Microstructural influence of Mn additions on thermoelastic and pseudoelastic properties of CuAlNi alloys

Acta metall, mater. Vol. 42, No. 5, pp. 1583-1594, 1994

~

Pergamon

0956-7151(93)E0050-D

Copyright © 1994ElsevierScienceLtd Printed in Great Britain.All fights reserved 0956-7151/94 $6.00+ 0.00

MICROSTRUCTURAL INFLUENCE OF Mn ADDITIONS ON THERMOELASTIC A N D PSEUDOELASTIC PROPERTIES OF Cu-A1-Ni ALLOYS

M. A. MORRISand T. LIPE Institute of Structural Metallurgy, University of Neuchfitel, Av. Bellevaux 51, 2000 Neuch~tel, Switzerland (Received 29 April 1993; in revised form 18 October 1993)

Abstract--The thermoelastic and mechanical properties of Cu-AI-Ni-B--Mn shape memory alloys have been studied as a function of manganese concentration and of heat treatment. Below a limiting value of manganese content, the loss of thermoelastic and pseudoelastic properties has been observed, in particular in the quenched specimens. The partial transformation and its degradation during thermal cycling observed in the low manganese content alloy has been attributed to the lower degree of B2 order achieved during the quench, leading to slower kinetics of DO 3 ordering. The accommodation of strains between martensite variants and between the martensite/austenite phases appear to need dislocation accumulation at their interfaces. The presence of dislocations observed during the reverse transformation seem responsible for the degradation of the transformation and the loss of pseudoelastic properties of this alloy.

INTRODUCTION Copper based shape memory alloys have received a renewed interest in recent years due to the possibility that they offer to achieve transformation temperatures above 100°C, making them attractive for high temperature applications. In particular, Cu-A1-Ni alloys, over a varied range of compositions, including additions of other elements, have been shown to be potential candidates as best alloy systems [1-3]. On the other hand, compositional changes can modify not only the transformation temperatures but also the deterioration of the ductility [4] and may even result in the loss of thermoelastic properties during thermal ageing [5]. If polycrystalline alloys with sufficiently high transformation temperatures (As ~ 150°C) are to be obtained, an aluminium concentration lower than 14 wt% is required [6]. Recent research carried out in alloys containing 12 wt% A1 with manganese and boron additions [7-9] has shown that both ductility and transformation temperatures are very much improved with respect to traditional C u - A I - N i alloys [4, 10]. The influence of increasing boron concentration on increasing transformation temperatures [8] was attributed to the increasing number of boride particles present in the alloys reducing the concentration of manganese and/or aluminium in solution. Also the influence of these particles on grain size together with the better cohesion strength at grain boundaries were made responsible for the large ductilites (4--6%) that can be attained at room temperature. However the role that manganese additions play on the properties of these alloys has not yet been understood. For this reason, the present study

has been made on similar alloys but varying the manganese concentration for a constant boron content. Alloys with nominal compositions Cu-12AI-4Ni-0.04B (wt%) and different manganese concentrations between 2 and 4% have been studied. The role that the latter has on modifying the microstructures responsible for specific thermoelastic and mechanical properties will be discussed.

EXPERIMENTAL The alloys used for this study have nominal compositions Cu-12AI-4Ni-0.04B (wt%) with additions of different amounts of manganese, namely 2% (M14) 3% (MI5) and 4% (M16). The chemical analysis obtained from these alloys is given in Table 1. They were prepared by induction melting under a helium atmosphere. They were subsequently heat treated at 820°C for 30 min prior to extrusion at this temperature. Different t-homogenization heat treatments were carried out on the extruded bars between 820 and 900°C for 30 rain followed by water quench. Also anneals between 30 rain and 6 h were given to the initially quenched specimens. The alloys were studied after the different quenched and annealed treatments by optical and transmission electron microscopy (TEM), X-ray diffraction and differential scanning calorimetry (DSC). The latter was carried out by thermal cycling between room temperature and 300°C in order to obtain Table 1. Chemicalanalysisfrom the three alloys studied Cu A1 Ni B Mn M14 82.57 11.86 3.92 0.033 1.62 MI5 81.51 11.83 3.97 0.033 2.66 MI6 80.56 11.83 3.95 0.033 3.63

1583

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MORRIS and LIPE: THERMOELASTIC AND PSEUDOELASTIC PROPERTIES OF Cu-AI-Ni

information on transformation temperatures and their stability together with the values of the transformation enthalpies. The latter provided understanding on the total fraction transformed and its evolution during thermal cycling. Up to eight heating/cooling cycles (denoted as H I - H 8 and C1-C8 respectively) were monitored in each case using a heating rate of 10°C/rain. Also, a first complete heating/cooling cycle was carried out to 750°C in each heat treated alloy to obtain information on other transformations due to different ordering stages or phase changes. In order to assess the difference of the time elapsed between the annealing treatments and the calorimetry experiments, comparisons were made between the transformation temperatures obtained from the first heating and cooling cycles of the specimens annealed in the furnace at 300°C with those obtained after an isothermal anneal within the calorimeter. No difference was detected between any of these values. X-ray diffraction analysis was performed with a Philips MPD 1880 diffractometer using monochromatic copper K~ radiation on specimens electropolished using a solution of 40% orthophosphoric acid in water. This same solution was used for polishing specimens for optical observation and also for electropolished surfaces of the tensile specimens that were used to evaluate the mechanical properties. Tensile tests were carried out from specimens with a gauge length of 20 and 3 mm in diameter machined from heat treated bars. The tests were performed at different temperatures between 20 and 300°C at a strain rate of 5 x 10 -5 s -~ using a Shenck Universal testing machine equipped for high temperature experiments. Some tests were performed to fracture to evaluate the total ductility of the materials while others were performed by stress-cycling (loading and unloading at the same strain rate and constant temperature). In situ loading and heating experiments were performed under an optical microscope using a specially constructed tensile machine equipped with load and elongation transducers as well as heating facilities. The tensile specimens used for these experiments were spark-machined and had a gauge length of 60 mm with a section 3 mm wide and 1 mm thick. The same electropolished surfaces were obtained before testing. TEM observations were performed from this foils using a CM12 Philips electron microscope equipped with a heating holder to carry out in situ heating experiments. The thin foils used for these observations were prepared by the jet electropolishing technique with a solution of 30% nitric acid in methanol at - 30°C.

effect of the different heat treatments on the martensitic structures obtained and on the grain size stability responsible for the mechanical and transformation properties of the materials. The average grain sizes measured were 100 + 20 #m in all cases and they have been confirmed to be independent of heat treatments for all the alloys. These small values have been attributed to the presence of small boride particles [9] preventing grain growth during the annealing treatments. Figure 1 shows examples of these fine structures observed by TEM where the presence of randomly oriented martensite variants produced during cooling appear to accommodate any elastic strain near grain boundaries and triple points. In Fig. l(b) we see a detailed tweed-like microstructure within the martensite laths typical of the manganeserich alloys after annealing at 300°C. The X-ray data obtained from the different alloys as a function of heat treatment have confirmed the presence of a martensite with the orthorhombic 18R structure, irrespective of the heat treatment given to the alloys. The lattice parameters measured indicate that the manganese atoms produce a slight dilatation of the lattice, in particular along the b-axis, as shown in Table 2. We also note that the a/b ratio is less that

RESULTS 1. Microstructures: effect o f heat treatments Optical and TEM observations made from the three alloys studied have provided evidence on the

Fig. 1. Typical martensitic structure observed by TEM in the quenched specimens: (a) MI4 alloy, (b) tweed-like structure observed within the martensite laths after annealing the alloy M15 at 300°C for 30min.

M O R R I S and LIPE:

T H E R M O E L A S T I C A N D PSEUDOELASTIC PROPERTIES OF Cu-A1-Ni

Table 2. Lattice parameters measured from the 18R structure analysed by X-ray diffraction from all the alloys Lattice parameters

M 14

M 15

M 16

a (nm) b (nm) c (nm)

0.443 0.525 3.807

0.444 0.526 3.812

0.443 0.529 3.811

~/3/2 due to the fact that the structure is ordered and the basal plane contains atoms of different sizes. A measure of the degree of order that the martensitic structure has inherited from the parent phase (austenite) has been obtained by measuring the difference in d-spacing (Ad) between pairs of planes (splitting) which would have identical peaks when the alloys have disordered structures but whose peaks would be split in the ordered state [11]. For those planes the spacing difference Ad would provide a qualitative indication of the degree of long range order: the higher the degree of order the larger Ad values will be obtained [11, 12]. Table 3 shows the measured Ad values corresponding to the pairs of peaks 12]-202, 12~-208 and 2010-1210 respectively for the three alloys studied after quenching from 820 and 900°C respectively. Here we note that although the Ad values increase with increasing manganese content, there is no appreciable difference in Ad by quenching from either 820 or 900°C. Therefore this implies an increase in degree of order with increase in manganese concentration which is independent on the fl-homogenization temperature used in this study. As will be discussed in the next section the stability of the transformation degrades during thermal cycling of the quenched specimens and for this reason the latter were subsequently annealed at 200 and 300°C between 30 min and 6 h. The measured Ad values do not show much variation during annealing at 200°C. However, during the 300°C heat treatment a different behaviour was observed between the alloy MI4 with low manganese content and the other two alloys richer in manganese. While the latter showed an increase in the degree of long range order during the first 60 min at 300°C (as evidenced by the increase in the Ad values seen in Fig. 2 for the two pairs of planes 12~-202 and 1210-20]-0), the low manganese content alloy (M14) shows a decrease in the values of Ad after any period of annealing time at that temperature. Also from Fig. 2 we note that after 60 min annealing, a decrease in these values was observed for the alloys containing 3 and 4% Mn. Besides, the alloy with higher manganese content shows the largest Ad values measured after annealing at 300°C while the

1585

lowest values are those for the alloy with lower Mn concentration. TEM high resolution images were taken from the martensitic structure of the quenched and annealed specimens such that the lattice fringes corresponding to the periodicity of the stacking order could be obtained when the beam direction was parallel to the close packed planes [13]. Figure 3 shows an example where the lattice fringes corresponding to two twinned martensitic laths are seen when imaged along the [010] direction together with the corresponding superimposed diffraction patterns from both laths. The fringe spacing measured was between 0.62-0.64 nm in all cases confirming the presence of the 18R martensite with a c-lattice parameter corresponding to that measured from the X-ray diffraction data and shown in Table 2. Other than the degree of order inherited by the martensitic structure, the type of order has also been analysed from TEM observations. Figure 4 shows examples of the ordered domains obtained from the lower manganese content alloy (M14). In the quenched state from 820 and 870°C [Fig. 4(a)] the

7

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$ Table 3, Measured Ad values from split peaks of the ordered martensitic structures obtained by quenching from the fl-homogeneization temperatures 820°C Q

900°C Q

Ad(10 3nm)

M14

M15

MI6

Ml4

Ml5

MI6

Adi2L2o2

4.8 2.2 6

5.4 2.3 6.9

6.5 3.3 7.9

4.4 2.1 5.7

5.3 2.5 6.9

6.4 3.4 7.9

Ad,21,208 Adi210,20T0

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I

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t(h) Fig. 2. Variation of the Ad values characteristic of the amount of splitting between pairs of X-ray peaks as a function of annealing time at 300°C for all the alloys (the initial value at time zero is that of the quenched specimens, Q).

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MORRIS and LIPE:

THERMOELASTIC AND PSEUDOELASTIC PROPERTIES OF Cu-AI-Ni

only ordered domains present are those corresponding to the B2 type structure as evidenced from the dark field images obtained using the 010 superlattice spot. We see in the same Figure (4b) that after annealing at 300°C for 30 min the presence of both DO3 and B2 ordered domains, obtained in dark field with the superlattice spot 111, has been detected in this alloy. For the other two alloys containing 3 and 4% Mn respectively, both B2 and DO~ ordered domains were already observed in the quenched specimens. Therefore, there appears to be a limiting manganese content which allows the transformation between B2 and DO3 order to occur during the quench (i.e. the kinetics leading to D O 3 order accelerate with increase in manganese concentration between 2 and 4%.

-

--

2. Transformation properties: thermoelasticity The reverse and martensitic transformations of our alloys have been evaluated from the DSC curves obtained during thermal cycling between 20 and 300°C by measuring the transformation-temperatures As, Af, Ms, Mfas well as the total enthalpies obtained from the transformations. Since the initial condition

Fig. 4. Examples of ordered domains obtained from the M14 alloys: (a) only B2 superlattice spots are seen in the quenched state and the corresponding B2 domains are observed in dark field using the superlattice spot 010 (marked by arrow). (b) After annealing at 300°C also DO 3 supedattice spots are seen and the corresponding DO3 domains are obtained using the 111 diffraction spots.

in all cases was martensitic, a thermal cycle started by a reversed transformation and was followed by the martensitic one during the cooling part of the same cycle. Figure 5 shows one example comparing the DSC curves obtained from the first and second

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Fig. 5. Typical DSC curves obtained from the first (HI,CI) and second (H2,C2) heating -)- cooling thermal cycles of the quenched MI5 alloy. Note that the heating part of the first cycle, direct from the quenched specimen, produces only a partial reverse transformation with lower enthalpy.

MORRIS and LIPE: THERMOELASTIC AND PSEUDOELASTIC PROPERTIES OF Cu-A1-Ni 250

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heating + cooling cycles (noted H1,C1 and H2,C2 respectively) in the case of the quenched M 15 alloy. One important characteristic of the quenched specimens is the different behaviour between the first and subsequent heating + cooling cycles. Figure 6 shows plots of the transformation temperatures as a function of the number of cycles for the alloys M14 and M15 respectively. We note that between the first and second thermal cycles the intervals M s - M r and As-A f decrease substantially due to the different shape of the curves, in particular in the case of the low manganese containing alloy. From the surface areas measured inside the curves, the total enthalpies of the transformations have been obtained and these have been shown to be rather low in these quenched specimens indicating that the transformations are only partial during the first thermal cycle. These enthalpies have values ranging between 2.6 and 3.9 J/g in the M14 alloy quenched from temperatures between 900 and 820°C respectively while for the other two alloys the transformation enthalpies were about 6 J/g irrespective of the quenching temperature. In order to understand the stability of the transformations as well as to make them be more

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complete, the different annealed specimens were studied during thermal cycling. Although the transformation of the alloy M14 (2% Mn) appeared more complete after annealing, the stability of the transformations was again rather poor and during cycling the transformation enthalpies decreased rapidly from

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MORRISand LIPE: THERMOELASTICAND PSEUDOELASTICPROPERTIES OF Cu-A1-Ni

Table 4. Transformation temperatures and enthalpy values (AH) measured after thermal cycling for all the alloys quenched from 820°C followed by annealing for 30 rain at 300°C M14

MI5

MI6

NO

cycles 1 2 4 6 8

Af

Ms

As

Mf

A H (J/g)

Af

Ms

As

Mf

AH (J/g)

Af

Ms

As

Mf

AH (J/g)

227 222 204 178 166

183 172 147 136 133

111 110 102 98 93

96 95 91 89 88

9.1 5.6 5.3 4.8 4.8

164 166 167 167 167

145 142 142 142 142

115 117 122 124 125

93 93 92 92 92

9.0 9.2 9.4 9.5 9.6

96 99 100 104 106

60 62 66 68 72

66 62 64 66 67

41 42 47 47 47

8.9 9.0 9.1 9.1 9.2

7

9 J/g obtained in the first cycle to about 5 J/g after five heating cycles respectively. Figure 7 shows this variation for different quenching and annealing conditions where we note that during the martensitic transformation (cooling) the decrease in enthalpy is more important in specimens quenched from 900°C followed by annealing at 300°C for 30 min. In a similar way the transformation temperatures of this alloy varied rapidly during cycling, in particular the values of Af and Ms decreased by as much as 50°C during the first eight cycles. Table 4 shows the transformation temperatures together with the measured enthalpies (AH) for all the alloys after quenching from 820°C followed by annealing for 30 min at 300°C. From this table we note the better stability of the transformation temperatures measured from the alloy M15 and M16 (containing 3% and 4% Mn respectively). In addition these alloys show high values of the reverse and martensitic transformation enthalpies, measured as 9J/g and showing little variation during thermal cycling (note that since similar enthalpies were measured during heating and during cooling only one value is given) and we conclude that these heat treatments produce the best conditions for the thermoelastic transformation of these alloys. Figure 8 shows the DSC curves obtained during the first (HI,C1) and eighth (H8,C8) thermal cycles from which the latter values were measured for the M 15 alloy. The different anneals at 200°C after quenching and at 300°C for periods of time other than 30 rain have shown either lower values of transformation temperatures or decrease of the transformation enthalpies during cycling. As an example Table 5 shows the transformation temperatures measured after the second thermal cycle for the M15 alloy (3% Mn) after annealing for 30 min at 200 and 300°C respectively. Here we note the slightly lower values in the specimens annealed at the lower temperature. Finally we remark that the alloy with higher manganese concentration (M16) has lower

transformation temperatures irrespective of the quenched or annealed state of the specimen.

3. Mechanical properties: ductility and pseudoelasticity The mechanical properties of these alloys have only been evaluated from the materials quenched from 820°C followed by annealing at 300°C for 30 min since under these conditions the alloys had the best transformation properties. The tensile curves shown in Fig. 9 for two of the alloys provide evidence of the ductility and strain hardening exhibited by these materials at all temperatures during the tests performed to failure. The values of flow stress at 0.2% strain, maximum tensile strength and ductility together with the values of Young's moduli measured, have been plotted in Fig. 10 as a function of testing temperatures for all the alloys. Here we note that although similar ductilites are achieved for all the alloys at room temperature, the lower manganese containing alloy has lower ductilities at higher testing temperatures.

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Af Ms A~ Mr

820°C Q +300°C, 0.5 h

820°C Q +200°C, 0.5 h

166 142 117 93

145 130 100 90

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E(%) Fig. 9. Examples of the tensilecurves obtained from the tests performed to failure from two of the a11oys.

MORRIS and LIPE: THERMOELASTIC AND PSEUDOELASTIC PROPERTIES OF Cu-A1-Ni

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1590

MORRIS and LIPE: THERMOELASTIC AND PSEUDOELASTIC PROPERTIES OF Cu-A1-Ni :

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Fig. 12. Stress induced crystallographic martensite variants observed during in situ stress-cycling at a high temperature (above Af) for the same grain seen in Fig. 13: (a) initial loading to a = 200 MPa, (b) increased deformation under the same stress as in (a). (c) Unloading after first loading cycle. (d) Unloading after third loading cycle (note permanent stress-induced variants marked by arrows). In situ tensile experiments have been carried out under the optical microscope to confirm the presence of new martensite variants after unloading the specimens deformed at room temperature. During high temperature deformation, i.e. in the austenitic state, the yield stress also defines the stress at which the formation of stress-induced martensite occurs. However, any strain produced by the latter should be recovered on unloading leading to the pseudoelastic effect. From the curves shown in Fig. 9, we note that there is an important hardening effect in our alloys which becomes more important at high temperatures, i.e. in the austenitic state, such that ductilities up to 10% can be achieved. In order to evaluate whether such large strains could be recovered on unloading, stress-cycling tests have been carried out at temperatures above Ar (Td ~ Af + 50°C) for the three alloys. The stress-strain curves obtained from these tests are shown in Fig. 11. Also in this figure we show for comparison one of the curves obtained after stresscycling at room temperature (below Mr). We note that the alloys M15 and M16 present a pseudoelastic effect of 3.5 and 4.5% strain respectively, i.e. irrespective of the total strain to which the specimens have been deformed the maximum recoverable strains are limited by these values. Also we note that there is a limiting stress which determines this maximum pseudoelastic strain for each alloy. Besides we see that the

low manganese containing alloy (M14) shows no pseudoelastic effect at all, the stress-strain curves at high temperature appear rather similar to those observed during unloading the different alloys at room temperature. During in situ stress cycling experiments under the optical microscope at temperatures above Ar we observed the formation of stress-induced crystallographic martensite variants that disappeared on unloading. However, after two or three cycles we note that some variants remain permanently within the austenite phase. This can be seen in Fig. 12 and we note in Fig. 12(b) that at 120°C (Td = A f + 50°C) and under a stress tr = 200 MPa (tr0.2 + 40 MPa), several systems of crystallographic martensite variants are observed. On unloading the first time only some residual contrast of the original martensite is seen in Fig. 12(c) while after the third unloading cycle we see (marked by the arrows) the presence of some permanent stress-induced martensite. Finally, Fig. 13 shows some examples of stressinduced martensite observed from the alloys M 14 and M 15 after deformation at 200°C to fracture. Since the total strain achieved was 5 and 8% strain for the M14 and M 15 alloys respectively, of which nothing was recovered on unloading in the M14 alloy and only 3.5% strain is pseudoelastic in the M15 material, the crystallographic stress-induced martensite observed

MORRIS and LIPE: THERMOELASTIC AND PSEUDOELASTIC PROPERTIES OF Cu-A1-Ni

+ii~!¸i~!~,i

Fig. 13. Examples of permanent stress-induced martensite variants observed from specimens deformed to failure at a temperature above Af: (a) M14 alloy, (b) MI5 alloy. in Fig. 13 represents the favourable variants that produce maximum and permanent elongation of both alloys, DISCUSSION

(a) Transformation properties and related microstructures The influence of heat treatments on microstructural changes as well as on the transformation properties of our alloys has made it possible to establish a dependence of manganese concentration on related microstructures and thermoelastic properties. In particular, the different behaviour of the low manganese containing alloy for all the heat treatments used indicate that there is a limiting manganese concentration characterising the thermoelastic behaviour of these alloys. Although in all cases the martensitic structure observed had the same crystallography, there were slight differences in terms of the amount of splitting, Ad, between the X-ray peaks characteristic of the degree of long range order that the martensite inherits from the parent phase. The low manganese containing alloy has smaller Ad values in all cases confirming a lower degree of distortion of the existing martensite. Also during annealing at 300°C, an increase in the Ad values was observed during the first hour for the alloys containing 3 and 4% manganese while only a

1591

decrease in these values was observed for the M14 alloy. At this same time, the quenched specimens of the latter alloy presented only B2 ordered domains while DO3 domains were observed by TEM (see Fig. 4) after annealing at 300°C for 30 min. However, both types of domains were observed in quenched and annealed specimens of the M15 and M16 materials. To have information on the B2 and DO 3 ordering temperatures from the different materials in the quenched and annealed state, these were submitted to a full thermal cycle up to 750°C. Figure 14 shows an example of the DSC curves obtained between 300 and 600°C where the relevant peaks were observed. All the alloys showed similar DO3 and B2 disordering/ordering peaks (with only slightly different temperature intervals but always close to those shown) during heating and cooling respectively. We note a sluggish disordering DO 3 transformation during heating whose peak overlaps with that of the B2 disordering reaction. The decrease in the Ad values indicates a decrease in distortion of the basal plane of the 18R structure and this can be interpreted as being due to a more disordered martensitic structure or to a higher degree of DO3 order [12, 14]. Indeed in the DO3 ordering case the symmetrical forces, due to atomic arrangements where one site is surrounded by six symmetrical sites, will keep the axial ratio b/a of the 18R structure equal to the ideally x/3/2 value of the closed packed structure [15]. However, in the case of B2 or L21 ordering arrangements, the asymmetrical forces between the lattice sites with different sizes will lead to the increase of the b/a ratio such that more distorted close packed layers than those corresponding to the DO 3 ordered configurations will be obtained [15]. Also it has been shown that the extent of the distortion in the close packed planes increases with increasing degree of B2 order [12, 14]. During

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MORRIS and LIPE: THERMOELASTIC AND PSEUDOELASTIC PROPERTIES OF Cu-AI-Ni

the quench of our alloys the B2 order can be reduced to a certain extent while the DO3 order can be completely suppressed [16]. During ageing at 300°C, ordering to the DO3 structure can occur (that is ordering of the next-nearest neighbours occurs), while at the same time the order parameter describing nearest-neighbour ordering can increase further. In the case of the M14 alloy with lower manganese content, it appears that the quench suppressed the D O 3 order since the diffraction spots corresponding to this superlattice were not present and DO3 ordered domains were not observed. This might have been due at the same time to a lower degree of B2 order and to a slower D O 3 ordering kinetics during further cooling. However D O 3 domains were present after annealing at 300°C. This is consistent with the decrease in the Ad values observed during annealing being produced, as a major effect, by the increase in DOa order of the alloy. However, the larger Ad values measured in the quenched manganese-rich alloys must be attributed to a higher degree of nearestneighbour ordering (as found in a B2-ordered material) since the quenched alloys were already DO3 ordered showing both nearest-neighbour and nextnearest°neighbours domain boundaries. Therefore, the lower Ad values measured from the M14 quenched alloy compared to those from the manganese-rich materials have been assumed to be due to a lower degree of B2 order. The increase in Ad values measured after annealing at 300°C up to I h in the M15 and MI6 alloys can only be consistent with an increase in the distortion produced by the re-arrangement of the aluminium and manganese atoms trapped during the first quench such that the more stable Cu2 Mn A1 ordered L2~ structure forms [16]. The presence of a tweed-like structure in these annealed alloys, much less evident in the MI4 alloy [see Fig. l(b)], confirms this assumption. Indeed, the Cu-A1-Mn system [16] presents a two-phase field due to a spinoidai decomposition from the DO3 into the DO3 + L21 phases at temperatures below 350°C. In our alloys this decomposition seems to be produced in the alloys where the higher manganese concentration accelerates the re-ordering process. Comparisons between the transformation properties of the quenched alloys show that the lower transformation temperatures As, Ms are obtained for the MI6 alloy will higher manganese content and where the larger Ad values measured indicate a higher degree of B2 order. This implies that the increase in degree of order facilitates the transformation since probably a lower misfit strain needs to be accommodated at the martensite/austenite interface. By contrast, the very high transformation temperatures observed in the MI4 alloy together with a very low enthalpy of transformation, that decreases fast during cycling, indicate that the lower degree of B2 order (leading to an incomplete DO 3 order) achieved in this alloy (low Ad value) hinders the transformation.

From the different DSC curves obtained between the first and second heating cycles (see Fig. 5) of the manganese-rich alloys, after quench, we note that a partial transformation is obtained after the first cycle with a large A~Af temperature interval. This implies that the first thermal martensite that formed by the quench is rather stabilized, probably due to the excess strain energy accumulated at the interfaces, and the reverse transformation requires higher temperatures to be achieved. Even at these higher temperatures, the low values of enthalpy measured (6 J/g) indicate that the massive transformation of all the material has not occurred. After the second cycle, the decrease in the As-A f interval together with the increase in total transformation enthalpy between 6 and 9.5 J/g indicate that the full thermoelastic behaviour of the alloys has been achieved. This could be attributed either to the exposure at 300°C during the first thermal cycle or to the effect that the transformation itself has on the microstructural changes of the alloys (i.e. an improved degree of order obtained by the annealing process after the heating part of the first cycle could help reduce the strain energy accumulated at the parent/martensite interface). A recent study made on Cu-AI-Ni alloys has shown that thermal cycling produces changes in the degree of DO3 order while the B2 order remains unaffected [17]. This has been related to a dislocation structure observed in the parent phase after thermal cycling. In our alloys, in situ heating experiments carried out in the TEM have shown that only the quenched MI4 alloy has dislocations in the austenitic phase during heating. This can be seen in Fig, 15(a) where dislocations are seen gliding as the reverse transformation is taking place. By contrast we show in Fig. 15(b) an example where typical dislocation-free structures are seen after transformation into the austenite phase of the M15 alloy. These dislocation-free structures are characteristic of the alloys that had been heated in the TEM in our previous studies [9] and therefore it does not seem possible that the constraining stresses of the sample holder during heating could cause plastic deformation of the foils. Only during the heating experiments of all the M14 foils the presence of dislocations was detected. Therefore it must be concluded that these defects are the result of stress relaxations occurring at the martensite/austenite interfaces. The better transformation properties of the annealed alloys can be understood in terms of the ideal degree of order in the structure of the parent phase such that a one to one correspondence between atoms exists and coherent interfaces between the martensite and austenite phases are created. If any loss of coherency exists such that dislocations are needed to accommodate any misfit strain this will lead to a loss of thermoelasticity and only partial transformations will occur.

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to the total strain produced in shape memory alloys as Etota I = Eelasti c + Epseudoelasti c + Epseudoplastl c -~- Eplasti c .

Fig. 15. Reverse transformations observed during in situ heating experiments in the TEM (a) M14 alloy (note the presence of dislocations gliding during the transformation (marked by the arrows) due to the release of internal stresses): (b) M15 alloy (dislocation free) austenitic phase.

(b) Mechanical properties: pseudoelasticity

The values of yield stress, ductility and those of the Young's modulus obtained from our alloys show that they decrease continuously with increasing testing temperature down to a value near As. However they increase for higher temperatures between A~-A r and above. Although these values are lower for the rich-manganese alloy (M16) at temperatures below As they are, however, higher at testing temperatures above As. In order to interpret these results it is convenient to describe the deformation modes according to the phases present at each temperature. Since our alloys can be classified as Class II [18] such that the transformation temperatures obey the relationship Ms > A~ and the interval M s - Mf necessary to complete the transformation is large, the microstructure for which the lower yield values are obtained corresponds to the presence of both martensite and austenite phases. Each one of these phases will deform by the formation of stress-induced martensite variants which will be created at different levels of the applied stress (since the yield stress values are different). In order to discuss the different effect caused by the deformation we introduce the different contributions

(1)

Each one of those terms contributes or not to the deformation process depending on the stress applied at a given temperature. At room temperature, below Mr, only martensitic structures exist and on applying a stress, the thermal martensite variants can sustain an elastic strain, Eelastic,defined by the Young's modulus of the alloy. At the stress defined by the elastic limit, a0.2, new re-oriented stress-induced martensite variants form that produce an "apparently" permanent strain, Epseudoplasti e. On heating, this strain is recovered as the transformation into the parent phase occurs and the one way shape memory effect is obtained. We note, then, that the lower Young's modulus of this alloy will cause the stress-induced martensite to form at a lower applied stress leading to a higher capacity to accumulate pseudoplastic strain. The total strain in the austenite phase at temperatures above Af has been shown to contain a contribution from three of the terms shown in equation (1) when the alloys are deformed to fracture. The elastic strain Ee~ c is defined by the Young's modulus of the austenite phase, the pseudoelastic strain defined by the maximum strain that can be recovered on unloading, and the truly plastic strain which is not recovered. The latter has been shown to be due to the stable stress-induced martensite variants that remain present on unloading as seen in Fig. 12. The alloy containing higher manganese concentration, M I6, has shown a maximum pseudoelastic effect, with a total 4.5% pseudoelastic strain that can be recovered compared to 3.5% in the M15 alloy. On the other hand the total plastic strain that the latter can sustain is higher at the equivalent temperature, Af + 50°C, indicating that less stable stress-induced martensite variants will remain on unloading the M 16 alloy. This has been confirmed by the few crystallographic variants observed in the microstructure of that alloy after deforming to fracture at 150°C. The low manganese containing alloy, M14, has shown no pseudoelastic effect, whatsoever, during deformation at temperatures above hf. The temperature Af measured by DSC represents the value at which the reverse transformation has finished. However, the low values of the transformation enthalpy measured in this alloy indicate that at this temperature not all the martensite has transformed. This has been confirmed by the observations made during in situ heating experiments under the optical microscope where some untransformed martensite is present at temperatures as high as 200°C above Af. After several cycles the fraction of transformed martensite decreases and so does the measured enthalpy. The partial transformation that occurs in this alloy leads to a partially austenitic structure at temperatures above A r. The Young's modulus of this alloy is lower

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at high temperature since only a partially austenitic structure is obtained and the contribution from the martensitic phase will still be present. During high temperature deformation of this material the stress will induce new re-oriented martensite variants in both phases. The compatibility between the different variants will introduce stress concentrations which will not have been accommodated elastically, leading to stabilization of the crystallographic variants (see Fig. 13) and to the non-recoverable strain which explains the lack of pseudoelasticity. The increase in manganese concentration in our alloys produces better ductility and higher yield stress and Young's modulus at high temperature (between As and At). This indicates that manganese additions improve the mechanical properties of the austenite phase as also it has been shown to produce a higher degree of order. This increases the pseudoelastic properties of the alloy as well as introducing good thermoelastic behaviour. The alloy with intermediate manganese content constitutes a best choice in terms of higher transformation temperatures being achieved together with a sufficiently good behaviour in terms of mechanical and thermoelastic properties. CONCLUSIONS The thermoelastic and mechanical properties of a series of C u - A I - N i shape memory alloys have been studied as a function of increase in manganese concentration. An increase in manganese concentration increases the thermoelastic and pseudoelastic behaviour as well as the ductility of the alloys while reducing the transformation temperatures. The alloy with lower manganese content (2%) exhibits a lack of thermoelasticity, in particular in the quenched condition. After annealing, a partial transformation is observed with low values of the transformation enthalpies that decrease rapidly during

thermal cycling. These properties have been interpreted by a decrease of the degree in B2 order and suppression of D O 3 order during quenching and the increase of both types of order during annealing with a major increase occurring for the DO3 order. The presence of dislocations observed during the reverse transformation has been made responsible for the degradation of the thermoelastic properties and the lack of pseudoelasticity in this alloy. REFERENCES

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