The mechanism of the failure of the dispersion-strengthened Cu–Al 2 O 3 nanosystem

Abstract

The method of “in situ tensile testing in SEM” is suitable for investigations of fracture mechanisms because it enables to
observe and document deformation processes directly, thank to which the initiation and development of plastic deformation
and fracture can be reliably described. The deformation and fracture mechanisms of Cu–Al2O3 nanomaterials with 5 vol.% of Al2O3 phase has been analyzed using technique of the “in situ tensile testing in SEM.” It has been shown that the deformation process
causes break-up of large Al2O3 particles and decohesion of smaller ones. The final fracture path is influenced also by boundaries of nanograins, through
which the principal crack propagates towards the sample exterior surface. Based on the experimental observations, a model
of damage and/or fracture mechanisms has been proposed.

Full text

The mechanism of the failure of the dispersion-strengthened
Cu–Al
2
O
3
nanosystem
Michal Besterci
•
Oksana Velgosova
´
•
Jozef Ivan
•
Tibor Kvac
ˇ
kaj
Received: 26 December 2009 / Accepted: 6 April 2010 / Published online: 22 April 2010
Ó Springer Science+Business Media, LLC 2010
Abstract The method of ‘‘in situ tensile testing in SEM’’
is suitable for investigations of fracture mechanisms
because it enables to observe and document deformation
processes directly, thank to which the initiation and
development of plastic deformation and fracture can be
reliably described. The deformation and fracture mecha-
nisms of Cu–Al
2
O
3
nanomaterials with 5 vol.% of Al
2
O
3
phase has been analyzed using technique of the ‘‘in situ
tensile testing in SEM.’’ It has been shown that the
deformation process causes break-up of large Al
2
O
3
par-
ticles and decohesion of smaller ones. The final fracture
path is influenced also by boundaries of nanograins,
through which the principal crack propagates towards the
sample exterior surface. Based on the experimental
observations, a model of damage and/or fracture mecha-
nisms has been proposed.
Introduction
We have analyzed, as reported in [1–3] the fracture of
Cu–Al
2
O
3
and Cu–TiC systems by direct monitoring of the
strain and fracture in a scanning electron microscope (SEM)
(in situ tensile test in SEM). Both systems were prepared by
different powder metallurgy technologies. The dispersed
oxides and carbides in the matrix were not coherent. Dif-
ferences in particle size and distribution caused differences
in fracture mechanism, although both factures were ductile
transcrystalline with dimples. Owing to the excellent high-
temperature properties and sufficiently high values of
electrical and thermal conductivity, the dispersion-
strengthened Cu–Al
2
O
3
materials, prepared by the methods
of powder metallurgy, have found use as conductors in
electrical machines employed at high temperatures, in
contacts, in electrodes and in vacuum technique parts.
The aim of work is to analyze fracture mechanism in the
Cu–Al
2
O
3
nanocomposite system and to propose a damage
model. The method of in situ tensile test in SEM was used
for other materials, too [4–10].
Experimental materials and methods
Reaction milling and mechanical alloying was used to
prepare the samples. Cu powder with the calculated addi-
tion of Al was homogenized by attrition in oxidizing
atmosphere. The distribution of the obtained CuO was
uniform. A subsequent treatment at 750 °C induced the
reaction of CuO with the added Al powder, and led to the
formation of Al
2
O
3
particles. The remaining CuO was
reduced by attrition in a mixture of H
2
? H
2
O (rate 1:100).
The powder was compacted by cold pressing and hot
extrusion at 750–800 °C.
M. Besterci
Institute of Materials Research, Slovak Academy of Sciences,
Watsonova 47, 043 53 Kos
ˇ
ice, Slovak Republic
e-mail: mbesterci@imr.saske.sk
O. Velgosova
´
(&)
Department of Non-ferrous Metals and Waste Treatment,
Faculty of Metallurgy, Technical University, Letna
´
9/A,
04200 Kos
ˇ
ice, Slovak Republic
e-mail: oksana.velgosova@tuke.sk
J. Ivan
Institute of Materials and Machine Mechanics, Slovak Academy
of Sciences, Rac
ˇ
ianska 75, 838 12 Bratislava, Slovak Republic
e-mail: ummsivan@savba.sk
T. Kvac
ˇ
kaj
Department of Metals Forming, Faculty of Metallurgy,
Technical University, Vysokos
ˇ
kolska
´
4, 04200 Kos
ˇ
ice,
Slovak Republic
e-mail: tibor.kvackaj@tuke.sk
123
J Mater Sci (2010) 45:4073–4077
DOI 10.1007/s10853-010-4493-5

Microstructured material with 5 vol.% Al
2
O
3
was
transformed by the equal channel angular pressing (ECAP)
method in two passes into a nanocomposite material. The
experimental material was pressed through two right
angled (90°) channels of a special die by route ‘‘C.’’ The
ECAP technology allows obtaining the very fine-grained
microstructure–nanostructure by multiple pressings
through the die.
This material with dimensions of [10 9 70 mm was
deformed by the ECAP technique in two passes at room
temperature in a hydraulic press in the equipment described
in [11].
For the purposes of investigation, very small flat tensile
test pieces, Fig. 1, with 0.15 mm thickness were prepared
by electroerosive machining, keeping the loading direction
identical to the direction of extrusion. They were ground
and polished mechanically to a thickness of approximately
0.1 mm. The final operation consisted in double-sided final
polishing of specimens with an ion thinning machine. The
test pieces were fitted into special deformation grips inside
the scanning electron microscope JEM 100 C, which
enables direct observation and measurement of the defor-
mation by ASID-4D equipment. From each one of the
systems five samples were prepared.
Results and discussion
The microstructure of the starting material with 5 vol.%
Al
2
O
3
was fine-grained (the mean matrix grain size was
1 lm) Fig. 2, heterogeneous, with Al
2
O
3
particles distrib-
uted in parallel rows as a consequence of extrusion and
ECAP. The experimental materials were deformed at 20 °C
at a strain rate of 6.6 9 10
-4
s
-1
in the elastic region. The
replica is shown in Fig. 3, and diffractogram is shown in
Fig. 4. Particles of size less than 0.25 lm were classified as
effective dispersion particles of category A, and particles of
size greater than 0.25 lm as particles of category B; the
latter are ineffective from the point of view of strength-
ening; although they affect the deformation process and
plastic properties. In addition to the particles mentioned,
Fig. 1 The shape and dimension of a specimen
Fig. 2 The microstructure of the starting material
Fig. 3 Replica of material with 5 vol.% Al
2
O
3
Fig. 4 Difractogram of Al
2
O
3
4074 J Mater Sci (2010) 45:4073–4077
123

the material also contained impurities, which were intro-
duced during the preparation process of the material.
The material after ECAP is on the border of nano-
structured materials. The TEM micrographs, Fig. 5,
showed that the mean grain size was 100–200 nm, small
amount of dislocations are present in nanograins, too, but
mostly on the boundaries. The nanostructure formation
takes place most probably by a three-stage mechanism,
described in [12–14]. This model has been experimentally
verified only for several specific materials but in our case it
seems to be probably usable. The model includes creation
of cell structure, then formation of transitory cell nano-
structure with large angle disorientation, and finally for-
mation of nanostructured grains with size of *100 nm.
However, here one has to consider the retarding effect due
to present dispersoid particles.
Deformation process of the loaded layer causes fracture
of large, B-type, particles in the middle of the specimen
(Fig. 6), which initializes fracture path roughly perpen-
dicular to the loading direction. The fracture path is
determined also by decohesion of smaller particles (type A)
(Fig. 7). Since the volume fractions of Al
2
O
3
particles are
small, their distribution in lines does not influence the
trajectory of fracture, which has low relative deformation
e = 0.1. Unlike the microstructured Cu-based composites,
in this case, it has been shown that the nanograin bound-
aries play an important role. In the final phase (Fig. 8), a
crack propagates along the nanograin boundaries, which
has been observed experimentally on the crack line
Fig. 5 The TEM micrographs of the material after ECAP
Fig. 6 Fracture of large (B) Al
2
O
3
particles in the middle of the layer
Fig. 7 Decohesion of smaller particles category A
Fig. 8 The final phase of the crack
J Mater Sci (2010) 45:4073–4077 4075
123

(profile), and it is documented also by the ductile fracture
surface with typical dimples in Fig. 9.
A detailed study of the deformation changes showed that
the crack initiation was caused by decohesion, and occa-
sionally also by rupture of the large particles. Decohesion is a
result of different physical properties of different phases of
the system. The Cu matrix has significantly higher thermal
expansion coefficient and lower elastic modulus (a =
17.0 9 10
-6
K
-1
, E = 129.8 GPa) than Al
2
O
3
(a = 8.3 9
10
-6
K
-1
and E = 393 GPa). Large differences in the
thermal expansion coefficients result in high stress gradients,
which arise on the interphase boundaries during the hot
extrusion. Since a
matrix
[ a
particle
, high compressive stresses
can be expected. However, because the stress gradients arise
due to the temperature changes, during cooling (which
results in increase of the stress peaks) their partial relaxation
can occur. Superposition of the external load and the internal
stresses can initiate cracking at interphase boundaries. This is
in accordance also with the dislocation theories, which argue
that the particles in composite may cause an increase in the
dislocation density as a result of thermal strain mismatch
between the ceramic particles and the matrix during prepa-
ration and/or thermal treatment. In our case, the coefficient of
thermal expansion of the matrix is much higher than that of
the secondary particles and the resulting thermal tension may
relax around the matrix–particle interfaces by emitting dis-
locations, whose density can be calculated according to
procedure described in [15].
Based on the microstructure changes observed in the
process of deformation, the following model of fracture
mechanism is proposed (Fig. 10a–d).
(a) The microstructure in the initial state is character-
ized by Al
2
O
3
particles, categorized as B.
(b) With increasing tensile load, local cracks, predom-
inantly on specimen side surfaces, are formed by
decohesion of smaller A particles.
(c,d) In further increasing deformation of nanocomposite
materials, the nanograin boundaries start to play an
important role. Since the volume fraction of these
boundaries is high and the size of the B and A
particles is equal to the matrix grain size, crack
propagates preferentially along the nanograin
boundaries in a 45° angle.
Fig. 10 a–d Model of the
fracture mechanism
Fig. 9 Surface morphology (dimples) of the material with 5 vol.%
Al
2
O
3
4076 J Mater Sci (2010) 45:4073–4077
123

Conclusion
The aim of the study was to evaluate the influence of
volume fraction of Al
2
O
3
(5 vol.%) particles on the frac-
ture mechanism by means of the method ‘‘in situ tensile
test in SEM.’’
Based on the microstructure changes obtained in the
process of deformation of the dispersion strengthened
Cu–Al
2
O
3
alloys, a model of fracture mechanism was
proposed. With increasing tensile load, the local cracks,
predominantly on specimen’s side surfaces, are formed
by rupture of large B and decohesion of smaller A
particles. The orientation of cracks tends to be perpen-
dicular to the loading direction, depending on the par-
ticle volume fraction. The final rupture, i.e.,
interconnection of the side cracks along the loading
direction, takes place at nanograin boundaries, depending
on the volume fractions of oxide (Al
2
O
3
) particles in a
45° angle.
Acknowledgement The work was supported by the Slovak National
Grant Agency under the Project VEGA 2/0105/08.
References
1. Besterci M, Ivan J (1997) Kovove Mater 35:278
2. Besterci M, Ivan J, Kovac L, Weissgaerber T, Sauer C (1998)
Kovove Mater 36:239
3. Besterci M, Ivan J, Kovac L (2000) Mater Lett 46:181
4. Besterci M, Ivan J (1996) J Mater Sci Lett 15:2071
5. Besterci M, Ivan J (1998) J Mater Sci Lett 17:773
6. Besterci M, Ivan J, Kovac L (2000) Kovove Mater 38:21
7. Besterci M, Ivan J, Velgosova
´
O, Pes
ˇ
ek L (2001) Kovove Mater
39:361
8. Besterci M, Velgosova
´
O, Ivan J, Hvizdos
ˇ
P, Kohu
´
tek I (2008)
Kovove Mater 46:139
9. Mocelin A, Fougerest F, Gobin PF (1993) J Mater Sci 28:4855.
doi:10.1007/BF00361147
10. Velı
´
sek R, Ivan J (1994) Kovove Mater 32:531
11. Besterci M, Velgosova
´
O, Ivan J, Hvizdos
ˇ
P, Kvac
ˇ
kaj T (2009)
Kovove Mater 47:221
12. Valiev RZ, Alexandrov IV (2000) Nanostrukturnyje materialy
poluc
ˇ
ennyje intensivnoj plastic
ˇ
eskoj deformaciej. Logos, Moskva
13. Valiev RZ (2003) In: Senkov ON, Miracle DB, Firstov SA (eds)
Proceedings of metallic materials with high structural efficiency.
NATO Science Series, Kijev. IOS Press, Amsterdam and Kluwer
Acad. Publishers, Dordrecht, p 79
14. Valiev RZ (1995) Nanostruct Mater 6:73
15. Luka
´
c
ˇ
P, Trojanova
´
Z (2006) Kovove Mater 44:243
J Mater Sci (2010) 45:4073–4077 4077
123