A novel aluminum based nanocomposite with high strength and good ductility

1. Accepted Manuscript
A novel aluminum based nanocomposite with high strength and good ductility
H. Ramezanalizadeh, M. Emamy, M. Shokouhimehr
PII: S0925-8388(15)30485-0
DOI: 10.1016/j.jallcom.2015.07.088
Reference: JALCOM 34779
To appear in: Journal of Alloys and Compounds
Received Date: 21 April 2015
Revised Date: 17 June 2015
Accepted Date: 10 July 2015
Please cite this article as: H. Ramezanalizadeh, M. Emamy, M. Shokouhimehr, A novel aluminum based
nanocomposite with high strength and good ductility, Journal of Alloys and Compounds (2015), doi:
10.1016/j.jallcom.2015.07.088.
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A novel aluminum based nanocomposite with high
strength and good ductility
H. Ramezanalizadeh*a
, M. Emamya
, M. Shokouhimehrb
a
School of Metallurgy and Materials Engineering, College of engineering,
University of Tehran, Tehran, Iran.
b
School of Chemical and Biological Engineering, College of Engineering, Seoul
National University, Seoul, Korea.
Email: hralizadeh@ut.ac.ir
Abstract:
Aluminum based nanocomposite containing nano-sized Al3Mg2 reinforcing was fabricated via
mechanical milling followed by hot extrusion techniques. For this, Al and Al3Mg2 powders were
mixed mechanically and milled at different times (0, 2, 5, 7, 10, 15 and 20 h) to achieve Al–10
wt. % Al3Mg2 composite powders. Hot extrusion of cold pressed powders was done at 400 º
C
with extrusion ratio of 6:1. Microstructures of the powders and consolidated materials were
studied using transmission electron microscopy, scanning electron microscope and X-ray
diffraction. Fracture surfaces were also investigated by scanning electron microscopy equipped
with EDS analyzer. The results showed that an increase in milling time caused to reduce the
grain size unlike the lattice strain of Al matrix. In addition, the fabricated composites exhibited
homogeneous distribution and less agglomerations of the n-Al3Mg2 with increasing milling time.
The mechanical behavior of these nanocomposites was investigated by hardness and tensile tests,
which revealed it has four times the strength of a conventional Al along with good ductility. It
was found that the ultimate tensile strength (UTS) and elongation of the nanocomposites were
significantly improved with increases in milling time up to 15 h. This improvement was
attributed to the grain refinement strengthening and homogeneous distribution of the n-Al3Mg2.
Fracture surfaces showed that the interfacial bonding between Al and Al3Mg2 could be improved
with increasing in milling time. Also HRTEM results from interface showed that a metallurgical
clean interface and intimate contact between matrix and second phase. By extending the milling
process up to 20 h, there was no significant improvement in mechanical behavior of materials,
due to the completion of milling process and dynamic and static recovery of composite at higher
milling times.
Keywords: Composite materials, nanostructures, powder metallurgy, tension test,
ductility, fracture.

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1. Introduction
Particulate-reinforced metal matrix composites (PRMMCs) have the potential to
provide improved mechanical properties, for example, high specific stiffness and
specific strength and creep resistance, which makes them attractive for applications
in the aerospace, defense and automotive industries [1]. Among the various
MMCs, Al-based composites are of interest because of their low density and good
formability [2]. Metal matrix nanocomposites (MMNCs) which are usually
mixtures of nanoparticles in a metal demonstrate new interesting properties [1].
The particles are supposed to act as obstacles to dislocation motion and to
effectively pin grain boundaries, conferring therefore microstructural stability to
the composite.
Nowadays, there are two methods to fabricate MMNCs. In one route, the melt
technology needs that nanoparticles are dispersed in the melt of the metal matrix.
This method is similar to the production of polymer nanocomposites and faces
similar problems. The most important one is the poor wettability of the particles by
the melt [1,3]. This limits the amount of reinforcing that can be properly dispersed
[4]. In addition, due to the high temperature of the melt, chemical reactions take
place at the interface between the second phase and the matrix, which in some
cases leads to the formation of a brittle interphase which decreases the material
performance. In second route, the most promising production technology is based
on powder metallurgy (PM) [5]. The composites fabricated by normal powder
metallurgy routes usually have weak interfaces between the second phase and the
matrix. These interfaces tend to debond when subjected to an applied load,
resulting in deterioration in mechanical behavior [6].
One type of the powder metallurgy techniques is mechanical alloying or milling
(MA/MM). The MA/MM process, using ball-milling techniques, has received
much attention as a powerful tool for the fabrication of several advanced materials
including equilibrium, nonequilibrium (e.g. amorphous, quasicrystals,
nanocrystalline), and composite materials [7,8,9]. Also, MA/MM has produced
notable improvements in the strength, toughness, fatigue life, and corrosion
resistance of aluminum alloys [10].

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This method requires mixing of two materials in powder form, mechanical alloying
by ball milling followed by pressing, sintering, and/or hot pressing or extrusion.
This method allows avoiding issues with miscibility and generally leads to sharp,
mechanically strong interfaces between nanoparticles and the matrix.
Uniaxial cold-pressing is commonly used to densify the powder particles, of varied
geometries, at low cost and with high productivity, providing cold-pressed
powders that can be consolidated by sintering or hot extrusion. The latter has the
advantage of promoting full densification [11]. Hot extrusion lets a high shear
strain rate, providing a high-strength bonding between particles, and a
microstructure very similar to that of a wrought product. In the case of aluminum
and its alloys, hot extrusion breaks the typical oxide layer that coats the powder,
providing better bonding of the particles [12]. In MMCs, hot extrusion tends to
eliminate the clustering of reinforcing particles and therefore a better distribution
through the metal matrix [13]. Various aluminum matrix composite systems with
reinforcing phases of B4C [14], SiC [1,4,5], Al2O3 [4], BN [15] and AlN [16], have
been developed via various techniques. Other possible candidates as reinforcing
agents in MMCs are complex metallic alloys (CMAs), intermetallic compounds
with giant unit cells, comprising up to more than a thousand atoms per unit cell
[17]. In fact, CMAs show many attractive properties for reinforcing applications,
such as high strength to weight ratio, good oxidation resistance and high-
temperature strength [17,18]. Among the various CMAs, the β-Al3Mg2 phase (with
1168 atoms per unit cell) [17,19] has been studied with special attention to its
structure as well as to its physical and mechanical properties [17]. In addition, it is
worthy to note that the Al3Mg2 has lower specific gravity of 2.25 g/cm3
[17] (less
than that of Al, 2.7 g/cm3
), compared to 3.95 g/cm3
for Al2O3, 3.21 g/cm3
for SiC,
2.51 g/cm3
for B4C and 3.51 g/cm3
for diamond [20], which normally have been
used as second phase in MMCs. These unique properties, along with other
attractive properties such as high hardness and wear resistance [17], high-
temperature strength (~300 MPa at 573 K [17]), and high capacity for hydrogen
absorption [21], make Al3Mg2 a likely candidate as the reinforcing in an Al matrix
composite.
Up to now, there is little work on the effect of novel reinforcing addition on the
structural and mechanical behavior of MMNCs. Scudino et al. propound an idea

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for applying CMAs as new reinforcements instead of conventional reinforcing
materials [19].
In this work, we study the effect of Al3Mg2 nanoparticles and of processing
conditions like milling time on the microstructural and mechanical behavior of the
composite. The powder is consolidated using cold press and then vacuum sintering
followed by hot extrusion. The mechanical behavior of the composites is probed by
performing uniaxial tensile tests and microhardness measurements. Moreover,
comparative study of hardness, UTS and elongation values of the nanocomposite
material will be done relative to the pure Al.
2. Experimental
Nominal composition of the selected Al3Mg2 was Al60Mg40 according to the phase
Al-Mg diagram [22], shown in Fig. 1. According to this image, Al3Mg2 single
phase could be formed as a stable phase from a congruent melt with melting point
of 451 º
C at 38.5-40.3 atomic percent range of Mg. In this study pure Al (>99.92%)
and Mg ingots (>99.90%), were used in a well atmosphere controlled furnace to
prepare Al3Mg2 ingots. Then powders of Al3Mg2 were fabricated by mechanical
milling of the broken ingot using an attrition ball mill with rotation speed of 400
rpm and ball to powder weight ratio of 12:1. The vial and balls were made of
hardened chromium steel. In all experiments stearic acid (2 wt. %) was also added
to the powder mixture as a process control agent (PCA). The milling operation was
pursued in pure argon atmosphere (99.999%) to avoid the oxidation of the
materials. Al powder, (63 µm, >99.90%), blended with 10 wt. % of β-Al3Mg2
powders were synthesized through mechanical milling technique. Table 1 shows
the material compositions and milling conditions for this study. The phase
compositions of milled powders were investigated by X-ray diffraction analysis
(XRD) using a ‘‘Philips PW 1730’’ X-ray diffractometer with CuKα filtered
radiation and 2 deg/min scanning rate.
Table 1. Materials composition and milling conditions used in this study.
Materials
label*
Matrix Al3Mg2 (wt %) Milling time (h)
Al Al 0 0
AC10-UM Al 10 0
AC10-2HM Al 10 2
AC10-5HM Al 10 5

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AC10-7HM Al 10 7
AC10-10HM Al 10 10
AC10-15HM Al 10 15
AC10-20HM Al 10 20
*The symbol: “A” represents Al; the second symbol: “C” represents
Composite, the third symbol: “H” represents Hour and the last
symbol: “M” represents Milling. The first two digits designate the
amount of Al3Mg2 and the next digits designate the milling time.
The UM means Unmilled.
XRD data was also used to determine the crystallite size (D) and lattice strain (ε).
The crystallite size was determined from the broadening (B) of the diffraction lines
1 1 1, 2 0 0, 2 2 0 and 3 1 1 using following Scherrer equation [23]:
‫ܦ‬ =
଴.ଽఒ
஻ ௖௢௦ ఏ
(1)
The lattice strain (ε) was also calculated for the same diffraction lines from the
following equation [23]:
ε = B/4tan θ (2)
where, ߣ is wave length = 1.54059 Å (CuKߙ radiation), B is the full width at half
maximum, θ is the angle in radians and d is d-spacing between the planes.
Particle size distribution was determined by laser particle size analyzer (CILAS
1064 Liquid), to determine the volume size distribution, D10, D50 and D90
automatically.
Grain structure of the milled powder was investigated by utilizing transmission
electron microscope (TEM, JEOL JEM-2010 and JEOL JEM-3010 microscopes)
operated at 200 kV. For the sample preparation, the powder particles were
dispersed in ethanol using an ultrasonic bath and subsequently a small drop of the
suspension was placed on a Cu grid coated with carbon film, and finally dried in a
vacuum desiccator.
High resolution transmission electron microscopy (HRTEM) was performed using
JEOL JEM-2100 and JEOL JEM-3000F microscopes operated at 200 kV and 300
kV, respectively. FIB-SEM hybrid system, SMI3050SE, SII NanoTechnology Inc

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was used for TEM sample preparation from extruded materials. Fine probe energy
dispersive X-ray spectroscopy (EDS) was employed to determine the distribution
of dispersoids within the matrix. Further compositional information was obtained
with Z-contrast imaging, also termed high-angle annular dark field (HAADF)
imaging. In this imaging mode, the brighter regions correspond to heavier atoms,
as the scattering cross-section is proportional to Z2
[24].
Field emission scanning electron microscope (FESEM) type Carl Zeiss, Sigma,
equipped with an energy dispersive spectrometer (EDS) was employed to
investigate the morphology and particle size of the milled powders after different
milling times. The microstructures of extruded samples and also fracture surfaces
were studied by FESEM.
Cold pressed specimens were prepared through powder metallurgy technique at a
compaction pressure of 500 MPa and then sintered at 448 º
C for 40 minutes in an
accurate electrical resistance furnace with a heating rate of 10 º
C per minute. After
sintering, the specimens were hot extruded at 400 º
C with extrusion ratio of 6:1
[17].
The relative density and porosity of extruded bodies were measured by
Archimedes method according to ASTM: B962-13. Theoretical density of
compacts was calculated using the simple rule of mixtures, considering the fully
dense values of aluminum and Al3Mg2 are 2.7 g/cm3
and 2.25 g/cm3
, respectively.
For mechanical property measurements, samples were cut into round dog-bone
shaped specimens, according to ASTM: E8/E8 M-11, along the extrusion direction
(Fig. 2). The tensile samples were mechanically polished to mirror-like images. In
order to examine reproducibility, three replicates were used for test. Tensile testing
was conducted in a tensile machine at ambient temperature and a strain rate of 0.1
mm/min. Microhardness was measured using a microhardness tester with a Vickers
indenter under a load of 100 g for 10s. An average of 30 indentations was
considered as the Vickers microhardness value.

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Fig. 1. Phase diagram of Al-Mg [22].
Fig. 2. Schematic drawing of tensile sample dimensions.
3. Results and discussion
3.1. Composition and morphology of composite powders
Fig. 3a shows XRD results obtained from Al3Mg2 ingot. As it can be seen from
Fig. 3a, all peaks are related to β-Al3Mg2. As mentioned before Al3Mg2 powder
was produced by mechanical milling of broken ingot. XRD pattern of milled
Al3Mg2 after 25 h milling is shown in Fig. 3b. Comparison between Fig. 3a and
Fig. 3b reveals that mechanical milling has significant influence on changing the
lattice of Al3Mg2 intermetallic. As expected, mechanical milling introduces strain
in the lattice of Al3Mg2 resulting in increased dislocation density, reducing space
between planes and finally decreasing particle or crystallite size of the intermetallic
[7]. Fig. 4 shows a FESEM image of the Al3Mg2 powders after 25 h milling. This

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image revealed that the high energy of milling imposed on the powder during the
25 h milling leads to the refinement of the micron-powders (broken ingot, not
shown here) down to ultrafine spherical shaped nanoscale powder (<100 nm in size
that were agglomerated in clusters ranging between 0.5 and 3 µm in size with
average substructure size of ∼30 nm) (Fig. 4b).
Fig. 3. XRD patterns of a) as-cast Al3Mg2 and b) milled Al3Mg2 for 25 h.
Fig. 4. FESEM image of milled Al3Mg2 for 25 h (a) and high magnification showing Al3Mg2
nanoparticles (b).
Bright field TEM image of Al3Mg2 particles after 25 h milling is shown in Fig 5.
This image reveals the presence of nanocrystalline grains and the corresponding
selected area diffraction (SAD) pattern (the image inset to Fig. 5) shows diffused

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rings suggesting the presence of nanostructured β-Al3Mg2 phase along with a
possible presence of amorphous phase. As expected, due to the high brittleness and
hardness of Al3Mg2 intermetallic, its size reduces vigorously during milling. Also,
when this compound embeds in Al matrix as reinforcing agent and nanocomposite
is formed, hardness and brittleness of Al matrix could be raised. So, the size of
nanocomposite particles during mechanical milling may decrease by extending
milling time [25].
Fig. 5. Bright field TEM image of milled Al3Mg2 after 25 h and corresponding SAD pattern
(inset image).
FESEM imaging and particle size distribution measurement of the powders before
and after milling revealed that the high energy of milling imposed on the powder
during the 15 h milling time led to the refinement of the micron-powders (63 µm
average particle size and 84 nm average grain size for pure Al (Fig 6. a, b) down to
ultrafine feathery shaped nanoscale powder (<100 nm in size that were
agglomerated in clusters ranging between 5 and 9 µm in size with average
substructure size of ∼70 nm) (Fig. 6c, d, e).

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Fig. 6. (a) FESEM image and (b) Laser particle size distribution graph of starting Al powder, (c)
low magnification and (d) high magnification FESEM image and (e) Laser particle size
distribution graph of milled AC10-15HM sample.
Fig. 7a shows a bright field TEM image of AC10-2HM powder sample, in which
the Al3Mg2 nanoparticles have been dispersed in the nanocrystalline Al matrix.
Also, a bright field TEM micrograph of AC10-15HM powder sample is shown in

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Fig. 7b, indicating the homogeneous dispersion of second phase in the matrix.
Typical analysis of the nanocomposite is also given in Fig. 7c, which corresponds
to the composition of Al/Al3Mg2. Also, the insert in Fig. 7b is the SAD pattern
from the interface indicated by arrow. The SAD rings indicate a random crystal
orientation of the Al and Al3Mg2 phases. The absence of any other phases in the
composite confirms that the interface between matrix and second phase is clean
[26]. Comparison between Fig. 7a and Fig. 7b demonstrates that by increasing the
milling time up to 15 h, not only the size of particles and grains decreases, but also
introduces more structural homogeneity with less particle agglomeration. The
mentioned factors could result in good mechanical properties of composite
materials after consolidation process.
Fig. 7. Bright field TEM micrographs of (a) AC10-2HM and (b) AC10-15HM powder samples,
(c) the typical EDS of nanocomposite (the insert image to Fig. 7b is corresponding SAD pattern
from detected area by arrow).

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3.2. XRD study of the extruded nanocomposite
Fig. 8 shows XRD patterns of extruded samples produced after different milling
times. It seems the present peaks are only related to Al phase. Although, it should
be noted that due to minor amount, small particle size and lower density of Al3Mg2
in comparison with Al, the intensity of Al3Mg2 peaks are weaker than Al ones [19].
The inset picture in Fig 8 shows the XRD peaks at 2θ = 38.4 for bulk samples
milled from 0 to 15 h. By camparision, one can see that the peak intensities
increase for AC10-UM sample, because of adding Al3Mg2 particles to Al matrix.
As mentioned before, Al and Al3Mg2 peaks overlap with each ohther in some
positions and this caused to increase the intensity of peaks. Obvious peak
broadening and intensity decline can be seen with increasing milling time which is
evidence of internal grain size reduction. During the mechanical milling process,
the milling balls impact the powder particles causing them to plastically deform
and subsequently work harden, which is then followed by fracture and re-welding
of the particles.
Fig. 8. XRD patterns of extruded AC10 composites produced after different milling times.

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Crystal size and lattice strain are important parameters, since they have significant
effect on both compacting of the powders during sintering process and properties
of the finely obtained aluminum matrix strengthened by fine particles. The average
crystal size of Al matrix in the composite was estimated using broadening of XRD
peaks. The effect of milling time on crystal size and lattice strain of examined
samples was presented in Fig. 9, indicating that the crystal size decreases with
increasing milling time, according to the Eq. 3 [10]:
D = Kt -2
(3)
Where, K is a constant, D is crystallite size (nm) and t is milling time (h) [10].
In this work, the most intensive crystal refinement occurs in the early stage of
milling, up to 10 h. With prolonged time, the crystal size of the milled composites
decreases slowly. The lattice strain increases while crystal size reduces with
increasing milling time due to distortion effect caused by dislocation in the lattice
[28]. With increasing milling time, severe plastic deformation brings about a
deformed lattice with high density of dislocations [29].
Fig. 9. Crystallite sizes and lattice strains of extruded AC10 samples at different milling times.

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3. 3. Relative density and porosity
The effect of milling time on relative density and porosity of the extruded samples
are shown in Table 2. As expected, the theoretical density of Al/10%wt. Al3Mg2
composites (2.655 g/cm3
) is lower than that of Al density (2.7 g/cm3
). The relative
density of the extruded samples increases with increasing milling time. This
increased density is high probably due to the large surface area of the fine particles,
which induces high reactivity, i.e., particles with small mean size values are easily
densified as bigger particles under the same compaction conditions. Increasing the
density with increasing the milling time has also been reported in several works
[30].
Table 2. Effect of milling time on relative density and apparent porosity of AC10 composite.
Sample Theoretical
density (g/cm3
)
Archimedes
density (g/cm3
)
Relative
density (%)
Apparent
porosity (%)
Al 2.7 2.68 99.25 0.75
Al3Mg2 2.25 - - -
AC10-UM 2.655 2.609 98.3 1.7
AC10-2HM 2.655 2.6119 98.38 1.62
AC10-5HM 2.655 2.6149 98.49 1.51
AC10-7HM 2.655 2.6188 98.64 1.36
AC10-10HM 2.655 2.6239 98.83 1.17
AC10-15HM 2.655 2.6297 99.05 0.95
Hot deformation process such as extrusion was found to be effective in reducing
the number of pores and improving the density of composite materials. In general,
the presence of oxide layers on the surface of the aluminum powders can
substantially degrade the solid-phase-sintering ability as a diffusion barrier (note
that it is solid phase-sintering at 448 º
C in the present study). However external
applied stress can break up the oxide layer, and thus improve the solid-phase-
sintering ability and decrease the volume fraction of the pores.
3.4. Microstructure of the hot extruded samples
Fig. 10a and Fig. 11a show FESEM micrographs taken from the cross-section of
the extruded nanocomposites along with EDS analysis (Fig. 10b and Fig. 11b) and

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map analysis of Al (Fig. 10c and Fig. 11c) and Mg (Fig. 10d and Fig. 11d)
elements. As seen in these figures, the microstructures consist of dark areas (the β-
Al3Mg2 CMA reinforcement) dispersed in the Al matrix (the gray regions). As
shown in Fig. 10a, some clustered and agglomerated of Al3Mg2 are observed in
AC10-UM sample. It can be seen that with increasing milling time, the distance
between β-Al3Mg2 particles in the final product reduces and more homogenous
structure is obtained (Fig. 11a). In such case, only few pores are visible, further
corroborating the high density of the consolidated specimens.
Fig. 10. (a) FESEM micrographs for the extruded AC10-UM material, (b) typical EDS analysis,
(c), map analysis of Al and (d) map analysis of Mg.

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It is well known that the first requirement for superior performance of a composite
material is the homogeneous distribution of the reinforcing phase. In particulate-
reinforced composites; any agglomeration of the reinforcement particles
deteriorates the mechanical properties. Differences in particle size, densities,
geometries, flow or the development of an electrical charge during mixing all
contribute to particle agglomeration. Reduction in reinforcement particle size
brings about an increase in the mechanical strength of the composite, but the
tendency of particle clustering also increases (Fig. 10a) [25]. In powder
metallurgy, the matrix and reinforcement mixing process is the critical step
towards a homogeneous distribution of reinforcement particles throughout the
matrix. One of the methods that can be used to achieve homogeneity of particle
distribution throughout the matrix and also reduce their size during the process is
MM (Fig. 11a) [10]. Also, in MMCs, hot extrusion tends to eliminate the clustering
of reinforcement particles and therefore provide a better distribution through the
metal matrix [31]. So, as both MM and hot extrusion processes were used in the
present study, the final distribution of Al3Mg2 nanoparticle is expected to be
uniform (Fig. 11c, d).

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Fig 11. (a) FESEM micrographs for the extruded AC10-15HM material, (b) typical EDS
analysis, (c), map analysis of Al and (d) map analysis of Mg.
3.5 Secondary phase sizes and their distribution
Different from diffraction contrast, the HAADF mode (STEM) provides incoherent
images and uses high-angle scattering and therefore leads to strong atomic number
(Z) contrast [24]. The intensity of the atom columns is approximately proportional
to the mean square of the atomic number; and as such, the contrast in such a mode
is strongly dependent on chemical composition [32].
Fig. 12a presents a HAADF image of AC10-15HM sample. The grains and grain
boundaries can be seen with dark dispersoids (with a diameter/length of about 10
nm). The dark color, i.e., intense scattering from the dispersoids, suggests that they
consist of low atomic number elements. EDX point scan results, a representative

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spectrum of which is shown in Fig. 12b, indicate that the composite includes Al
and Mg elements. The other elements like Cu, C, Ga and O is originated from
sample preparation for TEM. To verify this further, EDX element mapping
analysis was also conducted as shown in Fig. 12c–d. From these figures it could be
concluded that the homogenous distribution of Al3Mg2 particles thorough Al
matrix could be improved after 15 h milling. On the other hand, the size of Al3Mg2
particles seems to be around 10 nm which confirms the severe reduction in particle
size of Al3Mg2 after 15 h milling. These factors; i.e. homogenous distribution and
decreasing the particle size could result in good mechanical properties of material.
Fig 12. (a) STEM image of AC10-15HM sample revealing the distribution of the Al3Mg2
particles in the Al matrix, (b) the typical EDX result of (a) showing the Al and Mg elements, and
EDX elements mapping of (c) Al and (d) Mg.
3.6 Interface structure and dislocation density
The quality of the interface between a matrix and reinforcement determines the
performance of the consolidated composite material [33]. A strong interface allows
effective load transfer from the matrix to the reinforcement, leading to the
improved strength and stiffness.
To provide insight into the role of interfaces on the mechanical behavior, a region
for detailed TEM examination was selected. Fig. 13a shows a lower magnification
bright-field TEM image from the Al3Mg2/Al composite milled for 15 h. Fig 13b is
HRTEM image from the region that is enclosed by a rectangle in Fig. 13a. From

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this figure, it is evident that the Al lattices are directly linked with the lattices of
the Al3Mg2. On the other word, an Al3Mg2 particle is in intimate contact with the
Al lattice and no intermediate layers or defects are evident at this interface (Fig.
13b). The Moiré fringe structure that is evident in the Al region is thought to arise
from the misorientation from two overlapped Al grains. This type of interface,
which is most commonly observed, may be formed in cases where there is no large
lattice mismatch between Al3Mg2 and Al.
Fig 13. (a) Bright field TEM image of AC10-15HM sample showing an overall view of
nanograins with planar features within the grains and (b) HRTEM image from the area enclosed
by a rectangle in (a) showing the Al3Mg2 lattice linked with Al matrix directly and also
Al3Mg2/Al interface is metallurgically clean.
The accumulation of a high-density of dislocations in the matrix, second phase and
also around the second phase is believed to play an important role in determining
the strength of the matrix and second phase and therefore the overall strength;
hence, the dislocation density in the composite was measured by HRTEM methods
[34], as shown in Fig. 14(a). Fig. 14(b) is the corresponding inverse Fast Fourier
Transformation (FFT) pattern of Fig. 14(a), in which the edge dislocation
structures are marked by “T”. The dislocation density present in these regions was
calculated to be 7.54×1014
/m2
on the basis of statistical dislocation numbers
obtained from tens of HRTEM images. It should be noted that dislocation density
can also be measured by several techniques, including X-ray peak broadening
method [35] and electrical resistivity techniques [36]. However, for the current

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composite material, HRTEM provides a direct measurement of the dislocation
density in the samples.
Fig. 14. (a) HRTEM image taken from the AC10-15HM sample (b) Inversed Fast Fourier
transformation pattern of the image in (a), where dislocation cores are marked by “T”.
3.7 Mechanical properties of hot extruded samples
3.7.1 Microhardness
The variation of composite microhardness versus milling times is presented in Fig.
15. Microhardness values for extruded pure Al and unmilled sample were
measured to be 48 HV and 73 HV, respectively. As shown in Fig. 15, hardness
markedly increases with the addition of second hard particles as a combined effect
of load share between the matrix and reinforcing phase (due to their high module,
reinforcement particles can withstand higher amount of stress before they start
deforming plastically, helping the softer metallic matrix) and their effect on flow
obstruction of dislocation during deformation, acting like barriers.
The profound high hardness of composite powders produced by milling process is
due to work-hardening effect of the milling operation, the effect of reinforcing
phase [13] and nanostructured alloy with higher dissolved alloying element much
higher than the equilibrium limit [37].

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For AC10 samples, microhardness value increases from 73 HV to 164 HV after 10
h milling time and at longer milling times, the increasing trend continues with as
lower rate (200 HV for 20 h). The hardness increment is caused by the increase of
the dislocation density as well as the crystallite refinement, Fig. 15 [38]. Slower
rate in increasing the microhardness value at longer milling time may be attributed
to the completion of milling and dynamic recovery due to high work hardening
effects of deformed Al powders. It may be even ascribed to static recovery of high
deformed Al matrix with local increase of temperature in particles during
collisions.
Fig. 15. Micro hardness values for extruded samples.
3.7.2 Tensile strength of nanocomposite
The mechanical properties of MMCs are directly proportional to the distribution
type, amount and size of its reinforcements. As shown in Fig. 11 and Fig. 12, fine
Al3Mg2 particles are well dispersed and uniformly distributed in Al matrix. The
uniform distribution of Al3Mg2 particles can be attributed to the time and the
method of mixing. For the composite materials, it is very important to obtain
homogeneous reinforcement in the matrix to enhance the mechanical, electrical
and thermal properties [29].
Fig. 16 shows the stress–elongation curves from tensile tests of extruded samples
after applying different milling times. A considerable difference is seen between
the strength of pure aluminum sample and the composite specimens. From Fig. 16,
it is clear that the addition of second phase to Al matrix without milling increases

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the strength of composite in cost of reduction of elongation. Then, the strength and
elongation of the composite increases with increasing milling time and the sample
milled for 15 h exhibits comparatively exceptional properties. Considering these
plots, the sample containing 10 wt. % Al3Mg2 when milled for 15 h has the highest
UTS values (623 ± 12 MPa).
Fig. 16. Typical tensile strain–stress plots of Al/10 wt.% Al3Mg2 extruded nanocomposite.
Considering Fig. 16, increasing milling time will increase the strength
significantly. The reason is that Al3Mg2 nanoparticles prevent the movement of
dislocations in pure aluminum matrix through dispersion strengthening mechanism
[39]. This behavior is described by Eq. 4. According to Eq. 4, as the reinforcement
particle size decreases, the distance between the particles will also decreases [39].
ߣ =
ସ ሺଵି௙ሻ௥
ଷ௙
(4)

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Where ߣ is the distance between the reinforcement particles, f is the particles
volume fraction, and r is the particle radius, assuming them spherical. Lowering
the reinforcing particle size at a constant amount decreases the distance between
the reinforcing particles. In other word, according to Eq. 5 decreasing the distance
between the Al3Mg2 particles will increase the required tension for dislocations
movement between the Al3Mg2 particles leading to an increase in the composite
strength.
߬଴ =
ீ௕
ఒ
(5)
In Eq. 5, ߬0 is the required tension for forcing dislocations to move among
reinforcement particles. G is the material’s elastic modulus and b is the Berger’s
vector [39].
This phenomenon is also explained by the Hall–Petch relationship.
ߪ଴ = ߪ௜ + ‫ܦܭ‬ିଵ/ଶ
(6)
In this equation, σo is the flow stress, σi is the stress opposing the movement of
dislocations, K is constant and D is the grain size. According to Eq. 6, as the grain
size becomes smaller, flow stress also increases, leading to high strength in the
composite [39].
A close look at Fig. 15, it is revealed that both strength and elongation values
increase with increasing milling time. As mentioned before, this could be attributed
to the reduction in particle size and introducing low interparticle spacing as
documented by several investigators [40,41]. A significant amount of particle
cracking and fragmentation may take place during both milling and extrusion
process which could result in reduction in particle size and interparticle spacing. It
has also been proposed, that because of the higher plastic constraint imposed by
the lower interparticle spacing, that the nucleated voids are unable to coalesce as
easily [40]. A higher work hardening rate has also been observed with decreasing
particle size [41]. This is attributed to the formation of dislocation tangles around
the particles, due to plastic incompatibility between the reinforcement and matrix,
and the formation of a dislocation cell structure with a cell size inversely
proportional to the interparticle spacing [42].

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Fig. 17 illustrates the variation of ultimate tensile strength and elongation value of
Al/ 10 Wt. % Al3Mg2 nanocomposite with milling time.
Fig. 17. The variation of the ultimate tensile strength and elongation of the composites as a
function of milling time.
As shown in Fig. 15 and Fig. 17, mechanical properties of the composite increase
up to 15 h of milling. In fact the mechanical properties of the Al3Mg2/Al
composites are mainly affected by three factors: matrix strength, interface energy
and dispersion of fine Al3Mg2 reinforcing. Furthermore, uniform dispersion of fine
Al3Mg2 particles thorough the matrix and tighter bonding between the Al3Mg2 and
Al are also expected. So, mechanical behavior of the Al3Mg2/Al composites
improves with increasing the milling time.
As shown in Fig. 17, similar trend is observed for ductility enhancement with
increasing milling time. The elongation of the Al3Mg2/Al composites is mainly
affected by the following factors: interfacial characteristics of the

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matrix/reinforcement, dispersion of Al3Mg2, and impurities. Good dispersion of
Al3Mg2 can increase the elongation of the composites. The positive effect of 15 h
milling time on the composite ductility can be because of the lack of harmful
phases in the matrix or interface, less agglomerated Al3Mg2 nanoparticles and good
dispersion of Al3Mg2 in the matrix.
By looking at Fig 17 carefully, one can see no remarkable effect on the mechanical
properties of samples by extending the milling time after 15 h. In the other word, it
seems that the mechanical behavior of Al3Mg2/Al nanocomposite in this study
reaches to a stable state at 15 h milling. In addition, it should be considered this
important point that the level of contamination increases and some undesirable
phases form if the powder is milled for times longer than required [10]. Therefore,
it is desirable that the powder is milled just for the required duration and not any
longer. From these results, it is clear that the time of milling is a critical parameter
to achieve optimum properties of final products. According to this, the optimum
milling time in current study realized to be 15 h.
For a clearer view of potential of materials developed in the present study, a
comparison of the mechanical properties of these materials with results reported in
the literature has been done. Table 3 presents some mechanical properties of well-
known aluminum alloys [43].
Table 3. Mechanical properties of some aluminum alloys [43].
Sample UTS
(MPa)
Microhardness
(HB)
Elongation Density
(g/cm3
)
A1100-H18 165 44 15 2.71
A2024-T351 470 120 20 2.77
A5083-H34 345 75 9 2.66
A6063-T83 255 82 9 2.69
A7075-T6 572 150 11 2.8
Al/10%wt. Al3Mg2
nanocomposite (this study)
623 180 15 2.65
As can be seen, the mechanical properties of Al-10%wt. Al3Mg2 nanocomposite
developed in the present study is comparable well with aluminum alloys.
According to the properties of Al7075 alloys, their applications are in the aircraft
structural parts and other highly stressed structural applications where very high
strength and good ductility are required [43]. By considering the obtained results

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from mechanical properties of Al/10%wt. Al3Mg2 nanocomposite fabricated in
current research, it could be say that this material is a good candidate for
applications in the aerospace, defense and automotive industries.
3.6.3. Fractography
Fig. 18 shows fracture surfaces of as-extruded unreinforced Al (18a), AC10-UM
(18b), AC10-2HM (18(c, d)) and AC10-15HM (18(e, f)) samples after tensile test.
Comparison between fracture surfaces of AC10-UM and unreinforced Al samples,
illustrates different mode of fracture.

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Fig. 18. Fracture surfaces of: (a) unreinforced Al, b) AC10-UM, (c) AC10-2HM, (d) enlarged
image from circle region in (c), (e) AC10-15HM and (f) enlarged image from circle region in (e).
As expected, unreinforced Al specimen in the absence of hard and brittle phases
demonstrates a ductile fracture mode. In this condition, grains may slide on
consecutive planes so that finally the sample separate along planes with 45°
angle
[39]. Since the areas of ductile fracture are detected by dimples in the fracture
surface, more dimples present higher ductility of the material. This fracture
mechanism is called dimple rupture. Also, it could be noted that dimples initiation
in unreinforced Al is concentrated in existing inclusions in the base Al. If the
matrix is free of inclusions, dimples initiation occurs in grain boundaries [44].
Then these dimples grow and come together and ultimately lead to the fracture of
sample. In contrast, in fracture surface of AC10-UM sample large cracks (arrows
in Fig. 18b) with smooth surfaces and sharp edges are observed. Widespread
transverse cracking and transgranular fracture can also been seen. Although some
dimples are observed in the matrix, the fraction is generally shown to be brittle in
nature. These results confirm the reduction of ductility in AC10-UM sample in
contrast with unreinforced Al sample.
For the AC10-2HM specimen, some cracks can also be observed (arrows in Fig.
18c) but less large Al3Mg2 particles can be identified from the fracture surfaces.
Cracked surfaces are basically fine and uneven. Presence of layered fracture
surfaces and fine dimples (Fig. 18d) suggest relatively ductile mode of failure. The
layered fracture surfaces might be an indication of a number of activated slip
planes caused by the presence of nanoparticles within the grains. By looking at Fig.
18e carefully, fine and uneven features are seen on the surface. Many shallow
dimples can also be observed. Since rough and uneven surfaces occupy over a
large area, more energy may be absorbed in fracturing and hence introduce higher
ductility. It is believed that the mechanism of ductile rupture initiation in materials
containing second phase particles is generation of voids at the particles (Fig. 18f).
These cavities grow as the material is strained further, and ultimately they coalesce
by an internal necking mechanism to give the dimpled fracture surface
characteristic of ductile fracture [45]. Furthermore, the ductile fracture is decided
by the size of dimples, meaning more homogenous and dipper dimples resulting a
higher ductility. It is well established that the decreasing of particle size and
increasing fine dimples leads to enhanced elongation values [46]. This fact results

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in reaching about 15% of elongation value in the extruded AC10-15HM sample.
When stress is applied, the matrix plastically deforms, gradually transferring stress
to the particles. The behavior of the particle depends on the relative strength of the
interface and the matrix. When the interface is strong enough the load is
transferred to the particles and fracture occurs as soon as the threshold stresses
reached. This phenomenon is observed in Fig. 18f. As is shown in Fig. 18f, cracks
which can initiate from matrix due to loading are not able to propagate along the
interface, and so they go through the second phase and result in particle cracking in
fracture surfaces, confirming the strong interface (dash line in Fig. 18f) between
matrix and second phase. However, if the matrix is soft with a low work hardening
rate and a weak interface, decohesion prevails (Fig. 18b) [47]. Babout et al.
proposed a model to simulate the fracture behavior in MMCs suggesting that
regardless of the volume fraction of the particles, interface fracture was strength-
controlled, and that decohesion occurred when the average tensile stress in the
particle surpassed a critical value known as the interface strength [47]. When the
matrix is strong enough with high work hardening rate, the load is transferred to
the particles and fracture takes place (Fig. 18f). The work hardening rate of the
matrix in hot-extruded specimens considerably increases and stress transfers to the
particles gradually, which leads to cracked particles increment. As it is indicated in
Fig. 18d by arrows, some fractured particles are apparent which can be the other
characterization of ductile mode of fracture. Extrusion process leads to the
increased particle breakage and formation of fragmented tiny particles and more
uniformity distribution of Al3Mg2 particles in the matrix, as shown in Fig. 11a.
Hence, nucleation sites for cracking are reduced which leads to the improved
tensile properties.
One very important point about fracture surface of milled samples that must be
noted is that debonding between particle and matrix does not occur, as shown in
Fig. 18. This indicates that bonding between particle/matrix in composites
produced by mechanical milling is very strong. In other words, the interface
between particle and matrix has a good metallurgical quality (has no crack or
inclusion) and load transformation from matrix to particle occurs simply and fast
[26].

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Conclusions:
The results of this study are summarized as follows:
(1) The Al3Mg2 particles experienced a large reduction in size, less than 100 nm,
and were uniformly distributed in the Al matrix; no segregation or agglomeration
was observed after ball milling.
(2) Calculating the crystal size of matrix indicates that severe plastic deformation
of powder particles during mechanical milling decreases the crystal size of matrix
to nanometer, so that base grain size decrease to lower than 100 nm. However the
lattice strain of matrix increases during mechanical milling.
(3) The results of mechanical tests show that presence of Al3Mg2 particles
increases the strength and hardness of pure Al but decreases its ductility. The
reason is the increase of barriers across dislocations movement which limits their
movement.
(4) Mechanical milling of composite improved the strength and ductility up to 15 h
milling, owing to effect of ball milling on grain refinement, good dispersion of
Al3Mg2s in the Al matrix and good interfacial bonding between Al3Mg2s and Al
matrix.
(5) For AC10-15HM sample, microhardness increases to 194 HV due to
dislocation density as well as the decrease of the crystallite size. The acceleration
of the grain refinement process by adding Al3Mg2 particles leads to increase in
microhardness. From 15 to 20 h, slower rate of increasing in microhardness may be
attributed to the completion of milling process, and dynamic and static recovery of
powders.
(6) The relative density of the extruded samples was increased with increasing
milling time due to reduction in particle size and also the large surface area of the
fine particles, which induced high reactivity and easily densified. On the other
hand, apparent porosity was decreased because of decreasing the number of the
pores by extrusion.
(7) The fracture surface of unreinforced Al sample indicates the features of ductile
fracture in comparison to AC10-UM sample; in return, the fracture surface of
AC10-2HM and AC10-15HM samples are smoother and finer dimples are

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observed in their fracture surface. This shows that the fracture in these samples
occurred in a ductile manner.
(8) HRTEM results revealed the absence of discontinuities or voids at the
Al3Mg2/Al interface. It also showed Al3Mg2 particles are in intimate contact with
the Al lattice and no intermediate layers or defects are evident at this interface.
Finally it means a clean metallurgical interface between the Al3Mg2 particles and
Al matrix was achieved in this study.
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ACCEPTED MANUSCRIPT
• A novel aluminum-based nanocomposite with high strength and good
ductility was synthesized successfully.
• The distribution of second phase in matrix was improved by mechanical
milling.
• The grain size of matrix and second phase was decreased to nanometer
size during milling.
• The interfacial bonding between matrix and second phase was
improved.