Structural and magnetic properties of Co50Ni50 powder mixtures

Journal of Magnetism and Magnetic Materials 323 (2011) 3063–3070

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Structural and magnetic properties of Co50Ni50 powder mixtures ˜ ol b N. Loudjani a,n, N. Bensebaa a, L. Dekhil a, S. Alleg a, J.J. Sun a b

Laboratoire de Magne´tisme et Spectroscopie des solides, De´partement de Physique, Faculte´ des Sciences, Universite´ Badji Mokhtar, B.P. 12, 23000 Annaba, Algeria Dep. de Fisica, Universitat de Girona, Campus Montilivi, Girona 17071, Spain

a r t i c l e i n f o

abstract

Article history: Received 27 January 2011 Received in revised form 19 June 2011 Available online 28 June 2011

In the present work, morphological, structural, thermal and magnetic properties of nanocrystalline Co50Ni50 alloy prepared by high energy planetary ball milling have been studied by means of scanning electron microscopy, X-ray diffraction, and differential scanning calorimetry. The coercivity and the saturation magnetization of alloyed powders were measured at room temperature by a vibration sample magnetization. Morphological observations indicated a narrow distribution in the particle and homogeneous shape form with mean average particle size around 130 mm2. The results show that an allotropic Co transformation hcp-fcc occurs within the three first hours of milling and contrary to what expected, the Rietveld refinement method reveals the formation of two fcc solid solutions (SS): fcc Co(Ni) and Ni(Co) beside a small amount of the undissolved Co hcp. Thermal measurement, as a function of milling time was carried out to confirm the existence of the hcp phase and to estimate its amount. Magnetic measurement indicated that the 48 h milled powders with a steady state particles size have the highest saturation (105.3 emu/g) and the lowest coercivity (34.5 Oe). & 2011 Elsevier B.V. All rights reserved.

Keywords: Nanocrystalline material Mechanical alloying Co50Ni50 alloy Microstructure Calorimetry Magnetic measurement

1. Introduction Amongst the non-equilibrium materials (amorphous, quasicrystals and solid solution), nanocrystalline materials, which are defined as materials with new functionalities, show great promise for use in industrial applications because of their unique and special chemical, physical and mechanical properties. These properties are now recognized to be dependent on the chemical composition phase, structure, size and shape of individual crystalline particles and particles boundaries. Nanocrystalline materials have been synthesized by number of techniques starting from the vapor phase, liquid and solid state such as mechanical alloying (MA) [1]. The main advantage of using this technique is due to its ability to produce large quantities (from a few hundreds of mg up to tons) of material in the solid state with the same physical properties [2]. In addition, this route allows producing materials usually prepared by means of high temperature synthesis (as ceramic method). Consequently, MA has attracted much attention because of its promising results and its wide application and was thus developed as a technique for synthesis of a variety of alloy phase including solid solutions, quasicrystalline and crystalline intermetallics and amorphous phases [3–5]. The milling process leads to considerable size refinement of the powders to nanometer scales due to a high level of plastic deformation [1]. However, continued decreasing of particles sizes causes the change in magnetic properties from hard to soft; this is

explained by the called random anisotropy model (RAM) [6]. When the particles size is smaller than the magnetic exchange length, coercivity increases as a D6 (up to DCritical ¼20 nm). Then the coercivity goes through a maximum and decreases when the particles size is larger than the magnetic exchange length, as 1/D. The macroscopic magnetic properties of the Co–Ni alloys were studied by various methods. These properties were investigated for Co–Ni prepared by sol–gel routs [7,8], microemulsion method [9,10], polyol process [11–13], Co–Ni thick film [14,15] and by ball milling process [16,17]. These studies showed that Co–Ni has excellent magnetic properties such as low values of both squareness ratio Mr/Ms and the coercivity, Hc, and high Curie temperature, Tc. All this indicates that these alloys are well suited for high temperatures applications. In the present work, nanocrystalline Co50Ni50 powders were produced by the mechanical alloying in a high-energy planetary ball mill. The effects of the milling time on the microstructure and magnetic properties of the mechanically alloyed powder mixtures by means of the scanning electron microscopy (SEM), X-ray diffraction (XRD), differential scanning calorimetry (DSC) and magnetic measurements were carried out at room temperature using a vibrating sample magnetometer (VSM) with a maximum magnetic field of 10 kOe.

2. Experimental procedures n

Corresponding author. Tel.: þ21395543591. E-mail address: [email protected] (N. Loudjani).

0304-8853/$ - see front matter & 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.jmmm.2011.06.059

Pure elemental powders of cobalt (2–5 mm, 99.0%) and nickel (6–10 mm, 99.7%) were mixed to give a nominal composition of

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Co50Ni50 (wt%), using a planetary high energy ball mill (Fritsch P7). The milling process was performed at room temperature using hardened steel vials and balls (5 balls, diameter 12 mm). The ball to powder weight ratio was set a 35:4. The rotation speed was 400 rpm. The hardened steel vials and vials were sealed under argon atmosphere in glove box to prevent from oxidation. In order to avoid the local temperature rise inside the vials, milling was interrupted each 1/2 h for 1/4 h. Powder morphology and particle size evolution during the milling process were followed by scanning electron microscopy (SEM). The particle size distributions were made with image analyzer soft ware (ImageJ, WRNIH, USA). Structural and microstructural changes in the milled powders were characterized by XRD measurements on Brucker D8 Advance diffractometer in a (y–2y) geometry using Cu-Ka radiation (l ¼0.1540566 nm). Phase transformation and microstructural parameters (lattice parameters, crystallite size, microstrains and phase percentage) were obtained from the full pattern XRD Rietveld method [18]. Thermally activated phase evolution in the ball milled was examined by DSC were carried out using a DSC822 apparatus of Mettler Toledo at a heating rate of 10 1C/min under argon atmosphere, in the temperature range 35–700 1C. The second run was used as baseline to estimate the enthalpy output of the transformations from the integrated area of the peaks. The magnetic measurement was performed in a vibratory sample magnetometer (VSM) at 300 K in an external field up to 10 KOe. The magnetic characterization of the ball milled powders such as the saturation magnetization Ms and the coercivity Hc in the nanostructured Co50Ni50 alloys are discussed as a function of milling time.

3. Results and discussion 3.1. Morphology and particle size analysis Fig. 1 shows the change of the morphological shape and the corresponding particles size distribution of the powders obtained after 3, 12 and 24 h milling times. The particles size distributions were obtained from dark field SEM images by measuring the size of more 200 diffracting particles. Particles size distribution parameters (mean, range) were obtained for each of the three times. Knowing that, the difference between the mean particle size and the mean crystallite size established that the particles are made up of smaller primary particles. The unmilled Ni powder particles have roughly spherical and rounded shapes, whereas the Co particles have thin plate like morphology. It is known that cold welding and fracture are the two essential processes involved in the milling process [1,19]. It can be seen that the morphology changes greatly after few hours of milling (3 h, Fig. 1a). The particles are flattened by the plastic deformation caused by the progressive forces induced by the contacts between ball–powder–ball and/or ball–powder–vial. Thus leads to increasing the size of particles beside some small particles produced. These two types of particles are commonly seen in the early stage of milling of ductile materials, and are the product of the continuous cycle of plastic deformation cold welding and fracturing. As confirmed by regarding the particles size distribution histogram, a large distribution of particles sizes (50–2500) mm2. When the milling time reaches 12 h (Fig. 1b), fracturing becomes the most relevant event due to the inherently brittle nature of the Co particles, hence reducing the overall powder size and narrowing the particle size distribution (25–1000) mm2. A slight growth of particles is observed in Fig. 1c for milling up to 24 h, may be explained in terms of the fact that the fragments became welded together and the fine particles embedded in the formed layered matrix. As a result, narrower distribution of particle sizes as a balance is reached between the fracture and

welding processes. The histogram of the particle sizes reveals a bimodal distribution of particles sizes ranging in values from (75 to 1000) mm2 and (1000 to 3000) mm2. 3.2. Structural analysis X-ray diffraction patterns were taken for the elemental powders mixture before and after ball milling in order to follow the mixing process of Co and Ni into the following composition Co50Ni50 (wt%) and the structural changes that occur during mechanical alloying (Fig. 2). With increasing milling time, both the broadening of the diffraction peaks and the continuous decrease in the peak heights originate from a reduction in crystalline size, faulting and microstrains within the diffracting domains. Fig. 3 displays the XRD pattern of Co50Ni50 mixture before milling (0 h). One can observe the main diffraction peaks of the elemental fcc Ni, hcp Co and fcc Co powders. The three main peaks could be indexed to three plans /1 1 1S, /2 0 0S and /2 2 0S of fcc phase and /1 0 0S, /0 0 2S and /1 0 1S of hcp phase. It is important to emphasize the presence of additional peaks assigned to the oxide CoO and/or NiO. One does also mention that the fcc Ni and fcc Co have nearly lattice parameters which are of about 0.3523 and 0.3544 nm, respectively, preventing their resolution on the X-ray pattern. In fact, solid solubility levels have been generally determined from changes in the lattice parameter values deduced from peak positions in the XRD patterns [1]. Hume-Rothery rules applied to Co and Ni clearly demonstrate how these metallic elements form an fcc solid solution in a wide range of compositions, preventing them to form intermetallic compounds, particularly very similar electronegativity values for Co and Ni, 1.88 and 1.91, respectively [20]. The XRD patterns of the Co50Ni50 powder mixtures milled for 3 h (Fig. 4) clearly shows the decreasing of the hcp Co peaks, it seems to be absent, but the refinement of the XRD patterns using the Rietveld method reveals the existence of this phase with a small proportion (4%). This is assigned to the allotropic transformation of Co from hcp to fcc form and to the atomic diffusion of Co into the Ni matrix. The absence of solute peaks in the X-ray diffraction patterns is frequently taken as a proof of complete dissolution, and this has been interpreted as evidence for enhanced solid solubility limits. But, it has been revealed that the solid solubility limits cannot be exactly determined only by noting the absence of solute peaks in the X-ray diffraction patterns [1]. In fact, the allotropic transformation of Co from hcp to fcc form has been observed in several studies by ball milling [16,21,22]. Generally, the allotropic transformation of Co during the ball milling is strongly dependent on the milling intensity and the milling time [23,24].The progressive mutual dissolution of the elemental Co and Ni powders gives rise to the formation of two fcc solid solution Ni(Co) and Co(Ni). Indeed, it is well established that the diffusivity of Co into Ni is larger than that of Ni in Co [25]. Fig. 4 shows that the characteristic peaks of the two SS did not disappear, even after ball milling for 48 (Fig. 5). Though the characteristic peaks for the fcc Ni(Co) and fcc Ni(Co) phases became broadened and the intensity of their diffraction pattern decreased. This was due to the simultaneous decrement of particles size and increment of internal strain during the highenergy ball milling process. Except for 24 h of milling (Fig. 2), the peaks of the fcc SS increase in intensity. This suggests the reverse transformation of the hcp-Co to the fcc-Co. The hcp-fcc transformation was observed for pure ball-milled Co [21–24]. Cardellini and Mazzone [23] explain this allotropic transformation by the contamination of the powder by the iron coming from milling tools, which stabilizes the fcc phase. However, the others [21,22] relate this transformation to the amount of stacking faults accumulated in the hcp lattice during milling.

Number percentage of particules (%)

Number percentage of particules (%)

Number percentage of particules (%)

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16 14 12 10 8 6 4 2 0 0

500

1000

1500 2000 Particle size (µm2)

0

500

1000 1500 2000 Particle size (µm2)

0

500

2500

3000

30 25 20 15 10 5 0 2500

3000

16 14 12 10 8 6 4 2 0 1000 1500 2000 2500 3000 Particle size (µm2)

Fig. 1. Morphology and particles size distribution of Co50Ni50 milled for: 3 h (a) 12 h, (b) 24 h (c).

Fig. 6 shows that for Co(Ni) and Ni(Co) SS formed in Co50Ni50 powder mixtures, the lattice parameter increases with increasing milling time due to the heavily cold worked and plastically deformed state of powders [1] and to the particles expansion linked to the increase of dislocations density. Thus confirmed the formation of composition alloys of nickel and cobalt with two fcc solids solutions: Co(Ni) and Ni(Co). Generally, in nanostructured materials, the lattice expansion (0.2% and 0.4%, respectively) may be caused by defects introduced in interfaces. The less dense structure of interfaces can result in some negative pressure on the interfaces [26] and this can lead to an increase in lattice parameter. On the other hand, the lattice parameter decreases for Co(Ni) alloys after 3 h of milling, then increase again. Whereas the decrease of this parameter may be arise from the particles compression due to the presence of compressive fields within the non-equilibrium particles boundaries inside of crystallites, and as a result cause a decrease in the lattice parameter [27]. This latter may be as well due to allotropic phase transformation of

Co (hcp2fcc) and/or to the triple defect disorder [1,28]. This type of disorder has been observed in the Co alloys [28]. The variation of the crystallite size, /LS, and the lattice microstrains, /s2S1/2, as a function of milling time, for both solid solutions are plotted in Figs. 7 and 8, respectively. One can see an important decrease of the crystallite size and an increase of the microstrains during the first hours of milling. Within the first 3 h of milling, MA leads to a rapid decrease in the crystallite size (less than 32 nm). Then slightly increases (MA 12 h) and remains constant at a steady state value up to 24 h. Not that the /2 0 0S dense line of fcc structure has the more important decrease in the crystallite size ,10 and 12 nm for Co(Ni) and Ni(Co), respectively. The increase in the crystallite size at 12 h of milling with milling time is attributed to dislocation generation caused by severe plastic deformation [28]. Simultaneously, the lattice microstrains are found to exhibit significantly anisotropy behavior. Beginning by Co(Ni): one not that /s2111S1/2, /s2200S1/2 and /s2220S1/2 increases from about

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0 to 0.90%, 0.86% and 0.85%, respectively, after 48 h of milling, followed by Ni(Co), the microstrains decreases from about 0 to 0.29%, 0.34% and 0.31%, the decrease of the lattice microstrains can be due to severe plastic deformation which linked to the increasing dislocations density. The small decrease in microstrains associated with 24 to 48 h of milling for Co(Ni) SS is attributed to the temperature rise leading to the annihilation of some dislocations by their rearrangement [27]. According to result presidents, we summarize that Co(Ni) takes the smallest values of size and the greatest values of deformation because of the brittle nature of Co particles comparable to the Ni particles [12,13].

3.3. Thermal stability Fig. 9 shows the DSC scans of the different milled powders as a function of milling time, taken at a heating rate of 10 1C/min for Co50Ni50 powders mixture. No difference is observed between samples milled at different milling time.

Intensity (Arbitrary units)

44.5

44.48

44.46

44.56

44.7

51.68

51.72

51.72

76.12

92.36

76.24

92.52

48 h

24 h 3.4. Magnetic properties

92.58

76.26

51.78

76.22

92.66

52.04

76.52

93.02

12 h

3h

0h 40

These curves show a broad exothermic hump composed of three main peaks located at about 300, 450 and 650 1C. The first sharp exothermic peak, around 300 1C, may be attributed to the presence of cobalt oxide [29], and/or to the structural relaxation of the system. The second exothermic feature in the temperature range 450–500 1C can be attributed to the Co allotropic transformation from hcp to fcc [12,30]. We can determine the volume fraction for the transformed Co hcp phase from the activation enthalpy DH (0.04 J/g). The amount of the transformed phase, which is about 4%, is in good agreement with that found by the Rietveld analysis. Even after long milling time, hcp-Co still exist indicating that this quantity has not undergone neither an allotropic transformation nor not took part in the formation of the solid solutions. A similar broad peak was already observed [31] for the ball milled Cu–Ni–Co–Fe quaternary alloys. However, the analyze by means of XRD reveals fcc Co crystallite structure formed after heating [31]. The prominent peak appearing in the temperature range between 550 and 700 1C is not a result of a structural change, but originates from the relaxation of the heavily deformed powders, as proved by the XRD investigations of the powders after milling times (/s2S1/2  0.9% in Co(Ni) and 0.30% in Ni(Co)). Moreover, the relaxation under heat treatment also occurs in a large temperature interval and is often overlapped by growth or crystallization process of pre-existing nuclei [32]. This process becomes more energetic as the milling time is increased.

100

60 80 2 Theta [degrees]

Fig. 2. XRD patterns of the Co50Ni50 powders milled for various times.

A considerable change in the magnetic behavior have been observed for the nanocrystalline ferromagnetic materials compared to the conventional materials, which are strongly dependent on many parameters such as composition, crystallographic texture, internal stress, particles shape anisotropy and particles size distribution. Fig. 10 illustrates the milling time dependence of the hysteresis loops, at 300 K, of the milled Co50Ni50 powder. These sigmoidal hysteresis cycles are usually observed in nanostructured samples with small magnetic domains. This is due to the presence of structural distortions inside particles. The small hysteresis losses are properties generally desired in soft magnetic materials.

〈 002〉 H + 〈111〉 C

〈200〉 C 〈110〉 H + 〈220〉 C

〈110〉 H + 〈220〉 C 〈100〉 H

〈201〉 H + 〈311〉 C

〈101〉 H 〈102〉 H

〈103〉 H

Fcc-Ni CoO/NiO Fcc-Co Hcp-Co

Fig. 3. Rietveld refinement of the XRD pattern of the unmilled Co50Ni50 powder mixtures.

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Fig. 4. Rietveld refinement of the XRD pattern of the ball milled Co50Ni50 powder mixtures for 3 h.

48 h

Fcc-Co(Ni) Fcc-Ni(Co) CoO/NiO Hcp-Co

Fig. 5. Rietveld refinement of the XRD pattern of the ball milled Co50Ni50 powder mixtures for 48 h.

The coercivity value of the fabricated Co50Ni50 powders was measured at various milling times. Fig. 11 shows the changes in the coercivity force as a function of ball milling time. The coercivity curve of the Co50Ni50 alloy has been divided into three regions I, II and III for more accurate discussion. In the first region (between 0 and 3 h of milling), Hc decreases rapidly from 84.8 to 34.5 Oe. The decrease of Hc, during the first hours of milling, is related to the effect of the small crystallite size which overcomes the effect of microstrains. In other words, when the crystallite size is much smaller than the ferromagnetic exchange length, there are better soft magnetic properties. The increase in powder particle size, observed by SEM as milling time increases, may also be responsible for the decrease of Hc. Thus, as the powder particle increases, being multi-domain systems, the movement of the magnetic domain walls would become easier and the material

would become magnetically softer. Another probable reason was pointed by Chen [33] who found that the dislocation density is the main parameter affecting the coercivity. The second region corresponds to the increase of Hc to 45 Oe when the milling time increases to 12 h suggesting the effect of microstrains which can be the prevailing factor in the coercivity rather than the usual effect of crystallite size reduction. In the third region, the decrease of Hc to 33 Oe, followed by a steady state on prolonged milling time, can be related to the effect of the very small crystallite size. The values of Hc obtained after 48 h for milling is comparable to the value 34 Oe found in the submicrometer size crystalline Co50Ni50 powders prepared by polyol process [12]. As shown in the insert of Fig. 11, the evolution of the coercivity Hc and the crystallite size /LS as a function of the milling time relating to Co–Ni alloy. The curves obtained show

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the magnetic vector. Moreover, the higher value of Ms can be to the increase substitution of cobalt atoms in the predominate fcc Ni(Co) solid solution leading to an increase of the lattice parameter with increasing milling time. This decrease of Ms may also explained by the reduction of the crystallite size as suggested by previous work [34,35]. Amils et al. [36] have displayed that the ball milling process produces a high density of defects, particularly of antisite type, which causes a 0.8% lattice parameter expansion in Fe–Al alloys. Further milling to 12 h leads to the

a regular and a similar diminution and augmentation of the two parameters when the milling time is increased. After 24 h, Hc and /LS remained nearly unchanged. The variation of saturation magnetization, Ms, as a function of milling time of Co50Ni50 powders mixtures displays an antagonist behavior (Fig. 12). In effect, for Co50Ni50 alloy, the increase of Ms to 107.3 emu/g, during the early stage of milling, can be related to the reduction in magneto-crystalline anisotropy due to the crystallite size refinement, which leads to an easier rotation of

4.0

0.3552

633.25

3.5 Exoth. heat flow (mW)

Lattice parameters (nm)

0.3548 0.3544 0.3540 0.3536 Co (Ni) Ni (Co)

0.3532

3.0 2.5

463.18 575.05616.99 3h

2.0

458.99

1.5

305.93 290.21 259.83

1.0

48h

446.77

24h

0.5

0.3528

0.0 6

0

12

18 24 30 Milling time (h)

36

48

42

54

0

Fig. 6. Lattice parameter as a function of milling time for fcc Co(Ni) and Ni(Co).

200

Co (Ni)

111 200 220

Crystallite size (nm)

80 70

50 40 30

48

111 200 220

100 80 60 20

42

700

120

40

18 24 30 36 Milling time (h)

600

140

10 12

Ni (Co)

160

20

6

300 400 500 Tempearature (°C)

180

60

0

200

54

0

6

12

18 24 30 36 Milling time (h)

42

48

54

48

54

Fig. 7. Milling time dependence of the mean crystallite size, /LS, for different [h k l] directions.

0.35

1.0

0.30 Microstrains (%)

Microstrains (%)

0.8 0.6

111 200 220

0.4

0.25 111 200 220

0.20 0.15 0.10

0.2 Co (Ni)

0.05

0.0

Ni (Co)

0.00 0

6

12

18 24 30 36 Milling time (h)

42

48

54

800

Fig. 9. DSC scans of the ball milled Co50Ni50 powders mixtures for 3, 24 and 48 h.

Crystallite size (nm)

90

100

0

6

12

18 24 30 36 Milling time (h)

Fig. 8. Milling time dependence of the mean microstrains /s2S1/2, for different [h k l] directions.

42

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3h

0h 100

M (emu/g)

50

-100

100

-5000

-50

5000

0 H (Oe)

-100

0 100 200 H (Oe)

10000

-10000

12 h

100

-5000

5000

0 H (Oe)

10000

24 h

M (emu/g)

50

0 M (emu/g)

M (emu/g)

50

-50

20 15 12 h 10 5 0 -5 -10 -15 -20

0 -50 -100

-2 0 -1 0 -150 00 -5 0 0 50 10 0 15 0 20 0

-100

H (Oe)

-10000

20 15 3 h 10 5 0 -5 -10 -15 -20 H (Oe)

M (emu/g)

-10000

0

-2 0 -1 0 -150 00 -5 0 0 50 10 0 15 0 20 0

-50

20 15 0 h 10 5 0 -5 -10 -15 -20 -200 -100

M (emu/g)

0 M (emu/g)

M (emu/g)

50

-5000

5000

0 H (Oe) 100

20 15 10 5 0 -5 -10 -15 -20

24 h

-2 0 -1 0 -150 00 -5 0 0 50 10 0 15 0 20 0

100

3069

H (Oe)

10000

-10000

-5000

0 H (Oe)

5000

10000

48 h

0 M (emu/g)

M (emu/g)

50

-50 -100 -10000

-5000

20 15 48 h 10 5 0 -5 -10 -15 -20 -200 -100

0 H (Oe)

0 100 200 H (Oe)

5000

10000

Fig. 10. Typical hysteresis loops dependence on milling time, at T¼300 K of Co50Ni50. The inset is an enlargement of low-field region.

decrease of Ms due to the formation of the non magnetic intermetallic compound together with the slight effect of dead or inert layer by decrease in the crystallite size from 80 to 16 nm for fcc-Co(Ni) SS and from 185 to 17 nm for fcc-Ni(Co) along /1 1 1S dense line. On prolonged milling time, the increase in Ms up to 105.4 emu/g can be ascribed to the completion of alloying and the diminution in magneto-crystalline anisotropy due to the microstructure refinement, which leads to an easier rotation of the magnetic vector [37]. Since the magnetization process originates through domain wall movement and spin rotation, it can thus be affected by the crystallite refinement. In other words, at low crystallite sizes, each particle may be treated as a single magnetic domain eliminating the influence of the magnetic walls. One can conclude that the magnetization saturation increases with increase in milling time, while coercivity decreases with increases milling time. Correlating this with the XRD result, it can be attributed to the crystallite size consequent of milling. This variation with crystallite size is also explained on the basics of

domain structure, mean size of particles and crystal anisotropy. Since milling time causes changes by decomposition or transformation of phases, which results in the increase in the crystallite size increases, decreasing the structural strains and subsequently decreased values of coercivity [38]. The remanence to saturation ratio, Mr/Ms, which is another significant magnetic parameter, is plotted as a function of milling time in the Fig. 13. Mr/Ms decreases rapidly, during the first hours of milling, then increases slightly and decreases again reaching a value of about 0.046 after 48 h of milling. On the contrary of what expected in the same polycrystalline materials, the small values of Mr/Ms for both alloys, can probably be linked to the small magnetic particles which are typically single domains. In effect, Ferna´ndez et al. [8] studied the magnetic properties of Co–Ni alloys nanoparticles with composition Co80Ni20 and Co50Ni50 dispersed in a silica matrix using the sol-gel route and found that the squareness ratio, Mr/Ms, is of about 0.42 for both alloys. The values of Mr/Ms, obtained in the present work, are comparable to

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100 I

II

Crystallite size (nm)

90 80

Hc (Oe)

70 60

Co50Ni50

IV

III 90 80 70 60 50 40 30 20 10

Co (Ni)

0

Nanostructured Co50Ni50 powders mixture produced by high energy MA at different times and then examined by SEM, DRX, DSC and VSM. The milling for few hours leads to the Co hcp-fcc transformation and to the formation of two fcc SS: predominate Co(Ni) and Ni(Co). The crystallite size decreases down to the nanometer for both SS. While the fcc Co(Ni) has the highest microstrains and the lowest crystallite size. It is clearly shown that coercivity will be directly affected by the crystallite size and composition. Coercivity and the saturation magnetization of the reached alloy are of about 34.5 Oe and 105.3 emu/g, respectively.

111 200 220

6 12 18 24 30 36 42 48 54 Milling time (h)

I: Fracture II: Welding III: Fracture IV: Steady state

50 40

4. Conclusion

30

References

20 0

6

12

18 24 30 Milling time (h)

36

42

48

54

Fig. 11. Variation of coercivity as a function of milling time.

110 Co50Ni50

Ms (emu/g)

105

100

95

0

6

12

24 18 30 Milling time (h)

36

42

48

54

Fig. 12. Variation of saturation magnetization as a function of milling time.

0.11

Co50Ni50

0.10

Mr/Ms

0.09 0.08 0.07 0.06 0.05 0.04 0

6

12

18 24 30 Milling time (h)

36

42

48

54

Fig. 13. The remanence-to-saturation ratio Mr/Ms as a function of milling time.

those observed in the Fe50Co50 [3] and Ni10Co90 [35] binary alloys synthesized by high energy ball milling. But relatively higher compared to previously obtained results, 0.12, found by Elumalai et al. [12].

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