Influence of milling conditions on the FeAl intermetallic formation by mechanical alloying

MATERIALS SCIENCE & EMtlNEERlMt Materials Scienceand EngineeringA207 (1996)97-104

Influence of milling conditions on the FeAl intermetallic formation by mechanical alloying K. Wolski”, G. Le CaErb, P. Delcroixb, R. Fillit”, F. Th6venota, J. Le Cozea “SMS,

Ecole des Mines de Saint bLSG2M, Ecole des Mines

Erieme, 158 cows Fawiel, de Nmey, Pare de Saurupt,

42023 Saint Etienne, France 54042 Nancy, France

Received17 April 1995;in revisedform 10 Aumst 1995

Abstract The formation of FeAl inter-metallic compound by mechanical alloying has been investigated as a function of milling time. Mixtures of elemental powders of Fe and Al are progressively transformed into a disordered solid solution characterized by an average composition of Fe-35at.%Al. Changes in powder morphology, the degree of reaction advancement, as well as crystallite size and microstress evolution have been described. It has been found that the iron crystallite size tends to about 15 nm and that FeAl grains of the same size are formed. Mijssbauer spectroscopy (MS) study has provided additional information about intermediate solid solutions, the final composition and the influence of the milling velocity on FeAl formation. No free iron has been detected in powders processed for 24 h at the highest rotation speed (400 rev mm-‘). MS appears to be a unique technique that can provide unequivocal information about the state of mixed powders, especially for long milling times. Keywords: Mechanical alloying; Milling conditions; FeAl intermetallic compounds; Nanostructures

1. Introduction Mechanical

alloying

(MA)

[l],

which

was initially

conceived for the production of dispersion strengthened superalloys [2], is nowadays used for synthesizing a wide range of materials including intermetallics [3,4]. MA is a solid state, dry milling process that leads, through microsandwich morphology, to the ultimate mixing of elemental powders and eventually to alloy formation. Additionally, the repetitive action of cold welding, fracture and rewelding of microsandwiches yields a homogeneous distribution of oxidation products. These oxidation products can come from intentional in situ oxidation from milling gaseous atmosphere or as surface oxides from elemental powders. MA appears to be one of the best methods for preparing oxide dispersion strengthened materials with homogeneous dispersions of ceramic particles. The purpose of the present paper is to report results of a systematic investigation of FeAl formation by MA process, using four different techniques: scanning electron microscopy (SEM), differential thermal analysis 0921-5093/96/$15.000 1996- ElsevierScienceS.A. All rights reserved

(DTA), X-ray diffraction (XRD) and Mijssbauer spectroscopy (MS). This is a part of a wider study [5] aimed at improving high temperature slow plastic flow properties of FeAl intermetallics reinforced by ceramic dispersions. (As high temperature properties were measured in constant velocity compression tests at 727 “C and initial strain rate within the range 10-4-10-6 s-l, it is preferable to speak about slow plastic flow properties rather then creep properties [5].)

2. Experimental 2.1. Powders

Commercially available, very fine though highly oxidised Fe and Al powders were used in this study (Table 1). For each milling time, powders were mixed in the same proportions (Fe- 15.13 g, Al-4.87 g), initially calculated to reach the final composition Fe-40at.%Al. The presence of oxygen in the starting powders and further oxidation during milling decreased the alu-

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Table 1 Some characteristics of elemental powders used in this study Powder Producer Diameter Oxygen content (wt.%) WI

Carbon content h-W

Certified Measured Certified Measured Fe Al

Prolabo Koch

10 8-15

0.2 1.0

0.5 5.0

200 -

140 400

minium content to Fe - xat.%Al with x depending on the milling time and ranging from 32 to 37 at.%Al. 2.2. Psocessing MA was carried out in a Fritsh planetary ball mill, using hardened steel containers with a capacity of 0.75 dm3. The rotation speed of the containers equals the rotation speed on the disc. The representative series of millings was done at 200 rev min - ’ for 1 min, 1, 2, 4, 8, 16 and 24 h and additionally at 300 and 400 rev min- ’ for 24 h. Each milling was realised according to the same procedure: (i) loading under air of 20 g of elemental Fe, Al powders in the required proportions together with 1 kg hardened steel balls (diameters 4.0 or 4.7 mm, lOOC6 steel, ball to powder weight ratio 50: 1); (ii) milling up to the required total milling time (discontinuous milling with stops every 15 min for approximately 10 min to allow for air cooling of the whole container); (iii) opening under air and sieving. Stearic acid was systematically added in order to avoid sticking of the powders to the balls and to the wall of the container. The quantity of stearic acid, optimised to 2 wt.% (0.4 g) on the basis of minimum welding, allowed recovery of more than 90% of the initially-introduced powders. 2.3. Methods of analysis The morphology of the resulting powders was characterized by SEM, equipped with a standard micfoanalyser (JEOL JSM 840 + TRACOR). The degree of reaction advancement has been analysed by means of a NETZSCH AT 429 DTA. An additional series of powders milled 1 min, 4, 10, 20 and 30 h at 200 rev min- ’ were used in this analysis. The internal structure of powder particles was analysed by XRD using a Philips diffractometer (1730/ 10) equipped with a DOSOPHATEX system [6] with Co Ka (,?.= 0.1789 nm) radiation. The evolution of mean crystallite sizes and average microstresses has been quantified by deconvolution of Fe or FeAl (110) diffraction peaks. This deconvolution was based on the Warren-Averbach method wherein the broadening due to decreasing crystallite size results in a Cauchy function shape whereas increasing microstresses gives a

Gauss function shape. These calculations were performed with the help of a computer program described in Ref. [7]. Additional information has been obtained from S7Fe room temperature MS in transmission geometry using a source of 57Co in a rhodium matrix (10 mCi). Hyperfine field distributions (HFDs) were calculated from Messbauer spectra with a constrained Hesse-Riibartsch method described in Ref. [8].

3. Results Every analysis of the present study was performed on several powders as a function of increasing milling time and has given access to a specific aspect of the solid state formation of FeAl intermetallic compound. The powders synthesized by MA were first analysed by SEM (Fig. 1). Initially, spherical (Fe, Fig. l(a)) or irregular (Al, Fig. l(b)) powders achieved a lamellar morphology (Fig. l(c)) after 1 h of milling. At this stage there was still no microsandwich formation and the use of microanalysis readily allows the distinction between two populations of particles (pure Fe or pure Al). Increasing the milling time up to 4 h resulted in small and slightly angular particles (Fig. l(d)). A further increase up to 16 and 24 h brought back the spherical morphology, with the average particle size going through a minimum at 16 h (Fig. l(e)); it is no longer possible at this stage to distinguish separate Fe or Al particles. Milling for 24 h led to a final morphology (Fig. l(f)) that did not change, even for increased velocity (400 rev min - ‘, Fig. l(g)). Final powders are characterized by an average composition of Fe37at.%Al+ 2at.% (a semiquantitative result based on ten SEM microanalyses). DTA diagrams obtained from a series of powders milled for 1 min, 4, 10, 20 and 30 h are plotted in Fig. 2. The curve corresponding to the powder milled for only 1 min presents a sharp exothermic peak at 660 “C, the melting temperature of aluminium. With increased milling time this peak becomes weaker and expands to a range of temperatures around 660 “C. Further increase in milling time (20 or 30 h) causes this peak to completely disappear and might be indicative of complete FeAl formation. The results of XRD analysis, as a function of milling time and milling intensity are displayed in Fig. 3. Only a narrow range of diffraction angles is presented (28 =40-56”), but it contains the main lines of all components and is representative of the evolution that occurs. For a short milling time (1 min), only sharp diffraction lines of elemental components are detected. After 2 and 4 h milling these lines exhibit a gradual broadening xld finally the Al line disappears (8 h). In the same time period, the main Fe line becomes asym-

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(4

99

(b) f 64WCl

I

rr’C1

z

Fig. 2. DTA analysis as a function of milling time.

metric and very broad. With further increase of milling time a new broad peak arises on the low-angle shoulder of Fe peak (16 h) and eventually centres in the vicinity of the theoretical FeAl (110) position (24 h). Neither superstructure reflections nor peaks corresponding to DO, structure have been observed in these mechanically alloyed FeAl powders. Mijssbauer analysis yields additional information. Fig. 4 presents Miissbauer spectra (left column) and the associated HFDs (right column, P(H)dH is the probability of finding a field between H and H+ dH) for several milling times. Three main components can be seen on these spectra. For very short milling times the observed Mossbauer spectrum is composed of one main component, that of pure iron which is a sextuplet characterized dy a hyperfine field of H = 330 kG. For milling times up to 8 11

MP48-1

mn/v=2

MP50-2h/v=2

A

MP51-4h/v=2

(0

(d

Fig. 1. Powder morphology evolution during the MA process. Starting (a) iron and (b) aluminium powders; (c) 1 h, 200 rev min-‘; (d) 4 h, 200 rev min-‘; (e) 8-16 h, 200 rev min-‘; (f) 24 h, 200 rev min-‘; (g) 24 h, 400 rev min - ’

40

42

44

46

46

50

52

54

20 [degl Fig. 3. XRD patterns as a function of milling time.

56

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P (H) 0

100

200

300 H W)

-10

-5

0

5

wiches between two or more elemental components and subsequent breaking. With increasing milling time the layers become finer and finer and finally thicknesses of nanometric or even interatomic distances are achieved. This extremely low layer thickness together with substantial increase in temperature due to shocks between colliding balls helps solid state reactions to occur. This process applied to the formation of FeAl intermetallic from elemental Fe and Al powders can be shown schematically (Fig. 5) and is described as follows. The process appears to be composed of two steps: progressive refinement of Fe and Al grains within the sandwich type microstructure (step l), followed by FeAl formation (step 2), presumably at the interfaces between Fe and Al grains. Our analysis, based essentially on XRD for the first step and MS for the second step, showed that these two steps take place almost simultaneously. In fact FeAl formation tends to start after short milling times, at least in a fraction of the processed powders. As FeAl formation is the essential feature of the present MA investigation, we are going to start the discussion with step 2. Magnetic properties of Fe l0,,-,YAl,Y alloys, and consequently their Mtissbauer spectra, depend sensitively on the Al content and also on chemical order, particularly for B2-type alloys with Al contents close to those

10 mm/s

Fig. 4. 57Fe room temperature Miissbauer spectra and some associated HFD as a function of milling time and intensity (200 or 400 rev min - ‘).

the iron sextuplet is still predominant but a second magnetic component characterized by very broad lines arises with gradually increasing intensity. This component is very significant after 16 h of milling and eventually tends to be the sextuplet characterized by an average field H = 250 kG (24 h, 400 rev min - ‘), At the same time, the intensity of the iron sextuplet (first component) is progressively decreasing (16 h); therefore, after 24 h at 200 rev min- ’ only the two external cc-Fe lines can be clearly distinguished. An important increase in milling intensity (24 h, 400 rev min- ‘) is, however, needed to completely remove uncombined iron and achieve iron-aluminium solid-solution formation. Note also the temporary presence of a small peak in the centre of the spectrum (which is in fact a part of a paramagnetic doublet, here called the third component), whose intensity goes through a maximum at 8 h (Fig. 4).

4. Discussion MA is commonly described [j] as a repetitive action of plastic deformation under shocks, welding into sand-

4h

24h

FeAl with dlsordered

structure

@

Feetom

8

Aletom

Fig. 5. Scheme of MA process applied to the formation of Fe Al intermetallics.

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investigated in the present work [g-14]. The room temperature Mijssbauer spectrum of a disordered Fe,,Al,, alloy is composed of a broad magnetic sextuplet with an average hyperfine field of about 200 kG, while a paramagnetic spectrum with a single line is recorded for ordered alloys with a long-range order parameter close to the maximum value of 0.8 [9,14,15]. For Al contents greater than 50 at.%, paramagnetic room temperature spectra, with eventual quadrupole split doublets (x > SO), are observed [16,17]. Miissbauer spectra of powder mixtures of mechanically alloyed Fe,,Al,, with a container speed u = 200 rev min - ’ (Fig. 4) show three main contributions whose relative intensities change with milling time. The external narrow lines of the sextuplet associated with uncombined a-Fe (H= 330 kG) are clearly observed even after 24 h milling but no uncombined a-Fe is left after 24 h milling when the largest container rotation speeds are used (Fig. 4). A central paramagnetic doublet is associated with aluminium-rich alloys. Its intensity goes through a maximum as almost pure aluminium alloys do not contain enough iron to be detected for short milling times. The latter doublet is no longer observed in powders ground for 24 h. The last component is a broad six-line pattern with an average hyperfine field smaller than the field of K-Fe. For milling times less than or equal to 8 h, it is associated with a variety of Fe rich Fe-Al solid solutions as confirmed by an almost flat HFD. For milling times larger than or equal to 16 h, the hyperfine field which corresponds to the maximum amplitude of P(H) is about 250 kG, which is also the average field for samples ground for 24 h whatever the rotation speed (Fig. 4). The latter sextuplet can only be associated with a disordered FeAl solid solution [9] as also shown by XRD patterns without superlattice reflections. Moreover, the alloy composition can be deduced from the concentration dependence of the average hyperfine field at room temperature. Compositions Fe6*Al,, or Fe,,Al,, are respectively obtained from the experimental and calculated field values published by Huffman and Fisher [9]. The complete formation of FeAl solid solution has also been confirmed by XRD, where the asymmetry of (1 IO) peak, still present for 200 rev min- l, disappears for a container rotation speed of 400 rev min-‘. While MS is better adapted for the understanding of FeAl formation (step 2), the nanocrystallization (step 1) can be better assessed by XRD. Initial symmetric broadening of Al and Fe peaks (2 h) is significant because it represents progressive nanocrystallization. Asymmetric broadening of Fe peaks (4 h, 8 h), is mainly due to the formation of iron-rich solid solutions. The unit cell parameters of

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such b.c.c. solid solutions (according to the whole diffraction pattern) are higher than that of pure iron, as indicated by broadening on the left side of the (110) Fe peak. The fraction of milled powders in the form of iron-rich solid solutions increases with increasing milling time and finally (24 h, 400 rev min- ‘) leads to the formation of a broad peak that is close to the theoretical FeAl position. The analysis of the Al peak (cf. Fig. 3) indicates its rapid nanocrystallization, but no information about Al-rich solid solutions could be obtained. The theoretical positions of FeAl diffraction lines cannot be known with precision because they depend on composition. The main fraction of FeAl should have a composition in the vicinity of Fe-35at.%Al. If we take the range of compositions likely to form, as wide as Fe-30at.%Al to Fe-40at.%Al, we are still left with a very small variation in FeAl unit cell parameter, which remains between 0.2895 and 0.2897 nm. The parameter measured on our final product (a = 0.292 nm) is significantly higher ( + 0.8%) than the theoretical one and is attributed to the presence of 1 wt.% carbon coming from stearic acid, and retained in the solid solution. This result was verified by additional measurements with an internal standard (pure Al powder added to the powder MP47) and confirmed by a study of FeAl synthesized by ball milling without a process control agent (stearic acid) [18]. An increase in unit cell parameter was not observed for the latter sample. In order to calculate mean crystallite size (mcs) and microstress evolutions, the (110) peak was first dissociated into two components: the (110) line of iron and the (110) line of FeAl; the calculations based on shape analysis were than performed using the Fe line up to 8 h, and the FeAl line for longer milling times. These calculations permit an assessment of the increase of microstresses and the decrease of mcs. The mcs seems to tend asymptotically to 15 nm. This observation is apparently not consistent with the MA scheme (cf. Fig. 5), where distances of the order of a nanometre, or even of the size of a unit cell, should be achieved in order to form the FeAl compound by solid state reaction. Our explanation of this phenomena is based on the analysis of the XRD detection limit and on the MA process itself, including local inhomogeneities in milled powders. First, the crystallite sizes lower than 15 nm are difficult to assess as corresponding XRD lines become very broad (26’ >>1”). Note that the increasing intensity of the background, between 28 = 44 to 48” and symmetrically, suggests the existence of smaller crystallites (of the order of 4 nm) but that contribution originates from a small volume fraction of the formed FeAl. The evolution of the Al reflection, between 2 and 4 h of milling, gives evidence of nanocrystallization of aluminium. With further increase in milling time up to 8 h

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Table 2 Evolution of mean grain size and of microstresses as a function of milling time and container rotation speed; the calculations are performed on (110) reflection line of Fe up to 8 h and of FeAl for longer milling times

MBchmlcdlyalloy

a.

Zone of Al melting ,

Fig. 6. (a) Sandwich-type internal microstructure developed during MA; (b) influence of morphological texture on the shape of diffraction lines; (c) mechanism of local melting of aluminium.

the Al peak is no longer detected. This nanocrystallization of aluminium certainly precedes the formation of Al-rich and Fe-rich solid solutions as detected from MS analysis. Second, the MA process results in a particular, lamellar morphology (morphological texture), characterized by plate-like internal structure (microsandwiches) of spherical particles (d = 2-5 pm). These particles contain a number of Al and Fe grains for low milling times (2 h, 4 h) (Fig. 6(a)). The thicknesses of these grains certainly decrease down to the level of some nanometres [I] and result in very broad reflection lines, while the two other dimensions of such plate-like grains may still be of the order of some dozens of nanometres, giving sharper reflections (Fig. 6(b)). The overall contribution of Fe grains to the shape of reflection lines yields an effective value of 25 nm (Fe, 8 h). Every diffraction peak is, in fact, a sum of peaks having different widths (depending on grain orientation (and size)) and weighted by the volume fraction of grains in this partic‘ular state [19]. This analysis, based on morphological texture, is particularly helpful to explain the evolution of the iron mean crystallite size. One should note that the overall evolution of crystallite size (see Table 2),

Sample (powder)

Size W-4

Microstresses (x 10-3)

Milling time @)/speed (rev min-I)

MP49 MP50 MP51 MP52 MP37 MP38 MP46 MP47

981 111 65 25 24 22 18 15

0.7 0.6 1 3 1.7 2.4 9.4 10.6

l/200 2/200 4/200 s/200 161200 24/200 24/300 24/400

which was calculated from the (110) Fe reflection for milling times up to 8 h and from the (110) FeAl reflection for higher milling times, appears as regular. Finally, the mean crystallite size of formed FeAl is of the order of 15 nm. The explanation of this asymptotic evolution should be based on the physics of the MA process. Two contributions have to be taken into account. On the one hand, it has been reported [20,21] that the temperature locally achieved during the MA process may increase up to several hundreds of degrees. On the other hand, it is known that melting temperature decreases with decreasing grain size. The latter effect is particularly significant for nanograins. As the tip radius of aluminium grains in a sandwich is of the order of a dozen nanometres, and powder particles are shocked repeatedly, these two contributions can lead to the local melting of aluminium. This local melting results in a fast exothermic reaction for the formation of FeAl [22], probably within one particle, or within a group of strongly welded particles. This is followed by rapid cooling due to the very small volume affected and to the high temperature gradient between the reaction zone and the bulk environment (the average temperature of which was measured to be 80 “C). The crystallite size of FeAl formed by this mechanism of local reaction can be of the order of particle size up to 2-5 pm (cf. Fig. l(g)), but further milling is likely to decrease it. At a given time, only a small volume fraction of synthesized FeAl can be produced by this local mechanism, though this contribution cannot be detected on XRD patterns. With increasing milling time this reaction continues to take place in other fractions of powder while the FeAl crystallites initially formed are simultaneously reduced to nanocrystals. In the final MA stage, the FeAl formed by the latter mechanism and followed by a nanocrystallisation gives rise to the main contribution to the XRD pattern. This mechanism, based on local reactions, certainly needs to be confirmed by statistical observations of the

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internal structure of the milled powder, especially for milling times between 8 and 16 h. It provides us though, with a consistent explanation of the observed final crystallite size of 15 nm and suggests a new mechanism to explain the final step of MA process. A bulk exothermic reaction of FeAl formation has been observed on a global scale by DTA (cf. Fig. 2). During the heating, once the melting temperature of Al (660 “C) is reached, the strong exothermic peak is observed for the powder after a short milling time (1 min). This peak is typical of reactive sintering, that occurs at the melting temperature of aluminium (endothermic process). This endothermic peak (AH = + 2.6 kcal mol- ‘) is not observed because of the exothermic reaction of FeAl formation (AH = - 16.4 kcal mol- l of Al) [23], which takes place immediately after aluminium melting [24]. One can calculate the adiabatic temperature of this reaction, considering that all the heat released during the reaction (1Alf 1.5Fe-+2.5Fe,,,A10,,, AH = - 16.4 kcal mol-’ of Al) is used to increase the temperature of Fe,,,Al,,h formed. In these ideal conditions one can find AT= 1100 “C (the reaction starts between solid iron and melted aluminium at 660 “C, 1 mol Fe,,,Al,,, = 44.3 g, the specific heat has been taken equal to 3R = 24.9 J mol-’ K-i). In real situations, the increase in temperature will be lower than that calculated for adiabatic conditions; however, this will still be high enough to explain the strong exothermic peak. Increasing the milling time causes the solid state reaction of FeAl formation to occur, as proved by the gradually decreasing intensity of the DTA peak. From a mechanistic point of view, this solid state reaction is activated by high energy input from colliding balls. The nature of these collisions (frontal shocks between balls or balls against the container wall on the one hand, or shearing on the other hand), depends on the geometry of the milling device. The analysis of ball movement in a planetary mill (see Appendix A) has shown that the action of shearing is dominant. This observation has no direct consequence on the above discussion, but should be taken into account for MA modelling, where up to now [25,26] essentially only frontal shocks have been considered. It should also be pointed out that our analysis of the MA process has been carried out on powders containing approximately 1% carbon and 5% oxygen, but neither carbides nor oxides could be detected in mechanically alloyed powders. We have shown that carbon influences the unit cell parameter of mechanosynthesized FeAl. Note that this carbon can lead to complex carbide formation [18] and also modify oxide formation [27], both after heat treatment. In fact this heat treatment is necessary to achieve the B2-ordered structure. As far as oxygen is concerned, before heat treatment, it is certainly combined with Fe or Al atoms in the form of non-crystallized oxides. This

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contamination by oxygen explains why the final FeAl composition is significantly lower in Al than is the expected composition of Fe-40at.%Al. Once again, the sintering of this powder will lead to the formation of oxides [28], which are beneficial in improving FeAl creep properties [5]. These oxides might be detrimental for room temperature properties [29]; however, homogeneous dispersions of nanosized oxide particles have been shown to be beneficial in improving room temperature elongation of FeAl [30].

5. Conclusions MA of Fe and Al powders results in the formation of FeAl intermetallic in a chemically disordered state. This process occurs in two steps: a nanocrystallization step and an FeAl formation step. Nanocrystallization of Al powder is faster than that of Fe powder. FeAl formation starts rapidly by the creation of Al-rich and Fe-rich solid solutions; the latter become dominant after long milling times and tend to the composition Fe35 + 2at.%Al. The mean crystallite size decreases to an effective value of 15 nm. The explanation for the latter result relies on morphological texture (microsandwich microstructure) and on a mechanism of local melting of aluminium. MS has been shown to be a powerful tool for characterizing the FeAl formation and is also invaluable in detecting the eventual presence of uncombined iron in the as-milled powders. In the planetary mill geometry the MA process is mainly activated by the action of shearing.

Appendix A. Determination of the nature of ball collisions in a planetary mill Consider a typical planetary mill composed of a disk and a satellite. The satellite rotation speed o equals the disk rotation speed R but they rotate in opposite directions. Only a qualitative description of actual ball motion, based on a dynamic analysis, will be given. No simplifying assumptions are made. As an isolated ball is considered, its equilibrium position is that noted 1 in Fig. Al(a). The rotation of the satellite and the presence of friction between the ball and the satellite wall results in a friction force F,, that tends to move the ball to the position 2. As the ball can roll, the friction coefficient is small and the resulting force (FF) is also small. For each position of type 2, a tangential component FDT of the force due to the rotation of the disc FD appears and tends to bring the ball back to the equilibrium position 1. As a result, the ball is kept in the vicinity of this stable equilibrium position and there is no physical possibility for it to take off from the satellite wall. This ball movement has

K. Wolski et al. I Materials Science and Engineering A207 (1996) 97-104

[25]. In conclusion, the shearing action appears to be dominant in the planetary mill, and this observation should be taken into account in modelling of the MA process.

References a. Isolated ball

[1] J.S. Benjamin and T.E. Volin, Met. Tram., 5 (1974) 1929. [2] J.S. Benjamin, Met. Trnr~s., J (1970) 2943. [31 R. Sundaresan and F.H. Froes, J. hfer., S (1987) 22. [41 P.S. Gilman, Anrz. Rev. Muter. Sci., 13 (1983) 279. 151K. Wolski, Influence de la dispersion de phases ceramiques sur la resistance au fluage de I’intermetallique FeAI, P/r,D., Ecole des Mines de St-Etienne, France, Octobre 1994. PI R.Y. Fillit, A.J. Perry, J.P. Dodelet, G. Pcrrier and R. Philippe, in CO. Roud (ed.), Nordestrmtire Chmterisntiorl of hfnterinls IV, Plenum, New York, 1991, pp. l-8. [71 R.Y. Fillit, in preparation, PI G. Le Ca&r and J.M. Dubois, J. PhJ!s. E:, 12 (1979) 1083. PI G.P. Huffman and R.M. Fisher, J. App, PlrJ,s., 35 (1967) 735. 1101K. Sumiyama, Y. Hirose and Y. Nakamura, J. P/I~s. Sot. Jpj~.,

\-

59 (1990) b. Real situation

c. Vertical section

2963.

1111J.S. Kouvel, in A.I. Berkowitz and E. Kneller (eds.), hfaguetisrn and Metalhrrg~~, Academic Press, New York, 1969, p. 253.

WI E.P. Yelsukov, E.V. Voronina and V.A. Barinov, J. &fq.

Fig. Al. Forces acting on balls in a planetary mill and resulting stable ball configurations.

effectively been observed, for low rotation speed (us = q,), through a transparent top cover. During the MA process, a large number of balls are typically used, so the satellite is filled to 20-40% of its volume. The position of balls can be schematically described as in Fig. Al(b). In this configuration, frontal shocks can occur only when balls are moved from point 3 to 4. The main mechanical action is, however, done in the bulk by shearing between balls. One would also note that the configuration of balls changes within the height (Fig. Al(c)) and that Fs CCF,, because Fs (the force due to the rotation of the satellite) depends on the actual rotation speed of balls around the satellite centre. As all these features should be taken into account without any simplifying assumptions, modelling of ball movements in the planetary mill becomes very complex, and direct observations should be used in order to determine whether shearing or frontal shocks are dominant. Our low speed observations are in favour of shearing. Moreover, this shearing action of balls was recorded at high rotation speeds and several co/L2 ratios

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[I31 u41 [I51