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Thermochimica Acta 554 (2013) 71– 75

Contents lists available at SciVerse ScienceDirect

Thermochimica Acta

jo ur n al homepage: www.elsev ier .com/ locate / tca

ffect of Ag addition on phase transitions of theu–22.26 at.%Al–9.93 at.%Mn alloy

.A.G. Silvaa,∗, S. Gamaa, A. Paganotti a, A.T. Adornob, T.M. Carvalhob, C.M.A. Santosb

DCET, Universidade Federal de São Paulo, Campus Diadema, SP, BrazilDFQ, Instituto de Química – Unesp, Campus Araraquara, SP, Brazil

r t i c l e i n f o

rticle history:eceived 10 September 2012eceived in revised form 5 December 2012

a b s t r a c t

The phase transitions that occur in the Cu–22.26 at.%Al–9.93 at.%Mn and Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloys after slow cooling were studied using differential scanning calorimetry at different heat-ing rates, microhardness changes with temperature, magnetization changes with temperature, scanning

ccepted 22 December 2012vailable online 4 January 2013

eywords:etals and alloys

hase transitions

electron microscopy and energy dispersion X-ray spectroscopy. The results indicated that the presenceof Ag does not modify the transition sequence of Cu–Al–Mn alloy, introduces a new transition due to the(Ag-Cu)-rich precipitates dissolution at about 800 K, and changes the mechanism of DO3 phase dissolu-tion. This mechanistic change was analyzed and a sequence of phase transitions was proposed for thereaction.

hermal analysis

. Introduction

Alloys of the Cu–Al–Mn system are Hume–Rothery alloys. Thehase stability of these materials is largely dominated by theverage number of conduction electrons per atom, denoted as/a. Within a certain range of compositions around the Cu-richegion (far from the Cu2AlMn Heusler stoichiometry) the systemndergoes a martensitic transition associated with shape-memoryroperties [1]. The Cu-rich portion of the Cu–Mn–Al system haseen fully investigated at temperatures above 723 K. It was foundhat the addition of Mn expands the high temperature � phase fieldf the Cu–Al system. The � phase stability compositional rangencreases with temperature, and includes the composition of the

ell-known Heusler alloy Cu2MnAl at temperatures above 923 K.t 673 K the equilibrium phases in these alloys are: Cu3Mn2Al, �nd �Mn [2]. At temperatures lower than 600 K, a spinodal decom-osition can occur between a non-magnetic (DO3) Cu3Al-rich phasend a (L21) Cu2AlMn-rich phase, which orders ferromagnetically.he kinetics of this process is slow and this decomposition isvoided for typical cooling rates. Associated with this spinodalecomposition some authors have reported the appearance ofuperparamagnetism at low temperatures [3].

Silver additions to Cu–Al alloys increase its hardness [4], stressorrosion resistance [5] and modify the aging characteristics ofhe alloys [6], with no ternary intermediate phases being observed

∗ Corresponding author. Tel.: +55 11968369426.E-mail address: galdino.ricardo@gmail.com (R.A.G. Silva).

040-6031/$ – see front matter © 2013 Elsevier B.V. All rights reserved.ttp://dx.doi.org/10.1016/j.tca.2012.12.014

© 2013 Elsevier B.V. All rights reserved.

[7,8]. It is expected that Ag additions to Cu–Al–Mn alloys presenta similar effect, thus improving some properties of the systemmainly its corrosion resistance, apart from allowing the study ofthe effect of Ag presence on the characteristics temperatures andmechanism of phase transitions. In this work, the influence of addi-tion of 1.53 at.%Ag to the Cu–22.26 at.%Al–9.93 at.%Mn alloy wasstudied using differential scanning calorimetry (DSC) at differentheating rates, microhardness changes with temperature, magne-tization changes with temperature, scanning electron microscopy(SEM) and energy dispersion X-ray spectroscopy (EDXS). The pur-pose of this work is to study the phase transitions in the Cu–Al–Mnand Cu–Al–Mn–Ag alloys after slow cooling at about 5 K/h, cor-relating with the presence of Ag and its influence on the ternarysystem.

2. Experimental procedure

The Cu–22.26 at.%Al–9.93 at.%Mn and Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloys were prepared in an arc fur-nace under argon atmosphere using 99.95% copper, 99.97%aluminum, 99.995% manganese and 99.98% silver as startingmaterials. Cylindrical samples with 2.0 cm diameter and 6.0 cmlength were cut in disks of 2.0 mm thickness. The samples weremaintained for 120 h at 1123 K and then cooled at 5.0 K/h down toroom temperature. The DSC curves were obtained using a DSC Q20

TA Instruments at different heating rates. The temperature andheat flow were calibrated by measuring the melting temperaturesand fusion heat of pure In. Scanning electron microscopy imageswere obtained using a JSM LV5900 installed in the LME-LNNano,

72 R.A.G. Silva et al. / Thermochimica Acta 554 (2013) 71– 75

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-0,2

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1,6Cu-Al-Mn

Ma

gn

etic M

om

ent

(em

u/g

)

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Cu-Al-Mn-Ag

ig. 1. DSC curves obtained at different heating rates for the annealed alloys: (a)u–22.26 at.%Al–9.93 at.%Mn and (b) Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag.

razil. EDXS analyses were made using the micro analyzer installedn the LME-LNNano. Semi-quantitative measures were obtained inhe energies range from 0 to 10 keV. The Vickers microhardness

easurements were made with a SHIMADZU HMV-2T Microhard-ess Tester using a load of 9.8 N. The magnetic properties wereeasured using a vibrating-sample magnetometer (VSM) in the

emperature range from 300 K to 650 K.

. Results and discussion

Fig. 1 shows the DSC curves obtained at differenteating rates for the Cu–22.26 at.%Al–9.93 at.%Mn andu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloys. These alloysere initially maintained at 1123 K for 120 h and then cooled

t 5.0 K/h down to room temperature. The curves of Fig. 1a,orresponding to the Cu–22.26 at.%Al–9.93 at.%Mn alloy, showhree thermal events. The endothermic peak P1, at about 520 K,s shifted to higher temperatures with the increase of the heatingate. This thermal event is attributed to the DO3 + L21(f) → L21(p)ransition, in which the DO3 phase is dissolved into the matrixnd the ferromagnetic L21(f) structure produces the para-

agnetic L21(p) phase. This reaction corresponds to the inverse

f spinodal decomposition [3]. The magnetic transition wasonfirmed by measurements of magnetic moment change withemperature, as seen in Fig. 2. The Curie temperature for the

Fig. 2. Plots of magnetic moment change vs. temperature obtained for theCu–22.26 at.%Al–9.93 at.%Mn and Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloys,where (→) heating and (←) cooling.

Cu–22.26 at.%Al–9.93 at.%Mn alloy is at about 525 K. The endother-mic peak P2, at about 780 K, is related to the L21(p) → B2 and� + T3-Cu3Mn2Al + � → � + � + T3-Cu3Mn2Al reactions [1]. It isinteresting to notice that at heating rates lower than 10 K/min the� + T3-Cu3Mn2Al + � → � + � + T3-Cu3Mn2Al reaction is dominantwhile at higher heating rates the L21(p) → B2 transition is domi-nant, as seen in the enlarged portion inside Fig. 1a. For rates higherthan 10 K/min the dominant transition is not dependent on theheating rate. In Fig. 1a, the thermal event at about 855 K, associatedwith a baseline change, is due to the � + � + T3-Cu3Mn2Al → � + T3-Cu3Mn2Al and B2 → A2 reactions [1]. The curves of Fig. 1b,corresponding to the Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Agalloy, show the same thermal events observed in Fig. 1a withtransition temperatures shifted to higher values, as comparedto the Cu–22.26 at.Al%–9.93 at.%Mn alloy. This indicates thatthe presence of Ag does not modify the sequence of expectedtransitions, but changes the stability range of the phases. Themagnetic transition that occurs together with peak P1 in theCu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloy was confirmed bythe measurement of magnetic moment change with temperatureand the Curie temperature obtained is at about 535 K, as shown inFig. 2. The Curie temperature is little influenced by the presence ofAg, but at the beginning of the curves in Fig. 2 the maximum valueof magnetic moment is higher than that observed in the samplewithout Ag. This suggests a change in the degree of order of theferromagnetic Cu2AlMn(f) phase, probability due to the presenceof Ag-rich precipitates and its interaction with the ferromagneticphase.

For the Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloy theenthalpy value for the phase transitions that occur in P2 ishigher than that verified for the same thermal event in theCu–22.26 at.%Al–9.93 at.%Mn alloy. The mean value obtained forthe alloy with Ag addition was �H = (23.2 ± 0.5) J/g and for the alloywithout Ag addition was �H = (14.2 ± 1.6) J/g. This suggest thatanother thermal event can be occurring in the same temperaturerange of the L21(p) → B2 and � + T3-Cu3Mn2Al + � → � + � + T3-Cu3Mn2Al transitions, at about 790 K, in the curve obtained for theCu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloy. In the enlargedportion inserted in Fig. 1b it is possible to observe that at heating

rates lower than 2.5 K/min the � + T3-Cu3Mn2Al + � → � + � + T3-Cu3Mn2Al reaction is dominant and at higher heating rates theL21(p) → B2 transition and another phase transition are dominant.It is also important to notice that peak P2 depends on the heating

R.A.G. Silva et al. / Thermochimica Acta 554 (2013) 71– 75 73

F 1 at.%9

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ig. 3. Scanning electron micrographs (SEM) obtained for the Cu-22.49 at.%Al-10.070 K (c), respectively. (d) EDXS spectrum obtained from white precipitate in (b).

ate from 5.0 K/min, as seen in the enlarged portion in Fig. 1b. Thisonfirms that another thermal event is occurring together witheak P2.

Fig. 3a–c shows the scanning electron micrographs (SEM)btained for the Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloyuenched from different temperatures chosen from the DSC curves

n Fig. 1b. The SEM images were taken from a sample initiallynnealed and then quenched from 623 K and 970 K, respectivelyefore and after peak P2 in Fig. 1b. In the micrographs of Fig. 3a and

it is possible to notice the presence of precipitates (white region).

n the SEM image of Fig. 3c the precipitates were not detected.n Fig. 3d the EDXS spectrum taken from the precipitate (whiteegion) shows that it is formed mainly by Cu and Ag atoms. There-ore, the additional thermal event that occurs together with the

Mn-1.53 at.%Ag alloy initially annealed (a) and then quenched from 623 K (b) and

L21(p) → B2 and � + T3-Cu3Mn2Al + � → � + � + T3-Cu3Mn2Al tran-sitions, in the DSC curves of Fig. 1b at about 790 K, may be ascribedto (Ag–Cu)-rich precipitates dissolution.

Fig. 4 shows the curves of microhardness change withtemperature, obtained for the Cu–22.26 at.Al%–9.93 at.%Mnand Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloys initiallyannealed. In this figure one can observe that the curve obtainedfor the Cu–22.26 at.%Al–9.93 at.%Mn alloy presents microhard-ness values practically constant up to 423 K. From 423 K themicrohardness values decrease abruptly up to around 573 K.

This behavior was attributed to the DO3 phase dissolution thatoccurs from DO3 + L21(f) → L21(p) transition and the consequent� phase formation, as seen in Fig. 5a and b. From 573 to 723 Kone can observe an increase in the microhardness values. In

74 R.A.G. Silva et al. / Thermochimi

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923 K723 K

573 K

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423 K

taadpioCIimTdfipcaCisCD5t

Fig. 4. Plots of microhardness changes vs. quenched temperature.

his temperature range �, T3-Cu3Mn2Al, � and L21(p) phasesre stable and the increase in the microhardness values can bescribed to the rearrangement of Cu and Al atoms from DO3issolution and partial dissolution of unstable phases. In the tem-erature range between 723 and 923 K there is a slight decrease

n the microhardness values that is related to the occurrencef the L21(p) → B2, B2 → A2, � + T3-Cu3Mn2Al + � → � + � + T3-u3Mn2Al and � + � + T3-Cu3Mn2Al → � + T3-Cu3Mn2Al reactions.

t is important to notice that using microhardness measurementss not possible to separate the effects of these reactions on this

echanical property in the Cu–22.26 at.%Al–9.93 at.%Mn alloy.his seems to suggest that the atomic exchanges that occururing these phase transitions do not lead to intense modi-cations on the atomic distances and on interactions of theroduct phase constituents. The microhardness curve in Fig. 4,orresponding to the Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Aglloy, indicates a behavior similar to that observed for theu–22.26 at.%Al–9.93 at.%Mn alloy. Up to 573 K one can ver-

fy that the Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloyhows higher microhardness values than those obtained for the

u–22.26 at.%Al–9.93 at.%Mn alloy. This can be associated withO3 lattice stress, which is eliminated with its dissolution after73 K. This lattice stress can be due to an Ag atoms interaction withhe DO3 phase. From 573 to 723 K the Cu–22.26 at.%Al–9.93 at.%Mn

Fig. 5. Scanning electron micrographs (SEM) obtained for the Cu–22.26 at.%Al–9

ca Acta 554 (2013) 71– 75

and Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloys exhibitmicrohardness values quite similar. This suggests that the DO3phase and its interaction with Ag is determinant on the differencein the microhardness values observed between these alloys upto around 573 K. It is interesting to notice that between 573 Kand 723 K there are precipitates embedded in the matrix, butthey do not interfere on the microhardness values of the alloy(see Fig. 3b). This effect may be associated with great size ofthe precipitates, as seen in Fig. 3a and b. This decreases thecoherence between the precipitates and matrix, thus contribut-ing for decrease of the effect of Ag on microhardness of theCu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloy. For temperaturesabove 723 K the Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloyshows higher microhardness values due to the complete dissolu-tion of the precipitates and the consequent increase of the latticestress due to the accommodation of these atoms in the metallicmatrix at elevated temperatures.

Considering that the DO3 or Cu3Al phase in the presence ofAg is responsible for the shift in the microhardness values ofthe Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloy when com-pared with the alloy without Ag addition, a kinetic study and theproposition of a mechanism for the DO3 dissolution process couldcontribute for understanding of the DO3 phase stability. Peak P1 inthe DSC curves of Fig. 1 corresponds to the DO3 + L21(f) → L21(p)transition and it is shifted to higher temperatures with increaseof the heating rate, due to the diffusive processes associ-ated with DO3 phase dissolution. In order to analyze the DO3phase dissolution kinetics in the Cu–22.26 at.%Al–9.93 at.%Mnand Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloys, an isocon-versional method based on the Model-free kinetics was used [9,10].

The isoconversional method was used to obtain theactivation energy for the DO3 dissolution process in theCu–22.26 at.%Al–9.93 at.%Mn and Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloys, considering the first peak in Fig. 1 and Eq. (1)[11]:

ln

[�i

(dy

dT

)y,i

]= ln[Ayf (y)] − Ey

RTy,i(1)

where a subscript y designates values related to a given conversion,i is the number of non-isothermal experiments conducted at theheating rate �i, Ty,i is the absolute temperature, Ay is the Arrhenius

constant, Ey is the activation energy and R is the gas constant. Byplotting ln[�i·(dy/dT)y,i] against 1/Ty,i, the value of the Ey for a giveny can be obtained. These plots obtained from Eq. (1) were linear,with correlation coefficients no less than 0.9892.

.93 at.%Mn alloy initially annealed (a) and then quenched from 623 K (b).

R.A.G. Silva et al. / Thermochimi

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120

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y

Fig. 6. Plots of activation energy changes vs. conversion factor.

Fig. 6 shows the plots of Ey vs. y. The curve correspond-ng to the Cu–22.26 at.%Al–9.93 at.%Mn alloy indicates that theonversion dependence of the activation energy has an ascend-ng shape, indicating the presence of multi-step processes. Forhe Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloy the conver-ion dependence of the activation energy has a descending shape,ndicating that the reaction mechanism is related to multi-steprocesses involving reversible reaction [10].

It is known [12] that in Cu–Al–Mn alloys an intermediatepecie, DO3-L21(f) or [Cu3Al–Cu2AlMn(f)] phase, can be found. Forhe Cu–22.26 at.%Al–9.93 at.%Mn alloy the results indicate that theeaction mechanism could be:

u2AlMn(f)+Cu3Alk1−→[Cu2AlMn(f)–Cu3Al] (I) slow

Cu2AlMn(f)–Cu3Al]k2−→Cu2AlMn(p) + 3Cu + Al (II) fast

This mechanism suggests that the intermediate phase is slowlyormed in step I and then it is rapidly dissolved in step II, thusroducing a depletion of the [Cu3Al–Cu2AlMn(f)] phase and conse-uently an increase in the apparent activation energy values withonversion factor, as observed in Fig. 6. The step II is associatedith releasing of Cu for the system. This was confirmed by the scan-ing electron micrographs of the Fig. 5, where it is possible to seehe growing of the �-(Cu-rich) phase after the DO3 dissolution, asiscussed in Fig. 4.

In the presence of Ag atoms the [Cu3Al–Cu2AlMn(f)] phase isess stable and steps involving a reversible reaction (steps III andV) occur, as proposed:

u2AlMn(f) + Cu3Alk1−→[Cu2AlMn(f)–Cu3Al] (III) fast

k−1

Cu2AlMn(f)–Cu3Al]−→Cu2AlMn(f) + Cu3Al (IV) fast

Cu2AlMn(f)–Cu3Al]k2−→Cu2AlMn(p) + 3Cu + Al (V) slow

[

ca Acta 554 (2013) 71– 75 75

This process contributes for the increase in the DO3 phaserelative fraction in the Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Agalloy, as compared with the alloy without silver. In this way, theshift of microhardness changes to higher values is related to thedecrease in the intermediate phase stability and consequently tothe increase in the DO3 phase relative fraction, in the presence ofsilver.

4. Conclusions

The phase transitions in the Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloy were analyzed and a kinetics mechanism forthe DO3 phase dissolution was suggested. The results indicatedthat the DO3 phase interferes in the microhardness values ofthe Cu–22.49 at.%Al–10.01 at.%Mn–1.53 at.%Ag alloy. The � + T3-Cu3Mn2Al + � → � + � + T3-Cu3Mn2Al transition is dominant atlower heating rates while L21(p) → B2 transition is dominant athigher heating rates. The (Ag–Cu)-rich precipitates dissolutionoccurs in the same temperature range of the L21(p) → B2 and � + T3-Cu3Mn2Al + � → � + � + T3-Cu3Mn2Al transitions, at about 790 K.The proposed kinetics mechanism suggests a decrease in the inter-mediate phase stability in the presence of Ag.

Acknowledgments

The authors thank FAPESP (proc. 2011/11041-4), CNPq (proc.482348/2010-0) and LME/LNLS for technical support during elec-tron microscopy work (JSM-5900LV).

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