The Synthesis and Characterization of Arc Melted Fe1-xAlx Alloys

doi:10.4236/jamp.2013.12002

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Journal of Applied Mathematics and Physics, 2013, 1, 4-8

http://dx.doi.org/10.4236/jamp.2013.12002 Published Online May 2013 (http://www.scirp.org/journal/jamp)

The Synthesis and Characterization of Arc Melted

Fe1−xAlx Alloys

Sandeep Rajan1*, Rajni Shukla1, Anil Kumar2, Anupam Vyas2, Ranjeet Kumar Brajpuriya2

1Department of Physics, Deenbandhu Chhotu Ram University of Science & Technology, Sonipat, India

2Department of Physics, Amity University Haryana, Gurgaon, India

Email: *[email protected]

Received May 10, 2013; revised June 11, 2013; accepted June 17, 2013

which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

ABSTRACT

The paper presents correlation study on a series of Fe1−xAlx alloy samples prepared by arc melting. All the samples

show crystalline structure, irrespective of the Al content and are textured mainly along (110) direction. The particle size

decreases rapidly with x particularly after x > 0.3. The corresponding magnetic measurements were obtained at room

temperature using a VSM, with a maximum applied field of 14 kOe. The results show that the ferromagnetic state of the

samples disappears with x, and becomes paramagnetic for alloys with x ≥ 0.4. It is also found that coercivity (Hc) and

resistivity increase with x. The results were interpreted in terms of continuous change in their electronic structure i.e.

overlap of the electron wave functions of the magnetic atoms with the Al electron wave function.

Keywords: Transition Metal Aluminides; Magnetization and Electronic Properties

1. Introduction

During the last decade, enormous progress has been

made in manufacturing and processing of materials with

intermetallic matrices [1]. Mechanical alloying (MA) [2],

which was initially conceived for the production of dis-

persion strengthened superalloys, is nowadays used for

synthesizing a wide range of materials including inter-

metallics [3]. A great advantage of mechanical alloying

(MA) process is the fact that it modifies the structure and

solid solubility limits of alloys and solid solutions and

induces lattice strains and phase transformations [4,5].

The formation of metastable phases and disordering of

the lattice through alloying gives rise to interesting me-

chanical and magnetic properties [6]. This is particularly

evident in the case of Fe-Al intermetallic systems. The

complicated phase diagram and the dependence of mag-

netic properties on the state of disorder and the micro-

structure [6,7] make this an interesting system to study

through the MA process. Both from the point of view of

understanding the collective behaviour of the magnetic

phases and also possible applications, study of such in-

termediate phases is expected to be very rewarding. As

FeAl intermetallic compound offers a combination of

several attractive properties such as a high specific

strength, good strength at intermediate temperatures and

an excellent corrosion resistance at elevated temperatures

under oxidizing, carburing and sulfidizing atmospheres

[8,9]. So they have attracted considerable attention as

potential candidates for structural and coatings applica-

tions at elevated temperatures in hostile environments

and as promising substitutes for stainless steels at room

temperature [10-13]. Due to these reasons, mechanically

alloyed [3-5] and mechanically milled [14] Fe-Al sys-

tems of various compositions have been investigated in

recent years. With this objective, the authors have suc-

cessfully synthesized FeAl intermetallics by MA and

systematically investigated the structural, magnetic and

electronic properties of a series of arc melted Fe1−xAlx

(0.2 ≤ x ≤ 0.6) alloys using XRD, VSM and XPS.

2. Experimental Details

A series of arc melted Fe1−xAlx alloy samples were pre-

pared under Argon atmosphere using high purity (more

than 99.9%) constituent’s metals. Afterwards, the poly-

crystalline ingots were annealed at 600˚C for 120 hr un-

der the UHV condition.

The structure, average grain size and lattice constant a0

of each sample was inferred from the X-ray diffraction

(XRD). In the present study, the average crystallite size

<L> was calculated from the broadening of XRD peaks

*Corresponding author.

S. RAJAN ET AL.

according to Scherrer’s formula

cos









where k represented the shape factor which varied with

the crystal shape, λ is the wavelength of the CuKα radia-

tion, θ is the Bragg angle, and



is the Lorentzian part

of the integral breadth due to the crystallite size.

The lattice parameter calculated using the equation

2222

4sin hkl









The corresponding magnetic and resistivity measure-

ments were done using vibrating sample magnetometer

and standard four-probe resistivity method. The chemical

and electronic information of the samples at different

depth has been obtained from XPS technique. X-ray pho-

toelectron spectroscopy (XPS) study was carried out at a

base pressure better than 5 × 10−9 Torr. Hereafter, for

simplicity we have named the sample Fe80Al20 (x = 0.2)

as Sample A, Fe70Al30 (x = 0.3) as Sample B, Fe60Al40 (x =

0.4) as Sample C, Fe50Al50 (x = 0.5) as Sample D and

Fe40Al60 (x = 0.6) as Sample E.

3. Results and Discussion

Figure 1 displays the XRD patterns of Fe1−xAlx (0.2 ≤ x

≤ 0.6) samples as a function of x. Only a narrow range of

diffraction angles is presented (2θ =25˚ - 55˚), but it

contains the main lines of all components and is repre-

sentative of the evolution that occurs. All the samples

show crystalline structure and textured mainly along (110)

direction, irrespective of the Al content. Figure 1(a)

shows the diffraction pattern of sample A (Fe80Al20). The

peak at 31.04˚ can be attributed to a new phase which

possesses the same structure type as the bcc Fe (Al). The

peak at 2θ = 39.3˚ of Al (111) is absent shows complete

Figure 1. XRD patterns of Fe1−xAlx samples as a function of

dissolution of Al with Fe, which further supports the

formation of Fe (Al) solid solution. Further, it is seen that

the fundamental peaks are consistently broadened and

shifted to lower angles with increase in x. The broaden-

ing of the peaks is due to the rapid reduction in average

crystallite size and an increase in lattice strain. It is found

that that average crystallite size decreases rapidly with x

particularly after x > 0.3 (see Table 1). The shifting of

the fundamental peaks to lower angles indicates an ex-

pansion of the lattice and presence of internal strain oc-

curs as a result of non-uniform alloy formation during

sample preparation. The variation in lattice parameter (a0)

as a function of Al content is shown in Table 1. This can

be understood if we take into account that the Al atoms

enter the Fe lattice producing a local dilatation due to

their larger size.

The diffraction pattern of Sample C shows some mark-

ed differences. In addition to reduction in peak intensity

and shift in peak position, a small peak at 39.04˚ corre-

sponds to Al (111) is also present in the pattern indicat-

ing that whole Al does not mix completely with Fe to

form uniform FeAl alloy. As already mentioned, Arc

melting process is a non-equilibrium process and has no

sufficient time to form stable and ordered phases, so

meta-stable and disordered phases will be formed. Fur-

ther considerable reduction in the peak intensity of Sam-

ples D and E with x clearly indicates the structural trans-

formation and formation of Al rich FeAl phases. The

average crystallite size in these cases are reduces drasti-

cally to 28 nm and 26 nm, respectively. After seeing the

results one may expect the formation of an off stoichio-

metric Fe3Al phase in the Fe-rich samples, while the

Al-rich compositions have both Al-rich phases and clus-

tering of Fe and Al atoms, respectively.

The corresponding hysteresis loops M (H) for all the

samples were obtained at room temperature using a VSM,

with a maximum applied field of 14 kOe (see Figure 2).

From the inset of Figure 2(a) it is clearly seen that Sam-

ple A show saturation with applied magnetic field indi-

cating the ferromagnetic nature of the sample with a

strong anisotropy leading to in plane easy direction of the

Table 1. Parameters obtained from XRD and VSM data

analysis for Fe1−xAlx alloy samples as a function of x.

Parameter

Composition

L (nm)a0 (Å)



(cm) Hc (Oe)Ms (emu)

x = 0.2 1162.879 93 1.7 3.22

x = 0.3 1112.883 117 2.9 1.88

x = 0.4 78 2.889 178 33.3 0.81

x = 0.5 28 2.894 165 71.7 0.045

x = 0.6 26 2.894 144 46.4 0.0046

S. RAJAN ET AL.

Figure 2. M-H curves of Fe1−xAlx samples as a function of x, (a) Sample A; (b) Sample C; (c) Sample D; and (d) Sample E.

magnetization. However, magnetization and room tem-

perature ferromagnetic state gradually disappears with

increasing x, up to x ≤ 0.3. On further increasing x, the

ferromagnetic state disappears more rapidly, becoming

paramagnetic for alloys with x ≥ 0.4. The samples with

high Al content (x ≥ 0.4) do not show saturation even

after applying the maximum applied field of 14 kOe.

This characteristic reflects the gradual development of

alloying process of Fe with Al. Al reduces the direct fer-

romagnetic interaction between Fe-Fe sites and at the

same time increase in anti-ferromagnetic interaction

could takes place, which reduces the magnetic moment

of Fe. Further, there is also a shift in the hysteresis loop

indicating of moment pinning due to highly disordered

and non-uniform nature of the sample. This loop shift

can be explained as originating from the phenomena of

exchange bias [15] associated with the exchange anisot-

ropy created at the interface between AFM-FM phases. It

is also found that coercivity (Hc) increases with x (from

~1.7 Oe to ~71.7 Oe), see Table 1. The increase of Hc

with the decreasing particle size may indicate that pin-

ning at the grain boundaries may be important origin of

the coercivity; indeed smaller grain size corresponds to

larger number of grain boundaries results enhancement in

the anisotropy which act as barriers to the motion of the

domain walls, increasing thus the coercive field. Table 1

also shows the saturation magnetization value of the

Fe1−xAlx samples as a function of x. The progressive de-

crease of the saturation magnetization is attributed to

competition between nearest neighbour Fe-Fe ferromag-

netic exchange and an indirect FeAl anti-ferromagnetic

interaction and the alloy become nonmagnetic around

0.5% Al. The system can be considered to be consisting

of a ferromagnetic phase dispersed in a non-magnetic

medium. The presence of spontaneous magnetization in

the system indicates that the magnetic grains are coupled

through exchange interactions. The exchange interaction

can be a Ruderman Kittele Kasuyae Yosida (RKKY)

type of interaction mediated by the itinerant electrons

present in the system or through a super-exchange type

of interaction due to the overlap of the electron wave

functions of the magnetic atoms with the Al electron

wave functions.

The value electrical resistivity of Fe1−xAlx samples as a

function of x is shown in Table 1. It is seen that the re-

sistivity increases sharply with increase in x and show a

S. RAJAN ET AL.

maximum value of 178 cm at x = 0.4. The resistivity

initially increases with concentration as it should since

increasing disorder, brought about by rising Al concen-

tration, contributes to increasing electron scattering. In

addition, the formation of different phases of FeAl is

responsible for high value of resistivity. Further increase

of Al yields a steep decrease in the electrical resistivity,

and the resistivity of the FeAl alloy containing 60 at.%

Al approaches a value of 144 cm. It is easy to under-

stand from the electronic structure of Fe and Al atoms

that 3s orbitals of Al are very close in energy to the 3d

orbitals of Fe, thus providing the grounds for a strong

overlap between Fe-3d and Al-3s as these elements are

alloyed. But important to note that the 3p electrons of Al-

lie at an energy higher than the 3d and 4s electrons of Fe.

Thus, as these elements are alloyed, one would expect a

charge transfer from Al-3p to Fe-3d, as the 4s orbitals of

Fe are full and this charge transfer affect the resistivity as

well as magnetic moment of Fe due to filling of the 3d

orbitals. Further addition of Al above 40 at.% value es-

sentially adds more electrons to a nearly free electron

conduction band and the resistivity falls sharply.

Figure 3 shows the core level and valence band (VB)

spectra of Sample B and D recorded after 90 min of sput-

tering. The recorded spectrum of Fe and Al shows some

significant differences, which are characteristics of the

oxidized state (see Figures 3(a) and (b)). The Fe-2p

spectrum is split by the 2p-spin–orbit effect into the 2p3/2

and 2p1/2 regions and continuous tail is caused by elec-

tron-hole pair excitations and is a signature of the metal-

lic states. Similarly, the set of typical Al-2p spectra con-

Figure 3. Fe-2p, Al-2p core levels and VB spectra of Sample

B and Sample D recorded after 90 min sputtering.

sist of spin orbit doublet peaks corresponding to Al-2P3/2

and Al-2p1/2 level states.

However, spectrometer resolution (0.8 eV) does not

allow distinction of their structure. Another important

point is that the relative intensity of Al-Ox peak is higher

than pure Al-2p peak indicating that oxygen is mainly

reacted with Al and partially with Fe (see Figures 3(a)

and (b)). The corresponding VB spectra show three dis-

tinct features; a broad Fe-3d band around 2.3 eV near the

Fermi level and other broad band lying deeper in energy

at 6.3 eV are attributed to π and σ molecular orbitals,

respectively formed by the hybridization of Al-3sp and

O-2p orbitals and Al-3s band at 10.8 eV (see Figure

3(c)). However in case of sample D, Fe-2P3/2 peak is ob-

served at 707.1 eV and is shifted by 0.3 eV towards

higher binding energy and Al-2p peak by 0.35 eV but

towards lower binding energy as compared to their ele-

mental Fe-2p and Al-2p core line (see Figures 3(d) and

(e)).

If we compare Samples B and D, we find some differ-

ences in the binding energy positions of Fe-2p and Al-2p

core lines. The binding energy of the Fe-2p peaks is

shifted towards higher binding energy and Al-2p towards

lower binding energy as compared to their element peak

position; also the shift in Sample B is different from

Sample D. The different shift in each case is due to the

formation of different phases of FeAl, i.e. Fe rich in

Sample B and Al rich in Sample D. Further in the VB

spectra of Sample D, one can see apart from 3d photo-

emission band of Fe at 1.8 eV, strong feature around 9.6

eV is also observed corresponding to Al-3s as is very

small in case of Sample B. Due to this, Fe-3d density of

states (DOS) are drastically modified and the emission

band is shifted towards lower binding energy i.e. towards

Fermi level. This is due to the strong hybridization of

sp-d states at Fermi level as a result of charge transfer

from valance Al-3p and 3s electrons to minority 3d or-

bital of Fe. Our results were matches well with the earlier

theoretical results.

4. Conclusion

The results conclude the formation of an off stoichio-

metric Fe3Al phase in Fe-rich samples, while the Al-rich

compositions have both Al-rich phases and clustering of

Fe and Al atoms The corresponding magnetic measure-

ments show decrease in saturation magnetization and

increase in resistivity with x and is attributed to dilution

of Fe moment in the Al rich compositions due to sp-d

hybridization.

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