MCEN 5208 Introduction to Research, 2004
Colorado University at Boulder
Literature Survey Report: Nano-Dispersion Strengthening of Aluminum
Student: Luke Fischer
Email: Luke. Fischer@colorado.edu
Advisors: Professor Rishi Raj, Dr. Atanu Saha
1
Background and Motivation for Nano-Dispersion Particulate Strengthened
Metal Matrix Composites
Ceramic particulate Metal Matrix Composites (MMCs) offer significant performance
advantages over pure metals and metallic alloys. The desirable properties of these
materials give them many potential applications in area’s such as in the automotive,
aerospace and sporting goods industries. MMCs can be classified into one of three broad
categories, namely particle reinforced MMCs, short fiber reinforced MMCs and
continuous fiber reinforced MMCs. The cost of manufacture of reinforcing fibers and the
processing costs of fiber reinforced MMCs are relatively high, making them
uneconomical for the majority of applications. Particulate reinforced MMCs have been
found to be more economical in the many applications.
Particulate MMCs can involve particle’s ranging in size from around 10nm up to 500μm
or larger. Composites with a fine, uniform dispersion of particles in the range of 10nm 1μm are referred to as “nano-dispersion”, “nano-scale dispersion” or “nano-metric
dispersion” strengthened composites. Much research has been conducted on particulate
reinforced MMCs, however the majority of this work has focused on micro-metric
particle dispersions which are easier to achieve than nano-metric particle dispersions, but
less effective in strengthening. It is important to remember that this project is concerned
with nano-metric dispersions; however the production and use of micro-metric dispersion
strengthened aluminum alloys can contribute significantly to the understanding of the
factors involved in the production and use of particulate MMCs in general.
For many years particulate MMCs have been heralded as an area for great expansion; for
example Aerospace America published an article entitled “Particulates Promise
Affordable MMCs” in 1995 [1]. Economic considerations have meant that the properties
of particulate MMCs have not been exploited in many applications. There are notable
exceptions however, such as Specialized Bicycles who have been using micro-metric
aluminum matrix MMCs in production for some time. In January 2003 Kevorkijan
published an article in the American Ceramic Society bulletin [2] which focused on the
slow uptake of micro-metric aluminum matrix MMCs in the automotive industry as a
result of the high production costs in comparison with aluminum alloys, and the high
design costs of using a new material.
Dispersion hardened MMC systems were first investigated in the first half of the 20th
century, however much of the foundation work in nano-dispersion strengthening of
metals was performed for the Ni-ThO2 system and the Ni-Cr-ThO2 system in the 1960’s
and early 1970’s [3]. The high temperature strength of these dispersion strengthened
alloys was of particular interest, with several papers being written on the high
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MCEN 5208 Introduction to Research, 2004
Colorado University at Boulder
temperature creep resistance. This research led to the production of the high temperature
nickel based “super-alloys”. While the scope of this project does not include testing of
samples at temperature, the excellent high temperature properties of nano-dispersion
strengthened MMCs remain a key motivation for this project and continued research in
this field.
The dispersed phase used to strengthen the aluminum matrix in this project will be silicon
carbonitride (SiCN), which shall be derived from a commercially available liquid
polymer precursor with the trade-name CerasetTM. There are no examples of this
compound being utilized as the strengthening phase of an aluminum matrix MMC in the
literature. The nature of this precursor lends itself to new possibilities in the manufacture
of nano-dispersion strengthened MMCs, which have prompted this project.
2
Particulate MMCs versus Heat Treatable Aluminum Alloys
Heat treatable aluminum alloys such as the widely used 6000 and 7000 series alloys have
excellent room temperature tensile strength. This strength is imparted by solid solution
strengthening, Hall-Petch grain size strengthening and precipitation strengthening (also
known as age hardening). Solid solution strengthening and Hall-Petch grain size
strengthening can also be effectively utilized in particulate MMCs with an inter-metallic
aluminum alloy matrix [4]. The most effective precipitation strengthening is imparted by
the formation of the coherent, meta-stable θ” phase. This phase is arrived at by rapid
cooling from the melt to achieve a solid solution followed by artificial ageing at a
temperature around 200ºC for a period of about an hour. If the heat treatment is continued
beyond the optimal time the alloy becomes “over-aged”, and the θ” phase is succeeded by
the semi-coherent equilibrium θ phase, which is accompanied by a reduction in strength
and hardness.
Herein lies a key disadvantage of precipitation hardened alloys – use at elevated
temperature for any period of time will result in a significant reduction in mechanical
properties, even after the material is returned to room temperature. In contrast the
particulates in particulate MMCs are virtually insoluble in the matrix phase at
temperatures below the melting temperature of the matrix phase. This means aluminum
matrix particulate MMCs can be used in components which are subjected to high
temperatures, such as automotive brake discs and pistons in internal combustion engines
[2], whereas precipitation hardened alloys cannot. The inclusion of the dispersed phase
also has a beneficial effect on the high temperature creep strength of the MMC, and the
prevention of grain growth, as discussed later in this report.
3
Micro-Scale versus Nano-Scale Particulate MMCs
The mechanical properties of nano-metric dispersion strengthened MMCs are far superior
to those of micro-metric dispersion strengthened MMCs with a similar volume
composition of particulate. For example, the tensile strength of a 1 vol.%SI3N4(10nm)-Al
composite has been found to be comparable to that of a 15 vol.%SiCp(3.5 μm)-Al
composite, with the yield stress of the nano-metric MMC being significantly higher than
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MCEN 5208 Introduction to Research, 2004
Colorado University at Boulder
that of the micro-metric MMC [5]. Particles larger than 1.5 μm tend to act as microconcentrators and are susceptible to cleavage. Particles in the range of 200-1500nm have
been found to cause the formation of cavitities and pits caused by poor inter-phase
cohesion. Particles smaller than 200nm generally bond well with the matrix, which is key
to the excellent mechanical properties of nano-scale particulate MMCs [5].
Based on the above information a reasonable target mean particle size for this project
would seem be 100nm or less (with a reasonably narrow particle size distribution and a
uniform dispersion). Having said this, the finer the particle size the more effective the
dispersion hardening will be (Section 4, Strengthening Theory).
The strain hardening exponent of particulate MMCs is much higher than that of the
matrix material (discussed in more detail below). The rate of strain hardening has been
found to increase with increasing volume fraction of particulate and decreasing particle
size [6].
4
Strengthening Theory
The strengthening mechanism resulting from dispersion hardening is known as Orowan
strengthening, a theory which is well established and is covered in many texts such as
Dieters “Mechanical Metallurgy” [7]. The mechanism is shown schematically in Figure 1
below. According to this mechanism the yield stress is determined by the stress required
for a dislocation line to pass by the two particles shown. The dislocation line is bowed
around the two particles as the applied stress is increased until the dislocation line reaches
a critical curvature (stage 2). When this critical curvature is reached the dislocation line
can then move forward without increasing its curvature (stage 3). The segments of
dislocation line on either side of each particle then join, and a dislocation loop is left
around each particle. As each dislocation line moves past a particle the dislocation cell
structure around the particle builds up. It is this phenomenon that results in dispersion
hardened metals having a high rate of strain hardening, as investigated by Hansen [6].
Figure 1 - Dislocation Line Passing Particles (L = λ)
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MCEN 5208 Introduction to Research, 2004
Colorado University at Boulder
The basic equation for the Orowan shear stress is:
O 
Gb
(1)

Where G is the shear modulus of the matrix material, b is the burgers vector or lattice
parameter of the matrix material and λ is the inter-particle spacing. This relationship has
been refined by a number of people to incorporate the results of more accurate models of
the dislocation line tension and the energy incorporated in the dislocation segments on
either side of the particle. One common form is the Orowan-Ashby equation [7]:
 
0.13Gb

ln
r
b
(2)
In this equation r is the radius of the particle and the other parameters are as defined
above. Note that the particles are assumed to be spherical. The relationship between λ, r
and the volume fraction of dispersoid, f, can be found from the following expression
based on this assumption [7]:

4(1  f )r
3f
(3)
It is a reasonable to assume that the particles in the mechanically alloyed system to be
investigated in this project will be spherical, although other system’s which have been
investigated have plate-like strengthening as a result of the processing techniques used.
Hansen [6] used analytical methods to show that high aspect ratio plate-like particles give
calculated particle spacing’s which are not significantly different from those obtained
using spherical particles of the same volume. If the dispersed particles observed in the
samples deviate from the spherical assumption to a significant extent a more accurate
expression for λ can be obtained or derived based on the geometry observed.
Perhaps the most important assumption’s which must be made in utilizing the Orowan
equation or any equation based on the Orowan mechanism involve the particle size
distribution and the uniformity of the dispersion of particles. For any such equation to be
accurate the particle size distribution should be narrow (i.e. low standard deviation of
effective diameter) and the particles should be well dispersed in the matrix. Processing
techniques and thermodynamic considerations (involving interfacial energies in the melt)
can lead to agglomeration or grouping of particles which greatly reduces the dispersion
strengthening effect. This is often leads to discrepancies between the experimentally
observed strength and that calculated [5]. Any significant grouping or agglomeration in
the samples in this project should be detected using optical or electron microscopy and
hence can be accounted for.
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Colorado University at Boulder
Build-up of dispersed phase in the grain boundary regions will also reduce the magnitude
of the strengthening effect for a given volume fraction of dispersoid. Only those particles
in the grain can contribute to the Orowan strengthening mechanism.
Recent papers such as those written by Tong and Fang (1998) [8, 9] and Kang and Chan
(2004) [5] use different forms of the Orowan equation; however it should be noted that
the equations utilized in these papers were derived in 1972 and 1966 respectively. There
appears to be no agreement on which form of the equation is the most accurate model of
the Orowan mechanism. For this reason the Orowan-Ashby equation, which is the most
common [7], shall be used as a starting point.
In many systems Orowan strengthening is not the only mechanism at play; other
mechanism such as Hall-Petch strengthening due to grain size and solute strengthening
(in alloyed aluminum) also contribute to the strength of the resulting MMC. It appears to
be unanimously accepted that the strength of an MMC can be assumed to a linear
summation of the strength imparted by each mechanism [8, 10, 5, 11]. It is not expected
that the Hall-Petch mechanism will play a significant role in the strengthening of the
samples made in this project due to the slow rate of cooling from the melt, which will
mean that dispersion strengthening effect can be isolated.
Chawla describes the nature and importance of interfacial bonding in his book
“Composite Materials” [11]. The interface in the in the Al-SiCN system will be noncoherent and there will be some degree of interfacial bonding by means of mechanical
keying, depending on the roughness of the particle surface. Chemical bonding is caused
by reaction products forming at the particle matrix interface. It is conceivable that SiC
produced in the pyrolosis of the polymer particles could react with the aluminum melt to
form a small quantity of Al4C3 at the interface.
5
High Temperature Properties
As discussed earlier dispersion hardening is highly beneficial to the high temperature
properties of a MMC. Dispersed ceramic particles are stable at temperatures up to the
melting temperature of the matrix metal and do not tend to coarsen at elevated
temperatures, meaning there is very little drop off in the dispersion strengthening effect.
Weissgarber and Kieback [10] show that the elevated temperature strength of DISPAL
aluminum matrix MMCs with a micrometric dispersion of Al4C3 particles is considerably
better than that of aluminum alloys which have a higher strength at room temperature. In
addition to this the DISPAL MMCs lose very little hardness in annealing at temperatures
up to 500ºC, whereas age hardened aluminum alloys can decrease in strength by up to
50%.
The high temperature creep strength of metals is also greatly improved by the addition of
a high temperature stable dispersed phase, due to grain boundary pinning. Grain
refinement is not a desirable in high temperature materials as it results in a higher rate of
grain boundary sliding, which is a key mechanism of creep. For this reason dispersion
Nano-Dispersion Strengthening of Aluminum
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Colorado University at Boulder
strengthening is of particular importance in creep resistant metals such as the oxide
dispersion strengthened super alloys [3].
6
The Processing of CerasetTM and the Pyrolytic Conversion to Silicon Based
Ceramics
A practical understanding of how silicon based ceramics can be derived from the liquid
polymer precursor CerasetTM (manufactured by Allied Signal Ceramics Inc.) is necessary
to produce MMCs using SiCN as a dispersoid. The reactions and transformations
involved must characterized in order to determine processing techniques and parameters
to be used in the production of the MMC.
Li et al. [12] wrote a paper entitled “Thermal Cross-Linking and Pyrolytic Conversion of
Poly(ureamethylvinyl)silazanes to Silicon Based Ceramics” in 2001 which can be used to
gain the appropriate level of understanding of the nature of the transformations (note that
the chemical name of CerasetTM is poly(ureamethylvinyl)silazane, abbreviated as
PUMVS).
The soluble and meltable nature of PUMVS makes it suitable for the generation of
coatings. One technique which may be trialed in this project involves wet mixing
aluminum powder in a solvent containing PUMVS, taking advantage of the soluble
nature of PUMVS.
Liquid CerasetTM is cross-linked almost instantaneously at temperatures of 300ºC or
higher. The solid precursor undergoes pyrolytic conversion to SiCN ceramic in the
temperature range 600-800 ºC. The overall ceramic yield with respect to the starting
PUMVS liquid is 70 wt%, and the yield with respect to the cross-linked solid is 78 wt%.
The overall ceramic yield was found to be independent of the cross-linking temperature.
The ceramic which is formed is an amorphous silicon carbonitride (SiN0.82C0.86). The fact
that the ceramic is amorphous has important implications with regard to performing x-ray
diffraction on the MMC samples, because the amorphous ceramic phase will not
contribute any peaks to the x-ray diffraction results.
Formation of free carbon was evident in the temperature range 700-800ºC. Formation of
graphite was evident at temperatures over 1000 ºC when high heating rates were
employed. This information indicates that the cross-linked polymer powder should be
pyrolised in the aluminum melt in the temperature range of 800-1000 ºC. The time
required for the pyrolytic transformation to be completed will have to be determined
experimentally, however it is undesirable to hold the melt at temperature for
unnecessarily long time as this could have undesirable effects on the particle size and
distribution.
A considerable amount of work involving CerasetTM and SiCN has already been
conducted at Colorado University, Boulder.
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Colorado University at Boulder
7
Nanoscale Particulate MMC Production Processes in the Literature
Various processes have been used to develop dispersed particle MMCs, such as
mechanical alloying/powder metallurgy [5, 10, 13], “in situ” formation of dispersoids via
a chemical reaction within the matrix phase [8, 9, 14], spray deposition [11] and several
casting techniques [15]. The latter two techniques have been employed primarily in the
manufacture of micrometric particle reinforced MMCs.
The matrix metal, aluminum, is susceptible to oxidation as are some of the compounds
used as reinforcement particles, including CerasetTM. For this reason many of the
processes used to manufacture these MMCs must be carried out in a vacuum or an inert
gas environment, particularly when powders are involved, due to the high surface area
which is exposed to the atmosphere. A notable exception to this is cases where aluminum
powder is deliberately oxidized to create the aluminum oxide reinforcement phase [6].
The method of mechanical alloying was developed in the 1970’s for the production of
thoria dispersed nickel based super-alloys. A review of mechanical alloying was
published by Weissgarber and Kieback [10] in 2000. This method involves milling a
mixture of metal and non-metal powder in a high energy ball mill (several types are
available – the mill to be used in this project is of the vibratory type). The milling process
causes the powder particles to undergo severe plastic deformation and results in cold
welding of particles. The end result is a mixture of composite metal particles in which the
dispersoid particles are reduced in size and uniformly distributed in the metal matrix.
The resulting powder mixture is usually consolidated by hot extrusion or hot isostatic
pressing (HIP).
Powder metallurgy was used by Kang and Chan [5] to manufacture nano-metric Al2O3
particle reinforced aluminum. The starting Al2O3 mean particle size in this case was
50nm. The powder metallurgy process involved wet mixing (aluminum powder mixed
with varying volume fraction of Al2O3 powder in a pure ethanol slurry), following by
drying at 150 ºC then cold isostatic pressing (CIP, as opposed to HIP) to compact the
powders. The compacted powders were sintered in a vacuum at 620 ºC (approx. 60 ºC
below the melting temperature of aluminum).
Ductile materials are difficult to ball mill due to particle coarsening, resulting in the
advent of reaction milling in which dispersoid particles are formed by a chemical reaction
with the matrix element. This may occur during milling, or at a later stage during heat
treatment. The aluminum based MMCs with the trade-name DISPAL, which are
reinforced with Al4C3 particles, are manufactured using this method. Cu-TiC MMCs have
also been manufactured using this method [10].
There are a number of “in situ” reaction dispersoid formation techniques which have
proven successful in recent years. A key advantage of in situ methods is that they are not
limited by the starting powder size which often determines reinforcement particle size.
Cui et al. [14] published a review of this field in 2000 in a paper titled “Review of
Fabrication Methods of in situ Metal Matrix Composites”. These methods avoid the
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Colorado University at Boulder
problems created by poor particle wetting which often occur when fabricating particulate
MMC’s using conventional methods [14].
In situ methods are classified in five different categories. The first of these is powder
metallurgy, in which powders are blended in what appears to be a similar technique to
reaction milling as discussed above. The chemical reaction actually occurs after the
milling event when the powders are subjected to isostatic hot pressing at 1000ºC. Cui et
al. describe this method being used to create TiC-Al composite by blending Al powder,
Ti powder and graphite powder in what appears to be a very similar process to that
described by Weissgarber and Kieback [10] for a TiC-Cu composite.
The second method discussed by Cui et al. is liquid-gas reactions, in which gases
containing a reactive component are bubbled through a matrix metal melt, causing a
reaction to take place in which the dispersed phase particles are formed. The example
given is for a Al-AlN composite. The third method described is the slightly more
complex solid-gas-liquid reaction method. One example given by Cui et al. is that of an
aluminum matrix alloy containing TiN and AlN particles which are formed in a high
temperature (1573K) reaction between aluminum liquid, Al3Ti solid and N2/NH3 gas
mixture.
Tong and Fang [8] have done much research into the fourth method, ingot metallurgy.
This method would appear to be more simple, economical and flexible than many other
techniques. The experiments performed by Tong and Fang created an Al-TiC composite.
The technique used involved mixing micro-scale powders of Ti, Al and graphite,
subjecting the powder mixture to a dual stage heating process at ~1400K followed by
~1600K, then direct-chill casting into 12mm diameter ingot bars. The MMC then
underwent rapid solidification by chill block melt spinning, in which thin ribbons are
prepared by heating ingots using high frequency induction to 1623K then cooling on a
spinning copper wheel with a circumferential velocity of 30m/min. The resulting grain
structure is very fine hence the Hall-Petch type strengthening mechanism is significant.
Cui et al. found the fifth and final method of in situ dispersoid formation, plasma reaction,
to be the most effective, largely because it can be used in a large number of
matrix/reinforcement systems. The method involves entraining fine reinforcement
particles in ionized gas, usually argon or helium, at a very high temperature (in the range
2200~7000 ºC, but below the melting or sublimation temperature of the particles) then
injecting the gas stream into a matrix metal melt which is mechanically agitated to ensure
uniform mixing. The physical agitation is continued until the mixture is completely
solidified.
Spray deposition is a technique which has been used in the manufacture of micro-metric
particle MMCs as opposed to nano-metric MMCs [11]. This process uses a spray gun to
atomize a molten aluminum matrix into which ceramic particles are injected creating a
preform. This preform is then subjected to scalping, consolidation and finishing processes
meaning the end material is essentially a wrought material. The time the stream is in
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Colorado University at Boulder
flight in the atmosphere is extremely short meaning that unwanted reactions are not able
to take place.
Casting techniques have also been used in the creation of micro-metric particle reinforced
MMCs. Roghati et al. wrote a paper titled “Solidification, Structures and Properties of
Cast Metal–Ceramic Particle Composites” [15] which investigates in detail the issues
surrounding particle wetting/interfacial energies, density differences between the matrix
melt and the dispersed particles, and other factors which affect the processing and
properties of cast MMCs. Although this paper was concerned with micro-metric scale
particle reinforcement rather than nano-metric scale particles it is of interest because the
first method to be trialed in the manufacture of the MMCs in this project is essentially a
casting technique. One important point from this article is that a high density differential
between the melt and the dispersed particle will result in either settling or rising if time
allows, however it is not thought that this will be a problem with the Al-SiCN system. In
addition to this the dispersed phase will increase the viscosity of the melt, with a higher
volume fraction or a smaller particle size resulting in a high viscosity.
8
Summary of Nano-scale Particulate MMC Production Techniques to be
Trialed in this Project
A fine CerasetTM powder is to be created by cross-linking the liquid CerasetTM at 400ºC
in an inert gas environment, thus creating an infusible solid which is then to be ball
milled to give a fine powder.
The first technique to be trialed involves ball milling of the micro-metric aluminum
powder and cross-linked CerasetTM powder in a vibrating ball mill, in a similar manner to
that the initial stages of mechanical alloying [10]. This will ensure we have a fully mixed
composite powder with a fine particle size.
The resulting powder is then to be heated in a furnace in an inert gas environment to
approximately 800ºC (well above the melting temperature of aluminum) at the fastest
possible rate. The CerasetTM powder will undergo pyrolosis conversion to SiCN in the
range of 600-800ºC. It is hoped that the pyrolosis process will break up the precursor
particles into finer particles as the hydrocarbons are released in the reaction. The density
of SiCN is 2.4g/cm3, slightly lower than the density of pure aluminum which is 2.7 g/cm3.
Because these densities are relatively close it is not expected that there will be a problem
with settling or rising of the particles in the casting process.
The second technique to be trialed, if the first is unsuccessful, involves wet mixing and
drying similar to the method employed by Kang and Chan [5]. The aluminum powder
and the CerasetTM precursor will be mixed in a solvent solution which will then be dried.
The resulting mixture will be heated in a furnace to well above the melting temperature
of aluminum so that the CerasetTM particles undergo pyrolosis.
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Colorado University at Boulder
Both of the methods discussed above employ a relatively slow cooling rate, so the
resulting grain structure is expected to be relatively coarse, meaning it will have little
effect on the strength and hardness of the samples.
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