RESEARCH PAPER

Interparticle and collective states of interactionsin mechanically milled Fe/CoO nanocomposites

Satya Prakash Pati • Dipankar Das

Received: 19 September 2013 / Accepted: 13 January 2014 / Published online: 29 January 2014

� Springer Science+Business Media Dordrecht 2014

Abstract The effect of high-energy milling on

magnetic properties of Fe/CoO nanocomposites has

been studied. Formation of pure phase polycrystal-

line sample having size in the range 15–20 nm is

confirmed from micro-structural analysis. Moss-

bauer spectroscopy indicates the formation of a

broad distribution of particle sizes in the samples

milled for higher duration. Enhancement in coer-

civity as well as exchange field is confirmed from

magnetic hysteresis measurements. High-field bifur-

cation in temperature dependent magnetization

measurement points toward the formation of spin-

glass like state. Wait time dependent magnetization

relaxation confirmed this possibility and is sup-

ported by the field cooled memory effect. Com-

plexity arises when the absence of memory effect is

observed in time dependent magnetic relaxation

measurement. Polydispersity and defects are argued

to be possible causes for observed spin-glass

behavior.

Keywords Nanocomposites � Mechanical

milling � Exchange bias � Mossbauer

spectroscopy � Magnetic properties

Introduction

Nanoscale magnetism especially those associated with

an assembly of magnetic nanoparticles is one of the

most attractive topics of research in recent time for

material scientists due to its emerging potential

applications in many technological fields and as well

as novel physical properties at the fundamental level

(Bader 2006). In order to enhance the technological

applicability, it is required to synthesize magnetic

materials having size in the nanometer range. But

when the size of a magnetic particle approaches in the

nano regime, it often loses its stable magnetic

orientation and becomes Superparamagnetic (SPM)

after a critical size limit (Knobel et al. 2008). This

limits its technological applications. Nonetheless, the

magnetic ordering in the SPM particle can be restored

by coupling it with an antiferromagnet (AFM) to

activate the exchange interaction at the interfaces of

the SPM and AFM materials (Skumryev et al. 2003).

This may enable the use of the SPM material in

magnetic recording media, spin valves, permanent

magnets, domain stabilizer in recording head, etc.

(Bader 2006; Printz 1998; Nogues et al. 2005). The

S. P. Pati (&)

Department of Electronic Engineering, Tohoku

University, Sendai 980-8579, Japan

e-mail: phy_satya@yahoo.co.in

D. Das

UGC-DAE Consortium for Scientific Research, Kolkata

Centre III/LB-8, Bidhannagar, Kolkata 700 098, India

123

J Nanopart Res (2014) 16:2278

DOI 10.1007/s11051-014-2278-5

exchange interaction among a FM and an AFM

material may lead to some interesting physical phe-

nomena such as exchange bias (EB), memory effect,

super-spinglass, etc. EB is manifested by a shift of

magnetic hysteresis loop or enhancement of coercivity

in field cooling conditions (Nogues et al. 2005).

Although a lot of experimental and theoretical studies

are performed in last few years to predict the

mechanism and the factors that determine the EB

value (Nogues et al. 2005; Nogues and Schuller Ivan

1999; Berkowitz and Takano 1999; Stamps 2000,

2002; Pati et al. 2010, 2011, 2012), a complete

understanding of the microscopic coupling still is not

clear. Thus, a systematic investigation along this line

is always of immense importance. Interparticle inter-

action among the constituents of a magnetic assembly

and its impact on overall magnetic properties is

another emerging subject of interest for the material

scientist. An assembly of noninteracting spin clusters

should give SPM, whereas the collective interaction of

spin clusters should result in a spin-glass like state

termed as super-spin-glass (SSG). Sometimes, the

uncompensated spins or broken exchange bonds

associated with the surface or interface of nanoparti-

cles may lead to the formation of spin-glass like state.

Although the explanation of SSG behavior is a

controversial topic, some literature is available to

distinguish it from SPM (Dormann et al. 1988;

Bandyopadhyay and Dattagupta 2006; Raj Sankar

et al. 2007).

To pursue the above problems, we have performed

a systematic study on the Fe/CoO nanocomposite

systems prepared by mechanical milling. Among all

transition metal oxides, CoO has highest magneto-

crystalline anisotropy (K = 2 9 105 erg/cm3 at

4.2 K) (Sievers and Tinkham 1967) with Neel tem-

perature close to the room temperature (TN = 293 K)

making it more efficient for technological applica-

tions. On the other hand, iron nanoparticles have

attracted considerable attention because of their

interesting magnetic properties and applications in

high density recording media (Dormann and Fiorani

1992). There are several reports on Fe/CoO multilayer

(Radu et al. 2006; Gruyters 2005; Wu et al. 2010;

Abrudan et al. 2008), but as per our knowledge limited

references are available on Fe/CoO particulate system

(Pati et al. 2013).In this article, we address the

interface magnetism and magnetic relaxation behavior

in the Fe/CoO nanocomposite system. Interfacial

phenomena like EB, memory effect in the field cooled

(FC) dc magnetization measurement has been

observed. In addition, time dependent relaxation

behavior at zero FC (ZFC) and FC conditions has

been measured to explain the assumption of SSG like

behavior.

Experimental details

Nanocomposites of Fe/CoO were prepared in a Fritsch

Pulverisette 7 planetary ball mill, using an 80 cm3

stainless-steel vial charged with 10-mm diameter

stainless-steel balls. Analytical grade Fe (alpha aesar,

99.998 %) and CoO (99.9 %) powders were used as

precursors. Calculated amount of powders (having 10

atomic wt% of Fe) were milled in argon atmosphere

under closed milling condition without any additives

(dry milling). However, different conditions and

parameters of mechanical milling have been used by

different users (Pati et al. 2011; Bhoi et al. 2010). The

milling was carried out at 300 rpm for 2, 5, 10, 20, and

30 h keeping ball to powder mass ratio 10:1. Hence

onwards a particular nanocomposite sample is iden-

tified by its ball-milling duration i.e., C2 denotes the

sample that was ball milled for 2 h. The X-ray

diffraction (XRD) pattern of the milled product was

obtained at room temperature, using a Bruker D8

Advance diffractometer with Co-Ka (k = 1.78897 A)

radiation in the Bragg–Brentano geometry. A recent

version of Rietveld program MAUD 2.14 was used to

carry out the refinement analysis. This program was

prepared specially for the simultaneous refinement of

various structural and micro-structural parameters

through a least squares method (Lutterott and Scardi

1990). Morphological analysis of the samples was

performed by a JEOL 2100 model high-resolution

transmission electron microscope (HRTEM) operat-

ing at 200 kV. Room temperature Mossbauer mea-

surements were carried out in a standard personal

computer (PC) based multichannel analyser with 1024

channels working in the constant acceleration mode. A

10 m Ci 57Co in Rh matrix was used as the radioactive

source. The system was calibrated with a high-purity

iron foil of thickness 12 lm. The spectra thus obtained

were deconvoluted by using LGFIT 2 program

(Meerwall 1975). Positron annihilation lifetime spec-

troscopy (PALS) measurements were carried out using

a fast–fast coincidence system consisting of two 1 inch

2278 Page 2 of 12 J Nanopart Res (2014) 16:2278

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tapered off BaF2 scintillators coupled to XP2020Q

photomultiplier tubes. The prompt time resolution of

the system using a 60Co source with the 22Na gate was

247 ps. The lifetime spectra were deconvoluted using

the code PATFIT 88 (Kirkegaard and Eldrup 1972).

DC magnetization studies were performed by using a

SQUID magnetometer (Quantum Design, MPMS XL

7).

Results and discussions

Structural and morphological analysis

The XRD patterns obtained for the prepared Fe/CoO

nanocomposites are shown in Fig. 1. The diffracto-

grams confirm the presence of pure phase a-Fe and

CoO in the nanocomposites. Corresponding peaks

were matched with JCPDS file no 851410 for Fe and

780431 for CoO. The peak at 52.6� and 77.3� (value of

2h) corresponds to the (110) and (200) crystallo-

graphic plane of Fe, while peak at 42.9�, 49.9�, 73.2�,

and 88.4� corresponds to (111), (200), (220), and (311)

plane of CoO, respectively. There is no substantial

shift in peaks was observed, however, peak broaden-

ing occurred with milling duration. The average

crystallite size and strain evaluated by the Rietveld

refinement analysis is tabulated in Table 1. The size

and strain of the nanocomposites were determined by

assuming the distribution in particle size. It is

observed that with increasing milling duration crys-

tallite size decreases and strain increases. A linear

dependency was found when the variation of crystal-

lite size and strain plotted against the milling duration.

The calculated crystallite size and microstrain for 30 h

milled sample are 20.03 ± 0.10 nm and 2.23 ± 0.12

(910-3) for Fe and 14.89 ± 0.06 nm and 4.12 ± 0.18

(910-3) for CoO. No sign of any other iron oxide or

any intermediate phase is found in our sample which

confirms the formation of pure phase nanocomposite.

A typical TEM micrograph and HRTEM image of

sample C30 are shown in Fig. 2c, d, respectively.

Figure 2a represents TEM image of sample C10. It can

be seen from the figure that during early milling, Fe

nanoparticles (dark contrast) are covered by the

lamella of CoO (light contrast). In some cases

interparticle aggregation occurred. As the concentra-

tion of Fe nanoparticles (10 at.wt%) in the sample is

Fig. 1 X-ray diffraction pattern of all Fe/CoO nanocomposites

J Nanopart Res (2014) 16:2278 Page 3 of 12 2278

123

less and are jacketed with CoO, the chance of

oxidation is low. Subsequent milling causes further

decrease in the size of the components. It is clear from

the micrograph that the particle size of C30 varies

from 10 to 25 nm having average particle size nearly

18 nm which is in agreement with the XRD data. A

histogram showing distribution of particle size was

shown in the inset of Fig. 2c, d. TEM image of C30

shows that the particles are nearly monodisperse in

nature. Moreover, some coagulation among nanopar-

ticles is also found. The selected area diffraction

(SAED) pattern of C30 (Fig. 2b) confirmed the pure

phase and polycrystalline nature of the sample. The d-

values calculated from the ring diameters confirmed

Table 1 Crystallite size and microstrain as calculated from the rietveld refinement

Sample name Crystallite size

of Fe (nm)

Crystallite size

of CoO (nm)

Micro-strain

of Fe (910-3)

Micro-strain

of CoO (910-3)

C2 46.37 ± 0.48 31.43 ± 0.12 1.02 ± 0.06 2.13 ± 0.09

C5 41.02 ± 0.27 28.10 ± 0.08 1.06 ± 0.10 2.42 ± 0.02

C10 35.44 ± 0.34 21.13 ± 0.09 1.68 ± 0.05 3.06 ± 0.11

C20 24.21 ± 0.14 16.65 ± 0.07 1.73 ± 0.07 3.82 ± 0.08

C30 20.03 ± 0.10 14.89 ± 0.06 2.23 ± 0.12 4.12 ± 0.18

Fig. 2 a TEM image of sample C10, b–d SAED pattern (asterisks and hash represents CoO and Fe), TEM, HRTEM image of C30,

respectively. A histogram showing particle distribution of C30 is displayed in the inset of (c) and (d)

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the presence of both Fe and CoO phases, and the

corresponding crystal planes are labeled. The calcu-

lated d-values 2.06 and 1.46 A are assigned to (110)

and (200) plane of Fe (JCPDS file No: 851410), while

2.4, 2.1, and 1.5 A are assigned to (111), (200), and

(220) plane of CoO (JCPDS file no: 780431),

respectively.

Hyperfine analysis

The hyperfine properties of all mechanically milled

nanocomposites were studied by 57Fe Mossbauer

spectroscopy. As it is a sensitive technique for

different grain size distribution, it is expected to

provide a clearer picture of the magnetic microstruc-

ture of the nanocomposites. It is well known that with

reduction of particle size to the nanometer regime,

thermal stability as well as the symmetry around the

iron atom reduces resulting in an appearance of a

quadrupole doublet in the spectrum. Figure 3 shows

the room temperature Mossbauer spectra of all sam-

ples. It is seen from Fig. 3 that samples of prolonged

milling duration (C20 and C30) have a doublet in

addition to a sextet. The appearance of this doublet in

the spectrum is assigned to the Fe particles undergoing

SPM relaxation (Bomati-Miguel et al. 2005; Cho et al.

2006; Seshadri et al. 1994). Hyperfine parameters of

all the nanocomposites extracted from the Mossbauer

data are tabulated in Table 2. The obtained values of

isomer shift (I.S.), quadrupole splitting (Q.S.), and

hyperfine field (Hint) of the sextets signify Fe atoms are

in blocked state. However, the evolution of doublet

with increasing milling duration implies the formation

of more number of SPM iron particles as the percent-

age area of doublet increases with milling hour.

Positron annihilation lifetime analysis

A general consequence of high-energy milling is that

the milling products are rich in interface and boundary

defects. To investigate the presence of such defects,

we performed positron annihilation lifetime studies of

all the samples. To obtain the best fit, positron lifetime

spectra of mechanically milled Fe/CoO nanocompos-

ites were deconvoluted with two lifetime components

s1 and s2 with relative intensities I1 and I20 respec-

tively. The shorter lifetime s1 corresponds to the

positrons annihilating at structural defects in the grain

boundaries. The longer lifetime s2 has been assigned

to free volumes at the intersection of interfaces (e.g.,

Fig. 3 Mossbauer spectra of all Fe/CoO nanocomposites

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Table 2 Mossbauer parameters of all Fe/CoO nanocomposites

Sample

name

ISa (mm/s)

(±0.01 mm/s)

QSb (mm/s)

(±0.01 mm/s)

FWHM

(mm/s)

Hintc (T) (±0.3 T) % area

C2 0.00 0.01 0.29 32.9 100

C5 0.00 0.01 0.30 33.0 100

C10 0.00 0.01 0.33 33.0 100

C20

Sd 0.01 0.00 0.32 33.0 88.7

De 0.33 0.65 0.37 – 11.3

C30

S 0.01 0.00 0.32 33.2 75.4

D 0.41 0.82 0.89 – 24.6

a IS = Isomer shiftb QS = Quadrupole splittingc Hint = Internal magnetic fieldd S1 = First sextete S2 = Second sextet

Fig. 4 Variation of positron lifetime components and relative intensities with milling duration

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triple junction). The mean lifetime of positrons

annihilated in the sample can be calculated by using

the relation sm ¼ s1I1þs2I2

I1þI2. Each lifetime component

represents annihilations from a specific defect,

whereas mean lifetime (sm) represents overall defects

present in the sample. sm is the most statistically

reliable parameter that can be obtained from PALS

studies. Figure 4 shows the variation of lifetime

components and respective intensities. The high value

of mean life time for prolonged milled samples implies

the evolution of defects in the sample. During milling,

the grain boundary defects are created in the sample

due to collision of ball and powder. It was observed

that the individual lifetime component (both s1 and s2)

increases up to 20 h milling and then decreases at

30 h. However, the mean life time increases contin-

uously from 10 to 30 h milled samples. As during

calculation of mean lifetime, the individual lifetimes

and their intensities are taken into consideration, the

explanations regarding the variation of intensities are

required. If we compare the values of intensities, the

contribution of s2 (which is coming from the free

volumes at the intersection of interfaces) is more in sm

(as I2 is greater than I1). A close inspection shows that

during initial stage of milling (up to 10 h), I1 increases

while I2 decreases. At intermediate milling (10 and

20 h), both intensities remain almost constant. How-

ever, for higher duration milled sample I1 decreases

and I2 increases. This variation of intensities reflects

that during milling the defects related to grain

boundaries and free volume at the interfaces varies

inconsistently.

Magnetization analysis

The results of room temperature magnetic hysteresis

measurements are shown in Fig. 5. All the extrapo-

lated data were tabulated in Table 3. It was observed

that the saturation magnetization (MS) decreases with

milling hour. However, remanence (MR) and coerciv-

ity (HC) increase with milling duration which is due to

the reduction of particle size. The increase in grain

boundaries with decreasing grain size results in

domain wall pinning which causes the coercivity as

well remanence enhancement (Pati et al. 2011).

In order to investigate the exchange coupling at the

interface of Fe and CoO, the magnetic hysteresis loops

were measured in FC and ZFC conditions at 10 K in a

3 T cooling field (Hcool). From Fig. 6, it is clear that

there is a significant horizontal shift in the hysteresis

loop along the field axis when the sample C30 is

cooled in FC condition. Furthermore, an enhancement

of coercivity is also noticed in the FC measurement.

These observations confirm the presence of exchange

coupling at the interface of ferromagnetic Fe and

antiferromagnetic CoO. The values of coercivity (HC)

and EB (HE) field are quantified as HC = (Hright -

Hleft)/2 and HE = -(Hright ? Hleft)/2, respectively,

where Hright and Hleft are the points of intersection of

the loop with the field axis. Also, a substantial vertical

shift of the FC hysteresis loop has been observed for

the sample, i.e., the magnitude of magnetization in the

positive direction at zero field [M (?H0max)] is larger

than that for the magnetization in opposite direction

[M (-H0max)].This vertical shift can be quantified

through the exchange magnetization parameter

ME = [M (?H0max) - M (-H0

max)]/2. The values of

HC, HE, and ME are calculated to be *843 Oe,

*222.3 Oe, and 3.93 emu/g, respectively, for the

sample C30 cooled to 10 K in the presence of

Hcool = 3T. It was seen that the value of EB field

increases with increasing milling duration (inset of

Fig. 6). With increasing milling duration the surface to

area ratio of the nanoparticles increases and effective

coupling takes place between FM and AFM compo-

nent which results in the enhancement of HE (Pati et al.

2011).

Temperature dependence ZFC and FC magnetiza-

tion curves at different applied fields for sample C30

Fig. 5 Room temperature M–H loops of Fe/CoO nanocom-

posites, and the inset shows the enlarged central part

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are shown in Fig. 7. It can be seen that the ZFC and FC

magnetization curves show a distinct irreversibility

behavior in all cases. The divergence of the ZFC and

FC curves indicates that a part of the total moment is

frozen and cannot be reversed even in a high external

field indicating a strong FM/AFM exchange coupling

at the interface (Pati et al. 2010). With increasing

applied field we found that the bifurcation temperature

or irreversibility temperature (Tirr) of the ZFC–FC

curves shifts toward lower temperature. The high-field

divergence may have resulted from the spin-glass like

(SGL) phase as reported earlier (Martinez et al. 1998;

Kodama et al. 1996, 1997; Zheng et al. 2004).

Generally, in a sufficiently high external magnetic

field the spin-glass freezing takes place either at lower

temperature or not at all because such a field can

destroy the spin-glass phases (Binder and Young

1986; Mydosh 1993). But in our case, the SGL phase

persists even at 2 T applied field. It is not surprising

that the SGL state can exist at the interfaces of a

nanocomposite comprising a FM and AFM materials

as the competition between the spins of a FM and an

AFM at the interface may lead to the development of

an SGL state. Moreover, possibility of SGL phase due

to surface fracture, disorder, and breakage of bonds

cannot be ruled out as the sample was prepared by the

ball-milling technique. To explore the possibility of

Table 3 Magnetization data of all Fe/CoO nanocomposites

Sample name Room temperature Zero field cooled Field cooled @ 3T

MSa

(emu/g)

MRb

(emu/g)

HCc (Oe) MS

(emu/g)

MR

(emu/g)

HC (Oe) MS

(emu/g)

MR

(emu/g)

HC (Oe) HEd (Oe)

C2 90.37 0.41 9.66 93.06 0.94 20.49 93.39 1.64 18.49 0.71

C5 87.7 0.82 17.97 90.33 1.66 35.5 91.29 2.63 40.46 4.04

C10 83.41 2.1 55.8 85.76 3.71 105.09 85.87 5.33 129.46 16.84

C20 83.06 7.44 160.9 84.91 11.99 328.84 85.07 17.85 420.37 82.64

C30 80.07 9.38 268.72 81.59 15.85 550.0 81.59 27.06 883.58 215.57

a MS = Saturation magnetizationb MR = Remanent magnetizationc HC = Coercive fieldd HE = Exchange field

Fig. 6 Magnified magnetic hysteresis loop of the nanocom-

posite C30 recorded in ZFC and FC condition at 10 K in a

cooling field of 3 T and the inset shows the variation of HE with

milling duration

Fig. 7 Temperature dependence of ZFC–FC magnetization of

sample C30 in different applied field

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SGL state, we carried out time dependent relaxation

measurements (aging) in both ZFC and thermo

remanent magnetization (TRM) protocol (Fig. 8).

For ZFC protocol, sample is cooled to 30 K in zero

field. Then, magnetization is measured against time

with an applied field 200 Oe for various waiting time

(tw = 0, 1, and 2 h). For TRM protocol, sample is

cooled to 30 K in 200 Oe field. Then, magnetization is

recorded as a function of time after switching off the

field for different waiting times. It can be seen that the

magnetization increases and decays exponentially for

ZFC and TRM protocols, respectively. We also

observed a considerable wait time dependence mag-

netization in both ZFC and TRM protocols, which is in

support of SGL behavior of the sample. Superpara-

magnets are expected to exhibit no wait time depen-

dence of magnetization for ZFC and weak wait time

dependence in TRM. However, Spin glasses show

aging in both ZFC and TRM protocol (Sasaki et al.

2005; Jonsson et al. 1995).

In order to clarify the possibility of SGL phase, (as

we observed high-field bifurcation in the ZFC–FC data

and wait time dependence magnetization in ZFC,

TRM protocol) we have employed a method called

field-cooled memory effect for the sample C30. As per

the protocol of the experiment, magnetization was

measured during cooling the sample from 100 K down

to 5 K in the presence of 200 Oe field. During this

cooling process, the sample temperature was kept

constant for 1 h at three intermediate temperatures 60,

40, and 20 K. The magnetic field was switched off

during these halt periods. After completion of the halt

period (1 h), magnetic field (200 Oe) was again

applied and measurement was carried out. The

obtained curve after this measurement is termed as

‘‘FC-waiting’’ as showed in Fig. 9. In the second cycle

of the experiment, magnetization was measured in the

presence of 200 Oe field during heating of the sample

from 5 to 100 K without any halt in the presence of

200 Oe field. Interestingly, we observed a step-like

behavior at the temperatures where the halts were

made during the previous measurement cycle. The

obtained curve is designated as ‘‘FCH-memory.’’ This

observed memory effect is the characteristic of a SGL

behavior. A very weak rejuvenation, i.e., small regain

in magnetization when the field is switched on after

aging, was observed which is also considered as an

evidence for super-spin glass behavior in nanoparti-

cles (Jonsson et al. 2005).

To complement FC memory experiment, we have

investigated the effect of temperature and field change

on the time dependent ZFC magnetization relaxation

measurement as suggested by Sun et al. (2003) and

adopted by many authors. Figure 10a displays mag-

netic relaxation at 30 K in 200 Oe after zero-field

cooling from 300 K. Initially, measurement was

carried out for time t1 at 30 K. Then, the sample is

cooled to 20 K in zero field and subsequently magne-

tization was recorded for time t2. Finally, the sample

was heated back to 30 K in 200 Oe field, and

magnetization was further recorded for time t3. It

can be seen that, when temperature and field are raised

back (30 K, 200 Oe), relaxation does not start from the

Fig. 8 Wait time dependence a ZFC magnetization and b TRM

at 30 K Fig. 9 Temperature dependent field cooled memory effect of

the sample C30

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point at which it left off in the previous magnetization

at same condition (inset of Fig. 10a). Also, the

relaxation process cannot restore the previous history

when the sample is temporarily heated instead of

cooling in a constant field (200 Oe). The data are

demonstrated in Fig. 10b for the measurements in

ZFC and FC conditions. It is noticed that, after

returning back to the previous temperature, magneti-

zation does not revert back to the same level in which

it was before the temporary heating. This indicates that

memory effect was not observed in the case of both

temporary cooling and heating processes in the

relaxation measurements. Re-initialization of relaxa-

tion process occurs in the case of both temperature

cycles. The absence of memories involved in the

relaxation dynamics is interesting and yet to be

extensively investigated experimentally and theoreti-

cally on various systems for understanding the basic

mechanism. But the presence of temperature depen-

dent FC memory effect and high-field bifurcation of

ZFC–FC data point to the SGL nature of the sample.

There are several reports on memory effects and aging

phenomena in the literature and various explanations

have been given regarding observation of the FC

memory effect in dc magnetization. Effect of poly-

dispersity is one of the main reasons as reported by

Chakraverty et al. (2005, 2006). Noninteracting or

weakly interacting SPM system may show similar

results as reported by Sasaki et al. (2005) and Tsoi

et al. (2005). The possibility of hierarchical organiza-

tion of metastable states resulting from interparticle

interactions may be one of the reason of this ‘‘memory

effect’’ as suggested by Sun et al. (2003). Since the

hierarchical organization requires a large number of

degrees of freedoms to be coupled, the memory effect

may not arise due to the thermal relaxation of

independent particles. In the present case, the distri-

bution of energy barrier due to the distribution of

particle size (polydispersity) may be a reason for the

observation of SGL behavior. Also, the role of defects

or fractures created on the surface of the nanoparticles

during milling cannot be ruled out.

Conclusion

In conclusion, nanocomposites comprising ferromag-

netic a-Fe and antiferromagnetic CoO are prepared

successfully by the high-energy ball-milling method.

XRD and TEM measurements confirm the pure phase

polycrystalline nature of the nanocomposites having

average particle size 15 nm for 30 h milled sample.

Reduction of crystallite size and increase of micro-

strain are observed with increasing milling duration.

Mossbauer measurements show the presence of SPM

doublet in higher hour milled samples, which repre-

sents the presence of a broad distribution in particles

sizes. Positron annihilation lifetime spectra confirm

the existence of interfacial defects in the nanocom-

posites. A significant horizontal shift in the hysteresis

loop was observed along the field axis when the

sample is cooled in FC condition which confirms the

presence of EB in the system. A high-field bifurcation

in temperature-dependent ZFC–FC magnetization

Fig. 10 a Magnetic relaxation with temporary cooling and field

change in ZFC condition, and the inset shows the relaxation

curves at 200 Oe only, b magnetic relaxation with temporary

heating in ZFC and FC conditions at 200 Oe field

2278 Page 10 of 12 J Nanopart Res (2014) 16:2278

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indicates the presence of SGL phase. Waiting time

dependence TRM and ZFC magnetization confirmed

the possibility of SGL phase. Observed temperature

dependent FC memory effect is also an indication of

SGL phase. But memory effect is absent in the time

dependent relaxation measurements. Hence, an ambi-

guity arises in understanding the possible reasons of

the observed SGL phase. Polydispersity and defects

are argued to be possible causes of observed SGL

behavior.

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