Mechanically-Induced Solid-State Reaction for Fabrication of Soft Magnetic (Co75Ti25)100−xBx (x: 2, 5, 10, 15, 20, 25 at%) Metallic Glassy Nanopowders

Metallic glasses, with their short-range order structure, exhibit unique characteristics that do not exist in the corresponding crystalline alloys with the same compositions. These unusual properties are attributed to the absence of translational periodicity, grain boundaries, and compositional homogeneity. Cobalt (Co)-based metallic glassy alloys have been receiving great attention due to their superior mechanical and magnetic properties. Unluckily, Co-Ti alloys and its based alloys are difficult to be prepared in glassy form, due to their rather poor glass-forming ability. In the present work, the mechanical alloying approach was employed to investigate the possibility of preparing homogeneous (Co75Ti25)100−xBx starting from elemental powders. The feedstock materials with the desired compositions were high-energy ball-milled under argon atmosphere for 50 h. The end products of the powders obtained after milling revealed a short-range order structure with a broad amorphization range (2 at% ≤ B ≤ 25 at%). The behaviors of these glassy systems, characterized by the supercooled liquid region, and reduced glass transition temperature, were improved upon increasing B molar fraction. The results had shown that when B content increased, the saturation magnetization was increased, where coercivity was decreased.


Introduction
Life in the 21st century cannot depend on limited groups of materials; instead, it is dependent on unlimited families of advanced materials [1]. Despite the traditional categories of materials, which may not completely match with the modern industrial requirements, a rather newcomer so-called "metallic glass" [2] has found an important space in the functional classifications of metals and metal alloys [3]. Metallic glasses, with their short-range order structure, possess exciting properties which are of interest not only for basic solid-state physics, but also for metallurgy, surface chemistry, and technology [4,5]. These properties are quite different from crystalline materials (long-range order) metals, making them promising candidates for technical applications. The high mechanical ductility and yield strength, unusual corrosion resistance, high magnetic permeability, and low coercive forces are some of those desirable properties found in metallic glassy alloys [6][7][8][9].
There are different approaches used for preparing amorphous and metallic glassy alloys, such as the atomic disordering of crystalline lattices, solid-state amorphization reaction between pure elements, and solid-state transformations from metastable phase [10]; however, they are exclusively synthesized by rapid solidification (RS) of melts or vapors, using the melt spinning technique [2]. In the melt spinning

Morphology and Crystal Structure
The typical morphological characteristics of (Co 75 Ti 25 ) 100−x B x , exemplified by (Co 75 Ti 25 ) 75 B 25 powders, obtained after 6 h of MA time, is presented in Figure 1. The powder particles, trapped between Molecules 2020, 25, 3338 3 of 15 colliding grinding balls were experienced from excessive welding during the BM process, where metallic Co and Ti were severely plastically deformed and agglomerated to form larger powders of more than 100 µm in diameter, as presented in Figure 1a. These aggregated particles composited of Co, Ti, and B, as confirmed by the energy-dispersive X-ray spectroscopy (EDS)-elemental mapping, as shown in Figure 1b-d, respectively. At this early stage of mechanical alloying (MA), the composition of the milled powders significantly varied from particle to particle and within the individual particles themselves. The field-emission scanning electron microscope (FE-SEM) micrograph of the cross-sectional view for the powders presented in Figure 1a is displayed in Figure 1e. The powders revealed lamellar-like metallography, composited of thick Co, Ti metallic layers, where the fine B particles were dispersed into these layers, as indexed in Figure 1e. where metallic Co and Ti were severely plastically deformed and agglomerated to form larger powders of more than 100 μm in diameter, as presented in Figure 1a. These aggregated particles composited of Co, Ti, and B, as confirmed by the energy-dispersive X-ray spectroscopy (EDS)elemental mapping, as shown in Figure 1b-d, respectively. At this early stage of mechanical alloying (MA), the composition of the milled powders significantly varied from particle to particle and within the individual particles themselves. The field-emission scanning electron microscope (FE-SEM) micrograph of the cross-sectional view for the powders presented in Figure 1a is displayed in Figure  1e. The powders revealed lamellar-like metallography, composited of thick Co, Ti metallic layers, where the fine B particles were dispersed into these layers, as indexed in Figure 1e.   Figure 2a. Formation of these thin-layers accelerated the solid-state diffusion that occurred at their fresh interfaces to form a more homogeneous composition. Meanwhile, the powder particles were continuously subjected to severe lattice imperfections, indexed by plastic deformation, and lattice and point defected, as shown in Figure 2b. After this intermediate stage of MA, the powders consisted of polycrystalline grains of the starting materials with no evidence for the formation metallic glassy phase, as characterized by those sharp spots that appeared in the selected area diffraction pattern (SADP) (Figure 2c).   Figure 2a. Formation of these thin-layers accelerated the solid-state diffusion that occurred at their fresh interfaces to form a more homogeneous composition. Meanwhile, the powder particles were continuously subjected to severe lattice imperfections, indexed by plastic deformation, and lattice and point defected, as shown in Figure 2b. After this intermediate stage of MA, the powders consisted of polycrystalline grains of the starting materials with no evidence for the formation metallic glassy phase, as characterized by those sharp spots that appeared in the selected area diffraction pattern (SADP) (Figure 2c). With further MA time (50 h), the powders were disintegrated into finer particles (Figure 3a,b), where the internal structure of the powders continues to be refined. After 50 h of MA time, the lamella-like metallography completely disappeared, as shown in the high-magnification FE-SEM presented in Figure 3a. This implies the completion of the MA process and the formation of a single homogeneous phase. The powders of this final product possessed desirable morphological characteristics of being ultrafine (ranged from 120 nm to 250 nm in diameter) and had spherical-like morphology with smooth surfaces, as displayed in Figure 3b. The field emission high-resolution transmission electron microscope (FE-HRTEM) image of the final product obtained after 50 h of MA time ( Figure 3c) revealed featureless maze-like morphology of an amorphous phase. Besides, the corresponding nanobeam diffraction pattern (NBDP) showed a typical amorphous-like halo-diffuse pattern, as displayed in Figure 3d.  With further MA time (50 h), the powders were disintegrated into finer particles (Figure 3a,b), where the internal structure of the powders continues to be refined. After 50 h of MA time, the lamella-like metallography completely disappeared, as shown in the high-magnification FE-SEM presented in Figure 3a. This implies the completion of the MA process and the formation of a single homogeneous phase. The powders of this final product possessed desirable morphological characteristics of being ultrafine (ranged from 120 nm to 250 nm in diameter) and had spherical-like morphology with smooth surfaces, as displayed in Figure 3b. The field emission high-resolution transmission electron microscope (FE-HRTEM) image of the final product obtained after 50 h of MA time ( Figure 3c) revealed featureless maze-like morphology of an amorphous phase. Besides, the corresponding nanobeam diffraction pattern (NBDP) showed a typical amorphous-like halo-diffuse pattern, as displayed in Figure 3d.   The X-ray diffraction (XRD) patterns of MA-(Co75Ti25)100−xBx (x: 2, 10, and 15 at%) obtained after 50 h of BM time are displayed collectively in Figure 4. All the patterns exhibited only broad diffuse haloes, in scattering range of 2θ between 40-50°, with no indication of any unprocessed crystalline phases. This implying the capability of MA approach to fabricate the desired amorphous alloy powders in a wide amorphous range. The FE-HRTEM micrographs taken for the end-products of (Co75Ti25)90B10, and (Co75Ti25)85B15 are displayed together with their NBDPs in Figure 5a,b and Figure 5c,d, respectively. The powders had typical fine morphology with random atomic distribution (Figure 5a,c), indicating the formation of short-range order structure. Furthermore, the NBDPs displayed spot-free halo-diffuse pattern, suggesting the absence of long-range order (crystalline phases) and medium-range order (metastable phase), as presented in Figure 5b,d. The FE-HRTEM micrographs taken for the end-products of (Co 75 Ti 25 ) 90 B 10 , and (Co 75 Ti 25 ) 85 B 15 are displayed together with their NBDPs in Figure 5a,b and Figure 5c,d, respectively. The powders had typical fine morphology with random atomic distribution (Figure 5a,c), indicating the formation of short-range order structure. Furthermore, the NBDPs displayed spot-free halo-diffuse pattern, suggesting the absence of long-range order (crystalline phases) and medium-range order (metastable phase), as presented in Figure 5b,d. The X-ray diffraction (XRD) patterns of MA-(Co75Ti25)100−xBx (x: 2, 10, and 15 at%) obtained after 50 h of BM time are displayed collectively in Figure 4. All the patterns exhibited only broad diffuse haloes, in scattering range of 2θ between 40-50°, with no indication of any unprocessed crystalline phases. This implying the capability of MA approach to fabricate the desired amorphous alloy powders in a wide amorphous range. The FE-HRTEM micrographs taken for the end-products of (Co75Ti25)90B10, and (Co75Ti25)85B15 are displayed together with their NBDPs in Figure 5a,b and Figure 5c,d, respectively. The powders had typical fine morphology with random atomic distribution (Figure 5a,c), indicating the formation of short-range order structure. Furthermore, the NBDPs displayed spot-free halo-diffuse pattern, suggesting the absence of long-range order (crystalline phases) and medium-range order (metastable phase), as presented in Figure 5b,d. To understand the degree of homogeneity in the chemical composition for the end-products (50 h), intensive FE-HRTEM/EDS investigations were conducted for all the samples. The FE-HRTEM micrograph of (Co 75 Ti 25 ) 75 B 25 powders, which implies the formation of amorphous structure (Figure 6a) with the absence of the crystalline phase (Figure 6b), was classified into six local zones (~5 nm for each) to investigate the local composition beyond the atomic level. The results of the EDS analysis, which is presented in Table 2 have indicated a high degree of homogeneity in the elemental composition and the absence of compositional degradation. The average chemical composition of this sample was very close to the real composition of the starting Co 56.28 Ti 18.79 B 24.93 powders ( Table 1). The as-processed powders obtained after this end-point of MA (50 h) consisted of ultrafine spherical particles with sizes ranged between 200 nm to 500 nm in diameter, as shown in Figure 6c). To understand the degree of homogeneity in the chemical composition for the end-products (50 h), intensive FE-HRTEM/EDS investigations were conducted for all the samples. The FE-HRTEM micrograph of (Co75Ti25)75B25 powders, which implies the formation of amorphous structure ( Figure  6a) with the absence of the crystalline phase (Figure 6b), was classified into six local zones (~5 nm for each) to investigate the local composition beyond the atomic level. The results of the EDS analysis, which is presented in Table 2 have indicated a high degree of homogeneity in the elemental composition and the absence of compositional degradation. The average chemical composition of this sample was very close to the real composition of the starting Co56.28Ti18.79B24.93 powders ( Table 1). The as-processed powders obtained after this end-point of MA (50 h) consisted of ultrafine spherical particles with sizes ranged between 200 nm to 500 nm in diameter, as shown in Figure 6c).  Table 2.    Table 2. The characteristics of the powders obtained after this stage of milling show a clear contrast in structure when compared with the previous HRTEM images for those samples contained a low concentration of B (≤25 at%). The higher magnification FE-HRTEM image taken for the circular zone indexed in Figure 7a indicated the existence of grey-nano-lenses (~2 to 3 nm in diameter), embedded in the fine amorphous matrix (Figure 7b). The NBDP (Figure 7c) related to the image presented in Figure 7b revealed a halo-diffuse pattern corresponding to the amorphous matrix, overlapped with sharp spots related to unprocessed nanocrystalline B, as shown in Figure 7c. The results indicated the formation of heterogeneous powders, which fluctuated in composition from spot to sport, as presented in Table 3.
Molecules 2020, 25, x FOR PEER REVIEW 7 of 15 of (Co75Ti25)70B30 powders obtained after 50 h of MA time is presented in Figure 7a. The morphological characteristics of the powders obtained after this stage of milling show a clear contrast in structure when compared with the previous HRTEM images for those samples contained a low concentration of B (≤25 at%). The higher magnification FE-HRTEM image taken for the circular zone indexed in Figure 7a indicated the existence of grey-nano-lenses (~2 to 3 nm in diameter), embedded in the fine amorphous matrix (Figure 7b). The NBDP (Figure 7c) related to the image presented in Figure 7b revealed a halo-diffuse pattern corresponding to the amorphous matrix, overlapped with sharp spots related to unprocessed nanocrystalline B, as shown in Figure 7c. The results indicated the formation of heterogeneous powders, which fluctuated in composition from spot to sport, as presented in Table  3.  Table 3.

Thermal Stability
Differential scanning calorimetry (DSC) technique was used to characterize the crystallization behavior of as-MA (Co75Ti25)100−xBx amorphous alloys, indexed by their glass transition temperature (Tg), crystallization temperature (Tx), and supercooled liquid region (∆Tx = Tx − Tg). However, differential thermal analysis (DTA) technique was employed to investigate their corresponding melting behaviors, characterized by the melting temperature (Tm), liquids temperature (Tl), and  Table 3.

Thermal Stability
Differential scanning calorimetry (DSC) technique was used to characterize the crystallization behavior of as-MA (Co 75 Ti 25 ) 100−x B x amorphous alloys, indexed by their glass transition temperature (T g ), crystallization temperature (T x ), and supercooled liquid region (∆Tx = Tx − Tg). However, differential thermal analysis (DTA) technique was employed to investigate their corresponding melting behaviors, characterized by the melting temperature (T m ), liquids temperature (T l ), and reduced glass transition temperature (T rg = T g /T l ).  Figure 10a-e. The heating rates used for conducting DSC experiments were 40 • C/min, where it was 10 • C/min for all of DTA experiments. In the DSC measurements, all the samples were isothermally heated up to the desired temperatures before cooling down to room temperature. Then, second heating runs were carried out with the same heating rates to establish baselines.   Figure  10.
The DSC thermograms presented in Figure 8a-e exhibited two opposite thermal events, taking place at different temperatures. The onset temperatures for the first events were endothermic, appearing at a low-temperature side in the range between 412 °C to 520 °C , as presented in Figure 8. These endothermic events are related to the Tg, which is a unique feature of metallic glassy alloys. At this temperature, the solid-amorphous, which is extended from room temperature to Tg, transformed The corresponding differential thermal analysis (DTA) curves for these systems are presented in Figure 10.
The DSC thermograms presented in Figure 8a-e exhibited two opposite thermal events, taking place at different temperatures. The onset temperatures for the first events were endothermic, appearing at a low-temperature side in the range between 412 • C to 520 • C, as presented in Figure 8. These endothermic events are related to the T g , which is a unique feature of metallic glassy alloys. At this temperature, the solid-amorphous, which is extended from room temperature to T g , transformed into liquid-phase (glassy phase) without any structural or compositional changes. The second events, however, were characterized by sharp pronounced exothermic peaks, taken place at a higher temperature (Figure 8), known as crystallization temperature (T x ). At these T x , the metallic glassy phase is crystallized into the long-range order phase. The thermodynamics parameters of T g and T x are used to describe the thermal stability of metallic glassy materials, where ∆T x is used to characterize their GFA. Wide ∆Tx indicates that the system has a good GFA.
Depending on B concentrations, metallic glassy (Co 75 Ti 25 ) 100−x B x powders revealed different GFA and crystallization behaviors, as presented in Figures 8 and 9. For instance, (Co 75 Ti 25 ) 95 B 5 glassy powders possessed low values of ∆T x (63 • C) and T g (412 • C), as presented in Figure 8a. This indicates a low GFA and lowers thermal stability of this system when compared with (Co 75 Ti25) 90 B 10 , revealed wider T x (102 • C), and higher T x value (511 • C), as shown in Figure 8b or Figure 9a,b. Increasing the B content (20 at% to 25 at%) led to a monotonical increase in T x (Figure 9a), to be 602 • C and 633 • C, respectively (Figure 8c,d). It should be clarified that ∆T x of (Co 75 Ti 25 ) 70 B 20 has a lower value of (82 • C) when compared with that measured value of (Co 75 Ti 25 ) 75 B 25 (121 • C), as indexed in Figure 8a or Figure 9b. into liquid-phase (glassy phase) without any structural or compositional changes. The second events, however, were characterized by sharp pronounced exothermic peaks, taken place at a higher temperature (Figure 8), known as crystallization temperature (Tx). At these Tx, the metallic glassy phase is crystallized into the long-range order phase. The thermodynamics parameters of Tg and Tx are used to describe the thermal stability of metallic glassy materials, where ∆Tx is used to characterize their GFA. Wide ∆Tx indicates that the system has a good GFA. Depending on B concentrations, metallic glassy (Co75Ti25)100−xBx powders revealed different GFA and crystallization behaviors, as presented in Figures 8 and 9. For instance, (Co75Ti25)95B5 glassy powders possessed low values of ∆Tx (63 °C ) and Tg (412 °C ), as presented in Figure 8a. This indicates a low GFA and lowers thermal stability of this system when compared with (Co75Ti25)90B10, revealed wider Tx (102 °C ), and higher Tx value (511 °C ), as shown in Figure 8b or Figure 9a,b. Increasing the B content (20 at% to 25 at%) led to a monotonical increase in Tx (Figure 9a), to be 602 °C and 633 °C , respectively (Figure 8c,d). It should be clarified that ∆Tx of (Co75Ti25)70B20 has a lower value of (82 °C ) when compared with that measured value of (Co75Ti25)75B25 (121 °C ), as indexed in Figure 8a or Figure  9b.  Based on this comparison, we claim that (Co 75 Ti 25 ) 75 B 25 metallic glassy system revealed the best GFA and possessed high thermal stability. In contrast to these full-metallic glassy systems, nanocomposite (Co 75 Ti 25 ) 70 B 30 powders, which consisted of a glassy phase coexisted with nano-spheres of unprocessed crystalline B particles (Figure 7) revealed lower ∆T x (79 • C) and T x (513 • C) values, as indexed in Figure 8a or Figure 9a,b. The degradation in these thermodynamics parameters may be related to the heterogeneity seen in the chemical composition of (Co 75 Ti 25 ) 70 B 30 , as detected by HRTEM/EDS (Figure 7b, Table 3).
In parallel to ∆T x parameter, the GFA of any metallic glassy alloys produced by rapid solidification technique is described by measuring the melting (T m ) and liquids (T l ) temperature, and calculate the T rg = T g /T l . All the DTA curves of (Co 75 Ti 25 ) 100−x B x systems presented in Figure 10a-e display a single endothermic event for each composition. This endothermic peak revealed an obvious head and tail points, which related to T m and T l temperature, respectively ( Figure 10). The onset temperature of T m and T l values measured for the (Co 75 Ti 25 ) 95 B 5 system were 1119 • C and 1184 • C (Figure 9a or Figure 10a), respectively, where the corresponding T rg of this composition was calculated and found to be 0.35 (Figure 9c). However, increasing the B concentration to 10 at% (Figure 10b) led to a significant increase in both T m (1153 • C) and T l (1224 • C), but corresponding T rg of this composition has not been improved and remained at the level of (0.35), as presented in Figure 9c. Both of T m and T l tended to increase with increasing the B content to be 1181 • C and 1268 • C ((Co 75 Ti 25 ) 80 B 20 ), and 1271 • C and 1324 • C ((Co 75 Ti 25 ) 75 B 25 ), as presented in Figure 10c,d, respectively. Notably, no improvement in T rg value could be detected, as shown in Figure 9c. The DTA trace of (Co 75 Ti 25 ) 70 B 30 system, which consisted of a glassy phase coexisted with nano-spheres of unprocessed crystalline B particles showed the highest value of T m (1319 • C) and T l (1398 • C), as indexed in Figure 10e. However, this system revealed the lowest value of T rg (0.31), as presented in Figure 9c. Based on this comparison, we claim that (Co75Ti25)75B25 metallic glassy system revealed the best GFA and possessed high thermal stability. In contrast to these full-metallic glassy systems, nanocomposite (Co75Ti25)70B30 powders, which consisted of a glassy phase coexisted with nanospheres of unprocessed crystalline B particles (Figure 7) revealed lower ∆Tx (79 °C ) and Tx (513 °C ) values, as indexed in Figure 8a or Figure 9a,b. The degradation in these thermodynamics parameters may be related to the heterogeneity seen in the chemical composition of (Co75Ti25)70B30, as detected by HRTEM/EDS (Figure 7b, Table 3).
In parallel to ∆Tx parameter, the GFA of any metallic glassy alloys produced by rapid solidification technique is described by measuring the melting (Tm) and liquids (Tl) temperature, and calculate the Trg = Tg/Tl. All the DTA curves of (Co75Ti25)100−xBx systems presented in Figure 10a-e display a single endothermic event for each composition. This endothermic peak revealed an obvious head and tail points, which related to Tm and Tl temperature, respectively ( Figure 10). The onset temperature of Tm and Tl values measured for the (Co75Ti25)95B5 system were 1119 °C and 1184 °C (Figure 9a or Figure 10a), respectively, where the corresponding Trg of this composition was calculated and found to be 0.35 (Figure 9c). However, increasing the B concentration to 10 at% ( Figure  10b) led to a significant increase in both Tm (1153 °C ) and Tl (1224 °C ), but corresponding Trg of this composition has not been improved and remained at the level of (0.35), as presented in Figure 9c. Both of Tm and Tl tended to increase with increasing the B content to be 1181 °C and 1268 °C ((Co75Ti25)80B20), and 1271 °C and 1324 °C ((Co75Ti25)75B25), as presented in Figure 10c,d, respectively. Notably, no improvement in Trg value could be detected, as shown in Figure 9c. The DTA trace of (Co75Ti25)70B30 system, which consisted of a glassy phase coexisted with nano-spheres of unprocessed crystalline B particles showed the highest value of Tm (1319 °C ) and Tl (1398 °C ), as indexed in Figure  10e. However, this system revealed the lowest value of Trg (0.31), as presented in Figure 9c.  For excellent GFA systems prepared by rapid solidification technique, T rg should be far above 0.5. The current (Co 75 Ti 25 ) 100−x B x system in all range of B content showed low T rg values, laid in the range between 0.31 to 0.41, as displayed in Figure 9c. As a result, the present metallic glassy system is a challengeable system that cannot be easily prepared by the conventional rapid solidification approach.

Magnetic Properties
The saturation magnetization (B s ) and coercive force (H) of (Co 75 Ti 25 ) 100−x B x metallic glassy systems were obtained according to the measured hysteresis loops, as depicted in Figure 11a,b. The loops of all samples exhibited typical soft magnetic behaviors, as can be realized from Figure 11a. The dependence of Bs and H on the B (x content) is presented in Figure 12. Notably, the (Co 75 Ti 25 ) 100−x B x metallic glassy system revealed a rather modest value of B s (0.48 T), with a low coercivity value of about 15 kA m −1 , as presented in Figure 11b or Figure 12. When B content increased to 5 at%, Bs was increased to 0.66 T, where the H value approached a lower value (14.6 kA m −1 ), indicating an improvement in the soft magnetic characteristics. Further improvement in the increase in Bs was attained upon increasing the concentration B alloying element to 10 at% (0.73 T), 15 at% (0.82 T), and 20 at% (0.94 T), as depicted in Figure 12.
Molecules 2020, 25, x FOR PEER REVIEW 11 of 15 For excellent GFA systems prepared by rapid solidification technique, Trg should be far above 0.5. The current (Co75Ti25)100−xBx system in all range of B content showed low Trg values, laid in the range between 0.31 to 0.41, as displayed in Figure 9c. As a result, the present metallic glassy system is a challengeable system that cannot be easily prepared by the conventional rapid solidification approach.

Magnetic Properties
The saturation magnetization (Bs) and coercive force (H) of (Co75Ti25)100−xBx metallic glassy systems were obtained according to the measured hysteresis loops, as depicted in Figure 11a,b. The loops of all samples exhibited typical soft magnetic behaviors, as can be realized from Figure 11a. The dependence of Bs and H on the B (x content) is presented in Figure 12   In parallel, the coercivity force, indexed by H was monotonically decreased with increasing the B concentration to be (11.5 kA m −1 ) for (Co 75 Ti 25 ) 80 B 20 ( Figure 12). Metallic glassy (Co 75 Ti 25 ) 75 B 25 system possessed the best soft magnetic characteristics, as indexed by its high B s value (1.1 T) and a very low H (9.3 kA m −1 ), as presented in Figure 12. In contrast, the nanocomposite (Co 75 Ti 25 ) 70 B 30 system showed a lower B s value (0.87 T) and a relatively high H value (13.8 kA m −1 ), as indexed in Figure 12. The degradation shown here in the soft magnetic properties of this system was attributed to the heterogeneity in the chemical composition, where unprocessed B nanocrystalline particles have existed (Figure 7, Table 3). In parallel, the coercivity force, indexed by H was monotonically decreased with increasing the B concentration to be (11.5 kA m −1 ) for (Co75Ti25)80B20 ( Figure 12). Metallic glassy (Co75Ti25)75B25 system possessed the best soft magnetic characteristics, as indexed by its high Bs value (1.1 T) and a very low H (9.3 kA m −1 ), as presented in Figure 12. In contrast, the nanocomposite (Co75Ti25)70B30 system showed a lower Bs value (0.87 T) and a relatively high H value (13.8 kA m −1 ), as indexed in Figure 12. The degradation shown here in the soft magnetic properties of this system was attributed to the heterogeneity in the chemical composition, where unprocessed B nanocrystalline particles have existed (Figure 7, Table 3).
In addition to the current system of this study, there are a considerable number of recent reports which show the possibility of preparing metallic glassy soft magnetic materials through mechanical alloying and other approaches. For example, Co80−xTaxSi5C15 (x = 0, 5) glassy/nanocrystalline system possesses promising soft magnetic behavior, i.e., a minimum coercivity (Hc) of 1.2 kA m −1 , which is notably lower than a minimum value obtained for Co80Si5C15 (3.3 kA m −1 ) [19]. Recently, Matsui and Omura reported that Ni-Fe-P alloy exhibited a saturation magnetic flux density of 1.1 T and a coercivity of 8.4 A/m. [20]. They pointed out the impact of good coercivity on the grain refinement by the P alloying, which can result in a lower coercivity [20].
Furthermore, the effect of MA on the magnetic properties of mechanically alloyed Ni70Co30 has been reported by N. Loudjani et al. [21]. Their results have shown that both the saturation magnetization and coercivity decreased with milling time, attaining the values of 87 emu/g and 30 Oe, respectively, after 25 h of milling [21]. More recently, Zhao et al., have reported the formation of a novel (Fe0.25Co0.25Ni0.25Cr0.125Mn0.125)100-xBx system with B concentration ranged from 9 at% to 13 at%, using melt spinning approach [22]. They pointed out that the system exhibits good soft-magnetic properties: Hc = 2.5-7.0 A m −1 , μi = 4910-15,830, and Bs = 0.40-0.48 T [22]. An interesting study of preparing bulk nano Fe-Si-B-Cu-Nb based alloys by mechanical alloying has been recently reported [23]. This alloy was consolidated into bulk nanocrystalline objects, using spark plasma sintering (SPS) technique. Based on the milling time and composition, the Ms and Hc were varied from 159.6 to 166.8 emu/g, and 87 to 109.2 Oe, respectively [23].

Feedstock Materials
High purity elemental powders of Co (150 μm, >99.9 wt%, #266647: Sigma-Aldrich, St. Louis, MO, USA), Ti (45 μm, >99.99 wt%, #366994: Sigma-Aldrich), and B (45 μm, >98 wt%, #GF47837065: Sigma-Aldrich) were used as the starting feedstock materials. The powders were handled, balanced, and then mixed inside the He-atmosphere (99.99%) glove box (UNILAB Pro Glove Box Workstation, In addition to the current system of this study, there are a considerable number of recent reports which show the possibility of preparing metallic glassy soft magnetic materials through mechanical alloying and other approaches. For example, Co 80−x Ta x Si 5 C 15 (x = 0, 5) glassy/nanocrystalline system possesses promising soft magnetic behavior, i.e., a minimum coercivity (H c ) of 1.2 kA m −1 , which is notably lower than a minimum value obtained for Co 80 Si 5 C 15 (3.3 kA m −1 ) [19]. Recently, Matsui and Omura reported that Ni-Fe-P alloy exhibited a saturation magnetic flux density of 1.1 T and a coercivity of 8.4 A/m. [20]. They pointed out the impact of good coercivity on the grain refinement by the P alloying, which can result in a lower coercivity [20].
Furthermore, the effect of MA on the magnetic properties of mechanically alloyed Ni 70 Co 30 has been reported by N. Loudjani et al. [21]. Their results have shown that both the saturation magnetization and coercivity decreased with milling time, attaining the values of 87 emu/g and 30 Oe, respectively, after 25 h of milling [21]. More recently, Zhao et al., have reported the formation of a novel (Fe 0.25 Co 0.25 Ni 0.25 Cr 0.125 Mn 0.125 ) 100-x B x system with B concentration ranged from 9 at% to 13 at%, using melt spinning approach [22]. They pointed out that the system exhibits good soft-magnetic properties: H c = 2.5-7.0 A m −1 , µ i = 4910-15,830, and B s = 0.40-0.48 T [22]. An interesting study of preparing bulk nano Fe-Si-B-Cu-Nb based alloys by mechanical alloying has been recently reported [23]. This alloy was consolidated into bulk nanocrystalline objects, using spark plasma sintering (SPS) technique. Based on the milling time and composition, the M s and H c were varied from 159.6 to 166.8 emu/g, and 87 to 109.2 Oe, respectively [23].

Preparations of Metallic Glassy Alloy Powders
The mixed powders of each composition were individually charged into tool steel vials (500 mL in volume) and well-sealed together and with 75 tool steel balls (11 mm in diameter) in the glove box under He-atmosphere (99.99%), using a ball-to-powder weight ratio was 20:1. The vials were then fixed on a high-energy ball mill (Planetary Mill PULVERISETTE 5, Fritsch, Kitzingen, Germany), where the ball milling (BM) process was carried out for 15, 30, 45, and 60 h at ambient temperature. After each milling run, a small amount (<500 mg) was discharged from the vial for different analyses. Then, the BM process was resumed under the same operating conditions.

Sample Characterizations
The crystal structures of all samples were investigated by X-ray diffraction (XRD) with CuKα radiation, using 9kW Intelligent X-ray diffraction system, provided by SmartLab-Rigaku (Tokyo, Japan). The local structure of the synthesized materials was studied by 200 kV-field emission high-resolution transmission electron microscopy/scanning transmission electron microscopy (HRTEM/STEM) supplied by JEOL-2100F (Tokyo, Japan), and equipped with Energy-dispersive X-ray spectroscopy (EDS) supplied by Oxford Instruments (Abingdon, UK).
The morphological characteristics of the milled and consolidated samples were investigated through field-emission scanning electron microscope (FE-SEM), using 15 kV-JSM-7800F, JEOL (Tokyo, Japan). The local elemental analysis was investigated by energy-dispersive X-ray spectroscopy (EDS, Oxford Instruments, Abingdon, UK) system interfaced with the FE-SEM.
Differential scanning calorimeter (DSC) and differential thermal analysis (DTA), provided by Setaram-France(Caluire, France), using a heating rate of 40 • C/min and 10 • C/min, respectively, employed to investigate the glass transition temperature, glass-forming ability, and thermal stability indexed by the supercooled liquid region and crystallization temperature of the metallic glassy samples.
The magnetization (B s ) of the as-consolidated samples was measured at room temperature, using a vibrating sample magnetometer (VSM) with a maximum applied magnetic field of 670 kA m −1 . The coercive force was measured with a B-H loop tracer.

Conclusions
Due to the absence of deep eutectic compositions in the Co-Ti binary phase diagram, it is very difficult to obtain a metallic glassy phase for this system and its based alloys, using the melting and casting and melt spinning techniques. This is in contrast to the Co-Zr binary system, which possesses several deep eutectic compositions, allowing a wide glass-formation range. The present work has been addressed in part to employ the MA approach for preparing metallic glasses of ternary (Co 75 Ti 25 ) 100−x B x systems in a very wide range of (2 ≤ x ≤ 30 at%), using high-energy ball milling technique. Based on the results of the present work, we can conclude that: (1) The GFA of (Co 75 Ti 25 ) 100−x B x was improved upon increasing the B molar fraction in the range between 2 at% to 25 at%. (2) The effect of elemental B of enhancing the GFA is attributed to the chemical bonding between the alloying elements (Co, Ti and B) in the alloy. This can be realized upon considering that electronegativity, which is a very important factor for glass formation, is directly related to the chemical bonding. Electrons of metalloid B element, which transferred to metallic Co-Ti may lead to the formation of very strong covalent bonding and hence led to improve the GFA of (Co 75 Ti 25 ) 100−x B x systems. (3) Among the as-prepared metallic glassy alloys of (Co 75 Ti 25 ) 100−x B x , (Co 75 Ti 25 ) 75 B 25 system revealed excellent thermodynamics properties, as indexed by its high GFA as characterized by the widest ∆T x (121 • C), and high thermal stability, indicated by its highest T x value (633 • C). Moreover, this glassy system revealed the highest T rg value (0.41), which indicates its high GFA when compared with other compositions in the (Co 75 Ti 25 ) 100−x B x systems. (4) The results have shown that when B content increased, the saturation magnetization was increased to reach to 1.1 T for (Co 75 Ti 25 ) 75 B 25 system, where coercivity was decreased to a very low level of 9.3 kA m −1 .