Materials Sciences and Applicatio ns, 2011, 2, 1340-1348
doi:10.4236/msa.2011.29182 Published Online September 2011 (http://www.SciRP.org/journal/msa)
Evaluation of Characteristics of Interfacial Phases
Produced in Al/Ni3Al Composite during
Manufacturing
Seyed Abdolkarim Sajjadi1*, Maryam Abbasi2, Mazyar Azadbeh2
1Department of Materials Science and Metallurgical Engineering, Engineering Faculty, Ferdowsi University of Mashhad, Mashhad,
Iran; 2Faculty of Materials Science and Metallurgical Engineering, Sahand University of Technology, Tabriz, Iran.
Received February 24th, 2011; revised March 21st, 2011; accepted July 4th, 2011.
ABSTRACT
Metal matrix composites (MMCs) are currently being investigated because of their superior properties. The properties
are mainly attributed to the efficiency of the load transfer from the matrix to the reinforcements through the ma-
trix-reinforcement interface. The aim of this study is to investigate the effect of manufacturing parameters on the micro-
structure and morphology of the interface and the aluminide phases formed at the matrix-reinforcement interfaces. The
parameters are: milling time to fabricate Ni3Al, method of mixing of Ni3Al and Al powders, compaction pressure and
sintering temperature. The composite studied in this research was Al/5 Vol% Ni3Al made from two different types of
Ni3Al powders. The results showed that compacting and sintering at higher levels lead to the transformation of Ni3Al
particles to thin layers of Al3Ni. It was also shown that the prolonged milling time to produce Ni 3Al reinforcements an d
the prolonged ball milling procedu re for mixing the powders, both, p romote the diffusion process at reinforcement/ ma-
trix interface.
Keywords: Metal-Matrix Composites (MMCs), Interf ace , X-Ray Diffraction (XRD), Powder Processing
1. Introduction
Aluminum-matrix composites (AMC) reinforced with
intermetallic compounds have been proposed as substi-
tutes for ceramic reinforced composites due to their ade-
quate mechanical properties. The key challenge in tradi-
tional metal-matrix composite production is to obtain a
good reinforcement-matrix interfacial bonding. Ceramic
reinforcements are normally not well compatible with the
metal matrix, which results in poor reinforcement-matrix
interfacial bonding. Such interfaces reduce mechanical
properties and increase corrosion sensitivity. The new
class of reinforcements namely intermetallic compounds,
with their extremely good mechanical properties offer
new solutions to this problem.
There have been some studies on the properties of
aluminum-matrix composites reinforced with intermetal-
lic particles in the binary systems of Fe-Al [1], Ti-Al [2]
and Ni-Al [3-7]. They showed Ni3Al is one of the best
known and characterized reinforcements amongst the
intermetallics.
In composites, blending or mixing of reinforcement
and matrix powders is just as important. because it con-
trols the final distribution of reinforcement particles and
porosity in green compacts. This, in turn, strongly affects
the mechanical properties of the composites. Segregation
and clustering are the common problems associated with
the present state-of-the-art blending or mixing methods.
The segregation and clustering during blending can be
overcome by a technique developed during the 1960s
called mechanical alloying [8-10]. Mechanical alloying is
a dry, high-energy ball-milling process for producing
composite powders with a fine controlled microstructure
[11].
Ball milling causes a uniform distribution of particles
and metal powders with a satisfactory microstructure
after compaction [12]. The distribution of reinforcements
affects the compressibility of the composite powder. It
has been reported that when mechanical milling is ap-
plied on aluminum powder, powder compressibility is
reduced and the green density decreases with increasing
milling time [13,14].
It is difficult to manufacture highly reinforced com-

Evaluation of Characteristics of Interfacial Phases Produced in Al/NiAl Composite during Manufacturing1341
3
posites because the green density decreases with in-
creasing fraction of reinforcement particles. This often
leads to insufficient strength for supporting subsequent
processes such as sintering and machining. In other
words, mechanically alloyed Al powder containing a fine
and homogeneous distribution of submicroscopic inter-
metallics is very hard and, therefore, its compressibility
is low. Also, surface oxide films and hardness make
milled Al powders very difficult to sinter. On the other
hand, the properties of composites are very sensitive to
sintering temperature because the reinforcement-matrix
interface changes by changing the sintering temperature
[15-17]. The reinforcement-matrix interface has a strong
influence on the mechanical properties of any compos-
ite.
2. Experimental Procedure
Aluminum-matrix composites were produced using pure
aluminum as metal matrix and 5 Vol% Ni3Al particles as
reinforcement. Elemental Al powder (Merck-1056-99%,
<160 μm, flaky morphology) and Ni (Merck-112277-
99.5%, <10 μm, spherical morphology) were used as raw
materials. The reinforcement particles (Ni3Al) were pro-
duced under two different milling conditions of pure Ni
and Al as shown in Table 1. As the table indicates, all
parameters for production of the two types of intermetal-
lic compounds are the same except for the milling time.
In order to evaluate the effect of mixing conditions on
the properties of composites, two types of mixing proc-
esses such as blending and high energy ball milling
(planetary ball mill) were applied. The blending process
was carried out in a tumbling mixer for 30 minutes. The
ball milling of Al and Ni3Al powders was conducted us-
ing a planetary apparatus (model FP2) to produce com-
posite powder particles. The powders and hardened 20
mm diameter steel balls were sealed in a hardened
stainless steel vial at room temperature. The ball-to-
powder weight ratio and rotational speed were 20:1 and
300 rpm, respectively. To prevent excessive welding of
particles in the vial, 1 wt% stearic acid was added to the
system. To avoid oxidation, the entire process was per-
formed in argon atmosphere. After 12 hours of milling,
the powders were analyzed for morphology and struc-
ture.
Table 1. Milling condition parameters for producing Ni3Al
powders.
Speed of milling 550 rpm
Milling time 15 and 55 h
Weight ball/powder ratio 20:1
Ball diameter 20 mm
Ball material stainless steel
Process control agent Stearic acid (1 wt%)
Atmosphere Argon
The blended and milled powders were cold compacted
uniaxially in a floating die at two compacting pressures
of 400 and 800 MPa. After ejecting the specimens from
the die, their densities were measured by volumetric
method. Weight and dimensions of the compacts were
measured by an accurate balance (±0.1 mg) and a mi-
crometer (±0.1 mm). The error of measurement was
found to be less than 1%.
The compacted test specimens were sintered at three
temperatures of 580˚C, 620˚C and 650˚C in a vacuum
furnace for 30 minutes followed by furnace cooling.
Densities of the sintered parts were determined according
to Ar- chimedes principle (DIN ISO 3369).
Microstructure observations were made by an optical
microscope and a scanning electron microscope model
LEO 440 i. In addition, size distribution of powder parti-
cles was evaluated by Clemex Image Analyzer. In order
to study the distribution of elements within the interfacial
layers, chemical composition was measured by the en-
ergy dispersive X-ray (EDX) method. The interfacial
phases under different manufacturing conditions were
characterized by Bruker; Advance-D8 X-ray diffracto-
meter.
Vickers macro-hardness tests were performed on the
sectioned and polished specimens. The values reported
are the average of at least 10 indentations by applying 30
kg load for 10 s.
3. Results and Discussion
3.1. Production of Particulate Reinforcement
Diffusion of aluminum atoms into the nickel system dur-
ing ball milling process resulted in the formation of
Ni(Al) solid solution and then Ni3Al compound as a final
product after 15 hours of milling. X-ray diffraction pat-
tern of the milled powders is shown in Figure 1. The
point analysis of chemical composition of powders using
energy dispersive X-ray (EDX) method at random sites
indicated a 75:25 atomic percentage ratio of Ni to Al,
which corresponds to Ni3Al compound. It has been
claimed that the same composition has been produced after
Figure 1. X-ray diffraction of milled powders.

Evaluation of Characteristics of Interfacial Phases Produced in Al/Ni3Al Composite during Manufacturing
1342
prolonged ball milling treatment (up to 55 hours) [18].
XRD peak broadening represents the variation in the
crystallite size of nanostructure Ni3Al and accumulation
of lattice strain with increasing time of milling. The XRD
peaks were analyzed by Williamson-Hall method. The
size and lattice strain of 15-hour milled powders were
4.81 nm and 0.005%, respectively. While the values of
these parameters reach 19.19 nm and 0.88% at 55 hours
of ball milling time.
Figure 2 shows morphology of the two different types
of Ni3Al produced after 15 and 55 hours of milling. The
particle size distribution of the powders is illustrated in
Figure 3. It is observed that the size of 55-hour milled
powder is greater than that of the 15-hour milled powder.
During initial milling times, sequences of welding and
fracturing are perceived. After passing 15 hours of mill-
ing process, welding dominates fracturing. Cold welding
between powder particles occurs during long milling
time (55 hours) which causes increased average particle
size. Image analysis indicated that spherical diameter of
Ni3Al powders after 15 and 55 hours of milling were 6.7
and 10.9 μm and their aspect ratio were 1.96 and 2.61,
respectively.
Since Ni3Al particles are brittle, it is expected that
during milling fracture should predominate over cold
welding leading to particle refinement instead of particle
coarsening. However, as mentioned earlier, it was ob-
served that the particle size of 55-hour milled powder is
larger than 15-hour milled powder. Therefore, in the case
of Ni3Al particles the results showed to be the opposite.
This is due to the fact that mechanical milling causes
formation of nanocrystalline powder particles which re-
sults in grain boundary sliding or migration during mill-
ing. As a result, nanocrystalline structure develops duc-
tility in Ni3Al compound and encourages cold welding.
Such line of reasoning has been postulated by Kumaran
et al. [19] for coarsening of TiAl powders during high
milling process.
(a) (b)
Figure 2. Morphology of Ni3Al powders obtained by different milling time; (a) 15 hours, and (b) 55 hours.
(a) (b)
Figure 3. The particle size distribution of Ni3Al powders obtained by different milling time; (a) 15 hours, and (b) 55 hours.

Evaluation of Characteristics of Interfacial Phases Produced in Al/NiAl Composite during Manufacturing 1343
3
3.2. Distribution of Ni3Al Particles in Al
Powders
Blending of aluminum and reinforcement powders
(Ni3Al) results in heterogeneous size distribution of the
reinforcement particles (Figure 4(a)). The largest size is
about 50 μm which is dispersed heterogeneously between
aluminum powders. It is certain that the homogenization
brings about improved physical and mechanical proper-
ties. In the milling process, impact of milling balls on the
powders causes plastic deformation and formation of
flake-shaped soft aluminum powders and brittleness and
fracturing of Ni3Al particles. Thus the fine reinforce-
ment particles are cold welded to aluminum powder
flakes. Continuation of milling causes cold welding of
these flakes together. Consequently, a composite powder
consisting of aluminum and Ni3Al is formed. More uni-
form distribution of reinforcement particles in Al powder
is shown in Figure 4(b), which is the result of milling for
12 hours.
Another factor influencing composite properties is the
compaction pressure. To investigate the effect of this
parameter, two different compaction pressures were ap-
plied. Although milling enhances the distribution of re-
inforcements, it results in lower compressibility of ball
milled powder as compared to the blended powder.
Therefore, compaction of ball milled powders needs
more pressure. Therefore, compression of the powders at
400MPa was impossible.
As a general rule, due to the hardening effect of the
mechanical milling, green density of milled powders
decreases. The other reason for the incompressibility of
the high energy milled powder at the lower pressure is
the flat shape of the powder [13,14]. The applied pres-
sure (400 MPa) is not high enough to activate higher
plastic deformation due to the flattened morphology. At
higher compaction pressure (800 MPa), green density of
flattened composite (due to plastic deformation of pow-
ders) comes up to only 2.65 g/cm3 which equals to 90.5%
of the theoretical density. Whereas, compaction of
blended powders is easier and green density reaches 89%
and 95% of the theoretical density at 400 and 800 MPa,
respectively.
3.3. Effect of Manufacturing Parameters on the
Interfacial Bonding
The main parameters such as: duration of mechanical
alloying, mixing conditions, compacting pressure and
sintering temperature strongly affect the constituent
phases and thickness of multi-layers around reinforce-
ment particles.
Figure 5 shows microstructure of Al/5 Vol% Ni3Al
samples compacted at 400 MPa and sintered at 620˚C.
The powders were produced by milling at different times
(15 and 55 hours). As it is observed, the size of 55-hour
milled powder is greater than the powder milled for 15
hours due to cold welding of the flakes together.
In the case of Al/5 Vol% Ni3Al (15-hour milled)
composite sintered at 620˚C, the effect of strain accu-
mulation on the diffusion, phase transformation and in-
terfacial bonding is less pronounced as compared to the
composite with its reinforcement powder mechanically
alloyed for 55 hours and then sintered at the same sinter-
ing temperature. The comparison of microstructures of
the composites (Figure 5) shows that in 15-hour milled
specimen, Ni3Al particles are surrounded by diffusion
layers and Ni-aluminide phases. While, in the specimen
with 55-hour milled Ni3Al, diffusion and reaction rates
are so high that Ni3Al particles are mostly replaced by
Al3Ni and Al3Ni2. The stored elastic energy and higher
dislocation densities due to longer milling times are the
main parameters in creating the condition which leads to
an increase in the reaction rate by providing short diffu-
(a) (b)
Figure 4. Morphology of Al/5 vol% Ni3Al; the indicated phases show reinforcement particles (a) Blended powders, (b)
12-hour Milled powder s.

Evaluation of Characteristics of Interfacial Phases Produced in Al/NiAl Composite during Manufacturing
1344 3
(a) (b)
Figure 5. Microstructures of Al/5 vol% Ni3Al composite sintered at 620˚C for 30 min. (a) after 55 hours; and (b) after 15
hours milling.
sion paths . The results are in good agreement with find-
ings of Lieblich et al. [20]. They have claimed that the
reaction of Ni3Al and aluminum matrix and dissolution
of Ni3Al particles follows a parabolic law and occurs
during three steps. In the first step, nucleation and growth
of Al3Ni phase occurs which surrounds the dissolving
Ni3Al particles; in the second step and at longer times,
the Al3Ni2 phase nucleates and grows between the Ni3Al
and the Al3Ni layers; and in the third step, the Ni3Al and
Al3Ni2 phases completely dissolve and the stable Al3Ni
phase remains.
It has also been reported [21] that increasing milling
time generates more defects such as vacancies and stored
strain energy in the powders. The defects formed due to
the large plastic deformation result in an increase in reac-
tion rate by providing short circuit diffusion paths.
Figures 6(a) to (c) show microstructures of mixed
Al/Ni3Al powders compacted at 400 MPa and then sin-
tered at three different temperatures (580˚C, 620˚C and
650˚C). The microstructure of compacted powder at
400MPa and sintered at 580˚C shows that Ni3Al rein-
forcement particles remain almost separated, and the
bonding between matrix and the intermetallic particles is
not strong (Figures 6(a)). Also, the thickness of the dif-
fusion layers around the reinforcements is negligible.
Higher sintering temperatures facilitate the diffusion
of elements at matrix/reinforcement interface. Sintering
at 620˚C causes the formation of multilayer bonds and
conjunction between matrix and reinforcements (Figure
6(b)). Chemical analysis of samples by the energy dis-
persive X-ray (EDX) method on reaction layers at rein-
forcement/matrix interface of composite sintered at
620˚C has shown the formation of thin layers Ni2Al3 and
Al3Ni around the remaining Ni3Al nucleus.
At 650˚C sintering temperature, the Ni3Al and Al3Ni2
phases completely dissolve and the equilibrium Al3Ni
phase remains (Figure 6(c)). The reinforcement size
produced by sintering at 650˚C is larger than the one
created at 620˚C. Furthermore, the microstructure of re-
inforcements in sintered composite at 650˚C is lamellar.
Also, the voids in the sintered composite at 650˚C are
more or less rounded as compared to the other samples
sintered at lower temperatures (Figures 6(a) and 5(b)).
Because of the decrease in number of voids with in-
creasing sintering temperature, the density of the com-
posite sintered at 650˚C is 2.6 g/cm3 which is higher than
the other ones (2.57 and 2.59 g/cm3 corresponding to
580˚C and 620˚C sintering temperature, respectively).
X-ray diffraction patterns shown in Figure 7 confirm
the formation of the above mentioned phases at rein-
forcement/matrix interface at different sintering tem-
peratures.
Figures 6(d) to (f) illustrate the influence of compac-
tion pressure on microstructure of interface layers of the
blended composites, compacted at 800 MPa, and sintered
at different temperatures. As it was expected the increas-
ing number of contacts of the reinforcement to matrix,
and consequently their reaction with each other, facili-
tated phase transformation of reinforcement at the inter-
face. Comparing Figures 6(b) with Figure 6(d) shows
that at the higher compaction pressure, the microstructure
of composite sintered at 580˚C is similar to that com-
pacted at 400 MPa and sintered at 620˚C. Increasing the
compaction pressure to 800MPa for blended grade pow-
ders caused the highest sintered density of 2.7 g/cm3 after
sintering at 620˚C.
The effect of dispersion condition of Ni3Al particles in
composite powder on phase transformation of reinforce-
ments is noticeable. Microstructural homogeneity of re-
inforcement is strongly affected by the milling operation.
The comparison of the XRD patterns of blended and
milled starting powders sintered at different temperatures

Evaluation of Characteristics of Interfacial Phases Produced in Al/NiAl Composite during Manufacturing1345
3
Figure 6. Microstructures of the composites compacted and sintered at different conditions. (a)(d)(g) Sintered at 580˚C;
(b)(e)(h) Sintered at 620˚C; (c)(f)(i) Sintered at 650˚C.
Figure 7. X-ray diffraction pattern of as sintered blended grade samples compacted at 400 MPa.

Evaluation of Characteristics of Interfacial Phases Produced in Al/NiAl Composite during Manufacturing 1346 3
reveal that milling promotes phase transformation of
Ni3Al to Al3Ni with respect to the blended one so that the
two phases can be observed at all sintering temperatures
(see Figure 7 and Figure 8).
It is a well known fact that milling operation can cause
the decrease in density of composites. As expected, the
lower green density resulted in the drop of sintered den-
sity values. The maximum value of density in milled
grade was 2.65 g/cm3. The results presented so far indi-
cate that the density of blended and milled grade samples
sintered at 650˚C is 2.62 g/cm3. This can be attributed to
the composites with Al3Ni reinforcement particles.
Hardness measurement revealed sticking of reinfor-
ment to matrix at the interface. The results presented in
Table 2 followed the same pattern described above. The
data shows that the hardness values of sintered specimen
produced from ball milled powder are markedly higher
than those made of the blended powder grade. Work
hardening of powders and the uniform distribution of
reinforcement on Al powder introduced by ball milling
process are considered to be the main reasons for such a
trend, especially at temperature at which multilayer
forms at interface (620˚C).
In blended grade, increasing the sintering temperature
from 580˚C to 620˚C causes an increase in hardness val-
ues, while sintering at 650˚C results in the lower hard-
ness. This can be attributed to the formation of a softer
intermetallic phase such as Al3Ni. A similar drop in
hardness of the blended grade compacted at 400 MPa is
also expected. Approximately the same values of hard-
ness were observed in samples sintered at 620˚C and
650˚C, and is considered to be the consequence of void
elimina- tion that occurred at 650˚C. In other words,
hardness drop due to the formation of softer intermetallic
phase can be compensated by increasing hardness result-
ing from void elimination.
Figure 8. X-ray diffraction pattern of ball milled grade samples compacted at 800 MPa and sintered at three different tem-
peratures.
Table 2. Influence of manufacturing parameters on hardness of sintered Al/Ni3Al.
Sintering temperature
Mixing condition Compacting pressure 580 (˚C) 620(˚C) 650(˚C)
Blended grade 400 (MPa) 20 HV30 30 HV30 30 HV30
Blended grade 800 (MPa) 35 HV30 45 HV30 34 HV30
Ball milled grade 800 (MPa) 87 HV30 95 HV30 61 HV30
4. Conclusions
Interfacial bonding at matrix/reinforcement interface is
significantly affected by mechanical milling time in pro-
ducing Ni3Al particles, mixing condition of Al and Ni3Al
powders, compacting pressure and sintering tempera-
ture.
 Mechanical milling produces Ni3Al reinforcement
powders that promote the diffusion process owing to
large deformation of powder particles as well as more
stored strain energy.
 The 55-hour mechanically milled reinforcements pro-
vide an effective means of thicker diffusive layers
formation around the reinforcements during sintering,

Evaluation of Characteristics of Interfacial Phases Produced in Al/NiAl Composite during Manufacturing1347
3
as compared to the thinner interfaces in the composite
containing 15-hour mechanically milled intermetallic
particles at the same sintering temperature.
 It was shown that finer distribution of reinforcement
is significantly more pronounced in ball milled grade
than blended Al/Ni3Al. This apparently improved
mechanical properties of sintered composite and en-
couraged phase transformation of reinforcement.
 In blended grade samples compacted at 400 MPa, low
sintering temperature (e.g. 580˚C) cannot eliminate
the pores around intermetallic particles that generated
primarily during compacting by decohesion of the
matrix-particle interface. Therefore, the bonding be-
tween the matrix and intermetallic particles is not
strong and thickness of diffusion layer around the in-
termellic is negligible.
 Microstructure of the composite sintered at 580˚C
shows that Ni3Al particle reinforcements remain ap-
proximately intact. With increasing sintering tem-
perature, diffusive layers will be extended and reac-
tion phases improve the bonding.
 More contacts between reinforcements and matrix
caused by higher compaction pressure enhance the
reaction at interface; therefore, at higher compaction
pressures similar muiltilayers around reinforcements
form at relatively lower sintering temperature.
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