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Thermally induced fcc ↔ hcp martensitic transformation in Co–Ni
Acta Materialia 53 (2005) 3625–3634 www.actamat-journals.com
Thermally induced fcc M hcp martensitic transformation in Co–Ni Yinong Liu
a,* ,
Hong Yang a, Yan Liu a,b, Bohong Jiang c, Jun Ding d, Robert Woodward e
a
c
School of Mechanical Engineering, The University of Western Australia, Crawley, WA 6009, Australia b Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China Open Laboratory of the Ministry of Education for High-Temperature Materials and Testing, Shanghai Jiao Tong University, Shanghai 200030, China d Department of Materials Science, National University of Singapore, 10 Kent Ridge Crescent, Singapore 119260, Singapore e School of Physics, The University of Western Australia, Crawley, WA 6009, Australia Received 22 March 2005; accepted 12 April 2005 Available online 3 June 2005
Abstract This study investigated the thermal behaviour of the fcc M hcp martensitic transformation of Co–Ni alloys. The transformation was found to be incomplete in both directions upon cooling and heating. Upon thermal cycling, the transformation, as detected by differential scanning calorimetry, diminished rapidly, indicating that the transformation was highly irreversible. Optical microscopic observation revealed that surface reliefs appeared on pre-polished surfaces upon forward transformation, but did not disappear upon the reverse transformation, implying that the transformation also exhibited mechanical irreversibility. The irreversibility is explained on the basis of multiple choice of crystalline shuffling along 1=6h1 1 2if1 1 1gfcc for the fcc M hcp transformation. These observations suggest that the fcc M hcp martensitic transformation system in Co–Ni is not suitable for thermomechanical shape memory effect, which requires high reversibility of the transformation both crystallographically and mechanically. It is regarded that this rule is generic and applicable to other fcc M hcp transformation systems. 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Co–Ni; Martensitic phase transformation; Differential scanning calorimetry; X-ray diffraction; Transmission electron microscopy
1. Introduction Binary Co–Ni alloys of less than 35 wt.% Ni exhibit a face-centered cubic hexagonal close-packed (fcc M hcp) martensitic phase transformation [1–3]. The fcc phase at high temperature is commonly known as the c-phase whereas the hcp phase at low temperature is known as the e-phase. The transformation occurs in the ferromagnetic state, i.e., both phases are ferromagnetic within the vicinity of the transformation temperature. It is known that the transformation temperature decreases with increasing Ni content, whereas the thermal hysteresis of the transformation increases. When Ni content is *
Corresponding author. Tel.: +61 8 64883132; fax: +61 8 64881024. E-mail address: [email protected] (Y. Liu).
greater than 35 wt.%, no e-phase is generally formed upon cooling (to liquid nitrogen temperature). Being an alloy system exhibiting ferromagnetism and a thermoelastic martensitic transformation, Co–Ni alloys of less than 35 wt.% Ni content came to attention in the search for magnetically activated shape memory alloys [4]. Co–Ni alloys are characterised by high saturation magnetisation [5], high ductility stemming from their metallic solid-solution nature [6], and high corrosion resistance owing to the nobility of their constituent elements. Such a combination renders these alloys promising potentials to serve as functional shape memory materials in smart structure applications. A recent study reported the thermomechanical behaviour of the fcc M hcp martensitic transformation in a Co–32Ni alloy single crystal, and concluded that
1359-6454/$30.00 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2005.04.019
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the transformation system showed no thermally activated shape memory effect, due to multiplicity of the return path of the transformation [6,7]. This study continues the investigation by reporting on the thermal transformation behaviour of the Co–Ni alloys.
2. Experimental details
3.1. Single crystal Co–Ni alloy Fig. 1 shows the Co–Ni phase diagram. The two lines dividing c and e regions mark the start and the finish of the c ! e transformation on cooling. There is no information regarding the reverse transformation. It is seen that the temperature of the c ! e martensitic transformation decreases with increasing Ni content whereas the temperature window of the transformation in-
exothermic
creases. It is also seen that in the composition range of 26–34 wt.% Ni the phase diagram is tentative at below 200 C, although a mixture of c- and e-phases may be expected. Fig. 2 shows DSC measurement of the transformation behaviour of the Co–32Ni single crystal. The specimen was annealed at 1073 K. The measurement started by heating from room temperature. It is seen that no transformation was recorded upon the first heating (marked h1). Upon cooling, the specimen exhibited an exothermic transformation at 200 K (marked c1). The reverse endothermic transformation occurred at 460 K (peak temperature) on subsequent heating (h2), leaving a thermal hysteresis of 260 K. The forward transformation on cooling was abrupt whereas the reverse transformation occurred over a wider temperature range. The
Co-32Ni single crystal annealed at 1073 K
Heat Flow
3. Results
Fig. 1. Binary Co–Ni phase diagram.
c1 h1 h2
endothermic
Button-shaped ingots of Co–Ni alloys of 30 32 wt.% Ni were prepared from elemental Co bars of P99.95% purity and Ni flakes of P99.99% purity using a non-consumable electrode arc furnace in vacuum. The ingot was re-melted three times to ensure homogeneity. A cylinder-shaped single crystal of / 10 · 30 mm in dimension was grown from a Co–32Ni ingot by a crucibleless method using an optical floating-zone furnace (FZ-35WHV) in a high-purity Ar atmosphere. The asgrown single crystal was annealed at 1073 K for 7.2 ks. The single-crystallinity of the as-annealed specimen was confirmed by optical microscopic examination on the cross-section of the cylindrical ingot, using an etching solution of HF:HNO3:H2O = 1:4:5. The orientation of the single crystal was determined using X-ray back reflection Laue diffraction method. The polycrystalline buttons and single crystal cylinder were annealed in vacuum at 1073 K for 3.6 ks prior to cutting into specimens. Thermal transformation behaviour of the single crystal was studied using a Perkin-Elmer DSC4 differential scanning calorimeter (DSC). DSC measurement was performed in an Ar atmosphere with a heating/cooling rate of 10 K/min. X-ray diffraction (XRD) measurements were carried out on Co–30Ni polycrystalline samples using a Siemens D5000 diffractometer with a Cu Ka radiation. Optical microscopic observation was performed on polished samples without etching, so to better reveal surface relief effect of the martensitic transformation. Transmission electron microscopic (TEM) observations were made of the single crystal using a JEOL 3000F microscope operating at 300 kV. TEM samples were prepared by twin-jet electrochemical polishing in a 7 vol.% perchloric acid–methanol solution electrolyte at 253 K.
h3 h4 h5
200
300 400 Temperature (K)
500
600
Fig. 2. Thermal behaviour of c–e martensitic transformation of Co– 32Ni single crystal.
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heat effect of the reverse transformation was measured to be 1.25 J/g. It is interesting to note that the sharp exothermic peak on cooling was detected only in the first cycle in the as-annealed state and disappeared completely in subsequent thermal cycles. However, despite the absence of an obvious forward transformation on cooling, the endothermic peak associated with the reverse hcp ! fcc transformation on heating in subsequent thermal cycles was detected, as shown in the three curves labelled h2, h3 and h4. Prior to each of these measurements, the specimen had been cooled in liquid nitrogen. The heat intensity of the reverse transformation decreased progressively with cycling before disappearing completely after 5 cycles (h5). These measurements indicate that the fcc M hcp martensitic transformation is highly irreversible; it becomes stabilised after a small number of thermal cycles. To explore the possibility of recovery of the transformation after it had been stabilised by thermal cycling, the same single crystal sample shown in Fig. 2 was reheated to 1073 K for 3.6 ks for recovery and then its transformation behaviour was measured again. Fig. 3 shows the DSC measurement of the sample after three recovery
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anneals. It is seen that recovery anneal at 1073 K was effective in restoring the fcc M hcp martensitic transformation. In these figures, the curves marked ‘‘first heating’’ were recorded without cooling to liquid nitrogen. It is seen that a small endothermic event occurred at 400 K, indicating the existence of some hcp martensite at the ambient temperature. Similar to the original situation, the transformation was quickly stabilised after a few thermal cycles. In addition, there appeared to be a small increase in temperature for the reverse transformation with thermal cycling. 3.2. Polycrystalline Co–Ni alloy Similar thermal transformation behaviour was also observed in a Co–30Ni polycrystalline alloy. Fig. 4 shows DSC measurements of the reverse transformation on heating in consecutive thermal cycles. Prior to each of these measurements, the specimen had been cooled in liquid nitrogen. The absence of thermal event on cooling was confirmed by separate DSC measurement to 173 K. The dashed curve was the same measurement of the first curve labelled ‘‘first heating’’ for the as-annealed sample. Co-32Ni single crystal
Heat Flow
first heating after 1st LN2
Heat Flow
Co-32%Ni single crystal
(Y-scale enlarged 5 times for the dashed curve)
after 1st LN 2
after 3rd LN2
380
(a)
400
420
endothermic
endothermic
after 2nd LN2
(Y-scale enlarged 5 times for the dashed curve)
after 4th LN2
440 460 Temperature (K)
480
500
520
first heating
after 2nd LN 2
after 3rd LN 2 (Y-scale enlarged 5 times for the dashed curve)
380
400
420
(b)
after 4th LN 2
440 460 Temperature (K)
480
500
520
Co-32%Ni single crystal
Heat Flow
first heating after 1st LN2
endothermic
after 2nd LN2
after 3rd LN2 after 4th LN2
380
(c)
400
420
440 460 Temperature (K)
480
500
520
Fig. 3. Reverse e–c transformation of Co–32Ni single crystal: (a) after first recovery at 1073 K; (b) after second recovery at 1073 K; (c) after third recovery at 1073 K.
Y. Liu et al. / Acta Materialia 53 (2005) 3625–3634
Co-30Ni, Polycrystalline
(Y-scale enlarged 5 times for the dashed curve)
after 1st LN2 after 2nd LN2
endothermic
after 3rd LN2 after 4th LN2 after 5th LN2 after 6th LN2 after 7th LN2 300
350
400 450 500 Temperature (K)
550
Co-30Ni, Polycrystalline
(111)fcc (111)fcc (0002)hcp
X-ray Diffraction Intensity (arbitrary unit)
Heat Flow
first heating
(e) cooledin LN2 after2nd anneal
(200)fcc
(220)fcc
(d) after 2nd annealat 1073 K
(1010)hcp
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(1011)hcp
(c) annealedand cooledin LN2 (1012)hcp
(1120)hcp (1013)hcp
(b) as-annealed
600
(a) as-received
Fig. 4. Reverse e–c transformation of polycrystalline Co–30Ni. 40
A very mild reverse transformation was evident on the first heating from the room temperature. After cooling in liquid nitrogen, the reverse transformation emerged with a strong intensity at 485 K. In subsequent thermal cycles, the heat effect of the reverse transformation diminished with increasing number of the cycles whereas the temperature of the reverse transformation shifted slightly to higher temperatures. The transformation disappeared after seven cycles. The slightly higher transformation temperature compared to the single crystal specimen presented in Fig. 2 is consistent with the expectation for a lower Ni content according to the phase diagram. Fig. 5 shows XRD measurements of a Co–30Ni polycrystalline alloy. All measurements were carried out at the room temperature. Measurement (a) was conducted in the as-received condition, which was hot rolled. Spectrum (b) was recorded after anneal at 1073 K for 3.6 ks. Spectrum (c) was recorded after cooling in liquid nitrogen. Spectrum (d) was measured after a second anneal at 1073 K for 3.6 ks following eight thermal cycles between the liquid nitrogen temperature and 573 K, a temperature well above that of the reverse transformation (Fig. 4). Spectrum (e) was measured after a cooling in liquid nitrogen following the second anneal. Spectra (a), (b) and (c) are indexed to a mixture of the fcc c-phase and the hcp e-phase, whereas spectra (d) and (e) are fully indexed to the c-phase. The lattice parameter of the fcc c-phase is determined to be a = 0.353 nm. The lattice parameters of the hcp e-phase are determined to be a = 0.250 nm and c = 0.408 nm, giving a c/a ratio of 1.58. It is also easy to compute that this transformation is associated with nil volume change. The (1 1 1)fcc and (0 0 0 2)hcp diffraction peaks appear at the same location. Thus, the appearance and disap-
50
60
70
80
90
2θ (o)
Fig. 5. Effect of thermal cycling on structure of polycrystalline Co– 30Ni alloy.
pearance of the fcc c-phase and the hcp e-phase are best indicated by their second strongest peaks, (2 0 0)fcc and ð1 0 1 0Þhcp peaks, respectively. Spectra (a) and (b) indicate that at room temperature the alloy contained a mixture of c and e phases. It appeared that annealing at 1073 K encouraged the c ! e transformation upon cooling to room temperature, as evidenced by the relative increase of the intensity of the ð1 0 1 0Þhcp peak to that of the (1 1 1)fcc peak. Spectrum (c) indicates that after cooling in liquid nitrogen the specimen was also in a mixed state of c- and e-phases, although the relative fraction of the hcp e-phase had increased. It is to be noted that the room temperature is well below the temperature of the reverse transformation, as is evident in Fig. 3. Thus, the crystal structure at the liquid nitrogen temperature is expected to be preserved at room temperature. These measurements suggest that thermally induced transformation by cooling to liquid nitrogen temperature was incomplete. This implies that the heat effect measured by DSC cannot be taken as the latent heat of the transformation. An interesting observation is the phenomenon shown in spectra (d) and (e). After the eight prior thermal transformation cycles, the crystal structure is expected to be stabilised and the c M e martensitic transformation is expected to cease, as shown in Fig. 4. Annealing at 1073 K is expected to revert the stabilised structure to the fcc c-phase. In contrast to the mixed structure at room temperature after the initial anneal, the second anneal at the same temperature, 1073 K, after the eight thermal transformation cycles appeared to have stabilised the
Y. Liu et al. / Acta Materialia 53 (2005) 3625–3634
c-phase to the ambient temperature (spectrum (d)). Furthermore, the annealed c-phase was also able to survive the cooling in liquid nitrogen without transforming to the low temperature e-phase (spectrum (e)). This is apparently related to the prior thermal cycling of the specimen. This phenomenon is in sharp contrast to the case of single crystal specimen, where recovery anneals at 1073 K is effective in restoring the c M e martensitic transformation repeatedly. Fig. 6 shows XRD measurements of the effect of transformation cycling on the structure of the Co– 30Ni polycrystalline specimen. The spectra were all measured after cooling in liquid nitrogen, except the first measurement, which was in the as-annealed state, and the third measurement, which was recorded after heating to 573 K for 3.6 ks following the second measurement. Between each cooling in liquid nitrogen, the specimen was heated to 573 K for 3.6 ks. All the XRD spectra are indexed to a mixture of the fcc c-phase and the hcp e-phase, as labelled in the figure. The integrated intensities of the (1 1 1)fcc peak and the ð1 0 1 0Þhcp peak are shown in Fig. 7. The intensities have been normalised to the total combined integrated intensities of the two peaks. The two values on the dashed line at ‘‘1.5’’ cycle indicate the intensities of the sample
(111)fcc (0002)hcp
Co-30Ni, Polycrystalline annealed at 1073 K (1011)hcp (200)fcc
X-ray Diffraction Intensity (arbitrary unit)
8th LN2 7th LN2
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after the heating to 573 K. It is seen that the first cooling in liquid nitrogen caused a marked increase in the amount of the hcp e-phase at the expense of the fcc c-phase. Heating to 573 K after the first cooling reverted the forward c ! e transformation. The e-phase formed by the second cooling was less than that formed by the first cooling. Correspondingly, more residual c-phase was detected after the second cooling. The amounts of both phases appeared to be largely unchanged after each cooling in subsequent cycles, stabilising at values between the maximum and the minimum of the asannealed state and after the first cooling. These observations are consistent with the DSC measurement that the first thermal cycle was associated with a larger endothermic effect and the thermal effect diminished rapidly with cycling. The fact that the final amounts of the two phases after eight cycles appeared between their respective maximum and minimum of the as-annealed state and after the first cooling indicates that both the forward and the reverse transformations were incomplete during thermal cycling, leading to the gradual diminishing of the transformation. As evident in the results presented above, the c M e transformation is highly irreversible. To investigate the possibility of recovery of the transformation in the polycrystalline material, the stabilised samples were heated to elevated temperatures and their structure was determined by means of XRD. Fig. 8 shows the measurements for a Co–30Ni polycrystalline specimen in the as-received condition. All the measurements were conducted at room temperature after cooling in liquid nitrogen following each of the overheating treatments, except the first measurement, which was conducted in the asreceived condition. Overheating treatment was conducted for 3.6 ks in vacuum. The temperatures of the
6th LN2 5th LN2
3rd LN2 2nd LN2 573 K 1st LN2
X-ray Relative Integrated Intensity
4th LN2
0.7
0.6
(1011)hcp
0.5
(111)fcc 0.4
as-annealed 42
44
46
48
50
52
54
2θ (o) Fig. 6. Effect of thermal cycling on structure of polycrystalline Co– 30Ni alloy.
0.3 0
1
2
3 4 5 Transformation Cycle
6
7
8
Fig. 7. Effect of thermal cycling on structural evolution of polycrystalline Co–30Ni alloy.
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Co-30Ni, Polycrystalline, As-received (200)fcc
X-ray Diffraction Intensity (arbitrary unit)
(111)fcc (0002)hcp
1073 K 973 K 873 K 773 K 723 K (1011)hcp
648 K 623 K 573 K 1st cooling as-received
40
45
50
55
2θ (o) Fig. 8. Effect of overheating on the transformation behaviour of polycrystalline Co–30Ni alloy.
treatment are indicated in the figure. It is seen that the as-received specimen contained a mixture of and e phases. Cooling in liquid nitrogen resulted in an increase in the fraction of the e-phase. Overheating to 573 K, 70 K above the finishing temperature of the reverse transformation (Fig. 4), led to the formation of less e-phase upon cooling to liquid nitrogen. This is obviously due to the effect of transformation cycling (Fig. 6). One clear difference was observed upon heating to 623 K, 120 K above the finishing temperature of the reverse transformation. However, upon overheating to 648 K, the structure was apparently totally reverted to fcc and no hcp e-phase was formed at the liquid nitrogen temperature. It needs to be emphasised that these XRD spectra were recorded after cooling in liquid nitrogen following each overheating, instead of at the room temperature after the heating. This means that the overheating at 648 K not only caused a full reversion of the e-phase to the c-phase, but also stabilised the c-phase to the liquid nitrogen temperature. It is evident that even heating to 1073 K, the original annealing temperature, was ineffective in restoring the c M e transformation. 3.3. Microscopic evidences Fig. 9 shows a low-magnification optical image of the surface of a Co–32Ni single crystal specimen. The
Fig. 9. Surface relief of the c M e martensitic transformation in Co– 32Ni single crystal.
viewing surface was the (1 1 1)fcc plane of the single crystal. The specimen was pre-polished at room temperature in the as-annealed state and then cooled in liquid nitrogen. The micrographs were taken at room temperature after the cooling. The micrographs reveal clear surface reliefs. The surface reliefs appear to have three distinctive morphologies. Some parts exhibited long straight reliefs across the entire specimen, as indicated by the arrows in regions A. In other parts the surface reliefs exhibited equilateral triangular morphology of various sizes, as indicated by the triangle in region B. In region C, as indicated by the circle, the surface reliefs were small in dimension and less regular in shape. However, all these reliefs appeared to conform to three basic variants following the three sides of the triangle, which are indexed to Æ1 1 0æ directions of the fcc structure. This is clearly a manifestation of the expansion of stacking faults along h1 1 2if1 1 1gfcc systems, which is the mechanism of the fcc ! hcp martensitic
Y. Liu et al. / Acta Materialia 53 (2005) 3625–3634
transformation. It is also evident that there existed large areas, as large regions or between surface relief lines, which remained unaffected by the transformation, suggesting that the fcc ! hcp transformation was localised and incomplete. Fig. 10 shows a series of optical micrographs of region D, which is marked by the rectangle, in Fig. 9. All micrographs were taken at room temperature and are presented in their chronicle order. Micrograph (a) was taken after the pre-polish in the as-annealed state. Micrograph (b) was taken in the e-martensite state after the first cooling in liquid nitrogen (the same condition as the micrographs shown in Fig. 9). Micrograph (c) was taken after a heating to 533 K, some 50 K above the completion of the reverse transformation, following micrograph (b). It is seen that, whereas a reverse transformation had occurred upon heating to 533 K, the surface relief did not disappear; instead, it became more extensive. Micrograph (d) was taken after a re-polish in the c-austenite state following micrograph (c). Micrograph (e) was taken after a second cooling in liquid nitrogen after the re-polish. It is seen that the same martensite variants to those appeared in micrograph (b) were formed again, even the small irregularity at the lower left hand corner. At the same time, the extent of the surface relief was much reduced compared to that after the first cooling (micrograph (b)). Micrograph (f) was taken after a third polish in the e-martensite state. Micrograph (g) was taken after a heating to 533 K following micrograph (f). It is seen that exactly the same variants were formed by this reverse transformation compared to those formed on the forward transformation (Figs. (b) and (e)), but to a much less extent. Micrograph (h) shows the surface after a third polish in the c-austenite state and micrograph (i) shows the surface relief after cooling in liquid nitrogen flowing the polish. The number of surface relief lines was drastically reduced. This is obviously due to the effect of transformation cycling. Fig. 11 shows transmission electron micrographs of a Co–32Ni single crystal specimen, with (a) showing the specimen in the as-annealed state and (b) showing the microstructure after chilling in liquid nitrogen. Both micrographs were taken at ambient temperature. It is seen that the as-annealed state contained a large number of stacking faults. This is consistent with previous observations reported in the literature [8,9] as well as the XRD measurement of the co-existence of fcc and hcp phases in the as-annealed state. It is also evident in the image that the stacking faults were extensively interlinked between different fault systems. The microstructure of the as-chilled specimen showed large areas of untransformed fcc phase. This is consistent with the XRD observation of the incomplete transformation in either direction.
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Fig. 10. Effect of pre-polish and transformation cycling on the surface relief of the c M e martensitic transformation in Co–32Ni single crystal.
4. Discussions 4.1. Transformation mechanism The fcc M hcp martensitic transformation in Co–Ni alloy system occurs by the expansion and shrinkage of stacking faults via the movement of a6 h1 1 2i Shockley partial dislocations along the {1 1 1}fcc systems [3,8]. Activation of one such partial dislocation only creates or expands a stacking fault. For a hcp martensite of finite thickness to form a coordinated activation of
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Fig. 11. Transmission electron micrographs of the microstructure of Co–32Ni single crystal at room temperature: (a) in as-annealed state, (b) after chilling in liquid nitrogen.
Æ1 1 2æ gliding on every second (1 1 1) plane is required [8,10,11]. In the fcc structure there are 12 such systems, with three equivalent a6 h1 1 2i directions on each of the four {1 1 1} planes. It is important to note that the three a h1 1 2i partial dislocation systems on each (1 1 1) plane 6 produce exactly the same variant of the hcp martensite, despite the difference in siding direction. Therefore, there are in total four martensite variants possible from a fcc single crystal. Whereas a coordinated sliding on every second (1 1 1) plane is required, there is no restriction to the selection of sliding among the three h1 1 2i directions on each plane. There exist three possible combinations of sliding directions: (1) all mobile Shockley partial dislocations that contribute to the transformation have the same Burgers vector and slide in the same ½1 1 2 direction (bi = b1), (2) all three directions are equally selected, P P P giving the statistical condition that b1 + b2 + b3 = 0, (3) a random P or P P selection among the three so that b1 + b2 + b3 6¼ 0. Condition (1) has been referred to in the literature as mode A, condition (2) as mode B and condition (3) as mode C [3]. It is obvious that for mode B the transformation produces no crystal shape distortion beyond the scale of a few atomic planes. For mode A, it is easy to compute that, for a standard hcp structure of c/a = 1.63, the transformation will produce a maximum shear strain of 35% in the chosen ½1 1 2 direction on the chosen (1 1 1) plane. For condition (3), the crystal will experience some long-range shape distortion. It is important to note that all the three mechanisms produce exactly the same variant of the hcp martensite, the only difference in the product being the external shape of the crystal. The actual mechanisms of the fcc M hcp martensitic transformation in Co–Ni alloys have been studied previously by means of TEM [3,8,9,12], optical microscopy [2,3], atomic force microscopy [3], neutron scattering
spectrometry [10] and thermomechanical testing [6,7]. Waitz and Karnthaler [3] suggested that in short-range order the transformation occurs by mode A mechanism, producing a net shape change, whereas in long-range order internal stresses created by a mode A transformation in one location induce other mode A transformations in neighbouring regions to facilitate a self-accommodation of the martensite, yielding a nil net shape change. Hayzelden et al. [8] concluded from their TEM study that the hcp martensite does not form in self-accommodation mode and that the transformation, at the local environment of individual partial dislocations, is highly reversible. There is no specific determination of the Burgers vectors of the partial dislocations contributing to the transformation in this study. However, the optical microscopic evidences presented in Figs. 9 and 10 appear to suggest that the transformation occurs by the mechanism of mode C – there is no specific coordination of the Shockley partial dislocation movement. The sliding is highly random, both within one (1 1 1) variant system and among all Æ1 1 1æ variant systems; and both in the short spatial range and over large distances. It is also important to note that whereas the lattice distortion of the transformation can be perfectly self-accommodated within one Æ1 1 1æ variant, there exists no accommodation mechanism among the four different Æ1 1 1æ variants. 4.2. Latent heat of the transformation In the particular case of Co–Ni, the lattice distortion of the transformation is pure shear and involves no change in the lattice parameters. This renders the transformation of minimal lattice reshuffling. For this type of first order transformations the latent heat is expected to be small. In this work the heat effect measured by DSC is in the order of 1 J/g at its maximum during the initial transformation cycles. Even considering the fact that
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some hcp phase existed prior to the transformation and that the transformation associated with the thermal event was incomplete, the total heat of the transformation is not expected to be greater than several joules per gram. Hayzelden et al. [8] reported latent heats of 6.5–8.5 J/g for the martensitic transformation in a small-grained polycrystalline Co–10Ni alloy. The latent heat of the transformation in Co–32Ni alloy is expected to be lower. 4.3. Thermal reversibility It is interesting to note that despite the pre-existence of the stacking faults at ambient temperature, as evidenced by the XRD measurements and TEM observation, the forward transformation occurred in an abrupt manner during the first cooling. This implies that the pre-existing Shockley partial dislocations are not mobile and in a low free energy state. The abrupt occurrence of the fcc ! hcp transformation upon cooling is a manifestation of a sudden activation of the partial dislocations that had been pinned, in analogue to the case of activation of dislocations during the Lu¨ders deformation of mild steels, where a higher critical driving force is required to trigger the process than to propagate the process. This is consistent with the expectation that after full anneal at 1073 K the partial dislocations (and the expansion between each pair) stay the least energised state, or the most thermodynamically stable state. Another interesting aspect is that the forward c ! e transformation was observed only during the first cooling from the as-annealed state. In subsequent thermal cycles similar exothermic transformation peak was not detected by DSC down to 173 K. Evidenced by the detection of the reverse transformation on heating, apparently the forward transformation occurred over a wide temperature range at extremely low intensity, which is beyond the detection sensitivity of the DSC. The occurrence of the broad forward transformation is believed to be due to the liberation of the partial dislocations from the initial pinning effect of impurities and structural defects that renders the partial dislocations a wide spread in terms of mobility, or critical driving force for movement. The DSC measurements indicate clearly that the fcc M hcp martensitic transformation is highly irreversible. The structure stabilises and the transformation diminishes rapidly during thermal transformation cycling. This is confirmed also by the optical microscopic observation shown in Fig. 10. The gradual reduction in the transformation volume is an indication of the reduction of the population of mobile partial dislocations. This reduction is attributed to an arrest effect of the partial dislocations by structural defects during transformations. During transformations the mobile partial dislocations move forth and back in the matrix,
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and are arrested when fall into low-energy traps, such as structural defects or free surfaces. The XRD measurements confirm that the transformation stabilises during thermal transformation cycling towards a mixed state of fcc c-phase and hcp e-phase. This is consistent with the hypothesis of arrest of partial dislocations. This observation is in contradiction to the claim by Hayzelden et al. [8] that the transformation is microscopically reversible without significant hysteresis. Another interesting observation is the clear difference between the single crystal and the polycrystalline materials in terms of the recoverability of the transformation after stabilisation by thermal cycling. The single crystal specimen was revived repeatedly by recovery anneal at 1073 K whereas the polycrystalline materials was permanently stabilised after cycling. Given the fact that the polycrystalline material had a lower Ni content and thus its fcc phase is less stable than for the single crystal specimen, this observation is of higher certainty. This is probably related to the pinning effect of grain boundaries, which arrest the mobile partial dislocations beyond recovery at 1073 K. 4.4. Mechanical reversibility The optical microscopy observations show that transformation-induced surface relief prevails in subsequent reverse transformations. This is consistent with previous observations [2]. This demonstrates that, in addition to the thermodynamic irreversibility, the fcc M hcp martensitic transformation is also mechanically irreversible. This is schematically expressed in Fig. 12. The figure shows a (1 1 1)fcc plane of the fcc austenite, which coincides with the (0 0 0 1)hcp of the hcp martensite. The
d
b
a
c
[1 10]
[121]
(111)fcc – (0002)hcp Fig. 12. Schematic illustration of mechanisms of fcc M hcp martensitic transformation, demonstrating the multiplicity of transformation path for the forward and the reverse processes.
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edges of the triangles are the Æ1 1 0æfcc directions and the arrows indicate the Æ1 1 2æfcc directions. The open circles represent atomic positions in the fcc phase whereas the closed circle marks the position for the hcp martensite. Upon cooling, the fcc atoms have three equivalent possibilities to shift into hcp martensitic position by lattice shear along the Æ1 1 2æfcc directions. Upon the reverse transformation, the hcp atoms also have the same three possibilities to return to the fcc austenite by the same lattice shear, i.e. the constraint for the same return path is absent. This is obviously responsible for the prevalence of the surface relief upon transformation reversion. It ought to be pointed out despite the multiple choices of the return path for each martensite variant, they all return to the same original fcc crystal. This implies that the fcc single crystal remains a single crystal after thermal transformation cycles. It is also obvious that this non-uniqueness of the return path for the transformation disqualifies the alloy as a thermomechanical shape memory alloy.
5. Conclusions The experimental evidence generated by the different techniques of DSC, XRD, optical microscopy and TEM is highly consistent. Based on this evidence the following conclusions may be drawn regarding the fcc M hcp martensitic transformation in Co–Ni: (1) The alloys are in a mixed fcc–hcp state at ambient temperature after annealing, the microstructure contains a high population of stacking faults, and the stacking faults are extensively interlinked between different Shockley partial dislocation systems. (2) The fcc M hcp transformation in Co–Ni alloys of 30 wt.% Ni is highly irreversible. It is characterised by a large thermal hysteresis. It stabilises quickly through thermal transformation cycling. Both the forward fcc ! hcp transformation and the reverse transformation are highly irreversible and the transformation stabilises with a mixed structure of fcc and hcp. Annealing at 1073 K is effective in annihilating the stabilisation effect for Co–32Ni single crystal alloy, but ineffective for the Co–30Ni polycrystalline alloy. The latent heat of the transformation is very small.
(3) The transformation is crystallographically reversible, in that it always returns to the same fcc crystal upon the reverse transformation, despite the retention of some hcp phase at temperatures above the completion of the reverse transformation. In other words, no multiple variants may be generated through thermal transformation cycling. (4) The transformation is mechanically irreversible. Whereas the original fcc crystal structure and orientation is always re-achieved after a thermal transformation cycle, the shape of the fcc crystal can be changed. This is particularly the case when the influence of a mechanical bias is present during the process of transformation. This disqualifies the fcc–hcp martensitic transformation in Co–Ni from serving as a mechanism for the thermomehanical shape memory effect. It can obviously be postulated that all fcc–hcp martensitic transformations are invalid for the thermomechanical shape memory effect.
Acknowledgement This work is financially supported by ARC Discovery Project Grant DP0345880.
References [1] Hansen M, Anderko K. Constitution of binary alloys. New York (NY): McGraw-Hill; 1958. p. 485. [2] Gaunt P, Christian JW. Acta Metall 1959;7:529. [3] Waitz T, Karnthaler HP. Acta Mater 1997;45:837. [4] Liu Y, Zhou W, Jiang B, Qi X, Liu Y, Wang J, et al. J Phys IV 2003;112:1013. [5] Kong HZ, Ding J, Wee ATS, Liu Y, Li Y. Int J Nanosci 2004;3:615. [6] Liu Y, Yang H, Tan G, Miyazaki S, Jiang B, Liu Y. J Alloys Compd 2004;368:157. [7] Liu Y, Tan G, Miyazaki S, Liu Y, Jiang B. J Phys IV 2003;112:1025. [8] Hayzelden C, Chattopadhyay K, Barry JC, Cantor B. Philos Mag A 1991;63:461. [9] Hitzenberger C, Karnthaler HP, Korner A. Acta Metall 1985;33:1293. [10] Blaschko O, Krexner G, Pleschiutschnig J, Ernst G, Hitzenberger C, Karnthaler HP, et al. Phys Rev Lett 1988;60:2800. [11] Christian JW. Dislocations and properties of real materials. London: The Institute of Metals; 1985. p. 94. [12] Hitzenberger C, Karnthaler HP. Philos Mag A 1991;64:151.
Thermally induced fcc M hcp martensitic transformation in Co–Ni Yinong Liu
a,* ,
Hong Yang a, Yan Liu a,b, Bohong Jiang c, Jun Ding d, Robert Woodward e
a
c
School of Mechanical Engineering, The University of Western Australia, Crawley, WA 6009, Australia b Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China Open Laboratory of the Ministry of Education for High-Temperature Materials and Testing, Shanghai Jiao Tong University, Shanghai 200030, China d Department of Materials Science, National University of Singapore, 10 Kent Ridge Crescent, Singapore 119260, Singapore e School of Physics, The University of Western Australia, Crawley, WA 6009, Australia Received 22 March 2005; accepted 12 April 2005 Available online 3 June 2005
Abstract This study investigated the thermal behaviour of the fcc M hcp martensitic transformation of Co–Ni alloys. The transformation was found to be incomplete in both directions upon cooling and heating. Upon thermal cycling, the transformation, as detected by differential scanning calorimetry, diminished rapidly, indicating that the transformation was highly irreversible. Optical microscopic observation revealed that surface reliefs appeared on pre-polished surfaces upon forward transformation, but did not disappear upon the reverse transformation, implying that the transformation also exhibited mechanical irreversibility. The irreversibility is explained on the basis of multiple choice of crystalline shuffling along 1=6h1 1 2if1 1 1gfcc for the fcc M hcp transformation. These observations suggest that the fcc M hcp martensitic transformation system in Co–Ni is not suitable for thermomechanical shape memory effect, which requires high reversibility of the transformation both crystallographically and mechanically. It is regarded that this rule is generic and applicable to other fcc M hcp transformation systems. 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Co–Ni; Martensitic phase transformation; Differential scanning calorimetry; X-ray diffraction; Transmission electron microscopy
1. Introduction Binary Co–Ni alloys of less than 35 wt.% Ni exhibit a face-centered cubic hexagonal close-packed (fcc M hcp) martensitic phase transformation [1–3]. The fcc phase at high temperature is commonly known as the c-phase whereas the hcp phase at low temperature is known as the e-phase. The transformation occurs in the ferromagnetic state, i.e., both phases are ferromagnetic within the vicinity of the transformation temperature. It is known that the transformation temperature decreases with increasing Ni content, whereas the thermal hysteresis of the transformation increases. When Ni content is *
Corresponding author. Tel.: +61 8 64883132; fax: +61 8 64881024. E-mail address: [email protected] (Y. Liu).
greater than 35 wt.%, no e-phase is generally formed upon cooling (to liquid nitrogen temperature). Being an alloy system exhibiting ferromagnetism and a thermoelastic martensitic transformation, Co–Ni alloys of less than 35 wt.% Ni content came to attention in the search for magnetically activated shape memory alloys [4]. Co–Ni alloys are characterised by high saturation magnetisation [5], high ductility stemming from their metallic solid-solution nature [6], and high corrosion resistance owing to the nobility of their constituent elements. Such a combination renders these alloys promising potentials to serve as functional shape memory materials in smart structure applications. A recent study reported the thermomechanical behaviour of the fcc M hcp martensitic transformation in a Co–32Ni alloy single crystal, and concluded that
1359-6454/$30.00 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2005.04.019
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the transformation system showed no thermally activated shape memory effect, due to multiplicity of the return path of the transformation [6,7]. This study continues the investigation by reporting on the thermal transformation behaviour of the Co–Ni alloys.
2. Experimental details
3.1. Single crystal Co–Ni alloy Fig. 1 shows the Co–Ni phase diagram. The two lines dividing c and e regions mark the start and the finish of the c ! e transformation on cooling. There is no information regarding the reverse transformation. It is seen that the temperature of the c ! e martensitic transformation decreases with increasing Ni content whereas the temperature window of the transformation in-
exothermic
creases. It is also seen that in the composition range of 26–34 wt.% Ni the phase diagram is tentative at below 200 C, although a mixture of c- and e-phases may be expected. Fig. 2 shows DSC measurement of the transformation behaviour of the Co–32Ni single crystal. The specimen was annealed at 1073 K. The measurement started by heating from room temperature. It is seen that no transformation was recorded upon the first heating (marked h1). Upon cooling, the specimen exhibited an exothermic transformation at 200 K (marked c1). The reverse endothermic transformation occurred at 460 K (peak temperature) on subsequent heating (h2), leaving a thermal hysteresis of 260 K. The forward transformation on cooling was abrupt whereas the reverse transformation occurred over a wider temperature range. The
Co-32Ni single crystal annealed at 1073 K
Heat Flow
3. Results
Fig. 1. Binary Co–Ni phase diagram.
c1 h1 h2
endothermic
Button-shaped ingots of Co–Ni alloys of 30 32 wt.% Ni were prepared from elemental Co bars of P99.95% purity and Ni flakes of P99.99% purity using a non-consumable electrode arc furnace in vacuum. The ingot was re-melted three times to ensure homogeneity. A cylinder-shaped single crystal of / 10 · 30 mm in dimension was grown from a Co–32Ni ingot by a crucibleless method using an optical floating-zone furnace (FZ-35WHV) in a high-purity Ar atmosphere. The asgrown single crystal was annealed at 1073 K for 7.2 ks. The single-crystallinity of the as-annealed specimen was confirmed by optical microscopic examination on the cross-section of the cylindrical ingot, using an etching solution of HF:HNO3:H2O = 1:4:5. The orientation of the single crystal was determined using X-ray back reflection Laue diffraction method. The polycrystalline buttons and single crystal cylinder were annealed in vacuum at 1073 K for 3.6 ks prior to cutting into specimens. Thermal transformation behaviour of the single crystal was studied using a Perkin-Elmer DSC4 differential scanning calorimeter (DSC). DSC measurement was performed in an Ar atmosphere with a heating/cooling rate of 10 K/min. X-ray diffraction (XRD) measurements were carried out on Co–30Ni polycrystalline samples using a Siemens D5000 diffractometer with a Cu Ka radiation. Optical microscopic observation was performed on polished samples without etching, so to better reveal surface relief effect of the martensitic transformation. Transmission electron microscopic (TEM) observations were made of the single crystal using a JEOL 3000F microscope operating at 300 kV. TEM samples were prepared by twin-jet electrochemical polishing in a 7 vol.% perchloric acid–methanol solution electrolyte at 253 K.
h3 h4 h5
200
300 400 Temperature (K)
500
600
Fig. 2. Thermal behaviour of c–e martensitic transformation of Co– 32Ni single crystal.
Y. Liu et al. / Acta Materialia 53 (2005) 3625–3634
heat effect of the reverse transformation was measured to be 1.25 J/g. It is interesting to note that the sharp exothermic peak on cooling was detected only in the first cycle in the as-annealed state and disappeared completely in subsequent thermal cycles. However, despite the absence of an obvious forward transformation on cooling, the endothermic peak associated with the reverse hcp ! fcc transformation on heating in subsequent thermal cycles was detected, as shown in the three curves labelled h2, h3 and h4. Prior to each of these measurements, the specimen had been cooled in liquid nitrogen. The heat intensity of the reverse transformation decreased progressively with cycling before disappearing completely after 5 cycles (h5). These measurements indicate that the fcc M hcp martensitic transformation is highly irreversible; it becomes stabilised after a small number of thermal cycles. To explore the possibility of recovery of the transformation after it had been stabilised by thermal cycling, the same single crystal sample shown in Fig. 2 was reheated to 1073 K for 3.6 ks for recovery and then its transformation behaviour was measured again. Fig. 3 shows the DSC measurement of the sample after three recovery
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anneals. It is seen that recovery anneal at 1073 K was effective in restoring the fcc M hcp martensitic transformation. In these figures, the curves marked ‘‘first heating’’ were recorded without cooling to liquid nitrogen. It is seen that a small endothermic event occurred at 400 K, indicating the existence of some hcp martensite at the ambient temperature. Similar to the original situation, the transformation was quickly stabilised after a few thermal cycles. In addition, there appeared to be a small increase in temperature for the reverse transformation with thermal cycling. 3.2. Polycrystalline Co–Ni alloy Similar thermal transformation behaviour was also observed in a Co–30Ni polycrystalline alloy. Fig. 4 shows DSC measurements of the reverse transformation on heating in consecutive thermal cycles. Prior to each of these measurements, the specimen had been cooled in liquid nitrogen. The absence of thermal event on cooling was confirmed by separate DSC measurement to 173 K. The dashed curve was the same measurement of the first curve labelled ‘‘first heating’’ for the as-annealed sample. Co-32Ni single crystal
Heat Flow
first heating after 1st LN2
Heat Flow
Co-32%Ni single crystal
(Y-scale enlarged 5 times for the dashed curve)
after 1st LN 2
after 3rd LN2
380
(a)
400
420
endothermic
endothermic
after 2nd LN2
(Y-scale enlarged 5 times for the dashed curve)
after 4th LN2
440 460 Temperature (K)
480
500
520
first heating
after 2nd LN 2
after 3rd LN 2 (Y-scale enlarged 5 times for the dashed curve)
380
400
420
(b)
after 4th LN 2
440 460 Temperature (K)
480
500
520
Co-32%Ni single crystal
Heat Flow
first heating after 1st LN2
endothermic
after 2nd LN2
after 3rd LN2 after 4th LN2
380
(c)
400
420
440 460 Temperature (K)
480
500
520
Fig. 3. Reverse e–c transformation of Co–32Ni single crystal: (a) after first recovery at 1073 K; (b) after second recovery at 1073 K; (c) after third recovery at 1073 K.
Y. Liu et al. / Acta Materialia 53 (2005) 3625–3634
Co-30Ni, Polycrystalline
(Y-scale enlarged 5 times for the dashed curve)
after 1st LN2 after 2nd LN2
endothermic
after 3rd LN2 after 4th LN2 after 5th LN2 after 6th LN2 after 7th LN2 300
350
400 450 500 Temperature (K)
550
Co-30Ni, Polycrystalline
(111)fcc (111)fcc (0002)hcp
X-ray Diffraction Intensity (arbitrary unit)
Heat Flow
first heating
(e) cooledin LN2 after2nd anneal
(200)fcc
(220)fcc
(d) after 2nd annealat 1073 K
(1010)hcp
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(1011)hcp
(c) annealedand cooledin LN2 (1012)hcp
(1120)hcp (1013)hcp
(b) as-annealed
600
(a) as-received
Fig. 4. Reverse e–c transformation of polycrystalline Co–30Ni. 40
A very mild reverse transformation was evident on the first heating from the room temperature. After cooling in liquid nitrogen, the reverse transformation emerged with a strong intensity at 485 K. In subsequent thermal cycles, the heat effect of the reverse transformation diminished with increasing number of the cycles whereas the temperature of the reverse transformation shifted slightly to higher temperatures. The transformation disappeared after seven cycles. The slightly higher transformation temperature compared to the single crystal specimen presented in Fig. 2 is consistent with the expectation for a lower Ni content according to the phase diagram. Fig. 5 shows XRD measurements of a Co–30Ni polycrystalline alloy. All measurements were carried out at the room temperature. Measurement (a) was conducted in the as-received condition, which was hot rolled. Spectrum (b) was recorded after anneal at 1073 K for 3.6 ks. Spectrum (c) was recorded after cooling in liquid nitrogen. Spectrum (d) was measured after a second anneal at 1073 K for 3.6 ks following eight thermal cycles between the liquid nitrogen temperature and 573 K, a temperature well above that of the reverse transformation (Fig. 4). Spectrum (e) was measured after a cooling in liquid nitrogen following the second anneal. Spectra (a), (b) and (c) are indexed to a mixture of the fcc c-phase and the hcp e-phase, whereas spectra (d) and (e) are fully indexed to the c-phase. The lattice parameter of the fcc c-phase is determined to be a = 0.353 nm. The lattice parameters of the hcp e-phase are determined to be a = 0.250 nm and c = 0.408 nm, giving a c/a ratio of 1.58. It is also easy to compute that this transformation is associated with nil volume change. The (1 1 1)fcc and (0 0 0 2)hcp diffraction peaks appear at the same location. Thus, the appearance and disap-
50
60
70
80
90
2θ (o)
Fig. 5. Effect of thermal cycling on structure of polycrystalline Co– 30Ni alloy.
pearance of the fcc c-phase and the hcp e-phase are best indicated by their second strongest peaks, (2 0 0)fcc and ð1 0 1 0Þhcp peaks, respectively. Spectra (a) and (b) indicate that at room temperature the alloy contained a mixture of c and e phases. It appeared that annealing at 1073 K encouraged the c ! e transformation upon cooling to room temperature, as evidenced by the relative increase of the intensity of the ð1 0 1 0Þhcp peak to that of the (1 1 1)fcc peak. Spectrum (c) indicates that after cooling in liquid nitrogen the specimen was also in a mixed state of c- and e-phases, although the relative fraction of the hcp e-phase had increased. It is to be noted that the room temperature is well below the temperature of the reverse transformation, as is evident in Fig. 3. Thus, the crystal structure at the liquid nitrogen temperature is expected to be preserved at room temperature. These measurements suggest that thermally induced transformation by cooling to liquid nitrogen temperature was incomplete. This implies that the heat effect measured by DSC cannot be taken as the latent heat of the transformation. An interesting observation is the phenomenon shown in spectra (d) and (e). After the eight prior thermal transformation cycles, the crystal structure is expected to be stabilised and the c M e martensitic transformation is expected to cease, as shown in Fig. 4. Annealing at 1073 K is expected to revert the stabilised structure to the fcc c-phase. In contrast to the mixed structure at room temperature after the initial anneal, the second anneal at the same temperature, 1073 K, after the eight thermal transformation cycles appeared to have stabilised the
Y. Liu et al. / Acta Materialia 53 (2005) 3625–3634
c-phase to the ambient temperature (spectrum (d)). Furthermore, the annealed c-phase was also able to survive the cooling in liquid nitrogen without transforming to the low temperature e-phase (spectrum (e)). This is apparently related to the prior thermal cycling of the specimen. This phenomenon is in sharp contrast to the case of single crystal specimen, where recovery anneals at 1073 K is effective in restoring the c M e martensitic transformation repeatedly. Fig. 6 shows XRD measurements of the effect of transformation cycling on the structure of the Co– 30Ni polycrystalline specimen. The spectra were all measured after cooling in liquid nitrogen, except the first measurement, which was in the as-annealed state, and the third measurement, which was recorded after heating to 573 K for 3.6 ks following the second measurement. Between each cooling in liquid nitrogen, the specimen was heated to 573 K for 3.6 ks. All the XRD spectra are indexed to a mixture of the fcc c-phase and the hcp e-phase, as labelled in the figure. The integrated intensities of the (1 1 1)fcc peak and the ð1 0 1 0Þhcp peak are shown in Fig. 7. The intensities have been normalised to the total combined integrated intensities of the two peaks. The two values on the dashed line at ‘‘1.5’’ cycle indicate the intensities of the sample
(111)fcc (0002)hcp
Co-30Ni, Polycrystalline annealed at 1073 K (1011)hcp (200)fcc
X-ray Diffraction Intensity (arbitrary unit)
8th LN2 7th LN2
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after the heating to 573 K. It is seen that the first cooling in liquid nitrogen caused a marked increase in the amount of the hcp e-phase at the expense of the fcc c-phase. Heating to 573 K after the first cooling reverted the forward c ! e transformation. The e-phase formed by the second cooling was less than that formed by the first cooling. Correspondingly, more residual c-phase was detected after the second cooling. The amounts of both phases appeared to be largely unchanged after each cooling in subsequent cycles, stabilising at values between the maximum and the minimum of the asannealed state and after the first cooling. These observations are consistent with the DSC measurement that the first thermal cycle was associated with a larger endothermic effect and the thermal effect diminished rapidly with cycling. The fact that the final amounts of the two phases after eight cycles appeared between their respective maximum and minimum of the as-annealed state and after the first cooling indicates that both the forward and the reverse transformations were incomplete during thermal cycling, leading to the gradual diminishing of the transformation. As evident in the results presented above, the c M e transformation is highly irreversible. To investigate the possibility of recovery of the transformation in the polycrystalline material, the stabilised samples were heated to elevated temperatures and their structure was determined by means of XRD. Fig. 8 shows the measurements for a Co–30Ni polycrystalline specimen in the as-received condition. All the measurements were conducted at room temperature after cooling in liquid nitrogen following each of the overheating treatments, except the first measurement, which was conducted in the asreceived condition. Overheating treatment was conducted for 3.6 ks in vacuum. The temperatures of the
6th LN2 5th LN2
3rd LN2 2nd LN2 573 K 1st LN2
X-ray Relative Integrated Intensity
4th LN2
0.7
0.6
(1011)hcp
0.5
(111)fcc 0.4
as-annealed 42
44
46
48
50
52
54
2θ (o) Fig. 6. Effect of thermal cycling on structure of polycrystalline Co– 30Ni alloy.
0.3 0
1
2
3 4 5 Transformation Cycle
6
7
8
Fig. 7. Effect of thermal cycling on structural evolution of polycrystalline Co–30Ni alloy.
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Co-30Ni, Polycrystalline, As-received (200)fcc
X-ray Diffraction Intensity (arbitrary unit)
(111)fcc (0002)hcp
1073 K 973 K 873 K 773 K 723 K (1011)hcp
648 K 623 K 573 K 1st cooling as-received
40
45
50
55
2θ (o) Fig. 8. Effect of overheating on the transformation behaviour of polycrystalline Co–30Ni alloy.
treatment are indicated in the figure. It is seen that the as-received specimen contained a mixture of and e phases. Cooling in liquid nitrogen resulted in an increase in the fraction of the e-phase. Overheating to 573 K, 70 K above the finishing temperature of the reverse transformation (Fig. 4), led to the formation of less e-phase upon cooling to liquid nitrogen. This is obviously due to the effect of transformation cycling (Fig. 6). One clear difference was observed upon heating to 623 K, 120 K above the finishing temperature of the reverse transformation. However, upon overheating to 648 K, the structure was apparently totally reverted to fcc and no hcp e-phase was formed at the liquid nitrogen temperature. It needs to be emphasised that these XRD spectra were recorded after cooling in liquid nitrogen following each overheating, instead of at the room temperature after the heating. This means that the overheating at 648 K not only caused a full reversion of the e-phase to the c-phase, but also stabilised the c-phase to the liquid nitrogen temperature. It is evident that even heating to 1073 K, the original annealing temperature, was ineffective in restoring the c M e transformation. 3.3. Microscopic evidences Fig. 9 shows a low-magnification optical image of the surface of a Co–32Ni single crystal specimen. The
Fig. 9. Surface relief of the c M e martensitic transformation in Co– 32Ni single crystal.
viewing surface was the (1 1 1)fcc plane of the single crystal. The specimen was pre-polished at room temperature in the as-annealed state and then cooled in liquid nitrogen. The micrographs were taken at room temperature after the cooling. The micrographs reveal clear surface reliefs. The surface reliefs appear to have three distinctive morphologies. Some parts exhibited long straight reliefs across the entire specimen, as indicated by the arrows in regions A. In other parts the surface reliefs exhibited equilateral triangular morphology of various sizes, as indicated by the triangle in region B. In region C, as indicated by the circle, the surface reliefs were small in dimension and less regular in shape. However, all these reliefs appeared to conform to three basic variants following the three sides of the triangle, which are indexed to Æ1 1 0æ directions of the fcc structure. This is clearly a manifestation of the expansion of stacking faults along h1 1 2if1 1 1gfcc systems, which is the mechanism of the fcc ! hcp martensitic
Y. Liu et al. / Acta Materialia 53 (2005) 3625–3634
transformation. It is also evident that there existed large areas, as large regions or between surface relief lines, which remained unaffected by the transformation, suggesting that the fcc ! hcp transformation was localised and incomplete. Fig. 10 shows a series of optical micrographs of region D, which is marked by the rectangle, in Fig. 9. All micrographs were taken at room temperature and are presented in their chronicle order. Micrograph (a) was taken after the pre-polish in the as-annealed state. Micrograph (b) was taken in the e-martensite state after the first cooling in liquid nitrogen (the same condition as the micrographs shown in Fig. 9). Micrograph (c) was taken after a heating to 533 K, some 50 K above the completion of the reverse transformation, following micrograph (b). It is seen that, whereas a reverse transformation had occurred upon heating to 533 K, the surface relief did not disappear; instead, it became more extensive. Micrograph (d) was taken after a re-polish in the c-austenite state following micrograph (c). Micrograph (e) was taken after a second cooling in liquid nitrogen after the re-polish. It is seen that the same martensite variants to those appeared in micrograph (b) were formed again, even the small irregularity at the lower left hand corner. At the same time, the extent of the surface relief was much reduced compared to that after the first cooling (micrograph (b)). Micrograph (f) was taken after a third polish in the e-martensite state. Micrograph (g) was taken after a heating to 533 K following micrograph (f). It is seen that exactly the same variants were formed by this reverse transformation compared to those formed on the forward transformation (Figs. (b) and (e)), but to a much less extent. Micrograph (h) shows the surface after a third polish in the c-austenite state and micrograph (i) shows the surface relief after cooling in liquid nitrogen flowing the polish. The number of surface relief lines was drastically reduced. This is obviously due to the effect of transformation cycling. Fig. 11 shows transmission electron micrographs of a Co–32Ni single crystal specimen, with (a) showing the specimen in the as-annealed state and (b) showing the microstructure after chilling in liquid nitrogen. Both micrographs were taken at ambient temperature. It is seen that the as-annealed state contained a large number of stacking faults. This is consistent with previous observations reported in the literature [8,9] as well as the XRD measurement of the co-existence of fcc and hcp phases in the as-annealed state. It is also evident in the image that the stacking faults were extensively interlinked between different fault systems. The microstructure of the as-chilled specimen showed large areas of untransformed fcc phase. This is consistent with the XRD observation of the incomplete transformation in either direction.
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Fig. 10. Effect of pre-polish and transformation cycling on the surface relief of the c M e martensitic transformation in Co–32Ni single crystal.
4. Discussions 4.1. Transformation mechanism The fcc M hcp martensitic transformation in Co–Ni alloy system occurs by the expansion and shrinkage of stacking faults via the movement of a6 h1 1 2i Shockley partial dislocations along the {1 1 1}fcc systems [3,8]. Activation of one such partial dislocation only creates or expands a stacking fault. For a hcp martensite of finite thickness to form a coordinated activation of
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Fig. 11. Transmission electron micrographs of the microstructure of Co–32Ni single crystal at room temperature: (a) in as-annealed state, (b) after chilling in liquid nitrogen.
Æ1 1 2æ gliding on every second (1 1 1) plane is required [8,10,11]. In the fcc structure there are 12 such systems, with three equivalent a6 h1 1 2i directions on each of the four {1 1 1} planes. It is important to note that the three a h1 1 2i partial dislocation systems on each (1 1 1) plane 6 produce exactly the same variant of the hcp martensite, despite the difference in siding direction. Therefore, there are in total four martensite variants possible from a fcc single crystal. Whereas a coordinated sliding on every second (1 1 1) plane is required, there is no restriction to the selection of sliding among the three h1 1 2i directions on each plane. There exist three possible combinations of sliding directions: (1) all mobile Shockley partial dislocations that contribute to the transformation have the same Burgers vector and slide in the same ½1 1 2 direction (bi = b1), (2) all three directions are equally selected, P P P giving the statistical condition that b1 + b2 + b3 = 0, (3) a random P or P P selection among the three so that b1 + b2 + b3 6¼ 0. Condition (1) has been referred to in the literature as mode A, condition (2) as mode B and condition (3) as mode C [3]. It is obvious that for mode B the transformation produces no crystal shape distortion beyond the scale of a few atomic planes. For mode A, it is easy to compute that, for a standard hcp structure of c/a = 1.63, the transformation will produce a maximum shear strain of 35% in the chosen ½1 1 2 direction on the chosen (1 1 1) plane. For condition (3), the crystal will experience some long-range shape distortion. It is important to note that all the three mechanisms produce exactly the same variant of the hcp martensite, the only difference in the product being the external shape of the crystal. The actual mechanisms of the fcc M hcp martensitic transformation in Co–Ni alloys have been studied previously by means of TEM [3,8,9,12], optical microscopy [2,3], atomic force microscopy [3], neutron scattering
spectrometry [10] and thermomechanical testing [6,7]. Waitz and Karnthaler [3] suggested that in short-range order the transformation occurs by mode A mechanism, producing a net shape change, whereas in long-range order internal stresses created by a mode A transformation in one location induce other mode A transformations in neighbouring regions to facilitate a self-accommodation of the martensite, yielding a nil net shape change. Hayzelden et al. [8] concluded from their TEM study that the hcp martensite does not form in self-accommodation mode and that the transformation, at the local environment of individual partial dislocations, is highly reversible. There is no specific determination of the Burgers vectors of the partial dislocations contributing to the transformation in this study. However, the optical microscopic evidences presented in Figs. 9 and 10 appear to suggest that the transformation occurs by the mechanism of mode C – there is no specific coordination of the Shockley partial dislocation movement. The sliding is highly random, both within one (1 1 1) variant system and among all Æ1 1 1æ variant systems; and both in the short spatial range and over large distances. It is also important to note that whereas the lattice distortion of the transformation can be perfectly self-accommodated within one Æ1 1 1æ variant, there exists no accommodation mechanism among the four different Æ1 1 1æ variants. 4.2. Latent heat of the transformation In the particular case of Co–Ni, the lattice distortion of the transformation is pure shear and involves no change in the lattice parameters. This renders the transformation of minimal lattice reshuffling. For this type of first order transformations the latent heat is expected to be small. In this work the heat effect measured by DSC is in the order of 1 J/g at its maximum during the initial transformation cycles. Even considering the fact that
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some hcp phase existed prior to the transformation and that the transformation associated with the thermal event was incomplete, the total heat of the transformation is not expected to be greater than several joules per gram. Hayzelden et al. [8] reported latent heats of 6.5–8.5 J/g for the martensitic transformation in a small-grained polycrystalline Co–10Ni alloy. The latent heat of the transformation in Co–32Ni alloy is expected to be lower. 4.3. Thermal reversibility It is interesting to note that despite the pre-existence of the stacking faults at ambient temperature, as evidenced by the XRD measurements and TEM observation, the forward transformation occurred in an abrupt manner during the first cooling. This implies that the pre-existing Shockley partial dislocations are not mobile and in a low free energy state. The abrupt occurrence of the fcc ! hcp transformation upon cooling is a manifestation of a sudden activation of the partial dislocations that had been pinned, in analogue to the case of activation of dislocations during the Lu¨ders deformation of mild steels, where a higher critical driving force is required to trigger the process than to propagate the process. This is consistent with the expectation that after full anneal at 1073 K the partial dislocations (and the expansion between each pair) stay the least energised state, or the most thermodynamically stable state. Another interesting aspect is that the forward c ! e transformation was observed only during the first cooling from the as-annealed state. In subsequent thermal cycles similar exothermic transformation peak was not detected by DSC down to 173 K. Evidenced by the detection of the reverse transformation on heating, apparently the forward transformation occurred over a wide temperature range at extremely low intensity, which is beyond the detection sensitivity of the DSC. The occurrence of the broad forward transformation is believed to be due to the liberation of the partial dislocations from the initial pinning effect of impurities and structural defects that renders the partial dislocations a wide spread in terms of mobility, or critical driving force for movement. The DSC measurements indicate clearly that the fcc M hcp martensitic transformation is highly irreversible. The structure stabilises and the transformation diminishes rapidly during thermal transformation cycling. This is confirmed also by the optical microscopic observation shown in Fig. 10. The gradual reduction in the transformation volume is an indication of the reduction of the population of mobile partial dislocations. This reduction is attributed to an arrest effect of the partial dislocations by structural defects during transformations. During transformations the mobile partial dislocations move forth and back in the matrix,
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and are arrested when fall into low-energy traps, such as structural defects or free surfaces. The XRD measurements confirm that the transformation stabilises during thermal transformation cycling towards a mixed state of fcc c-phase and hcp e-phase. This is consistent with the hypothesis of arrest of partial dislocations. This observation is in contradiction to the claim by Hayzelden et al. [8] that the transformation is microscopically reversible without significant hysteresis. Another interesting observation is the clear difference between the single crystal and the polycrystalline materials in terms of the recoverability of the transformation after stabilisation by thermal cycling. The single crystal specimen was revived repeatedly by recovery anneal at 1073 K whereas the polycrystalline materials was permanently stabilised after cycling. Given the fact that the polycrystalline material had a lower Ni content and thus its fcc phase is less stable than for the single crystal specimen, this observation is of higher certainty. This is probably related to the pinning effect of grain boundaries, which arrest the mobile partial dislocations beyond recovery at 1073 K. 4.4. Mechanical reversibility The optical microscopy observations show that transformation-induced surface relief prevails in subsequent reverse transformations. This is consistent with previous observations [2]. This demonstrates that, in addition to the thermodynamic irreversibility, the fcc M hcp martensitic transformation is also mechanically irreversible. This is schematically expressed in Fig. 12. The figure shows a (1 1 1)fcc plane of the fcc austenite, which coincides with the (0 0 0 1)hcp of the hcp martensite. The
d
b
a
c
[1 10]
[121]
(111)fcc – (0002)hcp Fig. 12. Schematic illustration of mechanisms of fcc M hcp martensitic transformation, demonstrating the multiplicity of transformation path for the forward and the reverse processes.
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edges of the triangles are the Æ1 1 0æfcc directions and the arrows indicate the Æ1 1 2æfcc directions. The open circles represent atomic positions in the fcc phase whereas the closed circle marks the position for the hcp martensite. Upon cooling, the fcc atoms have three equivalent possibilities to shift into hcp martensitic position by lattice shear along the Æ1 1 2æfcc directions. Upon the reverse transformation, the hcp atoms also have the same three possibilities to return to the fcc austenite by the same lattice shear, i.e. the constraint for the same return path is absent. This is obviously responsible for the prevalence of the surface relief upon transformation reversion. It ought to be pointed out despite the multiple choices of the return path for each martensite variant, they all return to the same original fcc crystal. This implies that the fcc single crystal remains a single crystal after thermal transformation cycles. It is also obvious that this non-uniqueness of the return path for the transformation disqualifies the alloy as a thermomechanical shape memory alloy.
5. Conclusions The experimental evidence generated by the different techniques of DSC, XRD, optical microscopy and TEM is highly consistent. Based on this evidence the following conclusions may be drawn regarding the fcc M hcp martensitic transformation in Co–Ni: (1) The alloys are in a mixed fcc–hcp state at ambient temperature after annealing, the microstructure contains a high population of stacking faults, and the stacking faults are extensively interlinked between different Shockley partial dislocation systems. (2) The fcc M hcp transformation in Co–Ni alloys of 30 wt.% Ni is highly irreversible. It is characterised by a large thermal hysteresis. It stabilises quickly through thermal transformation cycling. Both the forward fcc ! hcp transformation and the reverse transformation are highly irreversible and the transformation stabilises with a mixed structure of fcc and hcp. Annealing at 1073 K is effective in annihilating the stabilisation effect for Co–32Ni single crystal alloy, but ineffective for the Co–30Ni polycrystalline alloy. The latent heat of the transformation is very small.
(3) The transformation is crystallographically reversible, in that it always returns to the same fcc crystal upon the reverse transformation, despite the retention of some hcp phase at temperatures above the completion of the reverse transformation. In other words, no multiple variants may be generated through thermal transformation cycling. (4) The transformation is mechanically irreversible. Whereas the original fcc crystal structure and orientation is always re-achieved after a thermal transformation cycle, the shape of the fcc crystal can be changed. This is particularly the case when the influence of a mechanical bias is present during the process of transformation. This disqualifies the fcc–hcp martensitic transformation in Co–Ni from serving as a mechanism for the thermomehanical shape memory effect. It can obviously be postulated that all fcc–hcp martensitic transformations are invalid for the thermomechanical shape memory effect.
Acknowledgement This work is financially supported by ARC Discovery Project Grant DP0345880.
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