhibited the most homogeneous and refined microstructure with a predominant Ce₂Fe₁₄B phase and minimal secondary phase formations. As a result of the refined crystallite size and well-distributed hard/soft phase configuration, remarkable improvements were achieved in the coercivity (H c ), remanence (M r ), and maximum energy product ((BH) max ) values. Accordingly, the highest H c , M r , and (BH) max were measured as 4129.91 Oe, 80.78 emu/g, and 13.62 MGOe, respectively. However, the reduction in the saturation magnetization (M s ) was observed due to the suppression of the soft magnetic α-Fe phase. At higher wheel surface speeds of 15 and 25 m/s, excessive amorphous and oxidized regions weakened the exchange-interaction mechanism, leading to weak in both the H c and (BH) max values. In conclusion, comparison to the conventional mechanical milling method, the hydrogen decrepitation (HD) approach provides an effective means of tailoring the magnetic properties through the formation of finer grain size and a reduced amount of rare earth-rich and α-Fe phase. Moreover, it can be said that optimizing the HD process by adjusting the wheel surface speed during the melt spinning is crucial for further modifying both these properties. Declarations Declaration of competing interest The authors state that they have no competing financial interests or personal relationships that could have appeared to influence the research presented in this article. Author Contribution All authors contributed to the study conception and design. Material preparation, data collection, and analysis were performed by M.F.K. and Y.Y. The first draft of the manuscript was written by Y.Y, and both authors commented on and revised previous versions of the manuscript. All authors read and approved the final manuscript. Acknowledgement This work was financed by the Scientific and Technological Research Council of Turkey (TUBİTAK) under Project Number 223M348, and Scientific Research Projects Coordinator of Sivas University of Science and Technology with Project Numbers of 2023-YLTP-Müh-0005, 2025-GENL-Müh-0001, 2023-ÖAP-Mühe-0002, and 2021-GENL-Müh-0005. References M.F. Kılıçaslan, B. Akgül, Y. Yılmaz, Evolution of the magnetic properties of melt-spun NdFeB alloys with the addition of waste NdFeB magnet. J. Mater. Sci. Mater. Electron. (2023). 10.1007/s10854-023-10571-y L. Zhao, C. Li, Z. Hao, X. Liu, X. Liao, J. Zhang, K. Su, L. Li, H. Yu, J.M. Greneche, J. Jin, Z. Liu, Influences of element segregation on the magnetic properties in nanocrystalline Nd-Ce-Fe-B alloys. Mater. Charact. (2019). 10.1016/j.matchar.2018.12.022 Y.P. Li, J.Y. Li, Z.Z. Wang, Y. Huang, B. Kong, Synthesis of nanocrystalline Ce₁₃Fe₈₁B₆ alloy through mechanically driven disproportionation–recombination process. Mater. Lett. 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(1999). 10.1016/S1359-6454(99)00088-9 Additional Declarations No competing interests reported. Cite Share Download PDF Status: Published Journal Publication published 24 Mar, 2026 Read the published version in Journal of Materials Science: Materials in Electronics → Version 1 posted You are reading this latest preprint version Research Square lets you share your work early, gain feedback from the community, and start making changes to your manuscript prior to peer review in a journal. As a division of Research Square Company, we’re committed to making research communication faster, fairer, and more useful. We do this by developing innovative software and high quality services for the global research community. Our growing team is made up of researchers and industry professionals working together to solve the most critical problems facing scientific publishing. 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16:32:22","extension":"png","order_by":35,"title":"","display":"","copyAsset":false,"role":"acdc-reference","size":306455,"visible":true,"origin":"","legend":"","description":"","filename":"Onlinefloatimage6.png","url":"https://assets-eu.researchsquare.com/files/rs-8354310/v1/ec9f7ea9de589a8fd150ed17.png"},{"id":99218537,"identity":"b3b83ede-9b8f-471b-a992-363ec1b28697","added_by":"auto","created_at":"2025-12-30 09:22:54","extension":"xml","order_by":36,"title":"","display":"","copyAsset":false,"role":"acdc-reference","size":165840,"visible":true,"origin":"","legend":"","description":"","filename":"661b9fb169e64c49b0c3253c8f13ed151structuring.xml","url":"https://assets-eu.researchsquare.com/files/rs-8354310/v1/5c2dfc07b8132e2f768503ec.xml"},{"id":99218541,"identity":"500395f1-764e-44f0-bd42-a0ad038779c6","added_by":"auto","created_at":"2025-12-30 09:22:54","extension":"html","order_by":37,"title":"","display":"","copyAsset":false,"role":"acdc-reference","size":175947,"visible":true,"origin":"","legend":"","description":"","filename":"earlyproof.html","url":"https://assets-eu.researchsquare.com/files/rs-8354310/v1/40284a2902dad54991ec557b.html"},{"id":99317947,"identity":"10ff6def-fd15-4d94-b23e-b6439a076ba3","added_by":"auto","created_at":"2025-12-31 16:30:59","extension":"png","order_by":1,"title":"Figure 1","display":"","copyAsset":false,"role":"figure","size":157626,"visible":true,"origin":"","legend":"\u003cp\u003eThe schematic illustration of the hydrogen decrepitation (HD) and jet milling (JM) processes used to convert the melt-spun CeFeB ribbons into fine alloy powders.\u003c/p\u003e","description":"","filename":"1.png","url":"https://assets-eu.researchsquare.com/files/rs-8354310/v1/71e3765378fd49cebe59cff1.png"},{"id":99218509,"identity":"32db2c32-094b-4778-b92b-b2568a249df7","added_by":"auto","created_at":"2025-12-30 09:22:53","extension":"png","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":143377,"visible":true,"origin":"","legend":"\u003cp\u003eXRD patterns of the hydrogen decrepitated (HDed) CeFeB ribbon specimens produced at different wheel surface speeds of 3, 5, 15, and 25 m/s, respectively.\u003c/p\u003e","description":"","filename":"2.png","url":"https://assets-eu.researchsquare.com/files/rs-8354310/v1/23d0d8de2eadf83eaf12d8e7.png"},{"id":99318171,"identity":"7ad4a1ba-79fd-4cf3-9e5d-3d673823879c","added_by":"auto","created_at":"2025-12-31 16:31:48","extension":"png","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":337358,"visible":true,"origin":"","legend":"\u003cp\u003eMagnetic hysteresis (M–H) loops (a)) and corresponding second-quadrant demagnetization curves (b)) of the HDed CeFeB specimens.\u003c/p\u003e","description":"","filename":"3.png","url":"https://assets-eu.researchsquare.com/files/rs-8354310/v1/b7dbd492cc65bc77d7b795a8.png"},{"id":99318014,"identity":"78825d8f-db29-4d40-a8e1-40b1ae549a7d","added_by":"auto","created_at":"2025-12-31 16:31:10","extension":"png","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":608631,"visible":true,"origin":"","legend":"\u003cp\u003eThe dependence of the magnetic properties of HDed CeFeB specimens on the wheel surface speed during the melt spinning process.\u003c/p\u003e","description":"","filename":"4.png","url":"https://assets-eu.researchsquare.com/files/rs-8354310/v1/57d96a1dd8d71ef1a59edd3f.png"},{"id":99317376,"identity":"c6f4970d-5818-4714-94d4-d987aefb2f32","added_by":"auto","created_at":"2025-12-31 16:30:05","extension":"png","order_by":5,"title":"Figure 5","display":"","copyAsset":false,"role":"figure","size":579478,"visible":true,"origin":"","legend":"\u003cp\u003eSEM-EDS results of the HD CeFeB specimens obtained at different wheel surface speeds: (a) MS3, (b) MS5, (c) MS15, and (d) MS25 samples, respectively.\u003c/p\u003e","description":"","filename":"5.png","url":"https://assets-eu.researchsquare.com/files/rs-8354310/v1/f28b872db9b2134513c724f8.png"},{"id":105755471,"identity":"72455495-47c7-4acf-aae1-96fef0f80043","added_by":"auto","created_at":"2026-03-30 16:27:25","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":2141758,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-8354310/v1/6007085f-ac84-4621-9480-b7e9466a9d0c.pdf"}],"financialInterests":"No competing interests reported.","formattedTitle":"Microstructural and magnetic characterization of hydrogen-decrepitated CeFeB melt-spun ribbons fabricated at various wheel surface speeds","fulltext":[{"header":"1. Introduction","content":"\u003cp\u003eRare earth-based NdFeB alloys have become a very important research topic in both academia and industry since their discovery in the 1980s [\u003cspan additionalcitationids=\"CR2\" citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e]. They are the strongest magnetic alloys at room temperature, with superior coercivity (H\u003csub\u003ec\u003c/sub\u003e) and maximum energy product ((BH)\u003csub\u003emax\u003c/sub\u003e) [\u003cspan citationid=\"CR4\" class=\"CitationRef\"\u003e4\u003c/span\u003e, \u003cspan citationid=\"CR5\" class=\"CitationRef\"\u003e5\u003c/span\u003e]. Therefore, they play a key role in the advancement of next-generation clean energy systems such as wind turbines, hybrid electric vehicles (HEVs), and other applications. Their widespread use is also increasing day by day [\u003cspan additionalcitationids=\"CR7 CR8 CR9\" citationid=\"CR6\" class=\"CitationRef\"\u003e6\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e]. However, production of NdFeB alloys is highly dependent on heavy rare earth elements (HREEs) such as Dy, Sm, and Tb [\u003cspan additionalcitationids=\"CR12 CR13 CR14\" citationid=\"CR11\" class=\"CitationRef\"\u003e11\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e]. This leads to the rapid depletion of the resources containing these elements [\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e, \u003cspan additionalcitationids=\"CR17\" citationid=\"CR16\" class=\"CitationRef\"\u003e16\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR18\" class=\"CitationRef\"\u003e18\u003c/span\u003e]. Nowadays, permanent magnetic alloys that not based on the HREEs have become increasingly important. The design and production of these alloys aim to overcome the limitations caused by these elements [\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e, \u003cspan citationid=\"CR19\" class=\"CitationRef\"\u003e19\u003c/span\u003e, \u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e].\u003c/p\u003e \u003cp\u003eAmong the rare-earth permanent magnetic alloys, CeFeB offers a promising low-cost alternative compared to their NdFeB counterpart. This is thanks to the lower cost and higher availability of the element Ce compared to the element Nd [\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e, \u003cspan additionalcitationids=\"CR21 CR22 CR23 CR24\" citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR25\" class=\"CitationRef\"\u003e25\u003c/span\u003e]. Nevertheless, these alloys suffer form notable disadvantages. The CeFeB alloys exhibits lower Curie temperature (T\u003csub\u003eC\u003c/sub\u003e), magneto-crystallin anisotrophy (H\u003csub\u003eA\u003c/sub\u003e), and saturation magnetisation (M\u003csub\u003es\u003c/sub\u003e), comapred to those of the NdFeB alloys. These drawbacks limits their applications [\u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e, \u003cspan citationid=\"CR22\" class=\"CitationRef\"\u003e22\u003c/span\u003e, \u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e]. Despite these drawbacks, the CeFeB alloys have potentials for the applications to fill the gap between ferrites and the NdFeB alloys [\u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e]. Over the past few years, the CeFeB alloys have become the focus of extensive studies. All of these studies aim to enhance the magnetic properties such as coercivity (H\u003csub\u003ec\u003c/sub\u003e), remanance (M\u003csub\u003er\u003c/sub\u003e), and maximum energy product ((BH)\u003csub\u003emax\u003c/sub\u003e) [\u003cspan citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e, \u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e, \u003cspan citationid=\"CR27\" class=\"CitationRef\"\u003e27\u003c/span\u003e, \u003cspan citationid=\"CR28\" class=\"CitationRef\"\u003e28\u003c/span\u003e]. The choice of manufacturing method plays a crucial role in obtaining the CeFeB magnetic alloys with improved magnetic properties [\u003cspan citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e, \u003cspan citationid=\"CR29\" class=\"CitationRef\"\u003e29\u003c/span\u003e, \u003cspan citationid=\"CR30\" class=\"CitationRef\"\u003e30\u003c/span\u003e]. Of these, melt spinning enables to manufacture the alloys with superior magnetic properties (H\u003csub\u003ec\u003c/sub\u003e, (BH)\u003csub\u003emax\u003c/sub\u003e, B\u003csub\u003er\u003c/sub\u003e) due to extremely high cooling rates of 10⁶\u0026ndash;10⁹ K/s that form amorphous and/or nanocrystalline structures [\u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e, \u003cspan additionalcitationids=\"CR32 CR33\" citationid=\"CR31\" class=\"CitationRef\"\u003e31\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR34\" class=\"CitationRef\"\u003e34\u003c/span\u003e]. Nanocrystalline CeFeB alloys have been produced successfully by the melt spinning in previous studies [\u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e, \u003cspan additionalcitationids=\"CR32 CR33 CR34 CR35\" citationid=\"CR31\" class=\"CitationRef\"\u003e31\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR36\" class=\"CitationRef\"\u003e36\u003c/span\u003e].\u003c/p\u003e \u003cp\u003eRecently, hydrogen decrepitation (HD) process has also been utilized to break down the rare-earth-based magnetic alloys through their reaction with the hydrogen gas. This leads to mechanisms of hydride formation, volume expansion, and a highly brittle microstructure, respectively. This intrinsic brittleness induced by the hydrogenation is considered one of the key advantages of the HD technique. Because it facilitates the powder manufacturing, while maintaining the microstructural properties of the alloys [\u003cspan additionalcitationids=\"CR38 CR39 CR40\" citationid=\"CR37\" class=\"CitationRef\"\u003e37\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR41\" class=\"CitationRef\"\u003e41\u003c/span\u003e]. The HD enables the manufacturing the fine powders with desirable particle sizes, and enhanced magnetic properties such as the H\u003csub\u003ec\u003c/sub\u003e and (BH)\u003csub\u003emax\u003c/sub\u003e. Therefore, it can offer significant advantages in preparing precursor powders for the obtainind the bulk magnetic alloys [\u003cspan citationid=\"CR38\" class=\"CitationRef\"\u003e38\u003c/span\u003e]. The HD process was first introduced in 1978 by Harris et al. for SmCo₅ alloys [\u003cspan citationid=\"CR38\" class=\"CitationRef\"\u003e38\u003c/span\u003e], and was later applied to Sm₂(Co,Fe,Cu,Zr)₁₇ alloys with rather high efficiency [\u003cspan citationid=\"CR42\" class=\"CitationRef\"\u003e42\u003c/span\u003e, \u003cspan citationid=\"CR43\" class=\"CitationRef\"\u003e43\u003c/span\u003e]. Following the discovery of the NdFeB alloys in 1983, the HD process was extended to such alloys, enabling the fabrication of the magnetic alloys with superior coercivity and maximum energy product values [\u003cspan citationid=\"CR38\" class=\"CitationRef\"\u003e38\u003c/span\u003e, \u003cspan citationid=\"CR44\" class=\"CitationRef\"\u003e44\u003c/span\u003e]. Comparative studies further revealed that the NdFeB alloys manufactured by HD exhibit the higher H\u003csub\u003ec\u003c/sub\u003e and a more favorable microstructure, compared to those obtained by conventional mechanical milling [\u003cspan citationid=\"CR45\" class=\"CitationRef\"\u003e45\u003c/span\u003e].\u003c/p\u003e \u003cp\u003eTo the best of our knowledge, no systematic investigations have yet been reported on the hydrogen decrepitation (HD) process to the melt-spun CeFeB alloys. Another gap in the literature is the lack of studies investigating how different wheel surface speeds during the melt-spinning process influence the performance of the HD process in terms of magnetic and microstructural properties. The cooling rate used in the melt spinning method determines the amount of a-Fe soft magnetic phase in the microstructure. During the solidification at the relatively low cooling rates, this phase becomes more prominent [\u003cspan citationid=\"CR25\" class=\"CitationRef\"\u003e25\u003c/span\u003e, \u003cspan citationid=\"CR46\" class=\"CitationRef\"\u003e46\u003c/span\u003e]. These soft α-Fe phases prevent the hard magnetic 2:14:1 phases from magnetically aligning, resulting in deterioration of hard magnetic properties [\u003cspan citationid=\"CR47\" class=\"CitationRef\"\u003e47\u003c/span\u003e]. To prevent the formation of these α-Fe phases, it is very important to use higher cooling rates during the rapid solidification [\u003cspan citationid=\"CR46\" class=\"CitationRef\"\u003e46\u003c/span\u003e, \u003cspan citationid=\"CR48\" class=\"CitationRef\"\u003e48\u003c/span\u003e]. This highlights an important research topic regarding the systematic investigation of the effect of different wheel surface speeds on the HD process for the melt-spun CeFeB ribbons. Therefore, in the present study, we aim to fill this gap by systematically investigating the influence of the hydrogen decrepitation (HD) process on the microstructural evolution and magnetic performance of melt-spun CeFeB ribbons fabricated at different wheel surface speeds. As a result, the highest coercivity (H\u003csub\u003ec\u003c/sub\u003e) and maximum energy product ((BH)\u003csub\u003emax\u003c/sub\u003e) were obtained as 4129.9 Oe and 13.62 MGOe, respectively, in the specimen fabricated at the wheel surface speed of 5 m/s. These results indicate that an optimum cooling rate promotes uniform crack distribution and effective grain refinement during the HD process, leading to enhanced hard magnetic properties. Through a comprehensive correlation between the processing parameters, microstructural features, and magnetic properties, this work provides new insights into the optimization of CeFeB-based magnetic materials as cost-effective alternatives to their NdFeB counterparts.\u003c/p\u003e"},{"header":"2. Experimental","content":"\u003cp\u003eCeFeB master alloys with a nominal composition of Ce\u003csub\u003e35\u003c/sub\u003eFe\u003csub\u003e64\u003c/sub\u003eB\u003csub\u003e1\u003c/sub\u003e (wt.%) were prepared by vacuum induction melting (VIM). High-purity Ce, Fe, and Fe\u0026ndash;B (18 wt.% B) raw materials were used to ensure compositional accuracy and homogeneity. The alloys were remelted and subsequently processed by the melt spinning at the wheel surface speeds of 3, 5, 15, and 25 m/s. The process was carried out under a vacuum of 10⁻⁴ bar, and a partial argon atmosphere of 470 mbar to minimize oxidation. Each of the melt-spun ribbons was subjected to hydrogen decrepitation (HD) in order to obtain fine and brittle powders suitable for subsequent milling stage. For each experiment, approximately 0.5 g of ribbon specimens were placed into hydrogen-resistant stainless steel reactors. Based on the previous studies in the literature, it can be said that the HD process is performed at an average hydrogen gas pressure of around 2 bar and at room temperature for 1 hour [\u003cspan citationid=\"CR38\" class=\"CitationRef\"\u003e38\u003c/span\u003e, \u003cspan citationid=\"CR39\" class=\"CitationRef\"\u003e39\u003c/span\u003e, \u003cspan citationid=\"CR45\" class=\"CitationRef\"\u003e45\u003c/span\u003e, \u003cspan citationid=\"CR49\" class=\"CitationRef\"\u003e49\u003c/span\u003e]. Therefore, the reactors were then purged and saturated with ultra-high purity hydrogen gas (99.999%) at a pressure of 4 bar, and room temperature (25\u0026deg;C) for 1 h. During the HD process, the rare-earth-rich phases in the CeFeB alloys readily absorbed hydrogen, leading to the formation of hydride phases. As a consequence, the melt-spun ribbon samples became mechanically brittle. After the HD process, a dehydrogenation step was carried out to remove the hydride phases formed in the powders, under a vacuum of 10⁻\u0026sup3; mbar at 300\u0026deg;C for 2 h. Then, the HD samples were milled into fine powders with jet-milling methods. A schematic representation of the hydrogen decrepitation (HD) apparatus employed in this study is presented in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003e. Phase identification was performed by X-ray diffraction (XRD, Bruker D8 Advance, Cu-Kα radiation, 20\u0026deg;\u0026ndash;90\u0026deg;, 2\u0026deg;/min) analysis. Magnetic properties including M\u003csub\u003es\u003c/sub\u003e, M\u003csub\u003er\u003c/sub\u003e, H\u003csub\u003ec\u003c/sub\u003e, and (BH)\u003csub\u003emax\u003c/sub\u003e were determined from hysteresis loops using vibrating specimen magnetometer (VSM) (PPMS DynaCool-9, \u0026plusmn;\u0026thinsp;4 T) analysis. The surface morphology of the samples was examined using a FEI Quanta FEG 250 scanning electron microscope (SEM) equipped with an energy-dispersive X-ray spectroscopy (EDS) detector. Before the microscopic examinations, the ribbon specimens were ground, mechanically polished, and then chemically etched in an HCl\u0026ndash;HNO₃ solution containing 25 vol.% HNO₃ to expose the grain boundaries and reveal the microstructural features.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e"},{"header":"3. Results and Discussion","content":"\u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003e presents the XRD patterns of the CeFeB-based ribbon specimens obtained by the melt-spinning method at the wheel surface speeds of 3, 5, 15, and 25, m/s followed by the HD process. As can be seen, all of the specimens contain paramagnetic CeFe\u003csub\u003e2\u003c/sub\u003e, soft magnetic a-Fe and hard magnetic Ce\u003csub\u003e2\u003c/sub\u003eFe\u003csub\u003e14\u003c/sub\u003eB phases. Here, the increase in the wheel surface speed corresponding to the higher cooling rates led to broadening and reduction in the intensity of the diffraction peaks. This trend indicates the microstructure, which crystallinity gradually decreases and amorphous structure predominates. The increase in the cooling rate effectively limits long-range atomic diffusion and suppresses crystallization. Therefore, it inhibits the nucleation and subsequent growth of the crystalline phases [\u003cspan citationid=\"CR50\" class=\"CitationRef\"\u003e50\u003c/span\u003e, \u003cspan citationid=\"CR51\" class=\"CitationRef\"\u003e51\u003c/span\u003e]. However, at lower cooling rates, sufficient time for the atomic diffusion facilitates the formation of the nanocrystalline regions with relatively coarser grain sizes [\u003cspan citationid=\"CR52\" class=\"CitationRef\"\u003e52\u003c/span\u003e, \u003cspan citationid=\"CR53\" class=\"CitationRef\"\u003e53\u003c/span\u003e]. From the perspective of the HD process, hydrogen absorption into the lattice induces a local volumetric expansion, generating intergranular and/or transgranular stresses that promote the formation of microcracks. These microcracks play a crucial role in facilitating the refinement of the grain size [\u003cspan citationid=\"CR39\" class=\"CitationRef\"\u003e39\u003c/span\u003e, \u003cspan citationid=\"CR40\" class=\"CitationRef\"\u003e40\u003c/span\u003e]. Apparently, the formation of such finer crystallites became more pronounced in the specimens obtained at the higher wheel surface speeds, where the initial grain size was already smaller due to the higher cooling rate. Therefore, the HD process further facilitates the microstructural refinement in these specimens, leading to a higher proportion of ultrafine grains [\u003cspan citationid=\"CR38\" class=\"CitationRef\"\u003e38\u003c/span\u003e]. However, the specimens obtained at the lower wheel surface speeds, characterized by relatively coarser grains, underwent less pronounced hydrogen-induced fragmentation. This resulted in a less reduction in the grain size following the HD process. It can be said that the combined effects of the rapid solidification and HD in favor of the evolution of finer, predominantly amorphous microstructure characterized by suppressed crystallization and refined crystallite size. This is consistent with the findings reported in a similar studies about the strip-casting and HD, before [\u003cspan citationid=\"CR54\" class=\"CitationRef\"\u003e54\u003c/span\u003e]. These nanocrystalline structures are recognized for their positive influence on the hard magnetic behavior of the rare earth permanent magnetic materials [\u003cspan additionalcitationids=\"CR56\" citationid=\"CR55\" class=\"CitationRef\"\u003e55\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR57\" class=\"CitationRef\"\u003e57\u003c/span\u003e]. The mean crystallite sizes (D) of the specimens were determined using the Debye\u0026ndash;Scherrer equation, as given in Eq.\u0026nbsp;1 [\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e, \u003cspan citationid=\"CR56\" class=\"CitationRef\"\u003e56\u003c/span\u003e]. As presented in Table I, the crystallite sizes corresponding to the hard, soft, and paramagnetic phases in the HD specimens generally exhibited a reduction trend with increasing wheel surface speed.\u003cdiv id=\"Equa\" class=\"Equation\"\u003e\u003cdiv format=\"TEX\" class=\"mathdisplay\" id=\"FileID_Equa\" name=\"EquationSource\"\u003e\n$$\\:D=\\frac{K.}{\\beta\\:.cos}\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\:\\left(1\\right)$$\u003c/div\u003e\u003c/div\u003e\u003c/p\u003e \u003cp\u003eIn this equation, \u003cem\u003eK\u003c/em\u003e represents the Scherrer constant (approximately 0.94 for cubic crystal systems), \u003cem\u003eβ\u003c/em\u003e denotes the full width at half maximum (FWHM) of the diffraction peak, \u003cem\u003eλ\u003c/em\u003e is the X-ray wavelength (1.5406 \u0026Aring; for Cu-Kα radiation), and \u003cem\u003eθ\u003c/em\u003e corresponds to the Bragg angle associated with the specific diffraction peak. The crystallite sizes corresponding to each identified phase are summarized in Table I. Based on the Debye\u0026ndash;Scherrer eqaluation, the HD CeFeB alloys exhibit crystallite sizes in the range of 24.46\u0026ndash;37.71 nm, 16.15\u0026ndash;28.37 nm, and 25.40\u0026ndash;58.06 nm for the Ce₂Fe₁₄B, α-Fe, and for CeFe₂ phases, respectively.\u003c/p\u003e\u003cp\u003e\u003cstrong\u003eTable I\u0026nbsp;\u003c/strong\u003eCrystallite size distribution of the soft magnetic (\u0026alpha;-Fe), hard magnetic (Ce₂Fe₁₄B), and paramagnetic (CeFe₂) phases identified in the HDed CeFeB specimens.\u003c/p\u003e\n\u003cdiv\u003e\n \u003ctable border=\"1\" cellspacing=\"0\" cellpadding=\"0\"\u003e\n \u003ctbody\u003e\n \u003ctr\u003e\n \u003ctd rowspan=\"2\" valign=\"top\" style=\"width: 113px;\"\u003e\n \u003cp\u003e\u003cstrong\u003eSpecimen\u003c/strong\u003e\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd colspan=\"3\" valign=\"top\" style=\"width: 484px;\"\u003e\n \u003cp\u003e\u003cstrong\u003eCalculated Crystallite Size (nm)\u003c/strong\u003e\u003c/p\u003e\n \u003c/td\u003e\n \u003c/tr\u003e\n \u003ctr\u003e\n \u003ctd valign=\"top\" style=\"width: 142px;\"\u003e\n \u003cp\u003eCe\u003csub\u003e2\u003c/sub\u003eFe\u003csub\u003e14\u003c/sub\u003eB\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" style=\"width: 161px;\"\u003e\n \u003cp\u003ea-Fe\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" style=\"width: 182px;\"\u003e\n \u003cp\u003eCeFe\u003csub\u003e2\u003c/sub\u003e\u003c/p\u003e\n \u003c/td\u003e\n \u003c/tr\u003e\n \u003ctr\u003e\n \u003ctd valign=\"top\" style=\"width: 113px;\"\u003e\n \u003cp\u003eHD3\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" style=\"width: 142px;\"\u003e\n \u003cp\u003e37.71\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" style=\"width: 161px;\"\u003e\n \u003cp\u003e28.37\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" style=\"width: 182px;\"\u003e\n \u003cp\u003e58.06\u003c/p\u003e\n \u003c/td\u003e\n \u003c/tr\u003e\n \u003ctr\u003e\n \u003ctd valign=\"top\" style=\"width: 113px;\"\u003e\n \u003cp\u003eHD5\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" style=\"width: 142px;\"\u003e\n \u003cp\u003e24.46\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" style=\"width: 161px;\"\u003e\n \u003cp\u003e23.73\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" style=\"width: 182px;\"\u003e\n \u003cp\u003e25.40\u003c/p\u003e\n \u003c/td\u003e\n \u003c/tr\u003e\n \u003ctr\u003e\n \u003ctd valign=\"top\" style=\"width: 113px;\"\u003e\n \u003cp\u003eHD15\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" style=\"width: 142px;\"\u003e\n \u003cp\u003e33.27\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" style=\"width: 161px;\"\u003e\n \u003cp\u003e25.24\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" style=\"width: 182px;\"\u003e\n \u003cp\u003e46.43\u003c/p\u003e\n \u003c/td\u003e\n \u003c/tr\u003e\n \u003ctr\u003e\n \u003ctd valign=\"top\" style=\"width: 113px;\"\u003e\n \u003cp\u003eHD25\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" style=\"width: 142px;\"\u003e\n \u003cp\u003e30.71\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" style=\"width: 161px;\"\u003e\n \u003cp\u003e19.15\u003c/p\u003e\n \u003c/td\u003e\n \u003ctd valign=\"top\" style=\"width: 182px;\"\u003e\n \u003cp\u003e35.08\u003c/p\u003e\n \u003c/td\u003e\n \u003c/tr\u003e\n \u003c/tbody\u003e\n \u003c/table\u003e\n\u003c/div\u003e\u003c/br\u003e\u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ea) and 3b) show the magnetic hysteresis curves along with the corresponding second-quadrant regions of the CeFeB HDed specimens. It can be seen that the experimental data obtained reveal a clear dependence of the magnetic properties on the wheel surface speed and the HD process. Here, the HD3 specimen exhibits the highest saturation magnetization (M\u003csub\u003es\u003c/sub\u003e) of 131.49 emu/g, although the lowest coercivity (H\u003csub\u003ec\u003c/sub\u003e) of 2113.82 Oe and maximum energy product ((BH)\u003csub\u003emax\u003c/sub\u003e) of 8.10 MGOe. This indicates a dominance of the soft magnetic α-Fe and paramagnetic CeFe₂ phases, which increase the M\u003csub\u003es\u003c/sub\u003e value and decrease the H\u003csub\u003ec\u003c/sub\u003e value. The M\u003csub\u003es\u003c/sub\u003e value of the permanent magnetic alloys is closely related to the amount of soft magnetic phases [\u003cspan citationid=\"CR56\" class=\"CitationRef\"\u003e56\u003c/span\u003e, \u003cspan citationid=\"CR58\" class=\"CitationRef\"\u003e58\u003c/span\u003e]. As seen, the XRD pattern of the HD3 specimen indicates more a-Fe and CeFe\u003csub\u003e2\u003c/sub\u003e phases that those of the other specimens. Therefore, the M\u003csub\u003es\u003c/sub\u003e value of the HD3 specimen was the highest. Unlike the α-Fe and CeFe\u003csub\u003e2\u003c/sub\u003e phases, the amount of the Ce\u003csub\u003e2\u003c/sub\u003eFe\u003csub\u003e14\u003c/sub\u003eB phase is quite low. Furthermore, according to the Table I, the average grain size of this phase is larger than (i.e., 37.71 nm) that of the other specimens. In the rare earth magnetic alloys, the content and grain size of the hard magnetic phase affect their permanent magnetic properties (i.e. the H\u003csub\u003ec\u003c/sub\u003e, M\u003csub\u003er\u003c/sub\u003e, and (BH)\u003csub\u003emax\u003c/sub\u003e) [\u003cspan citationid=\"CR59\" class=\"CitationRef\"\u003e59\u003c/span\u003e]. As the grain size in the microstructure coarsens, the number of grain boundaries per unit area decreases. This leads to weaken the pinning effect of the domain walls, which is a significant mechanism for the high coercivity [\u003cspan citationid=\"CR60\" class=\"CitationRef\"\u003e60\u003c/span\u003e]. Therefore, these factors cause the HD3 specimen to exhibit the weakest permanent magnetic behavior. When the wheel surface speed increased to 5m/s, significant enhancements in the H\u003csub\u003ec\u003c/sub\u003e and (BH)\u003csub\u003emax\u003c/sub\u003e values were provided, while the M\u003csub\u003es\u003c/sub\u003e value decreased. Here, the H\u003csub\u003ec\u003c/sub\u003e and (BH)\u003csub\u003emax\u003c/sub\u003e values ​​increased from 2113.82 Oe to 4129.91 Oe and from 8.10 MGOe to 13.62 MGOe, respectively. However, the M\u003csub\u003es\u003c/sub\u003e value decrease from 131.49 emu/g to 118.88 emu/g. This is closely related to the decrease in the grain size of the hard magnetic Ce\u003csub\u003e2\u003c/sub\u003eFe\u003csub\u003e14\u003c/sub\u003eB phase to 24.46 nm. The obtained coercivity (H\u003csub\u003ec\u003c/sub\u003e) and maximum energy product ((BH)\u003csub\u003emax\u003c/sub\u003e) values in this study are comparable to those reported in the literature for similar rare-earth magnetic alloys [\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e, \u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e, \u003cspan citationid=\"CR27\" class=\"CitationRef\"\u003e27\u003c/span\u003e, \u003cspan citationid=\"CR61\" class=\"CitationRef\"\u003e61\u003c/span\u003e, \u003cspan citationid=\"CR62\" class=\"CitationRef\"\u003e62\u003c/span\u003e]. This indicates the effectiveness of the used processing parameters. The decreased grain size provides more grain boundaries for the pinning mechanism of the domain walls [\u003cspan citationid=\"CR63\" class=\"CitationRef\"\u003e63\u003c/span\u003e]. The XRD pattern of the HD5 specimen reveals to enhance the crystallization of the Ce₂Fe₁₄B hard magnetic phase. This indicates a more ordered and refined microstructure. Small amounts of α-Fe and CeFe₂ phases are still observed. Their average crystallite sizes are relatively small, measured as 22.87 nm and 37.65 nm, respectively. The fine dispersion of these soft magnetic phases minimize their detrimental effect on the overall permanent magnetic properties. As a result, the microstructure becomes more favorable for the domain wall pinning, thus enhancing the H\u003csub\u003ec\u003c/sub\u003e and (BH)\u003csub\u003emax\u003c/sub\u003e values. The exchange-interaction mechanism between the hard and soft magnetic phases also affects the H\u003csub\u003ec\u003c/sub\u003e and (BH)\u003csub\u003emax\u003c/sub\u003e values [\u003cspan citationid=\"CR56\" class=\"CitationRef\"\u003e56\u003c/span\u003e, \u003cspan citationid=\"CR59\" class=\"CitationRef\"\u003e59\u003c/span\u003e]. The presence of finer-sized hard and soft magnetic phases in the microstructure provides a better exchange-interaction mechanism between them [\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e, \u003cspan citationid=\"CR55\" class=\"CitationRef\"\u003e55\u003c/span\u003e]. Accordingly, the HD5 specimen has a microstructure exhibiting the optimal exchange-interaction mechanism among the existing specimens. This allows for the optimum H\u003csub\u003ec\u003c/sub\u003e and (BH)\u003csub\u003emax\u003c/sub\u003e values ​​to be achieved. The specimens obtained at the higher wheel surface speeds (i.e., HD15 and HD25) showed a fluctuation in the H\u003csub\u003ec\u003c/sub\u003e and (BH)\u003csub\u003emax\u003c/sub\u003e values in the ranges of 2568.45\u0026ndash;3361.72 Oe, and 8.15\u0026ndash;10.94 MGOe, respectively. This can be explained by the more formation of the amorphous phases, i.e., deviation of the microstructure from the optimum state, due to the increased wheel speed [\u003cspan citationid=\"CR64\" class=\"CitationRef\"\u003e64\u003c/span\u003e, \u003cspan citationid=\"CR65\" class=\"CitationRef\"\u003e65\u003c/span\u003e]. This is consistent with the XRD patterns of the corresponding specimens. Because they show the decrease in the crystallinity and increase in the amorphism due to the increased cooling rates, indicating a decrease in the peak intensities of each phase [\u003cspan citationid=\"CR66\" class=\"CitationRef\"\u003e66\u003c/span\u003e]. Under the influence of these factors, the exchange interaction mechanism between the hard and soft magnetic phases begins to weaken, thus, the H\u003csub\u003ec\u003c/sub\u003e and (BH)\u003csub\u003emax\u003c/sub\u003e properties deteriorate. The HD process induces the expansion the crystal lattice and microcrack formation via the hydrogen absorption. This process further refines the grain size [\u003cspan citationid=\"CR37\" class=\"CitationRef\"\u003e37\u003c/span\u003e, \u003cspan citationid=\"CR38\" class=\"CitationRef\"\u003e38\u003c/span\u003e, \u003cspan citationid=\"CR42\" class=\"CitationRef\"\u003e42\u003c/span\u003e, \u003cspan citationid=\"CR43\" class=\"CitationRef\"\u003e43\u003c/span\u003e]. According to the XRD analysis, in the lower wheel surface speeds, this refinement was at the low level. When the wheel surface speed increased to 5 m/s, the HD process further refined the grain size. Thus, it can be said that the efficiency of the HD process increases due to the increasing wheel surface speed. The evolution of the microstructure and the related changes in the magnetic properties observed in this study are in strong alignment with the findings of Yang et al. Here, the increase in the wheel surface speed and the application of the HD process facilitated the grain refinement and provided the enhanced coercivity. However, the increase in remanance, and hence maximum energy product values ​​did not occured [\u003cspan citationid=\"CR54\" class=\"CitationRef\"\u003e54\u003c/span\u003e]. The magnetic parameters including H\u003csub\u003ec\u003c/sub\u003e, M\u003csub\u003er\u003c/sub\u003e, M\u003csub\u003es\u003c/sub\u003e, and (BH)\u003csub\u003emax\u003c/sub\u003e, of the HD CeFeB specimens were listed in Table II to facilitate a detailed comparison of their magnetic characteristics. Additionally, Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003e shows the efficiency of the HD process with respect to the magnetic properties, depending on the change in the wheel surface speed.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003e \u003cb\u003eTable II\u003c/b\u003e Magnetic properties of the various HDed CeFeB specimens.\u003c/p\u003e \u003cp\u003e \u003cdiv class=\"gridtable\"\u003e\u003ctable float=\"No\" id=\"Taba\" border=\"1\"\u003e \u003ccolgroup cols=\"5\"\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c1\" colnum=\"1\"\u003e\u003c/div\u003e \u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c2\" colnum=\"2\"\u003e\u003c/div\u003e \u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c3\" colnum=\"3\"\u003e\u003c/div\u003e \u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c4\" colnum=\"4\"\u003e\u003c/div\u003e \u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c5\" colnum=\"5\"\u003e\u003c/div\u003e \u003cthead\u003e \u003ctr\u003e \u003cth align=\"left\" colname=\"c1\" morerows=\"1\" rowspan=\"2\"\u003e \u003cp\u003eSpecimen\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colspan=\"4\" nameend=\"c5\" namest=\"c2\"\u003e \u003cp\u003eMagnetic Properties\u003c/p\u003e \u003c/th\u003e \u003c/tr\u003e \u003ctr\u003e \u003cth align=\"left\" colname=\"c2\"\u003e \u003cp\u003eCoercivity (H\u003csub\u003ec\u003c/sub\u003e), Oe\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c3\"\u003e \u003cp\u003eRemanance (M\u003csub\u003er\u003c/sub\u003e), emu/g\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c4\"\u003e \u003cp\u003eMaximum energy product ((BH)\u003csub\u003emax\u003c/sub\u003e), MGOe\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c5\"\u003e \u003cp\u003eSaturation magnetization (M\u003csub\u003es\u003c/sub\u003e), emu/g\u003c/p\u003e \u003c/th\u003e \u003c/tr\u003e \u003c/thead\u003e \u003ctbody\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eHD3\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e \u003cp\u003e2113.82\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e \u003cp\u003e99.37\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e \u003cp\u003e8.10\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c5\"\u003e \u003cp\u003e131.49\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eHD5\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e \u003cp\u003e4129.91\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e \u003cp\u003e80.78\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e \u003cp\u003e13.62\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c5\"\u003e \u003cp\u003e118.88\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eHD15\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e \u003cp\u003e2568.45\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e \u003cp\u003e74.84\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e \u003cp\u003e8.15\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c5\"\u003e \u003cp\u003e113.19\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eHD25\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e \u003cp\u003e3361.72\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e \u003cp\u003e79.06\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e \u003cp\u003e10.94\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c5\"\u003e \u003cp\u003e124.75\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003c/tbody\u003e \u003c/colgroup\u003e \u003c/table\u003e\u003c/div\u003e \u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eIn Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e5\u003c/span\u003e, SEM-EDS analysis of the HD specimens are shown. As observed, the variation in wheel surface speed, and the application of the HD process markedly influenced the microstructural evolution of the specimens. Because, increasing the wheel surface speed from 3 to 25 m/s significantly affected the HD behavior due to changes in the initial grain size and crack morphology. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e5\u003c/span\u003e(a), the HD3 specimen presents a coarse and heterogeneous morphology with irregular shaped particles distributed unevenly across the surface. The surface of the particles exhibits distinct cracks and fracture features formed during the HD process. As known, these microcracks are mainly intergranular in nature. Namely, they propagates along the grain boundaries due to the volume expansion associated with the hydrogen absorption and subsequent formation of brittle hydride phases [\u003cspan citationid=\"CR38\" class=\"CitationRef\"\u003e38\u003c/span\u003e, \u003cspan citationid=\"CR40\" class=\"CitationRef\"\u003e40\u003c/span\u003e, \u003cspan citationid=\"CR41\" class=\"CitationRef\"\u003e41\u003c/span\u003e, \u003cspan citationid=\"CR67\" class=\"CitationRef\"\u003e67\u003c/span\u003e, \u003cspan citationid=\"CR68\" class=\"CitationRef\"\u003e68\u003c/span\u003e]. During the dehydrogenation, these hydrides decompose, but the resulting microcracks remain and create a porous and fragmented structure [\u003cspan citationid=\"CR69\" class=\"CitationRef\"\u003e69\u003c/span\u003e]. The crack density is not homogeneous throughout the surface; instead, certain regions display more intense cracking formation. The coarse grain morphology, combined with the presence of extensive crack formation weakens the exchange-interaction between the neighboring magnetic grains [\u003cspan citationid=\"CR70\" class=\"CitationRef\"\u003e70\u003c/span\u003e]. The EDS spectrum shows that the HD3 specimen mainly consists of Fe (70.33 wt.%) and Ce (28.36 wt.%), together with small traces of oxygen (1.10 wt.%) and chlorine (0.51 wt.%). This composition reveals the existence of Fe-rich (α-Fe and CeFe\u003csub\u003e2\u003c/sub\u003e phases) regions, which aligns well with the XRD results. This indicates also insufficient proportion of the hard magnetic Ce₂Fe₁₄B phase for the hard magnetic behavior. The SEM micrograph of the HD5 specimen (Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e5\u003c/span\u003e(b)) reveals a significantly refined and homogeneous morphology compared to that of the HD3 sample. The particle surfaces are composed of finer grain structure with a more compact appearance. This indicates the formation of smaller crystallites during the melt spinning at the higher wheel surface speed, followed the HD process. The microcracks is are finer and more uniformly distributed than that in HD3 specimen. In contrast to the wide and irregular cracks observed in the HD3 specimen, those in the HD5 specimen appear narrower. This suggests that the hydrogen absorption\u0026ndash;desorption stresses during the HD process are more uniform due to the finer microstructure. The EDS spectrum confirms that the HD5 specimen consists mainly of Fe (64.48 wt.%) and Ce (32.67 wt.%), with oxygen (1.76 wt.%) and chlorine (1.09 wt.%) traces. The higher Ce content suggests an increased fraction of the hard magnetic Ce₂Fe₁₄B phase, reduced amount of α-Fe phase. The balanced Ce/Fe ratio also implies that the alloy is more uniform and showed minimum local compositional segregation at the existing conditions. The fine-grained microstructure, uniformly distributed microcracks and optimized Ce/Fe ratio confirm that the HD5 sample is the most suitable in microstructural and magnetic aspects. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e5\u003c/span\u003e(c), the SEM image of the HD15 specimen exhibits a highly refined nanocrystalline morphology, consisting of uniformly distributed equiaxed grains. Compared with the HD5 sample, the grain size distribution is narrower and the surface texture appears smoother in certain regions. This suggests more amorphous region formation due to the higher wheel surface speed of 15 m/s. As can be seen, the fracture morphology exhibits finer microcracks and their distributions are more uniform than those in the HD3 and HD5 specimens. According to the EDS analysis, the HD15 specimen contains Fe (70.59 wt.%), Ce (26.08 wt.%), and oxygen (3.33 wt.%). This oxygen enrichment indicates a stronger tendency for surface oxidation after the HD, likely promoted by the higher surface area of the nanocrystalline and amorphous regions. The increase in the amorphous regions and oxidation can weaken the exchange-interaction mechanism, leading to a decrease in the H\u003csub\u003ec\u003c/sub\u003e and (BH)\u003csub\u003emax\u003c/sub\u003e values [\u003cspan citationid=\"CR71\" class=\"CitationRef\"\u003e71\u003c/span\u003e]. At the highest wheel surface speed of 25 m/s, the microstructure is predominantly amorphous. It presents a smooth surface with dispersed nanograins and very fine and shallow cracks. This can be explained by the fact that the relatively high wheel surface speed of 25 m/s significantly suppresses the crystallization. This also confirms that rapid solidification limits the long-range atomic diffusion. This corresponds well with the XRD results showing a decrease in the diffraction intensity of crystalline phases. Here, the microcracks they are very fine and discontinuous, forming shallow grooves rather than deep intergranular cracks. This implies that the material fractured in a more ductile-like manner at the nanoscale, which is typical of amorphous metallic systems [\u003cspan citationid=\"CR72\" class=\"CitationRef\"\u003e72\u003c/span\u003e]. Although the nanograins can provide the coercivity enhancement, the excessive amorphous content prevented the formation of continuous Ce₂Fe₁₄B networks essential for the high coercivity in this specimen. The EDS analysis revealed a composition of Fe (68.54 wt.%), Ce (30.59 wt.%), and minor oxygen (0.87 wt.%). The relatively low oxygen content compared with the HD5 specimen indicates reduced surface oxidation, likely due to the lower defect density in the amorphous matrix. These microstructural and compositional results align with the resulting magnetic properties of the specimens.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003e \u003c/p\u003e"},{"header":"4. Conclusion","content":"\u003cp\u003eIn this study, the variation in the microstructural and magnetic properties of the hydrogen-decrepitated (HD) CeFeB magnetic alloys was systematically investigated as a function of wheel surface speed during the melt spinning process. The results demonstrated that the CeFeB alloys, which are considered cost-effective alternatives to their NdFeB counterparts, can be successfully obtained with tunable microstructural and magnetic features. The XRD analyses confirmed that all specimens consisted mainly of the hard magnetic Ce₂Fe₁₄B phase, along with minor α-Fe and CeFe₂ phases. With the increase in the wheel surface speed, the crystallite size of the Ce₂Fe₁₄B and α-Fe phases decreased, while the fraction of amorphous and regions gradually increased.\u003c/p\u003e \u003cp\u003eAmong all specimens, the HD5 sample exhibited the most homogeneous and refined microstructure with a predominant Ce₂Fe₁₄B phase and minimal secondary phase formations. As a result of the refined crystallite size and well-distributed hard/soft phase configuration, remarkable improvements were achieved in the coercivity (H\u003csub\u003ec\u003c/sub\u003e), remanence (M\u003csub\u003er\u003c/sub\u003e), and maximum energy product ((BH)\u003csub\u003emax\u003c/sub\u003e) values. Accordingly, the highest H\u003csub\u003ec\u003c/sub\u003e, M\u003csub\u003er\u003c/sub\u003e, and (BH)\u003csub\u003emax\u003c/sub\u003e were measured as 4129.91 Oe, 80.78 emu/g, and 13.62 MGOe, respectively. However, the reduction in the saturation magnetization (M\u003csub\u003es\u003c/sub\u003e) was observed due to the suppression of the soft magnetic α-Fe phase. At higher wheel surface speeds of 15 and 25 m/s, excessive amorphous and oxidized regions weakened the exchange-interaction mechanism, leading to weak in both the H\u003csub\u003ec\u003c/sub\u003e and (BH)\u003csub\u003emax\u003c/sub\u003e values.\u003c/p\u003e \u003cp\u003eIn conclusion, comparison to the conventional mechanical milling method, the hydrogen decrepitation (HD) approach provides an effective means of tailoring the magnetic properties through the formation of finer grain size and a reduced amount of rare earth-rich and α-Fe phase. Moreover, it can be said that optimizing the HD process by adjusting the wheel surface speed during the melt spinning is crucial for further modifying both these properties.\u003c/p\u003e"},{"header":"Declarations","content":"\u003ch2\u003eDeclaration of competing interest\u003c/h2\u003e \u003cp\u003eThe authors state that they have no competing financial interests or personal relationships that could have appeared to influence the research presented in this article.\u003c/p\u003e\u003ch2\u003eAuthor Contribution\u003c/h2\u003e\u003cp\u003eAll authors contributed to the study conception and design. Material preparation, data collection, and analysis were performed by M.F.K. and Y.Y. The first draft of the manuscript was written by Y.Y, and both authors commented on and revised previous versions of the manuscript. All authors read and approved the final manuscript.\u003c/p\u003e\u003ch2\u003eAcknowledgement\u003c/h2\u003e\u003cp\u003eThis work was financed by the Scientific and Technological Research Council of Turkey (TUBİTAK) under Project Number 223M348, and Scientific Research Projects Coordinator of Sivas University of Science and Technology with Project Numbers of 2023-YLTP-M\u0026uuml;h-0005, 2025-GENL-M\u0026uuml;h-0001, 2023-\u0026Ouml;AP-M\u0026uuml;he-0002, and 2021-GENL-M\u0026uuml;h-0005.\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\u003cli\u003e\u003cspan\u003eM.F. Kılı\u0026ccedil;aslan, B. Akg\u0026uuml;l, Y. Yılmaz, Evolution of the magnetic properties of melt-spun NdFeB alloys with the addition of waste NdFeB magnet. J. Mater. Sci. Mater. Electron. 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In this context, CeFeB alloys are considered a promising alternative because of the abundance and low cost of element Ce. In this study, CeFeB alloys (Ce: 35 wt%, Fe: 64 wt%, B: 1 wt%) were produced by the melt-spinning method at different wheel surface speeds of 3, 5, 15, and 25 m/s. Subsequently, the ribbons were subjected to a hydrogen decrepitation (HD) process to investigate the effects of cooling rate and hydrogen-induced cracking on their microstructural and magnetic properties. According to the X-ray diffraction results, all specimens exhibited a partially crystalline structure containing Ce₂Fe₁₄B as the main hard-magnetic phase, along with minor α-Fe and CeFe₂ phases. Increasing the wheel surface speed led to finer grains and more amorphization in the HD specimens. After the HD process, better microstructure and more homogeneous crack distribution were observed, in the specimen produced at the wheel surface speed of 5 m/s. The magnetic measurements revealed that coercivity (Hc), remanence (Mr), and maximum energy product ((BH)max) were strongly dependent on the cooling rate and the HD process. The highest Hc and (BH)max values were obtained as 4129.9 Oe and 13.62 MGOe, respectively, in the HD5 sample. However, formation of more amorphous regions, and oxidation at the higher cooling rates weakened the exchange-interaction mechanism, leading to a decrease in the Hc and (BH)max values.","manuscriptTitle":"Microstructural and magnetic characterization of hydrogen-decrepitated CeFeB melt-spun ribbons fabricated at various wheel surface speeds","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2025-12-30 09:22:48","doi":"10.21203/rs.3.rs-8354310/v1","editorialEvents":[{"type":"communityComments","content":0}],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"researchsquare","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":true,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"/submission","title":"Research Square","twitterHandle":"researchsquare","acdcEnabled":true,"dfaEnabled":false,"editorialSystem":"","reportingPortfolio":"","inReviewEnabled":false,"inReviewRevisionsEnabled":true}}],"origin":"","ownerIdentity":"ddd5e28a-99fb-40dc-be44-d870e297b328","owner":[],"postedDate":"December 30th, 2025","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"posted","subjectAreas":[],"tags":[],"updatedAt":"2026-03-30T16:23:29+00:00","versionOfRecord":{"articleIdentity":"rs-8354310","link":"https://doi.org/10.1007/s10854-026-16992-9","journal":{"identity":"journal-of-materials-science-materials-in-electronics","isVorOnly":false,"title":"Journal of Materials Science: Materials in Electronics"},"publishedOn":"2026-03-24 16:12:07","publishedOnDateReadable":"March 24th, 2026"},"versionCreatedAt":"2025-12-30 09:22:48","video":"","vorDoi":"10.1007/s10854-026-16992-9","vorDoiUrl":"https://doi.org/10.1007/s10854-026-16992-9","workflowStages":[]},"version":"v1","identity":"rs-8354310","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-8354310","identity":"rs-8354310","version":["v1"]},"buildId":"8U1c8b4HqxoKbykW_rLl7","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}