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Magnetic properties of nanocrystalline Fe–10%Ni alloy obtained by planetary ball mills
Journal of Alloys and Compounds 573 (2013) 157–162
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Magnetic properties of nanocrystalline Fe–10%Ni alloy obtained by planetary ball mills Rabah Hamzaoui a,⇑, Omar Elkedim b a b
Université Paris-Est, Institut de Recherche en Constructibilité, ESTP, 28 avenue du Président Wilson, 94234 Cachan, France Institut FEMTO ST, UMR 6174 CNRS, département MN2S, UTBM, Site de Sévenans, 90010 Belfort Cedex, France
a r t i c l e
i n f o
Article history: Received 10 January 2013 Received in revised form 16 March 2013 Accepted 20 March 2013 Available online 10 April 2013 Keywords: Fe–Ni alloy Ball milling Nanocrystalline Coercivity Magnetization Mechanical alloying
a b s t r a c t Planetary ball mill PM 400 from Retsch (with different milling times for X = 400 rpm, x = 800 rpm) and P4 vario ball mill from Fritsch (with different milling conditions (X/x), X and x being the disc and the vial rotation speeds, respectively) are used for obtaining nanocrystalline Fe–10wt% Ni. The structure and magnetic properties are studied by using X-ray diffraction, SEM and hysteresis measurements, respectively. The bcc-Fe(Ni) phase formation is identified by X-ray diffraction. The higher the shock power and the higher milling time are, the larger the bcc lattice parameter and the lower the grain size. The highest value of the coercivity is 1600 A/m for Fe–10 wt.%Ni (with shock mode (424 rpm/100 rpm) after 36 h of milling), while the lowest value is 189 A/m for (400 rpm/800 rpm) after 72 h of milling. The milling performed in the friction mode has been found to lead the formation of alloys exhibiting a soft magnetic behavior for nanocrystalline Fe–10%Ni. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Ball milling is an efficient and simple method for the fabrication of sub-micron or nanostructured powder materials [1–3]. There are different types of ball milling methods based on the movement of milling balls and vial, such as vibration mill, planetary mill and attritor [1]. In the case of planetary ball milling, main factors that affect the particle size reduction include rotation speed, size of balls, weight ratio of balls to powder, medium of milling, milling time, etc. Because of the variety of the powder materials, the selection of parameters varied substantially [4–8]. Planetary ball mill is characterized by the high energy of the milling media because of centrifugal forces, which can reach up to 20 times the gravitational acceleration. Centrifugal forces are a result of the rotation of a supporting disc and vial (45–500 ml) containing charge in opposite direction, causing the milling media and charge powder to roll off the inner wall of the vial and thereby thrown across the bowl. An important design parameter is that the turning direction of the disc and the vial being opposite to the other causes the centrifugal forces to synchronize and reverse in direction in an alternating manner [1–3]. A new planetary ball milling apparatus has been designed and realized by Gaffet [9] called G5 prototype of P4 from FRITSCH which allows the independent choice of the X and x values where X is the rotational speed of the disc on which the vial holders are fixed. The latter turn at a rotational speed x. In ⇑ Corresponding author. Tel.: +33 0149080334. E-mail address: [email protected] (R. Hamzaoui). 0925-8388/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2013.03.183
literature planetary ball mills have been used for obtaining different nanosructured alloys such as Ni [10], Co–Ni [11], NiTi [12], Mg2Ni [13], Fe–Ni [14–16], Fe–Co [17] and FeSiBCuNb [18–19]. The aim of this work is to study the effect of the ball milling conditions (rotation speed of the disc X, rotation speed of the vials x and the milling time) of two types of planetary ball mill PM 400 from Retsch and P4 vario mill from Fritsch on the structure and magnetic properties for nanocrystalline Fe–10 wt.%Ni. 2. Materials and methods Elemental Fe powder of purity 99.0+% with an average particle size of 7 lm, and elemental Ni powder of purity 99.5% with maximal particle size of 250 lm are used. The elemental powders used (Fe and Ni) are from GoodFellow Company. The milling is carried out using either a planetary high-energy ball mill (Retsch PM 400) and P4 vario ball mill (Fritsch). The planetary high energy ball mill (Retsch PM 400 and Fritsch P4) is constituted by four vials mounted on a planar disc. With the rotation of the disc, the vials move in a circular and in opposite direction compared to the disc rotation. Concerning a planetary ball mill (Retsch PM 400), the ratio of the rotation speed of the disc X and the vials x is two for high X. In our case, the rotation speed of disc is X = 400 rpm, the rotation speed of vials is x = 800 rpm and different milling times t where t = 8 h, 16 h, 24 h, 36 h, 72 h and 96 h. The planetary P4 vario ball mill is the Fritsch commercialized version of the G5 prototype. The planetary ball-milling equipment (G5) is designed by Gaffet [9], which allows the independent choice of the values of X and x (R = 75 mm, r = 20 mm, X 6 1000 rpm, |x| 6 1200 rpm). R is the radius of the disc and r is the internal radius of the vials. In this study, the milling duration is 36 h, while different rotation speeds of disc X where X = 212 rpm, 300 rpm and 424 rpm and different rotation speeds of vials x where x = 50 rpm, 100 rpm, 200 rpm and 400 rpm are considered. Each milling condition is thus characterized by two important parameters (X and x). In Table 1, both the shock power and the friction energy component (expressed in percentage of total injected power) are reported. These
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Table 1 Shock power and friction energy of P4 planetary ball mill for various ball milling conditions. Ball milling conditions, X (rpm)/x (rpm)
Shock power (W)
Friction energy (%)
212/400 300/400 424/50 424/100
1.4 2.6 8.4 8.8
50.1 20.8 2.2 6.1
values are based on the kinematic studies of the G5 ball mill developed by Abdellaoui and Gaffet [20]. According to the (X and x) rotation speed choice, the shock mode process (SMP) prevails when X > x, while the friction mode process (FMP) is stronger when X < x. For both of planetary ball mills used, twenty-millimeter diameter steel balls and 50 ml volume steel vials are used. The total weight of powders is 10 g and the ballto-powder weight ratio is 10:1. X-ray investigations are performed on a Philips X’Pert MPD diffractometer in continuous scanning mode using Cu Ka radiation (k = 0.1541 nm). The lines are measured in the range of 30–130° in steps of 0.01 for 10 s. The changes of the mean lattice parameters a are calculated based on the shift of the angle diffraction lines (several lines used (1 1 0), (2 0 0), (2 1 1), (2 2 0)) using Bragg law like shown in the following equation:
2dhkl sin ðhÞ ¼ k
ð1Þ
where h is the Bragg angle, k is the wavelength of the X-ray radiation and dhkl is distance between atomic layers and equal in our case:
a dhkl ¼ pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi 2 2 2 h þk þl
ð2Þ
a is lattice parameter and (h k l) are Miller indices denotes a plane orthogonal to a direction [h k ‘]. It is noticed that the accuracy of the lattice parameter is the mean of all standard deviations of each angle and each milling conditions. It is estimated to 5 105 nm. The crystallite size and lattice strain are calculated based on the approximation (Eq. (3)) that combines the Wilson formula and Scherrer formula following Williamson–Hall style plot [21].
b ¼ 2e tanðhÞ þ
0:9k D cosðhÞ
ð3Þ
where b is the full-width at half maximum intensity of a Bragg reflection excluding instrumental broadening, h is the Bragg angle, k is the wavelength of the X-ray radiation, e is the effective lattice strain and D the average crystallite size. On rearrangement, the Eq. (3) can be written as:
b cos ðhÞ ¼ 2e sinðhÞ þ
0:9k D
ð4Þ
The b cos (h) of each peak is plotted as a function of sin (h) and a straight line can be obtained by the least squares method with the intercept as 0.9 k/D and slope as 2e. From these, the crystallite size D and lattice strain e can be calculated. The powder microstructure are characterized by scanning electron microscopy (SEM) (stereo scan 120, Leo 120), with IDIFIX program (image analysis program associated to the SEM machine) coupled with energy dispersive spectrometer (EDS). The standard deviation of the chemical composition of samples from the nominal ones is less than 1%. A contamination of about 0.2 wt.% in Cr is found only in sample of Fe–10 wt.% Ni for 72 h of milling (attributed to the wear of Cr containing steel surfaces of balls). Taking into account the vial volume (50 ml) and the powder weight (10 g), this calculated method leads to constant oxygen to powder ratio of less than 0.14%. Transmission electronic Microscopy (TEM) observations are performed by JEOL JEM-2100 LaB6 microscope operating at 200 kV and equipped with a high tilt pole-piece achieving a point-to-point resolution of 0.25 nm. The magnetic measurements are realized at 300 K using a Hysteresimeter Bull M 2000/2100.
3. Results 3.1. Structure X-ray diffraction measurements are performed for all mixed powders at different stages of milling. The obtained diffraction patterns allowed controlling the process of alloy formation. Fig. 1 presents the X-ray patterns for Fe–10 wt.%Ni samples after different mechanical milling conditions (X, x, t). Concerning the milling
Fig. 1. X-ray diffraction patterns of mechanical alloying for Fe–10 wt.%Ni as a function of milling conditions (X (rpm)/x (rpm), milling time t).
time conditions by using planetary ball mill (RESTCH PM400), it is observed that the characteristic Ni lines decreased gradually and decay after 24 h and a bcc Fe(Ni) solid solution is formed. For planetary ball mill FRISTCH P4, it is observed that the higher ratio X/x (higher shock power) is, the lower the Ni peak intensity and decay for ratio X/x larger than 0.53 (equal to 1.4 W in shock power). Furthermore, a bcc Fe (Ni) solid solution is formed. In order to analyze the bcc Fe(Ni) formation, the evolution of the bcc lattice parameter versus the shock power and the milling time are shown in Fig. 2. It is observed that the mean value of the lattice parameter increases slightly from a = 0.28659 ± 0.00005 to 0.28695 ± 0.00005 nm for the shock power equal to 8.83 W and 0.28704 ± 0.00005 nm for 96 h milling time. Fig. 3 presents the evolution of mean crystallite size D and lattice strain e as a function of the shock power and of the milling time for nanocrystalline mechanically alloyed Fe–10 wt.%Ni. It is concluded that the higher the shock power is, the lower the crystallite size. It is also noticed that the lowest values of the crystallite size with the highest values of the lattice strain are found in SMP, whereas, the lowest values of the crystallite size with the lowest values of the lattice strain are found in FMP. Concerning the milling time, the reduction of the crystallite sizes is accompanied by the increase of the lattice strain level when milling time increases. The D value of the crystallite size has been found to decrease from D = 54 ± 5 nm for unmilled powder to D = 10.2 ± 2.0 nm for the shock power equal
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159
Fig. 3. Evolution of mean crystallite size D and lattice strain e for nanocrystalline Fe–10 wt.%Ni: (a) vs. milling time (X = 400 rpm, x = 800 rpm) (b) after 36 h of milling vs. shock power (SP expressed in W). Fig. 2. Evolution of lattice parameter a for nanocrystalline Fe–10 wt.%Ni: (a) vs. milling time (X = 400 rpm, x = 800 rpm), (b) after 36 h of milling vs. shock power (SP expressed in W).
to 8.83 W and stabilized to D = 10 ± 2 nm for 72 h and 96 h of the milling. The mean crystallite size D is equal to 9 nm for 72 h milling time have been found by Transmission electronic microscopy (see Fig. 4) and confirmed our results. The e value of the internal strain increases from e = 0.07 ± 0.05% for unmilled powder to e = 0.61 ± 0.05% for the shock power equal to 8.83 W and to e = 0.7 ± 0.05% for 96 h of milling for Fe–10 wt.%Ni alloy. The authors would like to precise that the initial value of the crystallite size and the lattice strain are only given by the instrumental width of the X-ray diffraction peaks. Fig. 4 shows a dark field image of the powder milled for 72 h. It can be seen that crystallites are aggregated and that their size distribution is between 5 and 14 nm. The majority of the crystallites taken in different regions have a size of about 10 nm which is in agreement with XRD. It is noticed that, depth microscopy study (SEM and TEM) concerning Fe–Ni is under investigation and it will be published later. Fig. 5a–e show the morphology of the powders after milling conditions (X, x, t). The alloyed particles formed by an assembly of nanocrystalline grains, (which size is equal to the coherence domain size given by the XRD patterns analyses) showed heterogeneous distributions. Their sizes varied from 4 lm to 80 lm for FMP and milling time. For SMP their sizes varied from 20 lm to 100 lm. During mechanical alloying, morphology and size distribution of powder particles have been changed because of the cold welding and fracturing of the particles.
Fig. 4. TEM dark field image for nanocrystalline Fe–10 wt.%Ni powders milled for 72 h (X = 400 rpm, x = 800 rpm).
3.2. Magnetic properties Fig. 6 presents the dependence of saturation magnetization Ms as function of the shock power and of the milling time. The higher
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(c)
(a)
Fe Ni
(d)
(b)
(e)
Fig. 5. SEM micrographs for nanocrystalline Fe–10 wt.%Ni for different milling conditions (X, x, t), (a) before milling. (b) Retsch PM400 (400 rpm, 800 rpm, 36 h). (c) Retsch PM400 (400 rpm, 800 rpm, 72 h). (d) Fritsch P4, FMP (212 rpm, 400 rpm, 36 h). (e) Fritsch P4, SMP (424 rpm, 50 rpm, 36 h).
the shock power and the higher milling time are, the higher the saturation magnetization. It is observed that the mean value of the saturation magnetization Ms. increases from 193 ± 2 Am2/kg to 226 ± 2 Am2/kg for the shock power equal to 8.83 W and 227 ± 2 Am2/kg for 96 h milling time. The evolution of coercivity Hc as a function of the shock power and of the milling time is shown in Fig. 7. It is observed that the coercivity decreases when milling time increases (Fig. 7a). The mean value of sharp decrease of the coercivity decreases from 1220 ± 10 A/m to a coercivity Hc, which lasts up to 72 h of milling and stabilized to 200 ± 10 after 96 h of milling. Concerning the evolution of the coercivity versus shock power, it has been found that shock mode process (SHP) gives the high mean values of Hc which equal to 1600 ± 10 A/m and 1500 ± 10 A/m for milling conditions (4 24/1 0 0) and (42 4/50) respectively. Whereas, friction mode process (FMP) gives low mean values of Hc which equal to 500 ± 10 A/m and 420 ± 10 A/m for milling conditions (3 0 0/4 0 0) and (2 1 2/4 0 0) respectively.
4. Discussions 4.1. Structure Considering the given results for Fe–10%Ni powders obtained by MA, a bcc solid solution Fe(Ni) is observed (see Fig. 1). Equilibrium solubility of nickel in the iron lattice is reported to be around 3.5 wt.% Ni [22,23] by conventional technique, and extended up to 30 at.% Ni by ball milling [24]. The diffraction peaks of the as milled alloys exhibit significant broadening due to the refinement of grain size and the increase of internal lattice strain. Such a solid solution Fe(Ni) formation have been previously observed by Kuhrt and Schultz [4]. Furthermore, a bcc Fe (Ni) solid solution is formed. Baldokhin et al. [25] and Tcherdyntsev et al. [26] have shown that milled Fe–10%wtNi alloyed powder contains single bcc phase. Annealing of mechanical alloyed Fe–10%Ni at 623 K did not lead to any change of crystalline structure. Annealing
R. Hamzaoui, O. Elkedim / Journal of Alloys and Compounds 573 (2013) 157–162 230
(a)
220
2
Magnetization Ms (Am /kg)
225
161
215 210 205 200 195 190
0
20
40
60
80
100
Milling time t (h) 230
(b)
220
2
Magnetization Ms (Am /kg)
225
215 210 205 200 195 190 0
2
4
6
8
10
Shock Power (W) Fig. 6. Evolution of saturation magnetization Ms for nanocrystalline Fe–10 wt.%Ni: (a) vs. milling time (X = 400 rpm, x = 800 rpm). (b) After 36 h of milling vs. shock power (SP expressed in W).
at 923 K for Fe–10 wt.%Ni has been found leading to the formation of two phases (bcc + fcc) and product. Kaloshkin et al. [15] have shown that concentration ranges of single bcc phase solid solution are larger in mechanical alloyed alloys compared with the equilibrium phase diagram. Such a minor change of the lattice parameter (see Fig. 2a and b) can be due to the small atomic size difference between the Ni and Fe atoms [27]. The lattice parameter of pure Ni is larger than that of pure Fe. As the milling time and shock power increases, the lattice parameter increases (Fe) due to Ni diffusion in Fe. Secondly, the rate of the increase in the lattice parameter decreased (Ni) due to the completion in the solid solution Fe(Ni). Generally, the increase of the lattice parameters has been attributed to the solid solution formation preceding the amorphous phase formation as the milling time increases [1]. However, the decrease of the lattice parameter has been attributed to the formation of the nanocrystalline intermetallic [1]. Pekala et al. [28] have found that the value of crystallite size and lattice strain reached 10 nm and 0.46%, respectively after 40 h of milling using high energy ball milling. Using low-energy milling, the same authors [28] have found that these parameters reach 9 nm and 0.47% after 400 h of milling. Concerning given results of the evolution of the crystallite size and of the lattice strain as a function as the shock power and milling time, it is observed for the mechanically alloyed Fe–10 wt.%Ni powders that when the shock power increases, the crystallite size decreases. Also, it is shown that with SMP we obtain the lowest values of the crystallite size and the highest values of the lattice strain. Whereas, with FMP
Fig. 7. Evolution of coercivity Hc for nanocrystalline Fe–10 wt.%Ni: (a) vs. milling time (X = 400 rpm, x = 800 rpm), (b) after 36 h of milling vs. shock power (SP expressed in W).
we obtain the lowest values of the crystallite size and the lowest values of the lattice strain. Using the G5 planetary ball milling for mechanically alloying Fe–Ni powders, Hays et al. [5] have shown that using FMP, the crystallite size decreases gradually as a function of the milling time with a particle size varying between 1 lm and 5 lm, whereas, in SMP, the crystallite size decrease sharply versus the milling time with a particle size varying between 20 lm and 100 lm. In Fig. 5, it is found that the as-milled Fe(Ni) alloy exhibits a typical microstructure of mechanically milled particles, which has a granular shape, rough surface, and a relatively broad distribution in particle size. The particles tend to agglomerate in the MA process because the grinding function in the stainless steel ball plays a leading role during the mechanical alloying, which leads to the formation of the granular shape.
4.2. Magnetic properties The increase of the magnetization is found to be linked to the crystallite size reduction to about 10 nm (see Fig. 6). Each grain may be considered as being a single magnetic domain eliminating the influence from magnetic walls [29,16]. On the other hand, the larger the lattice parameter, the higher the saturation magnetization. It has been shown by Amils et al. [6] that the ball milling pro-
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cess produces a high density of defects, particularly defects of antisite type, which causes a 0.8% lattice parameter expansion in the case of the Fe–Al alloys. The same authors [6] have previously reported on a good correlation between the saturation magnetization and the lattice parameter (increase of the lattice parameters and the saturation magnetization), which suggests that the observed magnetic transition may be partially related to the changes in the density of states at the Fermi level. Based on our results, it is obvious that the magnetization can be modified by the crystallite size and by the lattice parameter but not by the used mode (SMP, FMP). The increase of the coercivity in SMP can be understood as being the consequence of a considerable introduction of high internal strains into the material, which is inevitably related to the process of preparation (see Fig. 7b). Thus, the magnetostriction in combination with the high internal strain has been identified as the effect dominating the coercivity by the way of the magnetoelastic interactions [30]. The low values of the coercivity in FMP can be due to the low values of the crystallite size (in Fig. 7) where, the higher the milling time is, the lower the crystallite size. Also to the dependence of the coercivity and the crystallite size (the lower the coercivity is, the lower the crystallite size) prevails over the other predominant strain influence in this nanocrystallite size range [16]. Nevertheless, the resulting nanocrystalline structure led to the soft magnetic properties (low coercivity with low grain size) as suggested by Herzer [31]. However, the smaller the structural correlation lengths, the lower are the coercivity. Such a feature has been previously noticed in the case of amorphous alloys and of nanocrystalline alloys for grain size lower than 20 nm [31]. 5. Conclusions Nanostructured Fe-based soft magnetic could show some major technological advantages due to their predicted low coercivity and high saturation magnetization. In order to obtain such materials, it is necessary to produce Fe-based alloys, with an average crystallite size largely below 20 nm and most probably even below 10 nm. Mechanical alloying is one of the techniques that can synthesize such materials. The results presented in this work by using the PM 400 from Retsch and the vario mill (P4) planetary ball mill from Fritsch clearly shown great potential of this technique. High-energy ball milling of Fe and Ni powders results in the solid solution formation accompanied by a grain refinement for nanocrystalline Fe–10%Ni. The higher the shock power and milling time, the larger the lattice parameter (a = 0.28695 ± 0.00005 nm and 0.28704 ± 0.00005 nm for 8.83 W of the shock power and 96 h of the milling time, respectively) and the lower the crystallite size (10 ± 2 and 10 ± 2 nm for 8.83 W of the shock power and 96 h of the milling time, respectively). It has been found that the magnetization is not modified by milling performed in shock mode process (SMP) in friction mode process (FMP). Such a magnetization is only modified by the decrease of the crystallite size and by the increase of the lattice parameter. The highest values of coercivity have been found in SMP, i.e., 1600 A/m for Fe–10 wt.%Ni for the following milling condition (424 rpm/100 rpm). Following the classification of the magnetic materials (hard magnetic, memory materials and soft
magnetic materials), the FMP leads to the formation of alloys exhibiting a soft magnetic behavior (Hc between 380 A/m and 750 A/m) for Fe–10 wt.%Ni alloys and the SMP leads to the formation of alloys exhibiting a memory and hard magnetic behavior (Hc = 900–1600 A/m) for Fe–10 wt.%Ni alloys. Acknowledgments The authors wish to thank Dr E. GAFFET for his assistance concerning the values which are based on the kinematic studies of the G 5 ball mill. Also, the authors wish to thank Mr DE OLIVEIRA from FRITSCH Company and Mr VAXELAIRE from RETSCH Company which have given us more information concerning their machines (FRITSCH P4 and RETSCH PM400). References [1] C. Suryanarayana, Prog. Mater. Sci. 46 (2001) 1–184. [2] S. El-Eskandarany, Noyes Publications, William Andrew Publishing, Norwich, New York, USA, 2001, pp. 1–232. [3] E. Gaffet, G. Le Caër, in:Mechanical Processing for Nanomaterials, H.S. Nalwa (Ed.) Encyclopedia of Nanoscience and Nanotechnology, American Scientific Publishers, vol. 5 (2004), 91–129. [4] C. Kuhrt, L. Schultz, J. Appl. Phys. 75 (1993) 1975–1980. [5] V. Hays, R. Marchand, G. Saindrenan, E. Gaffet, Nanostruct. Mater. 7 (1996) 411–420. [6] X. Amils, J. Nogues, S. Surinach, M.D. Bar´o, J. Magn. Magn. Mater. 203 (1999) 129–131. [7] E. Jartych, J. Magn. Magn. Mater. 265 (2003) 176–188. [8] Yugo Nomura, Kazuo Fujiwara, Akihiko Terada, Satoshi Nakai, Masaaki Hosomi, Chemosphere 86 (2012) 228–234. [9] E. Gaffet, Mater. Sci. Eng. A 132 (1991) 181–193. [10] G.K. Rane, U. Welzel, E.J. Mittemeijer, Acta Mater V60 (2012) 7011– 7023. [11] N.E. Fenineche, R. Hamzaoui, O. Elkedim, Mater. Lett. 4493 (2003) 1–5. [12] S. Tria, O. Elkedim, R. Hamzaoui, X. Guo, F. Bernard, N. Millot, O. Rapaud, Powder Technol. 210 (2011) 181–188. [13] L.W. Huang, O. Elkedim, R. Hamzaoui, J.Allo.Comp. 509S (2011) 328–333. [14] Kh. Gheisari, S. Javadpour, J.T. Oh, M. Ghaffari, J.Allo.Comp. 20 (2009) 416–420. [15] S.D. Kaloshkin, V.V. Tcherdyntsev, Yu.V. Baldokhin, I.A. Tomilin, E.V. Shelekhov, J. Non-Cryst. Sol. 287 (2001) 329–333. [16] R. Hamzaoui, O. Elkedim, N. Fenineche, E. Gaffet, J. Craven, Mater. Sci. Eng. A 360 (2003) 299–305. [17] N.E. Fenineche, O. El Kedim, E. Gaffet, J. Metast, Nanocryst. Mater. 7 (2000) 41– 48. [18] R. Hamzaoui, O. Elkedim, J.Y. Rauch, M. Cherigui, D. Bassir, Recent Res. Devel. Mat. Sci. Engg. 4 (2011) 1–13. [19] Boqu Chen, Sha Yang, Xingxing Liu, Biao Yan, Lu Wei, J. Alloys Comp. 448 (2008) 234–237. [20] M. Abdellaoui, E. Gaffet, Acta. Metall Mater. 44 (1995) 1087–1098. [21] G.K. Williamson, W.H. Hall, Acta Metall (1) (1953) 22–31. [22] J. Yang, J.I. Goldstein, Lunar Planet. Sci. XXXIV (2003) 1156. [23] R.Hamzaoui, PHD Thesis, UTBM-UFC Belfort, France, 2004. [24] S.D. Kaloshkin, V.V. Tcherdyntsev, Yu.V. Baldokhin, I.A. Tomilin, E.V. Shelekhov, Physica B: Condensed Matter 299 (V) (2001) 236–241. [25] Yu.V. Baldokhin, V.V. Tcherdyntsev, S.D. Kaloshkin, G.A. Kochetov, Yu.A. Pustov, J. Magn. Magn. Mater. 203 (1999) 313–315. [26] V.V. Tcherdyntsev, S.D. Kaloshkin, I.A. Tomilin, E.V. Shelekhov, Yu.V. Baldokhin, Nanostruct. Mater. 12 (1999) 139–142. [27] M. Pekala, D. Oleszak, E. Jartych, J.K. Zurawicz, Nanostruct. Mater. 11 (1999) 789–796. [28] M. Pekala, D. Oleszak, E. Jartych, J.K. Zurawicz, J. Non-Cryst. Sol. 250 (1999) 757–761. [29] E. Jartych, J.K. Zurawicz, D. Oleszak, M. Pekala, J. Magn. Magn. Mater. 208 (2000) 221–230. [30] C. Kuhrt, L. Schultz, J. Appl. Phys. 73 (1993) 6588–6590. [31] G. Herzer, Nanocrystalline soft magnetic alloys, in:K.H.J. Buschow (Ed.), Handbook of Magnetic. Materials, Elsevier, Amsterdam, 1997, Vol. 10, pp. 415–462.
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Magnetic properties of nanocrystalline Fe–10%Ni alloy obtained by planetary ball mills Rabah Hamzaoui a,⇑, Omar Elkedim b a b
Université Paris-Est, Institut de Recherche en Constructibilité, ESTP, 28 avenue du Président Wilson, 94234 Cachan, France Institut FEMTO ST, UMR 6174 CNRS, département MN2S, UTBM, Site de Sévenans, 90010 Belfort Cedex, France
a r t i c l e
i n f o
Article history: Received 10 January 2013 Received in revised form 16 March 2013 Accepted 20 March 2013 Available online 10 April 2013 Keywords: Fe–Ni alloy Ball milling Nanocrystalline Coercivity Magnetization Mechanical alloying
a b s t r a c t Planetary ball mill PM 400 from Retsch (with different milling times for X = 400 rpm, x = 800 rpm) and P4 vario ball mill from Fritsch (with different milling conditions (X/x), X and x being the disc and the vial rotation speeds, respectively) are used for obtaining nanocrystalline Fe–10wt% Ni. The structure and magnetic properties are studied by using X-ray diffraction, SEM and hysteresis measurements, respectively. The bcc-Fe(Ni) phase formation is identified by X-ray diffraction. The higher the shock power and the higher milling time are, the larger the bcc lattice parameter and the lower the grain size. The highest value of the coercivity is 1600 A/m for Fe–10 wt.%Ni (with shock mode (424 rpm/100 rpm) after 36 h of milling), while the lowest value is 189 A/m for (400 rpm/800 rpm) after 72 h of milling. The milling performed in the friction mode has been found to lead the formation of alloys exhibiting a soft magnetic behavior for nanocrystalline Fe–10%Ni. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Ball milling is an efficient and simple method for the fabrication of sub-micron or nanostructured powder materials [1–3]. There are different types of ball milling methods based on the movement of milling balls and vial, such as vibration mill, planetary mill and attritor [1]. In the case of planetary ball milling, main factors that affect the particle size reduction include rotation speed, size of balls, weight ratio of balls to powder, medium of milling, milling time, etc. Because of the variety of the powder materials, the selection of parameters varied substantially [4–8]. Planetary ball mill is characterized by the high energy of the milling media because of centrifugal forces, which can reach up to 20 times the gravitational acceleration. Centrifugal forces are a result of the rotation of a supporting disc and vial (45–500 ml) containing charge in opposite direction, causing the milling media and charge powder to roll off the inner wall of the vial and thereby thrown across the bowl. An important design parameter is that the turning direction of the disc and the vial being opposite to the other causes the centrifugal forces to synchronize and reverse in direction in an alternating manner [1–3]. A new planetary ball milling apparatus has been designed and realized by Gaffet [9] called G5 prototype of P4 from FRITSCH which allows the independent choice of the X and x values where X is the rotational speed of the disc on which the vial holders are fixed. The latter turn at a rotational speed x. In ⇑ Corresponding author. Tel.: +33 0149080334. E-mail address: [email protected] (R. Hamzaoui). 0925-8388/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2013.03.183
literature planetary ball mills have been used for obtaining different nanosructured alloys such as Ni [10], Co–Ni [11], NiTi [12], Mg2Ni [13], Fe–Ni [14–16], Fe–Co [17] and FeSiBCuNb [18–19]. The aim of this work is to study the effect of the ball milling conditions (rotation speed of the disc X, rotation speed of the vials x and the milling time) of two types of planetary ball mill PM 400 from Retsch and P4 vario mill from Fritsch on the structure and magnetic properties for nanocrystalline Fe–10 wt.%Ni. 2. Materials and methods Elemental Fe powder of purity 99.0+% with an average particle size of 7 lm, and elemental Ni powder of purity 99.5% with maximal particle size of 250 lm are used. The elemental powders used (Fe and Ni) are from GoodFellow Company. The milling is carried out using either a planetary high-energy ball mill (Retsch PM 400) and P4 vario ball mill (Fritsch). The planetary high energy ball mill (Retsch PM 400 and Fritsch P4) is constituted by four vials mounted on a planar disc. With the rotation of the disc, the vials move in a circular and in opposite direction compared to the disc rotation. Concerning a planetary ball mill (Retsch PM 400), the ratio of the rotation speed of the disc X and the vials x is two for high X. In our case, the rotation speed of disc is X = 400 rpm, the rotation speed of vials is x = 800 rpm and different milling times t where t = 8 h, 16 h, 24 h, 36 h, 72 h and 96 h. The planetary P4 vario ball mill is the Fritsch commercialized version of the G5 prototype. The planetary ball-milling equipment (G5) is designed by Gaffet [9], which allows the independent choice of the values of X and x (R = 75 mm, r = 20 mm, X 6 1000 rpm, |x| 6 1200 rpm). R is the radius of the disc and r is the internal radius of the vials. In this study, the milling duration is 36 h, while different rotation speeds of disc X where X = 212 rpm, 300 rpm and 424 rpm and different rotation speeds of vials x where x = 50 rpm, 100 rpm, 200 rpm and 400 rpm are considered. Each milling condition is thus characterized by two important parameters (X and x). In Table 1, both the shock power and the friction energy component (expressed in percentage of total injected power) are reported. These
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Table 1 Shock power and friction energy of P4 planetary ball mill for various ball milling conditions. Ball milling conditions, X (rpm)/x (rpm)
Shock power (W)
Friction energy (%)
212/400 300/400 424/50 424/100
1.4 2.6 8.4 8.8
50.1 20.8 2.2 6.1
values are based on the kinematic studies of the G5 ball mill developed by Abdellaoui and Gaffet [20]. According to the (X and x) rotation speed choice, the shock mode process (SMP) prevails when X > x, while the friction mode process (FMP) is stronger when X < x. For both of planetary ball mills used, twenty-millimeter diameter steel balls and 50 ml volume steel vials are used. The total weight of powders is 10 g and the ballto-powder weight ratio is 10:1. X-ray investigations are performed on a Philips X’Pert MPD diffractometer in continuous scanning mode using Cu Ka radiation (k = 0.1541 nm). The lines are measured in the range of 30–130° in steps of 0.01 for 10 s. The changes of the mean lattice parameters a are calculated based on the shift of the angle diffraction lines (several lines used (1 1 0), (2 0 0), (2 1 1), (2 2 0)) using Bragg law like shown in the following equation:
2dhkl sin ðhÞ ¼ k
ð1Þ
where h is the Bragg angle, k is the wavelength of the X-ray radiation and dhkl is distance between atomic layers and equal in our case:
a dhkl ¼ pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi 2 2 2 h þk þl
ð2Þ
a is lattice parameter and (h k l) are Miller indices denotes a plane orthogonal to a direction [h k ‘]. It is noticed that the accuracy of the lattice parameter is the mean of all standard deviations of each angle and each milling conditions. It is estimated to 5 105 nm. The crystallite size and lattice strain are calculated based on the approximation (Eq. (3)) that combines the Wilson formula and Scherrer formula following Williamson–Hall style plot [21].
b ¼ 2e tanðhÞ þ
0:9k D cosðhÞ
ð3Þ
where b is the full-width at half maximum intensity of a Bragg reflection excluding instrumental broadening, h is the Bragg angle, k is the wavelength of the X-ray radiation, e is the effective lattice strain and D the average crystallite size. On rearrangement, the Eq. (3) can be written as:
b cos ðhÞ ¼ 2e sinðhÞ þ
0:9k D
ð4Þ
The b cos (h) of each peak is plotted as a function of sin (h) and a straight line can be obtained by the least squares method with the intercept as 0.9 k/D and slope as 2e. From these, the crystallite size D and lattice strain e can be calculated. The powder microstructure are characterized by scanning electron microscopy (SEM) (stereo scan 120, Leo 120), with IDIFIX program (image analysis program associated to the SEM machine) coupled with energy dispersive spectrometer (EDS). The standard deviation of the chemical composition of samples from the nominal ones is less than 1%. A contamination of about 0.2 wt.% in Cr is found only in sample of Fe–10 wt.% Ni for 72 h of milling (attributed to the wear of Cr containing steel surfaces of balls). Taking into account the vial volume (50 ml) and the powder weight (10 g), this calculated method leads to constant oxygen to powder ratio of less than 0.14%. Transmission electronic Microscopy (TEM) observations are performed by JEOL JEM-2100 LaB6 microscope operating at 200 kV and equipped with a high tilt pole-piece achieving a point-to-point resolution of 0.25 nm. The magnetic measurements are realized at 300 K using a Hysteresimeter Bull M 2000/2100.
3. Results 3.1. Structure X-ray diffraction measurements are performed for all mixed powders at different stages of milling. The obtained diffraction patterns allowed controlling the process of alloy formation. Fig. 1 presents the X-ray patterns for Fe–10 wt.%Ni samples after different mechanical milling conditions (X, x, t). Concerning the milling
Fig. 1. X-ray diffraction patterns of mechanical alloying for Fe–10 wt.%Ni as a function of milling conditions (X (rpm)/x (rpm), milling time t).
time conditions by using planetary ball mill (RESTCH PM400), it is observed that the characteristic Ni lines decreased gradually and decay after 24 h and a bcc Fe(Ni) solid solution is formed. For planetary ball mill FRISTCH P4, it is observed that the higher ratio X/x (higher shock power) is, the lower the Ni peak intensity and decay for ratio X/x larger than 0.53 (equal to 1.4 W in shock power). Furthermore, a bcc Fe (Ni) solid solution is formed. In order to analyze the bcc Fe(Ni) formation, the evolution of the bcc lattice parameter versus the shock power and the milling time are shown in Fig. 2. It is observed that the mean value of the lattice parameter increases slightly from a = 0.28659 ± 0.00005 to 0.28695 ± 0.00005 nm for the shock power equal to 8.83 W and 0.28704 ± 0.00005 nm for 96 h milling time. Fig. 3 presents the evolution of mean crystallite size D and lattice strain e as a function of the shock power and of the milling time for nanocrystalline mechanically alloyed Fe–10 wt.%Ni. It is concluded that the higher the shock power is, the lower the crystallite size. It is also noticed that the lowest values of the crystallite size with the highest values of the lattice strain are found in SMP, whereas, the lowest values of the crystallite size with the lowest values of the lattice strain are found in FMP. Concerning the milling time, the reduction of the crystallite sizes is accompanied by the increase of the lattice strain level when milling time increases. The D value of the crystallite size has been found to decrease from D = 54 ± 5 nm for unmilled powder to D = 10.2 ± 2.0 nm for the shock power equal
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159
Fig. 3. Evolution of mean crystallite size D and lattice strain e for nanocrystalline Fe–10 wt.%Ni: (a) vs. milling time (X = 400 rpm, x = 800 rpm) (b) after 36 h of milling vs. shock power (SP expressed in W). Fig. 2. Evolution of lattice parameter a for nanocrystalline Fe–10 wt.%Ni: (a) vs. milling time (X = 400 rpm, x = 800 rpm), (b) after 36 h of milling vs. shock power (SP expressed in W).
to 8.83 W and stabilized to D = 10 ± 2 nm for 72 h and 96 h of the milling. The mean crystallite size D is equal to 9 nm for 72 h milling time have been found by Transmission electronic microscopy (see Fig. 4) and confirmed our results. The e value of the internal strain increases from e = 0.07 ± 0.05% for unmilled powder to e = 0.61 ± 0.05% for the shock power equal to 8.83 W and to e = 0.7 ± 0.05% for 96 h of milling for Fe–10 wt.%Ni alloy. The authors would like to precise that the initial value of the crystallite size and the lattice strain are only given by the instrumental width of the X-ray diffraction peaks. Fig. 4 shows a dark field image of the powder milled for 72 h. It can be seen that crystallites are aggregated and that their size distribution is between 5 and 14 nm. The majority of the crystallites taken in different regions have a size of about 10 nm which is in agreement with XRD. It is noticed that, depth microscopy study (SEM and TEM) concerning Fe–Ni is under investigation and it will be published later. Fig. 5a–e show the morphology of the powders after milling conditions (X, x, t). The alloyed particles formed by an assembly of nanocrystalline grains, (which size is equal to the coherence domain size given by the XRD patterns analyses) showed heterogeneous distributions. Their sizes varied from 4 lm to 80 lm for FMP and milling time. For SMP their sizes varied from 20 lm to 100 lm. During mechanical alloying, morphology and size distribution of powder particles have been changed because of the cold welding and fracturing of the particles.
Fig. 4. TEM dark field image for nanocrystalline Fe–10 wt.%Ni powders milled for 72 h (X = 400 rpm, x = 800 rpm).
3.2. Magnetic properties Fig. 6 presents the dependence of saturation magnetization Ms as function of the shock power and of the milling time. The higher
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(c)
(a)
Fe Ni
(d)
(b)
(e)
Fig. 5. SEM micrographs for nanocrystalline Fe–10 wt.%Ni for different milling conditions (X, x, t), (a) before milling. (b) Retsch PM400 (400 rpm, 800 rpm, 36 h). (c) Retsch PM400 (400 rpm, 800 rpm, 72 h). (d) Fritsch P4, FMP (212 rpm, 400 rpm, 36 h). (e) Fritsch P4, SMP (424 rpm, 50 rpm, 36 h).
the shock power and the higher milling time are, the higher the saturation magnetization. It is observed that the mean value of the saturation magnetization Ms. increases from 193 ± 2 Am2/kg to 226 ± 2 Am2/kg for the shock power equal to 8.83 W and 227 ± 2 Am2/kg for 96 h milling time. The evolution of coercivity Hc as a function of the shock power and of the milling time is shown in Fig. 7. It is observed that the coercivity decreases when milling time increases (Fig. 7a). The mean value of sharp decrease of the coercivity decreases from 1220 ± 10 A/m to a coercivity Hc, which lasts up to 72 h of milling and stabilized to 200 ± 10 after 96 h of milling. Concerning the evolution of the coercivity versus shock power, it has been found that shock mode process (SHP) gives the high mean values of Hc which equal to 1600 ± 10 A/m and 1500 ± 10 A/m for milling conditions (4 24/1 0 0) and (42 4/50) respectively. Whereas, friction mode process (FMP) gives low mean values of Hc which equal to 500 ± 10 A/m and 420 ± 10 A/m for milling conditions (3 0 0/4 0 0) and (2 1 2/4 0 0) respectively.
4. Discussions 4.1. Structure Considering the given results for Fe–10%Ni powders obtained by MA, a bcc solid solution Fe(Ni) is observed (see Fig. 1). Equilibrium solubility of nickel in the iron lattice is reported to be around 3.5 wt.% Ni [22,23] by conventional technique, and extended up to 30 at.% Ni by ball milling [24]. The diffraction peaks of the as milled alloys exhibit significant broadening due to the refinement of grain size and the increase of internal lattice strain. Such a solid solution Fe(Ni) formation have been previously observed by Kuhrt and Schultz [4]. Furthermore, a bcc Fe (Ni) solid solution is formed. Baldokhin et al. [25] and Tcherdyntsev et al. [26] have shown that milled Fe–10%wtNi alloyed powder contains single bcc phase. Annealing of mechanical alloyed Fe–10%Ni at 623 K did not lead to any change of crystalline structure. Annealing
R. Hamzaoui, O. Elkedim / Journal of Alloys and Compounds 573 (2013) 157–162 230
(a)
220
2
Magnetization Ms (Am /kg)
225
161
215 210 205 200 195 190
0
20
40
60
80
100
Milling time t (h) 230
(b)
220
2
Magnetization Ms (Am /kg)
225
215 210 205 200 195 190 0
2
4
6
8
10
Shock Power (W) Fig. 6. Evolution of saturation magnetization Ms for nanocrystalline Fe–10 wt.%Ni: (a) vs. milling time (X = 400 rpm, x = 800 rpm). (b) After 36 h of milling vs. shock power (SP expressed in W).
at 923 K for Fe–10 wt.%Ni has been found leading to the formation of two phases (bcc + fcc) and product. Kaloshkin et al. [15] have shown that concentration ranges of single bcc phase solid solution are larger in mechanical alloyed alloys compared with the equilibrium phase diagram. Such a minor change of the lattice parameter (see Fig. 2a and b) can be due to the small atomic size difference between the Ni and Fe atoms [27]. The lattice parameter of pure Ni is larger than that of pure Fe. As the milling time and shock power increases, the lattice parameter increases (Fe) due to Ni diffusion in Fe. Secondly, the rate of the increase in the lattice parameter decreased (Ni) due to the completion in the solid solution Fe(Ni). Generally, the increase of the lattice parameters has been attributed to the solid solution formation preceding the amorphous phase formation as the milling time increases [1]. However, the decrease of the lattice parameter has been attributed to the formation of the nanocrystalline intermetallic [1]. Pekala et al. [28] have found that the value of crystallite size and lattice strain reached 10 nm and 0.46%, respectively after 40 h of milling using high energy ball milling. Using low-energy milling, the same authors [28] have found that these parameters reach 9 nm and 0.47% after 400 h of milling. Concerning given results of the evolution of the crystallite size and of the lattice strain as a function as the shock power and milling time, it is observed for the mechanically alloyed Fe–10 wt.%Ni powders that when the shock power increases, the crystallite size decreases. Also, it is shown that with SMP we obtain the lowest values of the crystallite size and the highest values of the lattice strain. Whereas, with FMP
Fig. 7. Evolution of coercivity Hc for nanocrystalline Fe–10 wt.%Ni: (a) vs. milling time (X = 400 rpm, x = 800 rpm), (b) after 36 h of milling vs. shock power (SP expressed in W).
we obtain the lowest values of the crystallite size and the lowest values of the lattice strain. Using the G5 planetary ball milling for mechanically alloying Fe–Ni powders, Hays et al. [5] have shown that using FMP, the crystallite size decreases gradually as a function of the milling time with a particle size varying between 1 lm and 5 lm, whereas, in SMP, the crystallite size decrease sharply versus the milling time with a particle size varying between 20 lm and 100 lm. In Fig. 5, it is found that the as-milled Fe(Ni) alloy exhibits a typical microstructure of mechanically milled particles, which has a granular shape, rough surface, and a relatively broad distribution in particle size. The particles tend to agglomerate in the MA process because the grinding function in the stainless steel ball plays a leading role during the mechanical alloying, which leads to the formation of the granular shape.
4.2. Magnetic properties The increase of the magnetization is found to be linked to the crystallite size reduction to about 10 nm (see Fig. 6). Each grain may be considered as being a single magnetic domain eliminating the influence from magnetic walls [29,16]. On the other hand, the larger the lattice parameter, the higher the saturation magnetization. It has been shown by Amils et al. [6] that the ball milling pro-
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cess produces a high density of defects, particularly defects of antisite type, which causes a 0.8% lattice parameter expansion in the case of the Fe–Al alloys. The same authors [6] have previously reported on a good correlation between the saturation magnetization and the lattice parameter (increase of the lattice parameters and the saturation magnetization), which suggests that the observed magnetic transition may be partially related to the changes in the density of states at the Fermi level. Based on our results, it is obvious that the magnetization can be modified by the crystallite size and by the lattice parameter but not by the used mode (SMP, FMP). The increase of the coercivity in SMP can be understood as being the consequence of a considerable introduction of high internal strains into the material, which is inevitably related to the process of preparation (see Fig. 7b). Thus, the magnetostriction in combination with the high internal strain has been identified as the effect dominating the coercivity by the way of the magnetoelastic interactions [30]. The low values of the coercivity in FMP can be due to the low values of the crystallite size (in Fig. 7) where, the higher the milling time is, the lower the crystallite size. Also to the dependence of the coercivity and the crystallite size (the lower the coercivity is, the lower the crystallite size) prevails over the other predominant strain influence in this nanocrystallite size range [16]. Nevertheless, the resulting nanocrystalline structure led to the soft magnetic properties (low coercivity with low grain size) as suggested by Herzer [31]. However, the smaller the structural correlation lengths, the lower are the coercivity. Such a feature has been previously noticed in the case of amorphous alloys and of nanocrystalline alloys for grain size lower than 20 nm [31]. 5. Conclusions Nanostructured Fe-based soft magnetic could show some major technological advantages due to their predicted low coercivity and high saturation magnetization. In order to obtain such materials, it is necessary to produce Fe-based alloys, with an average crystallite size largely below 20 nm and most probably even below 10 nm. Mechanical alloying is one of the techniques that can synthesize such materials. The results presented in this work by using the PM 400 from Retsch and the vario mill (P4) planetary ball mill from Fritsch clearly shown great potential of this technique. High-energy ball milling of Fe and Ni powders results in the solid solution formation accompanied by a grain refinement for nanocrystalline Fe–10%Ni. The higher the shock power and milling time, the larger the lattice parameter (a = 0.28695 ± 0.00005 nm and 0.28704 ± 0.00005 nm for 8.83 W of the shock power and 96 h of the milling time, respectively) and the lower the crystallite size (10 ± 2 and 10 ± 2 nm for 8.83 W of the shock power and 96 h of the milling time, respectively). It has been found that the magnetization is not modified by milling performed in shock mode process (SMP) in friction mode process (FMP). Such a magnetization is only modified by the decrease of the crystallite size and by the increase of the lattice parameter. The highest values of coercivity have been found in SMP, i.e., 1600 A/m for Fe–10 wt.%Ni for the following milling condition (424 rpm/100 rpm). Following the classification of the magnetic materials (hard magnetic, memory materials and soft
magnetic materials), the FMP leads to the formation of alloys exhibiting a soft magnetic behavior (Hc between 380 A/m and 750 A/m) for Fe–10 wt.%Ni alloys and the SMP leads to the formation of alloys exhibiting a memory and hard magnetic behavior (Hc = 900–1600 A/m) for Fe–10 wt.%Ni alloys. Acknowledgments The authors wish to thank Dr E. GAFFET for his assistance concerning the values which are based on the kinematic studies of the G 5 ball mill. Also, the authors wish to thank Mr DE OLIVEIRA from FRITSCH Company and Mr VAXELAIRE from RETSCH Company which have given us more information concerning their machines (FRITSCH P4 and RETSCH PM400). References [1] C. Suryanarayana, Prog. Mater. Sci. 46 (2001) 1–184. [2] S. El-Eskandarany, Noyes Publications, William Andrew Publishing, Norwich, New York, USA, 2001, pp. 1–232. [3] E. Gaffet, G. Le Caër, in:Mechanical Processing for Nanomaterials, H.S. Nalwa (Ed.) Encyclopedia of Nanoscience and Nanotechnology, American Scientific Publishers, vol. 5 (2004), 91–129. [4] C. Kuhrt, L. Schultz, J. Appl. Phys. 75 (1993) 1975–1980. [5] V. Hays, R. Marchand, G. Saindrenan, E. Gaffet, Nanostruct. Mater. 7 (1996) 411–420. [6] X. Amils, J. Nogues, S. Surinach, M.D. Bar´o, J. Magn. Magn. Mater. 203 (1999) 129–131. [7] E. Jartych, J. Magn. Magn. Mater. 265 (2003) 176–188. [8] Yugo Nomura, Kazuo Fujiwara, Akihiko Terada, Satoshi Nakai, Masaaki Hosomi, Chemosphere 86 (2012) 228–234. [9] E. Gaffet, Mater. Sci. Eng. A 132 (1991) 181–193. [10] G.K. Rane, U. Welzel, E.J. Mittemeijer, Acta Mater V60 (2012) 7011– 7023. [11] N.E. Fenineche, R. Hamzaoui, O. Elkedim, Mater. Lett. 4493 (2003) 1–5. [12] S. Tria, O. Elkedim, R. Hamzaoui, X. Guo, F. Bernard, N. Millot, O. Rapaud, Powder Technol. 210 (2011) 181–188. [13] L.W. Huang, O. Elkedim, R. Hamzaoui, J.Allo.Comp. 509S (2011) 328–333. [14] Kh. Gheisari, S. Javadpour, J.T. Oh, M. Ghaffari, J.Allo.Comp. 20 (2009) 416–420. [15] S.D. Kaloshkin, V.V. Tcherdyntsev, Yu.V. Baldokhin, I.A. Tomilin, E.V. Shelekhov, J. Non-Cryst. Sol. 287 (2001) 329–333. [16] R. Hamzaoui, O. Elkedim, N. Fenineche, E. Gaffet, J. Craven, Mater. Sci. Eng. A 360 (2003) 299–305. [17] N.E. Fenineche, O. El Kedim, E. Gaffet, J. Metast, Nanocryst. Mater. 7 (2000) 41– 48. [18] R. Hamzaoui, O. Elkedim, J.Y. Rauch, M. Cherigui, D. Bassir, Recent Res. Devel. Mat. Sci. Engg. 4 (2011) 1–13. [19] Boqu Chen, Sha Yang, Xingxing Liu, Biao Yan, Lu Wei, J. Alloys Comp. 448 (2008) 234–237. [20] M. Abdellaoui, E. Gaffet, Acta. Metall Mater. 44 (1995) 1087–1098. [21] G.K. Williamson, W.H. Hall, Acta Metall (1) (1953) 22–31. [22] J. Yang, J.I. Goldstein, Lunar Planet. Sci. XXXIV (2003) 1156. [23] R.Hamzaoui, PHD Thesis, UTBM-UFC Belfort, France, 2004. [24] S.D. Kaloshkin, V.V. Tcherdyntsev, Yu.V. Baldokhin, I.A. Tomilin, E.V. Shelekhov, Physica B: Condensed Matter 299 (V) (2001) 236–241. [25] Yu.V. Baldokhin, V.V. Tcherdyntsev, S.D. Kaloshkin, G.A. Kochetov, Yu.A. Pustov, J. Magn. Magn. Mater. 203 (1999) 313–315. [26] V.V. Tcherdyntsev, S.D. Kaloshkin, I.A. Tomilin, E.V. Shelekhov, Yu.V. Baldokhin, Nanostruct. Mater. 12 (1999) 139–142. [27] M. Pekala, D. Oleszak, E. Jartych, J.K. Zurawicz, Nanostruct. Mater. 11 (1999) 789–796. [28] M. Pekala, D. Oleszak, E. Jartych, J.K. Zurawicz, J. Non-Cryst. Sol. 250 (1999) 757–761. [29] E. Jartych, J.K. Zurawicz, D. Oleszak, M. Pekala, J. Magn. Magn. Mater. 208 (2000) 221–230. [30] C. Kuhrt, L. Schultz, J. Appl. Phys. 73 (1993) 6588–6590. [31] G. Herzer, Nanocrystalline soft magnetic alloys, in:K.H.J. Buschow (Ed.), Handbook of Magnetic. Materials, Elsevier, Amsterdam, 1997, Vol. 10, pp. 415–462.