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Vol. 8. Issue 5.
Pages 4995-5003 (September - October 2019)
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Vol. 8. Issue 5.
Pages 4995-5003 (September - October 2019)
Review Article
DOI: 10.1016/j.jmrt.2019.07.048
Open Access
Magnetic properties and crystal structure of elemental cobalt powder modified by high-energy ball milling
J.A. Betancourt-Canteraa,
Corresponding author

Corresponding author.
, F. Sánchez-De Jesúsb, A.M. Bolarín-Mirób, G. Torres-Villaseñorc, L.G. Betancourt-Canterab
a CONACYT-Corporación Mexicana de Investigación en Materiales, Ciencia y Tecnología, #790, Col Saltillo 400, CP 25290 Saltillo, Coahuila, Mexico
b Área Académica de Ciencias de la Tierra y Materiales, Universidad Autónoma del Estado de Hidalgo, 42184 Hidalgo, Mexico
c Departamento de Materiales Metálicos y Cerámicos, Instituto de Investigaciones en Materiales, Universidad Nacional Autónoma de México, México, DF 04510, Mexico
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Tables (1)
Table 1. Wt.% of cobalt hcp and fcc phases as a consequence of milling time.

Elemental cobalt powder was modified by high-energy ball milling for 0–15h, and the results obtained by X-ray diffraction (XRD) revealed a phase transformation. Initially, the unmilled powder consisted of two phases: hexagonal close-packed (hcp) and face centered cubic (fcc). From a milling time of 0–5h, the fcc phase underwent a phase transformation to hcp, and increased milling times promoted a transformation back to the fcc phase. Moreover, the crystallite size decreased from 88.9nm to 16.7nm for the samples at 0 and 15h, respectively. Conversely, the microstrain increased from 0.02% to 0.489% for the same times. On the other hand, the mean particle size increased from 6.8μm to 68.6μm for the first 5h, and this effect is related to the cold welding step in the milling process and cobalt ductility. The magnetic properties of the samples were analyzed by vibrating sample magnetometry (VSM) at room temperature. The hysteresis loops exhibited an increase in saturation polarization from 1.18 to 1.33T (145.5–163Am2/kg) for the first hour. On the other hand, the coercivity field and magnetic anisotropy presented a progressive reduction as the milling time increased, and this effect was attributed to particle size growth and the main content of the fcc phase. Magnetic–thermogravimetric analysis revealed that the Curie temperature was ∼1391K, which was associated with the fcc phase. In addition, other events between 700 and 900K were identified. These events were associated with the TC of the hcp phase.

High-energy ball milling
Allotropic transformation
Magnetic properties
Magnetocrystalline anisotropy
Curie temperature
Full Text

Cobalt is an important metal in areas where high temperature properties, energy storage, process efficiency and environmental benefits are essential requirements. Cobalt is used in the fabrication of materials required for diverse applications, ranging from production of magnets, hard metals, superalloys and gas turbine components to the manufacturing of lithium-ion batteries and industrial catalysts [1]. Cobalt is also widely used in cemented carbide, stainless steel, petrochemical, automobile manufacturing, machine tools and many other industries [2]. Cobalt is one of the transition metals used in electronics and magnetic recording [3], and its alloys are used in biomedical applications given their wear resistance, corrosion resistance, and heat resistance (strong even a thigh temperatures) [4,5]. Many of the properties of the alloys arise from the crystallographic nature of cobalt (in particular, its response to stress) [6]. When cobalt is subjected to MP, it undergoes an allotropic transformation of hcp and fcc to Co-fcc as well as hcp and fcc to hcp [7–12]. These transformations are related to the milling intensity, which generates an accumulation of structural defects [13,14], and to the ball-to-powder weight ratio [15].

During the milling process (MP), a high amount of energy from the milling balls is transferred to powder particles [16]. The MP can be used to produce alloys and compounds that are difficult or impossible to obtain by conventional melting and casting techniques [17]. The MP consists of repeated welding, fracturing and rewelding of powder particles, leading to particle size variations and changes in the powder particle shape [18,19]. The powder particles are actively deformed under a high-energy impact force. As a result of this energy, strain is introduced in the lattice, and the crystal is fractured into smaller pieces [20], promoting the continuous modification of the shape and degree of disorder in the lattice [21] that creates defects, such as vacancies, dislocations, stacking faults and grain boundaries [22,23]. These structural changes influence many of the physical properties [24] in alloys obtained by MP. In general, some magnetic properties can be improved when the grain size is reduced to the nanoscale, whereas the presence of stresses and defects introduced by MP impairs magnetic properties [25]. Nanostructured materials are important because they exhibit fascinating and novel properties, such as electrical, mechanical, optical and magnetic properties, which are superior to those found in conventional bulk materials given their quantum-size effect, small-size effect as well as large number of grain boundaries [26]. The magnetic properties that undergo changes include saturation magnetization, MS, which depends on the crystal structure, chemical composition and atomic distances [27]; Curie temperature (TC) and the coercivity (HC), which depends on grain size [28]; remanent magnetization (Mr); and the evolution of the microstructure of the sample [29,30]. The reduction of the crystallite size (D) affects parameters, such as the coercive field and the magnetocrystalline anisotropy, when D is less than the ferromagnetic exchange length, and better soft magnetic properties are produced [31]. On the other hand, grain size reduction increases the magnetic coercivity because it is an extrinsic property that is often used as the most crucial single criterion for determining whether a ferromagnet is soft or hard, and the stresses and defects that are generated during the milling process together with subsequent grain size reduction results in an increase in the magnetization and coercivity [32]. In addition, magnetic anisotropy and coercivity are also partially related to stacking faults [15]. The aim of this work is to demonstrate the effect of the milling time on the structural and magnetic properties of elemental cobalt powder.

2Materials and experimental procedures

Elemental cobalt powder (<2μm with >99.9% purity from Sigma–Aldrich) was used as the precursor. Briefly, 5g of powder was introduced with 6 hardened steel balls with a diameter of 12.7mm into a steel vial. The ball-to-powder weight ratio was 10:1. The mechanical milling process was conducted at room temperature in an argon atmosphere using a shaker mixer/mill machine.

The identification of the crystal structure and phase transformations of the milled samples as a function of the milling time was performed by X-ray diffraction (XRD) using a Siemens D5000 diffractometer. The powder diffraction patterns were collected and ranged from 30° to 120° of 2θ with a step size of 0.02 and Co Kα (λ=1.7902Å) radiation.

The crystallite size (D) and internal microstrain (ɛ) in the milled samples were calculated from the XRD line broadening using the Williamson–Hall method:

where B is the full-width at half maximum of the diffraction peak, θ is the Bragg angle, ɛ is the internal microstrain, λ is the wavelength of the X-ray, D is the crystallite size and K is a constant. In this method, Bcosθ is plotted against 2sinθ, ɛ is the slope and Kλ/D is the intercept [33].

Morphological characterization was performed using a Leica Stereoscan 440 electron microscope operated at 20kV. The magnetic properties of the samples, specifically the saturation magnetization (MS) and coercivity (HC), were measured using a LDJ9600 vibrating sample magnetometer (VSM) with a maximum applied field of 15kOe.

The magnetic anisotropy constant was determined for each sample using the law of approach to saturation. According to this law, the magnetization, M, can be expressed as a function of the magnetic field, H, in the saturation region, as follows:

where MS is the saturation magnetization, the term b/H2 is caused by uniform magnetocrystalline anisotropy, and a/H is attributed to the existence of structural defects and non-magnetic inclusions. Parameter b can be expressed in terms of the first order uniaxial magnetic anisotropy, K1, as follows:

For the case of a ferromagnet with a cubic crystal structure, the coefficient b is given as follows:

where K1 is the first-order cubic anisotropy constant [34].

The Curie temperature was determined using the magnetic thermogravimetric technique on a TGA/SDTA 851e Mettler-Toledo analyzer. The following test conditions were employed: a temperature range from 600K to 1500K under a pure argon atmosphere and a magnetic field of 0.36T (45Am2/kg).

3Results and discussion3.1Structural analysis

Fig. 1 shows the XRD patterns of cobalt samples submitted to different milling times. As noted, the un-milled cobalt powder has two allotropic phases that coexist at room temperature: hexagonal close packed (Co-hcp) and face centered cubic (Co-fcc). After 1h of milling time, the XRD shows that the main peak (101) of the hcp phase decreased drastically, whereas the unique peak (200) of the fcc phase, which does not overlap with any hcp peak, disappeared, indicating that is possible obtain the hcp phase after the first hour of milling. This transformation occurs because Co-hcp is metastable at room temperature and becomes unstable when an external mechanical or thermal energy is introduced [35]. With increased milling time, the accumulation of structural defects increases, and the absorbed energy in the materials is also increased. Therefore, the phase formation of cobalt is determined by the accumulation of structural defects, especially stacking faults, and not by the local temperature increase, as was remarked by Huang et al. [13]. The fcc phase is first transformed to the hcp phase and is then converted back to the fcc phase at longer milling times [36]. The refinement of the microstructure and the presence of microstrains caused by severe plastic deformation create a large number of defects in the fcc lattice, which cause modifications to the cell parameters and facilitate hcp phase formation [37]. However, this transformation is not complete, as reported in previous works [13–15], because the fcc phase is present at a lower amount as noted in the Rietveld refinements in Fig. 3a. Therefore, the structure achieved after 1h of milling time is composed of a mixture of hcp and fcc. The same behavior is observed for the sample at 3h given that the (100) and (101) peaks of the hcp phase are exclusively reduced. However, after 5, 7, 9, 12 and 15h, significant decreases in hcp phase peaks are observed in addition to an increased intensity of the plane (111) of the fcc phase and a small increase in the peak (200). For the samples obtained after 5h of milling time, a mixture of the same mixture of allotropic phases is present in different amounts.

Fig. 1.

XRD patterns of the unmilled cobalt and the cobalt samples milled for different times.


Fig. 2 shows the Rietveld refinement of the XR diffraction patterns of un-milled pure cobalt powder. The coexistence of two allotropic phases, Co-hcp and Co-fcc, were observed and quantified as 78±0.45wt.% of the hexagonal phase (hcp) and 22±0.55wt.% of the cubic phase (fcc) with a crystallite size of 88.9±2nm. The cell parameters were determined as a=3.5442Å for fcc and a=2.5074Å and c=4.0699Å for the hcp phase.

Fig. 2.

Rietveld refinement of unmilled cobalt powder.


The Rietveld refinement in Fig. 3 and Table 1 present the samples at 1, 5, 9 and 15h on a square-root scale for the intensity to clarify the low-intensity peaks from the background. The refinement shows the evolution of the hcp and fcc phases the milling time, and this reduction is corroborated by the table data where the wt.% for both phases is shown. In the first hour of milling, the composition obtained was 85.68wt.% hcp and 14.32wt.% fcc. Increasing the milling time, the hexagonal peaks diminish progressively as noted in Table 1; therefore, the cubic peaks exhibit increases in intensity and weight percentage. The same behavior is observed for the following milling times. Finally, when the cobalt powder is milled for 15h, a phase mixture composed of 61.77wt.% fcc and 38.23wt.% hcp is produced.

Fig. 3.

Rietveld refinement of the cobalt samples obtained at 1, 5, 9 and 15h of milling time.

Table 1.

Wt.% of cobalt hcp and fcc phases as a consequence of milling time.

Milling time  wt.% of hcp  wt.% of fcc 
185.68  14.32 
579.21  20.79 
968.37  31.63 
1561.77  38.23 

As a result of the plastic deformation, a diminution of the crystallite size and an increase in microstrain occurred. Fig. 4 shows the effect of the milling time on the crystallite size and microstrain. An increase in the milling time leads to a rapid decrease in the crystallite size from 88.95nm for the un-milled powder to 20nm in samples milled for more than 9h. The reduction in crystallite size is accompanied by an increase in the microstrain as a function of the milling time. Because of the structural defects introduced during the milling time, the microstrain increases rapidly from 0.02% to 0.22% in the first hour of milling and subsequently reaches a value of 0.489% for the powder milled for 15h.

Fig. 4.

Crystallite size and microstrain of the cobalt powder as a function of the milling time.


Fig. 5 shows the SEM micrograph of the Co powder for various milling times, demonstrating the morphology of the un-milled and milled powders after 5, 9, and 15h. The unmilled powder shown in Fig. 5a is between 1 and 2μm in width and 4 and 10μm in length with a ligament-like morphology. At the initial stage of milling, the particles tend to be agglomerated; after 5h of milling, the SEM micrographs show a large increase in particle size due to the formation of agglomerates with irregular shapes as a result of the cold welding caused by the MA (Fig. 5b). The increase is particle size is attributed to the ductile features of cobalt, which has a tendency to form agglomerates. When the milling time is increased to 9h (Fig. 5c), the agglomerates show a slight reduction with respect to the sample at 5h due to particle deformation, which induces strain hardening into the particles. Finally, at 15h (Fig. 5d), no further changes are observed with the exception of a decrease in particle size.

Fig. 5.

SEM micrographs of Co powders for different milling times: (a) 0, (b) 5, (c) 9 and (d) 15h.


The particle size and the particle-size distribution change as a result of the coalescence and fracture events occurring during the MP [38]. To corroborate the increases in particle size due to milling, Fig. 6 shows the plot of the cumulative volume % versus median diameter (Dm50) for samples obtained at different milling times. The un-milled powder has a median particle size of 6.8μm. After 1h of milling, the particle size increases to 38.24μm. The malleable samples of cobalt are deformed into long lamellae by the impact of the balls, which is followed by an increase in the number of lamellae due to cold welding. At 5h of milling time, the particle size increases to 68.6μm. At the subsequent milling times, there is a reduction in particle size, reaching a value of 43.7μm at 15h. The particle size variation as a function of the milling time can be explained by considering the stages of the milling process (flattened, cold welded, fractured and re-welded).

Fig. 6.

Variation of the particle size of the cobalt powder obtained at different milling times.


Fig. 7 shows the magnetic hysteresis loops (JH) of the unmilled powder and cobalt milled for 7h. The unmilled cobalt has a magnetic saturation polarization of 1.18vT (145.5Am2/kg) and a coercivity of 22.7kA/m (285.6Oe), whereas the cobalt milled for 7h presented a JS value of 1.33T (163Am2/kg) and a HC value of 9.17kA/m (115.3Oe). A reduction in the coercive field and an increase in the magnetic saturation polarization were observed as a function of the milling time. This effect can be appreciated more clearly in Fig. 8 where saturation polarization (Js) and coercivity (HC) are presented as functions of milling time. After 1h of milling, JS increases, after that the polarization value remains constant, independently of the milling time. This behavior is attributed to the reduction in the magnetocrystalline anisotropy due to the diminution in crystallite size, which leads to an easier rotation of the domain walls [28], reducing the directional dependence of a material's magnetic properties. In addition, according to Shokrollahi [25], a reduction in crystallite size promotes the averaging effect of magnetization over randomly oriented nano-size grains, which explains the reduction in the magnetocrystalline anisotropy. Furthermore, the reduction in the crystallite size reduces the HC from 22.7 to 9.17kA/m (285.6 to 115.3Oe) for the samples milled for 0 and 7h, respectively. However, for higher milling times, the HC increases.

Fig. 7.

Magnetic hysteresis loops (JH) for the samples milled for 0 and 7h at room temperature.

Fig. 8.

Saturation polarization (Js) and coercivity (Hc) as functions of the milling time.


In general, a combination of different factors affects the coercivity during milling, such as internal microstrain, impurities, pores and defects. In the studied material, the reduction in the coercive field is related to the decrease in magnetocrystalline anisotropy, which is promoted by the reduction of the crystalline size, and microstrain, which is promoted by the mechanical energy generated by the high-energy ball milling.

To understand this phenomenon produced by the high-energy ball milling over the Co powder, the magnetocrystalline anisotropy constant K1 is presented as a function of the milling time in Fig. 9. K1 tends to be reduced as the milling time increases. First, the un-milled sample presents a K1 value of 8.36×105J/m3. At low milling times, i.e., 1h, the anisotropy decreases significantly up to 6.22×105J/m3. For the samples milled at 3, 5, 7 and 9h, an oscillatory behavior between values of 6.58 and 6.06×105J/m3 is observed. Consequently, at 12h, the anisotropy of the sample is reduced to 5.5×105J/m3. Finally, at 15h, there is a slight increase in the K1. However, this value is very small compared with the un-milled sample. This general tendency for reduction is attributed to the effect of the reduction in crystallite size, which reduces the magnetocrystalline anisotropy caused by the average effect of magnetization over the randomly oriented nano-sized grains and is consistent with other studies [39].

Fig. 9.

Magnetocrystalline anisotropy constant, K1, as a function of the milling time.


Cobalt and its alloys have many applications given their soft magnetic properties, such as high saturation magnetization, high permeability and high Curie temperature [28]. Cobalt has two Curie temperatures due to it is allotropic characteristics. The Curie temperature associated with the fcc ferromagnetic phase is between 1404 and 1367K [40,41]. Moreover, the hcp phase is also ferromagnetic and has a TC; however, its value is not clear because direct measurements are difficult [40]. The Curie temperature of the hexagonal phase has been estimated by extrapolation as 1130K or between 1070K and 1150K [42]. Generally, when discussing the TC of cobalt-based alloys and elemental Co powder, the teperature taken into account is related to the fcc phase. Fig. 10 presents the thermal-magnetic characterization of the samples at 0, 3, 7 and 15h of milling. Fig. 10a shows a slope shift in the base line between 700 and 900K. Based on magnetic properties, this change may be associated with the TC of the hexagonal phase.

Fig. 10.

Curie temperature for cobalt at different milling times.


Fig. 10b shows the samples from 1000 to 1450K. A thermal-magnetic event is noted approximately 1391K for the unmilled cobalt, which is consistent with the results of previous works [32]. The same behavior is observed as milling time increases; only small decreases occurred for the samples milled for 15h with a TC of 1382K for cobalt. A thermal-magnetic event is noted at approximately 1390K, which increases in intensity as a function of milling time and is related to an increase in the quantity of the fcc phase. Conversely, the thermal-magnetic event associated with the hcp phase decreases in intensity based on the results obtained in the Rietveld refinements.


In summary, we analyzed the influence of the milling time on structural and magnetic properties of pure cobalt powder, which is composed of a mixture of two allotropic structures: hcp (85.68wt.%) and fcc (14.32wt.%). The cobalt powder undergoes a phase transformation as consequence of the energy received from the high-energy ball milling, which increases the amount the internal defects and microstrain, promoting the structural transformation. A rapid decrease in the crystallite size and increase in the microstrain were observed after 1h of milling and attributed to the generation of internal defects promoted by mechanical milling. The reduction in the crystallite size affects parameters, such as the coercive field and magnetocrystalline anisotropy, which decrease as a function of the milling time. This feature is attributed to intrinsic factors, such as allotropic transformation (hcpfcc), and extrinsic factors, such as reduced crystalline size and increased microstrain. In addition, thermal analysis revealed two thermal events at different temperatures ascribed to the Curie temperature of the Co-fcc (∼1391K) and Co-hcp (∼800K) phases. The TC of the fcc phase is consistent with the measurements in the literature; however, with respect to the TC of the hcp phase, no reference values were experimentally obtained. However, two events at different temperatures were related to the Curie temperature of the fcc (approximately 1391K) and hcp (approximately 800K) phases. The TC of the fcc phase matches the measurements in the literature; however, no reference values were experimentally obtained with respect to the TC of the hcp phase.


The authors acknowledge research support from CONACyT from México and Cátedras CONACyT with the following ID number: 2362.

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