7/28/2019 A study of solid-state amorphization in Zr30 at.% Al by mechanical attrition
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A study of solid-state amorphization in Zr30 at. % Al
by mechanical attrition
A. Biswas, G. K. Dey, A. J. Haq, D. K. Bose, and S. BanerjeeMetallurgy Division, Bhabha Atomic Research Centre, Bombay 400085, India
(Received 25 January 1995; accepted 6 November 1995)
Elemental powders of zirconium and aluminum in the atomic ratio of 70 : 30 were
mechanically alloyed in an attritor under argon atmosphere using zirconia balls as milling
media. Samples have been taken out for characterization after different durations of
milling. The process of alloying and resultant amorphization had been studied usingx-ray diffraction (XRD) and transmission electron microscopy (TEM). Scanning electron
microscopy (SEM) was carried out to study the morphological changes occurring
during repeated cold welding and breaking of the particles. Samples for TEM study
were prepared by dispersing the mechanically attrited particles in the nickel foil by
electrochemical codeposition. TEM study of the initial stages of milling revealed that
localized structural changes precede the bulk amorphization process during mechanicalalloying (MA). The sequence of phase evolution has been identified as (i) the formation
of nanocrystalline supersaturated solid solution of aluminum in a-zirconium, (ii)
amorphization of localized regions at powder interfaces, (iii) ordering of aluminum-richregions in the metastable Zr3Al (DO19) phase, and, finally, (iv) bulk amorphization
of the powders.
I. INTRODUCTION
Mechanical alloying (MA) was first reported1 in
1966 as a technique for the preparation of the oxide
dispersion-strengthened nickel alloys suitable for high
temperature applications. Solid-state amorphization bythis technique of high energy milling2,3 revived the
interest in this topic. Now, MA has become a versatile
processing method which is capable of preparing a
truly wide range of materials with unique properties,
namely intermetallics,4,5 alloys of immiscible metals,6
nanocrystalline phases,7 quasicrystals,8 amorphous,918
and other metastable phases19,20 in bulk quantities.
It has been shown that solid-state amorphization
can be achieved both from intermetallic powders and
mixtures of elemental powders. Usually, the amorphiza-
tion from intermetallic powders is termed as mechanicalmilling (MM),21 unlike mechanical alloying.
The mechanisms of the solid-state amorphization
and associated transformations in different systems are
not yet well understood. A number of probable theories
have been proposed so far. Yermakov et al.2 explainedthe process in terms of local melting and subsequent
rapid solidification. However, evidence of melting could
not be seen in any of the amorphization experiments.
Others hypothesized solid-state processes to be instru-
mental for this transformation. In early papers9,22 it has
been proposed that negative enthalpy of mixing and
widely differing diffusivities are two necessary condi-
tions for amorphization. However, many exceptions have
been reported later.23,24
Composition-induced destabilization of the crystal
lattice and resultant amorphization has been shown
recently25 in the Zr Al binary alloy system, where
a supersaturated aZr solid solution forms up to an
aluminum concentration of 15 at. % and amorphization
takes place in Zr1002xAlx when 15 , x , 40. At
x 50, a metastable nanocrystalline fcc phase (ZrAl)
evolves. Ma and Atzmon have studied this system
FIG. 1. Particle size distributions of initial and mechanically attrited
powders.
J. Mater. Res., Vol. 11, No. 3, Mar 1996 1996 Materials Research Society 599
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A. Biswas et al.: Study of solid-state amorphization in Zr 30 at. % Al
TABLE I. Median diameters of initial and mechanically attrited
powders.
Sample description Median diameter (mm)
Initial Al 27.0
Initial Zr 11.8
15 h 9.3
30 h 5.4
45 h 4.7
60 h 5.7
further and given calorimetric evidence for chemically
induced polymorphic transformation by determination of
enthalpies of both the supersaturated solid solutions and
the amorphous alloys of varying aluminum concentra-
tions. They have shown that the critical concentration
of aluminum required for amorphization is 17.5 at. %.
However, in both of these investigations it has been
found that the intermetallics, namely Zr3Al, Zr2Al,
and Zr3Al2, do not form due to kinetic constraints
associated with long-range ordering. In the current
investigation, MA of elemental powders of zirconium
and aluminum in the atomic ratio of 70 : 30 was taken
up for studying the structural evolution, as revealed
by TEM which has not been reported earlier. Wemade an attempt to look into the possibility of the
formation of any equilibrium or metastable ordered in-
termetallic phase in the evolutionary path of mechanical
alloying.
II. EXPERIMENTAL
Elemental powders of zirconium and aluminum of
the purity of 99.5% were alloyed in an attritor under
argon atmosphere. Five mm diameter zirconia balls were
used as milling media, and the ball-to-powder weight
ratio was kept at 10 : 1. The milling had been done in
FIG. 2. Change in particle morphology during milling. (a) Initial Zr, (b) initial Al, (c) 10 h milled powder, and (d ) 15 h milled powder.
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A. Biswas et al.: Study of solid-state amorphization in Zr 30 at. % Al
a water-cooled vessel to keep the milling environment
close to the ambient temperature. Milling was carried
out for up to 60 h at a constant milling speed of
550 rpm. Samples were taken out of the milling vessel
for characterization after different degrees of milling.Particle size distributions and corresponding median
particle sizes of the samples after different durationsof milling were determined in a Sedigraph. Changes in
the morphology of the attrited powders were studied by
SEM. Phases forming during alloying were characterized
by XRD analysis using Cu Ka radiation. TEM investiga-tions of selected samples were performed. Samples for
TEM were prepared by dispersing mechanically alloyed
particles in a nickel foil by electrochemical codeposition
from a nickel solution containing the attrited particles
in suspension.27 Subsequently, samples were thinnedinitially by ion milling followed by electropolishing
using either the jet or the window technique. This
technique ensured the absence of an ion-damaged area
in the electron transparent region.
III. RESULTS
In the course of milling, the particles first became
shiny and finally lost their luster and ended up as dark
fine powders with no visible heterogeneity. During thisprocess, the zirconia balls became coated, as was evident
from visual examination.
Particle size distributions of the initial unmixed
powders and that of the powders attrited for different
durations are given in Fig. 1. Median particle sizes of thestarting powders of zirconium and aluminum were found
to be 11.8 mm and 27 mm, respectively. The median size
of the mixed powder decreased to 9.3 mm after 15 h
milling and to 4.7 mm after 45 h (Table I). Still further
milling did not cause any refining; instead, a coarsening
effect was observed.Figures 2(a) and 2(b) show the initial morphologies
of aluminum and zirconium powders, while Fig. 2(c)
depicts the cold-worked morphology after sufficient time
of milling, which is a typical feature of this process.
Figure 2(d) clearly shows the broken pieces of mechani-cally attrited powder particles.
X-ray diffraction of samples taken from different
stages of milling provides information of the changesthat took place during the milling process. The (111) and
(200) peaks of fcc aluminum gradually diminished andhcp a-zirconium peaks broadened and shifted toward
high-angle values, as shown in Figs. 3(a) and 3(b). These
are in agreement with observations made earlier.25,26,28
After 15 h of milling, XRD results showed the presence
of a solid solution of aluminum in a-zirconium. Powders
obtained after the 20 h milling sample did not exhibit anysharp Bragg peak, except one broad peak correspond-
ing closely to (1010) of a-zirconium. Powders milled
FIG. 3. XRD patterns after different periods of milling. (a) Formation
of solid solution. (b) Ordering and bulk amorphization.
for 25 h sample showed some extra reflections that
disappeared during further milling. These were found to
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FIG. 4. TEM of 3 h milled sample. (a) Microstructure of powders embedded in nickel matrix (b) SAD pattern of Zr Al solid solution,
(c) unalloyed Al, and (d) microdiffraction from local amorphous region.
correspond to a lattice spacing of 5.4 nm, which matchesclosely to the superlattice reflections of the metastable
DO19(Zr3Al) structure. On further milling, the powders
transformed into amorphous phase. This bulk amorphous
phase remained unchanged up to 60 h of milling. Thewidth of the broad peak corresponding to the first near-
neighbor distance in the amorphous phase gradually
increased. Another interesting observation was that after
10 h of milling the (1010) peak ofa-zirconium became
the most intense, and remained so up to the time of bulkamorphization.
TEM studies were carried out with powder particlesembedded in a nickel matrix. The sequence of the
gradual transformation, as observed through XRD anal-
ysis, was investigated by TEM. Additional local features
could be observed that XRD was unable to reveal.The micrograph and corresponding electron diffraction
patterns shown in Fig. 4 are for a sample milled for
3 h. They show the presence of unalloyed aluminum,
partially alloyed a-zirconium conforming to XRD re-
sults, and, interestingly, some local amorphous regionsthat were not observed in the XRD pattern. Figure 5
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A. Biswas et al.: Study of solid-state amorphization in Zr 30 at. % Al
FIG. 5. TEM of 15 h milled sample. (a) Nanocrystalline structure containing both ( b) Zr Al solid solution and (c) amorphous phase.
shows the nanocrystalline structure of the 15 h milled
sample whose corresponding electron diffraction showed
the presence of both the amorphous phase and the
solid solution of aluminum in a-zirconium. The 20 h
milled sample showed a still finer structure of predomi-
nantly a-zirconium-aluminum solid solution, as shown
in Fig. 6. Ordering was noticed in the 25 h milledsample by the appearance of the weak innermost ring
(superlattice reflection-1010) of the diffraction pattern, as
shown in Fig. 7. This sample was found to be composed
of three phases, namely, DO19(Zr3Al), aZrAl solid
solution, and amorphous. The micrograph and diffraction
pattern in Fig. 8 corresponds to the 60 h milled sample,
demonstrating the presence of the bulk amorphous phase.
IV. DISCUSSION
The new findings of the present study are (i) local
amorphization at the interface of the particle in the early
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FIG. 6. Microstructure of 20 h milled sample. (a) Bright field. (b)
Dark field.
stage of milling, and (ii) formation of the metastable
DO19 structures.
Local amorphization results due to the compositional
heterogeneity at the interface, which is quite likely at
the initial stages of milling when a sharp composition
gradient is present from the core to the periphery of
each powder particle. Similar compositional gradients at
the particle interface have also been reported earlier forthe NiZr system, where the concentration depth profile
was measured by Auger Electron Spectroscopy.29
The appearance of the DO19 ordered phase was
observed after 25 h of milling. This was detected in
the selected area diffraction pattern (Fig. 7), where
the weak innermost diffraction ring corresponds to the
1010 d-spacing of the metastable Zr3Al phase of theDO19 structure. The sequence of structural evolutioncan be described in the following scheme: aZr 1
Al ! aZrAl solid solution 1 Al ! nanocrystalline
solid solution 1 amorphous! Zr3Al DO19 1 solid
solution 1 amorphous ! bulk amorphization.Free energy-composition (G-X) plots have been used
earlier by previous workers25,26,28 for explaining the
transformation that occurred during mechanical alloying
in the ZrAl system. In these investigations the free
energy values were either theoretically calculated orthe measured enthalpy values of the samples (havingdifferent aluminum concentrations) were used as an
approximation of the free energy. Prior to enthalpy
measurements in the aforementioned studies, samples
at different mixture compositions were milled for suf-
ficiently long times until they attained the steady state.The observed amorphization was explained in terms of
a concentration invariant polymorphic process (depicted
in G-X plot as vertical lines). Aluminum enrichment
ofaZrAl solid solution to a level of approximately
15 at. % Al made the polymorphic amorphization ther-
modynamically possible. It was envisaged25 that in the
first phase of milling, the aluminum concentration con-tinues to build up in the a-zirconium lattice to a level
of 15 at.% when the a-lattice becomes unstable with
respect to the polymorphic amorphization process. This
conclusion was reached as there was no evidence ofa two-phase structure (aZr Al solid solution and
amorphous) in XRD results.
The present work demonstrates that even after 3 h
milling amorphization can take place locally at the
particle interface and metastable ordering takes placeintermediately before bulk amorphization. Similar in-
termediate ordering was also reported in the Ti Al30
system. Although the solute concentration progressivelychanges during the course of milling, as alloying is
a gradual process, the G-X plot does not reflect the
transformations that occurred during the milling. More-over, as pointed out by Yavari et al.,31 it does not
differentiate between the amorphization from pure com-
ponents and that from any intermediate intermetallic
products. In order to explain the observed course of
transformation and phase evolution, we have consideredpartitioning of solute element between the competing
phases. A schematic G-X plot shown in Fig. 9 wasused, and the possibility of establishing local chemicalequilibrium was considered. With an increasing degree
of aluminum enrichment, the free energy at the interfaceregion gradually moves along the path 12. Once the
composition crosses point 2, it becomes thermodynami-
cally possible to nucleate the Zr3Al phase. Though the
equilibrium structure of Zr3Al is L12, there exists acompeting metastable DO19 structure, the latter being
a superlattice of the hcp a phase. As reported earlier,32
DO19 nucleation is kinetically favored when precipitation
occurs from the supersaturated a phase, presumably
because of a one-to-one lattice correspondence with
aZr and nearly equal spacings in the corresponding di-
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A. Biswas et al.: Study of solid-state amorphization in Zr 30 at. % Al
FIG. 7. TEM of 25 h milled sample. (a) Fine microstructure, ( b) SAD pattern showing DO19 ordering, (c) SAD pattern of solid solution,
and (d) microdiffraction from amorphous phase.
rections that ensures very good registry between the twophases.
With further aluminum enrichment, the composition
crosses point 3 when nucleation of the amorphous phasebecomes possible. It is to be emphasized that the com-
positional change occurs gradually from the interfaceand, therefore, the core of a particle remains crystalline
even when the amorphous phase starts appearing at the
interface. The present work points out that nucleation
of the DO19 phase and of the amorphous phase occurs
through alloy partitioning.As the composition of powder particles crosses
point 4, each particle as a whole can transform into an
amorphous phase by a polymorphic process.
The competition between the disordering process
by mechanical alloying and the reordering processdue to the thermodynamic tendency for the formation
of the more stable structure determines the steady-state structure in a given system. The evolution of
the steady-state structure in a driven system has been
studied theoretically by Haider et al.33 Experimental
observations reported here clearly point out that the
crystalline order is not sustainable in Zr30 at. %
Al alloy in steady state under the milling condition
employed.
It is also possible that the DO19 phase formed locally
had also undergone amorphization due to the damage
introduced by mechanical working.
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FIG. 8. TEM of 60 h milled sample. (a) Microstructure of amorphous
phase (dark region) in nickel matrix and (b) SAD pattern showing
bulk amorphization.
V. CONCLUSION
The sequence of phase evolution in mechanical
alloying of the ZrAl binary system is (i) the formation
of hcp zirconium-aluminum solid solution, (ii) interme-
diate formation of metastable intermetallic Zr3Al and an
amorphous phase in localized regions at the interfaces ofthe particles, and (iii) bulk amorphization. This has been
rationalized by a schematic free energy-composition plot.
FIG. 9. Schematic free energy versus composition diagram of Zr Al
alloy system.
Compositional inhomogeneity resulting from concentra-
tion gradients present at the initial stages of milling is
responsible for the formation of localized amorphous
regions.
ACKNOWLEDGMENTS
The authors wish to thank Mr. A. R. Biswas, Dr.
D. D. Upadhaya, Dr. N. C. Soni, and Mr. S. N. Athavale
for their useful services. We gratefully acknowledgethe valuable suggestions of Dr. S. K. Roy and Dr. P.
Mukhopadyay. The authors are indebted to Dr. C. K.Gupta, Director, Materials Group, B.A.R.C. for his sup-
port and encouragement during the course of this work.
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