Metastable and nanocrystalline polymorphs of magnetic intermetallics

MATERIALS SClEWCE & EMGIBIEERIWG Materials Science and Engineering A226228

Metastable

A

(1997) 491-497

and nanocrystalline polymorphs intermetallics

of magnetic

A.R. Yavari LTPCM-CNRS

lJRA29,

Institut

National

Polytechnique

de Gsenoble,

BP 75, Domaine

Univemitaire,

38402

St Martin

d’H&es,

France

Abstract Non-equilibrium processing of ordered magnetic intermetallic compositionsis found to lead to the formation of high temperaturephasesand structures.Examplesare given for transformationsto high temperature phasecrystal structuresand for suppressionof chemicallong-range order and the superlatticethat lead to disorderedsolid-solutionswithin the original crystal lattice. Saturation magnetization and Curie temperaturesin the metastablepolyrnorphs can be higher or lower than the ordered intermetallicsdependingon the crystal structure. Observation of low temperaturespin-glassbehavior in metastablepolymorphs of concentratedbinary intermetallicsare also considered.Finally, it is shownthat when magnetic contributions can be properly accounted for, the heats of formation of ordered intermetallicscan be derived from simpleDSC measurementsof the heat of reordering. 0 1997Elsevier ScienceS.A. Keywords:

Magnetic intermetallics; Metastable polymorphs; Nanocrystalline polymorphs

In this work we will give examples of each of the

1. Introduction Irradiation, heavy deformation (milling, filing, rolling) and rapid solidification (vapor deposition, liquid quenching) allow formation of metastable polymorphs of intermetallic and stoichiometric compounds [l-9]. These metastable polymorphs of different struc-

above from our research performed in the framework of a European Union research network including Cambridge, Grenoble, Barcelona, Trento and Nancy [ll], from the work performed at the Van der Waals-Zee-

man Laboratory in Amsterdam [7] and at the Institute de Magnetimo Aplicado in Madrid [12].

ture or chemical order possess previously unavailable atomic arrangements and modified electron band structures and properties. One of the most interesting classes of such metastable structures is that of intermetallics

containing one or more magnetic elements as magnetic order is highly sensitive to nearest-neighbor chemical order as well as to the crystal structure and magnetization and Curie temperatures vary sensitively with the

degree of order [lo] When an ordered intermetallic is subjected to nonequilibrium processing, three phenomena may occur: (a) Instead of the ordered intermetallic crystalline superstructure, an amorphous phase is obtained.

(b) Non-equilibrium processing of the intermetallic composition leads to a different crystalline structure with a different topological order.

(c) The topological

order and the basic crystal lattice

of the intermetallic is preserved but the chemical long-

range order (superlattice with sublattices) is eliminated or the order parameter reduced. 0921-5093/97/$17.00 0 1997 Elsevier Science S.A. All rights reserved. TlTTnnnnr rnn*,n,\,nF.T, m

1.1. T%e ordered intermetallic crystalline structure is replaced by an amorphous phase

This subject has been treated frequently in the literaamorphization of an alloy is obtained by rapid solidification, it is because of insufficient time for nucleation and growth of crystalline phases from the liquid state. Although this occurs more easily in near-eutectic compositions, the idea that liquid alloys of intermetallic compositions are not easy glass-formers is wrong. For example, both Fe,B and Fe14Nb2B can be amorphized by high-speed melt-spinning (see [13] and [14]). Polymorphous amorphization by heavy deformation occurs as increasing densities of defects are introduced until long range order crumbles by lattice instability [15]. In chemical terms and for an AB compound, we may consider that anti-site defects and APD boundaries ture and will be discussed only briefly. When

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generate ‘wrong bonds’ such as A-A and B-B bonds when A-B bonds have the lowest energy. The consequent enthalpy increases due to ‘elastic mismatch’ and other factors have been reviewed in the literature [16]. The subsequent replacement of such a high-energy crystalline state by an amorphous (liquid-like) phase results in the disappearance of such stresses and most anti-site defects. These contributions and others such as the lattice stability term for each crystalline polymorphs can be used to construct free energy versus concentration curves for the amorphous E17,18] and the increasingly deformed crystalline phases. However these considerations neglect the role of grain boundaries generated by precipitating dislocation cell walls that lead to nanocrystallization. Indeed, amorphization during milling is accompanied by nanocrystallization. The energy surplus due to grain boundary free energy per unit area ygb is approximately AGgb = 1.5 I/& V,/r while the free energy of polymorphous melting is approximately AG,, = AH&?, - T)/T, where AH,, is the enthalpy of melting at the melting temperature T, and V, is the molar volume. This effect alone would lead to spontaneous amorphization for grain radii Y, < 1.5 ySbV,/AGls as discussed elsewhere [19]. While the saturation magnetisation does not necessarily change with amorphisation, coercivity always drops with the disappearance of magnetocrystalline anisotropy. 1.2. Non-equilibrium crystalline structures

processing leading to different with a dzjjfeerent topological order

Rapid quenching from high temperatures (vapor deposition, meltspinning etc.) often freeze-in high temperature phases due to insufficient time for nucleation and growth of the equilibrium phases. Interestingly, generation of defects within the solid state by ball milling or ion irradiation also leads to the formation of high temperature phases. While it may appear paradoxical, this phenomenon is in fact expected, as high temperature phase formation is a form of absorption of the energy input. When annealed, quenched-in high-T phases transform to the equilibrium intermetallic structure exothermically thus releasing the quenched-in excess energy. Generation of defects within the solid state also corresponds to energy storage which is why it can lead the system back to high-temperature, high-energy phases and atomic configurations. A good example of such systems is Fe,Rh,-, which at lower temperatures takes a bee structure up to x = 0.55 but is fee for all compositions in a wide temperature range above [20]. FeRh solidifies as a paramagnetic fee solid-solution but upon cooling undergoes a first order transformation to paramagnetid ordered (B2) bee near 1600 K. Upon further cooling, the paramagnetic ordered-B2 or

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(a’) becomes ferromagnetic near 680 K with magnetic moment per atom ,!J= 1.75 fin at 403 K. At Still lower temperatures, a ferro-antiferromagnetic (a’ --f u”) transformation occurs in properly annealed samples [20]. Instead of a disordered bee phase, quenching from the liquid state by melt-spinning results in the retention of the high temperature y-fee solid-solution for Fe contents above 30% except near the equiatomic 50% Rh concentration ([21]) which orders to a’ during the quench. The metastable state of the high-T y-fee FeRh, which is paramagnetic, is obtained by filing [22] or milling [23]. After transformation to the metastable fee phase, the sample becomes paramagnetic. Consistent with [22], Fig. 1 shows that when the metastable fee FeRh of [23] is heated, it transforms to the ordered bee equilibrium phase near 525 K as signaled by an abrupt increase of magnetization (the antiferromagnetic transition in this sample is below 330 K). 1.3. Crystal lattice of the intermetallic is preserved but the chemical long-range order (supet+latticewith sublattices) is eliminated OYthe order parameter reduced

When the equilibrium crystal lattice is not destabilized by chemical disordering, various properties may evolve with the degree of chemical order as expressed by a chemical long-range-order (LRO) parameter. The degree of order is determined from the relative intensities of superstructure peaks using X-ray (or neutron) diffraction during disordering (for example by ball-milling) and reordering by annealing [24]. TEM observations during annealing in the microscope provide information on the reordering mechanism 1251.

H = 1000 Oe

400

I 500

I

I

600

700

T (K)

Fig. 1. Magnetization measureme& during heating and subsequent cooling of the metastable fee Fe-Rh phase (applied field H = 1 kOe).

A.R. Yauari /Materials

NI

7s

AL

12

Scierzce md Engineering A226228

Fe

M(T)

493

13

T IK)

Fig. 2. Magnetization

(1997) 491-497

T

(K)

versus temperature during heating to reorder and subsequent cooling of LI,-Ni,,Al,,Fe,,

disordered by milling [lo].

1.3.2. Bee disordered state of magnetic B2 irQermetaEEics 1.3.1. FCC disordered state of magnetic Ll, intevnetallics

Upon disordering, some Liz (ordered fee) intermetallies form amorphous alloys as in the case of Zr,Al [S]. In other cases such as for N&AI, the Ll, superlattice can be destroyed without significant amorphisation [5,7,26] and a disordered fee solid solution can be obtained. The y’-Ll, state of Ni,Al is known to be ferromagnetic only at very low temperatures but a few at. % Fe-addition raises the Curie temperature T, very rapidly to several hundred degrees K as giant moments around Fe-atoms polarize the surrounding nickel atoms [27]. Ferromagnetic N&Al + Fe alloys in their y’-ordered state usually become paramagnetic before undergoing the order-disorder (y’+ r) transition to their high-T disordered configuration. (T, < T,d> such that the equilibrium y-phase was thought to be paramagnetic [28,29]. However the disordered y-state produced by milling of the ordered state is ferromagnetic at room temperature [30]. Fig. 2 shows magnetization M(T) versus temperature curves for one such alloy. M(T) decreases for the metastable fee state before increasing with the onset of reordering during heating. The giant moment of Fe-atoms is responsible for the magnetization of neighboring Ni atoms and since in the ordered state Fe preferentially substitutes for Al sites which have only Ni nearest neighbors, the magnetization is significantly lower in the disordered state with a lower number of Fe-Ni nearest neighbors. On the other hand, ball-milled bee metastable polymorphs of ordered intermetallics with a single magnetic moment-carrying species such as FeAl [31] and many others [7] have been found to possess TC’s higher than those of the equilibrium ordered states because disordering increases the number of moment-carrying nearest neighbor atoms.

We consider disordering in the simplest of ordered bee intermetallics which are AB binary B2 structures where A and B atoms occupy cube corners and centers respectively. One of the most studied intermetallics of this type is FeAl. Contrary to isostructural FeRh which becomes fee upon disordering and CoZr which becomes amorphous, FeAl remains bee as its B2 superstructure is destroyed by heavy deformation (rolling, milling, filing). As mentioned earlier, disorder in these alloys increases magnetization and Curie temperature because it increases the number of pairs of magnetic atom nearestneighbors. While both ordered and disordered FeAl are paramagnetic at room temperature, Fe,,Al,, which is in the range of stability of the B2 structure becomes strongly ferromagnetic at room temperature upon disordering. Fig. 3 from [31] shows the increase in room temperature magnetization of ordered FehOAl,, during disordering with increasing milling time. X-ray diffraction which shows progressive decrease and final extinction of B2 ‘superstructure peak intensities during

1CO

Fig. 3. Increase in room temperature magnetization of ordered Fe,,AI,, during disordering with increasing milling time [31].

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2. Derivation of the formation enthalpy of B2-Fe,,AI,, from heat of reordering of disordered FeGOAl,, TGM: 72h 8M 20 Wmin -

mntreated

-

403Klh

______.

5=1(jh

The heat of formation of B2-FeAl intermetallics containing X, atom fraction of Al and X,, atom fraction of Fe from pure fee Al and bee Fe can be written as: AHformation = X,, nHz(bcc

+ AH,,(intermetallic) with AH,,,(intermetallic)

O-1 m

400

500

m

700

so0

- fee) + AH,&intermetallic)

9M)looo

TEMPERATURE (K) Fig. 4. Thermogravimetic (TGM) magnetization curves obtained during continuous heating of Fe,,Al,, disordered by heavy deformation (as-milled and after annealing) [31].

disordering can be used to determine the magnetization as a function of the order parameter. Fig. 4 from [31] shows the evolution of the magnetization of disordered Fe,,Al,, during reordering with increasing temperature. The sharp drop that begins below 500 K is due to ordering inside the grains as confirmed by the order parameter 1 2 Y 2 0, determined by X-ray diffraction, reaching its maximum possible value for this composition. The second small drop that follows and is completed below 800 K is due to grain growth because milling has also resulted in nanocrystallization and a significant contribution to magnetization, of residual disordered Fe,,Al,, at the grain boundaries disappears with grain growth. Thus, the Curie temperature T, of disordered Fe60A140 is not measurable as the transformation back to the ordered state during heating begins well below T,. However, T, should be near 800 K or higher while that of ordered Fe,,Al,, is below room temperature. We have recently proposed a new method for estimating the heats of formation of ordered intermetallics from the heat release measured during reordering of partially or fully disordered states of the intermetallics [32]. The resulting estimates of formation enthalpies are close to values obtained from other experimental data provided that care is taken to properly account for the magnetic contributions to the formation enthalpies in the ordered and disordereed states. We will now apply this method to the reordering of Fe,,Al,,.

= H,,,(intermetallic)

- X,, * H,r& (1)

where X,,+H~~(bcc - fee) is the so called lattice stability term associated with the enthalpy difference between the bee lattice and the equilibrium room-temperature fee lattice of Al, this contribution being zero for Fe. Lattice stability contributions are experimentally available or calculated with good precision and given in internationally compiled data bases such as [33]. The magnetic enthalpy change upon alloying AH maaneric(intermetallic) is calculated using experimentally measured magnetic moments B, (in ,u, Bohr magnetons per atom) and Curie temperatures T, of the intermetallic following a method proposed by Hillert and Inden [34,35] which fits well the data for pure elements (zero for Al) or alloys with a single magnetic component following a simple dilution law: H,,.,,,(T) = RT ln(B, + l)h(T/T,)

and h(T/T,)

= [ - OS64/p(T/T,) + 0.954(1/p - 1)((T/T,)3/2 + (T/TJg/15)]/D 69

where R is the gas constant, p = 0.40 and D = 1.5583 for bee structures. Compared to the other terms of Eq. (l), Hmas is negligible for T/T, > 1. Finally, AH,;, (intermetallic) is the equivalent of the mixing enthalpy AH,, (solid sol.) for solid-solutions but includes the contribution AHLRCO (intermetallic) of the chemical ordering of the solid-solution into sublattices to form the compound. Thus, AH,, (intermetallit) = AH,, (solid sol.) + AHLRCo (intermetallic). The can be measured for example, by comtotal AHformation parison of the enthalpy of solution, measured by dissolution calorimetry at infinite dilution, for the A,B, compound as compared to that of a mixture of nA and mB pure elements but such experiments are generally very difficult especially for high temperature melting compounds. Consider the formation enthalpy of the disordered alloy (solid-solution). Starting with the energy change 2AE = 2EMmFe- (EFeAFe+ EAITAI) that accompanies the replacement of an Al-Al ‘bond’ and an Fe-Fe ‘bond’ by two Al-Fe ‘bonds’, the regular solution model defines a

A.R. Yauari/Materinls Table 1 Contributions

495

to the formation enthalpy AHrormation of ordered intermetallic of Fe,,AI,,

Lattice stability X,,*Hky(bcc-fee) kJ/g.at. AffmagncticYintcrmetallic= 0.67 kJ/@. AfLagnctic

Scierzce and Engineering A226-228 (1997) 491-497

Ydisordcrcd

= 0.1’7

Experimental AH,,,,(Y) g.at.

4.0 4.3 1.0 & 0.1

Wwt.

of reordering kJ/

4.5 f 0.2 [31]

Interaction energy parameter Q [Eq. (6)] kJ/g.at. Afkmmtmn (intermetallic) [experimental] kJ/g.at. [36] (intermetallic) (Miedema) kJ/g.at. 1361 Wmnatmn AH,,,,,,, (intermetallic) from AHreordering [Eq. (7)] kJ/g.at.

-96” -26-i:2 -31 -25.3 + 0.3

“DSC results for sample with order parameter going from Y = 0.167 to Y = 0.67 as measured by x-ray diffraction.

nearest-neighbor interaction energy parameter R given by 12 = Z IcAv[EAleFe- (EFe _ Fe+ EA1-A1)/2] where Z is the number of nearest neighbors (n.n.) in the bee crystal structure. The excess enthalpy associated with formation of disordered Fe,Al,+ containing atomic fractions X,, and X, = 1 - X,, is then approximated by the expression AH,h,ti&solid

sol.) = R * X,, * X,

(3) The multiplication by XF,XA amounts to a bond counting procedure where the probability for each of the ZkAVXA1 nearest neighbors of Al atoms randomly distributed in one mole of Fe crystal to be an Fe atom will be equal to X,,. Eq.(3) is not appropriate for ordered Fe,Al, _ .y intermetallics where the probability of finding an Fe atom next to an Al atom is 1 instead of X,, (for X,, > 0.5). For the intermetallic with longrange chemical order (two sublattices), one can define an order parameter

where Xi is the fraction of the Al sublattice sites actually occupied by Al atoms, and Y goes from 0 for a randomly disordered state (equivalent to a solid solution) to 1 for fully ordered FeAl (intermetallic superstructure). The mixing enthalpy AH,, can then be written as AHmixing(Yinrermetauic) = AHmixing (solid Sol.) + MLRCO = ax,,+

(X,, + YX,,)

with AHmixing P’intermeta~~ic= 1) = fiX, and AH,,,, = R * XL. Y (4) Consider now the ordered intermetallic Fe,,Al,, in which the order parameter Y is reduced from its maximum value Y, = 0.67 to Y, = 0.1 by heavy deformation. The exothermic heat-release occuring during thermally induced reordering AHDSCereorder(Y)= 4.5 kJ g - ’ at.% from Yz back to Y1 as measured by DSC calorimetry [31] could then be approximated by:

= sZ.X2,~(Y, - YJ + H,,,,,(Y’,) - Kn,,P”,) (5) where Y is determined independently by X-ray diffraction and the magnetic contribution to the intermetallic’s enthalpy at a given degree of order defined by Y is calculated using Eq. (2), from experimental values of 3, and T,. The nearest-neighbor interaction energy parameter is then derivable from: Q

=

WI-‘)- Kna,W’,) - fLa,tW)I ix’, . WY,- Y2) W4xc

- reorder

(6)

with AH~m-mat~ion= X,, . H;:(bcc

- fee)

+ AH,,,(intermetallicY +Q2xX,.(X,,+Yl

1) X*J

(7) Table 1 gives the lattice stability, magnetic contributions to the intermetallic formation energy AHformation of Eq. (7), experimentally measured ordering energies AHDSC-reorcier(V as well as R derived from the these values using Eq. (6). The lattice stability terms have been calculated for T= 500 K using standard CALPHAD equations [33]. T = 500 K was selected because the transformations back to the ordered state during annealing were found to occur near 500 K as monitored by calorimetry. Magnetic contributions were calculated with the experimental magnetic moment B, = 1 /LB/atom of disordered Fe,,&,, [31]. The Curie temperature T, of disordered Fe,,Al,, T, 2 800 K as discussed in connection with Fig. 4 and we have used this value to calculate its Wnag Using the values of Table 1 in Eq. (7), we derive AH~omation and compare it in Table 1 to the estimation with the Miedema model and AHFormation due to derived from other experimental results [36]. A correction is needed to compensate for heat release from grain growth that occurs during ordering. In the case of Fe,,AI,, the grains were found to grow from a

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from the DSC heat release measured during ordering (4.5 kJ g-l at.%> used in Eq. (6) will lower the interaction energy n and heat of formation AHformntion by about 7%. This correction will bring the estimated value of AHfonnation for Fe,,AI,, even closer to the experimental value reported by others. It can thus, be seen that from the study of the reordering of disordered (metastable) intermetallics, it is sometimes possible to derive the formation energy of the stable ordered intermetallic provided that magnetic contributions are properly accounted for. 3. Low temperature behavior of disordered intermetallics

40

0

80

Historically, spin glass behavior, which is thought to be characterized by a frozen state of a disordered mixture of ferromagnetic and antiferromagnetic local interactions has been observed in dilute solutions of a magnetic element in a non-magnetic matrix such as Fe in Au or Mn in Cu [37,38] or Fe-based multicomponent metallic glasses [37]. More recently, two new spin glass states of a new type were discovered at the Van der Waals-Zeeman Laboratory in Amsterdam [39,40]. Spin glass behavior was observed in heavily deformed polymorphs of stoichiometric compounds, Co,Ge in amorphous state and Cl 5-type GdAI, in nanocrystalline disordered state. Fig. 5 from Zhou and Bakker [40], shows dc susceptibility versus temperature curves indicating maxima corresponding to spin glass freezing temperatures Tf near 50 K in nanocrystalline Cl5 (cubic) GdAIP.

120

T (K) Fig. 5. dc susceptibility versus temperature for nanocrystalline GdAl, spin glass obtained by ball milling [40].

diameter of 12.7 nm in the as-milled disordered state to 18.2 nm during ordering at 523 K (see [31]) corresponding to a grain-boundary specific surface reduction of 1.2 x lo7 cm2 g-’ at.% with an expected heat release of about 0.6 kJ g-l at.%. Deduction of this contribution

H = 5000 Oe * FC 7 ZFC

..

4+

:

L6

1 4 4

.

4

+

4

+

*

4

+

I,4

12

, 0

50

I I 1. 100

150

T (K>

1

‘t

+*

I

I

200

250

l +. 300

0

50

100

150

200

250

300

T 09

Fig. 6. Low temperature Magnetization 1M and dc suceptibility M/H of ball-milled fee FeRh during heating after field cooling (FC) and zero field cooIing (ZFC) showing field dependent maxima at a freezing temperature T r= 65 K. The irreversiblility is spin-glass like [41].

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Yavari/Materials

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and Engineering

The ordered bee (B2) phase FeRh and neighboring compositions present a ferro-antiferromagnetic (a’ + a’) transformation upon cooling near room temperature [20]. The low temperature magnetic properties of metastable fee FeRh were therefore studied during heating after field cooling (FC) and zero field cooling (ZFC). The resulting curves [41] for magnetization M and M( V)/H where M( y> is per unit volume shown in Fig. 6 are spin-glass like with the curves showing clear breaks near 65 K as freezing temperature Tf and may be compared to those of previously discovered nanocrystalline GdAl, spin-glass as reproduced in Fig. 5. Further work is needed to determine if this low T behavior is truely that of a spin glass or if it is superparamagnetic behavior of some residual ferromagnetic nanoparticle contaminants for example from the milling device made of stainless steel. In conclusion, the study of metastable polymorphs of ordered intermetallics is found to be an exciting and promising field of research that can provide positive feed-back for testing theoretical thermodynamic hypotheses about the various phases and leads to the discovery of new forms of magnetic behavior. It can also provide new information on the kinetics of phase transformations and lead to the development of new nanocrystalline materials with potential for applications but which are out of the scope of this paper.

References [I] [2] [3] [4]

CL. Corey and D.I. Potter, J. Appl. Phys., A.R. Yavari and B. Bochu, Philo. Mag. A, J.S.C. Jang and CC. Koch, J. Mater. Res., J.S. West, J.T. Manos and M.J. Aziz, Mater (1990)

38 (1967) 3894. 59 (1989) 697.

S(1990) 498. Res. Sot. Proc.,

213

859.

[5] T. Benameur and A.R. Yavari, J. Mater. Res., 7 (1992) 2971. [6] A.R. Yavari, S. Gialanella, T. Benameur, R.W. Cahn and B. Bochu, J. Mater. Res., 8 (1993) 242. [7] G. Zhou, Doctoral Thesis, Group of Prof. H. Bakker, Universiteit van Amsterdam, November 1994. IS] G.F. Zhou and H. Bakker, P/qx. Rev. Lett., 73 (1994) 344. [9] A.R. Yaveri (ed.), Ordering and Disordering in Alloys, Elsevier Applied Sciences, 1992. [lo] A.R. Yavari, Acta Metall. Mater., 41 (1993) 1391. [ll] European Union research network, Phase Transformations in Nanocrystaliine Materials, coordinated by R.W. Cahn.

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491

WI Collaborative work during the Betancourt-Perronet

Prize sabbatical of A.R. Yavari. and NanocrysEl31 A.R. Yavari, in M. Vasquez (ed.), Nanostmctwed talline Materials, Hernando, World Scientific, 1995, pp. 35-43. El41 D. Negri, A.R. Yavari, M. Vazquez, A. Hernando, A. Deriu and T. Hopfinger, Proceedings ISMANAM-96, Rome May 1996, M. Magini and D. Fiorani (eds.), Mate?. Sci. Fol~irn, in press. [151 P.R. Okamoto, L.E. Rehn, J. Pearson, R. Bhadra and H. Grimsditch, J. Less Coinm. Metals, 140 (1988) 231. 1161D.L. Beke, P.I. Leoff and H. Bakker, Acta Metall. Mater., 39 (1991) 1259 and 1267. El71 T. Benameur and A.R. Yavari, J. fifater. Res., 7 (1992) 2911. [I81 H. Bakker, G.F. Zhou and H. Yang, in A.R. Yavari (ed.), Mechanically

Alloyed

arid

h~anocrystalline

Materials,

IS-

MANAM-94, iMater. Sci. Fawn (1995) pp. 47-52. WI Yavari A.R., Trans. Japan Inst. Metals, 36 (1995) 228 Lw L.J. Swartzendruber, Bull. Alloy Phase Diagrams, 5 (1984) 1099. WI CC. Chao, P. Duwez and CC. Tsuei, J. Appl. Pilys., 42 (1971) 4282. WI J.M. Lommel and J.S. Kouvel, J. Appl. Phys., 38 (1967) 126. ~31 Navarro E., Yavari A.R., Hernando A., Martina C., Ibarra M.R., Solid State Conm., 100 (1996) 57. 1241 L. Lutterotti, S. Gialanella and R. Caudron, Proc. EPDIC Chester 1995, Mater. Sci. Forum, in press. VI S. Lay and A.R. Yavari, Acta Metall. Maler., 44 (1996) 35. WI M.D. Baro, S. Surinach, J. Melagelada, M.T. Clavaguera-Mora, S. Gialanella and R.W. Cahn, Acta Metail. IMater., 41 (1993) 1065. 1271 J.R. Nichols and R.D. Rawlings, Acta Metall. Mater., 25 (1977) 187. VI R. Kozubski, J. Soltys and M.C. Cadeville, J. Phys. C., 2 (1990) 3451. ~91 K.G. Efthemiadis, J.G. Antonopoulos and LA. Tsoukalas, Solid State Comm., 70 (1989) 903. [3Ol A.R. Yavari, M.D. Baro, G. Fillion, S. Surinach, S. Gialanella, M.T. Clavaguera, P. Desrk and R.W. Cahn, Mater Res. Sot. Proc., 2I3 (1990) 859. [3 11 S. Surinach, S. Gialanella, X. Amils, L. Lutterotti and M.D. Baro, R. Schulz (ed.) Proceedings ISMANAM-95, Mai 1995 Qu&bec, Mater. Sci. Forum, in press. ~321 A.R. Yavari, S. Gialanella, M.D. Baro and G. Le Caer, submitted to Phys. Rev. Lett., 1996. 15 (1991) 317-425. [331 A.T. Dinsdale, CALPHAD, 2 (1978) 227. [341 M. Hillert and M. Jarl, CALPHAD, [351 G. Inden, Physica, 103B (1981) 82. [341 F.R. de Boer, R. Boom, A.R. Miedema and A.K. Niessen, Cohesion in Transition Metal Alloys, North Holland. Introduction, Taylor [371 J.A. Mydosh, Spit1 Glasses: An Expeknental and Francis, London 1993. [381 V. Cannella and J.A. Mydosh, Phys. Rev., B9 (1972) 4220 t391 G.F. Zhou and H. Bakker, Phys. Rev. Lett., 72 (1994) 2290; 73 (1994) 344. 1401 G.F. Zhou and H. Bakker, Phys. Rev. Lett., 73 (1994) 344. r411 E. Navarro, M. Multigner, A,R. Yavari, A. Hernando, EwevAvs. Lett.. 35 119961 307.