Creep Resistance depending on Particle Reinforcement Size of Al-Alloys
produced by Powder Metallurgy
Bernd Bauer and Guillermo Requena
Institute of Materials Science and Technology
Vienna University of Technology
Marcela Lieblich
National Centre for Metallurgical Research, CENIM-CSIC
Spanish Research Council, Madrid
Abstract
The Young’s Modulus, the high temperature strength and the tensile creep behaviour are investigated
for unreinforced and particle reinforced AA2124 and AA6061 matrices with 25vol% of SiC particles
with different sizes. The materials were produced via powder metallurgy and subsequent extrusion.
The blending of the metallic powders and the SiC particles was carried out by wet blending (WB) and
ball milling (BM). The SiC particles size in WB samples seems less affected by the blending
operation, but the bigger particles are affected by the extrusion process. The BM process fractures the
SiC particles leading to sub-µm SiC particles and a significant amount of oxide dispersoids is
incorporated.
The Young’s Modulus is increased by about 50% compared to the unreinforced alloys due to the
introduction of the stiffer SiC particles regardless of the initial particle size. All BM samples and only
SiC particles <5µm in WB samples provide increased hot tensile strength. The stationary creep rate of
composites produced by BM is two orders of magnitude lower than that of the WB composites and the
unreinforced matrices. This enhancement of creep resistance of BM composites is achieved by the
presence of dispersoids as a result of BM. The creep resistance of the WB composites decreases the
larger the particle size acting as dislocation sources. A similar behaviour is observed for the high
temperature strength for which the dispersoids in the BM composites increase the strength, while the
larger SiC particles decrease the strength with respect to the matrices.
Introduction
Metal matrix composites (MMCs) are obtained when a continuous metallic matrix is reinforced by
means of a secondary stiffer phase, usually a ceramic [1]. The reinforcement of metals can have many
different objectives such as an increase in specific mechanical properties like creep resistance, tensile
strength, wear resistance and Young’s Modulus [1]. Accordingly, the requirements for the
reinforcement phases are low density, thermal stability, high compression strength, high tensile
strength and high Young’s Modulus [2].
Powder metallurgy (PM) techniques are solid phase processes that can be applied to produce MMCs.
These typically involve discontinuous reinforcement, due to the ease of mixing and blending. The
ceramic and metal powders are mixed, cold compacted and hot pressed. The compacted and pressed
material then typically undergoes a secondary operation such as extrusion. Powder processing is used
to fabricate primarily particle- or whisker-reinforced MMCs. The matrix and the reinforcement
powders are blended to produce a homogeneous distribution [2]. The blending can be carried out dry
or in liquid suspension. This is usually followed by cold compaction, canning, degassing and high
temperature consolidation. The relative size of metal and ceramic particles has been identified as
significant for homogeneous blending of powders to produce particulate reinforced metals (PRM) [3].
According to [4] many properties of powders– such as powder flow, apparent density, ejection stress,
delubrication behaviour as well as dimensional change and mechanical strength of extruded parts – are
quite sensitive to even small changes in particle size distribution and to fluctuations in the
concentrations of components within a powder mixture. In the present study the dependence of the
high temperature strength, the Young’s Modulus as well as the creep resistance of Al-based PRMs
produced by PM are investigated for different mixing routes and particle sizes. Regarding the creep
resistance controversial results were reported during the last decades. On the one hand, there should be
an increase of creep resistance of Al alloys due to the addition of SiC particles [5, 6], whereas it was
also reported that the SiC particles degrades the creep resistance [7, 8].
Experimental Procedures
The Materials
Two Al-based matrices reinforced with 25vol% of SiC particles of different sizes were produced by
PM. The matrices in the present study are AA2124 (AlCu4) and AA6061 (AlMg1SiCu) alloys. The
chemical compositions in Table 1 were determined by Inductively Coupled Plasma (ICP) - analysis
[9]. The matrix powder grains have sizes <75µm. Three different SiC particle sizes (denominated
<5µm, <10µm and <25µm) were used to reinforce the matrices. The shape of the utilized SiC particles
is irregular and with sharp edges. The initial SiC particle sizes can be seen in Table 2 [10]. Hereby, the
designations F360, F600 and F1000 correspond to the particle sizes <25µm, <10µm and <5µm,
respectively. The SiC powders were provided by ESK-GmbH [10]. Three values are presented to
characterize the powders: 1) the P97-value is the 97% percentile that means that 97% of the particle
number is below this value, 2) the median and 3) the P6-value which indicates that 6% of the particle
number is below this value.
The powders of the matrices and the SiC particulates were blended using two different routes: low
energy mixing was carried out within cyclohexane for a period of 24 hours and the liquid was
subsequently evaporated. Hereafter, this process is called wet blending (WB). During WB the powder
particles are neither deformed nor fractured. The blending time of 24 hours was chosen because of the
relatively high size ratio between the matrix powder particles and the SiC particles (3-15:1). The
smaller the ratio the easier the blending [11]. Taking into account the different SiC sizes and the fixed
blending time of 24 hours it was found empirically that the optimal amount of blended powder is 100g
and 200g in the given process. SiC particles with sizes <10µm were not used for the production of
PRM by WB.
High energy ball milling (BM) was the second blending technique. It was applied using a planetary
milling device. A period of 4 hours was necessary to obtain a homogeneous blending of powders as
determined by microscopic investigations of the blended powders after 2, 4, 6, 8 and 10 hours of BM.
The rotating velocity of the containers was 200 rpm. An amount of 60g of powder mixture of matrix
and SiC particles was blended in each of the containers. 12 steel balls with a diameter of 14.4mm and
31g weight were placed in each container to blend the alloy and the SiC powders. The powder
agglomerates are repeatedly deformed, fractured and cold welded during BM [4, 12].
Both types of blended powders were cold compacted within Al tubes and subsequently extruded. Non
lubricated forward extrusion [13] of the blended powders was conducted by means of a die with a
diameter of 11.8mm at 450°C. The inert gas fusion method was used to determine the oxygen of the
materials. The determination was carried out in the Institute of Chemical Technologies and Analytics
[14] using a LECO device. The principle of operation is based on fusion of a sample in a high-purity
graphite crucible at temperatures up to 3000°C in an inert gas such as helium. The oxygen in the
sample reacts with the carbon from the crucible to form carbon monoxide (CO) and/or carbon dioxide
(CO2), and these are detected using infrared (IR) detection. The results are given in Figure 1. The BM
matrix and PRM contain about 2-3 times more oxygen (i.e. oxides) than the unreinforced mixed
powder compacts and the WB PRM, respectively. The WB PRM contains about the same fraction of
oxygen as the unreinforced BM matrix. The oxygen content of the BM PRM with the small particles
as ingredients is about 40% higher than that of the BM PRM with the bigger ones.
Prior to testing all the materials were heat treated at 495°C for 20 minutes with subsequent water
quenching and overaged at 300°C for 1 hour in order to stabilize the precipitates by overaging. This
treatment is designated as T4S where S means stabilisation. Table 3 shows all the investigated
materials.
Si
[wt%]
Matrix Alloy
Fe
[wt%]
Cu
[wt%]
Mn
[wt%]
Mg
[wt%]
Cr
[wt%]
Ni, Zn,
Ti
3.9
0.27
0.6
≤ 0.01
1.45
0.96
≤ 0.01
0.16
≤ 0.01
-
PM 2124
0.03
0.08
PM 6061
0.45
0.15
Table 1 Chemical composition of the matrices
P6 percentiles
[µm] min.
F 360
12
F 600
3
F1000
1
Table 2 Initial SiC-particle sizes [10]
Grid designation
P97 percentiles
[µm] max.
40
19
10
Median [µm]
22.8 ± 1.5
9.3 ± 1.0
4.5 ± 0.8
Symbol
<25µm
<10µm
<5µm
0,7
0,6
Oxygen Fraction [wt%]
0,5
0,4
0,3
0,2
0,1
BM 2124/SiC<25µm/25p
BM 2124/SiC<5µm/25p
WB 2124/SiC<5µm/25p
2124 with BM
2124 without BM
0,0
Figure 1 Oxygen fraction of the unreinforced 2124 without and with BM, WB 2124/SiC<5µm/25p, BM
2124/SiC<5µm/25p, BM 2124/SiC<25µm/25p
Production route
Designation
0
SiC size [µm]
class
-
Extrusion
2124 Matrix without BM
0
-
BM + extrusion
BM 2124 Matrix
Matrix alloy
SiC [vol.%]
2124
2124
6061
0
<5
<10
2124
25
<25
<5
<25
<5
6061
25
<25
Table 3 List of processed and investigated materials
BM + extrusion
BM + extrusion
WB + extrusion
WB + extrusion
BM 6061 Matrix
BM 2124/SiC<5µm/25p
BM 2124/SiC<10µm/25p
BM 2124/SiC<25µm/25p
WB 2124/SiC<5µm/25p
WB 2124/SiC<25µm/25p
WB 6061/SiC<5µm/25p
WB 6061/SiC<25µm/25p
Light Optical Microscopy (LOM)
The samples were cut and embedded in order to study the microstructure in the longitudinal and
perpendicular directions with respect to extrusion. Images taken by LOM were analysed by means of
the software Axio Vision 4.7.1 in order to determine the distribution of the SiC particles and the
presence of porosity. Only the amount and distribution of SiC particles for the WB PRMs (>1µm)
were analysed by means of LOM because of the limited resolution. The existence of smaller ones can
be clearly observed but not analysed in terms of distribution and size.
Field Emission Gun – Scanning Electron Microscopy (FEG-SEM)
FEG-SEM was utilized to analyse the sub-µm SiC particles which could not be quantified by LOM as
well as the SiC particles >1µm for the BM PRMs. Images taken by FEG-SEM in CENIM, Madrid,
were used to analyse the SiC particles >1µm of the BM material. The applied device was a FEG-SEM
with detectors for secondary electrons as well as backscattered electrons achieving a resolution of 1.5
nm (15kV) or 5nm (1kV). Other FEG-SEM images were taken by means of FEI QUANTA 200 FEGSEM device provided by University Service for Transmission Elektron Microscopy (USTEM) at the
Vienna University of Technology [15] in order to investigate the SiC particles <1µm. The observed
particles were identified by energy dispersive microanalysis. Negligible differences were found for the
SiC particles <1µm between the longitudinal and the perpendicular directions. Therefore, only the
images taken in longitudinal direction were analysed.
EBSD measurements were carried out at the USTEM for all the materials in order to obtain their
initial grain size by means of FEI QUANTA 200 FEG.
Transmission Electron Microscopy (TEM)
The samples of the unreinforced matrices were prepared by electro polishing using a TenuPol-5
provided by Struers. The ion milling of the WB PRM was done by means of a Precision Ion Polishing
System (PIPS). The composition of the sub-µm particles was analysed by energy dispersive X-ray
spectroscopy (EDX) in a FEI TECNAI G20. TEM investigations as well as the samples preparation
were conducted by USTEM at the Vienna University of Technology [15].
Dynamical Mechanical Analysis (DMA)
The Young’s Modulus of all the investigated materials was measured as a function of the temperature
by means of DMA in order to evaluate the quality of the produced materials regarding porosity and the
processing routes. The investigations were conducted with a TA Instrument DMA 2980 equipment in
air atmosphere using a 3-point-bending clamp with a frequency of 1Hz, amplitude of 40µm and in a
temperature ranging from room temperature to 300°C. The thermal stability is given by ±0.1°C above
50°C and ±1°C below 50°C. The samples’ geometry was 55 x 4 x 2 mm³ (Length x Width x
Thickness). The accuracy of the results is within ±3%.
High temperature (HT) tensile tests
Hot tensile tests were carried out for all the investigated materials by means of a Zwick Z250 universal
test machine equipped with a heating chamber heated by electrical resistance. The tensile tests were
conducted in air atmosphere at a temperature of 300°C. The temperature was measured using 2
thermocouples, one on the upper end of the sample and another on the lower end of the sample.
The resulting temperature gradient was less than 7°C. The sample geometry is the same as for the
isothermal tensile creep tests.
Isothermal tensile creep tests
The tensile creep test rigs were loaded by weights on a lever system and equipped with a digital
temperature control system. The rigs allow to carry out simultaneously 10 tensile creep tests at
temperatures up to 800°C [16]. The creep tests were carried out at 300°C in air.
The temperature of the samples was measured using two thermocouples one at the bottom and another
at the top of the sample. The observed temperature gradients were smaller than 5°C. The strain during
the tests was measured using linear variable displacement transducers (LVDT) provided by MicroEpsilon [17] with a sensitivity of 56mV/V and a resolution of 0.01µm. The measured signals were
then amplified and recorded using a data acquisition system made up of a PC and external amplifiers
model Spider8 provided by Hottinger Baldwin Meßtechnik G.m.b.H. [18]. The materials were tested
under constant load conditions to determine the stationary creep rate. Once the stationary stage was
reached the load was increased in order to continue the test to the corresponding stationary stage. The
applied loads ranged from 15MPa to 70MPa resulting in test periods between 0.5h and 8000h.
Results
The microstructural differences between the BM and WB PRMs can be seen in Figure 2 and Figure 3.
The SiC particle size distribution in the WB materials (see Figure 2a) and b)) reflects the size class of
the ingredient powders. The SiC particles in the BM PRMs are fractured causing similar particle size
distributions for the different ingredient size classes and sub-µm SiC particles. This results in a
multimodal particle size distribution for the BM composites (see Figure 2 c) and d)). The SiC particles
of the WB PRMs maintain their sharp edges, whereas those of the BM PRMs are rounded. All the
images exhibit some elongated particle free zones which result from particle free matrix grains
deformed in the hot extrusion process. An alignment of the SiC particles in direction of extrusion is
evident in the materials mixed by WB, but hardly observed in the materials prepared by BM due to the
more spherical shape of the fractured particles.
The results of the quantitative analyses of the SiC particles distribution is shown in Table 4 for the
particles >1µm and in Table 5 for the particles <1µm. The 6 and 97 percentiles, the median as well as
the fraction of particles larger than 5µm in equivalent diameter (area >19.6µm²) for all the investigated
materials are given. These results reveal considerable differences between the WB PRMs with SiC
particles <5µm and the WB PRMs with particles <25µm. The WB PRMs <5µm contain about 4vol%
of particles bigger than 19.6µm² in cross section, while the WB PRMs <25µm more than 9vol%. The
amount of SiC particles >5µm is about 2vol% in the <10 and <25µm BM materials and only less than
1vol% in the <5µm BM PRMs. The median of the particle size is markedly reduced in the BM
composites with respect to WB. The median diameter of the SiC >1µm in BM decrease slightly with
increasing ingredient particle size, whereas the size of the SiC <1µm seems to be independent. No
sub-µm SiC particles were observed in the WB materials, whereas in the BM materials <1µm SiC
particles are present as shown in Figure 3a) as well as an increased amount of intermetallic phases and
oxides.
The EBSD measurements show that most of the grains of the unreinforced alloys (see Figure 3 b)) are
elongated in the extrusion direction. There are also chains of rounded grains of around 5µm in
diameter. The introduction of the SiC particles leads to a decrease in grain size down to a range of 35µm for the WB PRMs (Figure 3 c)) and even smaller due to the BM processing (1-5µm, see Figure 3
d)). Because of the very small grains in the BM material, it was not possible to analyse the grains
precisely, but the obtained images help to get an overview on the distribution of the grains.
a)
b)
SiC
SiC
20µm
c)
20µm
d)
Dispersoids
SiC
Dispersoids
SiC
20µm
20µm
Figure 2 LOM images: a) WB 2124/SiC<5µm/25p and b) WB 2124/SiC<25µm/25p. FEGSEM images: c) BM 2124/SiC<5µm/25p (BSE) and d) BM 2124/SiC<25µm/25p.
a)
b)
SiC
<1µm
50nm
c)
d)
Figure 3 a) TEM image of BM
2124/SiC<5µm/25p. EBSD images of b)
2124 without BM, c) BM PRM
2124/SiC<5µm/25p and d) WB PRM
2124/SiC<5µm/25p (SiC particles and not
clearly oriented regions are marked black)
Material
BM:
2124/SiC<5µm/25p
2124/SiC<10µm/25p
2124/SiC<25µm/25p
WB:
2124/SiC<5µm/25p
2124/SiC<25µm/25p
6061/SiC<5µm/25p
6061/SiC<25µm/25p
P6 percentiles [µm]
min.
Median [µm]
P97 percentiles [µm]
max.
Vol.fraction [%] of
particles in PRM
>5µm
1.1
1.1
1.1
2.1
1.9
1.7
5.0
8.0
11.0
0.6
2.0
2.0
1.8
1.2
1.3
1.3
3.2
3.4
3.2
3.9
7.0
14.0
7.0
11.0
4.0
9.7
3.8
9.1
Table 4 Quantitative analyses of the SiC particles >1µm
Material
BM:
2124/SiC<5µm/25p
2124/SiC<25µm/25p
WB: No SiC Particles <1µm detected
Median [µm]
P25-P75
percentiles [µm]
0.16
0.16
0.1-0.3
0.1-0.3
Table 5 Quantitative analyses of the SiC particles <1µm
Young’s Modulus
The 25vol% of SiC particles added to the 2124 and 6061 alloys cause an increase of the Young’s
Modulus by about 60% in the whole temperature range tested. The obtained curves of the Young’s
Modulus for a temperature range 25°C to 300°C are given in Figure 4. The results of the DMA tests of
the unreinforced 2124 matrices both with and without BM are very close. The unreinforced 6061
matrix shows a Young’s Modulus which is in good accordance with the unreinforced 2124 matrix at
lower temperatures. The higher the temperature the bigger the difference between 2124 and 6061
unreinforced matrices because of a faster decrease of the Young’s Modulus of 6061 at higher
temperatures.
For the BM PRMs with the different SiC particle sizes it can be observed that there are no differences
in the resulting temperature dependence of the Young’s Modulus. The curves are parallel to those of
the unreinforced matrix.
All in all, it can be concluded that the presence of SiC particles within the matrices increases the
Young’s Modulus. The Young’s Modulus of the 6061 PRM is independent of the initially introduced
SiC particle size. Regarding 2124 the different alloying routes and the different SiC particle sizes yield
little differences. The WB 2124 PRM with the smallest SiC particles yields the highest Young’s
Modulus, whereas the BM 2124 PRM with the biggest ingredient SiC particles represents the lower
limit. The difference in Young’s Modulus of the BM 2124 PRM is within the experimental scatter, but
the larger modulus of the WB 2124 PRM <5µm with respect to the other PRM seems significant. This
is possible to occur due to the shape factor of the SiC particles and a lack of porosity but will be
further investigated.
130000
120000
Young's Modulus [MPa]
110000
100000
90000
80000
70000
60000
50000
WB 2124/SiC<5µm/25p
WB 2124/SiC<25µm/25p
WB 6061/SiC<5µm/25p
WB 6061/SiC<25µm/25p
BM 6061 Matrix
BM 2124/SiC<5µm/25p
BM 2124/SiC<10µm/25p
BM 2124/SiC<25µm/25p
2124 Matrix without BM
BM 2124 Matrix
40000
30000
20000
0
20
40 60 80 100 120 140 160 180 200 220 240 260 280 300 320
Temperature [°C]
Figure 4 Curves of Young’s Modulus as a function of temperature of all the introduced materials
High temperature (HT) tensile strength
The results of the HT tensile tests conducted at 300° are given in Figure 5 and Table 6. The 2124
matrices have Rp0.2≈91±3MPa with elongation at fracture around 30% in the case of 2124 without BM
and reaching around 40% elongation at fracture in the case of 2124 subjected to BM. The unreinforced
2124 alloy shows HT strength about 60% higher than that of the unreinforced 6061 alloy. The
strengthening behaviour and the strain to failure of unreinforced and SiC particle reinforced Al alloys
were studied frequently during the last decades by means of hot tensile tests [19-23] which are in good
agreement with the obtained results.
The BM PRMs show the highest HT strength of the studied materials and it increases with smaller
initial SiC particle size combined with decreasing ductility. The same tendency is observed for the WB
PRMs. The WB PRMs <25µm show almost the same HT strength as the unreinforced matrix but with
a significant loss of ductility. This is observed for both the WB 2124 PRMs as well as for the WB
6061 PRMs. The ductility of all the investigated PRM is similar. Lower strength is observed for the
6061 based materials compared with those with 2124 matrix. The addition of SiC particles increases
the strength but decreases the strain to failure. A clear ranking in strength is obtained which can be
attributed to the particle size and thus to the interparticle distances. The BM PRMs contain as well
dispersoids.
140
Rp0,2 [N/mm²]
Rm [N/mm²]
120
Stress [N/mm²]
100
80
60
40
20
Figure 5 Results of the tensile tests at 300°C
6061 BM
WB 6061/SiC<25µm/25p
WB 6061/SiC<5µm/25p
2124 BM
2124 without BM
WB 2124/SiC<25µm/25p
WB 2124/SiC<5µm/25p
BM 2124/SiC<25µm/25p
BM 2124/SiC<10µm/25p
BM 2124/SiC<5µm/25p
0
Material
BM 2124/SiC<5µm/25p
BM 2124/SiC<10µm/25p
BM 2124/SiC<25µm/25p
WB 2124/SiC<5µm/25p
WB 2124/SiC<25µm/25p
2124 without BM
2124 BM
WB 6061/SiC<5µm/25p
WB 6061/SiC<25µm/25p
6061 BM
Rp0.2 [MPa]
123 ± 1
118 ± 1
113 ± 1
99 ± 2
89 ± 3
90 ± 2
92 ± 2
66 ± 1
56 ± 2
57 ± 1
Rm [MPa]
128 ± 1
122 ± 1
116 ± 1
109 ± 1
95 ± 2
91 ± 2
93 ± 2
70 ± 1
59 ± 1
58 ± 1
Rp0.2/Rm
0.96
0.97
0.97
0.91
0.94
0.99
0.99
0.94
0.95
0.98
A [%]
10 ± 1
12 ± 1
13 ± 1
10 ± 2
9±1
31 ± 1
39 ± 2
11 ± 1
12 ± 2
27 ± 1
Table 6 Results of HT tensile tests at 300°C of all the introduced materials in T4S condition
The creep resistance
The stationary creep rates are represented for all the investigated materials as a function of the applied
stress A shown in Figure 6 in a double logarithmic scale. The WB PRMs with SiC particles <5µm do
not show any reinforcing effect neither for the 2124 nor the 6061 matrix. The minimum creep rate of
these composites is in the range of their corresponding unreinforced matrices. The 2124 alloy
subjected to 4 hours BM results in a slightly better creep resistance than the same alloy without BM.
There is a detrimental effect of the <25µm SiC particles for the WB PRMs yielding a higher stationary
creep rate in the whole stress range. This shows that the creep resistance decreases with increasing
particle size for the WB PRMs.
The creep resistance is markedly increased for the PRMs prepared by BM. Here, it can be seen that the
BM PRMs reveal a decrease of the stationary creep rate of about 1-2 orders of magnitude for loads
between 30 to 35MPa and of about 3-4 orders of magnitude for loads ≥ 35MPa. No significant
differences in the creep resistance were found regarding the initial particle sizes for these materials.
Assuming power laws, the creep exponent for the unreinforced 2124 matrix with and without BM
n1=2 for tensile loads ranging from 15 to 30MPa increases sharply to n2=16 for higher applied tensile
loads. Very similar values of the creep exponents n1 and n2 were obtained for the BM PRMs. Here it
was found that the values for n1 result in about 2 for tensile loads ranging from 30 to 50MPa and n2
about 15 for tensile loads from 50 to 70MPa. As the corresponding unreinforced matrices give similar
values for the creep exponents n1 and n2, it can be concluded that the same mechanisms are responsible
for creep of these materials.
The WB PRMs with the <5µm SiC particles also show two power-law regions. The low stress levels
(15-30MPa) yield a creep exponent n1 of about 5, whereas in the high stress level (30-60MPa) again a
sharp increase of n2 up to values of about 15 is obtained. This is not the case for the WB PRMs <
25µm, which results with a value for n of about 8-10 in the entire range of loads tested.
0.01
2124 Matrix without BM
WB 2124/SiC<25µm/25p
WB 2124/SiC<5µm/25p
WB 6061/SiC<25µm/25p
WB 6061/SiC<5µm/25p
BM 6061 Matrix
BM 2124 Matrix
BM 2124/SiC<5µm/25p
BM 2124/SiC<10µm/25p
BM 2124/SiC<25µm/25p
-1
[s ]
1E-3
1E-4
stat
Stationary Creep Rate 
.
1E-5
1E-6
n1=2.0 / n2=14.8
n=8.3
n1=4.7 / n2=11.5
n=11.2
n1=5.6 / n2=15.5
n1=2.9 / n2=15.6
n1=2.0 / n2=19
n1=2.0 / n2=15
n1=2.0 / n2=15
n1=2.0 / n2=15
1E-7
1E-8
1E-9
1E-10
1E-11
1E-12
10
20
30
40
Applied Stress [MPa]
50
60
70 80 90 100
A
Figure 6 Summary of all the minimum creep rates of the unreinforced matrices and the WB PRMs
Discussion
Due to the addition of SiC particles the Young’s Modulus increases by about 50% for the whole
temperature range independently of the initially introduced SiC particle size. Estimations of the
Young’s Modulus by the inverse rule of mixture (ROM) results as follows:
Reuss: Ec 
1
=94GPa
vp 1 vp

Ep
EM
where Ec, Ep and EM are the Young’s Modulus of the composite, the SiC particles (400GPa) and the
matrix (75GPa), respectively, and vp is the volume fraction of the SiC particles (25vol%).
The increase of Young’s Modulus by addition of SiC particles is evidently bigger than ROM predicts.
The SiC particles in WB PRM exhibit a shape factor with particles oriented along the samples’ axis.
This strengthening effect might increase the elastic modulus somewhat but the measured difference to
the results according to ROM amounts to 18% at RT. In the case of BM PRM without any anisotropy,
the increase of volume fraction of dispersoids may cause a higher elastic modulus. The measured
Young’s Modulus at RT of BM PRM amounts to about 115GPa, more than 20% higher than the
prediction according to ROM. That would be the consequence of 30vol% ceramic reinforcement
instead of 25vol% SiC. In the case of WB 2124/SiC<5µm/25p maybe both contributions, shape factor
and increased reinforcing phase, yield the highest Young’s Modulus.
The mixing methods used in the present study lead to the following results:
1) In the case of WB PRMs the SiC particles are broken, but remain with their elongated shape having
more pointed corners than the particles which were subjected to mechanical alloying. It is concluded
that fracture of SiC particles occurs during extrusion.
2) By means of BM, the powders, especially the SiC particles, are rounded. During mechanical
alloying the size of SiC particles decreases significantly giving rise to a composite with finer SiC
particles than in the WB condition (see Figure 2). The difference in SiC particle sizes between the BM
and the WB PRMs can be seen in Table 4, which indicates the particle fragmentation when compared
with Table 2. In addition, BM produces sub-µm dispersoids, sub-µm SiC particles and oxides given in
Table 5.
The HT tensile strength due to the introduction of SiC particles also increases for the PRMs except for
those prepared by WB <25µm. The PRMs with the smaller SiC particles show an increase of the HT
tensile strength. A dependence of the HT tensile strength on the SiC particle size was found. The
smaller the initial SiC particle size the bigger the HT tensile strength. This is also the result for the BM
PRMs which show the highest HT tensile strength for initial SiC particle sizes <5µm. This can be
explained by the mobility of dislocations which is increased for the WB <25µm. Those materials show
the largest interparticle spaces and additionally the absence of the sub-µm SiC particles enhances the
mobility of dislocations. Dislocation movement is strongly decreased in the BM materials due to the
smaller SiC particles with its resulting smaller interparticle spaces and with increased content of subµm particles by pinning moving dislocations.
The BM PRMs show the highest creep resistance among all studied materials. The stationary creep
rate of the BM PRMs is about one order of magnitude lower than that of the BM 2124 unreinforced
matrix in the lower applied stress region (<50MPa for the BM PRMs). This difference increases up to
4 orders of magnitude for the higher applied stresses region (50-70MPa for the BM PRMs). The
stationary creep rate of the WB PRMs prepared with <5µm SiC particles is more or less the same as
for the corresponding unreinforced alloys. The creep resistance of the <25µm SiC WB PRMs is
considerably lower than that of the unreinforced matrices. The larger particle size present in the WB
composites is responsible for this behaviour [24-26]. The BM process reduces the reinforcement size
resulting in sub-µm SiC particles. The sub-µm SiC particles and the oxides act as dispersoids which
reduce the creep rate by pinning dislocations even if their motion is thermally activated [27].
All the unreinforced matrices, the 2124 without BM, BM 2124 and BM 6061 together with the WB
PRMs <5µm show a two region creep behaviour: the low stress region ranging from 15-30MPa
applied load and the high stress region from 30-60MPa applied load. The unreinforced 2124 alloy was
tested with and without BM process. The result shows that the matrix prepared by BM exhibits a
slightly higher creep resistance than the other unreinforced alloys. This comes from the higher volume
fraction of oxides formed around the Al particles during the BM Process (see Figure 1).
The WB PRMs reinforced with initial particle sizes <5µm contain a volume fraction of about 4vol%
of particles which are bigger than 5µm, whereas the WB PRMs with initial SiC particles <25µm result
in a total volume fraction of about 10vol% of particles which are bigger than 5µm. This may be an
indication that this larger amount of >5µm SiC provokes the lower creep resistance of the <25µm WB
PRMs. The absence of sub-µm SiC particles that could pin moving dislocations and the higher amount
of SiC particles >5µm are responsible for the low creep resistance shown by the WB PRMs.
For the WB PRMs <5µm the values of the creep exponent n1 is about 5. These values of n1 refer to the
low stress region (σA=15-30MPa) [28]. The WB PRMs <5µm show a change of the creep exponent n2
at σA=30MPa. The creep exponent n2 reaches values of >11 for the WB PRMs <5µm in the high stress
region similar to the unreinforced matrix.
The content of oxygen in the BM PRMs is 3 times higher than in the unreinforced BM matrix
suggesting a higher amount of oxides. Another consequence of BM is that during the blending process
continuously new matrix surfaces are formed due to repeated fracturing. In this way the oxide fraction
is also increased. The aluminide torn from the surface of the Al particles due to BM and extrusion act
as dispersoids increasing the creep resistance.
The reported controversial results of the last decades can be explained in this way: the addition of SiC
particles bigger than 5µm to an Al alloy have a detrimental effect on the mechanical properties by
acting as additional dislocation sources, whereas it can be seen that the introduction of smaller ones
enhances mechanical properties by acting as obstacles for the mobility of dislocations. The BM
blending route is the suggested method for producing MMCs since it provokes the rise of oxides
acting as dispersoids.
Conclusions




The introduction of 25vol% of SiC particles results in an increase of the Young’s Modulus of
about 50% independently of the initial particle size and the processing route.
There is an increase of HT tensile strength with a dependence of particle size introduced: the
smaller the particles the bigger the HT tensile strength. The ductility decreases with the
addition of SiC particles to an Al alloy independent of the size and of the interparticle
distances. The WB <25µm do not show higher strength due to early damage of the matrix by
big SiC particles with sharp edges.
The WB PRMs with particle size <5µm containing a small fraction of SiC particles (≈4vol%)
bigger than 5µm show approximately the same creep resistance as the unreinforced matrices.
Whereas 10vol% (more than 1/3 of the SiC particles) of SiC particles bigger than 5µm in WB
PRMs <25µm lead to a weakening of the creep resistance. The bigger the volume fraction of
the big particles the weaker the creep resistance because those act as dislocation sources [29].
A critical SiC particle size of about 5µm can be deducted below which the SiC particles act as
reinforcement at elevated temperatures. Bigger particles have a detrimental effect on the
mechanical properties due to the increasing interparticle spacing and their role as dislocation
sources [26, 30].
All the PRMs produced by BM show significantly higher elevated temperature strength and a
better creep resistance independent of the initial SiC particle size compared with those PRMs
produced by WB. This is the result of the broken SiC-particles of sizes <1µm and the
increased amount of oxides, both acting as dispersoids [12, 31].
Acknowledgements
The authors would like to thank the Austrian Science Fund (FWF) and the Austrian Agency for
International Cooperation in Education and Research (ÖAD) – Acciones Integradas for financial
support of this work. The University Service for Transmission Elektron Microscopy (USTEM) is
acknowledged for the provision of SEM and TEM facilities. ESK - SIC GmbH is acknowledged for
the provision of the SiC powders. The authors are also grateful to H.P. Degischer for supporting the
work by fruitful discussions and by assisting in the redaction of the manuscript.
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