physical and structural characterization of a monocrystalline cu-13.7al-4.2ni alloy subjected to...

Physical and Structural Characterization of a MonocrystallineCu-13.7Al-4.2Ni Alloy Subjected to Thermal Cycling Treatments

ELAINE CRISTINA PEREIRA, LIOUDMILA ALEKSANDROVNA MATLAKHOVA,ANATOLIY NIKOLAEVICH MATLAKHOV, CARLOS YUJIRO SHIGUE,and SERGIO NEVES MONTEIRO

A monocrystalline alloy with nominal 82wt pctCu-13.7wt pctAl-4.2wt pctNi composition andexhibiting reversible martensitic transformation (RMT) was subjected to multiple heating andcooling cycles within the RMT range of critical temperatures. Both untreated and cyclic treatedalloy samples were characterized by X-ray diffraction, optical microscopy, differential scanningcalorimetry, and Vickers microhardness. The results indicated that the alloy presents a complexRMT behavior disclosing a sequence of transformation steps: b1 M R and R M b¢1+ c¢1 as wellas possible b1 M b¢1 and b¢1 M c¢1 direct reactions. The thermal cycling treatment inhibits thedevelopment of c¢1 martensite without much change in both the physical and microstructurecharacteristics. This suggests a good resistance of the alloy to irreversible structural changes.

DOI: 10.1007/s11661-013-2153-5� The Minerals, Metals & Materials Society and ASM International 2014

I. INTRODUCTION

MONOCRYSTALLINE Cu-Al-Ni alloys based onthe Cu3Al intermetallic compound belong to an exten-sive family of metallic materials that suffer reversiblemartensitic transformation (RMT) associated with non-elastic effects including the shape memory effect(SME).[1–3]

In these alloys the RMT takes place within a relativelynarrow range of temperatures, usually close to roomtemperature (RT). This transformation is activated notonly by heating and cooling procedures but also by theapplication of an external stress. It has been observedthat the high temperature stable b1 phase, with anordered type DO3 cubic structure, may be transformedto both martensitic phases: c¢1 with type Cu3Ti orderand/or b¢1 with type 18R order, depending on thealuminum content.[1–4] The incorporation of nickel intothe monocrystalline Cu-Al alloys, in addition to theafore-mentioned martensitic phases, could also intro-duce another metastable phase with a type R-3mordered rhombohedric structure and Al7Cu4Nistoichiometric composition.[5] These RMT induced

microstructural changes affect the properties and con-sequently the practical applications, especially thoserelated to SME.SME alloys are currently used in engineering systems

and medical devices. The typical application requires arepetitive operational regimen in which the alloy issubjected to alternated heating and cooling cyclesenclosing the RMT critical temperatures of austenitestarting (As) and finishing (Af) interval, as well asmartensite starting (Ms) and finishing (Mf). Each cycle isassociated with a thermal hysteresis during which theparticipant phases remain either coherent or semi-coherent between themselves.[2–4] One important limita-tion for the practical use of SME alloys is the structuralchange resulting after a thermal cycling treatment(TCT). Indeed, during RMT the alloy experiencesmodifications in both its phase composition and micro-structural aspects that control the physical and mechan-ical properties.[3,6–9] Actually, every thermal cycleintroduces additional crystalline imperfections, such asdislocations and residual phases, that interfere with theRMT characteristics causing variation in the criticaltemperatures and degree of thermal hysteresis.[3,10–17]

It has been known that SME alloys, particularly thoseof the Cu-Al-Ni family, may undergo not only singlephase transformations, such as b1 fi b¢1 or b1 fi c¢1,but also a dual phase b1 fi c¢1+ b¢1, depending on thealloy content.[3,10–12] For instance, the literature[13]

indicates that the dual phase transformation occurs ina copper alloy containing 13.7 wt pct Al plus 4.0 wt pctNi. Therefore, a similar composition was selected for thepresent investigation. The objective of this work wasthen to perform physical and structural characteriza-tions of a monocrystalline copper alloy with 13.7 wt pctAl and 4.2 wt pct Ni (Cu-13.7Al-4.2Ni), which displaysRMT. This alloy was subjected to different level of TCTwithout applied load, i.e., stress-free.

ELAINE CRISTINA PEREIRA, Researcher, and LIOUDMILAALEKSANDROVNA MATLAKHOVA, Professor, are with theLaboratory for Advanced Materials (LAMAV), State University ofNorthern Rio de Janeiro (UENF), Avenida Alberto Lamego, 2000,Parque California, Campos dos Goytacazes, RJ CEP 28013-602,Brazil. Contact e-mail: [email protected], [email protected] NIKOLAEVICH MATLAKHOV, formerly Professorwith the Laboratory for Advanced Materials (LAMAV), StateUniversity of Northern Rio de Janeiro (UENF), is now deceased.CARLOS YUJIRO SHIGUE, Professor, is with the Scholl Engineer-ing Lorena (EEL-USP) of the University of Sao Paulo (USP), Lorena,Brazil. SERGIO NEVES MONTEIRO, Professor, is with the MilitaryInstitute of Engineering (IME-RJ), Rio de Janeiro, Brazil.

Manuscript submitted December 18, 2012.

METALLURGICAL AND MATERIALS TRANSACTIONS A

II. EXPERIMENTAL PROCEDURE

The basic material investigated was a high puritymonocrystalline Cu-13.7Al-4.2Ni alloy, obtained as a4 mm in diameter cylindrical bar from the MemoryCrystals Group, of the Technical University of SaintPetersburg, Russia.[18] Specimens with approximately5 mm in thickness were sectioned out, perpendicular tothe bar axis, by means of a Miniton cutter. Thesespecimens were subjected to 1, 100, 200, 300, 400, and500 stress-free thermal cycles, each one associated with acooling to 258 K (�15 �C) (below Mf) followed byheating up to 373 K (100 �C) (above Af), shown inFigure 1. The limiting temperatures of 258 K and 373 K(�15 �Cand+100 �C)per cyclewere exactly the same forall samples. The structural characterization of the alloywas conducted at RT, 293 K ± 5 K (20 �C ± 5 �C), byfinishing the treatment: (a) with a half cooling cycle to258 K (�15 �C), then up to RT or (b) with a half heatingcycle to 373 K (100 �C), then down to RT.

The phase identification was performed by X-raydiffraction (XRD) in a model DRON-3M diffractometerusing Cu Ka radiation, covering 2h angles from 25 to75 deg at scanning steps of 0.03 deg/3 seconds.

For microstructure characterization, the specimenswere polished to a mirror appearance with 0.1 lmalumina paste. No chemical attack was used since thesurface transformation relief was enough to reveal thephases. This relief was observed by optical microscopy(OM) in a Neophot-32 microscope with polarized light.Vickers microhardness (HV) tests were conducted in amodel HMV-2, SHIMADZU equipment.

The RMT critical temperatures as well as its intervalsof temperature and associated thermal effects after theTCT, were determined by differential scanning calorim-etry (DSC), in a model DSC Q10 V9.8 296, TAInstrument. These tests were conducted in the range oftemperature from 223 K to 373 K (�50 �C to +100 �C)at a heating and cooling rate of 10 K/min (10 �C/min)under an inert nitrogen atmosphere.

III. RESULTS

A. Characterization of the As-Received Alloy

Figure 2 shows the XRD pattern of the as-receivedCu-13.7Al-4.2Ni alloy. In this figure, two metastablephases: the b¢1 martensite, by its prominent peaks, andthe Al7Cu4Ni, denoted as the R phase[19,20] by its tinypeaks adjacent to b¢1 peaks are observed. The tiny(107)R peak indicates a corresponding smaller amountof the R phase, which may be understood as anintermediate during b1 M R M b¢1 transformation. Italso suggests that residual R phase could be formed witha coherent plane, b¢1|R, associated with the small(21�2�1)b¢1|(0213)R peak shown in Figure 2.

The DSC analysis of the as-received alloy in Figure 3shows that upon heating the reverse b¢1 fi b1 austenitictransformation occurs in the interval from 311.2 K(+38.2 �C) (As) to 341.8 K (+68.8 �C) (Af) by means ofan endothermic process with a phase transformation

enthalpy of 9 J/g. The onset of this process, by compar-ison with the baseline, was determined at 324.4 K(+51.4 �C) (Ao) and a peak corresponding to themaximum heat flow was found at 327.7 K (+54.7 �C)(Ap). The process ends (off-set) in the baseline at 333.4 K(+60.4 �C) (Aof). Upon cooling, the direct b1 fi b¢1martensitic transformation occurs from 324.4 K(+51.4 �C) (Ms) to 289.0 K (+16.0 �C) (Mf) with onsetat 320.7 K (+47.7 �C) (Mo), a peak at 315.0 K(+42.0 �C) (Mp) and off-set at 305.6 K (+32.6 �C)(Mof) in association with an exothermic process with anenthalpy of 8.8 J/g. These DSC results display criticalRMT intervals of 303.6 K (30.6 �C) (As-Af) and 308.4 K(35.4 �C) (Ms-Mf) similar to those reported by Silvaet al.[8] and Matlakhova et al.[14–16] in a Cu-13.5Al-4.0 Nialloy. However, the critical temperatures in the presentwork were comparatively higher than those As = 282 K(9 �C), Af = 325 K (52 �C), Ms = 298 K (25 �C), andMf = 264 K (�9 �C) shown in References 8, 14 through16. This is apparently an indication that small variationsin composition and structure may cause a sensible changein RMT parameters.Another point worth mentioning is that by consider-

ing the DSC curves, Figure 3, obtained at RT, i.e., thesame temperature at which XRD experiment wasconducted, the alloy is practically in its completemartensitic state. This is in agreement with Figure 2,which confirms that the as-received alloy structure ismainly composed of the martensitic b¢1 phase.The thermal hysteresis amplitude, DT, was evaluated

as the difference between the two critical peak temper-atures, Ap-Mp. According to Figure 3, this amplitudewas found as DT = 12.7 K (12.7 �C). The literatureindicates that the kind of RMT may be related to thethermal hysteresis amplitude.[1–4,6,8,14–16] For instance,in Cu-Al-Ni alloys the b1 M b¢1 transformation presentsDT around 10 �C,[11,12,14–16,21] which agree with thevalue of 12.7 K (12.7 �C) found in this work.TheOMobservation of the full circular cross section of

the as-received alloy is shown in Figure 4. This figuredisplays a characteristic geometrical pattern composed ofa central square block surrounded by four peripheralblocks. Details of these blocks reveal that the microstruc-ture in Figure 4 is composed of needles that are appar-ently thinner with a predominant single orientation in theperipheral blocks as compared to multiple orientations inthe central one. According to the literature,[1–3] theseneedles are b¢1 martensite variants. This result is con-firmed by the XRD analysis in Figure 2. The presence ofthe R phase is difficult to be observed by a specific surfacerelief inFigure 4, due to the fact thatR is coherentwithb¢1and intermediate between the high temperature (b1) andthe martensitic phases (b¢1).[7,14–16,19,20]The average value for more than ten microhardness

measurements was found as 302.9 ± 7.8 kgf/mm2 inthe cross section of the as-received alloy. This valueis comparable to others reported for similaralloys.[7,14–16,19,20] The relatively small standard devia-tion indicates that the martensitic microstructure isuniform in terms of mechanical properties throughoutthe different blocks seen in Figure 4.


B. Characterization After TCT

The following results shown in Figure 5 constitute thebasis of the present work regarding the Cu-13.7Al-4.2Nimonocrystalline alloy subjected to different TCT cycles(TCT’s) finishing with either: (a) half cooling (1/2 cool)or (b) half heating (1/2 heat) cycle, as shown in Figure 1.

1. XRD after TCT’s finishing with 1/2 cool cycleThe diffractograms shown in Figure 5(a) reveal the

alloy structural changes associated with TCT’s finishingwith 1/2 cool cycle, as the number of cycles is increased.After just one cycle (plus the 1/2 cool cycle), bycomparing to the XRD pattern of the as-received alloyin Figure 2, one notices the existence of other phasessuch as c¢1 with peaks (011)c¢1, (212)c¢1, and (111)c¢1 aswell as the high temperature b1 phase with its (311)b1peak. Already existing phases in Figure 2, such as R,with its (107)R peak as well as b¢1 with its (10�1�2)b¢1 and(20�2�4)b¢1 peaks are also seen in Figure 5(a). In addition,the coherent plane (21�2�1)b¢1|(0213)R peak, between b¢1and R phases is also observed in Figure 5(a) for one

Fig. 1—Schematic of the thermal cycling treatment.

Fig. 2—XRD patterns for the monocrystalline Cu-13.7Al-4.2Ni alloyin the initial state.

Fig. 3—RTM thermal effects during heating and cooling of themonocrystalline Cu-13.7Al-4.2Ni alloy for the initial state.

Fig. 4—Structural aspects of the monocrystalline alloy in the initialstate.


cycle. The b¢1 phase is the most evident in both Figures 2and 5(a).

After 100 cycles there is a tendency of increasing theparticipation of the b¢1 phase, owing to its greaterstability, as well as the coherent plane (21�2�1)b¢1 and(0213)b¢1 and the R phase. By contrast, a decreaseoccurred in the peaks corresponding to phases c¢1 andb1. Further accumulation of 200, 300, and 400 cyclesincreases even more the intensity of b¢1 peaks whiledecreasing all other peaks related to c¢1, R and b¢1|Rphases. In particular, the high temperature b1 peakalmost disappeared. After 500 cycles in Figure 5(a) asurprising decrease occurs in the intensity of the b¢1peaks, which causes its adjacent R peak to become moreevident. The reason for this apparent loss of stability ofthe b¢1 phase is not clear. One possibility is that aconsiderable number of cycles, such as 500, maypromote redistribution among the metastable phasesby increasing the relative participation of c¢1 and R atthe expense of b¢1. Moreover, an additional transfor-mation b¢1 M c¢1 might be occurring due to the stressaccumulated in the structure.[3]

2. XRD after TCT’s finishing with 1/2 heat cycleFigure 5(b) shows the sequence of XRD patterns with

the number of cycles of TCT’s finishing with 1/2 heatcycle. For one cycle (plus the 1/2 heat cycle), contrary tothe as-received alloy, Figure 2, and the TCT’s finishingwith 1/2 cool cycle, Figure 5(a), only a faint evidence ofb¢1, by its (10�1�2) peak, exists. No b¢1 (20�2�4) peak wasdetected. A relatively intense high temperature b1 peakis observed in Figure 5(b) together with c¢1 (011) and(111) peaks, not seen in the as-received alloy in Figure 2.The reason for the appearance of these phases in bothTCT’s finishing with 1/2 cool, Figure 5(a), or 1/2 heatcycle, Figure 5(b), after just one cycle is difficult toexplain. But it certainly indicates that even one RMTcycle is capable of significant microstructural changes inthe Cu-13.7Al-4.2Ni alloy. Another point worth men-tioning is the fact that the R phase, although with

relatively low intensity, is always present in theas-received alloy and in Figures 5(a) and (b) for anynumber of cycles. This could be an indication of itsstability during RMT, remaining as a residual phaseafter each cycle.With increasing number of cycles, from 100 to 400 in

Figure 5(b), the b¢1 peaks become increasingly intensewhile the high temperature b1 peak tends to vanish.However, for 500 cycles in Figure 5(b), similar to whathappened in Figure 5(a), the b¢1 peaks decrease inintensity probably due to phases redistribution. In fact,it is known that SME alloys, like the Cu-Al-Ni types, areable to undergo a single b1 M b¢1 or b1 M c¢1 ormultiple, b1 M c¢1+ b¢1 transformations depending onthe alloy concentration.[10–13] The difference in theenergy between these martensitic structures is very smalland the associated sliding of atomic layers easier tooccur. Consequently, the transformation from oneclose-packed martensitic plane into another is possibleto occur by simple redistribution.[3,22,23]

In previous works,[14–16,19–21,24–27] in which the influ-ence of TCT’s in monocrystalline Cu-Al-Ni was evalu-ated in an analogous way, it was reported that finishingwith either 1/2 cool or 1/2 heat cycle alters both thestability and the XRD intensity of existing phases. In thepresent work comparable results were obtained for theCu-13.7Al-4.2Ni alloy. Furthermore, after the first cycleup to 500 cycles, relatively few differences were found,Figures 5(a) and (b), depending on the way the TCT wasfinished. An important finding was the prominentparticipation of the b¢1 phase up to 400 cycles and itsreduction in intensity after 500 cycles.

C. DSC Analysis After TCT

Relevant results were obtained by means of DSCanalysis for TCT’s in comparison to the as-receivedCu-13.7Al-4.2Ni alloy. During the first DSC heating ofthe as-received alloy, Figure 3, only one peak associatedwith the RMT was observed. On the other hand,

Fig. 5—XRD patterns for the monocrystalline Cu-13.7Al-4.2Ni alloy after different thermal cycles, ending with 1/2 cycle of cooling from 258 Kto RT (�15 �C to RT) (a) and 1/2 cycle of heating from 373 K to RT (100 �C to RT) (b).


Figure 6 shows that TCT’s with more than one cycle(100 to 400) develop, after a first DSC run, a secondpeak around 373 K (100 �C), Figure 6(a), which tendsto disappear for the 500 cycles TCT. The reason for asecond DSC peak has not yet been discussed in theliterature. It is suggested that the different martensitictransformations associated with other metastable phasessuch as c¢1 and b¢1|R coherent plane are responsible forthe second DSC peak. In other words, the main DSCpeak around 327.7 K (54.7 �C) in Figure 3 is due to thereverse b¢1 M b1 austenitic transformation. However, afirst DSC run is not enough to completely transform theaustenitic alloy structure. Therefore, the partially trans-formed alloy presents a second peak slightly above373 K (100 �C), Figure 6(a), which could be related tob1 M c¢1 or b1 M b¢1|R. After 500 cycles, however, thedecrease in both b¢1 and b1, Figure 5(b), inhibits asecond DSC peak to take place in Figure 6(a). Thisbehavior can be understood by the analysis of theenthalpy values obtained from the DSC peaksin Figures 3 and 6(a). These values are plotted inFigure 6(b) as a function of the number of TCT cycles.In this figure the E1 curve corresponds to the lowertemperature [shortly over 323 K (50 �C)] main enthalpypeak; E2 the higher temperature [just over 373 K(100 �C)] second peak; and E1+2, the sum of both E1

plus E2 enthalpies. The reason for plotting E1+E2 is toshow that, for any number of cycles, this sum falls closeto the first heating enthalpy (9 J/g) of the as-receivedalloy. This is an indication that the second DSC peak inFigure 6(a) is actually a complement to the totalaustenitic transformation. Moreover, the absence of asecond peak for 500 cycles in Figure 6(a) can beattributed to a complete lower temperature (first mainpeak) transformation.

In case the DSC analysis is continued for each TCT (1to 500 cycles) following the first run, Figure 6(a), with a

cooling from 413 K to 223 K (140 �C to �50 �C), themartensitic peaks are observed for temperatures around313 K (40 �C) as shown in Figure 7(a). What came as asurprising result, was that still continuing the analysis bymeans of a second DSC run from 223 K to 413 K(�50 �C to 140 �C), only austenitic peaks are nowobserved around 328 K (55 �C) as shown in Figure 7(b).In other words, a second DSC run no longer displayspeaks slightly above 373 K (100 �C) that were observedin the first DSC heating run, Figure 6(a). It is suggestedthat this partial austenitic transformation, observed assecond peaks in the first heating run, gives rise to astable structure, which is no longer able to undergoRMT at higher [~373 K (100 �C)] temperatures.Based on the distinct phases existing in the as-received

alloy, Figure 2, and the transformations associated withTCT, Figure 5, together with DSC analyses in Figures 3and 6, it is proposed that a sequence of reaction steps:b1 M R and R M b¢1+ c¢1 is taking place. Moreover,additional direct transformations such as b1 M b¢1 andb¢1 M c¢1 might also occur, particularly at 500 cycles.

D. Critical RMT Temperatures After TCT

The influence of TCT’s on the RMT characteristics ofthe monocrystalline Cu-13.7Al-4.2Ni alloy was evalu-ated by means of the displacement of critical tempera-tures and intervals of transformations as shown inFigure 8. The results obtained in this figure are inagreement with those reported in the literature[3,9,17–21,26]

for alloys with comparable chemical compositions. Ingeneral, it was found that TCT’s with multiple RMTcycles, tend to decrease the critical temperatures asso-ciated with both cooling and heating transformations.These critical temperatures during cooling (Ms, Mo, Mp,Mof, Mf) and heating (As, Ao, Ap, Aof, Af) as well asrelated values of thermal hysteresis interval (DT) and

Fig. 6—RTM thermal effects during the first heating cycle (a) and enthalpy de RMT direct (b) of the monocrystalline Cu-13.7Al-4.2Ni alloy as afunction of the number of thermal cycles.


transformation enthalpy (DH) for 100 to 500 cycles ofTCT are presented in Figures 8 and 9.

In Figure 8, it is worth noticing that the peaktemperatures, Ap and Mp, are slightly displaced tolower values corresponding to reductions of up to 1.9 Kand 10.3 K (1.9 �C and 10.3 �C) after 500 cycles,respectively. As a consequence, the DT value which is12.7 K (12.7 �C) for the as-received condition isincreased to 21.1 K (21.1 �C) after 500 cycles as alsoshown in Figure 9(a). The literature[3,9,14–16,19–21,24–28]

indicates that defects, such as accommodation dislo-cations, stacking faults, and twin boundaries[17,28]

accumulated after each cycle, are responsible forincreasing DT by impairing the RMT. In particular,the density of dislocations increases with cycling until itreaches a limit.[28]

In addition to changes in the critical temperatures andDT, Figures 8 and 9, it was found an increase in the totalRMT interval, mainly after 500 cycles, as shown in

Figure 8. In fact, the critical intervals Mo-Mof andAo-Aof established after the first cycle were 15.1 K and9 K (15.1 �C and 9 �C), respectively. These intervalssignificantly increase with the number of cycles andreach 32.1 K and 15.9 K (32.1 �C and 15.9 �C), respec-tively, for 500 cycles.As for the value of DH, Figure 9(b) reveals only a

small decrease with increasing number of cycles for bothdirect, EM (upon cooling), and inverse EA (uponheating) transformation. This is an indication that thetotal enthalpy, i.e., the alloy thermodynamic state, ispractically not affected by either the number of TCTcycles or the direction of transformation, austenitic(heating), or martensitic (cooling).

E. OM Observation After TCT

The Cu-13.7Al-4.2Ni alloy structure after TCT isshown in Figure 10 finishing with 1/2 cool cycle. It is

Fig. 7—RTM thermal effects during cooling (a) and heating (b) in a second DSC run of the monocrystalline Cu-13.7Al-4.2Ni alloy for the initialstate and after thermal cycles.

Fig. 8—RTM critical temperatures obtained by DSC during cooling (a) and heating (b) of the monocrystalline Cu-13.7Al-4.2Ni alloy as a func-tion of the number of thermal cycles.


known that RMT introduces a martensite (plate, lath,lamella, or needle) relief on SME alloys mirror-polishedsurface.[1–3,14–16,19–21,24–27] For the present alloy this canbe clearly seen after one TCT cycle in Figure 10(a).Apparently not much difference exists between this onecycle relief structure and that observed by polarized lightfor the as-received alloy. In other words, the one cycleRMT is not sensibly changing the existing martensiticstructure. With minor differences, such as the appear-ance of c¢1, this one cycle OM result is in agreement withboth the XRD, Figures 2 and 5, and DSC, Figures 3and 6, results.

After 100 and 200 cycles, Figures 10(b) and (c), theTCT promotes the development of additional bandedrelief lines with a tendency of the right side, peripheralblock to disappear. This suggests that significantamounts of new phases were transformed and superim-posed to the existing ones, which also agree with XRD,Figure 5, and DSC, Figure 6, results. With largernumber of cycles, Figures 10((d) through (f)), a decreasein the type and amount of new transformed phases wasevidenced by the reduction of relief marks onto thepolished surface. After 300 cycles, Figure 10(d), thisreduction is particularly seen in the peripheral blocks, in

Fig. 9—Thermal hysteresis (a) and enthalpy of transformation direct (EM) and reverse (EA) (b) of the monocrystalline Cu-13.7Al-4.2Ni alloy as afunction of the number of thermal cycles.

Fig. 10—Structural aspects of the monocrystalline alloy after 1(a), 100(b), 200(c), 300(d), 400(e), and 500(f) thermal cycles that end with 1/2 cyclede cooling from 258 K to RT (�15 �C to RT).


which the relief lines, Figure 10(a) could be associatedwith the b¢1 martensite previously existing in the as-received alloy as shown in Figure 4. After 400 cycles,Figure 10(e), as well as 500 cycles, Figure 10(f), therelief caused by RMT barely appears onto the polishedsurface. This is somehow surprising since the XRDresults in Figure 5 reveal that, at least up to 400 cycles,the b¢1 martensite is increasing (peak intensity) with thenumber of cycles. A possible explanation could berelated to the interaction of multiple martensite variants.

It has been shown that multiple plate orientations fora given martensite phase may be obtained duringreversible transformation. Moreover, whenever an alloyis subjected to TCT, the self-accommodation betweenthe different martensite variants tends to cause bothshape modification and plate elimination with increasingnumber of cycles. As a consequence, relatively smallmacroscopic deformation is observed.[3,29,30] The presentresults in Figure 10 are in agreement with this macro-reduction-only mechanism since the surface relief for300, 400, and 500 TCT cycles are not as accentuated asfor smaller number of cycles, Figure 10, as well as forthe as-received structure in Figure 4, even though theamount of transformed martensite, Figures 5 and 6, wasnot reduced.

Another relevant result to be discussed is the multi-interface type of transformation, which occurs in theDSC changes observed in Figure 6(a) with thermalcycling. According to Ortın and Planes,[17] the numberand dimensions of interfaces changes all the time duringthe transformation. This is clearly seen in the sequenceof pictures with increasing number of cycles inFigure 10. Hence, the frictional force against interfacialmotion is almost continuously varying. A thermody-namic analysis, based on the chemical enthalpy changeand the frictional barrier opposing interfacial motion,indicates that a displacement towards lower valuesshould occur in the transformation temperatures.Indeed, it is apparently taking place in Figure 6(a),going from the initial to 500 cycles. Furthermore, anincrease in the frictional force during elastic accommo-dation of martensitic shape and volume is also expectedto occur with increasing number of cycles. This wouldcontribute to the broadening of DSC peaks, as alsoobserved in Figure 6(a). A physical rationale for thisbehavior might be related to the accumulation of defectswith each thermal cycle. Accommodation dislocations,for instance, not only retard the transformation to lowertemperatures but also extend the interval of thermalhysteresis.[3,9,14–16,19–21,24–27]

F. Microhardness After TCT

The microhardness evolution of the Cu-13.7Al-4.2Nialloy after increasing TCT cycles added relevant infor-mation for the understanding of related RMT.Figure 11 presents the average value of ten local testsrandomly performed in a cross section such as thoseshown in Figures 4 and 10, as a function of the numberof cycles. The microhardness tests were carried out insample after TCT finishing with 1/2 cool cycle, i.e., withthe XRD structure given in Figure 5(a). In Figure 11 it

is observed a tendency to decrease the microhardnessfrom the as-received alloy value of 303 kgf/mm2 down to250 kgf/mm2 after 200 cycles. Beyond this number ofcycles, there is a tendency of constant microhardnessvalues within the error (standard deviation) bars. Thisresult confirms that above 200 cycles, see Figure 10, themarks and interfaces are apparently reduced[3,9,14–16,18–21]

owing to self-accommodation of the numerous mar-tensite, most probably b¢1, variants. In fact, the rela-tively smoother surfaces in Figure 10 could indicate thatreduction of martensite plates was sensibly enough todecrease the microhardness as shown in Figure 11.Here it is important to mention that the micro-

hardness tests were performed at RT, i.e., within thecritical interval of RTM. This might induce transfor-mation upon application of the test load. As a conse-quence, individual values for a specific phase (b1, b¢1, c¢1,and R) may not represent the actual microhardnessowing to the phase metastability condition. The stressinduced transformation due to the applied load mightalter the indentation image, which then produces anunreliable result.

IV. FINAL REMARKS

A comparison between XRD patterns finishing with1/2 heating, Figure 5(b), i.e., RT fi 373 K fi RT(RT fi 100 �C fi RT) with the RMT critical tempera-tures, Figure 8(b), reveal a coherence of transformedphases. Indeed, the transformations occurring in theCu-13.7Al-4.2Ni alloy are complex and quite sensitive tostructural changes, Figure 10. An apparent indication ofthe coherence between existing phases in this alloy is theabsence of several DSC peaks, Figure 7, that werereported in TiNi and Cu-Al-Ni alloys[3,9–12,14–16,18–21,31]

The fact that only one RMT peak is observed after asequence of cycles, Figure 7, either cooling or heating,indicates that the monocrystalline Cu-13.7Al-4.2Nialloy is undergoing a complex sequence of transforma-tion, steps b1 M R and R M b¢1+ c¢1. It is also sug-gested that b1 M b¢1 and b¢1 M c¢1 reactions are alsooccurring with 500 cycles.

Fig. 11—Vickers microhardness of the monocrystalline Cu-13.7Al-4.2Ni alloy as a function of the number of thermal cycles.


Figure 8(a) shows that the RMT temperature Mof,determined after the first TCT cycle is above RT, thetemperature at which the XRD analysis was carried out.After this first heating cycle the presence of b1, R, b¢1,and c¢1 is revealed, Figure 5(b), with a more intense c¢1martensite. It is then suggested that the existence of c¢1in greater amounts occurs under the RMT finishing, i.e.,below Mof. With TCT accumulation up to 300 cycles,the value ofMof decreases, Figure 8, and approaches theroom temperature. This justifies a greater participationof the b¢1 martensite in the XRD patterns in Figure 5(b).For 400 and 500 cycles, the RMT finishes slightly aboveMof, which corresponds to the existence of b¢1 withlimited coherence with the R phase in Figure 5(b).

As shown in Figure 8(b), the room temperature,293 K ± 5 K (20 �C ± 5 �C), at which the XRD ana-lysis was conducted, does not reach the complete RMTcooling cycle. In fact, the critical values of As are wellbelow RT. For this reason, the high temperature b1phase is not found in considerable amount in the XRDpatterns in Figure 5.

In the present work, it was revealed that the mono-crystalline Cu-13.7Al-4.2Ni alloy may undergo a partialunstable transformation b1 M b¢1+ c¢1 with the appear-ance of others DSC peaks, Figure 6(a), slightly above373 K (100 �C). The literature[14–16,32–34] reported thatin other monocrystalline Cu-Al-Ni displaying ab1 M b¢1+ c¢1 RMT, the c¢1 martensite is graduallyinhibited as the number of cycles increases. Eventually,for a sufficiently larger number of cycles, c¢1 totallydisappears. In the present work, Figure 5 shows that,after the first cycle, the TCT tends to inhibit thedevelopment of the c¢1 martensite. In fact, the literatureindicates that a greater thermodynamic driving force(enthalpy) is needed for the b1 M c¢1 transformation ascompared to b1 M b¢1 transformation.[11,12,24] A possi-ble mechanism could be associated with structuralimperfections such as dislocations and twinning thataccumulates with the TCT cycles. These imperfectionsinterfere with the RMT completion not only by displac-ing Mf to lower temperatures, Figure 8(a), but also byinhibiting the development of the c¢1 martensite at RT,293 K ± 5 K, (20 �C ± 5 �C) which was the tempera-ture of the XRD patterns in Figure 5.

V. CONCLUSIONS

The RMT in the monocrystalline Cu-13.7Al-4.2Nialloy is a combined sequence of phase reaction stepsb1 M R and R M b¢1+ c¢1. Direct b1 M b¢1 andb¢1 M c¢1 transformations may also occur.

A TCT conducted for 1, 100, 200, 300, 400, and 500cycles displaces theDSCpeaks to lower temperatureswhileenlarges the total transformation interval. The TCT’s alsoinhibits the development of the c¢1 martensite. However,after 500 cycles, a b¢1 fi c¢1 transformationmay take placeand increase the relative participation of c¢1.

The structural modifications occurring in the alloyfinishing the TCT with half heating cycle,RT fi 373 K fi RT (RT fi 100 �C fi RT) revealedcoherence with RMT critical temperatures. The c¢1

martensite existence in relatively larger amounts occursat a temperature below Mof, associated with RTMfinishing on cooling. On the contrary, that for the b¢1martensite is close to Mof, which favors its largeramount, except after 500 cycles.The relatively small variation in both the critical

temperatures and the microhardness with the number ofTCT cycles, is an indication of the alloy resistance toirreversible structural changes caused by the RMT. Thisis a favorable condition for practical applications as ashape memory alloy.

ACKNOWLEDGMENTS

The authors thank the Brazilian agencies CNPq,CAPES, and FAPERJ for supporting this work. It isalso acknowledged the collaboration of Ivan Vahhi,Sergei Pulnev, and Luciana Lezira Pereira de Almeida.

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physical and structural characterization of a monocrystalline cu-13.7al-4.2ni alloy subjected to...

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