BALL MILLING OF INTERMETALLIC PHASES IN THE Nb-Al SYSTEM

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BALL MILLING OF INTERMETALLIC PHASES IN THE Nb-Al SYSTEM

M. Oehring, R. Bormann

To cite this version:

M. Oehring, R. Bormann. BALL MILLING OF INTERMETALLIC PHASES IN THE Nb-Al SYS- TEM. Journal de Physique Colloques, 1990, 51 (C4), pp.C4-169-C4-174. �10.1051/jphyscol:1990420�.

�jpa-00230780�

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COLLOQUE DE PHYSIQUE

Colloque C4, suppl6ment au n014, Tome 51, 15 juillet 1990

BALL MILLING OF INTERMETALLIC PHASES IN THE Nb-A1 SYSTEM

M. OEHRING and R. BORMANN

Institut far Werkstofforschung, GKSS-Forschungszentrum Geesthacht GmbH, Max-Planck-Strasse, 0-2054 Geesthacht, F.R.G.

Abstract - A powder consisting mainly of the intermetallic A15 phase Nb3A1 and a small amount of the o phase %,A1 was milled in a planetary mill. By X-ray diffractometry it is observed that the A15 phase can almost completely be transformed into a supersaturated bcc @A1 solid solution. The results are discussed with respect to the mechanism of the transformation. In order to estimate the enthalpy stored by ball milling the average crystallite size, the root mean square strain and the long-range order parameter are determined from X-ray diffractograms. It is concluded that the transformation of the A15 phase mainly originates from a decrease of its long-range order which rai- ses its free energy above the one of the bcc solid solution. This result will also be discussed with respect to a transformation of an intermetallic compound into an amorphous structure.

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INTRODUCTION

Recently the formation of a supersaturated bcc e A l solution was observed for Nb-rich com- positions up to 40 at% A1 by mechanical alloying of the elemental powders /l/. The results can be understood by considering free energy curves calculated by the CALPHAD method show- ing that the free energy of the bcc phase is lower than that of the amorphous phase in this concentration range /l/. For A1 concentrations higher than 40 at%, however, the amorphous phase is formed because the free energy curves cross at about 47 at%.

The difference between the A15 Nb3A1 and the bcc phase amounts to 2 kJ/mole for 25 at% Al.

Because an enthalpy of some kJ/mole can easily be stored in ordered intermetallic phases by ball milling as shown by DSc measurements /2,3,4/ we investigated the behaviour of the A15 phase under ball milling conditions. In this paper we report on the transformation of the A15 phase into the bcc structure and how this intermetallic phase is destabilized, e.g.

in which form its entha1py is raised with respect to the bcc phase. In particular we stu- died the chemical disordering of the A15 phase by X-ray diffraction for which NbAl intermetallic compounds are well suited due to the large difference in the atomic scattering fac- tors. By X-ray diffraction also the crystallite size and the root mean square (rms) strain can be obtained enabling one to estimate the enthalpy stored by grain boundaries and by the internal stress, both in the A15 phase as well as in the bcc solid solution.

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EXPERIMENTAL METHODS

As starting material we used 2N Nb-25at%A1 powder supplied by Alfa Ventron with a particle size of less than 45 pm. The oxygen and iron content was determined by chemical analysis and energy dispersive X-ray analysis, respectively, and amounted (0.3f0.1) wt% 0 and (0.3*0.2) wt% Fe. The milling and the handling of the powder was done in a glove box in an

Article published online by EDP Sciences and available at http://dx.doi.org/10.1051/jphyscol:1990420

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C4-170 COLLOQUE DE PHYSIQUE

atmosphere of purified Ar. During all experiments the measured oxygen and water content of the Ar atmosphere could be held below 1.5 ppm. For the milling we used a planetary mill (Fritsch pulverisette 5) with hardened steel vials and balls. The milling parameter were a ball to powder weight ratio of 5 and 10 and a milling intensity of 5 and 10, respectively, in the first and second milling series. It is estimated that the temperature of the vials did not exceed 50°C during the milling. After 62.5h of milling an oxygen content of

(0.3f0.17) wt% and a Fe content of (0.6f0.1) wt% was measured.

X-ray diffractometry was carried out on a Seifert 0-20-diffractometer MZ IV with graphite monochromator. Cu K, and in some cases CO K, radiation was used.

Fig. 1

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X-ray diffractograms observed with Cu-Ka radiation on powders milled for different times.

52.5 h

J,L

12.5 h '4 .- f .. - & d k

2.5 h

>I

.-

.W U) Q

Y c

-

0.5 h

l I

t I 1

Nb3AI

1 I I I , , I I 1 1 1

Nb2AI

I 1 1 a s - r e c .

1

^h

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3

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RESULTS

X-ray diffractograms of specimens milled in the first series for 0.5 h, 2.5 h, 12.5 h and 62.5 h and of the as-received material are shown in Fig. 1. In the untreated material only diffraction lines of the A15-%,A1 and a-Nb,Al intermetallic phases are found. The volume fraction of the o phase is low, in agreement with the phase diagram. With increasing milling time the lines broaden and after 2.5 h of milling Bragg-peaks of the supersaturated bcc &A1 solid solution can be detected. On further milling the intensity of the bcc peaks increases whereas that of the A15 phase peaks decreases. After 62.5 h of milling only tra- ces of Nb3A1 and Nb,Al are found as can be seen in Fig. 2. Considering the low Nb,Al content of the initial material %,A1 seems to be more stable against a transformation into the bcc structure.

In Fig. 3 the A1 concentration of the %,A1 and bcc &A1 phase is given for different milling times. The A1 concentration was calculated from the lattice parameters using the reference curves from Kammerdiner and Luo /5/ and Bormann et al. /6/ disregarding changes of the lattice parameters caused by the milling or by structural disorder. The obtained A1 concentration of 23.2 at% in the A15 phase of the untreated material agrees well with the value expected from the phase diagram. For shortly milled Nb3A1 relatively precise lattice parameters can be determined because sufficiently narrow diffraction lines allow the Nel- son-Riley correction to be applied. After 62.5 h of milling the formed bcc &A1 solid solution has nearly the same A1 concentration as the initial overall A1 content. Some values in this figure have large error bars due to weak and broad lines. Nevertheless, the results indicate that the A1 concentration of the initially formed bcc phase might be lower than the value of the Nb3Al phase.

Fig. 2 - X-ray diffractogram of a powder ball milled for 62.5 h.

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COLLOQUE DE PHYSIQUE

a s - 0.5 2.5 12.5 6 2 5

rec.

t ( h )

Fig. 3

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A1 concentration of the %,A1 and the bcc phase vs. milling time as determined from the lattice parameters.

The rms strain <EZ>'/~ and the average crystallite size L were determined from the broadening of the diffraction lines. To separate these two effects the procedure described by Hellstern et al. /2/ was employed. It is based on the fact that broadening deriving from small crystallite size is constant in k space whereas broadening due to internal strain of the lattice is proportional to k. The diffractograms were corrected for K, and instrumental broadening which was determined using pure silicon powder as a reference material. The method is rather inaccurate /7/ but separating crystallite size and strain broadening is una- voidable as both effects play an important role.

The results obtained from this analysis are plotted in Fig. 4b). The crystallite size and the rms strain were determined from diffraction lines of the Nb3A1 phase for milling times up to 2.5 h and of the bcc phase for longer milling times. The crystallite size of Nb,A1 decreases continuously with time down to a value of 43 m, whereas the crystallite size of the transformed bcc phase is approximately 10 nm, independent of the milling time. The rms strain is enhanced with milling time, both in the %,AI and the bcc phase up to a maximum value of <EZ>'/~ = 8 10-3. Similar values were determined by Hellstern et al. /2/ after extended milling of the B2-A1Ru intermetallic compound (L = 5-7 nm and <E~>'/' = 30 10-3) evaluating the line broadening in the same manner. Additionally, the crystallite size was confirmed in this case by HREM. Cocco et al. /8/ applied the more elaborate Warren-Averbach analysis to the line broadening of pure Ni powder resulting in a crystallite size of 16.5 nm and a strain of 1.2.10-3 after 25 h of ball milling.

From intensity ratios the relative degree of long-range order

was determined with Is, and If, being the intensities of superlattice and fundamental lines measured on the unmilled material and If and Is the intensities of the respective lines

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.

- - a - - - .

l I 1

a s - 0.5 2.5 12.5 6 2.5

r e c .

Fig. 4 - a) Relative degree of order S/So vs. milling time.

Fig. 4 - b) Crystallite size L and rms strain < E Z > ' / ~ as determined from the line broadening vs. milling time.

after milling. Due to their sufficient intensity the superlattice lines (110) and (220) and the fundamental line (211) were taken. In addition, the fundamental lines (200) and (400) could be considered for milling times up to 2.5h. The obtained relative degree of order could be affected by preferred orientations and a variation of the temperature factor with milling time. However, the latter does not change the intensity ratios drastically in Nb3A1 /9/ and texturing of the powder was not observed. From an overview by Fliikiger /9/ on measurements of the absolute order parameter it can be concluded that Sois about 0.97 for

%,A1 prepared by arc melting and crushing to a powder. As shown in Fig. 4a) the relative degree of order decreases with milling time reaching a value of approximately 0.8. A similar decrease to 0.7 was found by Hellstern et al. /2/ after extended milling of the intermetallic compound AlRu.

In the second milling series the milling intensity of our planetary mill was raised to 10.

The material transformed into the bcc phase, too, but on a shorter time scale. It was observed that a milling time of 1 h in the second series corresponds to 12.5 h in the first.

4 - DISCUSSION

The X-ray diffractograms show clearly that the A15 phase can be transformed by ball milling into a supersaturated bcc E A l solid solution which also forms by mechanical alloying of

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C4-174 _COLLOQUEDE PHYSIQUE

the pure elements in the appropriate ratio /l/. The results are in agreement with the free energy curves which exhibit a greater stability of the bcc solid solution than of the amorphous phase for a concentration of Nb-25at%A1 /l/. The NbzA1 phase which is also present in the as- received material seems to be more stable in agreement with its higher thermodyna- mical stability /l/ with respect to the bcc phase. Small amounts of the bcc solid solution could be detected already after 2.5 h of milling. Its volume fraction continuously increases while the amount of the A 15 phase decreases continuously. This behaviour could be ex- plained on the one hand by a nucleation and growth mechanism of the transformed phase, on the other hand it could be possible that the phase transformation occurs without a nucleation barrier. The latter has been proposed by Johnson/lO/ and would require an increase of the free energy of the intermetallic compound up to a critical value where the A15 phase becomes unstable. The experimentally observed continuous increase of the volume fraction of the bcc phase on further milling may then be due to a inhomogeneous distribution of defects in the intermetallic compound changing with the milling time.

The free energy required for the phase transformation can in principle be stored in the A15 phase by internal stresses due to a high dislocation density, by grain boundaries and by chemical disordering. The transformation of the A15 phase into the bcc solid solution is first observed after 2.5h of milling. At that stage the A15 phase has a mean crystallite size of 43 nm corresponding to a stored energy of 700 J/mole for a grain boundary energy of lJ/mZ. The rms strain amounts E: = 3 - 10-3 giving an elastic energy 1/2 E E ~ = 85 J/mole by taking into account an elasticity modulus of E = 180 GPa, a typical value for intermetallic compounds. The sum of these energies is lower than lkJ/mol and hence not sufficient for the phase transformation. Additionally, the transformation product has a smaller crystallite size and a larger strain than Nb3A1. Therefore we conclude that the free energy of the A15 phase must be raised by chemical disordering. Luzzi and Meshii /11/ have observed that the ordering enthalpy of an intermetallic phase typically is 21 -37 % of its heat of formation which amounts to 20 kJ/mole for Nb,Al. This indicates that an enthalpy change of 2 kJ/mole is possible by a decrease of the order parameter to 0.8, observed in the milled powder. In addition, it is postulated from these results that also the transformation of intermetallic compounds into an amorphous phase mainly originates from the chemical disorder introduced by the milling process and that the energy and a decreased grain size is usually not sufficient to destabilize the intermetallic compound with respect to the amorphous phase. This conclusion indicates that an intermetallic compound can only be amorphized if the amorphous phase has a lower free energy than a (disordered) solid solution. Further studies will aim for a test of this assumption.

We thank H. Bosse, U. Lorenz, G. Miillauer and C. Michaelsen for valuable discussions and experimental support and F. Schmelzer for chemical and EDX analyses.

REFERENCES

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(1988) 1399 /2/ E. Hellstern, H. J. Fecht, 2 . Fu, W. L. Johnson, J. Appl. Phys.

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/4/ R. B. Schwarz, R. R. Petrich, Proc. of the Conf. on Solid State Amorphizing

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/5/ L. Kammerdiner, H. L. Luo, J. Appl. Phys. 43 (1972) 4728

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/7/ A. Guinier, X-Ray Diffraction. Freeman, San Francisco, 1963

/8/ G. Cocco, S. Enzo, L. Schiffini, L. Battezzati, Proc. 6th Int. Conf. on Rapidly Quenched Metals, ed. R. W. Chochrane, J. 0. Strdm-Olsen, Mater. Sci. Engg.

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/9/ R. Fliikiger, Habilitation thesis. University of Geneva, 1987 /10/ W. L. Johnson, Progr. Mat. Sci.

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/11/ D. E. Luzzi, M. Meshii, Res. Mechanica 21 (1987) 207