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16
TH
INTERNATIONAL CONFERENCE ON COMPOSITE MATERIALS
1
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, by which the
initiation and development of plastic deformation
and fracture can be reliably described.
With increasing tensile load, local cracks are
formed by rupture of large particles and decohesion
of smaller particles. Further increase of load leads
to the crack growth by coalescence of cavities in the
direction from the surface to the specimen centre.
The cracks can be oriented parallel or
perpendicular to the loading direction in depending
on the particle volume fraction. The final rupture
takes place in variably dense rows, depending on the
volume fractions of carbide (Al
4
C
3
) and oxide
(Al
2
O
3
) particles.
1 Introduction
Dispersion strengthened materials belong to
the group of composite materials, which are made
mainly by the techniques of powder metallurgy.
Their structure is formed with a polycrystalline
matrix, in which dispersion particles are
incorporated, mainly of the oxide, carbide and
nitride types. The strengthening effect of dispersoids
is both direct, based on the braking of the matrix
dislocation movement, and indirect, resulting from a.
forming, dispersoids increase the density of
dislocations and refine the grain and subgrain
structure. The effectiveness of the strengthening
effect of dispersoids depends on their type, size,
morphology, volume fraction and distribution. Their
resistance against dissolution and coalescence is an
important factor in their strengthening effect, mainly
at high temperatures.
The polycrystalline matrix is formed from
both metallic and non-metallic powders. The most
important methods of powder production include
mechanical milling, reduction of oxides,
carbonization, water and gas spraying in air, in
a protective atmosphere or in vacuum, and also
electrolysis. The method of rotary electrode, rotary
disk and plasma spraying belong among the most
modern methods of powder production. The method
of evaporation and condensation is used for very fine
powders. The basic characteristics of powders
include their chemical composition, purity,
granulometric composition, shape, specific surface,
bulk density and density after jarring,
compressibility, sintering capacity and others.
Metallurgical purity, as well as changes taking place
during the treatment process is of key significance in
the production of dispersion-strengthened materials.
The oxides formed, foreign particles and gases,
which can be a cause of the origin of technological
defects of the microcrack type, separation of layers,
etc. plays the decisive role. It has been shown, from
theoretical considerations and experiments, that the
maximum effect of strengthening is achieved at the
following structural parameters:
a) the size of the strengthening particles of
secondary phases (dispersoids) should not exceed
50 nm. The particles of higher sizes are of little
effectiveness from the point of view of
strengthening.
b) the mean distance between strengthening
particles should be of the range from 0.1 to
0.5 µm, and their distribution should be uniform,
without any heterogeneities and clusters.
THE FRACTURE MECHANISM OF Al-Al
4
C
3
SYSTEM BY
“IN-SITU TENSILE TEST IN SEM”
Michal Besterci*, Oksana Velgosová**, Jozef Ivan***, Tibor Kvakaj****
*Institute of Materials Research, Slovak Academy of Sciences, Watsonova 47, 043 53 Košice,
Slovakia, **Dept. of Non-ferrous Metals and Waste Treatment, Faculty of Metallurgy,
Technical University, Letná 9/A, 04200 Košice, Slovakia, ***Institute of Materials and
Machine Mechanics, Slovak Academy of Sciences, Raianska 75, 838 12 Bratislava, Slovakia,
****Dept. of Metals Forming, Faculty of Metallurgy, Technical University, Vysokoškolská 4,
04200 Košice, Slovakia
Keywords:
composite material, “in-situ tensile testing in SEM”, fracture mechanisms
MICHAL BESTERCI, O.Velgosová, J.Ivan, T.Kvakaj
2
The real volume fraction of dispersoids
follows from the above-mentioned parameters. This
depends on the required properties of materials,
however, does not exceed, as a rule, 10 vol.%.
Mechanical alloying is the process of
production of macroscopically homogeneous
materials from heterogeneous mixtures. The process
is implemented in laboratories in attritors, and on the
industrial scale in high-energy ball mills. It is based
on deformation, repeated disintegration and welding
of powder particles during intensive dry milling.
Originally, this technology was developed for
production of Ni and Fe super alloys. It has since
been shown that mechanical alloying is able to
produce several stable and metastable phases
including solid solutions, metastable crystalline and
quasicrystalline phases, as well as amorphous
phases. It is remarkable that the crystalline phases
are usually of nanometric sizes. Because these
effects are similar to other ones, obtained by non-
equilibrium techniques, such as quick solidification
of the metal, the system is characterized by the
increased mechanical properties. The advantage of
preparing amorphous materials by mechanical
alloying in comparison with the technology of quick
solidification is represented by the fact that it makes
it possible to produce larger quantities of materials,
and extends the possibilities for alloying. The
technology of mechanical alloying is used for
metals, ceramic and polymers and considerable
attention is paid to it throughout the world. In the
literature, the term for mechanical alloying is
“mechanical alloying” if elementary powders are
used, and “mechanical milling”, if prealloyed
powder is used as the starting material. A series of
material systems has been developed using
mechanical alloying, mainly for high-temperature
applications on the basis of Ag, Al, B, Ba, Bi, Co,
Cr, Cu, Fe, Ge, Hf, La, Li, Mg, Mn, Nb, Nd, Pb, Re,
Si, Sm, Sn, Ta, Ti, V, W, Y and Zr. The materials
with low specific weight on the basis of Al, Li and
Mg represent an important group.
The dispersion strengthened alloys Al-Al
4
C
3
manufactured by mechanical alloying using powder
metallurgy technology are promising structural
materials enabling significant weight cut for use first
of all in aircraft and automobile industry and also at
elevated temperatures.
In our previous works [1-7] following [8-10]
we used “in-situ tensile test in SEM” to analyze
deformation processes in various types of Cu and Al
based composites. In works [1, 7] were studied the
strain and fracture on Al-Al
4
C
3
system. The
influence of Al
2
O
3
vol.% in Cu-Al
2
O
3
system was
analyzed in works [2, 5 and 6]. Deformation process
of Cu-TiC system was analyzed in works [3 and 4].
In works [8-10] were by “in-situ tensile test in SEM”
studied Al-Si-Fe and Al-Si systems. The result was a
design of several models of damage, which
considered physical parameters of matrix and
particles, as well as geometry and distribution of
secondary phases.
The aim of the present study is to evaluate the
influence of volume fraction of Al
4
C
3
particles
(8 and 12 vol. %) on the fracture mechanism.
2 Experimental materials and methods
The experimental materials were prepared by
mechanical alloying. Al powder of powder particle
size of <50 µm was dry milled in an attritor for 90
min with the addition of graphite KS 2,5 thus
creating 8 and 12 vol.% of Al
4
C
3
, respectively. The
specimens were than cold pressed using a load of
600 MPa the specimens had cylindrical shape.
Subsequent heat treatment at 550°C for 3 h induced
chemical reaction 4Al+3CAl
4
C
3
. The cylinders
were then hot extruded at 600°C with 94% reduction
of the cross section. Due to a high affinity Al to O
2
system also contains a small amount of Al
2
O
3
particles. The volume fraction of Al
2
O
3
phase was
low, 1-2 vol.%. Detailed technology preparation is
described in [11-15].
Experiments were considered also with
material Al-Al
4
C
3
with nano-matrix. It was tested by
SPD using ECAP (equal channel angular pressing).
The tested material was compact only after one pass
at the angle 90°, however at the next passes cracks
appeared, so we had to finish the testing. In (Fig.1) is
the dependence of mechanical properties on volume
fraction of Al
4
C
3
is shown the R
m
and A
5
values, too.
The mean grain size decreased for 1 µm to 0.6 µm.
For the purposes of investigation very small
flat tensile test pieces (7x3 mm) with 0.15 mm
thickness were prepared by electroerosive
machining, keeping the loading direction identical to
the direction of extrusion. The specimens were
ground and polished down to a thickness of
approximately 0.1 mm. Finally, the specimens were
finely polished on both sides by ion gunning. The
test pieces were fitted into special deformation grips
in the scanning electron microscope JEM 100 C,
which enables direct observation and measurement
of the deformation by ASID-4D equipment. From
each one of system (8 and 12 vol.% of Al
4
C
3
) was
prepared five samples.
THE FRACTURE MECHANISM OF Al-Al
4
C
3
SYSTEM BY “in-situ tensile test in SEM”
3
0
100
200
300
400
500
600
0 4 8 12 16 20
vol.% Al4C3
tensile strength [MPa] a
2
4
6
8
10
12
14
16
elongation [%]
Rm
A5 ECAP
Rm ECAP
A5
Fig. 1. The dependence of tensile strength and
ductility on Al-Al
4
C
3
system.
3 Results and discussion
The microstructures of the materials with 8
and 12 vol.% Al
4
C
3
were fine-grained (the mean
matrix grain size was 0.35 µm), heterogeneous, with
Al
4
C
3
particles distributed in parallel rows in
consequence of extrusion. The average distance
between the Al
4
C
3
particles, found in thin foils, was
1.1 µm.
Fig. 2. Al
4
C
3
particles identified by TEM.
When describing microstructures, one has to
consider geometrical and morphological factors.
According to the microstructure observations, the
particles in our materials can be divided into three
distinctive groups: A – small Al
4
C
3
particles,
identified by TEM, (Fig. 2), with mean size
approximately 30 nm which made up to 70% of the
dispersoid volume fraction; B – large Al
4
C
3
particles
with mean size between 0.4 and 2 m, found on
metallographic micrographs; and C – large Al
2
O
3
particles with mean size of 1 m. By morphology
Al
4
C
3
particles are elongated and Al
2
O
3
particles are
spherical. Let us assume that particles of all
categories during the high plastic deformation are
distributed in rows. Mean distance between the rows
is l and between the particles h. The particles are
spherical or have only a low aspect ratio, so that they
can be approximated as spherical. The experimental
materials were deformed at 20°C at a rate of
6.6x10
-4
s
-1
in the elastic region. In the material with
lower volume fraction (8 vol.%) of Al
4
C
3
with
increase of the deformation load the initiation of
microcracks on the large Al
4
C
3
particles (B) was
observed to occur by their rupture simultaneously
with decohesion of the smaller Al
4
C
3
and Al
2
O
3
particles (C and B - Fig.3).
Fig. 3. Fracture path in the material with 8 vol.%
Al
4
C
3
. Rupture of a large Al
4
C
3
grain and
decohesion of the smaller particles (ε=0.12).
50
µ
m 50
µ
m 50
µ
m
a)
b)
c)
Fig. 4. Propagation of the fracture toward the specimen interior: a) elongation 0.12 mm; b) elongation
0.18 mm; c) elongation 0.185 mm.
Al
4
C
3
660 nm
660
µ
m
MICHAL BESTERCI, O.Velgosová, J.Ivan, T.Kvakaj
4
The fracture may be initiated on the surface of
a specimen where large particles undergoing the
damage are located. Cases of crack initiation by
decohesion of large particles from the matrix and
propagation of cracks towards the interior of the
specimen was also observed, Fig.4a, b, c. The crack
then propagated from the surface into the bulk of the
specimen. On further deformation, as a result of
higher concentration of smaller Al
4
C
3
particles (A),
the perpendicular fracture trajectory partially
deviated toward the load direction (Fig.5) and
became irregular.
Fig. 5. Irregular fracture formed by a crack growing
alternatively along the particle rows and between
them in the material with 8 vol. % Al
4
C
3
(ε=0.15).
Fig. 6. Two cracks initiated on the opposite sides of
a specimen. Surface morphology and initiation of
cavities in the matrix–particle interphase in the
material with 12 vol% Al
4
C
3
(ε=0.04).
In the case of the higher volume fraction
(12 vol. %) of Al
4
C
3
the deformation process was
very rapid due to the low plasticity of the material.
Cracks were initiated on the surface and propagated
approximately perpendicularly to the tensile load
direction (Fig.6). Coalescence of the final fracture
progressed along densely populated rows of Al
4
C
3
(A, B) particles parallel to the load direction (Fig.7).
The morphology and size of the deformed surface
and three categories of particles on fracture surface
can be seen in Fig.8.
Fig. 7. Final fracture by interconnecting the two
opposite side cracks in the material with 12 vol%
Al
4
C
3
(ε=0.05).
Fig. 8. Surface morphology of the material with 12
vol% Al
4
C
3
(ε=0.05).
A detailed study of the deformation changes
showed that the crack initiation was caused by
decohesion, and occasionally also by rupture of the
large particles. Decohesion is a result of different
physical properties of different phases of the system.
The Al matrix has significantly higher thermal
expansion coefficient and lower elastic modulus
(from 23.5 to 26.5x10
-6
K
-1
, and 70 GPa) than both
Al
4
C
3
(5x10
-6
K
-1
, and 445 GPa) and Al
2
O
3
(8.3x10
-6
K
-1
, and 393 GPa), respectively. Large
differences in the thermal expansion coefficients
result in high stress gradients, which arise on the
interphase boundaries during the hot extrusion.
Since α
matrix
> α
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
50
µ
m
THE FRACTURE MECHANISM OF Al-Al
4
C
3
SYSTEM BY “in-situ tensile test in SEM”
5
external load and the internal stresses can initiate
cracking at interphase boundaries.
In view of the dislocation theories 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 preparation and/or thermal treatment.
The difference between the coefficients of thermal
expansion of the particles and the matrix may create
the thermal residual stresses after cooling from the
processing temperature to room temperature. The
coefficient of thermal expansion of the matrix is
much higher than that of the secondary particles.
The thermal tension may relax around the matrix-
particle interface by emitting dislocations. An
increase in the dislocation density reinforcement has
been calculated as [16].
( )
tfb
TBf 1
1⋅
−
∆
∆
=∆
α
ρ
(1)
where ∆α is the difference of the coefficient
of thermal expansion between matrix and particles,
∆T is a temperature change, t is the minimum size of
reinforcement, f is the volume fraction of particles, b
is the magnitude of the Burgers vector of
dislocations and B is a geometrical constant
(depending on the aspect ratio). The newly formed
dislocations are obstacles for the motion of
dislocations in the matrix. Therefore a higher stress
for the moving dislocations is necessary in
comparison to materials without secondary particles.
From Eq. 1 is obvious that the density of the
newly created dislocations increases with an
increase in the volume fraction. Therefore the
number of obstacles for the dislocation motion
increases and the stress necessary for the motion of
dislocations increases too. It should be mentioned
that the effect of different types of secondary
particles (Al
4
C
3
, Al
2
O
3
) on the reinforcement and
damage is depended not only from the coefficients
of thermal expansion difference but also from the
properties of the matrix/particle interface.
The fractures of the studied materials started
at the side-rims of the deformed samples. When
compared to the material with lower volume fraction
of Al
4
C
3
, in the present system the development of
slip bands in the bulk was inhibited. This fact, and
the absence of long-range slip in the matrix, implies
that the fracture is not inclined to the applied load
but is perpendicular to it. This is caused by the high
volume fraction of the strengthening particles and by
their short distance. Considering the sample width
(0.1 mm), the crack grew at 45° with respect to the
sample surface. The fracture was transcrystalline,
ductile.
a)
b)
c)
d)
Fig. 9. Model of the fracture mechanism.
Based on the microstructure changes observed
in the process of deformation, the following model
(it is not general model but it is a consequent model
on our experiments) of fracture mechanism is
proposed (Fig.9):
a) The microstructure in the initial state is
characterized by Al
4
C
3
and Al
2
O
3
particles,
categorized as A, B and C, whose geometric
A
C
h
d
l
B
MICHAL BESTERCI, O.Velgosová, J.Ivan, T.Kvakaj
6
parameters (l, h and d) depends on their volume
fraction.
b) With increasing tensile load local cracks,
predominantly on specimen side surfaces, are
formed by rupture of large (B, C) and decohesion
of smaller (A) particles.
c) Further increase of load leads to the crack growth
by coalescence of cavities in the direction from
the surface to the specimen centre. The cracks
can be oriented parallel or perpendicular to the
loading direction in depending on the particle
volume fraction.
d) The final rupture, i.e. interconnection of the side
cracks along the loading direction, takes place in
variably dense rows, depending on the volume
fractions of carbide (Al
4
C
3
) and oxide (Al
2
O
3
)
particles.
3 Conclusions
The aim of the study was evaluation of
volume fraction of Al
4
C
3
(8 and 12 vol.%) and
Al
2
O
3
(1-2 vol.%) particles on the fracture
mechanism of the method “in situ tensile test in
SEM”.
Based on the microstructure changes obtained
in the process of deformation the dispersion
strengthened Al-Al
4
C
3
alloys was model of fracture
mechanism proposed. With increasing tensile load
local cracks, predominantly on specimen side
surfaces, are formed by rupture of large (B, C) and
decohesion of smaller (A) particles. Further increase
of load leads to the crack growth by coalescence of
cavities in the direction from the surface to the
specimen centre. The cracks can be oriented parallel
or perpendicular to the loading direction in
depending on the particle volume fraction. The final
rupture, i.e. interconnection of the side cracks along
the loading direction, takes place in variably dense
rows, depending on the volume fractions of carbide
(Al
4
C
3
) and oxide (Al
2
O
3
) particles.
Acknowledgement
The work was supported by the Slovak National
Grant Agency under the Project VEGA 2/5142/25
and Project APVV-20-027205.
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