Preparation by mechanical alloying, characterization and sintering of Cu–20 wt.% Al2O3nanocomposites M.F. Zawraha,⇑, Hamdia A. Zayedb, Raghieba A…. [600942]

Technical Report
Preparation by mechanical alloying, characterization and sintering of Cu–20 wt.%
Al2O3nanocomposites
M.F. Zawraha,⇑, Hamdia A. Zayedb, Raghieba A. Essawyc, Amira H. Nassarc, Mohammed A. Tahac
aCeramics Department, National Research Center, Dokki, Cairo, Egypt
bPhysics Department, Faculty of Girls, Ain-Shams University, Egypt
cSolid State Physics Department, National Research Center, Dokki, Cairo, Egypt
article info
Article history:
Received 19 May 2012Accepted 22 October 2012
Available online 10 November 2012abstract
Metal-matrix nanocomposite, composed of copper/20 wt.% Al 2O3, was fabricated by mechanical alloying
method. The starting powders mixture was milled in planetary ball mill up to 20 h. The effect of milling
time on the properties of the obtained powders was studied. X-ray diffraction analysis (XRD) and trans-
mission electron microscopy (TEM) were used to investigate phase composition, crystal size and mor-phology of the milled powders. To study the sinterability, the milled nanocomposite powders were
cold pressed and sintered in argon atmosphere at different firing temperatures, i.e. 700 /C176, 800 /C176and
850/C176C, for 1 h. Physical properties, namely, bulk density and apparent porosity of sintered bodies were
determined by Archimedes method. Phase identification and microstructure of the sintered compositeswere investigated by using scanning electron microscope (SEM) as well as energy dispersive spectrome-
ter (EDS). Microhardness of sintered composite was also examined using Vickers hardness. The results
were discussed in terms of the effect of milling time on the properties of the prepared powders and sin-tered bodies. The results revealed that the grain size of milled powders was about 55 nm with a notice-
able presence of agglomerates. Uniform distribution of nano-sized alumina particles in the copper matrix
could be achieved with increasing milling time. The density of the sintered composites was affected bymilling time of the starting powders and firing temperature. It increased with increasing milling time
and firing temperature. Microhardness of the sintered bodies was found to be progressively increased
with increasing of milling time of starting powders.
/C2112012 Elsevier Ltd. All rights reserved.
1. Introduction
The fabrication of composite materials is a rational strategy to
design materials with properties that cannot be obtained for a
monolithic material. Metal–ceramic composites have received
extensive attention because they allow different combinations of
properties. Recently, the fracture toughness of ceramics has been
improved significantly by incorporating ductile metal phase
[1,2] ; while metal-matrix composites reinforced by ceramic partic-
ulates exhibit high specific strength and modulus as well as good
wear resistance compared to monolithic alloys [3,4] .
Mechanical alloying is a unique process in which a solid state
reaction takes place between fresh powder surfaces of the reactant
materials at room temperature [5]. Consequently, it can be used to
produce alloys and compounds which are difficult or impossible to
be obtained by the conventional melting and casting techniques
[6,7] . Mechanical alloying is a promising way for producing nano-
structured composites. Moreover, homogeneous distribution of
fine reinforcing particles and work hardening can be obtained bymechanical alloying. The repeated welding, fracturing and re-
welding of powder particles can result in intimate mixing of the
constituent powder particles on an atomic scale and lead to forma-
tion of stable and metastable supersaturated solid solutions, crys-
talline and quasi-crystalline intermediate phases, as well as
amorphous phases [8,9] . The most important advantage of this
method with respect to other alloying methods is the addition fea-
sibility of alloying elements to improve mechanical and physical
properties of alloys. Since mechanical alloying is a kind of high en-
ergy rate milling, thus all effective milling parameters can affect on
the process [10].
Nano-alumina particle-reinforced copper has received much
attention due to its outstanding high temperature properties com-
pared to pure copper. Nano-alumina particles dispersed in copper
particles improve copper matrix strength through impeding dislo-
cation movement. Moreover, it also can keep good electric conduc-
tivity and resistance to softening even at temperatures
approaching the melting point of copper. Nowadays, nano-alumina
particle-reinforced copper has been widely applied as spot welding
electrode, oxygen lance nozzle, contact material, etc. [11].
The goal of the current work is fabricating by mechanical alloy-
ing and sintering of Cu–20 wt.% Al 2O3nanocomposite. The influ-
0261-3069/$ – see front matter /C2112012 Elsevier Ltd. All rights reserved.
http://dx.doi.org/10.1016/j.matdes.2012.10.032⇑Corresponding author. Fax: +20 233370931.
E-mail addresses: mzawrah@stdf.org.eg ,mzawrah@hotmail.com (M.F. Zawrah).Materials and Design 46 (2013) 485–490
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ence of milling time on the properties of the prepared powders as
well as on the density, microstructure and microhardness of the
sintered composites was studied.
2. Materials and experimental procedures
Mixture of Cu and Al 2O3powders having 99.9 and 98.2% purity
andP10 and 1.4 lm average particle sizes, respectively, was used
as starting material to prepare Cu–20 wt.% Al 2O3nanocomposite.
Stearic acid was also used as process controlling agent to prevent
the agglomeration of the powder mixture during milling.
The powders mixture was milled by planetary ball mill up to
20 h using Al 2O3and ZrO 2balls having different diameters (6–
20 mm), at rotating speed 500 rpm and a ball-to-powder weight
ratio of 10:1.
The phase compositions of milled powders were investigated by
X-ray diffraction analysis (XRD) using a ‘‘Philips PW 1373’’ X-ray
diffractometer with Cu K-Ni filtered radiation and scanning rate
2 deg/min. XRD patterns were recorded in 2 hrange between 24
and 95.
The lattice parameter ‘‘ a’’ of the obtained phases was calculated
for the principle planes ( hkl), i.e. 111, 200, 220 and 311, from the
obtained XRD data according to the following equation [12]:
1
d2¼h2țk2țl2
a2ð1Ț
XRD data was also used to determine the crystallite size ( D) and
lattice strain ( e). The crystallite size was determined from the
broadening ( B) of the diffraction lines 111, 200, 220 and 311
using Scherrer equation as following [13–17] :
D¼0:9k
Bcoshð2Ț
The lattice strain ( e) was also calculated for the same diffraction
lines from the following equation [14–17] :
e¼B=4tan h ð3Ț
where k(wave length) = 1.54059 /C23A (Cu–Ni radiation), Bis the full
width at half maximum, his the angle in radians, and dis d-spacing
between the planes.
A transmission electron microscope (TEM), type JEOL JEM-1230,
operating at 120 kV and attached to a CCD camera, was employed
to investigate the morphology and particle size of the milled pow-
ders after different milling times.
The composite powders were uni-axially pressed into small
compacts of desired sizes at room temperature. Samples of
10 mm diameter and 4 mm height were pressed in a hardenedsteel die at 10 MPa using hydraulic pressing machine (SEIDNER,
Riedlinger, Germany, having a maximum loading capacity of
600KN).The compacted samples were sintered at different temper-
atures, i.e. 700 /C176, 800 /C176and 850 /C176C, in argon atmosphere for 1 h with
heating rate 8 /C176C/min.
The relative density and apparent porosity of sintered bodies
were measured by Archimedes method according to ASTM:
B962-08. Theoretical density of compacts was calculated using
the simple rule of mixtures, considering the fully dense values of
copper and Al 2O3are 8.96 and 3.95 g/cm3, respectively. Micro-
structure of sintered samples was examined by scanning electron
microscope (SEM), type ‘‘Philips XL30’’. The reinforcement of cop-
per matrix composites was estimated by microhardness measure-
ments (Vickers hardness) according to ASTM: B933-09. 1.961 N
Fig. 1. XRD of Cu–20 wt.% Al 2O3powder composites produced after different
milling times.0.40.50.60.70.80.91
0 5 10 15 20 25
Millin g time , hFull width at half-maximum , 2 θ
111311
200
220
Fig. 2. Full width at half maximum (FWHM) of Cu–20 wt.% Al 2O3versus milling
time.
00.20.40.60.811.21.4
0 5 10 15 20 25Millim g time , hLattice parameter , A°111
200
220
311
Fig. 3. Lattice parameter vs. milling time of Cu–20 wt.% Al 2O3powder mixtures.
1313.51414.51515.51616.51717.5
0 5 10 15 20 25
Milling time, hCrystal size , nm
0.40.420.440.460.480.50.520.54Lattice strain , %
Fig. 4. Crystal size and lattice strain of Cu–20 wt.% Al 2O3at different milling time.486 M.F. Zawrah et al. / Materials and Design 46 (2013) 485–490

load for 10 s. was applied during measuring hardness. The hard-
ness values of the investigated materials were measured as aver-
age of five readings along the cross section surface of the
specimens.
3. Results and discussion
3.1. Characterization of the prepared powders
3.1.1. Phase composition of the prepared powders
Fig 1 shows XRD patterns of Cu–20 wt.% Al 2O3composites pro-
duced after different milling times. Only two phases, i.e. Cu and
Al2O3were detected in the patterns of milled powders, following
to the card numbers (85-1326 & 88-0826) [18,19] . With increasing
milling time, the diffraction peaks of Cu and Al 2O3became broader
and their intensities became weaker. Full width at half maximum
(FWHM) measured from X-ray diffraction patterns for Cu–
20 wt.% Al 2O3powders is shown in Fig. 2 . Lines broadening were
observed with increasing milling time. This is due to severe lattice
distortion and grain size refinement [20,21] . Consequently, re-
duced crystal size and an elevated strain energy stored inside par-
ticles could be obtained because of the severe plastic deformation
introduced during ball milling.
The relationship between lattice parameters and milling time is
shown in Fig. 3 . During milling of Cu–20 wt.% Al 2O3composites,lattice parameters remain constant with milling time. It is well
known that the oxidation of Cu and formation of solid solution be-
tween metals and ceramics causes changes in the lattice parame-
ters. In the current work, the lattice parameters remain constant;
this means that neither oxidation of Cu nor formation of solid solu-
tion between Cu and Al 2O3has been done.
Crystal size and lattice strain are important parameters for
milled powders, since they have a significant effect on both com-
Fig. 5. TEM photomicrographs of Cu–20 wt.% Al 2O3composites after different milling time, (a) 1 h, (b) 5 h, (c) 10 h, (d) 15 h, and (e) 20 h.0100200300400500600700800900
0 5 10 15 20 25Millin g Time , hParticle size , nm
Fig. 6. Particle size of Cu–20 wt.% Al 2O3composites at different milling time.M.F. Zawrah et al. / Materials and Design 46 (2013) 485–490 487

051015202530
05 10 15 20 25Milling time, hApparent porosity, %
707580859095100
05 10 15 20 25
Millin g time, hRelative density , %a b
850 ș C
850 ș C800 ș C
800 ș C700 ș C700 ș C
Fig. 7. Effect of milling time and sintering temperature on (a) relative density and (b) apparent porosity.
Fig. 8. SEM photomicrographs of the sintered Cu–20 wt.% Al 2O3for 1 h at 850 /C176C prepared after different milling time. (a) 1 h, (b) 5 h, (c) 10 h, (d) 15 h, and (e) 20 h.488 M.F. Zawrah et al. / Materials and Design 46 (2013) 485–490

pacting of the powders during sintering process and properties of
the finely obtained copper matrix strengthened by fine dispersions.
The average crystal size of Cu matrix in the composites was esti-
mated using broadening of XRD peaks. The effect of milling time on
crystal size and lattice strain of examined powder particles was
presented in Fig. 4 . It is indicated that the crystal size ( D) decreased
with increasing milling time ( t) according to the equation:
D¼Kt/C02ð4Ț
where Kis a constant [22,23] .
In the present work, the most intensive crystal refinement oc-
curs in the early stage of milling up to 5 h. With prolonged time,
the crystal size of milled powders decreased slowly. The lattice
strain increased while crystal size reduced with increasing milling
time due to distortion effect caused by dislocation in the lattice
[22,24,25] . With increasing milling time, severe plastic deforma-
tion brings about a deformed lattice with high density of disloca-
tions [22,26] .
3.1.2. Morphology of the prepared powders
During high energy milling, the powder particles are repeatedly
deformed, cold welded and fractured by colliding balls. As a result,
composite powder particles with a characteristically layered
microstructure are formed during milling. The Al 2O3particles are
caught by the colliding copper particles and are mostly embedded
at interfaces between these particles [20]. The average particle size
increases until the welding process dominates the milling process,
and oppositely decreases after the fracture process becomes dom-
inant. In case of balance between fracture and welding processes,
the particles are rather uniform and equiaxed [27,28] .
Transmission electron microscopy (TEM) was used to gain bet-
ter insight into morphology and grain size of the composite pow-
der. Fig. 5 shows TEM images of milled composite powders for
different milling times. During high-energy milling, the powder
particles change their morphology and size as a consequence of re-
peated deformation, fracturing and welding processes. In the first
stage of milling, the ductile particles undergo deformation while
brittle particles undergo fragmentation. Then, when ductile parti-
cles start to weld, the brittle particles come between two or more
ductile particles at the instant of ball collision. As a result, frag-
mented reinforcement particles reside at the interfacial boundaries
of the welded metal particles, and the result is the formation of real
composite particles [14]. After 1 and 5 h milling time, the agglom-
erated alumina particles are extremely observed ( Fig. 5 a and b).
Such particle’s agglomeration could be removed by increasing mill-
ing time. After 10 h milling time, ( Fig. 5 c), alumina agglomerates
were observed, but they are smaller and less than that obtained
after 1 and 5 h milling. Also, a few amounts of alumina particles
have been welded with copper matrix. However, increasing millingtime caused uniform distribution of the reinforcement particles
and dissolution of agglomerated particles ( Fig. 5 d and e). Due to
the strong plastic deformation in the early stage of the milling pro-
cess which was followed by intensive welding, these particles were
not flat. Considering the balance between fracture and welding
processes, the particles are rather uniform in size and equiaxed
[29]. The particle size decreased with increasing milling time as
shown in Fig. 6 .
3.2. Sinterability of the prepared nanocomposites
3.2.1. Relative density and apparent porosity
The effect of milling time of composite powders on relative den-
sity and apparent porosity of the sintered samples at different sin-
tering temperature, i.e. 700 /C176, 800 /C176and 850 /C176C, for 1 h in argon
atmosphere, is shown in Fig. 7 . The theoretical density of Cu–20 wt.% Al
2O3composite was calculated and its value was
7.119 g/cm/C03. The relative density of the sintered samples in-
creased with increasing both milling time and sintering tempera-
ture. At high temperature range, where the diffusion mobility of
the atoms is sufficiently high, a complex diffusion mechanism of
mass transport responsible for the sintering process is occurred.
This means that with increasing sintering temperature, the actions
of complex diffusion mechanisms are more intense, directly affect-
ing the formation of surface contacts between the particles, forma-
tion of closed pores and grain growth. Therefore, the sintered
density has been increased with increasing sintering temperature
[30]. On the other hand apparent porosity was decreased with both
milling time of starting powder and sintering temperature.
3.2.2. Microstructure of the sintered composites
SEM images of sintered samples at 850 /C176C for 1 h in argon atmo-
sphere, prepared form starting material milled at different times,
are shown in Fig. 8 . Generally, homogeneous microstructure with
some alumina clusters (dark particles) is appeared in SEM photo-
micrographs. The amount of detected clusters decreased as we
go from the samples milled for shorter time into the samples
milled for longer time. However, increasing milling time caused a
uniform distribution of the reinforced particles, reduction in parti-
cles size (even to nanometer), disappearing of particle agglomera-
tion, reduction of distances between the composite particles and
consequently higher sinterability ( Fig. 8 e and f). EDX analysis of
the Cu–20 wt.% Al 2O3composite sintered at 850 /C176C for 1 h and pre-
Fig. 9. EDX spectra of sintered Cu–20 wt.% Al 2O3composite at 20 h milling time.
70075080085090095010001050110011501200
05 10 15 20 25
Milling time, hrMicrohardness , MPa
Fig. 10. Microhardness of sintered Cu–20 wt.% Al 2O3compacts at 850 /C176C versus
milling time of starting powder.M.F. Zawrah et al. / Materials and Design 46 (2013) 485–490 489

pared from milled powder for 20 h, is shown in Fig. 9 . The presence
of zirconium indicates contamination coming from milling balls.
3.2.3. Microhardness of sintered composite
The effect of milling time of starting powders on microhardness
of composites processed after sintering at 850 /C176C in argon for 1 h is
shown in Fig. 10 . The microhardness of the sintered composites in-
creased with increasing milling time of the starting powders
[17,20,27,31] . This increase in microhardness is a consequence of
fine Al 2O3particles dispersion in Cu matrix-compacts; and also
due to crystal size refinement of starting powders after milling
(Fig. 3 ) which led to increasing the sinterability and lowering the
porosity. The maximum value of microhardness of compacts pro-
cessed after 20 h-milling attains 1125 MPa.
4. Conclusion
The following remarks were concluded:
/C15Cu–20 wt.% Al 2O3nanocomposites have been fabricated using
mechanical alloying after different milling time up to 20 h, in
planetary ball mill. After milling, grain refinement was took
place and fine Al 2O3particles were regularly distributed in the
copper matrix.
/C15The crystal size has been decreased while the lattice strain was
found to be increased with increasing milling time due to dis-
tortion effect caused by dislocation in the lattice.
/C15The relative density of the sintered bodies was increased with
increasing both milling time and sintering temperature. On
the other hand, apparent porosity was decreased.
/C15Refinement of grains and dispersion of Al 2O3particles in the
composites have considerable effect on the hardness of the sin-
tered composites; the nano-sized Al 2O3particles have been act-
ing as a stronger strengthening parameter of the copper matrix.
The maximum hardness (1125 MPa) was for the sintered com-
pacts prepared after 20 h milling of its starting powder mixture.
References
[1] Chen Zhongchun, Takeda Takenobu, Ikeda Keisuke, Murakami Tadasu. The
influence of powder particle size on microstructural evolution of metal–
ceramic composites. Scripta Mater 2000;43:1103–9.
[2] Prielipp H, Knechtel M, Claussen N, Streiffer SK, Mullejans H, Ruhle M, et al.
Strength and fracture toughness of aluminium/alumina composites withinterpenetrating. Mater Sci Eng 1995;A197:19.
[3] Garcia DE, Schicker S, Janssen R, Claussen N. Nb- and Cr–A1
2O3composites
with interpenetrating networks. J Eur Ceram Soc 1998;18:601.
[4] Harrigan Jr WC. In: Everett RK, Arsenault RJ, editors. Metal Matrix
Composites. New York: Academic Press; 1991. p. 1.
[5] Zebarjad SM, Sajjadi SA. Microstructure evaluation of Al–A1 2O3composite
produced by mechanical alloying method. Mater Des 2006;27:684–8.[6] Davis RM, Mcdermott B, Koch CC. Formation of amorphous Fe–B alloys by
mechanical alloying. Metall Trans A 1988;19A:2867.
[7] Arik H. Production and characterization of in situ Al 4C3reinforced aluminum-
based. Mater Des 2004:25–31.
[8] Sheibani S, Heshmati-Manesh S, Ataie A. Influence of A1 2O3nanoparticles on
solubility extension of Cr in Cu by mechanical alloying. Acta Mater2010;58:6828–34.
[9] Suryanarayana C. Mechanical Alloying and Milling. New York: Marcel Dekker;
2004.
[10] Moshksar MM, Zebarjad SM. In: Proceedings of nonferrous metals, Iran,
Kerman; 1996. p. 913.
[11] Zheng Zhao, Xiao-jie Li, Gang Tao. Manufacturing nano-alumina particle-
reinforced copper alloy by explosive compaction. J Alloys Compd
2009;478:237–9.
[12] Sawaby A, Selim MS, Marzouk SY, Mostafa MA, Hosny A. Structure, optical and
electrochromic properties of NiO thin films. Physica B 2010;405:3412–20.
[13] Scherrer P. Bestimmung der Größe und der inneren Struktur von
Kolloidteilchen mittels Röntgenstrahlen. Nachrichten von der Gesellschaftder Wissenschaften zu Göttingen. Mathematisch-Physikalische Klasse
1918;2:98.
[14] Zawrah MF, Abdel-kader H, Elbaly NE. Fabrication of A1
2O3–20 vol.% Al
nanocomposite powders using high energy milling and their sinterability.
Mater Res Bull 2012;47:655–61.
[15] Sivasankarana S, Sivaprasadb K, Narayanasamya R, Satyanarayanac PV. X-ray
peak broadening analysis of AA 6061100- x-xwt.% A1 2O3nanocomposite
prepared by mechanical alloying. Mater Charact 2011;62:661–72.
[16] Danilchenko SN, Kukharenko OG, Moseke C, Protsenko IYU, Sukhodub LF,
Sulkio-Cleff B. Determination of the bone mineral crystallite size and latticestrain from diffraction line broadening. Cryst Res Technol
2002;37(11):1234–40.
[17] Klug HP, Alexander. Procedures for polycrystalline and amorphous. 2nd
ed. New York: John Wiley and Sons; 1974.
[18] Swanson HE, Tatge E. Natl. standard x-ray diffraction powder patterns. Bur.
Stand. (US), Circ. 539, 359, 1; 1953.
[19] Liu RS et al. Crystal structures and peculiar magnetic properties of alpha- and
gamma-(Al
2O3) powders. Mod Phys Lett B 1997;11:1169.
[20] Rajkovic Viseslava, Bozic Dusan, Jovanovic Milan T. Properties of copper matrix
reinforced with nano- and micro-sized A1 2O3particles. J Alloys Compd
2008;459:177–84.
[21] Lonnberg B. Characterization of milled Si 3N4powder using X-ray peak
broadening and surface area analysis. J Mater Sci 1994;29:3224–30.
[22] Razavi Tousi SS, Yazdani Rad R, Salahi E, Mobasherpour I, Razavi M. Production
of Al–20 wt.% A1 2O3composite powder using high energy milling. Powder
Technol 2009;192:346–51.
[23] Suryanarayana C. Mechanical alloying and milling. Prog Mater Sci
2001;46:1–184.
[24] Zhou F, Lee J, Lavernia EJ. Grain growth kinetics of a mechanically milled
nanocrystalline Al. Scripta Mater 2001;44:2013–7.
[25] Zhou Y, Li ZQ. Structural characterization of a mechanical alloyed Al–C
mixture. J Alloys Compd 2006;414:107–12.
[26] Nalwa HS. Nanocluster and nanocrystals. American Scientific Publishers; 2003.
[27] Maurice D, Courtney TH. Modeling of mechanical alloying: part III.
Applications of computational programs. Metall Mater Trans 1995;26A:2437.
[28] Rajkovic ´V, Eric ´O, Boz ˇic´D, Mitkov M, Romhanji E. Characterization of
dispersion strengthened copper with 3 wt%Al 2O3by mechanical alloying. Sci
Sintering 2004;36:205–11.
[29] Maurice D, Courtney TH. Applications of computational programs. Metall Met
Trans 1995;26A:2437–44.
[30] Korac ´Marija, Andic ´Zoran, Tasic ´Miloš, Kamberovic ´Zˇeljko. Sintering of Cu–
A12O3nano-composite powders produced by a thermochemical route. J Serb
Chem Soc 2007;72(11):1115–25.
[31] Rajkovic Viseslava, Bozic Dusan, Milan T, Jovanovic. Effects of copper and
A12O3particles on characteristics of Cu–A1 2O3composites. Mater Des
2010;31:1962–70.490 M.F. Zawrah et al. / Materials and Design 46 (2013) 485–490