MICROSTRUCTURAL STUDY OF THE MECHANICALLY ALLOYED AMORPHOUS Fe-8Nb-30B POWDER MIXTURES A. Hamouda

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International Conference on NDT and Metal Industry (IC-NDT-MI’08) Annaba 12-14 May 2008

MICROSTRUCTURAL STUDY OF THE MECHANICALLY ALLOYED AMORPHOUS Fe-8Nb-30B POWDER MIXTURES

A. Hamouda¹*, S. Alleg¹, R. Bensalem¹ and J. J. Sunõl²

1 Laboratoire de Magnétisme et de Spectroscopie des Solides (LM2S), Département de Physique, Faculté des Sciences, Université de Annaba

B. P. 12 (23000) Annaba-Algérie

2 Dep. Fisica, Universitat de Girona, Campus Montilivi, Girona 17017 Spain

*E-mail : hamouda_assia@yahoo.fr

ABSTRACT: In the metallurgical process of mechanical alloying (MA), powder particles are subjected to severe mechanical deformation from collisions with steel balls. High energy milling forces can be obtained by using high frequencies. During the continuous deformation, important microstructural changes such as grain refinement down to the nanometer scale and amorphization occur. Using X-ray diffraction (XRD), structural and microstructural changes of the mechanically alloyed Fe-8Nb-30B powder mixtures were investigated. Different deformation processes were observed. It was found that structural defects such as dislocations, grain boundaries (GB), vacancies and interstitials, play an important role in the nanostructure formation and amorphization processes.

As a result, increase in some physical parameters such as dislocation density is similar to that obtained during severe deformation processes.

1. INTRODUCTION

Mechanical alloying (MA) is a powder metallurgy processing technique involving repeated welding fracturing and rewelding of powder particles in a high energy ball mill. MA has been shown to be capable of synthesizing novel materials, such as amorphous alloys and nanomaterials. It results in severe mechanical deformation at high strain rates (10³- 10⁴ s^-1) similar to the deformation of railway tracks due to the pressure of the high speed trains. The process of MA consists of loading metal powders and hardened steel balls in a stainless steel container sealed under argon atmosphere (to avoid atmosphere contamination) for a desired length of time. The extreme deformation leads to different structural changes. Amorphization is one of the most frequently reported phenomena in mechanically alloyed powder mixtures. It has been reported that both Fe-B and Fe-Nb binary systems lead to amorphization under MA. The mechanism of α-Fe mixing with Nb at the atomic level show, (by scanning electron microscopy) [1], that the repeated fracture and cold welding of powder particles during milling generates a multilayer structure due to the ductility of the powders, this structure lead to amorphization. The mechanism of the amorphization of Fe-B is different [2]. The α-Fe alloying with B atoms of small radius leads to the formation of an amorphous phase in interfaces of the α-Fe nanostructure. The segregation of B atoms at the grain boundaries is the source for the amorphous phase formation. Amorphization is a solid state reaction which occurs by destabilization of the crystalline structure after accumulation of structural defects during the MA process [3]. In order to understand the amorphization mechanism related to the deformation processes and the role of defects such as GB, vacancies, interstitials and dislocations, XRD study of Fe-8Nb-30B powder mixtures prepared by ball milling were investigated.

2. EXPERIMENTAL

A mixture of elemental crystalline Fe, Nb and amorphous B powders with a purity of 99.9%

was performed in a high energy planetary ball mill Retsch PM 400 under argon atmosphere. The milling intensity was 350 rpm and the ball-to-powder weight ratio was about 41/5. The progress of milling was followed by XRD using a Brucker D8 Advance diffractometer with Cu Kα radiation (λ = 0.154056 nm). The MAUD program based on the Rietveld method combined with a Fourier analysis was used to analyse the XRD patterns with broadened peaks. Microstructure parameters such as crystallite size, residual micro strains, lattice parameters and phase proportions are obtained. The

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accuracy of the results is related to the value of the quality factor of fitting, goodness of fit χ which is close to unity.

3. RESULTS AND DISCUSSION

The XRD patterns of the Fe-8Nb-30B powder mixtures were taken for different milling time as shown in Fig. 1. We can see a continuous decrease in the heights of the Bragg peaks of the elements and an emergence of extremely broadened peaks centred at about 45 and 77° (2θ). As the milling process progresses one observes an increase of the diffracting volume of the broadened peaks and the disappearance of the crystalline peaks. Those features suggest the mixing of the initial crystalline elements and the formation of a highly disordered FeNbB amorphous-like structure.

Fig. 1: XRD patterns of the ball milled powders mixture for different milling times.

The lattice parameter of mixed Fe is shown in (Fig. 2). No change of this parameter has been observed during the first 16 h of milling. it has a value close to the initial value 0.2867 nm and then it is found to increases slightly to reache a value of 0.2876 nm after 50 h of milling.

According to Hume-Rothery rules the substitution type is hard to be formed with the radius ratio of solvent to solute atoms beyond 15 %. Taking the Fe-Nb-B alloys into consideration, the empirical radius values of Fe, Nb and B are 0.117 nm, 0.142 nm and 0.087 nm, respectively.

Regarding the Fe as the solvent, the value of radius ratio of Fe over B and Fe over Nb, are 35 % and 15 %, respectively. Then it's deduced that Nb is solved in Fe to form the substitution solid solution, while B and Fe combine the interstitial type. The slight increase of the lattice parameter indicates then that a solid solution, between Fe and Nb and B, is formed.

0 10 20 30 40 50

0 ,28 4 0 ,28 5 0 ,28 6 0 ,28 7 0 ,28 8

Lattice parameter (nm)

M illin g tim e (h )

Fig.2: Evolution of the lattice parameter of α-Fe(Nb,B) during milling.

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After instrumental broadening corrections and separation of grain size effects from internal strain effects, the evolution of the average coherently diffracting domains (grain size), <L>, and the average atomic level strain, <σ²>^1/2,(Fig. 3) can be studied as function of milling time.

0 10 20 30 40 50

0 200 400 600 800 1000

Milling time (h)

<L> (nm)

0,0 0,1 0,2 0,3 0,4 0,5

<σ 2> 1/2(%)

Fig.3: Evolution of the grain size, <L>, and the internal strain, <σ²>^1/2, of α-Fe(Nb,B) during milling.

The reduction of the α-Fe(Nb,B) grains down to nanometer scale is due to the large plastic deformation and repeated fracturing induced during milling. High density of dislocations is also generated with prolonged milling, leading to an increase in internal strain of the powder particles. The microstructural changes occur through three stages:

(i) Initially the internal refinement process with a reduction of the average grain size by a factor of 1000 results from the plastic deformation localized in shear bands consisting of an array of dislocations which organize to form sub-grains. At a certain strain level 0.36 % the deformation is stopped. The minimum yield stress σ required to deform plastically the metal powders is related to the grain size, <L>, by the Hall-Petch relationship σ = σ₀+ k d

-1/2. When the grain size, <L>, decreases at a few nanometers during the milling process, the yield stress increases, the plastic deformation is then stopped, and the grain size reaches a critical value, <L^*>.

(ii) The continuous deformation lead to the continuous reduction of the grain size but the rate of the reduction is slow. A steady state deformation is observed and the grain size decreases down to the critical value by a superplastic deformation process.

(iii) Finally we observe the nucleation of the amorphous structure at the GB, probably due to the chemical disordering and mixing of the elements. The low solubility of Nb and B in Fe generates structural defects such as (vacancies and interstitials) which lead to lattice distortion and instability of the crystalline structure at the GB. Amorphization is probably a strong relaxation mode of accumulation of defects at GB [6]. The continuous deformation of the powders enhances the amorphization with the increase of the GB area.

In fact the phase proportion of α-Fe(Nb,B) decreases rapidly.

Figure 4 exhibits the dislocation density as function of milling time. Two different regimes can be clearly distinguished. The first one (i) is related to the first stage when the dislocations annihilate and recombine to small angle grain boundaries. The grain size reaches a critical value estimated at, <L^*> = 15 nm, after 16 h of milling, separately the dislocation density achieve a value of 0.65 10¹⁶ similar to the maximum density reached in the case of severely deformed metals (10¹⁶) [4].

In this regime the deformation is conducted essentially by the dislocations. The second regime is related to the second stage when the role of dislocations become negligible, the deformation occurs via

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slip of grains, the orientations of the grains become completely random with high angle grain boundaries. The grain size is reduced down to 15 nm and reaches a value of about 7 nm after 50 h of milling. Localized deformation proceeds by the dilatation of the GB similar to superplastic behaviour.

After 100 h of milling the continuous accumulation of defects leads to the transformation of the GB in an amorphous intergranular layer. The increase of the thickness of this layer leads to the destabilization of the crystalline structure.

0 10 20 30 40 50

0,4 0,5 0,6

0,7 (ii)

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Dislocations density (1016 / m2 )

Milling time (h)

Fig.4: Evolution of the dislocations density of α-Fe(Nb,B) as a function of milling time.

4. CONCLUSION

It is important to note that the different deformation mechanisms versus different structural defects result in different microstructural changes controlling the physical properties of these materials. These deformation processes are important for fundamental studies of extreme mechanical deformation, similar processes control the deformation of technologically specific wear surfaces for example the effects of work hardening material transfer and erosion during wear situations result in microstructures comparable to those observed during MA.

REFERENCES

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[3] C. Suryanarayana, E. Ivanov, V.V. Boldyrev, The science and technology of mechanical alloying, Materials Science and Engineering, A304-306, 2001, pp151-158.

[4] H.J. Fecht, Nanostructure formation by mechanical attrition, Nanostructured Materials, Vol 6, 1995, pp33-42.

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[6]

X. Wu, N. Tao, Y. Hong, J. Lu, K. Lu, Localized solid-state amorphization at grain boundaries in a nanocrystalline Al solid solution subjected to surface mechanical attrition, Journal of Physics D:

Applied Physics, 38, 2005, pp4140-4143.