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Journal of Alloys and Compounds 473 (2009) 71–78

Contents lists available at ScienceDirect

Journal of Alloys and Compounds

journa l homepage: www.e lsev ier .com/ locate / ja l l com

haracterization of amorphous and nanocrystallinei–Ni-based shape memory alloys

. Inaekyana,b, V. Brailovskia,∗, S. Prokoshkinb,

. Korotitskiyb, A. Glezerc

Ecole de technologie superieure, 1100 Notre-Dame Street West, Montreal, Quebec H3C 1K3, CanadaMoscow Institute of Steel and Alloys, Leninskiy Prospect 4, Moscow 119 049, RussiaG.V. Kurdyumov Institute of Physical Metallurgy, 9/23, 2nd Baumanskaya Street, Moscow 107005, Russia

r t i c l e i n f o

rticle history:eceived 19 March 2008ccepted 12 May 2008vailable online 24 June 2008

eywords:morphous materials

ntermetallicsanostructuresrystal growthalorimetry

a b s t r a c t

Ti–50.0 and Ti–50.26 at.%Ni shape memory alloys are subjected to cold rolling (CR) with true strain varyingfrom moderate (e = 0.3) to severe (e = 1.5–2). The Ti–50.0%Ni alloy was also subjected to high-pressure tor-sion (HPT) with a true strain of e = 6.16. A melt-spun fully amorphous ribbon of Ti–25 at.%Ni–25 at.%Cu alloywas used as a reference sample for DSC analysis. The study of the material structure and its thermal stabil-ity was performed using TEM and DSC. CR with true strains higher than 0.5 initiate grain refinement andaustenite amorphization: the higher the cold work strain, the higher the grain refinement and the degree ofamorphization. Parallel TEM/DSC analyses show that CR- and HPT-processed Ti–Ni alloys are strongly het-erogeneous materials containing amorphous matrix with embedded nanocrystals. The volume fraction ofthe amorphous phase in such materials cannot be evaluated by DSC because exothermal peaks measuredcontain inseparable contributions from both crystallization and grain-growth phenomena. Evaluating

this quantity therefore requires direct TEM observations in addition to the DSC. It is suggested that fora heterogeneous material, its apparent thermal stability depends on a trade-off between volume frac-tions of the amorphous phase and deformation-induced nanocrystals and activation energies for theirrespective transformations: for CR (e = 2), the Avrami exponent n = 1.25, which corresponds to anneal-ing by grain-growth-dominated mechanism; for HPT (e = 6.16), n = 2.5, which corresponds to annealingby nucleation-and-growth mechanism. For fully amorphous melt-spun Ti–25 at.%Ni–25 at.%Cu alloy, theAvrami exponent reaches 4, which corresponds to the nucleation-dominated annealing.

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mfatfo

. Introduction

The cold rolling (CR) of Ti–Ni shape memory alloys can be used toailor specific material microstructures, ranging from dislocation-ardened substructure to ultrafine-grained/nanocrystalline struc-ures, and all the way to partially amorphized structures whenold work strains vary from moderate (true strain e = 0.3–0.5) toevere (e = 1.5–2). Following an adequate post-deformation anneal-ng (PDA), a dislocation-hardened substructure can be transformednto a polygonized dislocation substructure; through the same

rocess, a mixed amorphous/nanocrystalline structure can beransformed into a dislocation-free nanocrystalline structure. Its shown that a nanocrystalline austenite with a grain size ofbout 50–100 nm confers about 30–50% higher functional proper-

∗ Corresponding author. Tel.: +1 514 396 8800/8594; fax: +1 514 396 8530.E-mail address: vbrailovski@mec.etsmtl.ca (V. Brailovski).

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925-8388/$ – see front matter © 2008 Elsevier B.V. All rights reserved.oi:10.1016/j.jallcom.2008.05.023

© 2008 Elsevier B.V. All rights reserved.

ies (maximum recovery stress, �maxr , and maximum completely

ecoverable strain, εmaxr,1 ) than does the dislocation substructure

ith the same size of subgrains [1].It is generally observed that when the cold work strain of initially

artensitic Ti–Ni alloys increases above e = 0.3, a reverse trans-ormation takes place first, followed by a gradual refinement ofustenite grains to a nanosize scale (e = 0.3–0.5), and that when arue strain increases from 0.5 to 2, it is followed by a steadily intensi-ying material amorphization (e = 2 corresponds to the highest levelf cold work ever achieved through the conventional flat rolling ofuch materials).

To evaluate the quantity of the amorphized material and totudy the grain nucleation-and-growth processes during post-

eformation annealing, two analyses are usually performed: directbservations in the transmission electron microscope (TEM) andifferential scanning calorimetry (DSC). For coarse-grained mate-ials, it is the first technique that prevails because the release of therain boundary enthalpy is too small to be measured by DSC, while

72 K. Inaekyan et al. / Journal of Alloys and Compounds 473 (2009) 71–78

Table 1Ti–Ni and Ti–Ni–Cu alloys quenched from 800 ◦C and their transformation temperatures

Chemical composition Ti–Nix (x, at.%) Mf (◦C) Ms (◦C) As (◦C) Af (◦C) Manufacturer

50.0a 29 54 74 98 Central Research Institute of Materials, Russia50.26 43 58 77 95 Special Metals, USA25, 25 at.%Cub 56 59 62 69 Dr. A.V. Schelyakov, Moscow Engineering

Physics University, Russia

T ning

talliza

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2

w

asw

t

emperatures are measured by differential scanning calorimetry at a 10 ◦C/min scana Quenching from 700 ◦C.b The temperatures of B2 ↔ B19 transformation were determined after 1 min crys

or the homogeneous amorphous or fine-grained materials witharge grain interface-to-volume ratios, DSC becomes a convenientechnique for use in monitoring: (1) the crystallization of amor-hous alloys and (2) the grain growth in fine-grained materials.owever, the situation becomes more complicated when theaterial is heterogeneous and contains both nanocrystals and an

morphous matrix, in various proportions, depending on the coldork intensity. To better understand the real problems related to

he use of the DSC technique in studying the thermal stability ofeterogeneous compounds, parallel TEM and DSC studies are per-

ormed on severely deformed Ti–Ni and melt-spun Ti–Ni–Cu shapeemory alloys.

p(tn

Fig. 1. TEM images of the CR, e = 2 (a) and annealed for 1 h at 300 ◦C (b), 330 ◦C

rate.

tion at 700 ◦C.

. Materials and experimental techniques

This study involves an investigation of the Ti–Ni-based shape memory alloysith compositions listed in Table 1.

A quenched Ti–50.26 at.%Ni ∅1 mm wire was subjected to CR at room temper-ture. The maximum true plastic CR strain (e = ln(h0/h1), where h0 and h1 are thepecimen thicknesses before and after processing, respectively) was 2. The CR speedas 0.025 m/s, with no lubrication and pulling force applied.

Disk-shaped samples (initial diameter of 3 mm, thickness of 0.2 mm) ofhe Ti–50.0 at.%Ni alloy were subjected to high-pressure torsion (HPT): applied

ressure P = 4 GPa; number of revolutions N = 10. The maximum plastic straine = 6.16) imposed by the HPT was calculated according to the following equa-ion: e = ln(2�RN/h), where R and h are the sample initial radius and thick-ess.

(c), 350 ◦C (d), 400 ◦C (e) and HPT-processed, e = 6.16 (f) Ti–50.0%Ni alloy.

K. Inaekyan et al. / Journal of Alloys and Compounds 473 (2009) 71–78 73

%Ni a

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o2eb

Jts

3

aasta

Fa

Fig. 2. DSC thermograms of cold-rolled (e = 2) and annealed samples of Ti–50.26

A Ti–25 at.%Ni–25 at.%Cu (Ti50Ni25Cu25) ribbon was produced by the single-oller melt-spinning technique. Presynthesized Ti50Ni25Cu25 alloy was preparedrom the high purity metals by melting six times in argon arc furnace. The obtainedngots were re-molten in quartz crucibles under a purified helium atmosphere andjected onto the surface of a fast rotating copper wheel at the cooling rate of about06 K/s. The rapidly quenched Ti–Ni–Cu ribbon was of around 40 �m in thicknessnd about 1.5 mm in width.

A DSC analysis was performed using a PerkinElmer Pyris DSC, using two meth-ds: (a) non-isothermal (scanning) experiments upon heating at a scanning rate of0–100 ◦C/min in the 20–500 ◦C (700 ◦C) temperature range, and (b) an isothermal

xperiment at temperatures selected on the basis of scanning experiments, as wille explained below (Section 4).

The structure and substructure evolution was studied using JEM-100C andEM-200CX transmission electron microscopes, and thin foils were preparedhrough the “window” technique and electrolytic polishing in an HClO4 + CH3COOHolution.

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ig. 3. (a) DSC curves with different heating rates for Ti–50.0 at.%Ni alloy [7]; (b) the Kissinnd grain growth is obtained.

lloy (corresponding stress–strain diagrams shown in inserts (Ti–50.0%Ni alloy).

. TEM study of CR- and HPT-processed alloys

TEM observations show that for the Ti–50.0%Ni alloy, a log-rithmic thickness reduction of e = 2 leads to the formation ofmixed (approximately 50/50) nanocrystalline and amorphous

tructure: 2–8 nm B2-austenite nanograin zones are dispersed inhe amorphous matrix (Fig. 1a [1]). The diffraction patterns frommixed amorphous/nanocrystalline structure represent an amor-

hous halo ring with a superimposed narrow {1 1 0} ring from2-austenite nanograins. Austenite grains are so small that no sep-rate spots are visible from them. A post-deformation annealingt 300–330 ◦C for 1 h (Fig. 1b and c) leads to an almost completeanocrystallization of the amorphous structure and creates a spe-

ger [3] and (c) Ozawa [4] plots from which the activation energy for crystallization

74 K. Inaekyan et al. / Journal of Alloys and Compounds 473 (2009) 71–78

T (a) a

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atavadtstc[ttfis classically observed for moderately cold-worked alloys, whererecovery, polygonization and recrystallization onset temperaturesdecrease with increased deformation). It can therefore be assessedthat for a heterogeneous amorphous-nanocrystalline material, itsapparent thermal stability will depend on a trade-off between the

Fig. 4. Scanning DSC curves of CR, HP

ific austenite structure with a bi-modal grain size distribution: aoarser-grain population (nanograins of about 20 nm in size) grow-ng from nanocrystals formed during cold work, and a finer-grainopulation (5 nm nanograins) crystallized from the amorphoustructure (Fig. 1a). After annealing at 350 ◦C, the effect of the bi-odal grain size distribution is amplified with finer grains growing

p to 5–25 nm, and coarser grains, to 30–80 nm (Fig. 1d). Afternnealing at 400 ◦C, the structure contains normally distributedustenite nanograins with dimensions varying from 20 to 120 nm1] (Fig. 1e).

After HPT, e = 6.16, of the Ti–50.0%Ni alloy, a primarilymorphous structure is observed (see Fig. 1f), and melt-spuni50Ni25Cu25 alloy is completely amorphous, as verified in2].

. DSC study of CR-, HPT-processed and melt-spun alloys

Two DSC experiments are performed in a bid to cross-verify theesults of calculations of the activation energy for the crystalliza-ion and grain growth of the severely cold-worked Ti–Ni alloys:a) temperature scanning experiment (Kissinger–Ozawa analysis3,4]) and (b) isothermal experiment (Johnson–Avrami analysis5,6]).

.1. Introductory remarks: DSC of heterogeneous compounds

To study the influence of the post-deformation annealing onhe crystallization and grain growth of the severely cold-workedi–50.26%Ni alloy, a DSC analysis was performed at a constant ratef 50 ◦C/min (Fig. 2). It can be observed that annealing at tem-eratures of up to 300 ◦C (60 min) results in surprisingly identicalxothermal crystallization and grain-growth peaks. After annealingt 325 ◦C (10 min), a minor decrease is observed in the exothermalffect, and after 60 min, the exothermal effect vanishes completely,aking way for the endothermal effect of reverse martensitic trans-

ormation (Fig. 2).Obviously, the DSC peaks observed in the annealed alloy (300 ◦C,

0 min) do not reflect the same phenomena as those measuredn the as-cold-rolled material because the amorphous structure islmost absent after annealing at 300 ◦C. It becomes clear that scan-ing DSC experiments using heterogeneous samples such as those

ested cannot help in distinguishing between contributions fromhe crystallization and grain-growth processes, and thus in eval-ating the quantity of the amorphous phase in the cold-workedaterial. Note that microscopical evolutions occurring in the alloyhen its annealing temperature increases from 300 to 325 (330) ◦C,

FTa

nd melt-spun (b) Ti–Ni-based alloys.

uch as those captured by TEM (Fig. 1b and c) result in a drasticifference in terms of calorimetric and mechanical properties (see

nsert in Fig. 2).

.2. Thermal stability: scanning experiments (Kissinger/Ozawanalysis)

Two parameters derived from DSC analysis are generally seens indicators of the thermal stability of amorphous and nanocrys-alline materials: the onset (or peak) temperature for thermallyctivated crystallization or grain growth Tx (or Tp) and the acti-ation energy �E, with lower values of Tx (Tp) and �E indicatinglower thermal stability. Two concurrent mechanisms, which areifferent for the amorphous and nanocrystalline phases, influenceheir thermal stability. On the one hand, the higher the cold worktrain and the higher the volume fraction of the amorphous phase,he higher the thermal stability of the latter, because fewer retainedrystallites with high nucleation potency remain in the material7]. On the other hand, the finer the remaining crystalline struc-ure, the lower the onset temperature for the grain growth, becausehe higher level of accumulated elastic energy on grain boundariesacilitates grain-growth initiation (note that the latter phenomenon

ig. 5. Activation energy of crystallization and grain-growth processes fori–50.26%Ni alloy after CR, Ti–50.0%Ni alloy after HPT and melt-spun Ti50Ni25Cu25

lloy.

K. Inaekyan et al. / Journal of Alloys and Compounds 473 (2009) 71–78 75

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ig. 6. DSC isothermal thermograms for CR (e = 2) (a), CR (e = 2) + 300 ◦C (b) sampi50Ni25Cu25 alloy (d).

olume fractions of the amorphous and nanocrystalline phases andhe activation energies for their respective transformations.

.2.1. CR-processed alloysThe apparent activation energies (�E) for the crystallization

nd grain-growth processes are determined by Kissinger’s methodsing scanning experiments at different heating rates (Fig. 3a), andeasuring the slope of a [ln(sT2

p ) − T−1p ] plot [3] (Fig. 3b):

E = R

{d[ln(sT2

p ]

d(T−1p )

}, (1)

here R (J/K mol) is the gas constant, s (K/min) is the heating ratend Tp (K) is the peak temperature from a DSC plot. For verification,he activation energy is also calculated using Ozawa’s equation [4]Fig. 3c):

n s = −1.0516�E

RTp+ const. (2)

It is found that for the CR-processed Ti–50.26%Ni alloy,E = 262 kJ/mol (Kissinger) and 257 kJ/mol (Ozawa), which is about

0–40% lower than the values obtained for the rapidly solidifiedear-equiatomic Ti–Ni alloys in [8,9], but it is consistent with ourrevious results obtained with Ti–50.0 and 50.7%Ni alloys [7]. Thectivation energy for the Ti–50.26%Ni alloy after annealing at 200 ◦C1 h) is 244 kJ/mol (both methods); after annealing at 300 ◦C (1 h),63 and 260 kJ/mol (Kissinger versus Ozawa) and after annealingt 325 ◦C (10 min), 246 and 244 kJ/mol (Kissinger versus Ozawa).s we can see, the activation energy values for annealed samplesre closely approximated from the cold-rolled alloy. However, its obvious that the activation energies measured for the annealedamples correspond to the grain-growth process exclusively (Fig. 1bnd c), while for the as-CR samples, the activation energy measuredeflects both the crystallization and the grain-growth processes that

ollow only after TEM (Fig. 1a).

.2.2. HPT-processed alloysThe appearance of two exothermic peaks on the DSC thermo-

rams of the HPT-processed Ti–50.0%Ni samples can be seen in

ptm

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f Ti–50.26%Ni alloy, HPT-processed (e = 6.16) Ti–50.0%Ni alloy (c), and melt-spun

ig. 4a. This effect can be associated with the material heterogene-ty inherent to HPT, when a cold work strain varies from “zero,” inhe sample’s centre, to a maximum strain on its outer boundary,esulting in a different quantity of the remaining nanocrystals inhe amorphous matrix in different zones of the sample. This effect,nown as multistage crystallization, is also known for the non-SMA,nd generally explained by the heterogeneous nucleation of theanoscale crystals embedded in the amorphous matrix [10]. Usinghe Kissinger method for HPT samples (e = 6.16), the activation ener-ies that are separately assessed for the first and second DSC peaksre: �E1st = 210 kJ/mol and �E2nd = 197 kJ/mol.

.2.3. Melt-spun Ti50Ni25Cu25 summarizingThe crystallization activation energy measured for melt-spun

i50Ni25Cu25 alloy (Fig. 4b) is 350 kJ/mol, which is about 1.5 timesigher than that for the CR- and HPT-processed samples (Fig. 5),ut lower than that obtained by Buschow (Ti–Ni melt-spun rib-on: 531 kJ/mol) [8]) and Tong (Ti50Ni25Cu25: 406 kJ/mol [11]) andhen (Ti50Ni40Cu10: 388 kJ/mol [9]). This correlation between theE values of melt-spun and SPD-processed alloys also confirms the

resence of a heterogeneous structure in the second one (see Fig. 1and f).

Since the crystallization activation (amorphous/nanocrystallinetructure, CR-alloy) and grain-growth (nanocrystalline structure,R + annealing) energies are almost identical (Fig. 5), it is impossibleo estimate the amorphous fraction volume based solely on DSCates.

.3. Thermal stability: isothermal experimentJohnson–Mehl–Avrami analysis)

Through the isothermal calorimetric study, it is possible to qual-tatively assess the nature of the crystallization and grain-growth

henomena [12,13] and compare the activation energies (�E) forhe crystallization and grain-growth processes with that deter-

ined from scanning experiments (see above).Fig. 6 compares isothermal signals for CR (e = 2), CR + 300 ◦C

1 h), HPT (e = 6.16)-processed and melt-spun alloys. To obtain these

76 K. Inaekyan et al. / Journal of Alloys and Compounds 473 (2009) 71–78

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ig. 7. (a) Isothermal DSC curve of CR Ti–50.26%Ni alloy; (b) transformed volumealculated; (d) [ln k] versus [T−1] plot, from which the apparent activation energy fo

lots, the samples were heated at a 50 ◦C/min rate to a temperaturelose to that at the beginning of the DSC peak obtained from scan-ing experiments (around 350–360 ◦C for CR and HPT, and 470 ◦C

or melt-spun alloys (Fig. 4)), and then held until the signal varia-ions dropped below 1% of the maximum signal amplitude. Aftereing cooled down to room temperature, the samples were heatednce again in order to establish the baseline.

It can be seen that the melt-spun alloy shows one peak, and thePT-processed alloy, at least two peaks, while the CR-alloy signalontinuously decreases (Fig. 6a–c). It is known that the crystalliza-ion of a truly amorphous structure results in a peak-containing

alorimetric signal corresponding to the nucleation-and-growth ofnew crystalline phase, while the grain growth in fine-grainedaterials leads to a continuously decreasing calorimetric signal

13].

bmtg

ig. 8. (a) Isothermal DSC curve of CR + 300 ◦C Ti–50.26%Ni alloy, (b) transformed volumen] is calculated, and (d) [ln k] versus [T−1] plot, from which the apparent activation energ

n as a function of time; (c) the JMA plot from which the Avrami coefficient [n] istallization and growth [�E] is obtained.

In our study, the peak from the melt-spun alloy can be clearlyttributed to the crystallization of a truly amorphous materialFig. 6d), and the decreasing signal from the CR-alloy, to therain-growth phenomena; for the HPT-processed alloy, it cane attributed to a sequential realization of these two processes.iven that the amorphous phase is more thermally stable than

he ultrafine-grained material, the first peak of the HPT signalan be attributed primarily to the grain growth of the remain-ng nanocrystallites in the amorphous matrix, while the second,erives from the nucleation of a new crystalline phase and tohe grain growth of nanocrystals from both origins. It can also

e assumed that a deviation of the CR signal from the trulyonotonous (dotted line) results from a mixed contribution from

wo coexisting phenomena: crystallization and growth and grainrowth.

fraction as a function of time, (c) the JMA plot from which the Avrami coefficienty for crystallization and growth [�E] is obtained.

K. Inaekyan et al. / Journal of Alloys and Compounds 473 (2009) 71–78 77

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ig. 9. (a) Isothermal DSC curve of melt-spun Ti50Ni25Cu25 alloy; (b) transformed vn] is calculated; (d) [ln k] versus [T−1] plot, from which the apparent activation ene

Solid-state transformation kinetics involving both crystalliza-ion (nucleation-and-growth) and grain-growth phenomena isenerally studied using the Johnson–Mehl–Avrami (JMA) isother-al analysis [5,6]:

(t) = 1 − exp[−(kt)n], (3)

here x is the volume fraction of transformed material, n is theinetic (Avrami) exponent, which reflects the nucleation mecha-ism and growth morphology, k is the reaction rate constant, which

s a function of the absolute temperature (T) and the apparent acti-ation energy (�E):

= k0 exp(−�E

RT

). (4)

The values of k and n can be determined using the Eq. (1) rewrit-en in the following form:[ (

1)]

n ln1 − x

= n ln(k) + n ln(t − �), (5)

here � is the incubation time.By plotting ln[ln(1/1 − x)] versus ln(t − �), it is possible to deter-

ine n.

Fig. 10. Avrami exponent of CR, HPT and melt-spun alloys.

T

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4

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cttci

fraction as a function of time; (c) the JMA plot from which the Avrami coefficientr crystallization and growth [�E] is obtained.

The plot ln(k) versus 1/T gives the value of the activation energyrom the equation rewritten as follows:

n(k) = ln(k0) − �E

RT. (6)

.3.1. CR-processed alloysSamples of the Ti–50.26%Ni alloy after e = 2 were annealed

sothermally at temperatures of 200, 300, 325, 345, 360, and 375 ◦C.ssuming then that the volume fraction of the transformed mate-ial is proportional to the heat measured under the DSC peakselected diagrams are presented in Fig. 7a), the volume fractionsf the transformed material are then plotted according to timeFig. 7b). The Avrami plot ln[ln(1 − x)−1] versus [ln(t)] then yieldsstraight line with the intercept [n ln k] from which the slope [n]

an be calculated (Fig. 7c). Finally, [ln k] is plotted versus [T−1] toxtract parameter [k] and evaluate the apparent activation energyor transformation �E (Fig. 7d).

A similar isothermal experiment was performed on thei–50.26%Ni alloy after CR and annealing at 300 ◦C (1 h) (Fig. 8).

The Avrami coefficient and the activation energy calculated forhe annealed alloy are almost equal to that of the as cold-rolledlloy (see Figs. 7d and 8d), making it difficult to fully distinguishrystalline and partially amorphous structures using just the DSC-ethod. The Avrami exponent indicates that diffusion-controlled

rain growth in both cases passes with the decreasing of theucleation rate according to the Christian crystallization processeslassification [14].

.3.2. HPT-processed alloysThe Avrami–Johnson analysis is repeated for an HPT (e = 6.16)-

rocessed Ti–50.0 at.%Ni alloy separately for the first and the secondeaks of the isothermal DSC thermogram (Fig. 6c). The analysisroduces n = 2.1 and n = 2.95, respectively.

While the activation energy values are close for both pro-

esses, the Avrami exponent n after HPT is 2.5 times higher thanhat after CR, which reflects differences in their respective crys-allization and grain-growth mechanisms. Applying the Christianlassification [14], we can describe the transformation kineticsn the CR- and HPT-processed binary Ti–Ni alloys as follows: the

78 K. Inaekyan et al. / Journal of Alloys and Compounds 473 (2009) 71–78

Table 2Collected data on DSC results for partially and completely amorphized Ti–Ni-based alloys

Ti1−x–Nix (x, at.%) Processing DSC rate (◦C/min) Tx (Tp) (◦C) �H (kJ/mol) Ea (kJ/mol) Source

50 CR (e = 1.9) 10 265 (350) 0.11 – Ewert [16]50 Mechanical alloying 40 460 – – Eckert [17]50 CR (1.9) 50 354 (376) 1.07 262 Brailovski [7]50 HPT (e = 6.16) 50 1st peak 365 (387), 2nd peak 413 (433) 2.5 210 (1st peak),

197 (2nd peak)This study

50.26 CR (e = 2.0) 50 350 (375) 1.2 261 This study50.3 HPT (e = 6.7) 40 352 (374) 1.4 – Waitz [18]50.3 HPT (e = 7.3) 40 362 (379) 1.7 – Waitz [18]50.62 HPT (e = 7) 40 200 (375) – – Sergueeva [19]50.7 CR (e = 1.55) 50 342 (373) 0.76 262 Brailovski [7]49.93 Sputtering 50 502 (507) – 416 Chen [9]5TTT

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[[16] J.C. Ewert, I. Bohm, R. Peter, F. Haider, Acta Mater. 45 (5) (1997) 2197–2206.

0 Melt spinning 50 506i50Ni25Cu25 Melt spinning 40 472 (473)i50Ni25Cu25 Melt spinning 50 471 (474)i50Ni40Cu10 Sputtering 50 518 (529)

sothermal annealing of the CR (e = 2)-processed alloy (n = 1.25 andE = 262 kJ/mol) is dominated by an interface-controlled grain-

rowth mechanism, while that of the HPT (e = 6.16)-processed alloyn = 2.1–2.95 and �E = 210 kJ/mol (1st peak) and 197 kJ/mol (2ndeak)) is dominated by a nucleation-and-growth mechanism [14].

.3.3. Melt-spun Ti50Ni25Cu25 summarizingThe activation energy value for the melt-spun alloy is higher

han for SPD alloys (Figs. 9d and 10) as well as it was obtained bycanning DSC experiment (Fig. 5). The n values obtained for theelt-spun alloy (Figs. 9c and 10), which are greater than those for

R- and HPT-processed samples, indicates that the crystallizationrocess passes with the increasing nucleation rate, and the grain-rowth process does not dominate [14], unlike with CR-processedlloys.

Table 2 presents a comparison of the data obtained from the DSCnalyses of the CR- and HPT-processed binary Ti–Ni alloys and theelt-spun Ti–Ni–Cu alloy with data in the literature. It should be

oted that the activation energy values in Table 2 represent the acti-ation energy for the overall crystallization and growth processes,hich is a combination of two energy components: the nucle-

tion barrier (En) and the activation energy for growth (Eg). Thesewo energy components are related as follows: �E = (aEn + bEg)/n,here a and b are positive constants related to the Avrami expo-ent as n = a + b), and can be estimated separately using in situicroscopic investigations [15].

. Conclusions

1) CR- and HPT-processed Ti–Ni alloys contain amorphous matri-ces with embedded nanocrystals; exothermal peaks measuredby DSC contain contributions from both crystallization andgrain-growth phenomena, and so DSC experiments can there-fore not be used to evaluate the volume fraction of theamorphous phase in such materials; it can only be done bydirect TEM observations.

2) For CR (e = 2), the Avrami exponent n = 1.25, which corresponds

to annealing by a grain-growth-dominated mechanism, andfor HPT (e = 6), n = 2.5, which corresponds to annealing bynucleation-and-growth.

3) The apparent thermal stability of CR- and HPT-processed Ti–Nialloys depends on a trade-off between the volume fractions of

[

[[

3.5 531 Buschow [8]– 406 Tong [11]2.33 351 This study– 388 Chen [9]

the amorphous and nanocrystalline phases and activation ener-gies for their respective transformations—aspect to be studiedby direct TEM observations.

4) The crystallization and grain-growth processes in CR samplestake place in the same temperature interval as the grain-growthprocess for CR + annealing (until 300 ◦C (1 h)). The heat flow, theactivation energy and the Avrami coefficient are close for CR andCR + annealing alloys.

cknowledgements

The authors would like to thank the Natural Sciences and Engi-eering Research Council of Canada and the Federal Agency forducation of the Russian Federation for their financial support, Dr.. Shelyakov for supplying the authors with Ti–Ni–Cu melt-spunlloy, and Mr. V. Demers for performing tensile testing within theramework of this study.

eferences

[1] V. Brailovski, S. Prokoshkin, I.Yu. Khmelevskaya, K.E. Inaekyan, V. Demers, S.V.Dobatkin, E.V. Tatyanin, Mater. Trans. JIM 47 (2006) 795–804.

[2] H. Rosner, A.V. Shelyakov, A.M. Glezer, K. Feit, P. Scholoßmacher, Mater. Sci. Eng.A273–275 (1999) 733–737.

[3] H.E. Kissinger, Anal. Chem. 29 (1957) 1702–1706.[4] T.J. Ozawa, Therm. Anal. 2 (1970) 301–324.[5] W.A. Johnson, K.F. Mehl, Trans. AIME 135 (1939) 416–442.[6] M.J. Avrami, J. Chem. Phys. 7 (1939) 1103–1112;

M.J. Avrami, J. Chem. Phys. 8 (1940) 212;M.J. Avrami, J. Chem. Phys. 9 (1941) 177.

[7] V. Brailovski, S.D. Prokoshkin, E. Bastarash, V. Demers, K.E. Inaekyan, I.Yu.Khmelevskaya, Mater. Sci. Forum 539–543 (2007) 1964–1970.

[8] K.H.J. Buschow, J. Phys. F: Met. Phys. 13 (1983) 563–571.[9] J.Z. Chen, S.K. Wu, J. Non-Cryst. Solids 288 (2001) 159–165.10] A. Inoue, Acta Mater. (2000) 279–306.11] Y. Tong, Y. Liu, J. Alloys Compd. 449 (2008) 152–155.12] A.P. Zhilyaev, G.V. Nurislamova, S. Surinach, M.D. Baro, T.G. Langdon, Mater.

Phys. Mech. 5 (2002) 23–30.13] L.C. Chen, F. Spaepen, J. Appl. Phys. 69 (2) (1991) 679–688.14] J.W. Christian, The Theory of Transformations in Metals and Alloys, 2nd ed.,

Pergamon, New York, 1975.15] X. Wang, J.J. Vlassak, Scripta Mater. 54 (2006) 925–930.

17] J. Eckert, J.C. Holzer, C.E. Krill III, W.L. Johnson, J. Mater. Res. 7 (1992) 1751–1761.

18] T. Waitz, V. Kazykhanov, H.P. Karnthaler, Acta Mater. 52 (2004) 137–147.19] A.V. Sergueeva, C. Song, R.Z. Valiev, A.K. Mukherjee, Mater. Sci. Eng. A339 (2003)

159–165.