Advertisement

Journal of Thermal Analysis and Calorimetry

, Volume 118, Issue 1, pp 111–122 | Cite as

Influence of Ti additions on the martensitic phase transformation and mechanical properties of Cu–Al–Ni shape memory alloys

  • Safaa N. Saud
  • E. Hamzah
  • T. Abubakar
  • Mohd Zamri
  • Masaki Tanemura
Article

Abstract

The effect of Ti additions on the microstructure and mechanical properties of Cu–Al–Ni shape memory alloys (SMA) was studied by means of a differential scanning calorimeter, field emission scanning electron microscopy, transmission electron microscopy, X-ray diffraction (XRD), a tensile test, a hardness test, and a shape memory effect test. The experimental results show that the Ti additions have an effective influence on the phase transformation behavior through generating a new phase into the microstructure, which is known as X-phase and/or controlling the grain size. The results of the XRD confirmed that the X-phase is a combination of two compounds, AlNi2Ti and Ti3·3Al. Nevertheless, it was found that with 0.7 mass% of Ti, the best phase transformation temperatures and mechanical properties were obtained. These improvements were due to the highest existence of the X-phase into the alloy along with a noticeable decrement of grain size. The Ti additions to the Cu–Al–Ni SMA were found to increase the ductility from 1.65 to 3.2 %, corresponding with increasing the strain recovery by the shape memory effect from 50 to 100 %; in other words, a complete recovery occurred after Ti additions.

Keywords

Cu–Al–Ni SMA Martensitic transformation X-phase Shape memory effect 

Introduction

As a result of their inexpensive, wide-ranging transformation temperatures, being easy to produce, their high thermal stability, and their small hysteresis with an affordable shape memory effect, Cu-based shape memory alloys have attracted lots of attention from scientists and researchers. Among the Cu-based alloys, Cu–Al–Ni shape memory alloys (SMAs) are the most used alloys, and in particular whenever high temperatures are required; so they are assigned to their high transformation temperatures, which are able to work at 200 °C [1], which is usually difficult for Cu–Zn–Al and NiTi alloys [2, 3], whose maximum temperatures are around 100 °C [4]. The shape memory effect of these alloys is able to display at a specific composition of about 11–14 mass%, for aluminum and 3–5 mass% for nickel. Even so, those mainly responsible for these effects are the properties of the β-phase. Additionally, to control/slow down the diffusivity of Cu and Al, nickel was added. Meanwhile, this addition may result in suppressing the decomposition of the β-phase during the cooling process [5]. There are some ordered structures caused by the transformation of A2 type disorder in the Cu-based SMAs over the course of the quenching process from high temperatures. Moreover, this further cooling leads to a transformation of these types of ordered structures into various popular structures (6R, 18R, and 2H) [6, 7, 8], whereas the R refers to rhombohedral and H hexagonal structures. For Cu–Al–Ni SMAs, two martensite structures are typically produced, 2H ( \(\gamma^\prime_1\) ) and 18R ( \(\beta^\prime_1\) ) , of thermally induced martensite, as well as the formation of these structures being affected by Al and Ni compositions, along with the heat treatment. Throughout the cooling treatment, and also at a low availability of Al, the martensite formed into a structure of \(\beta^\prime_1\) , while \(\gamma^\prime_1\) formed at an elevated availability of Al. Each of those structures has a different transformation feature, and the difference between their hysteresis temperatures is about 20 °C [9, 10]. From this perspective, the amount of Al in these types of alloys has a considerable influence on the structure, along with effects on the produced properties. While Cu–Al–Ni SMAs have a high transformation temperature and better thermal stability, their applications are limited as a consequence of their low workability and suspicion of having a brittle intergranular crack [11], which could be a result of an extreme elastic anisotropy and large grain size. Consequently, it is designed to reduce the mechanical properties, and contributes to form a higher order degree in the parent phase with a different structure (B2, DO3, and/or L21). To overcome these particular problems, two ways have been identified by the researcher so far; by adding the alloying elements or heat treatment [12, 13, 14, 15], by which they are able to control grain size, and consequently having an effect on the mechanical properties [4, 16]. Adachi et al. [17] revealed that the Ti-doped in Cu–Al–Ni SMA can reduce grain size and leads to enhancing the mechanical properties, while Morris and Gunter [12] refer in their study to the fact that the addition of boron and zirconium can lead to an improvement in the mechanical properties. This paper aims to investigate the effects of various additions of Ti on the phase transformation characteristics, microstructures, and mechanical properties of Cu-11.9 % Al-4.5 % Ni (in mass%) SMA using a differential scanning calorimeter (DSC), field emission scanning electron microscopy (FESEM), transmission electron microscopy (TEM), energy dispersive spectroscopy (EDS), a tensile test, a shape memory test, and Vicker’s hardness.

Experimental

Material preparation

The alloy was produced by melting the high purity metals of Cu, Al, Ni, and Ti, using an induction furnace. These metals were melted in a silicon carbide crucible at a temperature of about 1,300 °C with continuous stirring, and were then poured into a cast iron mold with dimensions of 270 × 50 × 20 mm3. The ingot was cut into the required sample dimensions using EDM wire, then homogenized at 900 °C for 30 min, and then quenched in water, which led to the formation of martensite. The chemical composition analysis for the Cu–Al–Ni–Ti SMAs was investigated using inductively coupled plasma mass spectrometry (ICP-MS). These results are shown in Table 1.
Table 1

Chemical compositions of the alloys (mass fraction, mass%)

Alloy

Cu

Al

Ni

Ti

A

84.1

11.9

4

B

83.7

11.9

4

0.4

C

83.4

11.9

4

0.7

D

83.1

11.9

4

1

Material characterization

Flat specimens were cut from the aged samples with dimensions of 10 × 10 × 2 mm3 for the microstructural and X-ray diffraction (XRD) characteristics. TEM characterization in a JEM2010 operated at 200 kV was also used to investigate the microstructural changes. Filings of the alloys removed were about 2–6 mg, and were taken for the differential scanning calorimeter measurements using a Mettler Toledo DSC 822e; the heating and cooling rate was 10 K min−1 within the 323–573 K range. The phase identifications and crystal structure determinations were carried out at the room temperature using a D5000 Siemens X-ray diffractometer fitted with a CuKα X-ray source with a locked couple mode, a 2θ range of between 30°–75°, and a 0.05o s−1 scanning step. The quenched samples were ground and polished, then etched in a solution containing 2.5 g ferric chloride acid (FeCl3·6H2O) and 48 mL methanol (CH3OH) in 10 mL HCl for 4 min [9, 18, 19].

Mechanical tests

Tensile and hardness tests

The tensile test was performed using an Instron 5982-type universal testing machine operated at a constant strain rate of 0.1 mm min−1. The tests were carried out at room temperature until a failure occurred, then the fracture stress–strain was determined under the tensile load. The purpose of this test was to optimize the elastic and plastic area ranges before performing the shape memory effect test. Vicker’s microhardness test with 10 kg for 25 s was performed to measure the hardness of these alloys.

Shape memory effect test

The shape memory effect test was carried out using a specially designed machine as shown in Fig. 1, whereas the specially designed contents were analyzed using an Instron 5982-type universal testing machine, operated with special program parameters according to the shape memory test, which was connected with a heater tape and digital thermocouple to control the applied temperature, as well as an external extensometer to measure the shape extension and recovery. The tests were carried out at a temperature below Mf, which was about 100 °C, where the alloys would be able to obtain shape recovery. Then, the deformed sample that still had an unrecoverable shape was subsequently heated above the austenite finish temperature (A f + 60 °C), i.e., 300 °C for 10 min, followed by water quench to recover the residual strain (ε r). The recovered shape was attributed to the transformation of the detwinned martensite to the austenite phase, which had been termed as a transformation strain (ε t). After the cooling process, the martensite again formed into a self-accommodated structure.
Fig. 1

Schematic illustration of the shape memory effect test

Results and discussion

Microstructural observations

In order to investigate the effect of Ti additions on the microstructure of Cu–Al–Ni SMA, field emission scanning electron microscopy was used, as shown in Fig. 2. It was found that all the microstructures obtained a martensite phase, whereas the main difference between the captured micrographs was that the martensite formed in different morphologies and types along with many precipitations which occurred with the Ti additions. On the other hand, it was observed that grain size reduced with Ti additions; the grain size of the base alloy Cu–Al–Ni was about 1,350 µm, and refined to 900, 400, and 650 µm, corresponding with the 0.4 mass% Ti, 0.7 mass% Ti, and 1 mass% Ti additions, respectively. From this point of view, it was proven that grain size reduced by about 70 % after the Ti added, which was obtained with the alloy of 0.7 mass% Ti. This attribution may show that the Ti is able to diffuse very fast into the microstructure and accumulate at the grain boundaries, and then restrict grain growth. From Fig. 2a, it can be noticed that the microstructure consisted of two types of martensite, which were \(\beta^\prime_1\) with an 18R structure and \(\gamma^\prime_1\) with a 2H structure. The \(\gamma^\prime_1\) appearing as parallel martensite morphologies are known as a lamella structure [20]; these types of lamella morphology have also grown into grain, while the \(\beta^\prime_1\) phase is typically formed with self-accommodating groups in two different morphologies, plates and needles. Fig. 2b shows the microstructure of the Cu–Al–Ni with a 0.4 Ti addition, which consists of a \(\beta^\prime_1\) only along with a new precipitation that formed in a regular and irregular shape, and in which the regular shape looks like a flower shape, and the distribution of these precipitations is either between the \(\beta^\prime_1\) plates and the needles, or in an individual area. The fine flower shape precipitates as seen in alloys B, C, and D are known as the X-phase [21, 22], as indicated by the white dots. Recently, Ratchev et al. [23] confirmed the existence of this phase after the addition of Ti, and they mentioned that this phase has an equivalent composition and structure of the X phase; it has also been noticed that this phase has a strong influence on the stacking sequence of the martensite plates. According to an EDS analysis of the X-phase area, the X-phase precipitations are Ti-rich, as has been indicated in Fig. 3. It has also been proven that the microstructure still consists of a Ti; in other words, no shrinkage occurred for the Ti element in the microstructure. It can be noticed that with increasing the percentage of the Ti addition to 0.7 mass%, the volume fraction of the X-phase increased, and its shape changed from a regular flower to an irregular one and a circular shape as shown in Fig. 2c. However, this implies that the precipitation density increased remarkably with an increase in the percentage of Ti. When the percentage of Ti addition is further increased to 1 mass%, the X-phase starts to dissolve and creates visualized boundaries. Moreover, these boundaries impeded the growth of \(\beta^\prime_1\) as shown in Fig. 2d. When the reduction of both precipitates’ dispersion and grain sizes occurred, the interface mobility can lead to a reduction in the martensite plate sizes [24]. Even though the precipitates are capable of restricting the growth of the martensite plates, they are also able to widen the transformation hysteresis [25, 26].
Fig. 2

FESEM micrographs showing the microstructures of the Cu–Al–Ni SMA with different concentration of Ti additions: a Cu–Al–Ni (alloy A), b Cu–Al–Ni–0.4 mass% Ti (alloy B), c Cu–Al–Ni–0.7 mass% Ti (alloy C), d Cu–Al–Ni–1 mass% Ti (alloy D)

Fig. 3

EDS analysis of the Cu–Al–Ni–1 mass% Ti a Micrograph of scanned area, bspectrum 1, c spectrum 2

Figure 4 shows TEM bright images, HRTEM images, and selected area diffraction patterns of alloys A and C. The TEM image (Fig. 4a) clearly reveals the existence of \(\gamma^\prime_1\) and \(\beta^\prime_1\) martensite in the microstructure, whereas the HRTEM images Fig. 4b show the lattice fringes of the martensite phase. The SADP images show the diffraction patterns of the base alloy that presented a monoclinic structure, as shown inset in Fig. 4c, which is identical with the XRD results. Figure 4d–f shows the TEM bright images along with the SADP images and the HRTEM images of alloy C, where both \(\gamma^\prime_1\) and \(\beta^\prime_1\) phases were observed, corresponding with the existence of the X-phase (AlNi2Ti and Ti3·3Al compounds) on the planes (332) and (400); in which a similar structure with alloy A was confirmed. The HRTEM images (Fig. 4e) show the fringe structure and orientation of alloy C, and it was observed that the distance between the fringes increased and orientated regularly parallel with the matrix after the Ti addition.
Fig. 4

TEM images corresponding with selected area diffraction patterns of Alloy A and Alloy C: a Bright field of TEM image of Alloy A. b HRTEM image of the Alloy A. c SADP of the alloys A. d Bright field of TEM image of Alloy C, e HRTEM image of the Alloy C. f SADP of the alloys C

With the purpose of showing the distribution of the alloying element, the FESEM along with elemental mapping was taken for alloy D, as shown in Fig. 5a–e. It was noticed that the element distribution of alloy D (Cu, Al, and Ni) was homogenized, and the Ti was distributed on the grain boundaries as shown in Fig. 5c, especially the phenomena obtained after the Ti amount increased up to 1 mass%, so the Ti started to diffuse and dissolve at grain boundaries.
Fig. 5

Elemental distribution map showing the distribution of Cu, Al, Ni, and Ti in the Cu–Al–Ni–1 mass% Ti SMA

To investigate the effect of the Ti addition on the phase transformation characteristics along with the crystal size, the XRD diffraction patterns were analyzed, as shown in Fig. 6 and Table 2. The peak patterns of the base alloy confirmed the existence of two metastable phases, \(\gamma^\prime_1\) and \(\beta^\prime_1\). The (200) peak represented the \(\gamma^\prime_1\) , and the other peaks represent the \(\beta^\prime_1\) phase. However, it was observed that a significant change happened in the present patterns after the Ti was added; these changes are represented by shifting the peaks and/or increasing the value of intensity, depending on the percentage of the Ti additions. The indexing patterns of alloy B showed patterns almost similar to alloy A, along with a disappearance of the (200) γ1', where two additional peaks were added to the scanning patterns (1210) and (2010), which are associated with \(\beta^\prime_1\) . This observation seems to be as a result of the Ti addition. Furthermore, it was shown that with increasing the percentage of Ti additions, the (128) and (0018) peaks shifted to the right, while (202) and (122) shifted toward the left. In addition, there were two peaks, (332) and (400), observed at 2 theta of 63.5 and 64.4o, which represented AlNi2Ti and Ti3·3Al compounds, which in turn represented the X-phase. By increasing the percentage of Ti to 1 mass%, these peaks disappeared progressively. This is attributed to a low density of X-phases in the microstructure.
Fig. 6

X-ray diffraction patterns of a Cu–Al–Ni (alloy A), b Cu–Al–Ni–0.4 mass% Ti (alloy B), c Cu–Al–Ni–0.7 mass% Ti (alloy C), d Cu–Al–Ni–1 mass% Ti (alloy D) at the room temperature

Table 2

Lattice parameters and crystallite size of Cu–Al–Ni SMA with and without additions

Alloy

a

b

c

a/b

Crystallite size/nm

A

4.425

5.250

38.055

0.842

25.5

B

4.421

5.247

37.912

0.842

27.4

C

4.440

5.237

38.043

0.847

25.2

D

4.439

5.240

38.032

0.847

32.3

The lattice parameters and crystallite size of the Cu–Al–Ni SMA with and without addition were determined from the XRD patterns and recorded in Table 2. The lattice parameters were evaluated in accordance with an orthorhombic 18R structure, which was proven by the XRD indexing patterns. Thus, the lattice parameters were determined using the following relation [27]:
$$ \frac{1}{{d^{2} }} = \frac{1}{{a^{2} }}\left[ {\frac{{h^{2} }}{{{ \sin }^{2} \beta }}} \right] + \frac{{k^{2} }}{{b^{2} }} + \frac{1}{{c^{2} }} + \left[ {\frac{{l^{2} }}{{{ \sin }^{2} \beta }}} \right] - \frac{2hl\cos \beta }{{ac{ \sin }^{2} \beta }}. $$
(1)

The ratio a/b was also calculated and indicated in Table 2. Whereas this ratio for the base alloy and modified alloys was less than \( \sqrt 3 /2 \) in an ordered case due to the atomic sizes of the constituent atoms of the 18R martensite [28, 29], and which depended on the obtained results. The XRD of the alloys having a monoclinic structure, which was similar to other work done earlier [30, 31]. Although the XRD results showed similar characteristics, some of the diffraction pattern planes shifted in the location during the addition process. One of the most important factors that has a significant effect on martensitic transformation is the structure ordering [31, 32].

In a Cu-based SMA, the martensitic transformation in a β-phase form is evaluated according to (110)β planes of the austenite phase, which is well known as a basal plane for the martensite phase [33]. The basal plane in the original case has a rectangular structure with a DO3/L21 order, and during the martensite transformation it transforms into a hexagonal structure corresponding to a hexagonal distortion [31, 33]. In general, the equilibrium phases of the Cu–Al–Ni SMA are formed from α, β, and γ phases. During the alloying element additions, some atomic migration may occur, thus allowing precipitation of any of the aforementioned phases. These diffusion processes are mainly controlled by the chemical compositions of the austenite phases, and would be important for this study since we varied the composition of the addition element. Based on these phenomena, the explanation of phase transformation behavior variations would be more realizable. The crystallite size was determined by a Scherrer equation [34, 35] as follows:
$$ {\text{Crystallite size}} \,\left( d \right) = \frac{0.9*(\lambda )}{\beta * \cos \theta } $$
(2)
where λ is the XRD wavelength, β is the full width at half maximum, and θ is the Bragg’s angle. The crystallite sizes tend to increase/decrease during the additions. However, it was indicated that the highest crystallite size was with a 1 mass% Ti addition.

Transformation temperatures

Figure 7 shows the DSC results of alloys A, B, C, and D for heating and cooling processes. The transformation temperatures of the austenite \( \rightleftarrows \) martensite transformation were determined from the DSC curves and shown in the Table 3. During the heating process, there was only one peak observed, and the reverse of this peak was observed during the cooling process. The observed peaks shifted toward high transformation temperatures, and varied in shape after the Ti additions. In addition, it is observed that small exothermic peaks occurred after the main peak on the cooling curve but they are absent in the DSC curve for the heating process. These peaks are mainly attributed to the inter-martensitic transformations which have been suppressed upon heating. It is well known that martensite transformation is a diffusionless transformation that occurs between a high transformation temperature phase and a low transformation temperature phase corresponding to a first-ordered-phase structure in a crystalline solid; in other words, the structured atoms are able to move into the structure cooperatively.
Fig. 7

DSC curves of the Cu–Al–Ni SMA with different percentage of Ti addition: a Heating. b Cooling

Table 3

Transformation temperature of Cu–Al–Ni SMA with and without Ti additions

Alloy

Transformation temperatures

A s/K

A f/K

M s/K

M f/K

T o/K

A

504.76

515.14

503.98

496.94

509.56

B

506.62

518.56

506.91

499.14

512.735

C

509.35

520.1

508.78

502.35

519.44

D

505.3

515.88

504.74

498.2

510.31

As seen from the DSC graphs, the curves of both the heating and cooling processes became wider and smoother after the Ti additions, which may be attributed to the presence of the X-phase in the microstructure. Martensite stabilization may also obtain an important role for the phenomena of shifting, which was altered after the Ti was added. The highest transformation temperatures (A s, A f, M s, and M f) were observed with alloy C due to the high density of X-phase presences. From this point of view, it was proven that the existence of an X-phase can enhance the transformation temperatures. However, one of the typical properties of thermoelastic martensite transformation is hysteresis.

The following equation was used to express the equilibrium temperature (T o) between the martensitic and austenitic phases:
$$ \begin{gathered} \Delta G^{{\text{M}} \to {\text{A}}} \left( {T_{\text{o}} } \right) = G^{{\text{A}}} \left( {T_{\text{o}} } \right) - G^{{\text{M}}} \left( {T_{\text{o}} } \right) = \hfill \\ \left( {H^{{\text{A}}} - T_{\text{o}} S^{{\text{A}}} } \right) - \left( {H^{{\text{M}}} - T_{\text{o}} S^{{\text{M}}} } \right) = \Delta H^{{\text{M}} \to {\text{A}}} - T_{\text{o}} \Delta S^{{\text{M}} \to {\text{A}}} \hfill \\ \end{gathered} $$
(3)
or
$$ \begin{gathered} \Delta G^{{\text{A}} \to {\text{M}}} \left( {T_{\text{o}} } \right) = G^{{\text{M}}} \left( {T_{\text{o}} } \right) - G^{{\text{A}}} \left( {T_{\text{o}} } \right) = \hfill \\ \left( {H^{{\text{M}}} - T_{\text{o}} S^{{\text{M}}} } \right) - \left( {H^{{\text{A}}} - T_{\text{o}} S^{{\text{A}}} } \right) = \Delta H^{{\text{A}} \to {\text{M}}} - T_{\text{o}} \Delta S^{{\text{A}} \to {\text{M}}} . \hfill \\ \end{gathered} $$
(4)
In addition T o for the alloy, when the \( \Delta G^{M \to A} \left( {T_{o} } \right) \; \text{and}\; \Delta G^{A \to M} (T_{o} ) = 0 \) at T = T o is expressed as [36, 37]
$$ T_{o} = \frac{{\Delta H^{{\text{M}}\,\rightarrow\,{\text{A}}} }}{{\Delta S^{{\text{M}}\,\rightarrow\,{\text{A}}} }}\; {\rm or}\; T_{o} = \frac{{\Delta H^{{\text{A}}\,\rightarrow\,{\text{M}}} }}{{\Delta S^{{\text{A}}\,\rightarrow\,{\text{M}}} }} $$
(5)
The equilibrium temperature between the austenite and martensite phase transformation can be expressed as:
$$ T_{\text{o}} = \frac{1}{2}(M_{\text{s}} + A_{\text{f}} ) $$
(6)
The T o was calculated before and after the additions and recorded in Table 3, and it was found that the maximum value of T o was observed with alloy C. Moreover, the thermodynamics parameters of forward and reverse transformation were calculated and indicated in Table 4. The enthalpy and entropy of both transformations increased gradually with the Ti additions, and obtained a maximum value with alloy D. As thermal stability is inversely proportional to enthalpy, it was found that the martensite thermal stability was decreased gradually with increase the Ti addition, which may be related to the percentage of precipitation, the behavior, and the morphology of the martensite phase.
Table 4

Thermodynamic parameters of Cu–Al–Ni SMA with and without Ti additions

Alloy

ΔH forward/J g−1

ΔH reverse/J g−1

ΔS forward/J kg−1 K−1

ΔS reverse/J kg−1 K−1

A

1.947

1.424

3.821

2.795

B

2.267

1.439

4.422

2.807

C

2.659

1.432

5.170

2.785

D

3.892

2.180

7.627

4.271

Mechanical properties

Tensile and hardness test

Figure 8 shows the stress–strain curve obtained at room temperature with a constant strain of 0.1 mm min−1. The measured values from the stress–strain curves are given in Table 5. The Cu–Al–Ni with different percentage additions of the Ti exhibited a classical feature of shape memory alloys through obtaining a distinct elastic region followed by a linear plastic deformation region, as shown from the σ − ɛ curves in the Fig 8. The generation of latter regions is mainly attributed to the induced martensite–martensite deformation region and the reorientation of the martensite detwinning/variants.
Fig. 8

Stress-strain curves obtained from the tensile test performed at the room temperature

Table 5

Results obtained from the tensile, hardness, and shape memory tests on the Cu–Al–Ni SMA with and without Ti additions

Alloy

Fracture stress (σ F)/MPa

Fracture strain (ε F)/%

Hardness/MPa

Strain recovery by SME (ε SME)/%

A

270 ± 10

1.65 ± 0.07

253.6 ± 4

50

B

500 ± 10

2 ± 0.095

245 ± 2

66

C

710 ± 11

3.2 ± 0.1

228.2 ± 1

100

D

650 ± 6

2.8 ± 0.2

236.4 ± 2.4

80

On the other hand, the fracture stress–strain of the base alloy showed a significant variation after the Ti added. Alloys B, C, and D possessed considerably higher fracture stress and strain than alloy A; in other words, the ductility of the addition alloys improved from 1.65 to 3.2 %, which is almost twice than a base alloy. The increase in ductility with the addition of the Ti element may be attributed to the significant microstructural changes and presence of X-phase associated with decrease in the degree of order of the alloys, whereas this decrement can cause an increase in the ductility and workability of the alloys [38, 39].

Figure 9 shows the fracture surfaces of the Cu–Al–Ni SMAs with and without the Ti additions. Alloy A exhibited a typical characteristic of brittle fractures as shown in Fig. 9a. Alloy B with 0.4 mass% of Ti demonstrated an intergranular fracture along with a very small transgranular fracture, as shown in Fig. 9b. While alloys C and D showed a mixed mode of transgranular and intergranular fractures, obtaining a significant increase in the value of ductility compared with alloy B, as shown in Fig. 9c, d. This ductility improvement was due to the decrement of grain size, which occurred after the Ti additions. The fracture mode of alloys C and D completely agreed with the behavior of the stress–strain (Fig. 8). The large grain size and large elastic anisotropy of Cu–Al–Ni SMA, along with difference in crystal orientations, easily increased the stress concentration at the grain and grain boundaries, thus causing a brittle fracture in the alloy. This study indicates controlling grain size during the addition of Ti elements to reduce the stress concentration in grain and grain boundaries, thus enhancing the ductility of the alloys. With improved ductility, the workability of these alloys can be enhanced, which can be beneficial for engineering applications. The hardness values of the Cu–Al–Ni SMAs with and without additions are shown in Table 5. It was observed that different concentrations of Ti additions had a significant effect on the hardness property. As the addition of the Ti to Cu–Al–Ni SMA can lead to a decrease in grain size and create a new phase (X-phase) into the microstructure, so this decrement is accompanied by a reduction in the values of hardness.
Fig. 9

Fracture surface of the alloys: a Cu–Al–Ni (alloy A), b Cu–Al–Ni–0.4 mass% Ti (alloy B), c Cu–Al–Ni–0.7 mass% Ti (alloy C), d Cu–Al–Ni–1 mass% Ti (alloy D)

Shape memory effect

The strain recovery by the shape memory effect (ε SME) of the Cu–Al–Ni SMAs with and without the Ti additions was determined by using a specially designed tensile test at T < M f, whose values are given in Table 5. The ε SME varied with the variations in the amount of Ti at the initial strain value of 1.55 %, as shown in Fig. 10. It was observed that the addition of Ti with different mass percentages exhibited an increase in the values of strain recovery by the SME. Alloy B with a percentage of 0.7 mass% Ti exhibited almost a complete shape recovery after being preheated to T > A f. In addition, the strain recovery rose from 50 to 100 %. These enhancements in the strain recovery were attributed to the existence of the X-phase that was brought about by the Ti addition in the parent phase. Furthermore, the percentages of shape recovery were mainly dependent on the extent of martensitic transformation of the alloys as noticed from the XRD peaks, which indicated a higher intensity after the Ti was added.
Fig. 10

Shape memory effect curves of the alloys performed at T < M f, then preheated to T > A f to obtain the shape recovery

Conclusions

  1. 1.

    The microscope observation demonstrates that the addition of Ti to Cu–Al–Ni SMA contributes to producing a new phase along with possessing a grain refinement. Grain size was reduced by up to 70 % after an addition corresponding to varying the morphology, structure, and order of the martensite phase. Moreover, it was discovered that after increasing the Ti addition to 1 mass%, the X-phase begins to dissolve while producing visualized grain boundaries.

     
  2. 2.

    The Ti addition to Cu–Al–Ni SMA influences the phase transformation temperatures, and which tends to increase these particular temperatures after its addition. Furthermore, it was found that the highest transformation temperatures came with a 0.7 mass% of Ti as a result of the presence of the high percentage X-phase in the microstructure. Nevertheless, the thermodynamic parameters, such as enthalpy and entropy, increase after the Ti addition.

     
  3. 3.

    When the Ti is added to Cu–Al–Ni SMA, the fracture stress–strain curve increases and reaches its maximum values with alloy C (0.7 mass% Ti). Furthermore, it was observed that these values begin to decrease after increasing the Ti addition to 1 mass%, which may occasionally be attributed to the dissolution of the X- phase into the microstructure.

     
  4. 4.

    Shape recovery improves after the Ti addition, while strain recovery increases from 50 to 100 %. These improvements are due to the presence of the X-phase together with the variations in the structures and morphologies of Cu–Al–Ni SMA after addition. The highest strain recovery was realized with alloy C (0.7 mass%), which obtained a complete shape recovery.

     
  5. 5.

    With the Ti additions, the values of hardness lead to a decrease due to the existence of the X-phase along with a decrement in grain size.

     

Notes

Acknowledgements

The author(s) would like to thank the Malaysian Ministry of Higher Education (MOHE) and Universiti Teknologi Malaysia for providing the financial support and facilities for this research, under Grant No. R.J130000.7824.4F150.

References

  1. 1.
    Ibarra A, Juan JS, Bocanegra EH, Nó ML. Thermo-mechanical characterization of Cu–Al–Ni shape memory alloys elaborated by powder metallurgy. Mater Sci Eng A. 2006;438–440:782–6.CrossRefGoogle Scholar
  2. 2.
    Recarte V, Pérez-Landazábal JI, Nó ML, San J. Juan, Study by resonant ultrasound spectroscopy of the elastic constants of the β phase in Cu–Al–Ni shape memory alloys. Mater Sci Eng A. 2004;370:488–91.CrossRefGoogle Scholar
  3. 3.
    Font J, Cesari E, Muntasell J, Pons J. Thermomechanical cycling in Cu–Al–Ni-based melt-spun shape-memory ribbons. Mater Sci Eng A. 2003;354:207–11.CrossRefGoogle Scholar
  4. 4.
    Pérez-Landazábal JI, Recarte V, Sánchez-Alarcos V, Nó ML, Juan JS. Study of the stability and decomposition process of the β phase in Cu–Al–Ni shape memory alloys. Mater Sci Eng A. 2006;438–440:734–7.CrossRefGoogle Scholar
  5. 5.
    Saburi T, Wayman CM. Crystallographic similarities in shape memory martensites. Acta Metall. 1979;27:979–95.CrossRefGoogle Scholar
  6. 6.
    Van Humbeeck J, Van Hulle D, Delacey L. two-stage martensite transformation in a Cu-13. 99 mass Al-3. 5 mass% Ni alloy. Trans Jpn Inst Met. 1987;28:383–91.CrossRefGoogle Scholar
  7. 7.
    Gastien R, Corbellani C, Alvarez Villar H, Sade M, Lovey F. Pseudoelastic cycling in Cu–14.3 Al–4.1 Ni (wt%) single crystals. Mater Sci Eng A. 2003;349:191–6.CrossRefGoogle Scholar
  8. 8.
    Gastien R, Corbellani C, Sade M, Lovey F. Thermal and pseudoelastic cycling in Cu–14.1 Al–4.2 Ni (wt%) single crystals. Acta Mater. 2005;53:1685–91.CrossRefGoogle Scholar
  9. 9.
    Sarı U, Aksoy İ. Electron microscopy study of 2H and 18R martensites in Cu–11.92 wt% Al–3.78 wt% Ni shape memory alloy. J Alloy Compd. 2006;417:138–42.CrossRefGoogle Scholar
  10. 10.
    Recarte V, Pérez-Sáez RB, San Juan J, Bocanegra EH, Nó ML. Influence of Al and Ni concentration on the Martensitic transformation in Cu–Al–Ni shape-memory alloys. Metall Mater Trans A. 2002;33:2581–91.CrossRefGoogle Scholar
  11. 11.
    Wei ZG, Peng HY, Yang DZ, Chung CY, Lai JKL. Reverse transformations in CuA1NiMnTi alloy at elevated temperatures. Acta Mater. 1996;44:1189–99.CrossRefGoogle Scholar
  12. 12.
    Morris MA, Gunter S. Effect of heat treatment and thermal cycling on transformation temperatures of ductile Cu–Al–Ni–Mn–B alloys. Scripta Metallurgica et Materiala. 1992;26(11):1663–8.CrossRefGoogle Scholar
  13. 13.
    Sarı U, Kırındı T. Effects of deformation on microstructure and mechanical properties of a Cu–Al–Ni shape memory alloy. Mater Charact. 2008;59:920–9.CrossRefGoogle Scholar
  14. 14.
    Yildiz K, Kok M. Study of martensite transformation and microstructural evolution of Cu–Al–Ni–Fe shape memory alloys. J Therm Anal Calorim. 2014;115:1509–14.CrossRefGoogle Scholar
  15. 15.
    Karagoz Z, Canbay CA. Relationship between transformation temperatures and alloying elements in Cu–Al–Ni shape memory alloys. J Therm Anal Calorim. 2013;114:1069–74.CrossRefGoogle Scholar
  16. 16.
    Sutou Y, Omori T, Yamauchi K, Ono N, Kainuma R, Ishida K. Effect of grain size and texture on pseudoelasticity in Cu–Al–Mn-based shape memory wire. Acta Mater. 2005;53:4121–33.CrossRefGoogle Scholar
  17. 17.
    Adachi K, Shoji K, Hamada Y. Formation of (X) phases and origin of grain refinement effect in Cu–Al–Ni shape memory alloys added with titanium. ISIJ Int. 1989;29:378–87.CrossRefGoogle Scholar
  18. 18.
    Sari U. Influences of 2.5 wt% Mn addition on the microstructure and mechanical properties of Cu–Al–Ni shape memory alloys. Int J Miner Metall Mater. 2010;17:192–8.CrossRefGoogle Scholar
  19. 19.
    Chang SH. Influence of chemical composition on the damping characteristics of Cu–Al–Ni shape memory alloys. Mater Chem Phys. 2011;125:358–63.CrossRefGoogle Scholar
  20. 20.
    Nishiyama Z, Fine ME, Wayman CM. Martensitic transformation. New York: Academic Press; 1978.Google Scholar
  21. 21.
    Hurtado I, Ratchev P, Van Humbeeck J, Delaey L. A fundamental study of the X-phase precipitation in Cu–Al–Ni–Ti–(Mn) shape memory alloys. Acta Mater. 1995;44:3299–306.CrossRefGoogle Scholar
  22. 22.
    Lee JS, Wayman CM. Grain refinement of a Cu–Al–Ni shape memory alloy by Ti and Zr additions. Trans Jpn Inst Met. 1986;27:584–91.CrossRefGoogle Scholar
  23. 23.
    Ratchev P, Van Humbeeck J, Delaey L. On the formation of 2H stacking sequence in 18R martensite plates in a precipitate containing CuAlNiTiMn alloy. Acta Metall Mater. 1993;41(8):2441–9.CrossRefGoogle Scholar
  24. 24.
    Malarría J, Elgoyhen C, Vermaut P, Ochin P, Portier R. Shape memory properties of Cu-based thin tapes obtained by rapid solidification methods. Mater Sci Eng A. 2006;438–440:763–7.CrossRefGoogle Scholar
  25. 25.
    Dutkiewicz J, Pons J, Cesari E. Effect of γ precipitates on the martensitic transformation in Cu–Al–Mn alloys. Mater Sci Eng A. 1992;158:119–28.CrossRefGoogle Scholar
  26. 26.
    Bouabdallah M, Baguenane-Benalia G, Saadi A, Cheniti H, Gachon J-C, Patoor E. Precipitation sequence during ageing in β1 phase of Cu–Al–Ni shape memory alloy. J Therm Anal Calorim. 2013;112:279–83.CrossRefGoogle Scholar
  27. 27.
    Xuan JBQ, Hsu TY. The effect of martensite ordering on shape memory effect in a copper–zinc–aluminium alloy. Mater Sci Eng. 1987;93:205–11.CrossRefGoogle Scholar
  28. 28.
    Salzbrenner RJ, Cohen M. On the thermodynamics of thermoelastic martensitic transformations. Acta Metall. 1979;27:739–48.CrossRefGoogle Scholar
  29. 29.
    Adigüzel O. Martensite ordering and stabilization in copper based shape memory alloys. Mater Res Bull. 1995;30:755–60.CrossRefGoogle Scholar
  30. 30.
    Sampath V. Studies on the effect of grain refinement and thermal processing on shape memory characteristics of Cu–Al–Ni alloys. Smart Mater Struct. 2005;14:S253–60.CrossRefGoogle Scholar
  31. 31.
    Balo ŞN, Sel N. Effects of thermal aging on transformation temperatures and some physical parameters of Cu-13.5 wt% Al-4 wt% Ni shape memory alloy. Thermochim Acta. 2012;536:1–5.CrossRefGoogle Scholar
  32. 32.
    Yang G-S, Lee J-K, Jang W-Y. Effect of grain refinement on phase transformation behavior and mechanical properties of Cu-based alloy. Trans Nonferrous Met Soc China. 2009;19:979–83.CrossRefGoogle Scholar
  33. 33.
    Aydogdu A, Aydogdu Y, Adiguzel O. Long-term ageing behaviour of martensite in shape memory Cu–Al–Ni alloys. J Mater Process Technol. 2004;153–154:164–9.CrossRefGoogle Scholar
  34. 34.
    Cullity BD, Stock SR. Elements of X-ray diffraction. Upper Saddle River: Prentice Hall; 2001.Google Scholar
  35. 35.
    Patterson AL. The Scherrer Formula for X-ray Particle Size Determination. Phys Rev. 1939;56:978–82.CrossRefGoogle Scholar
  36. 36.
    Tatar C. Gamma irradiation-induced evolution of the transformation temperatures and thermodynamic parameters in a CuZnAl shape memory alloy. Thermochim Acta. 2005;437:121–5.CrossRefGoogle Scholar
  37. 37.
    Ortín J, Planes A. Thermodynamics of thermoelastic martensitic transformations. Acta Metall. 1989;37:1433–41.CrossRefGoogle Scholar
  38. 38.
    Mallik US, Sampath V. Influence of aluminum and manganese concentration on the shape memory characteristics of Cu–Al–Mn shape memory alloys. J Alloy Compd. 2008;459:142–7.CrossRefGoogle Scholar
  39. 39.
    Saiji Matsuoka MH, Oshima R, Eiichi Fujita F. Improvement of ductility of melt spun Cu–Al–Ni shape memory alloy ribbons by addition of Ti or Zr. Jpn J Appl Phys. 1983;22:528–30.CrossRefGoogle Scholar

Copyright information

© Akadémiai Kiadó, Budapest, Hungary 2014

Authors and Affiliations

  • Safaa N. Saud
    • 1
  • E. Hamzah
    • 1
  • T. Abubakar
    • 1
  • Mohd Zamri
    • 1
    • 2
  • Masaki Tanemura
    • 2
  1. 1.Faculty of Mechanical EngineeringUniversiti Teknologi MalaysiaJohorMalaysia
  2. 2.Department of Frontier MaterialsNagoya Institute of TechnologyNagoyaJapan

Personalised recommendations