Materials Science and Engineering A 417 (2006) 110–113

Polymorphic phase transformation in Ti50Zr50 binary alloyby mechanical alloying

S. Bera, I. Manna∗Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Kharagpur, West Bengal 721302, India

Received in revised form 4 October 2005; accepted 1 November 2005

Abstract

The present study reports a reversible hexagonal close packed (hcp) to face centered cubic (fcc) polymorphic phase transformation in Ti50Zr50

elemental powder blend in the course of mechanical alloying at ambient temperature in a planetary ball mill. This transformation is monitoredat appropriate stages by X-ray diffraction and high-resolution transmission electron microscopy. It is suggested that structural instability duetoplastic strain, increasing lattice expansion, and negative (from core to boundary) hydrostatic pressure is responsible for this reversible hcp–fcctransformation. Isothermal annealing at 1000◦C leads to partial transformation of the milled product from single phase fcc to hcp. Thus, the presentstudy confirms that Ti50Zr50 binary alloy undergoes precisely the same metastable hcp to fcc polymorphic transformation that was earlier reportedi©

K

1

plm(aptoaanonhciph

uchrlierhcpdentre,lat-

on),ssuretudy,

ationns-

ion;par-the

ental

boutmic

0d

n elemental Ti and Zr.2005 Elsevier B.V. All rights reserved.

eywords: Mechanical alloying; Polymorphism; Nanocrystalline material; Ti–Zr alloy

. Introduction

Nanocrystals are ultrafine single or multiphase mono/olycrystalline materials with nanometric grain or coherence

ength size in at least one dimension[1,2]. Due to its attractiveechanical (modulus of elasticity, hardness, etc.) and functional

chemical reactivity, sinterability, heat capacity, magnetic, cat-lytic, electrical resistivity, etc.) properties, nanocrystals (orolynanocrystalline aggregates) continue to draw the atten-

ion of the scientific community. Among the possible routesf synthesis of nanocrystalline metallic and ceramic materi-ls, mechanical alloying or attrition offers an easy, flexiblend inexpensive option capable of producing materials withovel microstructure in large quantity[3,4]. Besides noveltyf the microstucture produced by mechanical alloying (e.g.ano-intermetallic dispersed amorphous matrix Al-alloy[5]),igh-energy ball milling is reported to yield a body-centered-ubic (bcc) to face-centered-cubic (fcc) polymorphic changen elemental niobium due to nanocrystalization[6–8]. Similarolymorphic change in elemental zirconium and titanium fromexagonal close packed (hcp) to face centered cubic (fcc) was

also recently reported[9,10]. Besides elements, evidence of spolymorphic transformation by mechanical attrition was eareported in Ti–Al alloy[11]. These studies revealed that theto fcc polymorphic transformation was irreversible, depenon the milling intensity and not impurity driven. Furthermochange in crystal structure was primarily necessitated bytice instability due to nanocrystalization (grain size reductilattice expansion or strain, and negative hydrostatic prefrom the core to the surface of the grains. In the present swe present evidences of polymorphic hcp to fcc transformin Ti50Zr50 binary alloy and analyze the genesis of this traformation in terms of systematic microstuctural investigatmodel based thermodynamic calculation and suitable comison with earlier relevant studies. In general, we concludepresent results are identical as that reported earlier for elemNb, Ti or Zr.

2. Experimental procedure

Elemental titanium sponge and zirconium powders of a50–100�m size and >99.5 wt.% purity were mixed at an ato

∗ Corresponding author. Tel.: +91 3222 283266; fax: +91 3222 282280.E-mail address: imanna@metal.iitkgp.ernet.in (I. Manna).

ratio of Ti50Zr50 and milled in a Fritsch P6 planetary mono-mill with ZrO2 vials and balls up to a cumulative period of20 h. The ball to powder weight ratio and speed of the millwere kept constant at 10:1 and 300 rpm, respectively. To avoid

921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2005.11.012

S. Bera, I. Manna / Materials Science and Engineering A 417 (2006) 110–113 111

agglomeration or clogging of powders during milling, about50 ml toluene was added to the vials. The samples, taken atdifferent stages of milling were examined by X-ray diffraction(XRD) analysis using Cu K� radiation (0.1542 nm). Averagegrain size (dc) was calculated through peak broadening analysisusing Scherrer equation[12] after elimination of the respec-tive influence of strain and instrumental error by Lorentziancurve fitting[13]. Selected samples were studied by a 200 kVtransmission electron microscope (TEM) to obtain bright anddark field images including high-resolution lattice imaging atappropriate stages of milling. Besides microstructural investi-gation, the impurity content in the samples was also assessedby energy dispersive spectroscopy (EDS) attached to the TEM.The substitutional and interstitial impurity contents in the milledproduct at appropriate stages were further estimated by X-rayfluorescence spectroscopy (XRF) and dedicated LECO analyz-ers, respectively. Finally, the alloyed powders were annealed at1000◦C in vacuum for 2 h and re-examined with XRD to studythe reversibility of the polymorphic change in the present system.

3. Results and discussion

Fig. 1 shows the phase evolution sequence as a function ofmilling time, monitored by XRD of the Ti50Zr50 powder blend.It is evident that the elemental peaks of titanium and zirco-n tiona ralt icedoB andrT them 1 1)p -ttfi theh f theo rs asp

ucto n-

F -i

Fig. 2. XRD pattern of fcc Ti50Zr50solid solution obtained by 20 h of mechanicalalloying.

ing peaks are due to ZrO2. The average lattice parameter (a)of the alloy at this stage calculated from different peaks variesbetween 0.4490 and 0.4508 nm. The precise lattice parameter(a0) is determined asa0 = 0.4519 nm by suitable extrapolationof the variation ofa as a function of the Nelson–Relay parameter[12] (Fig. 3). A similar calculation based on the variation ofa asa function of sin2θ yields a value ofa0 = 0.4432 nm. However,the precise lattice parameter calculated fromFig. 3(a0 = 0.4519)is considered more appropriate because inadequate number ofpeaks at the back reflection region (2θ > 90◦) compelled us toconsider peaks in the low-2θ region in our extrapolation exercise.

Accordingly, the nearest neighbor distance (2r, wherer is theatomic radius) and volume per atom (Va) in the fcc unit cell arecalculated to be 0.3203 nm and 0.0231 nm3/atom, respectively.In comparison, the respective values of 2r andVa for hcp zirco-nium are 0.3179 nm and 0.0233 nm3/atom, and for hcp titaniumare 0.2896 nm and 0.0177 nm3/atom[12]. Thus, the hcp to fccphase transformation seems to be associated with approximately12% increase in (Va).

Significant reduction indc and increase inVa (or molar vol-ume Vm) may lead to a substantial increase in Gibbs energy(�G) in ultrafine crystals due to the Gibbs–Thompson or cap-illary effect: �G = 4γVm/dc, whereγ is the interfacial energy[14]. This increase in�G may eventually cause a high orderof lattice instability and compel the hcp titanium and zirco-nium to adopt a new polymorphic state with greater topologicalo ack-i that

Fc

ium disappear by 10 h of milling due to mutual dissolund formation of an fcc binary alloy. Earlier, similar structu

ransformation in elemental zirconium and titanium was notnly beyond 20–30 h of high-energy mechanical attrition[9,10].eyond 10 h, further milling causes significant broadening

eduction in intensity of the peaks in the XRD profile (Fig. 1).he milling product (beyond 10 h) is single phase fcc withost intense peak lying in between the earlier reported (1eaks of the fcc-zirconium and fcc-titanium[9,10]. The addi

ional peaks in the XRD profile are due to ZrO2 picked up fromhe milling media since 15 h of milling. Grain size (dc) of thenal milled products at 20 h is calculated to be 7 nm withelp of peak broadening analysis after careful separation overlapping peaks, and elimination of the instrumental erroer[12,13].

Fig. 2shows the XRD pattern of the final (20 h) milled prodf the alloy. Besides the fcc Ti50Zr50 alloy peaks, the remai

ig. 1. XRD pattern showing evolution of fcc Ti50Zr50 solid solution by mechancal alloying at different stages of milling.

r surface density of atoms without altering the crystal png density (=74%). In this regard, it is earlier reported

ig. 3. Variation ofa as a function of the Nelson–Relay parameter (cos2θ/sinθ +os2θ/θ).

112 S. Bera, I. Manna / Materials Science and Engineering A 417 (2006) 110–113

continued reduction indc below 10 nm leads to a substantialincrease in negative (from core to periphery) hydrostatic pres-sure (−Ph) on the grain boundary of nanocrystalline niobiumthat ultimately causes a bcc to fcc polymorphic change in thelatter [6–8]. In the present study the reduction indc results anincrease in excess free volume per atom (�VF), which devel-opsPh at the grain boundary. As stated by Chattopadhyay et al.[7,15] �VF is related todc as

�VF = (dc − ∆/2)2 − d2c

d2c

, (1)

where the grain boundary width,∆ = 1 nm assumed for the grainsof nanometric dimension. The hydrostatic pressure (Ph) as perthe isothermal equation of state for grain boundary, as adoptedby Fecht[16] from the original formulation of Rose et al.[17]is

Ph = −3B0[(V/V0)1/3 − 1]

(V/V0)2/3

× exp(−a∗)(1 − 0.15a∗ + 0.05a∗2), (2)

whereB0 is the equilibrium bulk modulus andV0 is the equi-librium volume of the atom with equilibrium Wiegner–Seitzcell radius (rWSE), and a* is expressed asa∗ = rWS−rWSE

l,

whererWS is the real Wiegner–Seitz cell radius of the atomwith a volumeV characteristic the length scale l of the givenm :VZa WithtT .T�

i� ngei turec5 ert sidess inG sone uet poly-mh

f sistso ize isa 1 1 1)o parew RDa sente –Zrb

hicc eale

Fig. 4. Nanocrystalline grain formation of Ti50Zr50 milling for 20 h.

Fig. 5. XRD pattern of Ti50Zr50 milled powder after annealing at 1000◦C.

at 1000◦C for 2 h in vacuum and slow cooled. Besides investi-gating the reversibility of the transformation, heating to 1000◦Cis also likely to dissociate hydride, nitride or oxide type of com-pounds of Ti or Zr, if any formed during milling.Fig. 5shows thatannealing at 1000◦C leads to re-appearance of the hcp-titaniumpeaks indicating that the hcp–fcc polymorphic change, unlikethat observed in elemental Zr[9] and Ti[10], is reversible.

It is relevant to point out that EDS analysis at several ran-domly chosen regions of the milled samples at intermediate andfinal stages shows no significant or detectable presence of proba-ble impurities like W, Co, Hf, Fe and O2 retained from the initialpowder or picked up during milling. Thus, the present hcp to fcctransformation of Ti50Zr50 cannot be attributed to the influenceof impurity. This is in accordance with our earlier observationwith regards to the influence of impurity on a similar polymor-phic change in elemental Nb[6–8], Zr [9] and Ti[10].

4. Conclusion

Planetary ball milling of Ti50Zr50 elemental powder blendleads to mutual dissolution and formation of a single-phase fccsolid solution. On heating to 1000◦C under vacuum, the equi-librium hcp solid solution reappears. This reversible hcp to fccpolymorphic transformation is associated with significant grainrefinement, lattice expansion and about 12% increase in volumep bilityd rainr tticee -core

aterial. Now�VF is related toV0 and V by the relationF = [(V − V0)/V0]. The values ofB0, rWSE and l for Ti andr are obtained from Ref.[17]. B0, rWSE and l of the alloyre calculated using the rule of mixture (weighted mean).

he help of Eq.(1), �VF is calculated to be 0.14 fordc = 7 nm.he corresponding−Ph calculated from Eq.(2) is 5.32 GPahis combined effect of significant increase in−Ph andVa [=Va,fcc− Va,hcp= (0.0231− 0.0205) = 0.0026 nm3/atom]

s equivalent to an elastic strain energy increment of−PhVa = 1.38× 10−20 J/atom. The theoretically predicted cha

n enthalpy in the event of a hcp to fcc crystal struchange in titanium and zirconium are 5.5× 10−21 J/atom16 and.6× 10−21 J/atom, respectively[18]. These values are low

han the strain energy value obtained in this study. Betrain energy, reduction indc causes significant increaseibbs energy (�G) of the system due to the Gibbs–Thompffect[14]. Thus, this structural/thermodynamic instability d

o strain and grain size reduction may ultimately cause aorphic change from hcp to fcc phase that has a lower�G origher structural stability.

Fig. 4exhibits the HRTEM images of Ti50Zr50 sample milledor 20 h. It is evident from the image that the sample conf several nanocryatalline grains and the average grain sbout 5–10 nm. Some of these crystallites seem to have (rientation with an interplanar spacing of 0.29 that comell with the same (=0.27 nm) obtained by corresponding Xnalysis (Fig. 2). It may be mentioned that no fcc phase is preither in elemental Ti or Zr (as an allotrope), or in the Tiinary equilibrium diagram.

In order to examine the reversibility of the polymorphange, selected as-milled sample, were isothermally ann

d

er atom. Perhaps, the thermodynamic and structural instaue to grain size reduction (capillary effect), high strain/state deformation and elastic strain energy (product of laxpansion and negative hydrostatic pressure from grain

S. Bera, I. Manna / Materials Science and Engineering A 417 (2006) 110–113 113

to surface/boundary) is responsible for this hitherto unknownhcp–fcc polymorphic change in Ti50Zr50 binary alloy.

Acknowledgments

Partial financial support from the CSIR (Grant No.: 70 (0048)03-EMRII) and MNES (Grant No.: 103/08/2001-NT) is grate-fully acknowledged. Useful technical discussion with Profs. F.Banhart, H.J. Fecht and S. Bhattacharyya, is deeply appreciated.

References

[1] H. Gleiter, Prog. Mater. Sci. 33 (1989) 223.[2] R.W. Siegel, in: R.W. Cahn, P. Haasen, E.J. Kramer (Eds.), Processing

of Metals and Alloys, vol. 15, VCH Publishers, Weinheim, 1991, 583pp.

[3] C.C. Koch, in: R.W. Cahn, P. Haasen, E.J. Kramer (Eds.), Processing ofMetals and Alloys, vol. 15, VCH Publishers, Weinheim, 1991, 193 pp.

[4] H.-J. Fecht, in: A. Edelstein, R. Camarata (Eds.), Nanomaterials: Syn-thesis, Properties and Applications, IOP, Philadelphia, 1996, 89 pp.

[5] P.P. Chattopadhyay, R.N.R. Gannabattula, S.K. Pabi, I. Manna, ScriptaMater. 45 (2001) 1191.

[6] P.P. Chattopadhyay, S.K. Pabi, I. Manna, J. Appl. Phys. 86 (1999)5912.

[7] P.P. Chattopadhyay, P.M.G. Nambissan, S.K. Pabi, I. Manna, Phys. Rev.B 63 (2001) 054107.

[8] P.P. Chattopadhyay, S.K. Pabi, I. Manna, Mater. Sci. Eng. A 304–306(2001) 424.

[9] I. Manna, P.P. Chattopadhyay, F. Banhart, H.-J. Fecht, Appl. Phys. Lett.81 (2002) 4136.

[10] I. Manna, P.P. Chattopadhyay, P. Nandi, F. Banhart, H.-J. Fecht, J. Appl.Phys. 93 (2003) 1520.

[11] E. Bonetti, G. Cocco, S. Enzo, G. Valdre, Mater. Sci. Technol. 6 (1990)1258.

[12] B.D. Cullity, Elements of X-Ray Diffraction, Addition-Wesley, Reading,MA, 1995.

[13] T.H. Keijser, I.L. Langford, E.J. Mittemeijer, B.P. Vogel, J. Appl. Crys-tallogr. 15 (1982) 308.

[14] R.B. Schwarz, C.C. Koch, Appl. Phys. Lett. 49 (1986) 146.[15] P.P. Chattopadhyay, I. Manna, Key Eng. Mater. 227 (2002) 87.[16] H.-J. Fecht, Acta Metall. Mater. 38 (1990) 1927.[17] J.H. Rose, J.R. Smith, F. Guinea, J. Ferrante, Phys. Rev. B 29 (1984)

2963.[18] L. Kaufman, in: P.S. Rudman, V. Stringer, R.I. Jaffe (Eds.),

Phase Stability in Metals and Alloys, McGraw-Hill, London, 1966,p. 125.