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Effect of γ2 phase evolution on mechanical properties of continuous columnar-grained Cu–Al–Ni alloy
Materials Science and Engineering A 532 (2012) 536–542
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of 2 phase evolution on mechanical properties of continuous columnar-grained Cu–Al–Ni alloy Z. Wang, X.F. Liu, J.X. Xie ∗ Key Laboratory for Advanced Materials Processing (MOE), University of Science and Technology Beijing, Beijing 100083, China
a r t i c l e
i n f o
Article history: Received 21 July 2011 Received in revised form 2 November 2011 Accepted 3 November 2011 Available online 11 November 2011 Keywords: Shape memory alloy Continuous columnar grain 2 phase Heat treatment Martensitic transformation
a b s t r a c t Effect of 2 phase evolution on mechanical properties of Cu–14%Al–3.8%Ni (mass fraction) alloy wires fabricated by continuous unidirectional solidification technology was investigated before and after heat treated at 700–780 ◦ C. Mechanism for the improvement of mechanical properties of the alloy was analyzed. It was found that the alloy retained continuous columnar grains after heat treatment. With the heat treatment temperatures increasing from 700 ◦ C to 770 ◦ C, the coarse dendrite 2 phase evolved into the fine polygonal, ellipsoidal and spherical particles in the grains, while the long-banding, discontinuous block and ellipsoidal particles at the grain boundaries. The average size of the 2 phase decreased from 20 m before heat treatment down to 2 m, and its amount reduced from 50.0% down to 0.8%. The 2 phase was full dissolution at 780 ◦ C. The tensile strength of the alloy treated at 700–780 ◦ C ranged from 577 MPa to 710 MPa. The elongation of the alloy ranged from 7.5% to 23.8% and had a peak value at 760 ◦ C. Excellent balance between strength and elongation could be obtained when the spherical 2 phase was distributed throughout the alloy with the size smaller than 5 m and the amount in the range of 11.4–16.6%. The nano-hardness and elastic modulus of the 2 phase decreased gradually with the reduction of the amount, leading to the improvement of mechanical properties. © 2011 Elsevier B.V. All rights reserved.
1. Introduction In recent years, Cu–Al–Ni shape memory alloys (SMA) have attracted an increasing attention due to their low cost, excellent aging stability and thermal stability as well as higher operating temperature [1,2]. However, the engineering application of the alloy has been restricted seriously because equiaxed polycrystalline Cu–Al–Ni alloy is prone to intergranular fracture during plastic deformation due to the large grain size and the existence of brittle 2 phase [3]. Among the manufacturing methods available for improving mechanical properties of the alloy, the continuous unidirectional solidification technology is generally accepted as one of the most effective routes, which is based on the special competitive growth mechanism of crystal and can generate cast products with a beneficial texture [4]. Therefore, the highly oriented grains are deformed coherently under the axis load, which is helpful to avoid severe stress concentration at the grain boundaries. At present, columnar-grained Cu–Al–Ni alloy wires and pipes have been successfully fabricated by continuous unidirectional solidification technology [5–7].
∗ Corresponding author. Tel.: +86 10 62332254; fax: +86 10 62332253. E-mail address: [email protected] (J.X. Xie). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.11.019
In the study of grain refinement for Cu–Al–Ni alloy after hightemperature deformed, it was noted that the equiaxed grain growth during subsequent hot-working or heat treatment was retarded by the spherical 2 particles (about 5 m), which is distributed throughout the microstructure. Meanwhile, the fine 2 phase would not be expected to have a harmful effect on the mechanical properties of the alloy [8]. Nevertheless, when the solidification rate of the alloy was slow, namely the alloy had a poor cooling ability, numerous coarsely block and dendritic 2 phase (about 10–15 m) would be precipitated easily in the grains or at the grain boundaries of the alloy. This leads to the enlargement of the grain boundary embitterment and the degradation of mechanical properties [9]. It is demonstrated that the variation of the distribution, amount, size, and shape of the 2 phase obtained under different conditions plays distinct roles on mechanical properties of Cu–Al–Ni alloy. In order to accurately control the 2 phase and mechanical properties of the continuous columnar-grained Cu–Al–Ni alloy, it is necessary to study the evolution of the 2 phase and clarify the influence of this evolution on the mechanical properties of the alloy. This paper focuses on the evolution of the 2 phase and corresponding mechanical properties of continuous columnar-grained Cu–14%Al–3.8%Ni (mass fraction) alloy wires before and after heat treatment (HT). The mechanism for the improvement of mechanical properties of the alloy was analyzed, thus to provide basis for accuracy control and reasonable utilization of the 2 phase in the Cu–Al–Ni alloy.
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Fig. 1. Microstructures in longitudinal cross sections of the alloy wires, in which the black arrows show the columnar-grained boundaries: (a) as-cast, (b) 700 ◦ C, and (c) 780 ◦ C.
2. Experimental procedure Cu–14 wt.%Al–3.8 wt.%Ni alloy wires with a diameter of 6 mm and smooth surface were prepared from pure Cu, Al and Ni of 99.99% purity by using the vacuum melting and argon-shield vertical continuous unidirectional solidification equipment [7]. The investigated alloy was heated to a given temperature in a boxtype resistance furnace and then quenched into room temperature water. Since phase in the alloy is difficult to be adjusted at a low temperature and the alloy oxidizes easily at a high temperature, the HT temperatures were set at 700, 730, 750, 760, 770 and 780 ◦ C, respectively. In order to improve the efficiency of heat treatment and realize the on-line heat treatment in the future, the holding time was kept for 1 min. Phase composition of the alloy was determined by an Xray diffractometer (XRD) with an operated voltage of 40 kV, a current of 150 mA and a scanning speed of 10◦ /min. The spatial group, the lattice parameters and the atom positions ¯ a = 0.8720 nm, ˛ = ˇ = = 90◦ ), of the 2 phase (Cu9 Al4 , P 43m, ¯ ˇ1 phase (DO3 , Fm3m, a = 0.5822, ˛ = ˇ = = 90◦ ), ˇ1 martensite phase (18R1 , C2/m, a = 3.8190 nm, b = 0.5190 nm, c = 0.4490 nm, ˛ = = 90◦ , and ˇ = 89.7◦ ), 1 martensite (2H, Pmmn, a = 0.5342 nm, b = 0.4224 nm, and c = 0.4390 nm) and ˛1 martensite (DO22 , I4/mmm, a = 0.3590 nm, and c = 0.7550 nm) were determined referring to Refs. [10–12]. The martenstic transformation (MT) temperatures (Mf , Ms , As , Af ) of the alloy treated at different HT temperatures were measured using a differential scanning calorimeter (DSC) with a cooling rate and heating rate of 10 ◦ C/min. Optical microscopy (OM), transmission electron microscopy (TEM) and scanning electron microscopy (SEM) equipped with an energy dispersive spectrometer (EDS) were employed to characterize the microstructures, the fracture morphologies and the composition in the remaining ˇ1 phase of the alloy. Uniaxial tensile tests at room temperature were performed along the axial direction of the alloy
wire at a rate of 0.02 mm/s, using a MTS810 testing machine. The hardness and elastic modulus of the 2 phase under different conditions were determined by a Nano Indenter II micro-mechanical probe. 3. Results 3.1. Effect of heat treatment temperature on evolution of the 2 phase of Cu–Al–Ni alloy Fig. 1 shows the microstructures of the continuous columnargrained Cu–Al–Ni alloy wires before and after heat treatment. It can be seen that the alloys treated at 700 ◦ C and 780 ◦ C retain columnar grains, and the grain boundaries (marked by the black arrows in Fig. 1) are clear and straight. The average grain size of the heattreated alloy is about 440 m, which is identical with that of the as-cast alloy. This implies that the grain size has no effect on the mechanical properties of the alloy in this study. From Fig. 2, the ascast alloy consists of mainly ˇ1 phase and 2 phase. The 2 dendrite phase is in a dense distribution with the amount of 55.7%, and the average dendrite-arm span is about 20 m (Fig. 1a). Fig. 3 shows the morphologies of 2 phase in Cu–Al–Ni alloy wires at different HT temperatures. When the HT temperature is 700 ◦ C, one can see that most of the coarse 2 dendrite phase in the grains evolves into the fine ellipsoidal particles, usually in the size of 2–5 m, while a small amount of it is transformed into the polygonal structure with a size range of 5–9 m. At the grain boundaries, the dendritic 2 phase are accumulated to the long-banded structure and their longitudinal axis is in accordance with casting direction, which may cause significant anisotropic properties. With the HT temperature increasing from 730 ◦ C to 770 ◦ C, the polygonal 2 phase in the grains is further changed into the ellipsoidal and spherical particles, and the long-banded ones at the grain boundaries evolve into the discontinuous block and
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Fig. 2. XRD patterns of the alloy wires.
ellipsoidal particles gradually. The average size of the ellipsoidal and spherical 2 phase is about 4 m at 760 ◦ C, and that of the spherical 2 phase is about 2 m at 770 ◦ C. As a result, the 2 phase becomes finer and more uniform in size and shape with the HT temperature increasing. Moreover, the amount of 2 phase decreases continuously with the HT temperature increasing, which is about 50.0% at 700 ◦ C, 29.2% at 730 ◦ C, 16.6% at 750 ◦ C, 11.4% at 760 ◦ C, 0.8% at 770 ◦ C, and full dissolution at 780 ◦ C, respectively, as shown in Fig. 3. Such behavior can also be demonstrated in Fig. 2 that the peak intensity of the 2 phase decreases, but that of the ˇ1 phase increases correspondingly with an increase of the HT temperature. At 780 ◦ C, the 2 phase cannot be detected in the XRD examinations, so the alloy is composed of only ˇ1 phase. The effect of heat treatment on the morphology evolution of 2 phase is essential the result of the breaking-up of the dendritic 2 structure at valleys due to spheroidizing and Ostwald ripening during heating and holding at elevated temperature. The dendritic structure firstly transforms to an ellipsoidal structure, and
Fig. 3. Morphologies of 2 phase in longitudinal cross sections of the alloy wires treated at different heat treatment temperatures: (a) 700 ◦ C, (b) 730 ◦ C, (c) 750 ◦ C, (d) 760 ◦ C, (e) 770 ◦ C, and (f) 780 ◦ C.
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Table 1 Aluminum concentration in the remaining ˇ1 phase under different conditions. Sample
As-cast
700 ◦ C
730 ◦ C
750 ◦ C
760 ◦ C
770 ◦ C
780 ◦ C
Al% (mass fraction)
12.79
13.08
13.41
13.78
13.96
14.07
14.1
subsequently spontaneous contracts to a spherical structure for the surface tension effect [13]. In addition, a higher temperature leads to a faster diffusion of the alloying element and breaking-up process. Therefore, the 2 phase is more uniform and finer in size and shape as well as fewer in amount at a higher temperature than that at a lower temperature.
The tensile stress-stain curves of the Cu–Al–Ni alloy wires under various conditions are shown in Fig. 6. It can be seen that the stressstain curve of the as-cast alloy is steep with high strain hardening rate and shows a typical characteristic of brittle fracture without
(a) 3.2. Effect of 2 phase evolution on transformation temperatures of Cu–Al–Ni alloy
3.3. Effect of 2 phase evolution on mechanical properties of Cu–Al–Ni alloy It is well known that distinct type of the MT may be stressinduced depending upon test temperature, when the Cu–Al–Ni alloy is composed of mainly ˇ1 phase [14,15]. In the temperature range between Mf and Ms , three distinct stress-induced MT, ˇ1 → 1 , 1 → ˇ1 and ˇ1 → ˛1 , were shown on the stress-strain curves. On the other hand, in the temperature range a little higher than Af , two distinct stress-induced MT, ˇ1 → ˇ1 and ˇ1 → ˛1 , may occur successively.
Heat flow (mW/mg)
730°C 750°C 760°C 770°C 780°C -100
(b)
Heat flow (mW/mg)
Fig. 4 shows the evolution of the characteristic MT temperatures with HT temperatures. It can be seen that the transformation temperatures first decrease rapidly with the HT temperature increasing from 700 ◦ C to 730 ◦ C and Mf < As < Ms < Af . Then they decrease slightly when the HT temperature is up to 750 ◦ C and Mf < Ms < As < Af . This behavior can be ascribed to the dissolution of 2 phase increasing the aluminum concentration in the remaining ˇ1 phase, as shown in Table 1, hence decreasing the MT temperature in the alloy. The more the aluminum concentration increases, the faster the MT temperature decrease. DSC data obtained around the martensitic phase transformation are plotted in Fig. 5. It can be found that the peaks during the heating or cooling process are very low and wide at 700 ◦ C, but become sharp and narrow with increasing the HT temperatures. Since the amount of the 2 phase decreases continuously with increasing temperature (Fig. 3), it can be indicated that the shape of the MT cycle becomes sharp and narrow with the reduction of the amount of the 2 phase.
Heating
700°C
-75
-50
-25
0
25
50
75
100
25
50
75
100
Temperature (ºC)
Cooling
700°C 730°C 750°C 760°C 770°C 780°C
-100
-75
-50
-25
0
Temperature (ºC) Fig. 5. DSC thermograms of the MT cycle for different heat treatments: (a) the heating process; (b) the cooling process.
900
100
Ms As Mf
50 0
700
RT=25°C
-50 -100
σt
800
Af
600
Stress (MPa)
Transformation temperature (ºC)
150
500
300 200 εt
100
700
720
740
Temperature (ºC)
760
780
Fig. 4. Transformation temperatures of the alloy wires treated at different heat treatment temperatures.
as cast 700°C × 1min 730°C × 1min 750°C × 1min 760°C × 1min 770°C × 1min 780°C × 1min
400
0
0
5
10
15
20
Strain (%) Fig. 6. Tensile stress-stain curves of the alloy wires.
25
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Fig. 8. Effect of heat treatment temperature on mechanical properties of the alloy wires.
Fig. 7. XRD patterns of the alloy wires after tensile fracture.
remarkable yield limit. Since the Mf , Ms , As , Af temperatures of the as-cast alloy are 17.5 ◦ C, 30.0 ◦ C, 62.8 ◦ C and 99.4 ◦ C, respectively, ˇ1 → 1 → ˇ1 → ˛1 MT may take place during the tensile deformation at room temperature of 25 ◦ C. From Fig. 6, the critical ˇ1 →
1 associated with ˇ1 → 1 MT is so high that starting stress s ˇ1 → 1 occurs difficultly before the as-cast alloy fractures. The critical stress t was determined as illustrated in the insert figure in Fig. 6 [16]. Therefore, the alloy is composed of a few 1 martensite but lots of 2 and ˇ1 phases after the tensile test, indicating that the stress-induced MT is difficult to proceed in the as-cast alloy, as shown in Fig. 7. From Fig. 6, the tensile stress-strain curve of the alloy treated at 700 ◦ C is similar to that of the as-cast alloy, which suggests the
ˇ1 →
1 difficulty with ˇ1 → 1 stress-induced MT owing to the high s ◦ of 450 MPa (Mf < 25 < Ms , as shown in Fig. 4). Therefore, the peak intensity of 1 martensite is still much lower than that of the 2 phase and ˇ1 phase in the alloy treated at 700 ◦ C (Fig. 7). With the HT temperature increasing from 730 ◦ C to 780 ◦ C, the tensile stress-strain curves of the Cu–Al–Ni alloy wires are composed of two stress-induced MT platforms, and the plasticity of the alloy is improved significantly with a decrease in tensile strength, as shown in Fig. 6. The critical stress t and the slope of the stressinduced MT platform decrease as well as the transformation strain εt increases gradually with the temperature increasing. Since the test temperature is higher than the Af temperature (25◦ < Af , as shown in Fig. 4), ˇ1 → ˇ1 → ˛1 MT may occur in the samples treated at 730–780 ◦ C. So the εt associated with ˇ1 → ˇ1 MT is about 7.8% at 780 ◦ C, which is close to the calculated elongation value of 8.5% of the [0 0 1] oriented single-crystalline alloy [17]. From Fig. 7, new peaks of ˇ1 and ˛1 martensites can be found in the XRD patterns and their peak intensity strengthens gradually with the HT temperature increasing, while the peak intensity of the 2 and ˇ1 phases weakens gradually. This indicates that the ˇ1 → ˇ1 MT occurs easily in the alloy due to a significant reduction in the
ˇ1 →ˇ 1 s , and the ˇ1 → ˛1 MT takes place following the accomplishment of ˇ1 → ˇ1 MT. There are two MT platforms formed in the tensile curves. In addition, the intensity of (0 0 1 8)ˇ1 and (0 0 4)˛1
diffraction peak weakens with increasing the HT temperature,
demonstrating that the ˇ1 and ˛1 martensites may undergo the reorientation during the MT proceeding [18]. Based on the above analysis, the heat-treated alloy has better plasticity, but lower tensile strength than that of the as-cast alloy. The change in tensile strength and elongation of the alloy with the heat treatment temperature is shown in Fig. 8. It can be seen that the tensile strengths of the alloy first a little increase to a maximum of 710 MPa at 730 ◦ C, and then decrease to a minimum of 577 MPa at 780 ◦ C. Meanwhile, the alloy displays a rapid increase in the elongation value from 7.5% to 23.8% with the temperature increasing from 700 ◦ C to 760 ◦ C, and then a slight reduction to 21.4% at 780 ◦ C. Due to the metastable character of the ˇ1 phase, the precipitation of 2 phase strongly influenced the characteristics of the martensitic transformation in Cu-based SMA by modifying the energy balance between the ˇ1 phase and martensites in different ways, that is, by creating new martensite nucleation points, changing the matrix composition or impeding the movement of the martensite–autensite or the martensite–martensite interface [19–22]. In this sense, when the elastic constant of the 2 phase is different from those of the matrix, an elastic interfacial force between obstacle and interface appears and the interfacial mobility is affected [23]. Therefore, the nano-hardness and elastic modulus of the 2 phase measured by micro-mechanical probe are listed in Table 2. One can see that the 2 phase in the as-cast alloy has the highest hardness and elastic modulus. With the HT temperature increasing from 700 ◦ C to 780 ◦ C, the hardness and elastic modulus of the 2 phase decrease first rapidly and then slightly. From the analysis of the amount evolution of the 2 phase, it can be found that the hardness and elastic modulus of the alloy decrease with the reduction of the amount of 2 phase, which is agreement with the conclusion of Ref. [22]. 4. Discussion The mechanical properties of the Cu–Al–Ni alloy before and after heat treated at 700–780 ◦ C are mainly determined by the evolution of the 2 phase, such as distribution, amount, shape and size. Fig. 1a shows that numerous coarsely dendritic 2 phase particles form in the as-cast alloy. Such microstructure generally induces very low plasticity, high tensile strength and high hard rate of
Table 2 Nano-hardness and elastic modulus of the 2 phase under different conditions. Sample
As-cast
700 ◦ C
730 ◦ C
750 ◦ C
760 ◦ C
770 ◦ C
780 ◦ C
Nano-hardness (GPa) Elastic modulus (GPa)
7.63 170.0
5.57 128.4
5.24 116.3
5.11 109.9
5.04 105.1
5.01 104.5
5.02 106.7
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Fig. 9. Fractured morphologies of the alloy wires: (a) as-cast, (b) 700 ◦ C, (c) 730 ◦ C, and (d) 760 ◦ C.
the sample, because the hard and brittle 2 phase acts as effective obstacles to the motion of the phased interfaces making the martensite transformation more difficult during the tensile deformation. In addition, the hardness and elastic modulus of 2 phase in the as-cast alloy is so high (Table 2) that the 2 phase is separated from the ˇ1 phase easily when the alloy is deformed. So micro-crack initiates at the 2 /ˇ1 interfaces due to the dimples expands quickly with increasing stress. This behavior gives rise to the intergranular fracture of the alloy, as shown in Fig. 9a. Therefore, the tensile strength of the as-cast alloy is up to 856 MPa, and its elongation to failure is only 5.2%. When the HT temperature is 700 ◦ C, the impediment effect for the movement of phase interface during the deformation is weakened due to the dissolution and size refinement of the 2 phase (Fig. 4a), which induces the reduction of the tensile strength [24,25]. In addition, the hardness and elastic modulus of the 2 phase in the sample are much lower than that in the as-cast alloy (Table 2), leading to an increase in the bonding force between 2 phase and ˇ1 phase. So, it is necessary to generate larger strain for forming crack. From Fig. 9b, there are some small and shallow dimples on the fracture morphology of the alloy, indicating that the alloy has a certain plastic deformation during the fracture. However, due to the low HT temperature and short holding time, only a small amount of the 2 phase is dissolved, and the 2 phase is inhomogeneous in size and shape. As a result, microcrack is still easily to be produced at the interfaces between ˇ1 phase and coarsely polygonal and long-banding 2 phase, which causes the intergranular fracture of the alloy. Therefore, compared with the as-cast alloy, the tensile strength of the alloy treated at 700 ◦ C decreases by 166 MPa, but the elongation only increases by 2.3%. When the alloy is heat treated at 730–750 ◦ C, a great reduction of the amount and size of the 2 phase due to the elevated temperature (Fig. 3b and c) leads to a remarkable reduction of the t of the alloy, which facilitates the ˇ1 → ˇ1 → ˛1 MT proceeding.
Such behavior favourably improves the plasticity of the alloy. Meanwhile, it can be seen that the tensile strength of the alloy treated at 730 ◦ C is a little higher than that of the alloy treated at 700 ◦ C, although the amount of the 2 phase decreases greatly from 50.0% to 29.2%. This can be ascribed to the dispersion strengthening effect of the fine 2 phase in the alloy. From Fig. 9c, the fracture morphology consists of many larger and deeper dimples, which generate, expand and connect to each other constantly until the alloy fractures. This is because the bonding force between the 2 /ˇ1 interfaces is further improved owing to the further decrease of the hardness and elastic modulus of the 2 phase at 730 ◦ C (Table 2). Although the transgranular cracks can be seen on the tensile fractured morphology, the intergranular cracks caused by the rupture of the discontinuous block 2 phase at the grain boundaries are still the main cause for the alloy fracture. Similar fracture morphology can also be obtained at 750 ◦ C. With the HT temperature is increased up to 760 ◦ C, the amount of the 2 phase goes on decreasing until full dissolution at 780 ◦ C, and the 2 phase distributed throughout the alloy are almost ellipsoidal and spherical (Fig. 3d–f). Since the amount of 2 phase is so small, the softening effect takes the leading role compared with the dispersion strengthening effect in the alloy. This induces that the tensile strength of the alloy decreases with the reduction of the 2 phase, and has a lowest value of 577 MPa at 780 ◦ C. In addition, the slope of the MT is descending gradually with the HT temperature increasing, which is useful for inducing the ˇ1 → ˇ1 → ˛1 MT at a lower applied stress. Such behavior favourably increases the elongation of the alloy. It can be seen from Fig. 9d that the fracture morphology of the alloy after treated at 760 ◦ C is composed of deep and well-distributed dimples, which is produced by the instantaneous transgranular toughness fracture. Therefore, the fracture mode of the alloy after heat treated at 760–780 ◦ C is dominated by the transgranular fracture, and the 2 phase was no longer regarded as the main cause for the alloy fracture. The fracture morphologies at 770–780 ◦ C are similar to that at 760 ◦ C.
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(2) The evolution of the 2 phase, including distribution, amount, shape and size, strongly affected the mechanical properties of the Cu–Al–Ni alloy. The tensile strengths of the heat-treated alloy increase in the beginning to a maximum of 710 MPa at 730 ◦ C, and then decrease to a minimum of 577 MPa at 780 ◦ C. The elongation of the alloy ranges from 7.5% to 23.8% and has a peak value at 760 ◦ C. Excellent balance between strength and plasticity can be obtained when the spherical 2 phase is distributed throughout the alloy with the size smaller than 5 m and the amount in the range of 11.4–16.6%. (3) The hardness and elastic modulus of the 2 phase decrease gradually with the reduction of its amount, which is benefit to obtain good mechanical properties of the Cu–Al–Ni alloy. Acknowledgments
Fig. 10. TEM Morphology of the martensite and 2 phase.
Fig. 10 shows the TEM morphology of the martensite and 2 phase obtained from the alloy treated at 770 ◦ C after the tensile test. It can be seen that the martensite variants are able to overcome the impediment of the fine 2 phase and pass through the 2 phase during the stress-induced MT process. A lot of parallel dislocation walls have been left in the martensite variants and blocked on both sides of the 2 phase. As a result, when the 2 phase is refined to a certain size, it will not have a harmful effect on the martensite growth and the plasticity of the alloy, which is agreement with the expectation of Sure [8]. In addition, the tensile strength of the alloy can be improved by a small amount of fine 2 phase distributed homogenously. Therefore, excellent balance between strength and plasticity can be obtained when the spherical 2 phase is distributed throughout the alloy with the size smaller than 5 m and the amount in the range of 11.4–16.6%. 5. Conclusions In the present study the following conclusions are drawn: (1) With the HT temperature increasing from 700 ◦ C to 770 ◦ C, the coarsely dendritic 2 phase evolves into the fine polygonal, ellipsoidal and spherical particles in the grains, while the long-banding, discontinuous block and ellipsoidal particles at the grain boundaries. The average size of the 2 phase decreases from 20 m before heat treatment down to 2 m, and its amount reduces from 50.0% down to 0.8%. The 2 phase is full dissolution at 780 ◦ C.
This work is supported by the Fundamental Research Funds for the Central Universities (FRF-TP-10-002B), the National Key Technology R&D Program (2011BAE23B02), the NSFC (50674008 and 51174027) and the 973 Program of China (2011CB606300). References [1] V. Recarte, J.I. Pérez-Landazábal, A. Ibarra, M.L. Nó, J. San Juan, Mater. Sci. Eng. A 378 (2004) 238–242. [2] A. Creuziger, W.C. Crone, Mater. Sci. Eng. A 498 (2008) 404–411. [3] Z.G. Wei, R. Sandstrom, S. Miyazaki, J. Mater. Sci. 33 (1998) 3743–3762. [4] A. Ohno, H. Sode, A. Mclean, Adv. Mater. 28 (1990) 161–168. [5] G. Motoyasu, M. Kaneko, H. Soda, A. Mclean, Metall. Mater. Trans. A 32 (2001) 585–593. [6] L.S. Cai, Y.Q. Yu, Z.H. Yin, D.F. Wang, W.G. Li, Heat Treat. Met. 31 (2006) 53–56 (in Chinese). [7] X.F. Liu, W.H. Li, J.X. Xie, Chin. J. Nonferrous Met. 18 (2008) 1248–1253 (in Chinese). [8] G.N. Sure, L.C. Brown, Metall. Trans. A 15 (1984) 1613–1621. [9] G. Lojen, I. Anzel, A. Kneissl, A. Krizman, E. Unterweger, B. Kosec, M. Bizjak, J. Mater. Process. Technol. 162–163 (2005) 220–229. [10] M.A. Dvorack, N. Kuwano, S. Polat, H. Chen, C.M. Wayman, Scripta Metall. 17 (1983) 1333–1336. [11] Z. Nishiyama, S. Kajiwara, Jpn. J. Appl. Phys. 2 (1963) 478–486. [12] J. Ye, M. Tokonami, K. Otsuka, Metall. Trans. A 21 (1990) 2669–2678. [13] C.P. Chen, C.Y.A. Tsao, J. Mater. Sci. 30 (1995) 4019–4026. [14] C.M. Wayman, Prog. Mater. Sci. 36 (1992) 203–224. [15] H. Sakamoto, K. Shimizu, K. Otsuka, Trans. Jpn. Inst. Met. 26 (1985) 638–645. [16] R.V. Krishnan, L. Delaey, H. Tas, H. Warlimont, J. Mater. Sci. 9 (1974) 1536–1544. [17] K. Otsuka, H. Sakamoto, K. Shimizu, Acta Metall. 27 (1979) 585–601. [18] M.S. Liu, Z. Zhang, W.L. Lu, Funct. Mater. 23 (1992) 27–32. [19] N. Kennon, D. Dunne, L. Middleton, Metall. Mater. Trans. A 13 (1982) 551–555. [20] P. Rodriguez, G. Guenin, Mater. Sci. Eng. A 129 (1990) 273–277. [21] N. Zárubova, A. Gemperle, V. Novák, Mater. Sci. Eng. A 222 (1997) 166–174. [22] V. Recarte, J.I. Pérez-Landazábal, V. Sánchez-Alarcos, Appl. Phys. Lett. 86 (2005) 231903–231913. [23] M. Grujicic, G.B. Olson, W.S. Owen, Metall. Trans. A 16A (1985) 1723–1734. [24] K.V. Sapozhnikov, V.V. Vetrov, S.A. Pulnev, S.B. Kustov, Scripta Mater. 34 (1996) 1543–1548. [25] M.I.A.E. Aal, Mater. Sci. Eng. A (2010), doi:10.1016/j.msea.2011.05.072.
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of 2 phase evolution on mechanical properties of continuous columnar-grained Cu–Al–Ni alloy Z. Wang, X.F. Liu, J.X. Xie ∗ Key Laboratory for Advanced Materials Processing (MOE), University of Science and Technology Beijing, Beijing 100083, China
a r t i c l e
i n f o
Article history: Received 21 July 2011 Received in revised form 2 November 2011 Accepted 3 November 2011 Available online 11 November 2011 Keywords: Shape memory alloy Continuous columnar grain 2 phase Heat treatment Martensitic transformation
a b s t r a c t Effect of 2 phase evolution on mechanical properties of Cu–14%Al–3.8%Ni (mass fraction) alloy wires fabricated by continuous unidirectional solidification technology was investigated before and after heat treated at 700–780 ◦ C. Mechanism for the improvement of mechanical properties of the alloy was analyzed. It was found that the alloy retained continuous columnar grains after heat treatment. With the heat treatment temperatures increasing from 700 ◦ C to 770 ◦ C, the coarse dendrite 2 phase evolved into the fine polygonal, ellipsoidal and spherical particles in the grains, while the long-banding, discontinuous block and ellipsoidal particles at the grain boundaries. The average size of the 2 phase decreased from 20 m before heat treatment down to 2 m, and its amount reduced from 50.0% down to 0.8%. The 2 phase was full dissolution at 780 ◦ C. The tensile strength of the alloy treated at 700–780 ◦ C ranged from 577 MPa to 710 MPa. The elongation of the alloy ranged from 7.5% to 23.8% and had a peak value at 760 ◦ C. Excellent balance between strength and elongation could be obtained when the spherical 2 phase was distributed throughout the alloy with the size smaller than 5 m and the amount in the range of 11.4–16.6%. The nano-hardness and elastic modulus of the 2 phase decreased gradually with the reduction of the amount, leading to the improvement of mechanical properties. © 2011 Elsevier B.V. All rights reserved.
1. Introduction In recent years, Cu–Al–Ni shape memory alloys (SMA) have attracted an increasing attention due to their low cost, excellent aging stability and thermal stability as well as higher operating temperature [1,2]. However, the engineering application of the alloy has been restricted seriously because equiaxed polycrystalline Cu–Al–Ni alloy is prone to intergranular fracture during plastic deformation due to the large grain size and the existence of brittle 2 phase [3]. Among the manufacturing methods available for improving mechanical properties of the alloy, the continuous unidirectional solidification technology is generally accepted as one of the most effective routes, which is based on the special competitive growth mechanism of crystal and can generate cast products with a beneficial texture [4]. Therefore, the highly oriented grains are deformed coherently under the axis load, which is helpful to avoid severe stress concentration at the grain boundaries. At present, columnar-grained Cu–Al–Ni alloy wires and pipes have been successfully fabricated by continuous unidirectional solidification technology [5–7].
∗ Corresponding author. Tel.: +86 10 62332254; fax: +86 10 62332253. E-mail address: [email protected] (J.X. Xie). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.11.019
In the study of grain refinement for Cu–Al–Ni alloy after hightemperature deformed, it was noted that the equiaxed grain growth during subsequent hot-working or heat treatment was retarded by the spherical 2 particles (about 5 m), which is distributed throughout the microstructure. Meanwhile, the fine 2 phase would not be expected to have a harmful effect on the mechanical properties of the alloy [8]. Nevertheless, when the solidification rate of the alloy was slow, namely the alloy had a poor cooling ability, numerous coarsely block and dendritic 2 phase (about 10–15 m) would be precipitated easily in the grains or at the grain boundaries of the alloy. This leads to the enlargement of the grain boundary embitterment and the degradation of mechanical properties [9]. It is demonstrated that the variation of the distribution, amount, size, and shape of the 2 phase obtained under different conditions plays distinct roles on mechanical properties of Cu–Al–Ni alloy. In order to accurately control the 2 phase and mechanical properties of the continuous columnar-grained Cu–Al–Ni alloy, it is necessary to study the evolution of the 2 phase and clarify the influence of this evolution on the mechanical properties of the alloy. This paper focuses on the evolution of the 2 phase and corresponding mechanical properties of continuous columnar-grained Cu–14%Al–3.8%Ni (mass fraction) alloy wires before and after heat treatment (HT). The mechanism for the improvement of mechanical properties of the alloy was analyzed, thus to provide basis for accuracy control and reasonable utilization of the 2 phase in the Cu–Al–Ni alloy.
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Fig. 1. Microstructures in longitudinal cross sections of the alloy wires, in which the black arrows show the columnar-grained boundaries: (a) as-cast, (b) 700 ◦ C, and (c) 780 ◦ C.
2. Experimental procedure Cu–14 wt.%Al–3.8 wt.%Ni alloy wires with a diameter of 6 mm and smooth surface were prepared from pure Cu, Al and Ni of 99.99% purity by using the vacuum melting and argon-shield vertical continuous unidirectional solidification equipment [7]. The investigated alloy was heated to a given temperature in a boxtype resistance furnace and then quenched into room temperature water. Since phase in the alloy is difficult to be adjusted at a low temperature and the alloy oxidizes easily at a high temperature, the HT temperatures were set at 700, 730, 750, 760, 770 and 780 ◦ C, respectively. In order to improve the efficiency of heat treatment and realize the on-line heat treatment in the future, the holding time was kept for 1 min. Phase composition of the alloy was determined by an Xray diffractometer (XRD) with an operated voltage of 40 kV, a current of 150 mA and a scanning speed of 10◦ /min. The spatial group, the lattice parameters and the atom positions ¯ a = 0.8720 nm, ˛ = ˇ = = 90◦ ), of the 2 phase (Cu9 Al4 , P 43m, ¯ ˇ1 phase (DO3 , Fm3m, a = 0.5822, ˛ = ˇ = = 90◦ ), ˇ1 martensite phase (18R1 , C2/m, a = 3.8190 nm, b = 0.5190 nm, c = 0.4490 nm, ˛ = = 90◦ , and ˇ = 89.7◦ ), 1 martensite (2H, Pmmn, a = 0.5342 nm, b = 0.4224 nm, and c = 0.4390 nm) and ˛1 martensite (DO22 , I4/mmm, a = 0.3590 nm, and c = 0.7550 nm) were determined referring to Refs. [10–12]. The martenstic transformation (MT) temperatures (Mf , Ms , As , Af ) of the alloy treated at different HT temperatures were measured using a differential scanning calorimeter (DSC) with a cooling rate and heating rate of 10 ◦ C/min. Optical microscopy (OM), transmission electron microscopy (TEM) and scanning electron microscopy (SEM) equipped with an energy dispersive spectrometer (EDS) were employed to characterize the microstructures, the fracture morphologies and the composition in the remaining ˇ1 phase of the alloy. Uniaxial tensile tests at room temperature were performed along the axial direction of the alloy
wire at a rate of 0.02 mm/s, using a MTS810 testing machine. The hardness and elastic modulus of the 2 phase under different conditions were determined by a Nano Indenter II micro-mechanical probe. 3. Results 3.1. Effect of heat treatment temperature on evolution of the 2 phase of Cu–Al–Ni alloy Fig. 1 shows the microstructures of the continuous columnargrained Cu–Al–Ni alloy wires before and after heat treatment. It can be seen that the alloys treated at 700 ◦ C and 780 ◦ C retain columnar grains, and the grain boundaries (marked by the black arrows in Fig. 1) are clear and straight. The average grain size of the heattreated alloy is about 440 m, which is identical with that of the as-cast alloy. This implies that the grain size has no effect on the mechanical properties of the alloy in this study. From Fig. 2, the ascast alloy consists of mainly ˇ1 phase and 2 phase. The 2 dendrite phase is in a dense distribution with the amount of 55.7%, and the average dendrite-arm span is about 20 m (Fig. 1a). Fig. 3 shows the morphologies of 2 phase in Cu–Al–Ni alloy wires at different HT temperatures. When the HT temperature is 700 ◦ C, one can see that most of the coarse 2 dendrite phase in the grains evolves into the fine ellipsoidal particles, usually in the size of 2–5 m, while a small amount of it is transformed into the polygonal structure with a size range of 5–9 m. At the grain boundaries, the dendritic 2 phase are accumulated to the long-banded structure and their longitudinal axis is in accordance with casting direction, which may cause significant anisotropic properties. With the HT temperature increasing from 730 ◦ C to 770 ◦ C, the polygonal 2 phase in the grains is further changed into the ellipsoidal and spherical particles, and the long-banded ones at the grain boundaries evolve into the discontinuous block and
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Fig. 2. XRD patterns of the alloy wires.
ellipsoidal particles gradually. The average size of the ellipsoidal and spherical 2 phase is about 4 m at 760 ◦ C, and that of the spherical 2 phase is about 2 m at 770 ◦ C. As a result, the 2 phase becomes finer and more uniform in size and shape with the HT temperature increasing. Moreover, the amount of 2 phase decreases continuously with the HT temperature increasing, which is about 50.0% at 700 ◦ C, 29.2% at 730 ◦ C, 16.6% at 750 ◦ C, 11.4% at 760 ◦ C, 0.8% at 770 ◦ C, and full dissolution at 780 ◦ C, respectively, as shown in Fig. 3. Such behavior can also be demonstrated in Fig. 2 that the peak intensity of the 2 phase decreases, but that of the ˇ1 phase increases correspondingly with an increase of the HT temperature. At 780 ◦ C, the 2 phase cannot be detected in the XRD examinations, so the alloy is composed of only ˇ1 phase. The effect of heat treatment on the morphology evolution of 2 phase is essential the result of the breaking-up of the dendritic 2 structure at valleys due to spheroidizing and Ostwald ripening during heating and holding at elevated temperature. The dendritic structure firstly transforms to an ellipsoidal structure, and
Fig. 3. Morphologies of 2 phase in longitudinal cross sections of the alloy wires treated at different heat treatment temperatures: (a) 700 ◦ C, (b) 730 ◦ C, (c) 750 ◦ C, (d) 760 ◦ C, (e) 770 ◦ C, and (f) 780 ◦ C.
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Table 1 Aluminum concentration in the remaining ˇ1 phase under different conditions. Sample
As-cast
700 ◦ C
730 ◦ C
750 ◦ C
760 ◦ C
770 ◦ C
780 ◦ C
Al% (mass fraction)
12.79
13.08
13.41
13.78
13.96
14.07
14.1
subsequently spontaneous contracts to a spherical structure for the surface tension effect [13]. In addition, a higher temperature leads to a faster diffusion of the alloying element and breaking-up process. Therefore, the 2 phase is more uniform and finer in size and shape as well as fewer in amount at a higher temperature than that at a lower temperature.
The tensile stress-stain curves of the Cu–Al–Ni alloy wires under various conditions are shown in Fig. 6. It can be seen that the stressstain curve of the as-cast alloy is steep with high strain hardening rate and shows a typical characteristic of brittle fracture without
(a) 3.2. Effect of 2 phase evolution on transformation temperatures of Cu–Al–Ni alloy
3.3. Effect of 2 phase evolution on mechanical properties of Cu–Al–Ni alloy It is well known that distinct type of the MT may be stressinduced depending upon test temperature, when the Cu–Al–Ni alloy is composed of mainly ˇ1 phase [14,15]. In the temperature range between Mf and Ms , three distinct stress-induced MT, ˇ1 → 1 , 1 → ˇ1 and ˇ1 → ˛1 , were shown on the stress-strain curves. On the other hand, in the temperature range a little higher than Af , two distinct stress-induced MT, ˇ1 → ˇ1 and ˇ1 → ˛1 , may occur successively.
Heat flow (mW/mg)
730°C 750°C 760°C 770°C 780°C -100
(b)
Heat flow (mW/mg)
Fig. 4 shows the evolution of the characteristic MT temperatures with HT temperatures. It can be seen that the transformation temperatures first decrease rapidly with the HT temperature increasing from 700 ◦ C to 730 ◦ C and Mf < As < Ms < Af . Then they decrease slightly when the HT temperature is up to 750 ◦ C and Mf < Ms < As < Af . This behavior can be ascribed to the dissolution of 2 phase increasing the aluminum concentration in the remaining ˇ1 phase, as shown in Table 1, hence decreasing the MT temperature in the alloy. The more the aluminum concentration increases, the faster the MT temperature decrease. DSC data obtained around the martensitic phase transformation are plotted in Fig. 5. It can be found that the peaks during the heating or cooling process are very low and wide at 700 ◦ C, but become sharp and narrow with increasing the HT temperatures. Since the amount of the 2 phase decreases continuously with increasing temperature (Fig. 3), it can be indicated that the shape of the MT cycle becomes sharp and narrow with the reduction of the amount of the 2 phase.
Heating
700°C
-75
-50
-25
0
25
50
75
100
25
50
75
100
Temperature (ºC)
Cooling
700°C 730°C 750°C 760°C 770°C 780°C
-100
-75
-50
-25
0
Temperature (ºC) Fig. 5. DSC thermograms of the MT cycle for different heat treatments: (a) the heating process; (b) the cooling process.
900
100
Ms As Mf
50 0
700
RT=25°C
-50 -100
σt
800
Af
600
Stress (MPa)
Transformation temperature (ºC)
150
500
300 200 εt
100
700
720
740
Temperature (ºC)
760
780
Fig. 4. Transformation temperatures of the alloy wires treated at different heat treatment temperatures.
as cast 700°C × 1min 730°C × 1min 750°C × 1min 760°C × 1min 770°C × 1min 780°C × 1min
400
0
0
5
10
15
20
Strain (%) Fig. 6. Tensile stress-stain curves of the alloy wires.
25
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Fig. 8. Effect of heat treatment temperature on mechanical properties of the alloy wires.
Fig. 7. XRD patterns of the alloy wires after tensile fracture.
remarkable yield limit. Since the Mf , Ms , As , Af temperatures of the as-cast alloy are 17.5 ◦ C, 30.0 ◦ C, 62.8 ◦ C and 99.4 ◦ C, respectively, ˇ1 → 1 → ˇ1 → ˛1 MT may take place during the tensile deformation at room temperature of 25 ◦ C. From Fig. 6, the critical ˇ1 →
1 associated with ˇ1 → 1 MT is so high that starting stress s ˇ1 → 1 occurs difficultly before the as-cast alloy fractures. The critical stress t was determined as illustrated in the insert figure in Fig. 6 [16]. Therefore, the alloy is composed of a few 1 martensite but lots of 2 and ˇ1 phases after the tensile test, indicating that the stress-induced MT is difficult to proceed in the as-cast alloy, as shown in Fig. 7. From Fig. 6, the tensile stress-strain curve of the alloy treated at 700 ◦ C is similar to that of the as-cast alloy, which suggests the
ˇ1 →
1 difficulty with ˇ1 → 1 stress-induced MT owing to the high s ◦ of 450 MPa (Mf < 25 < Ms , as shown in Fig. 4). Therefore, the peak intensity of 1 martensite is still much lower than that of the 2 phase and ˇ1 phase in the alloy treated at 700 ◦ C (Fig. 7). With the HT temperature increasing from 730 ◦ C to 780 ◦ C, the tensile stress-strain curves of the Cu–Al–Ni alloy wires are composed of two stress-induced MT platforms, and the plasticity of the alloy is improved significantly with a decrease in tensile strength, as shown in Fig. 6. The critical stress t and the slope of the stressinduced MT platform decrease as well as the transformation strain εt increases gradually with the temperature increasing. Since the test temperature is higher than the Af temperature (25◦ < Af , as shown in Fig. 4), ˇ1 → ˇ1 → ˛1 MT may occur in the samples treated at 730–780 ◦ C. So the εt associated with ˇ1 → ˇ1 MT is about 7.8% at 780 ◦ C, which is close to the calculated elongation value of 8.5% of the [0 0 1] oriented single-crystalline alloy [17]. From Fig. 7, new peaks of ˇ1 and ˛1 martensites can be found in the XRD patterns and their peak intensity strengthens gradually with the HT temperature increasing, while the peak intensity of the 2 and ˇ1 phases weakens gradually. This indicates that the ˇ1 → ˇ1 MT occurs easily in the alloy due to a significant reduction in the
ˇ1 →ˇ 1 s , and the ˇ1 → ˛1 MT takes place following the accomplishment of ˇ1 → ˇ1 MT. There are two MT platforms formed in the tensile curves. In addition, the intensity of (0 0 1 8)ˇ1 and (0 0 4)˛1
diffraction peak weakens with increasing the HT temperature,
demonstrating that the ˇ1 and ˛1 martensites may undergo the reorientation during the MT proceeding [18]. Based on the above analysis, the heat-treated alloy has better plasticity, but lower tensile strength than that of the as-cast alloy. The change in tensile strength and elongation of the alloy with the heat treatment temperature is shown in Fig. 8. It can be seen that the tensile strengths of the alloy first a little increase to a maximum of 710 MPa at 730 ◦ C, and then decrease to a minimum of 577 MPa at 780 ◦ C. Meanwhile, the alloy displays a rapid increase in the elongation value from 7.5% to 23.8% with the temperature increasing from 700 ◦ C to 760 ◦ C, and then a slight reduction to 21.4% at 780 ◦ C. Due to the metastable character of the ˇ1 phase, the precipitation of 2 phase strongly influenced the characteristics of the martensitic transformation in Cu-based SMA by modifying the energy balance between the ˇ1 phase and martensites in different ways, that is, by creating new martensite nucleation points, changing the matrix composition or impeding the movement of the martensite–autensite or the martensite–martensite interface [19–22]. In this sense, when the elastic constant of the 2 phase is different from those of the matrix, an elastic interfacial force between obstacle and interface appears and the interfacial mobility is affected [23]. Therefore, the nano-hardness and elastic modulus of the 2 phase measured by micro-mechanical probe are listed in Table 2. One can see that the 2 phase in the as-cast alloy has the highest hardness and elastic modulus. With the HT temperature increasing from 700 ◦ C to 780 ◦ C, the hardness and elastic modulus of the 2 phase decrease first rapidly and then slightly. From the analysis of the amount evolution of the 2 phase, it can be found that the hardness and elastic modulus of the alloy decrease with the reduction of the amount of 2 phase, which is agreement with the conclusion of Ref. [22]. 4. Discussion The mechanical properties of the Cu–Al–Ni alloy before and after heat treated at 700–780 ◦ C are mainly determined by the evolution of the 2 phase, such as distribution, amount, shape and size. Fig. 1a shows that numerous coarsely dendritic 2 phase particles form in the as-cast alloy. Such microstructure generally induces very low plasticity, high tensile strength and high hard rate of
Table 2 Nano-hardness and elastic modulus of the 2 phase under different conditions. Sample
As-cast
700 ◦ C
730 ◦ C
750 ◦ C
760 ◦ C
770 ◦ C
780 ◦ C
Nano-hardness (GPa) Elastic modulus (GPa)
7.63 170.0
5.57 128.4
5.24 116.3
5.11 109.9
5.04 105.1
5.01 104.5
5.02 106.7
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Fig. 9. Fractured morphologies of the alloy wires: (a) as-cast, (b) 700 ◦ C, (c) 730 ◦ C, and (d) 760 ◦ C.
the sample, because the hard and brittle 2 phase acts as effective obstacles to the motion of the phased interfaces making the martensite transformation more difficult during the tensile deformation. In addition, the hardness and elastic modulus of 2 phase in the as-cast alloy is so high (Table 2) that the 2 phase is separated from the ˇ1 phase easily when the alloy is deformed. So micro-crack initiates at the 2 /ˇ1 interfaces due to the dimples expands quickly with increasing stress. This behavior gives rise to the intergranular fracture of the alloy, as shown in Fig. 9a. Therefore, the tensile strength of the as-cast alloy is up to 856 MPa, and its elongation to failure is only 5.2%. When the HT temperature is 700 ◦ C, the impediment effect for the movement of phase interface during the deformation is weakened due to the dissolution and size refinement of the 2 phase (Fig. 4a), which induces the reduction of the tensile strength [24,25]. In addition, the hardness and elastic modulus of the 2 phase in the sample are much lower than that in the as-cast alloy (Table 2), leading to an increase in the bonding force between 2 phase and ˇ1 phase. So, it is necessary to generate larger strain for forming crack. From Fig. 9b, there are some small and shallow dimples on the fracture morphology of the alloy, indicating that the alloy has a certain plastic deformation during the fracture. However, due to the low HT temperature and short holding time, only a small amount of the 2 phase is dissolved, and the 2 phase is inhomogeneous in size and shape. As a result, microcrack is still easily to be produced at the interfaces between ˇ1 phase and coarsely polygonal and long-banding 2 phase, which causes the intergranular fracture of the alloy. Therefore, compared with the as-cast alloy, the tensile strength of the alloy treated at 700 ◦ C decreases by 166 MPa, but the elongation only increases by 2.3%. When the alloy is heat treated at 730–750 ◦ C, a great reduction of the amount and size of the 2 phase due to the elevated temperature (Fig. 3b and c) leads to a remarkable reduction of the t of the alloy, which facilitates the ˇ1 → ˇ1 → ˛1 MT proceeding.
Such behavior favourably improves the plasticity of the alloy. Meanwhile, it can be seen that the tensile strength of the alloy treated at 730 ◦ C is a little higher than that of the alloy treated at 700 ◦ C, although the amount of the 2 phase decreases greatly from 50.0% to 29.2%. This can be ascribed to the dispersion strengthening effect of the fine 2 phase in the alloy. From Fig. 9c, the fracture morphology consists of many larger and deeper dimples, which generate, expand and connect to each other constantly until the alloy fractures. This is because the bonding force between the 2 /ˇ1 interfaces is further improved owing to the further decrease of the hardness and elastic modulus of the 2 phase at 730 ◦ C (Table 2). Although the transgranular cracks can be seen on the tensile fractured morphology, the intergranular cracks caused by the rupture of the discontinuous block 2 phase at the grain boundaries are still the main cause for the alloy fracture. Similar fracture morphology can also be obtained at 750 ◦ C. With the HT temperature is increased up to 760 ◦ C, the amount of the 2 phase goes on decreasing until full dissolution at 780 ◦ C, and the 2 phase distributed throughout the alloy are almost ellipsoidal and spherical (Fig. 3d–f). Since the amount of 2 phase is so small, the softening effect takes the leading role compared with the dispersion strengthening effect in the alloy. This induces that the tensile strength of the alloy decreases with the reduction of the 2 phase, and has a lowest value of 577 MPa at 780 ◦ C. In addition, the slope of the MT is descending gradually with the HT temperature increasing, which is useful for inducing the ˇ1 → ˇ1 → ˛1 MT at a lower applied stress. Such behavior favourably increases the elongation of the alloy. It can be seen from Fig. 9d that the fracture morphology of the alloy after treated at 760 ◦ C is composed of deep and well-distributed dimples, which is produced by the instantaneous transgranular toughness fracture. Therefore, the fracture mode of the alloy after heat treated at 760–780 ◦ C is dominated by the transgranular fracture, and the 2 phase was no longer regarded as the main cause for the alloy fracture. The fracture morphologies at 770–780 ◦ C are similar to that at 760 ◦ C.
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(2) The evolution of the 2 phase, including distribution, amount, shape and size, strongly affected the mechanical properties of the Cu–Al–Ni alloy. The tensile strengths of the heat-treated alloy increase in the beginning to a maximum of 710 MPa at 730 ◦ C, and then decrease to a minimum of 577 MPa at 780 ◦ C. The elongation of the alloy ranges from 7.5% to 23.8% and has a peak value at 760 ◦ C. Excellent balance between strength and plasticity can be obtained when the spherical 2 phase is distributed throughout the alloy with the size smaller than 5 m and the amount in the range of 11.4–16.6%. (3) The hardness and elastic modulus of the 2 phase decrease gradually with the reduction of its amount, which is benefit to obtain good mechanical properties of the Cu–Al–Ni alloy. Acknowledgments
Fig. 10. TEM Morphology of the martensite and 2 phase.
Fig. 10 shows the TEM morphology of the martensite and 2 phase obtained from the alloy treated at 770 ◦ C after the tensile test. It can be seen that the martensite variants are able to overcome the impediment of the fine 2 phase and pass through the 2 phase during the stress-induced MT process. A lot of parallel dislocation walls have been left in the martensite variants and blocked on both sides of the 2 phase. As a result, when the 2 phase is refined to a certain size, it will not have a harmful effect on the martensite growth and the plasticity of the alloy, which is agreement with the expectation of Sure [8]. In addition, the tensile strength of the alloy can be improved by a small amount of fine 2 phase distributed homogenously. Therefore, excellent balance between strength and plasticity can be obtained when the spherical 2 phase is distributed throughout the alloy with the size smaller than 5 m and the amount in the range of 11.4–16.6%. 5. Conclusions In the present study the following conclusions are drawn: (1) With the HT temperature increasing from 700 ◦ C to 770 ◦ C, the coarsely dendritic 2 phase evolves into the fine polygonal, ellipsoidal and spherical particles in the grains, while the long-banding, discontinuous block and ellipsoidal particles at the grain boundaries. The average size of the 2 phase decreases from 20 m before heat treatment down to 2 m, and its amount reduces from 50.0% down to 0.8%. The 2 phase is full dissolution at 780 ◦ C.
This work is supported by the Fundamental Research Funds for the Central Universities (FRF-TP-10-002B), the National Key Technology R&D Program (2011BAE23B02), the NSFC (50674008 and 51174027) and the 973 Program of China (2011CB606300). References [1] V. Recarte, J.I. Pérez-Landazábal, A. Ibarra, M.L. Nó, J. San Juan, Mater. Sci. Eng. A 378 (2004) 238–242. [2] A. Creuziger, W.C. Crone, Mater. Sci. Eng. A 498 (2008) 404–411. [3] Z.G. Wei, R. Sandstrom, S. Miyazaki, J. Mater. Sci. 33 (1998) 3743–3762. [4] A. Ohno, H. Sode, A. Mclean, Adv. Mater. 28 (1990) 161–168. [5] G. Motoyasu, M. Kaneko, H. Soda, A. Mclean, Metall. Mater. Trans. A 32 (2001) 585–593. [6] L.S. Cai, Y.Q. Yu, Z.H. Yin, D.F. Wang, W.G. Li, Heat Treat. Met. 31 (2006) 53–56 (in Chinese). [7] X.F. Liu, W.H. Li, J.X. Xie, Chin. J. Nonferrous Met. 18 (2008) 1248–1253 (in Chinese). [8] G.N. Sure, L.C. Brown, Metall. Trans. A 15 (1984) 1613–1621. [9] G. Lojen, I. Anzel, A. Kneissl, A. Krizman, E. Unterweger, B. Kosec, M. Bizjak, J. Mater. Process. Technol. 162–163 (2005) 220–229. [10] M.A. Dvorack, N. Kuwano, S. Polat, H. Chen, C.M. Wayman, Scripta Metall. 17 (1983) 1333–1336. [11] Z. Nishiyama, S. Kajiwara, Jpn. J. Appl. Phys. 2 (1963) 478–486. [12] J. Ye, M. Tokonami, K. Otsuka, Metall. Trans. A 21 (1990) 2669–2678. [13] C.P. Chen, C.Y.A. Tsao, J. Mater. Sci. 30 (1995) 4019–4026. [14] C.M. Wayman, Prog. Mater. Sci. 36 (1992) 203–224. [15] H. Sakamoto, K. Shimizu, K. Otsuka, Trans. Jpn. Inst. Met. 26 (1985) 638–645. [16] R.V. Krishnan, L. Delaey, H. Tas, H. Warlimont, J. Mater. Sci. 9 (1974) 1536–1544. [17] K. Otsuka, H. Sakamoto, K. Shimizu, Acta Metall. 27 (1979) 585–601. [18] M.S. Liu, Z. Zhang, W.L. Lu, Funct. Mater. 23 (1992) 27–32. [19] N. Kennon, D. Dunne, L. Middleton, Metall. Mater. Trans. A 13 (1982) 551–555. [20] P. Rodriguez, G. Guenin, Mater. Sci. Eng. A 129 (1990) 273–277. [21] N. Zárubova, A. Gemperle, V. Novák, Mater. Sci. Eng. A 222 (1997) 166–174. [22] V. Recarte, J.I. Pérez-Landazábal, V. Sánchez-Alarcos, Appl. Phys. Lett. 86 (2005) 231903–231913. [23] M. Grujicic, G.B. Olson, W.S. Owen, Metall. Trans. A 16A (1985) 1723–1734. [24] K.V. Sapozhnikov, V.V. Vetrov, S.A. Pulnev, S.B. Kustov, Scripta Mater. 34 (1996) 1543–1548. [25] M.I.A.E. Aal, Mater. Sci. Eng. A (2010), doi:10.1016/j.msea.2011.05.072.