Mechanism of magnetic recovery in the disorder-order transformationof Fe70Al30 mechanically deformed alloys

D. Martín Rodríguez,1,* E. Apiñaniz,1 F. Plazaola,1 J. S. Garitaonandia,2 J. A. Jiménez,3 D. S. Schmool,4 and G. J. Cuello5

1Elektrika eta Elektronika Saila, Euskal Herriko Unibertsitatea, 644 P. K. 48080 Bilbao, Spain2Fisika Aplikatua II Saila, Euskal Herriko Unibertsitatea, 644 P. K. 48080 Bilbao, Spain

3CENIM, Avda. Gregorio del Amo 8, 28040 Madrid, Spain4IFIMUP & Departamento de Física, Universidade do Porto, Rua do Campo Alegre 687, 4169-007 Porto, Portugal

5Institute Laue Langevin, 6 rue Jules Horowitz, B. P. 156, 38042 Grenoble, France�Received 30 July 2004; revised manuscript received 22 December 2004; published 30 June 2005�

The degree of order in Fe-Al intermetallic alloys has an important influence on their magnetic properties.Moreover, the deformation of ordered alloys causes a dramatic increase of magnetization. If deformed alloysare heated, their magnetic properties decrease again. The reordering process was monitored by neutron dif-fraction, Mössbauer spectroscopy, and calorimetric measurements on the Fe70Al30 crushed alloy. This indicatesthat the reordering process occurs in two stages. In the first �150–200 °C� new small B2 phase domains arenucleated due to vacancy migration. A second reordering stage occurs between 300 and 450 °C, where dislo-cation motion induces B2 domain growth and A2 phase elimination. The main mechanism responsible for thisdecrease of magnetization during the reordering process is the decrease of the disordered A2 phase content inthe alloy.

DOI: 10.1103/PhysRevB.71.212408 PACS number�s�: 75.50.Bb, 64.60.Cn, 61.12.Ld

Room temperature �RT� mechanical deformation of well-annealed Fe-Al alloys1–4 induces structural disorder and alarge magnetization increase. There is a close relation be-tween the amount of deformation and the saturation magne-tization in Fe70Al30 ball-milled alloy, which increases signifi-cantly with increasing plastic strain.5,6 There have beenvarious models proposed for this phenomenon, such as an-tiphase boundary �APB� tubes3,7 and APB ribbons.8 Reorder-ing can be used to study the magnetization increase withdeformation in Fe-Al alloys. The study of this binary alloyprovides a model system in which the magnetism is localizedto one atom. While A2 and D03 structures are ferromagneticat RT, the B2 structure is paramagnetic.1 We note that the A2structure has a larger magnetization than that of the D03phase.9

Most of the reported studies have been performed in or-dered B2 alloys, and in particular on Fe60Al40.

3,4,7,10,11 Thereare very few data available for alloys within the range ofD03 structure in the phase diagram.12,13 Calorimetric mea-surements show an exothermic peak at about150–200 °C3,10–12 �first reordering stage�. Just above thispeak a large drop of the magnetic susceptibility has beenmeasured.10,12 However, it is necessary to anneal to tempera-tures above 400 °C in order to obtain a well-ordered alloy.11

Many of the proposed mechanisms for explaining the firstreordering stage are related to dislocation removal8 or APBtubes removal.3,7 Only Amils et al.11 have suggested that forFe60Al40 ball-milled alloy, the first reordering stage at about200 °C might be due to point defect removal, while the sec-ond at about 400 °C arises from planar defect removal.

In the present paper we present in situ neutron diffractionmeasurements, performed at Institut Laue-Langevin �ILL�,Grenoble �France�, that allow an understanding of the recov-ery of magnetization in mechanically deformed Fe-Al alloys.

The Fe70Al30 alloy was prepared at the Max Planck Insti-

tut für Eisenforschung �Düsseldorf, Germany� by inductionmelting. The ingot was powdered by crushing in order toachieve the disordered state, as verified by the absence ofsuperlattice reflections in x-ray diffraction �XRD� scans.Neutron diffraction measurements were performed at ILL onthe D20 instrument using an incident wavelength of 1.3 Å.The sample was heated from RT up to 600 °C at a heatingrate of 1 °C min−1, while neutron diffraction spectra werecollected continuously. A second heating cycle was per-formed up to 600 °C with a heating rate of 2 °C min−1. ASiO2 pattern was used to calibrate the instrumental broaden-ing of the diffraction peaks.

Calorimetic measurements were carried out in a differen-tial thermal analyzer �DTA� �SETARAM TGDTA92� at sev-eral heating rates �5, 10, 20, and 50 K min−1�. 57Fe Möss-bauer spectroscopy measurements were obtained at RT intransmission geometry using a conventional spectrometerwith a 57Co-Rh source.

XRD data of the sample do not show any superstructurepeak.14 However, superstructure peaks are evident in the neu-tron diffraction spectrum of the as-crushed alloy taken at RT�see Fig. 1�. This means that the order has not been com-pletely destroyed. The order-disorder transformation duringcrushing causes a continuous reduction of ordered domainsizes and an increase of volume fraction of the disorderedphase.6 This process broadens and decreases the intensity ofsuperlattice peaks in the XRD pattern until it becomes simi-lar to that of the background. Using neutron diffraction, theintensity of diffracted lines can be increased by a factor of103, which means that the minimum detectable amount ofordered domains is much reduced.

The peaks corresponding to A2 �fundamental reflections�,B2, and D03 structures have been indexed following the D03structure. The superstructure �111� reflection belongs only toD03 structure while �200� reflections belong to both B2 andD03 superstructures.

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Two stages in the reordering process can be seen in thefirst heating cycle �Fig. 2�. In the temperature range150–200 °C, the �200� superstructure peak starts to increaseand broadens, while the �111� peak remains unchanged. Atabout 200 °C the increase plateaus off and remains un-changed from 200 to 300 °C. At a temperature of about300 °C, the �200� becomes even more intense and narrowerup to around 450 °C. Upon reaching 500 °C, which coin-cides with the D03→B2 transition,15 the �111� peak �corre-

sponding only to the D03 phase� disappears. The temperatureevolution of the neutron diffraction spectra in the first heat-ing cycle is completely irreversible. In the second heatingcycle �Fig. 2� only the D03→B2 transition at 500 °C isobserved.

Rietveld refinements using the Fullprof program16 and thepseudovoigt function fitting of selected peaks were per-formed in order to analyze the structural evolution of thesample. Rietveld refinements �see Fig. 3� indicate that in thefirst reordering stage the B2 content increases and the A2decreases, while the amount of D03 remains unchanged. At200 °C the first reordering stage ends and up to 300 °C theB2 phase remains essentially constant. In the range300–400 °C only the B2 phase increases and at about400 °C the D03 phase content shows some small signs ofchange. The D03 phase content reaches the maximumamount when A2 disappears, at about 450 °C, and then van-ishes at the D03→B2 phase transition.

Figure 4 shows the evolution with temperature of the fullwidth at half maximum �FWHM� of the analyzed peaks. TheFWHM of �111� reflection �D03 phase� is much broader thanthe instrumental linewidth and remains constant up to400 °C with an almost ideal Lorentzian shape. This is a clearindication that small and unstrained domains form the D03phase. The Scherrer formula17 can be applied to the Lorent-zian peak, giving an average size of the D03 domains in the

FIG. 1. �Color online� The neutron diffraction scan, with corre-sponding Rietveld fitting �red� and its difference curve �blue� for theFe70Al30 as-crushed sample.

FIG. 2. �Color online� The evolution of thesuperstructure peaks �111� and �200� with tem-perature in both heating cycles.

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as-crushed state of 56.0±0.8 Å. Above 400 °C the �111�FWHM decreases until the D03→B2 phase transition oc-curs. The FWHM behavior of the �200� reflection �commonto both B2 and D03 phases� is completely different. In therange from 150 to 200 °C the FWHM increases and theshape of the peak evolves from a mixed pseudovoigt to aLorentzian one. From 200 °C to 400 °C the linewidth re-mains constant and then decreases from 400 °C to 500 °C,where it reaches the instrumental value. The behavior shownby the FWHM and �200� peak shape evolution in the150–200 °C range must be related to the rapid growth of theB2 phase in the same temperature range �see Fig. 3�. Thelinewidth and its evolution to a Lorentzian shape indicatethat in the first reordering stage, instead of an increase in thesize of the B2 grains, new B2 domains are nucleated. Fur-thermore, the growth of the B2 phase plateaus after the firststage of reordering �see Fig. 3� at about 200 °C. Thiscoupled with the FWHM suggests that the B2 phase nucle-ation occurs in the most strained regions of the sample. Thestrain hinders the growth of the newly nucleated and alreadyexisting ordered phases. Therefore, the results indicate thatthe first reordering stage is caused by the nucleation of smallnew B2 domains in the most strained zones of the disorderedFe70Al30 alloy. This is also a clear indication that most of thestrain remains in the deformed regions of the sample; there-fore the first reordering stage is not caused by dislocationmotion.

The nucleation of new B2 domains, in the first stage ofreordering, can be explained by atomic jumps to nearestneighbor vacancies. Here both vacancy concentration andmigration energy are important parameters in mobility in cu-bic metals.18 In the present case, where the A2 disorderedstate is produced by lattice deformation, the number of va-cancies is large enough to drive atomic mobility. This indi-cates that the activation of the first reordering stage is drivenby the migration of the vacancies present in the A2 lattice.

Calorimetric measurements show an exothermic peak atthe temperature of the first reordering stage. By increasingthe heating rate, the leading slopes of the exothermic peakchange and the peak shifts to higher temperatures. The acti-vation energy calculated using the Kissinger method19 givesa value of 1.11±0.08 eV.

The second reordering stage starts above 300 °C. At thispoint the B2 phase starts to grow again �Fig. 3� and the

FWHM of the studied peaks decreases �Fig. 4�. The behaviorof the �200� and �111� peaks and Rietveld calculations indi-cate that B2 domains grow in this stage from 300 to 400 °C,while the behavior of the �220� peak suggests that a generalstrain removal occurs in the alloy. At about 450 °C the Ri-etveld refinements show that the reordering process is com-plete �Fig. 3�.

Therefore, in the second reordering stage dislocationsstart to move, allowing the growth of B2 domains thus re-leasing strain. It is also important to note that while B2 do-mains start to grow, the D03 phase content remains constantup to 400 °C �see Fig. 3�. This means that the disorder-ordertransformation from A2→D03 �stable phase in the phasediagram of this alloy� needs a transient B2 phase �order tofirst nearest neighbors� first and after D03 phase �order tosecond nearest neighbors� is created.

It has been reported that in the reordering process themagnetization shows a jump at the first reordering stage11,12

in Fe-Al alloys. Figure 5 shows three Mössbauer spectrawith their correspondent magnetic hyperfine field distribu-tions. The spectrum corresponding to the as-crushed sample�top� shows a well-defined sextet with wide peaks, indicatingthe strong magnetic character of the sample. The as-crushedsample was heated up to 300 °C in the DTA at a rate of20 °C min−1, just above the large exothermic peak present inthese alloys around 200 °C �Refs. 7 and 12� and at the end ofthe first reordering stage. The decrease of the ferromagneticcontribution is evident in its Mössbauer spectrum �centralgraph of Fig. 5�. The third spectrum, given in Fig. 5, showsthe changes in magnetic character after the second heatingcycle. The ferromagnetic contribution is much weaker thanin previous spectra and is identical to that observed in thefully ordered Fe70Al30 alloy.6

FIG. 3. The temperature evolution of the A2, B2, and D03 phasecontents in the first heating cycle.

FIG. 4. The evolution of FWHM for �a� superlattice reflections�111� �circles� and �200� �squares� and �b� �220� fundamental reflec-tion in both heating cycles �filled symbols for first heating cycle andopen symbols for the second one�. The FWHM of the SiO2 patterncorresponding to angles similar to the ones of the sample’s diffrac-tion peaks are marked. Lines are drawn as a guide to the eye.

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The reported abrupt decrease of magnetization12 and thechange in the Mössbauer spectra in the first reordering stageclearly must be related to the B2 domains nucleation withinthe A2 disordered phase. The B2 phase is known to benonmagnetic1 and self-consistent electronic calculations9 in-dicate that A2 disordered phase is the most magnetic phase

of Fe-Al alloys. The lattice parameter decrease is less than0.01 Å up to 300 °C; self-consistent electronic calculations20

show that the decrease in lattice parameter is so small that itwill not cause any appreciable change to the Fe magneticmoment. Therefore, the decrease of magnetic propertiesarises mainly from the decrease of A2 phase content.

The reordering process is very different from the disorder-ing one. In the disordering process the sextet present in theMössbauer spectra has the same hyperfine field from the be-ginning of the deformation, and only the area of most of thesubspectra changes with deformation.6 However, after thefirst reordering stage, the maximum field of the hyperfinefield distribution decreases by about 10%, indicating that va-cancy migration does not only induce B2 domains nucleationbut also a slight reordering in the neighborhood of Fe atoms,decreasing the number of nearest Fe neighbors.

Summarizing, the reordering process of mechanically de-formed Fe70Al30 alloys occurs in two stages. In the first, from150 to 200 °C, B2 phase domain nucleation occurs due tovacancy migration. A2 phase removal is the main cause ofthe magnetization reduction of the alloy. In the second reor-dering stage, between 300 and 450 °C, dislocations becomemobile and nucleated B2 domains start to grow. This causesthe decrease of A2 phase content and the further weakeningof the alloy’s magnetism. It was found that for alloys in theD03 phase field it is necessary to first nucleate a B2 phase inorder to create D03 phase afterwards.

This work was undertaken under Project Nos. MAT2002-4087-C02-01 and UPV 224.310-14553/2002. We gratefullyacknowledge J. Torregrossa during the neutron diffractionexperiment and Dr. M. T. Fernández-Díaz for very helpfuldiscussions. D. M. R. wishes to thank the Basque CountryUniversity �Euskal Herriko Unibertsitatea� for support.

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FIG. 5. The Mössbauer spectra and magnetic hyperfine fielddistribution taken �from top to bottom� in the as-crushed sample,after heating it up to 300 °C and after two heating cycles.

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