Materials and Design 32 (2011) 3263–3271

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Materials and Design

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Fabrication and evaluation of mechanical and tribological properties of boron

carbide reinforced aluminum matrix nanocomposites
E. Mohammad Sharifi ⇑, F. Karimzadeh, M.H. Enayati
Department of Materials Engineering, Nanotechnology and Advanced Materials Institute, Isfahan University of Technology, Isfahan 84156-83111, Iran

a r t i c l e i n f o a b s t r a c t

Article history: In this study, fabrication and characterization of bulk Al–B4C nanocomposites were investigated. B4C
Received 15 November 2010 nanoparticles were mixed with pure Al powder by ball milling to produce Al–B4C powder. Al–B4C pow-
Accepted 12 February 2011 ders containing different amounts of B4C (5, 10 and 15 wt.%) were subsequently hot pressed to produce
Available online 17 February 2011
bulk nanocomposite samples. Consolidated samples were characterized by hardness, compression and
wear tests. Results showed that the sample with 15 wt.% B4C had the optimum properties. This sample
Keywords: had a value of 164 HV which is significantly higher than 33 HV for pure Al. Also, ultimate compressive
A. Nano materials
strength of the sample was measured to be 485 MPa which is much higher than that for pure Al
C. Powder metallurgy
E. Wear
(130 MPa). The wear resistance of the nanocomposites increased significantly by increasing the B4C con-
tent. Dominant wear mechanisms for Al–B4C nanocomposites were determined to be formation of
mechanical mixed layer on the surface of samples.
Ó 2011 Elsevier Ltd. All rights reserved.

1. Introduction posite was greater than that of the Al–SiC composite. Moreover,
Al–B4C composites have been used in nuclear industries due to
Aluminum matrix composites (AMCs) are attractive materials the specific ability of the B10 isotope to capture neutrons [9].
for structural applications in aircraft, automotive and military Different techniques have been used for fabrication of Al–B4C
industries. High strength to weight ratio, environmental resistance, composites such as liquid phase methods [7,8,10] and solid-state
high stiffness and good wear resistance are characteristics that consolidation (powder metallurgy) [5,6]. However, due to the poor
have spurred more research to develop their applications by fur- wetting between Al and B4C, fabrication of bulk Al–B4C composites
ther improvement in the properties [1–3]. through liquid phase methods (such as casting) is difficult. It has
Ceramic particles such as SiC and Al2O3 are the most widely been reported that Al requires a temperature as high as 1100 °C
used materials for reinforcement of aluminum [1]. Boron carbide to wet the B4C surface completely [11,12]. Processing at such high
(B4C) is one of the most promising ceramic materials due to its temperatures leads to the formation of a series of undesirable com-
attractive properties, including high strength, low density pounds such as Al3BC, AlB2 and Al4C3 due to the chemical reactions
(2.52 g/cm3), extremely high hardness (the third hardest material between Al and B4C. These reaction products degrade the mechan-
after diamond and boron nitride), good chemical stability and neu- ical properties of the composite [13].
tron absorption capability [4–6]. Due to its high hardness, B4C Powder metallurgy processing is another approach to fabricate
could be an alternative to SiC and Al2O3 as a reinforcement phase of Al–B4C composites by mixing the powders of the Al with B4C
in AMCs for applications where a good wear resistance is a major particles, followed by consolidation. Avoiding detrimental interfa-
requirement. Shorowordi et al. investigated wear behavior of Al– cial reaction because of a lower manufacturing temperature and
B4C and Al–SiC composites fabricated by stir casting method under possibility of adding higher amounts of reinforcement particles
the same conditions [7]. They observed that the wear rate and fric- are some advantages of this process [5,6]. Also, in this process
tion coefficient of Al–B4C was lower than those of Al–SiC. Lee et al. the problem of non-wettability of B4C with molten aluminum does
fabricated aluminum matrix composite reinforced with B4C parti- not arise.
cles and SiC particles through the same route (pressureless infiltra- Increasing interest has recently focused on the nanostructured
tion method) and under the same conditions in order to compare AMCs due to their superior properties in comparison to the
the effect of reinforcement type on the tensile properties of the conventional microstructured composites [14]. Moreover, a de-
composites [8]. They reported that the strength of the Al–B4C com- crease in the reinforcement particle size to the nanometer range
can improve mechanical and tribological properties of the AMCs
⇑ Corresponding author. Tel.: +98 312 5201676; fax: +98 311 3912752. [15,16]. But, nanometer particulates are prone to agglomeration
E-mail address: e.mohamadsharifi@ma.iut.ac.ir (E. Mohammad Sharifi). and clustering, resulting in poor dispersing effects in composites.

0261-3069/$ - see front matter Ó 2011 Elsevier Ltd. All rights reserved.
doi:10.1016/j.matdes.2011.02.033
3264 E. Mohammad Sharifi et al. / Materials and Design 32 (2011) 3263–3271

Mechanical alloying (MA) is an attractive powder metallurgy tech- After appropriate milling time, the milled powders were poured
nique that produces uniform dispersion of the reinforcement par- in a uniaxial die made of X40CrMoV51 (AISI H13). Then, as-milled
ticles in the matrix through a repeated process of cold welding, powders were heated to 450 °C and pressed at constant pressure of
fracturing, and rewelding, giving rise to the reinforcement particles 300 MPa. The duration of hot pressing was 30 min. In order to
being well embedded into the matrix particles [15,17]. Moreover, avoid pores formation, the pressure on each specimen was not re-
the high degree of deformation involved may reduce the matrix leased until the specimen cooled down.
grain size to nanometer level, and as a result nanostructured com- X-ray diffractometry (XRD) was used to follow the structural
posite powders can be obtained by using this method [18]. changes of powders after milling and hot pressing. A Philips dif-
Few attempts have so far been made to fabricate B4C nanopar- fractometer (40 kV) with Cu Ka radiation (k = 0.15406 nm) was
ticulate reinforced Al matrix composites [19]. The aim of Khakbiz used for XRD measurements. The grain size of Al was estimated
and Akhlaghi work [19] was only synthesis and characterization from the broadening of XRD peaks using Williamson–Hall method
of Al–B4C nanocomposite powders and consolidation of nanocom- [20].
posite powders was not performed. Moreover, in Khakbiz et.al. The Archimedes technique was used to measure the density of
work, studies on the mechanical and tribological properties of bulk samples. Cylindrical specimens with length-to diameter ratio of
B4C nanoparticulate reinforced aluminum matrix composites were 2:1 (ASTM E9-89a [21]) were prepared from consolidated Al-B4C
not carried out. In the present work, bulk aluminum matrix com- nanocomposites and used for compression tests. A uniaxial com-
posites reinforced with various amounts of B4C nanoparticles were pression testing was performed with an Instron-type machine
produced via MA and hot pressing. Mechanical and tribological (Hounsfield H50-KS model) at the room temperature. The ultimate
properties of bulk nanocomposites were also investigated. More- compression strength (Mpa) and ductility (%) were measured. For
over, in this research, wear behavior of Al–B4C nanocomposites bulk samples, Hardness test was done with HV at a load of 10 kg.
as the function of the characteristics of the mechanical mixed layer Friction and wear properties of the samples were investigated
formed on the worn surfaces were discussed in detail. using a pin-on-disk wear test machine according to the ASTM
G99-05 Standard [22], where AISI 52100 steel with the hardness
of 63 RC was used as the pin. The disk specimens of 50 mm in
2. Materials and experimental procedure diameter were cut from the hot pressed nanocomposites. The tests
were conducted at room temperature at a sliding speed of 0.08 m/s
High purity aluminum powders and boron carbide nanoparti- under an applied load of 20 N under an unlubricated condition. The
cles were used as starting materials. The morphology of as-re- mass loss of the disk specimens was measured at a 25 m interval in
ceived Al and B4C powder particles are shown in Fig. 1. As can be sliding distance, with an analytical balance of 0.1 mg precision. The
seen, as-received Al powders had a random morphology, and their friction coefficients were continuously recorded with sliding
average particle size was about 60 lm (Fig. 1a). The sizes of B4C distance. The worn surfaces were examined using a Philips XL30
nanoparticles were between 10 to 60 nm (Fig. 1b).
Boron carbide nanoparticles were mixed with aluminum pow-
ders and mechanically milled to produce Al–B4C nanocomposites
with 5, 10 and 15 wt.% of reinforcement content, hereafter termed
as the A5, A10, and A15, respectively. To minimize the extreme
cold welding of aluminum powders, 0.3 wt.% of stearic acid was
used as a process control agent (PCA).
Ball milling was executed at a rotation speed of 600 rpm and
the ball-powder mass ratio was 10:1. In this work, hardened chro-
mium steel vial (125 ml) containing five steel balls (high chro-
mium–carbon steel) with a diameter of 20 mm was used. The
vial was evacuated and then filled with pure argon gas to prevent
oxidation during the milling process.
The cross-section of powder particles was prepared by mount-
ing a small amount of powder in a resin followed by conventional
grinding and polishing methods. The hardness of cross-section of
powders was determined by microhardness test using a Vickers
indenter at the load of 100 g. Ten indentations were made on each Fig. 2. XRD patterns taken from Al–10 wt.% B4C mixture milled for different hours.
sample to obtain an average value of hardness.

Fig. 1. Morphology of as-received powder particles: (a) SEM micrograph of Al, and (b) TEM picture of B4C.
 E. Mohammad Sharifi et al. / Materials and Design 32 (2011) 3263–3271 3265

scanning electron microscope (SEM) with an energy dispersive X-

ray spectrometer (EDS).

3. Results and discussion

3.1. Nanocomposite samples preparation

In the present work, variation of Al matrix grain size during

milling and hardness behavior of the composite powders were
evaluated to determine the optimum time for MA. Fig. 2 shows
the XRD patterns of A10 powders after different milling times up
to 40 h. As can be seen in Fig. 2, the significant phenomenon is peak
broadening which occurs due to a decrease in grain size and an in-
crease in lattice microstrain [18]. Due to the low content, the fine
Fig. 3. Variation in the Al average grain size during milling of the Al–10 wt.% B4C particle size (10–60 nm) and a limited scattering factor of B4C, its
composite powder, as derived from XRD peak broadening. XRD peaks were not apparent.
The Al grain size was estimated from broadening of XRD peaks
using Williamson–Hall method. Fig. 3 plots the Al grain size as a
function of milling time. The initial grain size of as-received alumi-
num was about 340 nm. As a result of milling, Al readily achieved a
nanocrystalline structure. The Al grain size after 2 h of milling time
was about 74 nm, which changed to 52 nm as milling time in-
creased to 5 h. The Al grain size appeared to approach a constant
value, saturated value, as milling proceeded.
MA is characterized by intensive plastic deformation of powder
particles at extremely high strain rate, creating a high density of
lattice defects, mainly dislocations. At the same time, recovery
phenomena can take place, reducing the density of lattice defects
[18]. The progressive accumulation and interaction of dislocations
lead to formation of dislocation cell structure which subsequently
creates low-angle grain boundaries [23]. With prolonged process-
ing, this structure is transformed to a fully nanocrystalline struc-
Fig. 4. Microhardness values as a function of milling time for Al–10 wt.% B4C ture [23,24]. The observed initial sharp decrease in Al grain size
composite powder. during the initial milling stage and the ultimate grain size achieved

Fig. 5. SEM micrographs of Al–10 wt.% B4C powder particles after (a) 15 min, (b) 2 h, and (c) 5 h of milling time.
3266 E. Mohammad Sharifi et al. / Materials and Design 32 (2011) 3263–3271

Fig. 6. The XRD patterns of A5, A10, and A15 samples after milling and hot pressing.
Fig. 9. Variations of wear mass loss as a function of sliding distance for A5, A10, and
A15 nanocomposite samples.

Table 1
Compressive properties of nanocomposite samples.

Samples Ultimate compressive strength (MPa) Ductility (%)

A5 371 16.3
A10 433 14.1
A15 485 12.1

Fig. 4 presents the change in Vickers microhardness of the A10

powders as a function of milling time. As can be seen, increasing
the milling time causes a raise in microhardness of powders. As
mentioned earlier, MA leads to extreme refinement of the micro-
Fig. 7. Average hardness value of consolidated nanocomposites as a function of structure, finally resulting in nanocrystalline structure with high-
B4C wt.%. lattice microstrain. Moreover, the presence of hard reinforcing par-
ticles might enhance the work hardening rate of the matrix, result-
ing in an increase in hardness. The hardness of the pure Al powder
before milling was approximately 33 HV. At the beginning of mill-
ing, the hardness of composite powder increased rapidly to over
90 HV. But the rate of increase in hardness value decreased as mill-
ing time increased gradually. After 5 h of milling time, the microh-
ardness value of composite powder was about 120 HV. Increasing
milling time from 5 to 40 h increased the hardness value at a much
slower rate. It can be assumed that these hardness results are
attributable to not only Al grain refinement and large strain intro-
duced by MA, but also to the uniform dispersion of B4C nanoparti-
cles [26].
Fig. 5 shows SEM images of A10 powder particles after differ-
ent milling times. During the early stages of MA, B4C particles
adhere to aluminum powder particles (Fig. 5a). Al powders with
B4C particles sticking onto them are cold welded with other sur-
faces. As a result, the B4C particles are embedded in the Al ma-
trix. The work hardening induced during MA process intensifies
the initiation and propagation of cracks within powder particles
(Fig. 5b). Propagation of these cracks through the matrix leads to
the fracture of aluminum particles. These new fractured surfaces
Fig. 8. Compressive stress–strain curves for A5, A10, and A15 nanocomposite
samples. with B4C particles on them would weld with other surfaces.
With the repeated fracturing and cold welding processes that
take place during MA, B4C particles are finally distributed uni-
formly within the Al matrix [19]. Fig. 5c shows the morphology
can be discussed with respect to the competing rates of creating of A10 powder particles after 5 h of milling. Average particles
dislocations due to plastic deformation and work hardening and size in this stage was about 10 lm.
recovery phenomena [23–25]. The former has a higher rate at ini- The observed grain size variations and hardness behavior in the
tial stage of MA, leading to a rapid decrease in grain size. Thereby, case of A5 and A15 powders were largely similar, and hence, not
the ultimate grain size achieved by MA is determined by the bal- reproduced here. So, According to the grain size measurement
ance between creation and annihilation of dislocations [24]. results and hardness behavior of the milled powders, 5 h is
 E. Mohammad Sharifi et al. / Materials and Design 32 (2011) 3263–3271 3267

Fig. 6 shows XRD patterns of A5, A10, and A15 samples after
milling and hot pressing at 450 °C under 300 MPa pressure. No so-
lid-state reaction between aluminum and B4C in all different com-
positions during hot pressing was detected. Moreover, the grain
size of aluminum after hot pressing was calculated from broaden-
ing of XRD peaks taken from samples using the Williamson–Hall
method. In all the samples, no significant grain growth observed
after hot pressing and aluminum grain size remained in the nano-
meter range (about 80 nm). In this case, the B4C nanoparticles pin
Al grain boundaries and prevent significant grain growth during
hot pressing. The relative density value of consolidated nanocom-
posites was measured to be about 98% (considering standard devi-
ation of 0.5%).

3.2. Mechanical properties evaluation

Fig. 10. Comparison between wear rate of A5, A10, and A15 nanocomposite
samples. The results of hardness tests of consolidated nanocomposites
and pure aluminum are presented in Fig. 7. The average hard-
considered as an optimum time for MA, and the powders milled for ness values of A5, A10 and A15 samples were measured to be
5 h are used for hot pressing. 112 HV, 136 HV and 164 HV, respectively. These obtained hard-

Fig. 11. Low magnification SEM micrographs showing the worn surface of nanocomposite samples: (a and b) A5, (c and d) A10, and (e and f) A15.
3268 E. Mohammad Sharifi et al. / Materials and Design 32 (2011) 3263–3271

ness values for nanocomposite samples are significantly higher

than those for pure aluminum (33 HV). The enhancement of
hardness by increasing the weight percentage of the B4C nano-
particles mainly results from: (a) the presence of extremely
harder B4C nanoparticles in the Al matrix, and (b) a higher con-
straint to the localized matrix deformation during indentation
[27].
Stress–strain curves for nanocomposite samples are illus-
trated in Fig. 8 and their compressive properties are summa-
rized in Table 1. The ultimate compressive strength for the
nanocomposites increased with an increase in the content of
the B4C nanoparticles, and it was much higher than that for
pure Al (130 MPa). As mentioned earlier, MA produces uniform
dispersion of the reinforcement particles in the matrix because
repetitive fracturing and cold welding process cause the rein-
forcement particles to be well embedded into each aluminum
particle. So, improving the compression strength by introducing
B4C nanoparticles could be explained by the homogeneous dis-
tribution of the B4C nanoparticles in the aluminum matrix.
These particles prevent the movement of dislocations in pure
aluminum matrix through dispersion strengthening mechanism.
Moreover, increasing the amount of B4C nanoparticles leads to
a decrease in the distance between them which causes an in-
crease in the required stress for dislocations movement between
the B4C nanoparticles. It consequently increases the material
strength and causes decrease in the ductility (see Fig. 8 and
Table 1) [28,29].

3.3. Wear measurements Fig. 12. (a) High magnification SEM micrograph of darker layer formed on the worn
surface of A5 sample, and (b) its EDS analysis.

Fig. 9 shows the wear mass loss of the nanocomposite samples

as a function of sliding distance. For any given specimen, the mass
loss continuously increases with increasing sliding distance. The
steady state values of the wear rate in the case of A5, A10 and
A15 samples are compared in Fig. 10. It can be clearly seen that
the wear rate of the nanocomposites is considerably improved by
the addition of the reinforcement particles and decreased by
increasing B4C weight fraction from 5 to 15 wt.%.
To discover the governing wear mechanism on each nanocom-
poite samples, SEM micrographs of worn surfaces in addition to
the EDS analyses were employed. The low magnification SEM
micrographs from the wear tracks of the A5, A10 and A15 samples
are shown in Fig. 11. As can be seen, all the worn surfaces are cov-
ered by a darker layer in most regions and the extent of coverage
provided by this layer is increased as the amount of reinforcement
is increased.
Higher magnification SEM micrograph and corresponding EDS
analysis of the darker layer (marked A in Fig. 11b) and bright area
(marked B in Fig. 11b) on worn surface of A5 sample are shown in
Fig. 12 and Fig. 13, respectively. SEM micrograph of the darker
layer shows formation of narrow wear grooves and some micro-
cracks (Fig. 12a), whereas SEM micrograph of bright area shows
heavy flow of materials along sliding direction which indicates
greater degree of wear and localized adhesion between the speci-
men surface and counter body (Fig. 13a). EDS analysis confirmed
that the darker area on A5 sample worn surface beside aluminum
element contained a considerable amount of oxygen and iron
(Fig. 12b). In contrast, EDS analysis of the bright area on A5 sample
worn surface showed only the presence of Al (Fig. 13b). The pres-
ence of Fe in the darker layer implies the transfer of Fe from the Fig. 13. (a) High magnification SEM micrograph of bright area on the worn surface
steel pin to the worn surface while the oxygen element predicates of A5 sample, and (b) its EDS analysis.
the oxidation reaction. These results imply that transfer and
mechanical mixing of materials have taken place between the
two sliding surfaces, and a mechanically mixed layer (MML) has researchers [30–34]. MML can act as an effective insulation layer
formed on the darker areas on the worn surface. The formation between the pin and the disk, which prevents metal to metal con-
of such a layer has been observed previously by some other tact [31,34].
 E. Mohammad Sharifi et al. / Materials and Design 32 (2011) 3263–3271 3269

Fig. 14. (a) High magnification SEM micrograph of darker layer formed on the worn
surface of A10 sample, and (b) its EDS analysis.

Fig. 16. Cross-sectional SEM micrographs of the wear track of: (a) A5, (b) A10, and
(c) A15 samples, showing presence of the MML.

layers. As a result, it seems that a mechanical mixing/oxidation

process was the controlling wear mechanism in the case of all
composites. Moreover, the presence of the MML can be also
visualized in cross sectional SEM micrographs taken from the
wear track of the A5, A10 and A15 samples (Fig. 16). The thick-
ness of the MML formed on the worn surface of A5, A10 and A15
samples was measured and the average values of at least ten
measurements obtained by SEM observation are reported in Ta-
ble 2. It was shown that the MML with higher thickness effec-
tively reduced wear rate [35].
As can be observed, the darker layer formed on the worn surface
Fig. 15. (a) High magnification SEM micrograph of darker layer formed on the worn of A15 sample (Fig. 15a) has less damaged regions and seems to be
surface of A15 sample, and (b) its EDS analysis. more stable compared to A10 (Fig. 14a) and A5 samples (Fig. 12a).
By comparing EDS analyses of the MML formed on the worn sur-
faces of A5, A10 and A15 samples, it is revealed that the A15 spec-
Higher magnification SEM micrograph and corresponding EDS imen had the highest Fe and O contents. Moreover, according to
analysis of the darker layer formed on A10 and A15 worn sur- the MML thickness measuring results (Table 2), it is clear that
faces are shown in Fig. 14 and Fig. 15, respectively. Similar to the thickness of the MML increases with increasing reinforcement
A5 sample, EDS analyses of A10 and A15 worn surfaces showed content. Thus, these results show that higher reinforcement con-
the presence of Fe and O in chemical composition of the darker tent in the nanocomposite promotes stronger material transfer
3270 E. Mohammad Sharifi et al. / Materials and Design 32 (2011) 3263–3271

Table 2 the surface. MML can be treated as a layer of solid lubricant since
The average thickness of the MML formed on the worn surface of nanocomposite it contains some oxide compounds [33,36]. So, formation of such a
samples.
layer on the surface of nanocomposite samples during sliding wear
Samples Average thickness of MML (lm) reduces the COF values. It should be noted that A15 sample exhib-
A5 3 its the lowest COF values (Fig. 17c). This can be explained by the
A10 5 presence of higher amount of oxide compounds content on the
A15 9 worn surface of A15 sample compared to A10 and A5 samples
(See EDS analysis results of the MML formed on the worn surface
from the counterface and oxidation reaction, and consequently of nanocomposites).
causes faster formation of more protective MML with higher thick- Moreover, there were some sudden fluctuations in the COF, par-
ness and higher amount of oxide compounds content on the worn ticularly in the case of A5 sample (Fig. 17a). As pointed out earlier,
surface, leading to the lower wear rate. This has also been con- the MML formed on the worn surface of A5 sample was less stable
firmed by other research findings [32–34]. and had less coverage in comparison with A10 and A15 samples.
The variation of coefficient of friction (COF) for the nanocom- These fluctuations in the COF can be due to delamination of weakly
posite samples as a function of the sliding distance is presented bonded and thin MML from the worn surface, which leads to leav-
in Fig. 17. In all the samples, the COF values gradually increased ing behind the fresh surface. This phenomenon was reflected as the
up to about 50 m sliding distance, and then decreased to reach rel- jump in the fluctuations at some points in the plots [37].
atively steady values (0.3–0.4). This trend in the variation of COF The main conclusion drawn from the above discussions is that
may be due to the formation of the mechanically mixed layer on the wear performance of worn samples is largely determined by

Fig. 17. The variations in coefficient of friction of the nanocomposite samples: (a) A5, (b) A10, and (c) A15.
 E. Mohammad Sharifi et al. / Materials and Design 32 (2011) 3263–3271 3271

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