A study of solid-state amorphization in Zr–30 at.% Al by mechanical attrition

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.

600 J. Mater. Res., Vol. 11, No. 3, Mar 1996


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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

J. Mater. Res., Vol. 11, No. 3, Mar 1996 601


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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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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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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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A study of solid-state amorphization in Zr–30 at.% Al by mechanical attrition