Correlation between microstructural evolution during high-pressuretorsion and isothermal heat treatment of amorphousAl85Gd8Ni5Co2 alloy

Peter Henits and Adam Revesza)

Department of Materials Physics, Eotvos University, Budapest, H-1518 Hungary

Erhard SchaflerPhysics of Nanostructured Materials, Faculty of Physics, University of Vienna,A-1090 Vienna, Austria

Peter J. SzaboDepartment of Materials Science and Engineering, University of Technology and Economy,Budapest, H-1111 Hungary

Janos L. LabarResearch Institute for Technical Physics and Material Science, Hungarian Academy of Sciences,Budapest, H-1121 Hungary

Lajos K. VargaResearch Institute for Solid State Physics and Optics, Hungarian Academy of Sciences,Budapest, H-1525 Hungary

Zsolt KovacsDepartment of Materials Physics, Eotvos University, Budapest, H-1518 Hungary

(Received 21 December 2009; accepted 23 February 2010)

Al85Gd8Ni5Co2 metallic glass was subjected to partial devitrification by high-pressuretorsion, continuous heat treatment, and isothermal annealing. The fully amorphous alloyexhibits a well-defined transition in its first devitrification product during isothermal heattreatments from tm þ a-Al phase mixture to primary a-Al by increasing the annealingtemperature above 555 K. This thermal sensitivity predestinates the composition toidentify the controversial thermal contribution of the plastic deformation in metallicglasses. Thermal stability and structure of the partially devitrified samples weresystematically analyzed and compared by calorimetry, x-ray diffraction, and electronmicroscopy. It seems that the effect of severe deformation cannot be singled out by asimple isothermal heat treatment; i.e., high-pressure torsion acts as a spectrum of heattreatments performed at different annealing temperatures.

I. INTRODUCTION

Amorphous Al-based aluminum–transition-metal–rare-earth alloys (Al–TM–RE) have successfully beensynthesized over a wide composition range in the lastdecades.1 Special compositions of these metallic glassescan produce primary precipitation of a-Al in heat treat-ments, where the combination of Al nanocrystals and theresidual amorphous matrix ensures superior mechanicalproperties.1,2

The ease of formation of crystalline phase from anamorphous precursor is primarily determined by thenucleation barrier and subsequent kinetics.3 The domi-nance of a-Al formation in the majority of the Al-basedalloys indicates that the combined nucleation and growthprocess for a-Al formation is the most competitive for

nearly all compositions. However, thermodynamic cal-culations for the Al–Gd–Ni system reveal that the a-Alphase exhibits lower Gibbs free energy than the com-pound phase; therefore, the low interfacial energybetween a-Al embryos and the amorphous matrix isessential in reducing the nucleation barrier.4 The lowinterfacial energy relies on the fact that the amorphousmatrix coordination number is fairly close to that of theface-centered cubic (fcc) structure obtained by atomicstructure studies using synchrotron and neutron diffrac-tion.5 After the nucleation of a-Al embryos, the slowlydiffusing large atoms, such as Ce, Gd, or Nd that areinsoluble in fcc-Al, block the rapid growth of primarycrystals and set up a concentration gradient that reducesthe thermodynamic driving force for intermetallic phasenucleation at the primary nanocrystallization front.6

Other elemental additions (Fe, Ni, Co, Cu) are needed tokeep the residual matrix in the amorphous compositionrange at the temperature of annealing.4,6

a)Address all correspondence to this author.e-mail: reveszadam@ludens.elte.hu

DOI: 10.1557/JMR.2010.0164

J. Mater. Res., Vol. 25, No. 7, Jul 2010 © 2010 Materials Research Society1388

In addition to heat-induced crystallization, deformation-induced crystallization is also able to promote a-Alnanodispersion in some Al-based metallic glasses. Recentexperimental reviews have shown that deformation-induced crystallization is a common phenomena of metal-lic glasses.7–15 To date no consensus has been reachedconcerning the microscopic mechanism of mechanicallyinduced crystallization. The diffusion profile around Alprecipitates in the Al90Fe5Gd5 amorphous alloy afterdeformation is nearly identical to that of obtained afterisothermal heat treatments, suggesting significant tempera-ture rise as a consequence of deformation.16 Additionally,the observed a-Al nanocrystals induced by deformationusually locates in the vicinity of shear bands; therefore,one could assume that the temperature rise and the atomicmobility in the shear band trigger the nucleation.17 How-ever, as Demetriou and Johnson have pointed out thetemperature rise solely cannot account for the crystalliza-tion, because the driving force for nucleation decreaseswith increasing temperature.18 Interestingly, not just theprimary precipitation of a-Al can be facilitated by defor-mation, but simultaneous crystallization of intermetallicphases and a-Al can also be observed for instance duringmechanical milling.19 It seems that in this case the mechan-ical nanocrystallization is strongly connected to growingconditions of crystalline phases and to the evolving concen-tration gradients. During the intense deformation the con-tinuous breakup and mixing results in the elimination of thediffusion layers and concentration gradients, which usuallyinhibits the nucleation of intermetallic phases. Recently,high-pressure torsion (HPT) originally invented for produc-ing highly dense, bulk ultrafine-grained materials20 hassuccessfully been applied to produce massive amorphousnanocomposites from melt-quenched fully amorphousribbons by consolidation at room temperature.21–26

The temperature rise, the stress state, and the compo-sition of the alloy have crucial roles in deformation-induced crystallization. Just a few attempts have beenmade so far to classify these effects.27,28 In this work, aspecial composition has been selected to elucidate therole of temperature rise during deformation-inducedcrystallization. With increasing isothermal annealingtemperatures, amorphous Al85Gd8Ni5Co2 alloy presentsa clear transition in its devitrification products,29 whichcan be used as an indicator for the estimation of temper-ature rise during deformation. In our experiments, theamorphous sample has been subjected to extreme sheardeformation applying HPT.

II. EXPERIMENTAL

Ingot of Al85Gd8Ni5Co2 was synthesized by inductionmelting of a mixture of high-purity (99.9%) Al, Gd, Ni,and Co metals (Sigma Aldrich, Milwaukee, WI). Fullyamorphous ribbon sample was obtained using a single-

roller melt-spinning technique in inert atmosphere. Someportion of the as-quenched ribbon was cut into smallpieces (flakes) and then placed between anvils of theHPT device resulting in several porosity free disks witha radius of R ¼ 5 mm and thickness of L ¼ 0.7 mm. Thetorsion straining was performed under 0.2 rev/min for N¼ 5 whole turns. Shear strain for torsion deformation at aradius r can be represented by:

g ¼ 2rpL

: ð1Þ

At the perimeter of our sample g(R) � 220. After thedeformation, the disk was fragmented into small piecesof about 1 mm3, which then were sorted into two groups,according to their distance measured from the center ofthe disk [sector A (r¼ 0 to 2.5 mm) and sector B (r¼ 2.5to 5 mm)].

The crystalline phase analysis was carried out ondiffractograms obtained by a Philips (PW1130) x-raygenerator (Almelo, The Netherlands) with a Guinierchamber setup. The chamber has a diameter of 100 mm,and the patterns were recorded on image plates.

For transmission electron microscopy (TEM), partsof the HPT disks were prepared by twin-jet electro-polishing (Tenupol-5) in a solution of water, ethanol,butoxyethanol and 1/78 mL perchloric acid (Struers A2Electrolyte, Willing, Germany). The thin part of the sam-ple around the perforation was investigated using a JEOLFX 2010 TEM (Tokyo, Japan), operated at 200 keV.Selected-area electron-diffraction patterns (SAED) andcorresponding dark-field (DF) images were recorded on“plane view” section of the HPT disk.

Morphology studies were performed using a Philips XL30 scanning electron microscopy (SEM) in backscatteredelectron (BSE) mode. The compositional changes of thesurface related directly to contrast differences in the BSEimage were revealed and quantitatively determined byenergy dispersive x-ray (EDX) analysis with a relativeaccuracy of 3%. For these cross-sectional microstructuralinvestigations, the disk was mechanically polished.

Continuous thermal scans were carried out in a Perkin-Elmer power-compensated differential scanning calorime-ter (DSC; Waltham, MA) on the as-quenched ribbon andHPT disks at 2.5, 5, 10, 20, and 40 K/min, using purifiedargon atmosphere. The temperature and enthalpy werecalibrated by melting pure Al and In. From the shift ofDSC peaks with increasing heating rate, the apparentenergy of crystallization process was determined usingthe Kissinger analysis.30

III. RESULTS

A. As-quenched ribbon

As seen in Fig. 1(a), linear heating DSC scan (ST) of theas-quenched Al85Gd8Ni5Co2 ribbon shows a significant

P. Henits et al.: Correlation between microstructural evolution during HPTand isothermal heat treatment of amorphous Al85Gd8Ni5Co2 alloy

J. Mater. Res., Vol. 25, No. 7, Jul 2010 1389

glass transition (Tg ¼ 541 K) followed by three exother-mic peaks (T1 ¼ 563 K, T2 ¼ 607 K, T3 ¼ 672 K, DH1 ¼50 � 5 J/g, DH2 ¼ 18 � 2 J/g, and DH3 ¼ 59 � 6 J/g).The width of the supercooled liquid range (T1 � Tg)is 22 K, and it strongly depends on the electronegativityof the RE component.31 To determine the structuralchanges associated with each crystallization stage, XRDmeasurements were carried out after linear heatings upto ST1 ¼ 580 K, ST2 ¼ 640 K, and ST3 ¼ 780 K. TheXRD pattern of the as-quenched ribbon reveals a broadhalo, typical for fully amorphous alloys, while at ST1

some definite peaks emerge corresponding to the dif-fraction peaks of fcc-Al [Fig. 1(b)]. During the secondcrystallization step, sharpening of the Bragg peaks takesplace, indicating a coarsening process and the relaxationof the matrix at ST2.32 Finally, Al3Gd (hexagonal,Ni3Sn-type, a ¼ 0.623 nm, c ¼ 0.460 nm) and t1(Al19Ni5Gd3, orthorhombic, a ¼ 0.409 nm, b ¼ 1.599 nm,c ¼ 2.709 nm) phases appear simultaneously on com-pletion of the devitrification (ST3). The activationenergies are Q1 ¼ 5.5 � 0.5 eV, Q2 ¼ 3.8 � 0.5 eV, and

Q3 ¼ 2.2 � 0.5 eV for the T1, T2, and T3 transformations,respectively. The activation energy of the first crystal-lization step is quite high compared with that of the amor-phous Al85Ce8Ni5Co2 (3.9 � 0.5 eV)33 and Al85Y8Ni5Co2(3.6 � 0.5 eV)34 alloys.

B. Isothermal annealing

To compare the crystallization products under isother-mal conditions with the microstructure formed after con-tinuous heatings, the as-quenched Al85Gd8Ni5Co2 alloywas subjected to isothermal heat treatments below andnear Tg in the Tann ¼ 535 to 560 K range for 25 min[these temperatures are indicated by ticks in Fig. 1(a)].The corresponding XRD patterns in Fig. 2(a) obtained

after isothermal annealing show important changes in thestructure. After heat treatments below Tann ¼ 555 K (LT:low-temperature heat treatment) the Bragg peaks of a-Al

FIG. 2. (a) XRD patterns obtained after isothermal heat treatments

and (b) the corresponding subsequent linear DSC scans.

FIG. 1. DSC scan of the as-quenched Al85Gd8Ni5Co2 ribbon obtained

at 40 K/min (a) with the XRD patterns corresponding to the amor-

phous alloy and (b) continuous heatings up to ST1, ST2, and ST3 tem-

peratures. Ticks denote temperatures Tann ¼ 535 to 560 K, where

isothermal annealing has been carried out.

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and tm phases appear simultaneously. The tm phase has acomposition of Al23Ni6Gd4 and a monoclinic crystalstructure (a ¼ 1.586 nm, b ¼ 0.408 nm, c ¼ 1.829 nm,b ¼ 113.01�).35 On the contrary, heat treatments aboveTann ¼ 555 K (HT: high-temperature heat treatment)results in the precipitation of primary Al, without anyother detectable phases, in accordance with Ref. 29.

Figure 2(b) summarizes the effect of isothermal heattreatment on the thermal stability of the partially crystal-lized alloy. As seen, continuous scans following LTannealing (e.g., at Tann ¼ 535 K) reveals only a slighteffect on the thermal behavior. However, as the isother-mal temperature increases, drastic change occurs in theexothermic peak shapes. After isothermal annealing atTann ¼ 540 K the intensity of T1 peak decreases, while ittotally diminishes after Tann ¼ 550 K. Additionally, theT3 peak splits into two peaks (T3 and T4), and with in-creasing annealing temperature the separation betweenT3 and T4 peaks decreases and at Tann ¼ 560 K (HT) itcompletely vanishes. Interestingly the intensity of T3peak increases, while that of the T4 peak decreases withannealing temperature.

As the devitrification products of the Al85Gd8Ni5Co2alloy change drastically with the applied isothermalannealing temperature around Tg, this composition is aperfect candidate to explain the controversial interpreta-tion of the thermal effects occurring during severe plasticdeformation (e.g., high-pressure torsion) of metallicglasses.17

C. High-pressure torsion

Figure 3(a) presents the effect of torsional straining onthe structure of the amorphous Al85Gd8Ni5Co2 alloy. TheXRD patterns measured in the center (sector A) and atthe edge of the HPT disk (sector B) exhibit faint Braggpeaks, superimposed on the amorphous halo. Thesepeaks indicate the precipitation of fcc-Al in the deformeddisk. The slightly larger intensity of Al peaks for sector Bis indicative of some deformation dependence of thetransformed nanocrystalline fraction.

The DSC thermogram of sector A shows the same T1,T2, and T3 exothermic peaks, similar to the as-quenchedribbon [Fig. 3(b)]. With increasing deformation, theintensity of T1 peak decreases significantly and the T3peak splits into T3 and T4 reminiscent of the LT heattreatment (sector B).

D. Devitrification stages and products

To identify the microstructural changes associatedwith the T1, T2, and T3 transformations correspondingto the HPT (sector B), LT and HT samples, linearheatings were carried out up to HPT2 ¼ 625 K, HPT3 ¼690 K, HPT4 ¼ 780 K [indicated in Fig. 3(b)]; LT2 ¼615 K, LT3 ¼ 680 K, LT4 ¼ 780 K, and HT2 ¼ 630 K,

HT3 ¼ 700 K [Fig. 2(b)], respectively. The XRD patternof the HPT2 state presents the amorphous halo and intensediffraction peaks of fcc-Al accompanied by some faintBragg peaks related to tm and Al11Gd3 [Fig. 4(a)]. Afterheating up to HPT3 the amorphous halo decomposes andthe mixture of Al11Gd3, Al9Co2, and Al3Ni phases forms,which finally transforms to a-Al, Al3Gd and t1 (HPT4).Similarly, continuous heating up to LT2 results in a mixedstate, where the tm and Al11Gd3 phases are present simul-taneously. During T3, tm disappears and Al11Gd3, Al9Co2and Al3Ni phases can be identified (LT3) in the XRDpattern. Finally, these phases transform to Al3Gd and t1phases at LT4 [Fig. 4(b)]. Continuous heating up to HT2 ¼640 K results in Bragg-peak sharpening, corresponding tothe growth of a-Al particles during T2 [Fig. 4(c)]. Subse-quently, the amorphous halo decomposes and the mixtureof a-Al, Al3Gd, and t1 emerges in a single-step transfor-mation during T3.

FIG. 3. (a) XRD patterns obtained after HPT corresponding to sectors

A an B of the disk and (b) subsequent linear DSC scans.

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To summarize the similarities and differences betweenthe ST, HT, LT, and HPT devitrification routes, the dif-ferent metastable states are depicted in Fig. 5 accordingto their enthalpy content. Note that the enthalpy contentof a given state is set to the base level of the correspond-ing box, which is drawn with thick line. As seen, both theLT and the HPT treatments exhibit a four-stage devitrifi-cation route; ST and HT exhibit only a three-stage crys-tallization process. Comparing the different treatments, itis seen that the HPT state exhibits substantially higherenthalpy content than the ST1, LT, and HT states.

IV. DISCUSSION

A. Thermal-induced devitrification

Depending on the isothermal temperature, the firstdevitrification step of the Al85Gd8Ni5Co2 metallic glasscan be summarized as

Process P : amorphous ! a-Alþ amorphous�

ðTann ¼ 560 KÞ ;

Process E : amorphous ! tmþða-AlÞðTann¼ 542.5 KÞ :

In the range between these two annealing temperatures,both process P and E operate [see Fig. 2(a)]. Examining

the heat flow signal ð _HÞ recorded during isothermalannealing [Fig. 6(a)], it is clear that the first devitrifica-tion step has been fully completed only aboveTann ¼ 542.5 K and apparently processes P and E ceaseduring the same single isothermal exothermic peak.

Using the peak positions of the isotherms, an isothermactivation energy Qiso ¼ 4.6 eV has been obtained [seethe inset of Fig. 6(a)], which represents approximatelythe activation energy of process E in our case.To compare the crystallization kinetics of the first

devitrification step, the transformed volume fractioncurves, f(t) ¼ R

_Hdt/DHiso have been calculated fordifferent Tann isothermal temperatures. Here DHiso

denotes the area of the first isothermal peak. The f(t)curves of the 542.5 and 545 K isotherms can ade-quately be fitted by JMAK kinetics fJMA(t) ¼ 1 �exp{�[(t � t0)/t0]

n}36 with exponent of n � 3.2, is in

FIG. 4. Devitrification routes and crystalline phases obtained after (a) HPT, (b) LT, and (c) HT.

FIG. 5. Crystallization stages attained by linear scan and devitrifica-

tion routes obtained after HPT, LT, and HT with the different metasta-

ble states and corresponding enthalpy contents.

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accordance with literature data obtained on Al–Y–Ni–Co metallic glass, which corresponds to decreasingnucleation rate under interface-limited growth of nucleiat eutecticlike phase transformation.37

Because the shape of the f(t) signal presents a slowlydecaying long tail at higher isothermal temperatures theJMAK fit is not satisfactory. Such behavior correspond-ing to sluggish kinetics for Tann ¼ 555 K and Tann ¼560 K is caused by the soft impingement of the solute-rich zones around primary forming a-Al crystals.38,39 Totake into account the effect of soft impingement, a moregeneral function

fGMðtÞ ¼ 1� 1þ ½ðt� t0Þ=t0�nw

� �w

; ð2Þ

was considered based on a general model (GM), whichdescribes the kinetics as an initially increasing transforma-

tion rate finishing in a slowly decaying tail.40 It is notedthat fGM(t) ! fJMA(t) in the w ! 1 limit. Evidently, thegeneral model can be fitted well on the f(t) curve corre-sponding to Tann ¼ 560 K in the whole transformationinterval, whereas the JMAK fit is not satisfactory at thestart of the event [compared in Fig. 6(b)]. The variation ofthe soft-impingement w parameter as a function of temper-ature indicates that, indeed a transition occurs in the kinet-ics by changing Tann (Fig. 7). For LT annealing w > 1,whereas the role of solute gradient is small; on the otherhand, solute gradients build up in the amorphous matrix,and the soft impingement becomes the main rate control-ling process for HT annealing (w � 0).

Comparing the temperature dependence of the crystal-lization processes, the a-Al formation, albeit it startslater, has larger activation energy (Q1) than that of thetm phase (Qiso). These results indicate that the a-Al for-mation should become the main devitrification route athigher temperatures where the transition between the twoprocesses occurs at about 550 to 555 K. Both processesstart with nucleation and growth of the correspondingphase (i.e., a-Al or tm), which lead to concentrationgradients around the nuclei. These concentration gradi-ents are substantially larger for the a-Al, because itsformation requires larger amount of solute rejection fromthe nucleus due to the larger concentration differencebetween the matrix and the crystal. As a consequence,the formation of the tm phase at lower temperatures ispreferred. For higher temperatures where the growth rateof the primary a-Al increases, the nucleation rate of thetm phase is reduced because of the sensitivity of theintermetallic phases to the concentration gradients.41 Incontrast, if the devitrification is dominated by the tmformation (Tann ¼ 540 to 550 K), a-Al crystals nucle-ation is preferred in Al-rich zones. This difference inthe two process leads to a change in the composition ofthe remaining amorphous matrix as a function of theannealing temperature. As the amount of a-Al increasesin the devitrification product, the composition of theremaining matrix shifts toward higher Gd contents,

FIG. 6. (a) Isothermal DSC signals measured at different Tann and the

corresponding Kissinger plot in the inset. (b) The effectiveness of

fitting is represented for JMAK and generalized model (GM), using

the Tann ¼ 560 K curve.

FIG. 7. The variation of w parameter as the function of annealing

temperature.

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which results in the increase of the corresponding crys-tallization temperature [see T3 peak in Fig. 2(b)].

To support this scenario we separate the nucleation ofthe tm phase from its consequent growth in a two-stagecourse by applying a short isothermal preanneal (2 and3 min) in the tm forming temperature range (Tann¼ 550 K)and a subsequent isotherm treatment (23 and 22 min) inthe a-Al forming range (Tann ¼ 560 K). XRD patterns inFig. 8 indicate no detectable crystallization after the pre-treatments. The fully treated sample following the 2 minpreannealing [2 min (550 K) þ 23 min (560 K)] containsonly a-Al, similarly to the HT treatment [25 min (560 K)].Notwithstanding, 3 min pretreatment [3 min (550 K)þ 22min (560 K)] results in the formation of both a-Al and tmphases, in line with an LT annealing [25 min (550 K)].Consequently, the tm phase grows but does not nucleate at

560 K due to the interaction with the significant solutegradients (w > 1), in accordance with the concentrationgradient sensitivity model41 for nucleation of intermetallicphases.

B. Severe plastic deformation induceddevitrification

Devitrification products of the Al85Gd8Ni5Co2 amor-phous alloy obtained after severe plastic deformation andheat treatments (linear heating up to ST1 and HT isother-mal treatments) exhibit similarities, for example the pri-mary precipitation of a-Al [compare Figs. 1(b), 2(a), and3(a)]. The first transformation peak (T1) is partially pres-ent in the subsequent linear DSC scan of the outer part ofthe HPT disk [see Figs. 2(b) and 3(b)], indicating thepresence of a small amount of untransformed volumefraction after the heavy shear deformation. At the sametime, the shift of the T3 peak to lower temperatures andthe appearance of the T4 peak [Fig. 3(b)] correspond tothe nucleation of a large amount of tm during the HPTdeformation alike in LT treatment, although this phase isnot visible in the XRD pattern [Fig. 3(a)]. The similarphase mixtures of the HPT2 and HPT3 states comparedwith the LT2 and LT3 states, respectively, also supportthis presumption (Fig. 4). Conclusively, the thermal andmicrostructural behavior of the HPT sample indicatesthat the effect of severe plastic deformation cannot besingled out by a single isothermal heat treatment.Because calorimetry is sensitive to small changes in thedevitrification product, the shape of the thermograms canroughly provide the main Tann temperature component ofHPT as 550 to 555 K; however, a minor contributionoriginates from Tann � 540 K [see Figs. 2(b) and 3(b)].To confirm the formation of tm phase, which was not

detected by XRD [Fig. 3(a)], electron microscope experi-ment was carried out on sector B of the as-deformed HPTdisk. As seen in the SAED pattern, spare diffraction spotsof the tm phase are present along with fcc-Al reflectionssuperimposed on the amorphous halo [Fig. 9(a)]. Byexploiting good statistics of the SAED pattern, the intensehalo and the slowly undulating background were filteredout by eliminating reflections wider than 2.2 nm�1. Thetypical width of the diffraction spots and the amorphoushalo is about 1.5 and 10 nm�1, respectively. As a result,faint and quasi-continuous diffraction rings of solely thefcc-Al can be resolved in the inset of Fig. 9(b). Theabsence of the tm phase in the one-dimensional patternobtained by integrating the filtered image around the beamcenter [Fig. 9(b)] is due to small contribution of the indi-vidual spare diffraction spots, in accordance with the cor-responding XRD pattern in Fig. 3(a). In the DF imageobtained by selected part of (111) fcc-Al ring of theelectron diffraction pattern, a homogeneous distributionof nanocrystals (4–6 nm) can be observed [Fig. 9(c)].

FIG. 8. XRD patterns obtained after two-stage heat treatments.

(a) Isothermal annealing have been performed at 550 K for 2 min

followed by 23 min long heat treatment at 560 K and (b) at 550 K for

3 min followed by 22 min long heat treatment at 560 K.

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Besides, spare larger crystals with size of 20–100 nm arealso visible.

Interestingly, elongated blocks with much larger size(5–40 mm) homogenously dispersed in the matrix can beobserved in SEM BSE images taken on the cross sectionof the outer part of the HPT disk (Fig. 10). Domains insimilar size range were also observed in heavilydeformed Cu-based metallic glasses.42 QuantitativeEDX analysis taken at several points confirmed that thecomposition of the contrastless matrix is equal to that ofthe as-quenched alloy; the average composition of thelight and dark gray blocks is about Al71Gd13Ni11Co5

and Al93Ni5Co2, respectively, which is close to the com-position of the phases (tm and a-Al) observed in theSAED pattern [Fig. 9(a)]. The contribution of thesephases is hardly visible in the XRD pattern [Fig. 3(a)],since the volume fraction of the domains is small. On aside note we observed the presence of similar blocks inanother disk processed by smaller deformation (N ¼ 1).The origin of these large blocks is not unambiguous, butprobably they form during the initial stage of the HPTprocess when the global temperature of the sample isconsiderably lower than Tg. Nevertheless, the local tem-perature rise caused by the inhomogeneous deformationcan be significant,43 especially at the interface of thesmall pieces of the amorphous ribbon.

As the deformation continues during HPT process, theglobal temperature can reach Tg according to a thermo-plastic model involving thermal balance between theplastic work input and heat conduction, proposed byHobor et al.44 The presence of the tm and a-Alnanocrystals [Figs. 9(a) and 9(c)] indicates that thehomogeneous temperature increase is substantial duringthe HPT deformation, and it exceeds Tg as is apparentfrom the DSC analysis [Fig. 3(b)]. Nevertheless, thelarger spare nanocrystals (20–100 nm) cannot form onlyby diffusion in the 540 to 560 K temperature range;therefore, the effect of shear deformation also has to betaken into account. Above Tg, the glass behaves like asupercooled viscous liquid; thus, the forming crystalsgrow in the sheared liquid. Above a critical size, thesecrystals grow faster45 than in the case of an isothermalannealing, since deformation can easily distort thesolute-rich zones around the nuclei. As a consequence,the higher concentration gradients at the nucleus/matrixinterface promote the growth of the larger nanocrystals.

At the same time, concentration gradients evolve inthe matrix as well. Until these concentration gradientsare relaxed in the matrix, the nucleation of tm phase issuppressed. During shear deformation, the shape of theinitially spherical solute shells around the crystal nucleibecome ellipsoidal with aspect ratios of (g, 1, 1/g). In thedirection of the smallest axis the mixing of the soluteatom become enhanced, which decreases the relaxationtime of the concentration fluctuations by a factor of g2/3.This shorter relaxation time leads to a reduced relaxationtemperature in linear heating. The reduced relaxationtemperature (T2

*), which is equal to (T2 ¼ 607 K) in theabsence of shear deformation, can be calculated by theQ2 activation energy of the relaxation process as

T2� ¼ Q2=kB

ðQ2=kBT2Þ þ lnðg2=3Þ ; ð3Þ

which is about 538 to 548 K for sector B; kB is theBoltzmann constant. In reality, T2

* is somewhat higherbecause concentration gradients on length scale smallerthan an atom are unlikely present in the glass. Apparently,

FIG. 10. BSE image taken on the cross section of sector B of the HPT

disk.

FIG. 9. (a) SAED pattern taken on sector B of the HPT disk with the

diffraction spots of the a-Al and tm phases. (b) One-dimensional

diffraction pattern obtained after integration of the filtered SAED

image. Inset shows the filtered image with the fcc-Al rings.

(c) Corresponding dark field image.

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T2* is comparable to the temperature reached in HPT

deformation (550–555 K), indicating that the relaxationof concentration gradients has already started during theHPT process; however, while shear deformation persists,the gradients cannot be erased completely. Therefore,nucleation of tm phase is suppressed in the supercooledliquid region until the end of the deformation process. In asubsequent linear DSC scan, however, these concentrationgradients relax immediately above Tg at T2

* < T2 leadingto fast nucleation of the tm. Consequently, the experimen-tal results on the heat treated HPT samples (Figs. 4 and 5)support the presence of tm phase.

V. CONCLUSIONS

In this work, direct correlations have been foundbetween the structure and crystallization of isothermallyheat treated and high-pressure torsioned amorphousAl85Gd8Ni5Co2 alloy. The conclusions of this investiga-tion can be summarized:(1) Isothermal heat treatments result in different crys-

tallization sequences of amorphous Al85Gd8Ni5Co2 alloyas a function of the Tann annealing temperature. BelowTann ¼ 560 K a mixture of a-Al and tm phases can beobserved, whereas above this limit only a-Al forms in theresidual amorphous matrix. In accordance, the subsequentlinear heating treatments exhibit two completely differentcrystallization sequences. The change of the devitrifica-tion products is related to the different apparent activationenergy of the underlying phase formation processes.(2) Severe plastic deformation by high-pressure torsion

results in partial devitrification of the Al85Gd8Ni5Co2alloy. According to SEM and TEM investigations, nano-crystals (�5 nm) and blocks (�50 mm) of these crystalsappear in the matrix. The largest blocks initiate in the earlystage of the HPT when the deformation is inherently inho-mogeneous and the average temperature is well below Tg.On the other hand, the smallest crystals grow homoge-neously, which is reminiscent to the microstructureobtained by heat treatments in the supercooled liquid state.(3) Comparing x-ray diffraction patterns of heat

treated samples, the devitrification products of HPTdeformation are similar to those obtained by isothermalannealing, although the effect of severe plastic deforma-tion cannot be singled out by a simple heat treatment. Asa good approximation, HPT can be considered as a spec-trum of isothermal heat treatments performed at differentannealing temperatures.(4) Considering growing a-Al nanocrystals in a

sheared viscous liquid, the DSC spectrum of the HPTsample can qualitatively be explained by taking intoaccount the T2

* reduced relaxation temperature ofsheared solute concentration gradients. The immediatenucleation of the tm phase above Tg is possible becauseT2

* is comparable to Tg.

ACKNOWLEDGMENTS

The authors acknowledge the Erich Schmid Instituteof Material Science, Austria Academy of Science forproviding the HPT facility. We appreciate the support ofthe Hungarian Scientific Research Fund under GrantNos. 67893 and 67692. A.R. is indebted for the JanosBolyai Research Scholarship of the Hungarian Academyof Sciences. Zs.K. is grateful for the support of theMagyary Zoltan Fund and of the EEA Grants and Nor-way Grants.

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