# Exchange Bias Demonstrated in Bulk Nanocomposites Processed by High Pressure
Torsion

M. Zawodzki  Erich Schmid Institute of Materials Science of the Austrian
Academy of Sciences, 8700 Leoben, Austria  L. Weissitsch  Erich Schmid
Institute of Materials Science of the Austrian Academy of Sciences, 8700
Leoben, Austria  H. Krenn  Institute of Physics, University of Graz, 8010
Graz, Austria  S. Wurster  Erich Schmid Institute of Materials Science of the
Austrian Academy of Sciences, 8700 Leoben, Austria  A. Bachmaier  Erich Schmid
Institute of Materials Science of the Austrian Academy of Sciences, 8700
Leoben, Austria

(January 2023)

###### Abstract

Ferromagnetic (Fe or Fe20Ni80) and antiferromagnetic (NiO) phases were
deformed by high pressure torsion, a severe plastic deformation technique, to
manufacture bulk sized nanocomposites and demonstrate an exchange bias, which
has been reported predominantly for bilayer thin films. High pressure torsion
deformation at elevated temperatures proved to be the key to obtain
homogeneous bulk nanocomposites. X-ray diffraction investigations detected
nanocrystallinity of the ferromagnetic and antiferromagnetic phases.
Furthermore, an additional phase was identified by X-ray diffraction, which
formed during deformation at elevated temperatures through the reduction of
NiO by Fe. Depending on the initial powder composition of Fe50NiO50 or
Fe10Ni40NiO50 the new phase was magnetite or maghemite, respectively.
Magnetometry measurements demonstrated an exchange bias in the high pressure
torsion processed bulk nanocomposites. Additionally, a tailoring of magnetic
parameters was demonstrated by the application of different strains or post-
process annealing. A correlation between the amount of applied strain and
exchange bias was found. The increase of exchange bias through applied strain
was related to the microstructural refinement of the nanocomposite. The
nanocrystalline maghemite was considered to have a crucial impact on the
observed changes of exchange bias through applied strain.

Keywords: severe plastic deformation; high pressure torsion; nanocomposite;
superior hardness; microstructural characterization; magnetic properties;
hysteresis; exchange bias

##  1 Introduction

The introduction and application of nanomaterials is indispensable for
progress in technological development. The use of nanomaterials is already
well advanced and ranges from applications in medicine (targeted drug delivery
for cancer treatment) [39], catalysis (optimisation of the catalysts
nanostructure to improve the selectivity and efficiency of catalytic processes
[29], such as the enhanced reduction of volatile organic compounds[52]) , gas
sensors [30], microelectronics [44] to engineering [32].
Metallic nanomaterials have outstanding physical properties, such as superior
strength, superior hardness and enhanced wear resistance, and attracted the
attention of engineers to reduce weight to improve efficiency for
transportation vehicles [47, 32].
The importance of magnetic materials for power efficiency is commonly known.
Material scientist focus here on the improvement of the nanostructure of soft-
or hard magnetic materials to enhance the magnetic properties [20, 10].The
important role of nanocrystalline microstructure can be demonstrated for hard
magnetic materials. Here, highly textured nanocrystals are considered to be
the key to push the energy product ($BH$) of a hard magnetic material to its
theoretical maximum of
$BH\textsubscript{max}=1/4\mu\textsubscript{0}M\textsubscript{S}\textsuperscript{2}$
($\mu\textsubscript{0}$…vacuum permeability and MS…saturation magnetisation)
[10, 43]. An increase in $BH$ of the hard magnetic material leads to a
reduction of weight and size of permanent magnet motors/generators and this
addresses directly the need for power efficiency [19].
Although synthesis of nanocrystalline material has been challenging, a variety
of methods have been established to obtain the desired nanocrystalline state,
(e.g. inert gas condensation, electrodeposition, mechanical alloying,
crystallisation out of an amorphous state and severe plastic deformation
(SPD)) [28]. Regarding synthesis, SPD methods, especially high pressure
torsion (HPT), offer unique advantages. For example, it is possible to process
a bulk nanocrystalline sample based on powder blends, which allows a large
variety of phase combinations. With HPT it is possible to synthesise
supersaturated solid solutions [23] or influence the magnetic properties of
the nanocrystalline sample through the application of strain [51, 2, 48, 27].
The focus of this study was on the use of rare-earth free phases due to their
abundant availability. Therefore, the choices for a suitable ferromagnetic
(FM) material were limited to Fe, Ni or FeNi-alloys. Fe was chosen, because of
the vast base of experience available concerning HPT-deformation [34]. A
suitable alternative FM-phase is the $\gamma$-Fe20Ni80-alloy, which possesses
lower MS compared to Fe and far larger domain wall width than Fe or Ni. Those
two properties of $\gamma$-Fe20Ni80 are regarded as beneficial to the
enhancement of exchange bias (Heb).
The Heb was first observed by Meiklejohn and Bean on Co-CoO nanoparticles and
introduced as ‘new unidirectional anisotropy’ [26]. Heb originates from a FM
spin exchange coupling mechanism between the surface spins of adjacent
antiferromagnetic (AFM) and FM phases and allows the preservation of an
external magnetic field direction through the alignment of AFM surface spins
during field cooling (FC) below the Neél temperature (TN). This phenomenon
causes a biased shift of the hysteresis in the FC direction.
A further aim of this study was to use a material, which exhibits AFM
properties at room temperature (RT). NiO has the benefit of possessing a face-
centred cubic crystal structure with a smaller lattice parameter mismatch to
Fe or $\gamma$-Fe20Ni80-alloy compared to other possible AFM materials (e.g.
FeS or $\alpha$-Fe2O3), and is also AFM far above RT; it was therefore the
material of choice.
Until recently, such material combinations have mainly been realised by thin
film deposition techniques to study the interfacial phenomenon of Heb [6], due
to its technical application in magnetic read heads for hard-disc drives [9].
Consequently, previous investigations have been limited to 2D structures. A
first successful synthesis by HPT of bulk composites, which possess an Heb,
has been reported for the Co-NiO system [27].
The aim in this study was to obtain a bulk nanocomposite possessing enhanced
magnetic properties, Heb and be able to demonstrate a tailoring of magnetic
parameters via applied strain. Furthermore, to investigate in an initial trial
the new bulk nanocomposite regarding their nanostructure and the influence of
the nanostructure on the magnetic properties.
A material synthesis with a metallic and an oxide phase can be a challenging
task, especially for HPT, if a homogeneous microstructure is desired. Although
oxide ceramics (i.e. Al2O3, ZrO2, TiO2 and Y2O3) have been processed
previously with HPT [13], the present study has investigated for the first
time the deformation process in combination with nanostructural
characterisation and magnetic characterisation systematically to provide a
detailed description of the FM-AFM bulk nanocomposite.
A further aim was to attain deeper insight into the HPT-deformation process of
NiO itself and its interaction with the mechanically softer FM-phase. The
importance of HPT processing at elevated temperatures was demonstrated to
synthesise homogeneous bulk nanocomposites. In addition to a tailoring of
magnetic properties, an unusually high Vickers microhardness was detected,
when the NiO-phase was preserved during deformation.

##  2 Materials and Methods

Commercially available powders (Fe - Mateck Fe-99,9% 100+200 mesh, Ni - Alfa
Aesar Ni 99,9% 100+325 mesh, NiO - Alfa Aesar NiO 99,8 % 325 mesh) were used.
All powder compositions are labelled in –at%, if not otherwise noted. Powders
were stored and prepared inside a glove box filled with Ar-atmosphere. For
powder consolidation with HPT an airtight capsule was used to prevent the
powder from oxidation. The capsule itself enclosed the HPT-anvils, which can
move freely only in axial direction for powder blend compaction. The obtained
pellet was used later on in a second step for HPT-deformation Figure 9. The
processed sample discs had dimensions of ($\varnothing 8\times 0.6$) $mm$
[21]. The process parameters were the following: applied hydrostatic pressure
of 6 $GPa$, rotation speed $\omega=1.25$ $min\textsuperscript{-1}$ and
deformation temperatures (Tdef) of 200-300°C. The Tdef was provided by
inductive heating of both anvils during the processing, and the temperature
was controlled by a pyrometer [34]. Samples deformed at RT were cooled with
pressurised air to ensure temperature stability of the sample. Equivalent von
Mises strain ($\epsilon$vM) was calculated with
$\epsilon\textsubscript{vM}=(2\pi Nr)/(\sqrt{3}t)$ ($N$…applied rotations,
$r$…radius and $t$…sample thickness)[21].
Scanning electron microscopy (SEM; LEO-1525 Carl Zeiss GmbH) was used in
backscattered electron mode (BSE; Model 7426, Oxford Instruments plc) with an
acceleration voltage of 20 $kV$. Energy dispersive X-ray spectroscopy (EDX;
XFlash 6–60, Bruker) analysis with SEM was performed at 5 $kV$ to maximise
lateral resolution. In order to assign the corresponding radial position to
the SEM images, the ends of the examined HPT half discs were determined during
the SEM examination. $r$ $\sim$0 $mm$ was defined as the half distance between
both ends, and $r$ $\sim$3 $mm$ was determined by the relative distance to $r$
$\sim$0 $mm$. $r$ $\sim$3 $mm$ was cross-checked by measuring the relative
distance to the end of the disc, which should be $1$ $mm$. Hardness
measurements were performed with a Micromet 5104 from Buehler with an
indention load of 500 $g$ along the radial direction every 250 $\mu m$.
For X-ray diffraction (XRD), samples were polished on the top side and
analysed with a Bruker D2 Phaser (Co-K$\alpha$ source). XRD data from
synchrotron measurements were collected in transmission mode at Deutsches
Elektronen-Synchrotron (DESY) at the beamline P07B (high energy materials
science) with a beam energy of 87.1 $keV$ and a beam size of $0.5\times 0.5$
$mm\textsuperscript{2}$. All wide angle X-ray scattering (WAXS) measurements
were done in axial direction [48] and the data were gained with a Perkin Elmer
XRD 1621 detector. Peak analysis was done with a self-written script in Octave
based on the pseudo-Voigt method to determine integral breadth and coherent
scattering domain size (CSDS) [46] via Scherrer-relation for spherical
crystallites. As XRD-reference pattern the American Mineralogist Crystal
Structure Database (AMCSD) is used (Fe: AMCSD 0012931, Ni: AMCSD 0012932, NiO:
AMCSD 0017028, $\gamma$-Fe2O3: AMCSD 0020516, $\alpha$-Fe2O3: AMCSD 0000143
and Fe3O4: AMCSD 0009109).
The annealing experiments were done within a vacuum furnace (XTube, type-1200,
Xerion Advanced Heating). For sample preparation, a deformed HPT-disc was cut
in quarters. One quarter of the sample material was used to obtain probing
material for magnetometry measurements from the same radial position, and the
other quarters for microstructural characterisation. For the annealing
experiment itself, one full quarter and one sample for magnetometry
measurements was heated up to 450°C or 550°C in a vacuum of 10-6 $mbar$ or
better and kept at the specific temperature for one hour.
Magnetometry samples were prepared with a diamond wire saw and samples were
cut out at desired radial position, having a volume of $\sim$2 $mm$3 or less
[48]. Due to the sample dimensions, the applied strain, as mentioned for the
magnetometry measurements, is an average value and related to the centre of
mass of the measured sample.
Field cooling was performed with the applied magnetic field parallel to the
radial direction from above the corresponding Néel-temperature from NiO of TN
$\sim$524 $K$ [45] to RT inside a vacuum chamber at 10-4 $mbar$ or better. An
electro magnet provided a magnetic field of $\sim$10 $kOe$. The field cooling
was subsequently continued below RT, realised by a superconducting quantum
interference device (SQUID) from Quantum Design MPMS-XL-7 from RT to 8 $K$ at
10 $kOe$, which was further used to determine magnetic hysteresis loops for
the Fe10Ni40NiO50 composition. Measurements were executed at 8 $K$ with a
maximum magnetic field strength of $\pm$70 $kOe$ to ensure magnetic
saturation. After determining HC1 and HC2, the Heb can be calculated with
$H\textsubscript{eb}=(H\textsubscript{C1}+H\textsubscript{C2})/2$. The
coercivity of the symmetric hysteresis is determined by
$H\textsubscript{C}=((-)H\textsubscript{C1}+H\textsubscript{C2})/2$. MS was
approximated with a linear fit from the saturated branches of hysteresis to
obtain MS at the intercept with the ordinate.

##  3 Results

###  3.1 Hardness and microstructural evolution

####  3.1.1 Fe50NiO50 composition

The first composition investigated was the binary Fe50NiO50 (Fe42.8NiO57.2 in
-wt%) blend. Figures 1a and b present SEM micrographs at $r$ $\sim$0 $mm$ and
$r$ $\sim$3 $mm$. Both samples were processed with 100 rotations and
deformation temperature was increased from RT (Figure 1a) to 300°C (Figure
1b). This yielded an applied strain for Tdef = RT of $\epsilon$vM $\sim$1760
and for Tdef = 300°C a $\epsilon$vM $\sim$1490 at $r$ $\sim$3 $mm$, due to a
slight difference in height. For the sample processed at RT, EDX measurements
confirmed that the dark phase seen in BSE-contrast consisted of NiO, while the
bright phase was identified as Fe-phase.
When processed at RT, both phases formed lamella structures (Figure 1a).
Although the applied strain was high, the expected refinement of the two-phase
microstructure was not observed. Within the phases, a grain refinement was
evident at $r$ $\sim$3 $mm$. Vickers microhardness measurements detected an
increase of microhardness, (Figure 1c) and confirmed the ongoing refinement,
which was seen in SEM within the phases. This observation indicated that the
applied strain had acted in the composite to a certain degree. However,
whether the total amount of applied strain had been absorbed by the composite
through deformation was in question. The deformation at RT of samples with NiO
of 50-at% (Figure 1a) or higher was considered to be unsuccessful because the
samples had inhomogeneous microstructure and the FM-phase dimensions were too
large for observing an Heb.
Deformation at elevated temperatures improved the deformation behaviour of the
nanocomposites and led to a homogeneous microstructure (Figure 1b). SEM
micrographs indicated that a homogeneous two-phase microstructure had been
formed, having microstructural refinement from $r$ $\sim$0 $mm$ to $r$ $\sim$3
$mm$. The large NiO structures, which were observed for Tdef = RT, were not
observed for Tdef = 300°C in the SEM micrographs along the sample radius.
Here, the dark phase, which is shown in BSE-contrast, was determined by EDX
measurements in combination with XRD-results (Figure 2) as Fe3O4 and the
bright phase was identified as a FeNi-phase.
Vickers microhardness measurements detected a large difference in hardness
between samples deformed at RT and 300°C (Figure 1c). For the Tdef = 300°C
composite, hardness values were considerably higher than for the RT sample,
varying from 6 $GPa$ at $r$ $\sim$0 $mm$ ($\epsilon$vM $\sim$0) to 6.8 $GPa$
at $r$ $\sim$3 $mm$ ($\epsilon$vM $\sim$1740). Above $r$ $\sim$1.5 $mm$ the
Tdef = 300°C composite became microstructurally saturated according to Vickers
microhardness measurements.

Figure 1: SEM micrographs in BSE mode depict samples deformed at RT (a) and
300°C (b) at $r=0$ $mm$ and $r=3$ $mm$. EDX spectroscopy identified the dark
phase as NiO for (a) or Fe3O4 for (b). (c) Vickers microhardness of Fe50NiO50
samples along radial direction.The sample processed at 300°C had a higher
saturation microhardness compared to the sample deformed at RT. (d) Vickers
microhardness values form $r=2.5$ $mm$ to $r=3.75$ $mm$. Higher deformation
temperatures led to a homogeneous microstructure and therefore higher hardness
values. (e) SEM micrographs in BSE mode gained of Fe50NiO50 samples at the
outer radii ($r$ $\sim$3 $mm$). An onset of homogenisation of microstructure
could be seen at 225°C and gradually improved towards 275°C. The micron bar
applies to all micrographs in the same row.

Further deformation experiments were conducted at deformation temperatures
between RT and 300°C (Tdef = 200°C, 225°C, 250°C and 275°C) to gain insight
into the onset of the microstructural homogenisation. Figure 1e displays SEM
micrographs at $r$ $\sim$3 $mm$. At Tdef = 200°C, the microstructure was
identical to that at RT; however, at Tdef = 225°C, the large NiO structures
had become more and more refined through the application of strain. This
process continued for Tdef = 250°C and Tdef = 275°C, gradually extending from
outer radius to inner radius. At Tdef = 275°C, the microstructure is
comparable to the Tdef = 300°C samples at $r$ $\sim$3 $mm$, but still large
NiO structures were observed for Tdef = 275°C at $r\leq 2$ $mm$ (not shown).
The Vicker microhardness values that were obtained, reflected the observations
from SEM. In Figure 1d hardness results at the outer radius are summarised.
The inhomogeneous microstructure of the sample deformed at Tdef = 200°C had a
lower microhardness (5 $GPa$) than those of samples deformed at temperatures
Tdef $\geq$225°C, which were approximately 6 to 7 $GPa$ at $r$ $\sim$3 $mm$
and similar to the hardness values of the sample deformed at Tdef = 300°C. All
samples processed at elevated temperatures exhibited significantly higher
Vickers microhardness values than the sample processed at RT.
Phase analysis with XRD detected Fe and NiO phases for the samples deformed at
RT and 200°C with an onset of Fe3O4 formation at Tdef =200°C (Figure 2). An
increase of deformation temperature from Tdef = 225°C to Tdef = 300°C yielded
a continuously growing amount of Fe3O4, caused by the reduction of NiO through
Fe. Ni became available to build the $\gamma$-FeNi phase, which was detected
in the XRD-scans. The content of Fe within the $\gamma$-FeNi phase could be
approximated through the shift to lower q-values. This shift is caused by the
substitution of Fe-atoms within the Ni lattice, resulting in an enlargement of
lattice parameter. This enlargement is caused by the higher atomic radius of
Fe compared to Ni [33]. The shifted peaks of $\gamma$-FeNi phase for Tdef =
300°C implied a composition of approximately $\gamma$-Fe39Ni61.
The quality of the XRD patterns and similarity of Fe3O4 and maghemite
($\gamma$-Fe2O3) patterns allowed a determination of FexOy phase only at
higher qz $\geq$3.5 Å-1. The peaks at qz = 3.889 Å-1 and qz = 4.235 Å-1, for
example, are in accordance with the Fe3O4 pattern. Presence of $\gamma$-Fe2O3
could not be excluded, but apparently Fe3O4 dominated the XRD-pattern;
$\gamma$-Fe2O3 was therefore omitted in the following consideration for the
Fe50NiO50 nanocomposite.
The assumption that the $\alpha$-Fe phase had been completely consumed by the
formation of Fe3O4 or $\gamma$-FeNi, allowed estimation of the residual NiO.
Using the known composition of $\gamma$-Fe39Ni61 phase for the sample deformed
at Tdef = 300°C, yielded a residual NiO-phase of maximum 16-wt%, which should
have been left after synthesis inside the sample, but could not be resolved by
the laboratory XRD-equipment due to overlapping XRD-peaks (Figure 2). Although
the sample deformed at Tdef = 300°C had a promising microstructure, this
nanocomposite was considered unsuitable for observing a significant Heb in
subsequent magnetometry measurements because it did not contain large amounts
NiO according to the XRD results.

Figure 2: XRD-data is displayed of the Fe50NiO50 sample, which was measured
with laboratory equipment. The formation of Fe3O4, through the reduction of
NiO by Fe, was seen for samples deformed at Tdef $\geq$225°C. $\alpha$-Fe was
not detectable at Tdef $\geq$275°C.

The CSDS was calculated via Scherrer-relation of Fe50NiO50 sample at RT and
Tdef = 300°C displayed nanocrystallinity, whereby Fe3O4 had the largest CSDS.
Considering residual strain with the Williamson-Hall (WH) method [46] and
comparing results, strain in the Fe3O4-phase is almost half compared to
$\gamma$-FeNi. Strain-free CSDS estimations yielded approximately 45 $nm$ for
both analysed phases at Tdef = 300°C (Table 1).

Table 1: Summary of peak analysis executed for Fe50NiO50 sample deformed at RT and 300°C. <Dx> is the average crystallite size or CSDS, Dx-WH represents the strain-less CSDS result and $\epsilon$x-WH residual strain in the crystallite, both values are obtained via the WH method. |  <DFe> / [$nm$] |  DFe-WH / [$nm$] |  $\epsilon$Fe-WH |  <DNiO> / [$nm$] |  DNiO-WH / [$nm$] |  $\epsilon$NiO-WH
---|---|---|---|---|---|---
|  |  |  |  |  |
Fe50NiO50 |  11$\pm$2 |  40 |  2$\cdot$10-2 |  11$\pm$2 |  30 |  2$\cdot$10-2
@Tdef =RT |  |  |  |  |  |
|  <DFeNi> / [$nm$] |  DFeNi-WH / [$nm$] |  $\epsilon$FeNi-WH |  <DFe3O4> / [$nm$] |  DFe3O4-WH / [$nm$] |  $\epsilon$Fe3O4-WH
|  |  |  |  |  |
Fe50NiO50 |  12$\pm$2 |  45 |  3$\cdot$10-2 |  19$\pm$7 |  43 |  1$\cdot$10-2
@Tdef=300°C |  |  |  |  |  |

####  3.1.2 Fe10Ni40NiO50 composition

To reduce the loss of NiO at 300°C, another composition of Fe10Ni40NiO50
(Fe8.4Ni35.4NiO56.2 in wt%) was tested with the idea to conserve the NiO-phase
by alloying Ni initially with Fe to prevent the reduction of NiO.
The microstructures depicted in Figure 3a were collected from two samples; one
has been deformed with 30 rotations the other with 125 rotations; all other
parameters were equal. For the micrographs shown, the applied strain was
$\epsilon\textsubscript{vM}$ $\sim$90 ($\triangleq$ $r$ $\sim$0.5 $mm$) and
$\epsilon\textsubscript{vM}$ $\sim$610 ($\triangleq$ $r$ $\sim$3.5 $mm$) for
the sample with 30 rotations. For the sample with 125 rotations, the applied
strain was equivalent to: $\epsilon\textsubscript{vM}$ $\sim$270 ($\triangleq$
$r$ $\sim$0.375 $mm$), $\epsilon\textsubscript{vM}$ $\sim$1280 ($\triangleq$
$r$ $\sim$1.75 $mm$) and $\epsilon\textsubscript{vM}$ $\sim$2560 ($\triangleq$
$r$ $\sim$3.5 $mm$). Due to the homogeneous microstructure, a uniform
deformation of the sample material was assumed and applied strain was
therefore used rather than the radius. The dark phase was identified as NiO
and the bright phase as Ni. The microstructure was different from the
Fe50NiO50 sample deformed at Tdef = 300°C (for comparison Figure 1b). Through
the preservation of NiO, a fragmentation process of NiO was realised, leading
to a highly diverse Ni and NiO nanocomposite structure. Additionally, SEM
micrographs demonstrated a refinement of the nanocomposite microstructure with
an increase of applied strain.
Vickers microhardness measurements along the radial direction demonstrated a
microstructural saturation at $\epsilon\textsubscript{vM}$ $\sim$700 of 9.3
$GPa$ in microhardness for the Fe10Ni40NiO50 sample after 125 rotations,
(Figure 3b). Beyond $\epsilon\textsubscript{vM}$ $\sim$700, further
deformation caused only a small increase in microhardness of approximately 0.6
$GPa$ to 9.9 $GPa$, which correlated with the minor refinement seen in SEM
between $\epsilon\textsubscript{vM}$ $\sim$1280 to
$\epsilon\textsubscript{vM}$ $\sim$2560\. Vickers microhardness values of 9.3
to 9.9 $GPa$ were significantly higher than previously reported results from
similar nanocrystalline materials [42, 1].

Figure 3:  (a) BSE micrographs of the Fe10Ni40NiO50 samples. Microstructural
refinement was detected from left to right. Two samples were deformed to
obtain the depicted micrographs. First samples was deformed with 30 rotations
and micrographs of an applied strain of $\epsilon\textsubscript{vM}$ $\sim$90
to $\epsilon\textsubscript{vM}$ $\sim$610 were taken. Second samples was
deformed with 125 rotations and micrographs of an applied strain of
$\epsilon\textsubscript{vM}$ $\sim$270, $\epsilon\textsubscript{vM}$
$\sim$1280 and $\epsilon\textsubscript{vM}$ $\sim$2560 were taken. The micron
bar applies to all micrographs in the same row. (b) Vickers microhardness of
both Fe10Ni40NiO50 samples deformed at 300°C.

Synchrotron WAXS investigations were conducted along the axial direction for
following applied strains: $\epsilon\textsubscript{vM}$ $\sim$270,
$\epsilon\textsubscript{vM}$ $\sim$1280, and $\epsilon\textsubscript{vM}$
$\sim$2560\. Detected phases were NiO, Ni and $\gamma$-Fe2O3. However, the
$\gamma$-Fe2O3 peaks could not be directly identified. The comparison of WAXS
patterns of the as-deformed and post-process annealed samples (Section 3.1.3)
showed a slight shift from the Fe3O4 to the $\gamma$-Fe2O3 reference pattern
(Figure 4). This shift became more evident through regarding the calculated
peak centres.
According to the WAXS measurements, the FM-phase consisted of solely Ni. An
expected substitution of Fe was thus not detectable with WAXS and implied that
the main part of Fe was consumed by the formation of $\gamma$-Fe2O3. The WAXS
results allowed estimation of residual amount of NiO within sample and yielded
the following composition: Ni48.6NiO39.4 and ($\gamma$-Fe2O3)12, all in wt%.
Integral peak breadth analysation with the Scherrer-relation showed minor
changes of CSDS for Ni and NiO, but the CSDS of $\gamma$-Fe2O3 reduced by 50%
with applied strain (Table 2). Residual strain was approximated via the WH
method and for NiO was approximately a factor of 10 higher than for Ni.

Figure 4:  Synchrotron WAXS results are shown for different applied strains of the Fe10Ni40NiO50 sample. Starting from bottom: $\epsilon\textsubscript{vM}$ $\sim$270, $\epsilon\textsubscript{vM}$ $\sim$1280 and $\epsilon\textsubscript{vM}$ $\sim$2560\. The two annealed samples are at the top: atop the sample annealed at 550°C and below the sample annealed at 450°C. At low q-values the development of $\gamma$-Fe2O3 peaks were visible. $\gamma$-Fe2O3-peaks became weaker with high applied strains due to a reduction of CSDS.  Table 2: A summary of integral peak breadth analysis is shown, which was done for the Fe10Ni40NiO50 sample deformed at Tdef = 300°C. <Dx> is the average crystallite size or CSDS, Dx-WH represents the strain-less CSDS result and $\epsilon$x-WH residual strain in the crystallite both values were obtained by the WH method. |  <DNi> / [$nm$] |  DNi-WH / [$nm$] |  $\epsilon$Ni-WH |  <DNiO> / [$nm$] |  DNiO-WH / [$nm$] |  $\epsilon$NiO-WH |  <D$\gamma$-Fe2O3> / [$nm$]
---|---|---|---|---|---|---|---
$\epsilon$vM$\sim$270 |  18$\pm$2 |  21 |  3$\cdot$10-3 |  11$\pm$3 |  25 |  2$\cdot$10-2 |  12$\pm$2
$\epsilon$vM$\sim$1280 |  15$\pm$5 |  19 |  2$\cdot$10-3 |  11$\pm$2 |  20 |  1$\cdot$10-2 |  9$\pm$4
$\epsilon$vM$\sim$2560 |  15$\pm$4 |  18 |  2$\cdot$10-3 |  11$\pm$2 |  19 |  1$\cdot$10-2 |  6$\pm$2
ann.450°C |  21$\pm$2 |  22 |  1$\cdot$10-3 |  14$\pm$3 |  22 |  9$\cdot$10-3 |  8$\pm$2
ann.550°C |  24$\pm$2 |  25 |  0.5$\cdot$10-3 |  17$\pm$3 |  26 |  7$\cdot$10-3 |  14$\pm$2

####  3.1.3 Annealing of Fe10Ni40NiO50 composition

Because of the complex microstructure of the nanocomposite after deformation
the effects of post-process annealing were investigated to understand the
influence of phase growth and phase interface morphology on the magnetic
properties of the Fe10Ni40NiO50 sample better. Hence, sample material with a
large amount of applied strain $\epsilon\textsubscript{vM}$ $\sim$1150 ($r$
$\sim$2.75 $mm$) was used to have .
The SEM micrographs are displayed in Figure 5 for the following states: as-
deformed, annealed in vacuum at 450°C and 550°C. For the as-deformed state,
the complex phase interface was assumed to have large variations in topography
at a nanometre scale, which was expected to simplify through annealing. The
contrast in BSE mode has been enhanced for the annealed samples. This
enhancement indicated a phase growth of the bright and dark phases. The
annealing led to a reduction of phase interface, which was expected to be
predominantly driven by the optimization of Gibbs free energy.
XRD-peak analysation affirmed the rise of CSDS for all three detected phases
(Table 2), whereby $\gamma$-Fe2O3 had the highest increase in relative
crystallite size of approximately two times to its initial CSDS in the as-
deformed state. The relative CSDS growth of Ni and NiO-phase during annealing
was significantly smaller.

Figure 5:  SEM micrographs were done in BSE mode of as-deformed and annealed
samples of the Fe10Ni40NiO50 composition. The annealing initiated a phase
growth and change of interface morphology for the samples annealed at 450°C
and 550°C. The micron bar applies to all micrographs in the same row.

###  3.2 Magnetic characterisation

For SQUID investigations, the promising microstructures of the Fe10Ni40NiO50
samples were chosen. The hysteresis loops shown in Figure 6a have a
continuously growing Heb, which could be correlated to the refinement of
microstructure through the application of strain. The applied strain changed
the hysteresis loop characteristic in general. As MS decreased, coercivity
rose. The coercivity nearly doubled from $\epsilon\textsubscript{vM}$ $\sim$90
with HC = 460 $Oe$ to $\epsilon\textsubscript{vM}$ $\sim$2560 having HC = 842
$Oe$. Magnetic remanence (MR) rose from MR = 18.5 $emu/g$ to MR = 23.3 $emu/g$
at $\epsilon\textsubscript{vM}$ $\sim$610 and became independent of the
applied strain for further deformation. Heb started at
$\epsilon\textsubscript{vM}$ $\sim$90 with Heb = -52 $Oe$ and ended at
$\epsilon\textsubscript{vM}$ $\sim$2560 with Heb = -243 $Oe$ (Table 3). As
mentioned previously, MS reduced from MS,$\epsilon\textsubscript{vM}$ $\sim$90
= 37.2 $emu/g$ to MS,$\epsilon\textsubscript{vM}$ $\sim$2560 = 27.5 $emu/g$.
These changes in MS and Heb appeared to be interrelated and might have had
their origins in the detected changes in CSDS, enlargement of phase interfaces
and the generation of $\gamma$-Fe2O3.
The slight increase in Vickers microhardness detected at $r$ $\leq$1.25 $mm$
(Figure 3b) was therefore confirmed through the continuously growing Heb for
such high amounts of applied strain and was in accordance with the impression
of an ongoing refinement gained by SEM and XRD investigations.
Figure 6b shows hysteresis loops of the annealed samples at 8 $K$. The
annealing treatment caused a drastic reduction of Heb from
Heb,$\epsilon\textsubscript{vM}$ $\sim$1150 = -200 $Oe$ to Heb,ann. 550°C =
-23 $Oe$ and MR (Table 3), whereas MS became larger, changing from
MS,$\epsilon\textsubscript{vM}$ $\sim$1150 = 32.3 $emu/g$ to MS,ann. 550°C =
36.6 $emu/g$. The rise of MS with annealing treatment correlated with the
detected growth of CSDS for all involved phases (Table 2).
The annealing initiated topological changes of the phase interface morphology,
leading to a reduction of interface roughness. This would be considered to
affect the Heb as well [24, 31], but its impact would be minor compared to the
growth of phase size, especially in a polycrystalline sample [31]. The
enlargement of FM-phase dimensions is qualitatively evident in the SEM
micrographs shown in Figure 5 and is presumed to have been the main reason for
the decline of Heb by facilitating magnetic nucleation within the FM-phase,
after annealing had been applied.

Figure 6: The magnetic hysteresis loops of the Fe10Ni40NiO50 sample were
measured at 8 $K$. The graph displays a zoom-in at $\pm$ 2000 $Oe$. Inlet in
the left corner provides the whole hysteresis loops. (a) Hysteresis loop of
as-deformed samples (b) Hysteresis loops of the annealed samples. (c) Heb for
different amounts of applied strain. The used fit is not based on a physical
model and serves only to approximate the expected saturation of Heb for an
infinite amount of applied strain.

The detected change in Heb versus applied strain for the Fe10Ni40NiO50
nanocomposite displayed a saturation behaviour similar to the microstructural
saturation in single-phase systems previously reported for several phases
deformed by HPT [34, 36]. A simple fit model based on an asymptotic
approximation yielded, for an asymptotic limit, an Heb,max of 300 $Oe$ (Figure
6c) for the Fe10Ni40NiO50 composite. The fit model is not based on a physical
theory and was only used here to approximate an assumed saturation of Heb
through a saturation of microstructural refinement.

Table 3: Summary of magnetic hysteresis loops of Fe10Ni40NiO50 samples measured at 8 $K$, which are shown in Figure 6a and b. Exchange bias (Heb), symmetric coercivity (HC), coercivity left side (HC1), coercivity right side (HC2), saturation magnetisation (MS) and magnetic remanence (MR).  |  Heb / [$Oe$] |  HC / [$Oe$] |  HC1 / [$Oe$] |  HC2 / [$Oe$] |  MS / [$emu/g$] |  MR / [$emu/g$]
---|---|---|---|---|---|---
$\epsilon$vM$\sim$90 |  -52 |  460 |  -512 |  408 |  37.3 |  18.6
$\epsilon$vM$\sim$270 |  -90 |  625 |  -715 |  535 |  32.6 |  21.2
$\epsilon$vM$\sim$610 |  -124 |  587 |  -710 |  463 |  34.1 |  23.3
$\epsilon$vM$\sim$1280 |  -186 |  786 |  -971 |  600 |  28.8 |  23.7
$\epsilon$vM$\sim$2560 |  -243 |  843 |  -1085 |  599 |  27.6 |  23.7
$\epsilon$vM$\sim$1150 |  -200 |  697 |  -894 |  494 |  32.3 |  26.8
ann. 450°C |  -67 |  515 |  -581 |  449 |  35.8 |  27.0
ann. 550°C |  -23 |  373 |  -396 |  349 |  36.6 |  19.6

##  4 Discussion

###  4.1 Deformation behaviour

The deformation with HPT of composites containing ’-ductile-brittle’ phases
has been reported for several phase systems including, Cu-Co and Cu-W [3, 38].
For those systems, a fracturing of the ’-brittle’ phase has been proposed to
have a crucial influence on the refinement and homogenisation of the sample’s
microstructure. Considering the Fe50NiO50 sample at RT, such a fragmentation
was not observed to an extensive degree. However, that does not imply that NiO
cannot deform at RT. As shown in Figure 1a at RT the Fe and NiO deformed to
lamella structures, but then reached a ’steady state’ without fragmentation of
the NiO phase in reasonably large amounts. This effect is clear when the SEM
micrographs at $r$ $\sim$0 $mm$ and $r$ $\sim$3 $mm$ are compared (Figure 1a).
In contrast, Vickers microhardness increased with radius for the sample
deformed at RT, indicating an ongoing grain refinement within the phases
(Figure 1c), which appeared to occur predominantly in the more ductile
$\alpha$-Fe phase.
With a body-centred cubic crystal structure, $\alpha$-Fe has {110}<-111> as
the primary slip system at RT. The {110}<-111> slip system family fulfils the
von Mises criterion of five independent slip systems for the absorption of
general strain via slip [16]. The mechanical properties of $\alpha$-Fe are
therefore different compared to ceramics such as NiO, which has an NaCl
crystal structure with the primary slip systems {110}<1-10> at RT [16]. Basic
consideration reveals, that such NaCl structures have two independent slip
systems and that the von Mises criterion is not fulfilled; the absorption of
plasticity by NiO is consequently limited. Hence, NiO cannot absorb general
strain through slip, and the likelihood of the fragmentation of
polycrystalline NiO is increased by the lack of matching slip systems of
adjacent grains [16].
These considerations are important for better understanding of the deformation
process in such complex systems. Whether the initial phase is $\alpha$-Fe or
Ni, both phases are more ductile at RT than NiO. It could be that, especially
at RT, the applied strain is absorbed by the $\alpha$-Fe to a large extent, as
the activation of slip is assumed to be easier in $\alpha$-Fe than in the NiO-
phase. It is plausible that through the elevated deformation temperatures, the
activation of the additional slip systems is eased in NiO. The increased
likelihood of slip in NiO could enhance fragmentation and consequently phase
mixing. The eased activation of slip systems in NiO could be one reason for
the drastic change of deformation behaviour at Tdef $\geq$225°C, which was
seen for the Fe50NiO50 samples.
The discrete change of deformation behaviour between 200°C and 225°C (Figure
1e), could not be related to a specific transition temperature. Although NiO
has a TN of $\sim$251°C [45], the expected change of plasticity around the TN
was not reported in [18]. The brittle-to-ductile transition temperature (BDTT)
of NiO has been reported to be $\sim$0.3 of the melting temperature or
approximately $\sim$600°C [14], which is still significantly above 225°C. In
the same study, a brittle-to-ductile transition region was noted, ranging from
300°C to $\sim$550°C. Above 550°C, NiO is considered to be ductile due to the
activation of additional slip systems [14].
A further important point is the slip transfer across phase boundaries, which
was observed in [17, 41]. The likelihood for a glide or dislocation transfer
between phases is higher, if the angle mismatch of adjacent slip systems is at
a minimum [41], in addition to other conditions mentioned in [41]. HPT
processing at elevated temperatures could lower the activation energy of
additional slip systems within NiO. Therefore, the likelihood of dislocation
transfer between adjacent phases is increased, because more slip systems are
available inside NiO, if processed at elevated temperatures. This assumption
could also explain the abrupt change in deformation behaviour observed in the
Fe50NiO50 samples, when Tdef was elevated from Tdef = 200°C to Tdef = 225°C.
Furthermore, the possibility cannot be excluded, that the application of high
hydrostatic pressure reduced the activation energy of additional slip systems
in NiO and caused a shift to lower BDTT. Activation of additional slip systems
through hydrostatic pressure has been discussed in the literature [49, 22] and
has been demonstrated for a Mg-alloy, in which non-basal glide systems were
activated with an applied pressure of $\sim$125 $MPa$ [22].
Eased slip in NiO can lead to a roughening of interface morphology on a
nanometre scale; this effect has been observed in other two-phase systems
[17]. The absorption of dislocations from local pile-ups, could cause local
plastic instabilities, which are believed to support fragmentation of the more
brittle phase, as discussed in a recent study [23].
When the Fe10Ni40NiO50 composition was deformed at Tdef = 300°C, the NiO-phase
was conserved. In addition to the previously discussed impact of slip system
activation and dislocation transfer in NiO, a work-hardening of the softer
phase could also result in an enhanced fragmentation of NiO. Such a phenomenon
has been reported for the Cu-Co-system during HPT-deformation [3]. The
conservation of NiO facilitates a prolonged fragmentation, which could lead to
the fine distribution of NiO imbedded in Ni (Figure 3a).
NiO can break off from larger agglomerates, which would ease the reduction by
Fe through the creation of additional phase interfaces. It is likely that the
accumulation of $\gamma$-Fe2O3 occurs at phase interfaces or the grain
boundary (GB). These small and dispersed clusters of $\gamma$-Fe2O3 could
inhibit GB-motion and accelerate grain refinement. In addition, through the
efficient fragmentation of NiO, a highly diverse Ni and NiO-phase structure is
created on a microscale, significantly hindering GB-motion and plastic
deformation. Both phenomena could have led to the small observed CSDS of
approximately 20 $nm$ according to the WH method at
$\epsilon\textsubscript{vM}$ >1280 (Table 2) and could have caused the unusual
high Vickers microhardness of 9 to 9.9 $GPa$.
The mismatch of slip systems and lattice parameters of participating phases,
especially between $\gamma$-Fe2O3 and the Ni- and NiO phases, could have led
to a dislocation pile-up at the $\gamma$-Fe2O3 phase, resulting in the
creation of a non-crystalline region at the phase interface adjacent to
$\gamma$-Fe2O3. Previous research has cited the build-up of a non-crystalline
region at the phase interface of the Cu-Nb system as creating an impenetrable
barrier for dislocation [35].
Comparing the observed Vickers microhardness values in Figure 3b to the
results obtained for nanocrystalline Ni [42] or slightly oxidized Ni-powder
processed by HPT [1] shows the influence of a nanostructured NiO phase on
mechanical properties. Despite the reported high hardness values in of
approximately 7 $GPa$ for nanocrystalline Ni or Ni with NiO clusters [42, 1]
the bulk nanocomposite of Fe10Ni40NiO50 surpassed these results by 30-40%.

###  4.2 Phase formation during deformation

The efficient reduction of NiO and the observed formation of Fe3O4 and
$\gamma$-Fe2O3 for both investigated compositions was unanticipated.
Commercially available Ni powder is reported to oxidize at $>$350°C in ambient
atmosphere [7]. In general, NiO is considered to be a stable oxide, and
therefore a reduction of Fe was not expected. This reduction can be explained
by comparing the heat of formation, which is significantly less for NiO than
for Fe3O4 and $\gamma$-Fe2O3 [11, 15]. The difference in heat of formation
could also explain the one-way characteristic of the reaction.
According to the data, the onset of NiO reduction by Fe during HPT processing
was initiated between 200°C and 225°C (Figure 2). This reaction occurred much
more efficiently at Tdef = 300°C in the Fe50NiO50 material system. The
formation to the thermodynamically more favourable Fe3O4 appears to have been
enhanced by the extreme conditions during sample synthesis combined with the
fragmented and nanocrystalline NiO, which provided a very large amount of
interfacial phase area for the observed reduction of NiO. A slip in the
{110}<1-10> system in the NiO phase could facilitate the oxidation of Fe,
because the occurrence of unsaturated oxygen bonds would be more likely. These
two mentioned factors are considered to be the main cause for the observed
two-phase system of Fe3O4 and $\gamma$-Fe39Ni61 for the Fe50NiO50 material
system, when deformed at Tdef = 300°C.
NiO could not be resolved in the XRD scan for the Fe50NiO50 samples deformed
at Tdef = 300°C (Figure 2), although 16-wt% NiO should have been left after
synthesis inside the nanocomposite. The possibility that some residual NiO
remained after deformation is supported by the analysis of the $\gamma$-FeNi
phase XRD-pattern. If complete reduction of NiO through Fe and complete
incorporation of the residual Ni-atoms in the $\gamma$-FeNi phase is assumed,
the composition of the $\gamma$-FeNi phase would change to $\gamma$-Fe20Ni80.
The peak pattern of $\gamma$-Fe20Ni80 had its main peak at qz = 3.073 Å-1 and
was therefore distinguishable from the peak position of the observed
$\gamma$-Fe39Ni61 phase.
The XRD results of Fe50NiO50 and Fe10Ni40NiO50 samples in Figures 2 and 4,
respectively, detected different formations of the FexOy phase through the
reduction of NiO. The formation of FexOy is influenced by the amount of
available O [50]. In the Fe10Ni40NiO50 sample, the Fe content was much lower
compared to Fe50NiO50, and therefore sufficient O was available to form
$\gamma$-Fe2O3.
Moreover, $\gamma$-Fe2O3 is thermodynamically more stable than Fe3O4 [11] and
therefore is considered more likely to be formed, provided that sufficient O
is available during deformation. Another difference from the Fe50NiO50 samples
at Tdef = 300°C is the neglectable substitution of Fe inside the Ni. After the
deformation of Fe10Ni40NiO50 composition, WAXS detected a $\gamma$-phase
consisting solely of Ni.

###  4.3 Magnetic properties

Regarding the magnetometry results, the aforementioned Fe-oxide phases (i.e.
Fe3O4 and $\gamma$-Fe2O3) were both FM, but distinguished themselves in their
magnetic properties, for example bulk-MS,Fe3O4 = 98.6 $emu/g$ [8] and bulk-
MS,$\gamma$-Fe2O3 = 85.6 $emu/g$ [37] at 8 $K$. Based on the results from WAXS
measurements, the MS,$\gamma$-Fe2O3 is used in the following discussion of the
results for Fe10Ni40NiO50 nanocomposite.

Figure 7: Relative change of Heb related to Heb,$\epsilon\textsubscript{vM}$
$\sim$270 and comparison of the ratio to the relative changes of MS and CSDS
size of t$\gamma$-Fe2O3 and tNi.

The Heb can be approximated for thin films with the simple theoretical model
of $H\textsubscript{eb}\sim\sigma/(M\textsubscript{S}\cdot
t\textsubscript{FM})$($\sigma$ = interfacial coupling energy density, tFM =
FM-thickness) [4]. Utilising this model, with the assumption that it is valid
to some extent for nanocomposites measured at $8K$ as well, allows estimation
of the origin of the observed change of Heb with applied strain.
The increase of Heb,$\epsilon\textsubscript{vM}$$\sim$90 = -52 $Oe$ to
Heb,$\epsilon\textsubscript{vM}$$\sim$2560 = -242 $Oe$ cannot be explained by
the decline of MS alone. As shown in Table 3, MS decreased by approximately
26% from MS,$\epsilon\textsubscript{vM}$$\sim$90 = 37.3 $emu/g$ to
MS,$\epsilon\textsubscript{vM}$$\sim$2560 = 27.6 $emu/g$, whereas Heb
increased almost fivefold from Heb,min to Heb,max. Considering the observed
changes of CSDS (Table 2) in addition to the decrease of MS, yields a more
accurate estimation (Figure 7). The increase of
Heb,$\epsilon\textsubscript{vM}$$\sim$270 = -90 $Oe$ to
Heb,$\epsilon\textsubscript{vM}$$\sim$2560 = -242 $Oe$ can be explained to a
great extent by the combined decreases of MS and CSDS of $\gamma$-Fe2O3 using
the approximation that $\sigma$ was the same for both applied strains. This
explanation implies, that NiO was exchange biased to the $\gamma$-Fe2O3 as
well as to the Ni-crystallites. From this calculation it can be deduced that
$\gamma$-Fe2O3 crystallites dominated the magnetic characteristic of the
Fe10Ni40NiO50 samples and that the Ni-crystallites contributed partially.
The annealing experiments allow a deeper insight into the dependency of Heb on
the nanocomposite microstructure. The variable MS recovered after deformation,
from MS,$\epsilon\textsubscript{vM}$ $\sim$1150 = 32.3 $emu/g$ to MS,ann.
550°C = 36.9 $emu/g$, and Heb decreased to $\sim$10% of its initial value of
Heb,as def = -200 $Oe$ through annealing at 550°C. The annealed samples
exhibited continuous CSDS growth for all three phases (Table 2) as well as
phase size growth (Figure 5). Changes in CSDS in $\gamma$-Fe2O3 were
approximately twofold as well, but the decrease of Heb was more significant
and could be described only partially by the previous estimation.
Although CSDS apparently had a crucial influence on Heb, magnetic properties
in annealed samples can be influenced by other effects as well. The
application of temperature can initialize optimisation of phase interface area
to lower surface energies and cause a reduction of phase interface topology.
Those effects should also transform the morphology of the FM-AFM interface to
a sharper phase gradient. The decreased interfacial area could affect the
enclosure of FM-grains by the AFM phase and reduce the overall magnetic
stiffness inside the FM-phase. Phase size growth could enable magnetic
nucleation during field reversal, which could further weaken Heb.
The variation of MS with applied strain or annealing temperature appears to
have been related to some degree to the rise of CSDS of $\gamma$-Fe2O3, due to
its large relative changes in size. Previous studies have found that FM-
nanocrystallites possess a size-dependent MS in general. The decline of MS is
related to a non-magnetic region at the crystallite boundary, which begins to
influence MS non-linearly below approximately 20 $nm$ of crystallite dimension
through a growing surface-to-volume ratio [5].

Figure 8: The influence of crystallite size on MS, according to [5], assuming
a non-magnetic layer between adjacent crystallites of 5.7 Å for $\gamma$-Fe2O3
and 3 Å for Ni. Due to the smaller CSDS for $\gamma$-Fe2O3 the surface-to-
volume ratio becomes larger, causing a stronger decrease of MS$\gamma$-Fe2O3.

To obtain a better insight, the simple approximation from [5] was used for MS,
based on the relationship
$M\textsubscript{S}(D\textsubscript{Cry})=M\textsubscript{S}(bulk)\cdot(1-3\cdot
t\textsubscript{non-mag}/D\textsubscript{Cry})$(DCry…crystallite size; for the
calculation, the CSDSs from Table 2 and tnon-mag…thickness of a non-magnetic
layer between adjacent nanocrystallites were used) for nanocrystalline
material. Utilising this model for the Fe10Ni40NiO50 nanocomposite, which
consisted of 12-wt% $\gamma$-Fe2O3 and 48.6-wt% Ni; yields with a non-magnetic
layer of 5.7 Å for $\gamma$-Fe2O3 [5], 3 Å for Ni (estimated for Ni) and
assuming a bulk MS,Ni = 58.6 $emu/g$ [12] and MS,$\gamma$-Fe2O3 = 85.6
$emu/g$[37]: Mcal.S,$\epsilon\textsubscript{vM}$$\sim$270 = 35.7 $emu/g$,
Mcal.S,$\epsilon\textsubscript{vM}$$\sim$1280 = 35.1 $emu/g$ and
Mcal.S,$\epsilon\textsubscript{vM}$$\sim$2560 = 34.2 $emu/g$, far more than
measured (Table 3). These results suggest that the assumed thickness of the
non-magnetic layer was underestimated for the considered nanocomposite and
should have been much higher to match the experimental results. Regarding the
different contributions to MS of $\gamma$-Fe2O3 and Ni, crystallites that
possess different CSDS allow the conclusion, that the decrease of MS was
mainly caused by the shrinking CSDS of $\gamma$-Fe2O3 crystallites through the
deformation (Figure 8).
The results from the annealed samples support the previous argument, that the
$\gamma$-Fe2O3 phase mainly influenced MS. Growth of $\gamma$-Fe2O3 from 8
$nm$ to 14 $nm$ CSDS would lead to the following calculated MS:
Mcal.S,$\epsilon\textsubscript{vM}$ $\sim$1150 = 34.4 $emu/g$,
Mcal.S,ann.450°C = 35.0 $emu/g$ and Mcal.S,ann.550°C = 36.0 $emu/g$. The
estimation for MS,$\epsilon\textsubscript{vM}$ $\sim$1150 is certainly higher
than the measured MS value; however, MS,ann.450°C and MS,ann.550°C are in good
accordance with the values presented in Table 3.
The discrepancy between MS,$\epsilon\textsubscript{vM}$ $\sim$1150 and
MS,$\epsilon\textsubscript{vM}$$\sim$2560 could be rooted in a larger number
of non-magnetic regions than assumed. If complete oxidation of Fe to
$\gamma$-Fe2O3 is assumed and crystal size effects are neglected, the
calculated MS for the annealed Fe10Ni40NiO50 sample yields, a value of MS,cal
= 38.3 $emu/g$. The MS,ann. 550°C = 36.9 $emu/g$ is strikingly similar to the
theoretical calculation. This comparison suggests that the non-magnetic phase
was predominantly composed of Ni and $\gamma$-Fe2O3 and was reduced to a
minimum by the annealing treatment.
There have been reports of a phase transition from FM $\gamma$-Fe2O3 to the
thermodynamically more favourable AFM $\alpha$-Fe2O3 at temperatures of
370-600°C [11]. While this transition could have occurred during annealing, an
accumulation of non-crystalline regions through the application of strain is
more likely. On one hand, the XRD peak pattern of $\alpha$-Fe2O3 is distinct
from $\gamma$-Fe2O3; on other hand, a formation of $\alpha$-Fe2O3 cannot
explain the rise in MS through annealing. The plausible existence of non-
crystalline regions at interfaces, presumably containing $\gamma$-Fe2O3, would
have a fraction of bulk MS, MS,amorphous $\sim$2-5% [25], and the transition
from crystalline $\gamma$-Fe2O3 phase to a phase with non-crystalline regions
at the crystallite boundary could explain the changes in MS, detected when
large amounts of strain had been applied. Composites, such as Nb-Cu, deformed
by SPD have been reported to contain non-crystalline regions at phase
interfaces [35, 40]. Crystallisation of these non-crystalline regions was
observed when annealing was applied [35]. These results suggest that similar
behaviour occurred in the Fe10Ni40NiO50 nanocomposite.

##  5 Conclusions

This study presents an insight into the deformation behaviour of phases
possessing very distinct mechanical properties. Although NiO has, due to its
NaCl-structure, different primary slip systems compared to Fe or FeNi, it was
possible to synthesise nanocomposites with a homogeneous microstructure on the
sub-micrometre regime. The importance of thermal HPT processing was
demonstrated on the Fe50NiO50 composition. Through the deformation at 300°C, a
homogeneous two-phase microstructure evolved. XRD investigations confirmed a
reduction of NiO through Fe leading to the formation of Fe3O4 and
$\gamma$-Fe39Ni61.
To conserve the AFM NiO in respect to enhance the Heb, a composition of
Fe10Ni40NiO50 was synthesised by HPT at 300°C and nanocomposites with a unique
microstructure were obtained. SEM micrographs showed a very homogeneous
microstructure, which primarily contained of NiO and Ni according to WAXS
investigations. The formation of an additional phase was detected by WAXS,
which consisted of $\gamma$-Fe2O3. Vickers microhardness measurements detected
a microstructural saturation behaviour with a slight increase to outer radii
to an unusual high Vickers microhardness of $\sim$9.9 $GPa$. WAXS, SEM and
magnetometry results conveyed an ongoing refinement of the nanocomposite’s
microstructure.
It was possible to demonstrate the existence of Heb in such bulk sized
nanocomposite and the continual increase of Heb with applied strain was
correlated to the evolving microstructure. The rise of Heb was mainly
attributed to a simultaneous decrease in MS, reduction of FM-phase size
dimensions and crystallite size of the FM-phases, predominantly
$\gamma$-Fe2O3.
Annealing experiments of Fe10Ni40NiO50 samples showed a decrease of Heb and an
increase of MS. Those changes of Heb and MS were contributed to the growth of
FM-phase dimensions and the subsequent reduction of phase interfaces.

##  6 Appendix A. Schematic Illustration of the HPT - Synthesis

Figure 9: Schematic description of the sample synthesis. From left to right:
Powder blends were compacted with HPT inside an airtight capsule. The obtained
pellet was used for HPT-deformation experiments. An inductive heating system
provided the needed processes temperature for the HPT-anvils. HPT-disc was cut
into the desired size for further investigations. The SQUID samples were cut
out at the desired radius. For the annealing experiments a HPT-disc was cut
into quarters. One quarter provided the SQUID-samples, which were cut out at
the same radius. The SQUID sample was simultaneously annealed with a quarter
of the HPT-disc.

Funding:
This project has received funding from the European Research Council (ERC)
under the European Union’s Horizon 2020 research and innovation programme
(grant agreement No:757333).

Conflicts of Interest:
The authors declare no conflict of interest.

Acknowledgement:
We acknowledge DESY (Hamburg, Germany), a member of the Helmholtz Association
HGF, for the provision of experimental facilities. Parts of this research were
carried out at (PETRA III) and we would like to thank (N. Schell and E.
Maawad) for assistance in using (P07B- High Energy Materials Science).

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