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Materials Chemistry and Physics
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Solid state amorphization transformation in the mechanically alloyed Fe 27.9 Nb 2.2 B 69.9 powders
S. Alleg
a,∗, A. Hamouda
a, S. Azzaza
a, R. Bensalem
a, J.J. Su ˜ nol
b, J.M. Greneche
caLaboratoire de Magnétisme et Spectroscopie des Solides, Département de Physique, Université de Annaba B.P. 12, Annaba 23000, Algeria
bDep. de Fisica, Universitat de Girona, Campus Montilivi, Girona 17071, Spain
cLaboratoire de Physique de l’Etat Condensé, CNRS UMR 6087, Université du Maine, 72085 Le Mans Cedex 9, France
a r t i c l e i n f o
Article history:
Received 6 November 2009
Received in revised form 15 February 2010 Accepted 5 March 2010
Keywords:
Nanostructures Powder metallurgy X-ray diffraction Mössbauer effect
a b s t r a c t
Mössbauer spectrometry and Rietveld analysis of X-ray diffraction patterns were used to follow the solid state amorphization transformation during the milling process of the Fe27.9Nb2.2B69.9powders. The reaction between elemental Fe, Nb and B powders leads to the formation of the Nb(B) and Fe(B) solid solutions after 1 and 10 h of milling, respectively. A mixture of␣-Fe, Nb(B) and highly disordered Fe(Nb, B) solid solution is found after 25 h of milling. An amorphous structure is obtained on further milling time (100 h). From the Mössbauer spectrometry results, it is observed that the total mixing of the elemental powders, at the atomic level, is achieved after 50 h of milling and a stationary state corresponding to a full paramagnetic amorphous phase is reached after 100 h of milling. The amorphization process can be described by an Avrami parameter close ton= 1.
© 2010 Elsevier B.V. All rights reserved.
1. Introduction
Nanocrystalline (NC) materials are interesting for several practi- cal applications because of their improved physical, mechanical and magnetic properties. NC alloys can be obtained by different meth- ods starting from liquid, solid or vapour states. Mechanical alloying (MA) process is an effective way to produce metastable microstruc- tures such as amorphous, NC and supersaturated solid solutions from pure elemental or compound powders[1–6]. The repeated fracturing and cold welding, during the milling process, lead to the destabilisation of the crystalline structure through the accumula- tion of structural defects such as vacancies, dislocations and grain boundaries and hence, to the formation of NC and/or amorphous structures[2,7].
NC Fe–Nb–B alloys (“Nanoperm”), which are attractive for their soft magnetic properties in the NC state after subsequent thermal treatment from the amorphous precursor, are used extensively in commercial applications such as micro-devices, telecommunica- tions, power electronics[8,9]. The increase of the Nb and B contents in the Fe-based amorphous alloys plays a role in the thermal sta- bility, as does B in the generation of a strong magnetic coupling, of the amorphous phase[10]. Depending on both the starting com- position and the milling conditions, ball milling of the Fe–Nb–B powder mixtures leads to the formation of either an amorphous
∗Corresponding author. Tel.: +213 38 87 53 99; fax: +213 38 87 53 99.
E-mail address:safia alleg@yahoo.fr(S. Alleg).
phase or NC metastable supersaturated bcc Fe(Nb, B) solid solu- tion[11–13]. Thus, MA seems to be suitable for the synthesis of supersaturated solid solutions in binary Fe–B and Fe–Nb as well as in ternary Fe–Nb–B alloys since the solubility of both B and Nb into Fe is very small (<1 at.%) under the equilibrium conditions. It has been reported that the solid solubility limit increases with ball milling but does not exceed 4 and 5% for B and Nb in Fe, respectively [2]. Hasegawa and Ray pointed out that Fe-based alloys with less than 12 at.% B are preferred to form an Fe(B) metastable supersat- urated solid solution with a bcc structure, while those with higher B concentration (>12 at.%) to form an amorphous phase[14]. Those features suggest that the high B content (30 at.%) enhances the amorphization process after segregation of B at grain boundaries.
In the present work, the solid state amorphization during the milling process of the Fe27.9Nb2.2B69.9powder mixtures was inves- tigated by X-ray diffraction, scanning electron microscopy and Mössbauer spectrometry.
2. Experimental details
Pure elemental powders of Fe (6–8m, 99.7%), Nb (74m, 99.85%) and amor- phous B (>99%) were mixed to give a nominal composition of Fe27.9Nb2.2B69.9(at.%).
The milling process was performed in a planetary ball mill Retsch PM400/2 under argon atmosphere using hardened steel vials and balls (10 balls with a diameter of 8 mm). The ball-to-powder weight ratio was about 8:1 and the rotation speed was 350 rpm. In order to avoid the increase of the temperature inside the vials, the milling process was interrupted each 1/2 h for 1/4 h. Morphological changes of the powder particles during the milling process were followed by scanning electron microscopy (SEM) in DSM960A Zeiss equipment. The phase identification, microstructural and structural evolutions were investigated by X-ray diffraction (XRD) by means of a Bruker D8 Advance diffractometer in a (–2) Bragg Brentano geometry using 0254-0584/$ – see front matter© 2010 Elsevier B.V. All rights reserved.
doi:10.1016/j.matchemphys.2010.03.006
Fig. 1.SEM micrographs of un-milled (0 h) and powders milled for 1, 16 and 100 h.
Cu–K␣radiation (Cu= 0.15406 nm). The structural parameters were obtained from the Rietveld refinement of the XRD patterns by using the MAUD program[15]. The local57Fe environment was studied by Mössbauer spectrometry in transmission geometry with a conventional constant acceleration spectrometer, at room temper- ature, using a57Co source diffused into a Rh matrix. The Mössbauer spectra were computer fitted with a least-square iteration Mosfit program[16]. The isomer shift, IS, values are given with respect to␣-Fe at 300 K.
3. Results and discussions 3.1. Evolution of particle morphology
Fig. 1shows the morphologies of the Fe27.9Nb2.2B69.9 powder particles before (0 h) and after different ball milling durations. The changes in morphology during the milling process are due to the competition between fracturing, cold welding, agglomeration and de-agglomeration of the powder particles. Since the powders are soft at the early stage of milling, they tend to agglomerate by cold welding and form bigger particles of about 40m (1 h). At the intermediate stage of milling (16 h), the agglomerated powder particles are subject to continuous disintegration with fragmenta- tion to form relatively fine powders with size less than 6m in diameter. The fracture is the main process at this stage of milling.
Further milling time (100 h) leads to roughly spherical agglomerate particles with a distribution of size in the range (0.1–5m).
3.2. Structural changes
The induced crystal defects such as dislocations, grain bound- aries vacancies and interstitials during the milling process, through
the heavy plastic deformation into the powder particles, promote the solid state reaction at ambient temperature[2]. Depending on the initial mixture, structural changes of mechanically alloyed pow- ders can occur as follows: grain refinement, solid solution diffusion and/or formation of new phases[17].
Fig. 2displays the XRD patterns of the Fe27.9Nb2.2B69.9powders milled for various times. The broadening of the diffraction peaks and the decrease of their intensities with increasing milling time
Fig. 2.XRD patterns of the Fe27.9Nb2.2B69.9powders as a function of milling time.
Fig. 3.Rietveld refinement of the XRD pattern of the powders milled for 1 h.
can be due to the crystallite size refinement down to the nanometer scale and the increase of the internal level strain. Furthermore, the appearance of some Bragg peaks is assigned to the formation of new phases.
The XRD pattern of the un-milled powder mixture shows the diffraction peaks of bcc ␣-Fe and bcc Nb. The B peaks are not detected owing to its low atomic scattering factor and amorphous state. The appearance, after 1 h of milling, of new diffraction peaks on the lower angle side of Nb ones is due to the formation of the bcc Nb(B) solid solution with a lattice parameter close to a= 0.3429(4) nm and a mean crystallite size of about 37 nm (Fig. 3).
The reaction between B and Nb, in the early stage of milling, can be related to the negative enthalpy of mixing (−39 kJ/mol)[18]. It has been reported that ball milling of Nb and amorphous B leads to the splitting of Nb peaks and the formation of the Nb(B) solid solution [19]. The reaction between Fe and B gives rise to the formation, after 10 h of milling, of the FeB boride (Fig. 4). A mixture of␣-Fe, Nb(B) and highly disordered Fe(Nb, B) solid solution[20]is obtained after 25 h of milling (Fig. 5). As the milling process progresses, the diffraction peaks of␣-Fe and Nb(B) solid solution disappear while those of the Fe(Nb, B) solid solution show considerable broaden- ing indicating the complete amorphization after 100 h of milling as evidenced by the broad diffuse halo centred at 2≈44◦(Fig. 6).
This is also confirmed by the evolution of the phase proportions as a function of milling time (Fig. 7).
The lattice constant of Nb increases significantly from a0= 0.3303(4) nm toa= 0.3429(4) nm after 1 h of milling. The rel- ative deviation of the lattice parameter from that of the perfect
Fig. 4.Rietveld refinement of the XRD pattern of the powders milled for 10 h.
Fig. 5.Rietveld refinement of the XRD pattern of the powders milled for 25 h.
Fig. 6.XRD patterns of the powders milled for 100 and 125 h.
crystal, which is defined bya= (a−a0)/a0, reaches as much as a= 3.8%. The lattice parameter of the␣-Fe nanophase is enhanced by a= 0.20% with a crystallite size of about 10 nm after 25 h of milling. The lattice distortion which is characterized by lattice expansion or contraction and increased atomic displacement can be attributed to the supersaturation of point defects or vacancies inside the nanometer crystallites due to their higher energetic solu- tion[21]. The solution of point defects in the crystal lattice will
Fig. 7.Evolution of the deduced phase proportions, from the Rietveld refinement, as a function of milling time.
Fig. 8.Variation of the Mössbauer spectra, taken at 300 K, of the Fe27.9Nb2.2B69.9
powders against milling time.
disturb the lattice structure around the vacancies, and hence results in a distorted crystal lattice. The smaller the crystallite size, the higher the solubility of vacancies and, consequently, the lattice dis- tortion is more significant. Different thermodynamic states lead to different lattice structure characteristics. It is evident that a reduc- tion in crystallite size enlarges the free energy of the crystallites and then raises the equilibrium solute solubility in the crystal lattice [22]. Crystal lattice distortion in nanophases seems to be a gen- eral feature because similar lattice distortion behaviour has been reported in other systems[3,23].
3.3. Local Fe environment
57Fe Mössbauer spectrometry is a non-destructive and effective tool to probe the Fe environment through the quantitative eval- uation of the local structure and composition changes during the milling process[24]. The room temperature Mössbauer spectra of the powders milled for selected times are shown inFig. 8. The mix- ing of the elemental powders, at the atomic level, is evidenced by the line shape changes of the Mössbauer spectra with increasing milling time. The Mössbauer spectrum of the powder milled for 1 h is fitted with two components: a magnetic sextet with sharp lines and a central paramagnetic doublet. The magnetic sextet with a hyperfine magnetic field,B= 32.8 T, isomer shift IS = 0 mm/s and quadrupole shift, 2ε= 0.02 mm/s is related to the pure␣-Fe. The paramagnetic doublet with IS = 0.18 mm/s, QS = 1.00 mm/s and a relative area of about 6% is assigned to the B-rich Fe(Nb, B) solid
are summarized in Table 1. The sextet is related to the NC␣- Fe phase, whereas the broad spectral component is assigned to a highly disordered Fe(Nb, B) structure where the Fe atoms are largely and diversely surrounded by non-magnetic Nb and B atoms.
The changes in the hyperfine parameters (Band IS) of the Fe(Nb, B) solid solution with increasing milling time can be related to the composition change. For example, the increase of IS from 0.22 to 0.15 mm/s after 10 and 25 h of milling, respectively, shows that electrons are transferred to the d-band of Fe with increasing boron concentration[26].
The hyperfine field distribution,P(B), of the powders milled for 10, 25 and 25 h (Fig. 10) can be roughly divided into three regions related to amorphous (B< 20 T), interface (20 <B< 32 T), and NC (B= 33 T) contributions[27]. The observed low hyperfine magnetic field component at 18.6 T with a relative area of about 8%, after 10 h of milling, can be ascribed to Fe atoms in environment close to FeB boride nanophase, in agreement with the presence of the FeB boride in XRD pattern of the powders milled for 10 h (Fig. 4).
In the NC alloys, the important contribution of the interface which is about 14–20% of the total Fe can be related to the small crys- tallite size (∼5 nm)[28]. The interface thickness can be estimated by the ratio between the Fe content in the pure NC phase and the Fe percentage in the combined NC and interface[29]. Hence, the estimated interface thickness of the milled powders for 10, 25 and 50 h is about 0.6, 0.8 and 0.9 nm, respectively. The increase of the interface thickness with increasing milling time can be related to more crystallite size refinement.
Further milling time gives rise to a stationary state correspond- ing to a full paramagnetic amorphous phase as evidenced by the broad paramagnetic doublet in the Mössbauer spectrum (Fig. 11) which may be attributed to a weakly magnetic B-rich Fe(Nb, B) solid solution. This can be assigned to the effects of plastic defor- mation during the milling procedure and therefore, the reduction of the domain size leading to a disordered structure and the enrich- ment in B and Nb of the matrix. So, the lattice defects introduced during MA lead to a wide variety of the number, spacing and angu- lar distributions of near neighbours around Fe atoms. Evidently, the progressive fragmentation of nanosized particles and then
Table 1
Hyperfine magnetic fieldB, isomer shift IS, quadrupole splitting QS, quadrupole shift2ε, line width, magnetic momentand relative area of the powders milled for 10, 25 and 50 h.
Milling time (h) Phase IS (mm/s)±0.02 QS or 2ε(mm/s)±0.02 (mm/s)±0.02 B(T)±0.2 (B) Relative area (%)±1
10 ␣-Fe 0.00 0.00 0.34 33.1 2.20 54
Fe(Nb, B) 0.15 0.27 0.34 18.2 1.20 25
Doublet 0.16 0.74 0.80 – – 21
25 ␣-Fe 0.00 0.02 0.32 33.1 2.20 11
Fe(Nb, B) 0.22 −0.26 0.34 16.3 1.10 18
Doublet 0.20 0.66 0.74 – – 71
50 ␣-Fe 0.08 −0.02 0.34 33.6 2.24 5
Fe(Nb, B) 0.01 0.26 0.34 18.2 1.20 17
Doublet 0.21 0.66 0.73 – – 78
Fig. 9.Fitted Mössbauer spectra of the Fe27.9Nb2.2B69.9powders milled for 10, 25 and 50 h.
the break down of magnetic ordering are responsible of the exis- tence of paramagnetic phase[30]. Hence, the electric quadrupolar interactions should be treated by the use of quadrupole splitting distribution with a linear correlation with IS. The obtained hyper- fine parameters are summarized inTable 2. The distribution curve P(QS) (see inset ofFig. 11) shows a broad peak situated in the range 0.2–1.9 mm/s reflecting the different crystalline symmetries as obtained in amorphous alloys prepared by ball milling[31].
Fig. 10.Hyperfine field distribution curves of the powders milled for 10, 25 and 50 h.
Fig. 11.Room temperature Mössbauer spectrum of the Fe27.9Nb2.2B69.9powders milled for 100 h. The corresponding quadrupole splitting distribution curve is shown in the inset.
The alloying process can be also followed by the evolution of both the average hyperfine magnetic field [B] and average iso- mer shift [IS], as a function of milling time (Fig. 12). The average hyperfine magnetic field decreases with increasing milling time to a minimum value of [B] = 0 T. Simultaneously, the average isomer shift increases with increasing milling time, reaches a maximum value of about 0.27 mm/s after 50 h of milling, then decreases to about 0.21 mm/s after 125 h of milling. The decrease of [B] can be attributed to the presence of B and Nb atoms in the vicinity of Fe ones leading to the reduction of the Fe magnetic moment and there- fore, to the decrease of the hyperfine magnetic field. The increase of [IS] can be ascribed to the diffusion of B atoms into bcc␣-Fe and consequently, to the Fe(B) solid solution. However, the slight decrease of [IS] can be related to the presence of Nb atoms in the Fe(B) solid solution. Since Nb atoms reduce the isomer shift, one can conclude that the effect of Nb is masked by that of B. There-
Table 2
Isomer shift IS, quadrupole splitting QS, line widthand relative area of the powders milled for 100 h.
Subspectra IS (mm/s)±0.02 (mm/s)±0.02 QS (mm/s)±0.02 Relative area (%)±1
1 0.31 0.28 1.91 3.0
2 0.23 0.28 1.27 12.0
3 0.23 0.30 0.86 32.0
4 0.25 0.28 0.55 31.0
5 0.23 0.28 0.21 22.0
Fig. 12.Evolution of the average hyperfine magnetic field [B] and average isomer shift [IS], during the milling process.
Fig. 13.JMA plot versus milling time.
fore, the amorphous matrix can be considered as B-rich Fe(Nb, B) phase.
The amorphization kinetics of the mechanically alloyed Fe27.9Nb2.2B69.9powders can be deduced from the Mössbauer spec- trometry results through the transformed fraction of ␣-Fe as a function of milling time. Since the milling process occurs at room temperature, one can suppose that the temperature is constant. In addition, the milling time can be considered as the necessary time for phase transformation. Consequently, the amorphization kinet- ics can be described by the Johnson-Mehl-Avrami formalism[32]
in which the fraction transformed exhibits a time dependence of the form:
X=1−exp [−(kt)n]
wherenis the order of reaction or Avrami parameter,Xis the vol- ume of transformed fraction,tis the milling time andkis the rate constant. The kinetics parameters can be obtained from the double logarithmic plot ln(−ln(1−X)) versus lnt (Fig. 13). The amor- phization process can be described by one stage with an Avrami parameter of aboutn= 1. This value is comparable to those obtained for transformations controlled by the diffusion at the interface and dislocations segregation. This might be correlated to the existence of a high density of dislocations and various types of defects as well as to the crystallite size refinement. Comparable values of the Avrami parameter were obtained for the Mo dissolution into the
Acknowledgements
The authors are grateful to Xavier Fontrodona from the Physics Department of the Girona University, Spain for the XRD measure- ments. One of the authors (Prof. Safia Alleg) is grateful to the University of Girona for financial support during her visit as invited Professor in the Physics Department.
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