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Process optimization, microstructures and mechanical properties of a Cu-based shape memory alloy fabricated by selective laser melting
Journal of Alloys and Compounds 785 (2019) 754e764
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Process optimization, microstructures and mechanical properties of a Cu-based shape memory alloy fabricated by selective laser melting Jian Tian, Wenzhi Zhu**, Qingsong Wei*, Shifeng Wen, Shuai Li, Bo Song, Yusheng Shi State Key Laboratory of Materials Processing and Die & Mould Technology, School of Materials Science and Engineering, Huazhong University of Science and Technology, Wuhan 430074, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 23 April 2018 Received in revised form 3 December 2018 Accepted 12 January 2019 Available online 14 January 2019
The Cu-13.5Al-4Ni-0.5Ti copper-based shape memory alloys (SMAs) were fabricated by selective laser melting (SLM). The parameters were optimized to obtain almost fully dense copper-based SMAs samples. The phases and microstructures were characterized and the tensile properties at room temperature and 200 C were evaluated. The XRD results showed that only the b10 phase could be detected in the Cu-based SMAs, which was attributed to the extremely short solidification time for the precipitation of a and g2 phases during SLM process. The grains were well refined and the average grain size was approximate 43 mm, which was much smaller than that of the cast copper-based SMAs. The X-phase (Cu2TiAl) is observed, which is granular and size 20e50 nm. The Cu-13.5Al-4Ni-0.5Ti copper-based SMAs fabricated by SLM exhibited excellent mechanical properties at room temperature, which was higher than that of the cast copper-based SMAs. This is attributed to two aspects: (1) grain refinement, (2) suppress of brittle g2 phase. Remarkably, the tensile test results at 200 C showed both higher strength and elongation, which is attributed to the bcc structure of the parent phase and the stress-induced martensitic transformation at 200 C. Meanwhile, the difference of microstructures and properties between SLMfabricated Cu-based SMAs and casting Cu-based SMAs were analyzed and discussed in detail. © 2019 Elsevier B.V. All rights reserved.
Keywords: Copper-based shape memory alloy Selective laser melting Process optimization Relative density Microstructure Mechanical properties
1. Introduction Shape memory alloys (SMAs) are recently developed as advanced materials due to their specific properties such as superelasticity, shape memory effect (SME), biocompatibility, high specific strength, excellent corrosion resistance, good wear resistance and high anti-fatigue property [1]. Among them, Ni-Ti and Cubased alloys are two types of the most widely used SMAs [2]. Although Ni-Ti SMAs show better shape memory property and super-elasticity than do Cu-based SMAs [3], they are only used at a low operating temperature not exceeding 100 C [4], and they have poor machinability [5]. However, the substitution of Ni with noble metal elements such as Au, Pt, and Pd or the substitution of Ti with Zr and Hf can improve the transformation temperature of Ni-Ti SMAs [6e8]. And this type of Ni-Ti-based high-temperature shape memory alloys (HT-SMAs) is currently being studied extensively. In contrast, Cu-based SMAs can reach higher transformation
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (W. Zhu), [email protected] (Q. Wei). https://doi.org/10.1016/j.jallcom.2019.01.153 0925-8388/© 2019 Elsevier B.V. All rights reserved.
temperatures and are preferential candidates for high-temperature applications (such as in thermal actuators and thermal sensors) up to 240 C [9] due to their super-elastic or pseudo-elastic properties, two-way shape memory effect and high damping ability [10]. Moreover, the excellent processing performance of Cu-based SMAs enables the manufacture of complex-shaped parts [4]. Cu-based SMAs are currently divided into two series, namely, Cu-Al-Ni and Cu-Zn-Al [11], whose martensite transformation temperatures can be adjusted within a wide range [12]. Since the transformation temperature of Cu-based SMAs is sensitive to the variation of composition [3], it can be tailored by adjusting the ratios of elements in the alloy to satisfy various temperature requirements in different application environments. More significantly, Cu-Al-Ni exhibits a higher martensite transformation temperature than do Cu-Zn-Al alloys [13]. For example, Sugimoto et al. [14] developed a Cu-Al-Ni-Mn alloy with good thermodynamic stability and anti-ageing behaviour at temperatures above 100 C, and this alloy showed the capability for high-temperature applications. However, Cu-Al-Ni SMAs are brittle in nature due to the intergranular cracking, which is caused by the high elastic anisotropy
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[15]. For example, the elastic anisotropy factors of the Cu-Al-Ni, CuZn-Al, and Ni-Ti alloys are 13 [16], 15 [17], and 2 [18], respectively. The large elastic anisotropy makes the intergranular elasticity and plastic deformation of the copper-based SMAs extremely uncoordinated, which causes the grain boundary stress concentration and intergranular cracking. Currently, grain refinement and plasticity enhancement are two main methods to improve the mechanical properties of Cu-based SMAs. Researchers found that the addition of B, Ce and V elements could result in the grain refinement of CuZn-Al alloys [19], whereas the addition of Ti and Zr elements similarly affected Cu-Al-Ni alloys [20,21]. Alternative routes to increase the performance of Cu-Al-Ni alloys are rapid solidification methods such as melt-spinning [22], melt extraction [23], spray forming [24] and selective laser melting [25]. As an additive manufacturing (AM) technique, selective laser melting (SLM) can create bulk parts by melting specific predefined areas of a powder bed layer by layer [26]. The processing of a thin powder layer (20e50 mm) on massive substrate plates in combination with small laser spot diameters (approximately 100 mm) results in high intrinsic cooling rates of 2.13e2.97 106 C/s [27]. Fine grain sizes within approximately one micron can be obtained because of the high cooling rates during SLM. If the parts prepared by SLM have a high relative density, the yield and ultimate tensile strength are comparable to (even better than) those of conventional manufactured samples, whereas the fracture strain remains lower [26]. Recently, several studies investigated the fabrication and characterization of SLM-fabricated Ni-Ti alloys. Fabricated parts using selective laser melting (SLM) technology show low impurity [28], microstructure control [29], high temperature properties [30] and satisfy the requirements in the ASTM F2063-05 for medical NiTi devices [31]. However, those of copper-based SMAs are notably limited. Gustmann et al. [32,33] studied the effect of SLM processing parameters on the microstructure and mechanical properties of Cu-11.85Al-3.2Ni-3Mn SMAs and obtained nearly dense samples with higher mechanical properties than those of the casting counterparts. The SLM-fabricated Cu-11.85Al-3.2Ni-3Mn alloys have an average ultimate tensile strength of 620 ± 50 MPa and elongation of 8.2 ± 0.9% at room temperature (RT). The average hardness is 245 ± 20 HV. However, the mechanical properties of Cubased SMAs must be improved, and there is a lack of hightemperature performance studies. This study examines the SLM of a new Cu-based SMA, Cu13.5Al-4Ni-0.5Ti. According to the phase diagram of Cu-Al-Ni ternary alloy, the alloy with 4 wt % Ni and 13.5 wt % Al exhibits a good shape memory effect. The addition of Ti element is expected to improve the material strength because of the grain refinement effect [20]. The objectives are to obtain the optimal SLM processing parameters to fabricate fully dense samples (r 99%), clarify the phase formation, microstructures and tensile properties at RT and 200 C, and understand their performance characteristic and difference from casting alloys. The findings provide a basic insight on SLM-fabricating high-performance Cu-based SMAs with tailored microstructure and elevated mechanical properties. 2. Experimental 2.1. Materials Gas-atomized Cue13.5Ale4Nie0.5Ti powder was produced by using high-purity elements (>99.9 wt %) in an induction furnace in argon atmosphere. The density (7.11 g/cm3) of the produced alloy was determined by a fully automatic true density analyser (AccuPyc 1330, USA) as the reference value to calculate the relative densities of the SLM-fabricated samples. The morphology of the powder was observed by scanning electron microscopy (SEM JSM-7600F, Japan),
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as shown in Fig. 1(a). The powder particle shows a nearly spherical shape. Small satellite particles were observed to adhere to the surface of large particles. The powder size distribution was measured by a laser particle size analyser (Mastersizer 3000, UK), as illustrated in Fig. 1(b). The powder size was 19.8e46.7 mm with an average value of 30.5 mm. 2.2. Selective laser melting and samples preparation The samples were fabricated by an SLM 125HL machine (Solutions GmbH, Germany). The SLM machine was equipped with a 400-W Gaussian beam fibre laser. The laser beam diameter (F90%) is approximately 80 mm. The process was performed in an argon atmosphere (approximately 500 ppm oxygen). Based on previous extensive SLM forming experiments, the laser power and scanning speed have greater effect on the formed parts than the layer thickness and scan spacing. Therefore, the laser power and scanning speed were chosen as the experimental variables of SLM process optimization experiments. Some cube samples (in 8 8 5 mm3) for process optimization were fabricated by SLM. The processing parameters with fixed layer thickness (0.04 mm), fixed scanning spacing (0.09 mm), varying laser power (250e340 W) and varying scanning speed (600e900 mm/s) were chosen. The long bidirectional scanning strategy was used, as shown in Fig. 2. To reduce the residual thermal stress during SLM, a 67 rotation of the scanning direction between two consecutive layers N and N þ 1 was set [34]. The substrate was preheated to 200 C before processing. Fig. 3(a) illustrates the Cu-13.5Al-4Ni0.5Ti alloy cubes (8 8 5 mm3) fabricated by SLM using various processing parameters. In addition, the hardness testing samples (cubes of 8 8 5 mm3) and tensile samples were fabricated with the optimized processing parameter (maximum relative density), as shown in Fig. 3. 2.3. Materials characterization The density of the cube samples was measured using a highprecision balance (Sartorius MC210P, Germany), whose measurement principle is Archimedes method. The phase identification was performed by XRD (XRD-7000S, Shimadzu, Japan) with a Cu target in a conventional X-ray tube, which operated at an accelerating voltage of 40 kV and an electron beam current of 30 mA. The SLMfabricated samples for microstructure observations were cut from the substrate by wire-electrode cutting, ground and polished with automatic grinding and polishing machine (Ecomet300/Automet300, Buehler, USA), and chemically etched with a solution (vol. 50% H2O and vol. 50% HNO3) at room temperature for approximately 10 seconds. The microstructures were investigated by optical microscopy (OM VHX-1000C, Keyence, Japan), scanning electron microscopy (SEM JSM-7600F, Japan) and transmission electron microscopy (TEM JEM-2100, Japan). The Vickers microhardness of the samples was determined by a hardness test machine (430-SVD, Wilson Hardness, USA) with a 5kg load and 30-s dwell time. Five points were chosen and measured at different positions of each surface to obtain an average value. The tensile samples were cut to the required dimension by a CNC machine (WDW3200, China), as shown in Fig. 3(b). Because the working environment of Cu-Al-Ni alloys at high temperature is mainly at 100e200 C [35], the tensile tests were performed on a mechanical testing machine (Zwick/Roell Z020, Zwick, Germany) at RT and 200 C. The tests were performed using a cross speed of 0.5 mm/min and a constant tension load. The strain was recorded by a laser extensometer. The fracture morphology of the tensile samples was observed by scanning electron microscopy (SEM JSM7600F, Japan).
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Fig. 1. Cu-13.5Al-4Ni-0.5Ti powder: (a) The micromorphology under SEM, (b) Powder particle size distribution.
A), the density substantially decreases with the increase in laser power but increases with the increase in scanning speed. The combination of the processing parameter, a high laser power and a low scanning speed easily makes the metal solution boil in molten pools due to an extremely high temperature, which causes the splashing of small metal droplets and finally the formation of small pores in the samples and a lower density [36]. When the scanning speed is improved to 800e900 mm/s (Fig. 4(b), part B), the relative density exceeds 99%. In this case, the density only slightly varies with the laser power and scanning speed. Thus, the laser energy is sufficient and appropriate to completely melt powder to form a continuous molten pool. The input energy density is usually used to evaluate the effects of the processing parameters on the density of SLM-fabricated samples, which can be described as [37]:
E¼
Fig. 2. Schematic of laser scanning strategies during SLM.
3. Results and discussion 3.1. Effects of the SLM processing parameters on the relative density Fig. 4 illustrates the relative densities of Cue13.5Ale4Nie0.5Ti SMAs fabricated by SLM with various laser powers and scanning speeds. When the scanning speed is 600e700 mm/s (Fig. 4(a), part
P HDV
(1)
where E is the input energy density (in J/mm3), P is the laser power (in W), V is the scanning speed (in mm/s), H is the scanning space (in mm), and D is the layer thickness (in mm). Fig. 5 illustrates the relationship between the input energy density and the relative density of SLM-fabricated samples. When the input energy density is 77e110 J/mm3, the relative density exceeds 99% and slightly increases with the increase in energy density. However, the relative density significantly decreases with the increase in energy density when the input energy density is over 110 J/mm3. Fig. 6 displays the morphology in SLM-fabricated samples at three input laser energy densities values of points A, B and C in
Fig. 3. The Cu-13.5Al-4Ni-0.5Ti samples fabricated by SLM: (a) Cubes (8 8 5 mm3) fabricated by various processing parameters (layer thickness: 0.04 mm, scanning space: 0.09 mm); (b) Tensile samples (laser power: 310 W, scanning speed: 800 mm/s, layer thickness: 0.04 mm, scanning space: 0.09 mm).
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Fig. 4. (a) The relationship between the laser power and the scanning speed and the sample density; (b) Top view of Fig. 4a: Part A is the density value of the sample at the lower scanning speed (600e700 mm/s); Part B is the density value of the sample at the higher scanning speed (800e900 mm/s).
many pores to form in the SLM-fabricated sample, which decreases the density of the sample. Simultaneously, spheroidization will increase the molten pool surface roughness, affect the powder spreading and follow-up melt spread and accumulation, and reduce the relative density. Based on the density analysis of the test samples fabricated with different parameters, when the laser energy density is approximately 110 J/mm3 (the power is 310 W, the scanning speed is 800 mm/s, the layer thickness is 0.04 mm, and the scanning space is 0.09 mm), the sample has the highest density and best formability. 3.2. Phases and microstructures
Fig. 5. Relationship between input energy density and relative density.
Fig. 5. When the input energy density is extremely low, a discontinuous spread of melting tracks is formed because some of the powder has not been completely melted [36]. Meanwhile, the irregular boundaries of molten pools are produced even without fusion overlapping, as shown in Fig. 6(a). Consequently, irregular micro pores result from the insufficient melt (Fig. 6(a)), and a small amount of round gas pores (Fig. 6(b)) is found. With the increase in input energy density, the melting tracks become more continuous and fully overlap one another, as shown in Fig. 6(c). No pores are found on the overlapping boundaries (Fig. 6(d)). Fig. 6(e) shows that the samples that were fabricated with a very high energy density have nearly linear overlapping boundaries, and macro cracks are generated because of excessive residual stress. Moreover, many micro metal balls can be found, as shown in Fig. 6(f). The extremely high energy density can result in molten pool boiling due to the high metal temperatures and splashing of small powder particles [38]. In addition, the large laser energy density input produces large spheroidal particles that separate from one another, as shown in Fig. 7, which will create more void (Fig. 7(c)) formation during the powder-spreading process. The voids will finally cause
Fig. 8 shows the phase composition of the Cu-Al-Ni alloy from the phase diagram and XRD measurements. The theoretical solidification pathway of Cu-13.5Al-4Ni alloy is: L/Lþb/b/aþNiAlþg2. Thus, during the slow cooling process, the Cu-13.5Al-4Ni alloy is composed of a, NiAl and g2 phases at room temperature (RT), where a is a solid solution with the FCC structure of the Al element incorporated into Cu element, NiAl is an intermetallic compound, and g2 is a solid solution based on an electron compound [5]. However, the XRD results in Fig. 8(b) indicate only the b10 phase (the martensite phases that inherit the composition and order of austenite) in the SLM-fabricated Cu13.5Al-4Ni-0.5Ti samples, which is attributed to the rapid cooling and solidification process of SLM. Because the cooling time is insufficient, the transition of b1 phase into a phase and g2 phase can be suppressed. When the cooling rate is greater than 5e6 C/min, the eutectoid transition is suppressed [5]. The cooling rate of SLM is up to 2.13 106e2.97 106 C/s [27]. Therefore, the solidification route at a high cooling rate is: L/Lþb/b/b1/b1'. Commonly, the direct casting Cu-Al-Ni SMAs must be quenched to complete the thermoelastic martensitic transformation. Therefore, more thermo-elastic martensite (b10 ) remain at RT, which improves the shape memory properties of SLM-fabricated Cu-13.5Al-4Ni-0.5Ti SMAs. Fig. 9 illustrates the SEM micro morphologies of the SLMfabricated samples (laser power: 310 W; scanning speed: 800 mm/s; layer thickness: 0.04 mm; scanning space: 0.09 mm; laser energy density: 107 J/mm3). Fig. 9(a) and (b) show the grain morphology on the X-Y plane. The SLM-fabricated
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Fig. 6. Micro morphologies by OM: (a), (b) Point A (77 J/mm3, 250W and 900 mm/s); (c), (d) Point B (107 J/mm3, 310 W and 800 mm/s); (e), (f) Point C (147 J/mm3, 310 W and 600 mm/s).
Fig. 7. After spheroidization the schematic diagram of the spreading process of powder during SLM [39].
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Fig. 8. Phase composition of Cu-Al-Ni alloys: (a) Cu-Al-Ni (4% Ni fraction) ternary phase diagram (The red arrow shows the solidification path and phase transition of Cu-13.5Al-4Ni) [40]; (b) X-ray diffraction patterns of SLM-fabricated Cu-13.5Al-4Ni-0.5Ti samples (laser power: 310 W; scanning speed: 800 mm/s; layer thickness: 0.04 mm; scanning space: 0.09 mm; laser energy density: 107 J/mm3). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Cue13.5Ale4Nie0.5Ti sample exhibits a ‘‘bimodal’’ grain size distribution [41] with equiaxed grains having diameters in the range of 10e30 mm in the regions where the melt tracks overlap. In the centre of the melting tracks, the grains are elongated perpendicular to the scanning direction and have typical sizes of approximately 30e80 mm. Fig. 9(c) with a high magnification in Fig. 9(b) presents a typical lath martensite and in the internal cross grain growth, which is consistent with the XRD test. The average grain size can be estimated as follows [42],
d¼2
rffiffiffiffiffiffiffi s pN
(2)
where s is the area of a certain circle, N is the effective number of grains inside this circle, and d is the average diameter of grains. According to the high magnification morphology of grains, as shown in Fig. 9(d), the average diameter of grains calculated by Eq. (2) is approximately 43 mm. Saud et al. [43] fabricated a Cu-11.9Al4Ni-0.7Ti SMA by a conventional casting method with an average grain size of 400 mm. The grain size of SLM-fabricated samples is only 1/10 of that of the casting samples, which is the combined effect of the high cooling rate during SLM and the grain refining of the Ti element. Fig. 9(d) shows a micro pore with the diameter of approximately 10 mm. The reason is the fast cooling and solidification speed during SLM processing, so the gas carried in molten pools is too late to overflow [36]. The presence of these pores may
Fig. 9. SEM microstructure of the Cu-13.5Al-4Ni-0.5Ti alloy fabricated by SLM: (a) Microstructure on X-Y plane; (b) High magnification microstructure of the circle area marked in (a); (c) Martensite morphology; (d) Micro pore.
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prevent the SLM-fabricated samples from achieving full density. However, the so-called X-phase (Cu2TiAl) [44] described in the reference is not observed by SEM, which we will discuss next. The result of TEM will prove the existence of the X-phase. To prove the existence of the X-phase (Cu2TiAl), transmission electron microscopy (TEM) is applied. Fig. 10(a) shows the brightfield image of the SLM-fabricated sample (laser power: 310 W; scanning speed: 800 mm/s; layer thickness: 0.04 mm; scanning space: 0.09 mm; laser energy density: 107 J/mm3), which indicates that the X-phase (Cu2TiAl) is randomly distributed on the surface of the matrix. Fig. 10(b) and (c) present the selected area diffraction pattern (SADP) of Fig. 10(a). The X-phase (Cu2TiAl) and twin martensite can be identified from SADP. The X-phase is granular with a size of 20e50 nm, which indicates that the selective laser melting does not affect the structure of the precipitates but merely causes size refinement. The X-phase is dispersed on the substrate and hinders the growth of the grains, resulting in grain refinement [44]. Thus, the Ti element in the alloy can refine the grains. 3.3. Mechanical properties 3.3.1. Hardness Fig. 11 shows the hardness values of the Cu-13.5Al-4Ni-0.5Ti alloy fabricated by SLM (laser power: 310 W; scanning speed: 800 mm/s; layer thickness: 0.04 mm; scanning space: 0.09 mm; laser energy density: 107 J/mm3). The hardness values at different measurement points are similar, which indicates the uniform composition distribution in the sample. The average hardness on the X-Z plane (289.1 ± 16.9 HV, Fig. 11(b)) is slightly higher than that on the X-Y plane (267.1 ± 24.2 HV, Fig. 11(b)) because the X-Z plane has a slightly higher molten pool boundary density than the X-Y plane. However, there is no obvious difference in phases between X-Y and X-Z planes. Silva et al. [45] fabricated a Cu-11.85Al3.2Ni-3Mn SMA alloy by SLM (SLM Solutions GmbH) with an average hardness of 249.3 HV, which exceeds the casting standard by approximately 20.9 HV because of grain refinement. The average grain size of the Cu-11.85Al-3.2Ni-3Mn alloy exceeds 123 mm, which is 3 times higher than that of the alloy fabricated in this study. Therefore, the grain refinement effect of the Ti element
contributes to the high hardness of the SLM-fabricated Cu-13.5Al4Ni-0.5Ti alloy. The excessive cooling rate of SLM also inhibits the precipitation of the brittle g2 phase and increases the hardness of the alloy. 3.3.2. Tensile properties at RT Fig. 12 illustrates the room temperature stress-strain curves of the SLM-fabricated Cue13.5Ale4Nie0.5Ti SMAs for three process parameters: (a) 77 J/mm3 (250 W and 900 mm/s); (b) 107 J/mm3 (310 W and 800 mm/s); (c) 147 J/mm3 (310 W and 600 mm/s). The average ultimate tensile strength of the SLM-fabricated samples for the three process parameters are 420 ± 16 MPa, 541 ± 26 MPa and 476 ± 14 MPa, respectively. The average ultimate elongation of the SLM-fabricated samples for the three process parameters are 5.72 ± 0.31%, 7.63 ± 0.39% and 6.32 ± 0.23%, respectively. The results show that the SLM-fabricated sample has the best tensile properties at moderate energy densities (107 J/mm3, 310 W and 800 mm/s), which is consistent with the microscopic analysis results of the forming quality in Fig. 6. At low energy density, the powder are not fully melted to form continuous melting tracks and overlap zones, which results in the formation of voids and reduces the mechanical properties of the alloy. When the laser energy is too high, the temperature of the molten pool is too high to cause the spatter of the molten pool to form pores, which reduces the mechanical properties of the alloy. Compared with the as-cast Cu11.9Al-4Ni-0.7Ti SMAs [43], the ultimate tensile strength decreases by 23%, and the elongation increases by 163%. The high elongation is related to the grain refinement and phase composition in the alloy. The grain size of the as-cast Cu-11.9Al-4Ni-0.7Ti by Saud et al. [43] is approximately 400 mm, which is approximately 9e10 times larger than that of the alloy fabricated in this study. Additionally, the absence of the g2 brittle phase in samples should benefit the ductility improvement of SLM-fabricated Cue13.5Ale4Nie0.5Ti alloys. However, the ultimate tensile strength is 79 MPa lower than that of the as-cast Cu-11.9Al-4Ni-0.7Ti SMAs, which may be attributed to the presence of pores in Fig. 9(d). The reason is that the stress easily concentrates where pores are under tensile stress or compressive stress, so cracks are first formed from the pores and finally decreases the ultimate strength value [46].
Fig. 10. (a) TEM microstructure of the Cu-13.5Al-4Ni-0.5Ti alloy fabricated by SLM; (b), (c) Corresponding selected area diffraction pattern (SADP) data.
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Fig. 11. (a) Microhardness test specimen of the Cu-13.5Al-4Ni-0.5Ti alloy fabricated by SLM: (b) Average hardness of X-Y and X-Z planes.
Fig. 12. Room temperature stressestrain curves of the SLM-fabricated Cue13.5Ale4Nie0.5Ti SMAs under three process parameters: (a) 77 J/mm3 (250W and 900 mm/s); (b) 107 J/ mm3 (310 W and 800 mm/s); (c) 147 J/mm3 (310 W and 600 mm/s).
Fig. 13 illustrates the room temperature tensile fracture morphologies of the SLM-fabricated Cue13.5Ale4Nie0.5Ti alloys. The macroscopic fracture (Fig. 13(a)) is composed of many facets with irregular orientations, from which radial stripes (Fig. 13(b)) and bright facets (Fig. 13(b)) can be observed. Thus, the macroscopic fracture features show that the tensile samples have no obvious plastic deformation and exhibit a brittle fracture. Moreover, many cleavage steps, cleavage planes and river-like patterns can be found
on the microscopic fracture, as shown in Fig. 13(c). This appearance also indicates that the tensile samples are a cleavage fracture. Remarkably, gas pores (Fig. 13(d)) are observed in the fracture, around which there is a crack distribution. Thus, the cracks first appear at the pores due to stress concentration in the tension process. Then, the crack propagations interconnect to form steps, which will merge or disappear during propagation and eventually form a river-like pattern.
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Fig. 13. Room temperature tensile fracture morphologies of the SLM-fabricated Cue13.5Ale4Nie0.5Ti alloys (laser power: 310 W; scanning speed: 800 mm/s; layer thickness: 0.04 mm; scanning space: 0.09 mm): (a) SEM at low magnification; (b), (c) and (d) SEM at high magnification.
3.3.3. Tensile properties at 200 C Fig. 14 illustrates the DSC curve and tensile curves at 200 C of the SLM-fabricated Cue13.5Ale4Nie0.5Ti alloys (laser power: 310 W; scanning speed: 800 mm/s, layer thickness: 0.04 mm; scanning space: 0.09 mm; laser energy density: 107 J/mm3). Fig. 14(a) shows the martensitic transformation starting temperature of the alloy is approximately 83 C, the end temperature of the martensite transformation is approximately 40 C, the starting temperature of the austenite transformation is approximately 63 C, and the end temperature transition temperature of the austenite is approximately 117 C, which indicates that the alloy is in a martensite state at room temperature and has high-
temperature application potential. Because the working environment of Cu-Al-Ni alloys at high temperature is mainly at 100e200 C [35], the tensile tests were performed at 200 C, as shown in Fig. 14(b). The average ultimate tensile strength is 611 ± 9 MPa, which is nearly 70 MPa higher than that obtained at RT, whereas the average ultimate elongation is 10.78 ± 1.87%, which is nearly 3.15% higher than that obtained at RT. The strength is improved because the tensile samples are in the austenitic state at 200 C. The austenitic parent phase is a bcc structure, which is not prone to plastic slip under external force, whereas martensite is prone to twin deformation and martensite reorientation, so the strength of the parent phase is higher than that of the martensite
Fig. 14. (a) DSC curve of the SLM-fabricated Cu-13.5Al-4Ni-0.5Ti; (b) Tensile curves of the SLM-fabricated Cu-13.5Al-4Ni-0.5Ti alloys at 200 C.
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Fig. 15. High temperature (200 C) tensile fracture morphologies of the SLM-fabricated Cue13.5Ale4Nie0.5Ti alloys: (a) SEM at low magnification; (b), (c) and (d) SEM at high magnification.
phase. Fig. 14(b) shows that the samples undergo two stages of deformation during the stretching process at 200 C. Stage I is the elastic deformation stage, where the sample is in the austenite state. In stage Ⅱ, a stress-induced martensitic transformation occurs after the stress reaches a certain level [47], and a strain platform appears in the curve. Then, martensite deformation and brittle fracture occur after a certain degree of deformation. Thus, the plasticity of the alloy increases at 200 C. However, the stressinduced martensite is not stable. When the external force disappears, the martensite transforms into the austenite parent phase. Fig. 15 shows the high-temperature (200 C) tensile fracture morphologies of the SLM-fabricated Cue13.5Ale4Nie0.5Ti alloys. Apparent cleavage characteristics are observed on the fracture surface, as shown in Fig. 15(a). However, unlike the room temperature tensile, there are obvious fibrous regions and dimples in the high-magnification SEM morphologies, as shown in Fig. 15(b), (c) and (d). As a result, the fracture is characterized by the combination of brittle fracture and ductile fracture, so the strength and plasticity of the alloy are enhanced at 200 C, and there is obvious yielding. The study shows that the SLM-fabricated Cue13.5Ale4Nie0.5Ti alloys have good mechanical properties and application potential at high temperatures. 4. Conclusions The Cu-13.5Al-4Ni-0.5Ti copper-based SMA was fabricated by SLM from pre-alloy powder. The effect of the SLM processing parameters on the relative density was investigated, and the optimized parameters were determined. In addition, the microstructure, phase composition, and tensile properties at RT and 200 C of the SLM-fabricated Cu-13.5Al-4Ni-0.5Ti alloy were studied by OM, SEM, XRD, TEM, and tensile measurements. The main findings and conclusions can be summarized as follows. (1) When the scanning speed is 800e900 mm/s, the relative density is more than 99%, and the density slightly varies with the laser power and scanning speed. The alloy density is relatively high when the input energy density is approximately 110 J/mm3. When the energy density is too low, a
discontinuous spread of melting tracks will be formed because the powder is not completely melted. Meanwhile, the irregular boundaries of molten pools are produced; there may even be no fusion overlapping. Consequently, there are irregular micro pores as a result of insufficient melt and a small amount of round gas pores. A high input laser energy density will produce large, separate spheroidal particles, which form more voids during the powder-spreading process. Finally, the voids cause many pores to form in the SLMfabricated sample and decrease the density of the sample. (2) Due to the rapid cooling rate during the SLM process, the aphase and g2-phase are suppressed, and the b10 -phase and Xphase (Cu2TiAl) are generated in the SLM-fabricated Cu13.5Al-4Ni-0.5Ti alloy. Elongated strip-like grains traverse the melting tracks. Large-size grains are found at the centre of the melting tracks, whereas the grain size in the overlap zone is obviously refined. The average grain size is approximately 43 mm. (3) The hardness values of the SLM-fabricated Cu-13.5Al-4Ni0.5Ti alloy are 267.1e289.1 HV. The SLM-fabricated sample has the best tensile properties at moderate energy densities (107 J/mm3, 310 W and 800 mm/s). The average ultimate tensile strength and elongation are 541 ± 26 MPa and 7.63 ± 0.39% at RT, respectively. The elongation is approximately 5.41% higher than that of as-cast Cu-11.9Al-4Ni-0.7Ti SMAs. It should be related to the grain refinement, which is the combined effect of the high cooling rate during SLM and grain refining of the Ti element. The average ultimate tensile strength is 611 ± 9 MPa at 200 C, which increases by nearly 70 MPa compared to that obtained at RT. Accordingly, the elongation is 10.78 ± 1.87%, which increases by nearly 3.15% compared to that obtained at RT. The parent phase with the bcc structure is higher strength than martensite, and the stress-induced martensite transformation at 200 C increases the plasticity of the alloy. This study preliminary verifies the feasibility of SLM to prepare high-performance copper-based SMAs, but the effects of the element ratios, thermal and shape memory properties have not been considered. These topics may be the focus of our future study.
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Acknowledgments This work was funded by the academic frontier youth team at Huazhong University of Science and Technology (HUST), Hubei Science and Technology Support Program (2014AAA020), National Natural Science Fund Youth Fund (51701078) and Wuhan Key Technology Research Project (201501020201008). Meantime, thanks for the supporting tests of State Key Laboratory of Material Processing and Die & Mold Technology and Analysis and Test Centre at HUST. References [1] J.M. Jani, M. Leary, A. Subic, M.A. Gibson, Review of shape memory alloy research, applications and opportunities, Mater. Des. 56 (2014) 1078e1113. [2] N.B. Morgan, Medical shape memory alloy applicationsdthe market and its products, Mater. Sci. Eng. A 378 (2004) 16e23. [3] D.C. Lagoudas, Shape memory alloys, J. Model. Eng. Appl. (2008). [4] U. Sari, Influences of 2.5wt.% Mn addition on the microstructure and mechanical properties of Cu-Al-Ni shape memory alloys, Int. J. Miner. Metall. Mater. 17 (2010) 192e198. [5] X.J. Liu, R. Kaimuma, I. Qnuma, K. Ishida, Phase equilibria in the Cu-rich portion of the Cu-Al binary system, J. Alloys Compd. 264 (2000) 201e208. [6] L. Kovarik, F. Yang, A. Garg, D. Diercks, M. Kaufman, R.D. Noebe, M.J. Mills, Structural analysis of a new precipitate phase in high-temperature TiNiPt shape memory alloys, Acta Mater. 58 (2010) 4660e4673. [7] B. Kockar, K.C. Atli, J. Ma, M. Haouaoui, I. Karaman, M. Nagasako, R. Kainuma, Role of severe plastic deformation on the cyclic reversibility of a Ti33.716 high temperature shape memory alloy, Acta Mater. 58 (2010) 6411e6420. [8] T. Sawaguchi, M. Sato, A. Ishida, Grain-size effect on shape-memory behavior of Ti 35.0 Ni 49.7 Zr 15.4, thin films, Metall. Mater. Trans. A 35 (2004) 111e119. [9] F. Dalle, E. Perrin, P. Vermaut, M. Masse, R. Portier, Interface mobility in Ni 49.8 Ti 42.2 Hf 8, shape memory alloy, Acta Mater. 50 (2002) 3557e3565. rez-Landaza bal, V. Recarte, V. Sa nchez-Alarcos, M.L. No , Study of the [10] J.I. Pe stability and decomposition process of the b phase in CueAleNi shape memory alloys, Mater. Sci. Eng. A 438 (2006) 734e737. rez-Landaza bal, M.L. No , J.S. Juan, Study by resonant ultra[11] V. Recarte, J.I. Pe sound spectroscopy of the elastic constants of the b phase in Cu-Al-Ni shape memory alloys, Mater. Sci. Eng. A 370 (2004) 488e491. [12] S.N. Saud, T.A.A. Bakar, E. Hamzah, M.K. Ibrahim, A. Bahador, Effect of quarterly element addition of cobalt on phase transformation characteristics of CuAl-Ni shape memory alloys, Metall. Mater. Trans. A 46 (2015) 1e15. [13] J. Font, E. Cesari, J. Muntasell, J. Pons, Thermomechanical cycling in Cu-Al-Nibased melt-spun shape-memory ribbons, Mater. Sci. Eng. A 354 (2003) 207e211. [14] J.V. Humbeeck, R. Stalmans, M. Chandrasekaran, L. Delaey, On the stability of shape memory alloys, Eng. Asp. S. M. A. 35 (1990) 96e105. [15] K. Sugimoto, K. Kamei, M. Nakaniwa, Cu–Al–Ni–Mn: a new shape memory alloy for high temperature applications, Asp. S. M. A. 94 (1990) 89e95. [16] G. Motoyasu, M. Kaneko, H. Soda, A. Mclean, Continuously cast Cu-Al-Ni shape memory wires with a unidirectional morphology, Metall. Mater. Trans. A 32 (2001) 585e593. [17] M. Suezawa, K. Sumino, Behaviour of elastic constants in Cu-Al-Ni alloy in the close vicinity of Ms-point, Scr. Metall. 10 (1976) 789e792. [18] G. Guenin, M. Morin, P.F. Gobin, W. Dejonghe, L. Delaey, Elastic constant measurements in beta Cu-Zn-Al near the martensitic transformation temperature, Scr. Metall. 11 (1977) 1071e1075. [19] O. Mercier, K.N. Melton, G. Gremaud, et al., Single-crystal elastic constants of the equiatomic NiTi alloy near the martensitic transformation, J. Appl. Phys. 51 (1980) 1833e1834. [20] J.M. Long, S.Y. Zhou, Effect of heat treatment on Ms point of Cu-26.23Zn3.92Al-0.033B shape memory alloy, Nonferrous Met. 10 (1989) 1e7. [21] R. Elst, J.V. Humbeeck, L. Delaey, Grain refinement of CueZneAl and CueAleNi by Ti addition, J. Mater. Sci. Technol. 4 (2013) 644e648. [22] J. Dutkiewicz, T. Czeppe, J. Morgiel, Effect of titanium on structure and martensic transformation in rapidly solidified CueAleNieMneTi alloys, Mater. Sci. Eng. A s 273e275 (1999) 703e707. [23] G. Lojen, M. Goji c, I. An zel, Continuously cast CueAleNi shape memory alloy e properties in as-cast condition, J. Alloys Compd. 580 (2013) 497e505. [24] R.D. Cava, C. Bolfarini, C.S. Kiminami, E.M. Mazzer, W.J.B. Filho, P. Gargarella, J. Eckert, Spray forming of Cue11.85Ale3.2Nie3Mn (wt%) shape memory
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Process optimization, microstructures and mechanical properties of a Cu-based shape memory alloy fabricated by selective laser melting Jian Tian, Wenzhi Zhu**, Qingsong Wei*, Shifeng Wen, Shuai Li, Bo Song, Yusheng Shi State Key Laboratory of Materials Processing and Die & Mould Technology, School of Materials Science and Engineering, Huazhong University of Science and Technology, Wuhan 430074, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 23 April 2018 Received in revised form 3 December 2018 Accepted 12 January 2019 Available online 14 January 2019
The Cu-13.5Al-4Ni-0.5Ti copper-based shape memory alloys (SMAs) were fabricated by selective laser melting (SLM). The parameters were optimized to obtain almost fully dense copper-based SMAs samples. The phases and microstructures were characterized and the tensile properties at room temperature and 200 C were evaluated. The XRD results showed that only the b10 phase could be detected in the Cu-based SMAs, which was attributed to the extremely short solidification time for the precipitation of a and g2 phases during SLM process. The grains were well refined and the average grain size was approximate 43 mm, which was much smaller than that of the cast copper-based SMAs. The X-phase (Cu2TiAl) is observed, which is granular and size 20e50 nm. The Cu-13.5Al-4Ni-0.5Ti copper-based SMAs fabricated by SLM exhibited excellent mechanical properties at room temperature, which was higher than that of the cast copper-based SMAs. This is attributed to two aspects: (1) grain refinement, (2) suppress of brittle g2 phase. Remarkably, the tensile test results at 200 C showed both higher strength and elongation, which is attributed to the bcc structure of the parent phase and the stress-induced martensitic transformation at 200 C. Meanwhile, the difference of microstructures and properties between SLMfabricated Cu-based SMAs and casting Cu-based SMAs were analyzed and discussed in detail. © 2019 Elsevier B.V. All rights reserved.
Keywords: Copper-based shape memory alloy Selective laser melting Process optimization Relative density Microstructure Mechanical properties
1. Introduction Shape memory alloys (SMAs) are recently developed as advanced materials due to their specific properties such as superelasticity, shape memory effect (SME), biocompatibility, high specific strength, excellent corrosion resistance, good wear resistance and high anti-fatigue property [1]. Among them, Ni-Ti and Cubased alloys are two types of the most widely used SMAs [2]. Although Ni-Ti SMAs show better shape memory property and super-elasticity than do Cu-based SMAs [3], they are only used at a low operating temperature not exceeding 100 C [4], and they have poor machinability [5]. However, the substitution of Ni with noble metal elements such as Au, Pt, and Pd or the substitution of Ti with Zr and Hf can improve the transformation temperature of Ni-Ti SMAs [6e8]. And this type of Ni-Ti-based high-temperature shape memory alloys (HT-SMAs) is currently being studied extensively. In contrast, Cu-based SMAs can reach higher transformation
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (W. Zhu), [email protected] (Q. Wei). https://doi.org/10.1016/j.jallcom.2019.01.153 0925-8388/© 2019 Elsevier B.V. All rights reserved.
temperatures and are preferential candidates for high-temperature applications (such as in thermal actuators and thermal sensors) up to 240 C [9] due to their super-elastic or pseudo-elastic properties, two-way shape memory effect and high damping ability [10]. Moreover, the excellent processing performance of Cu-based SMAs enables the manufacture of complex-shaped parts [4]. Cu-based SMAs are currently divided into two series, namely, Cu-Al-Ni and Cu-Zn-Al [11], whose martensite transformation temperatures can be adjusted within a wide range [12]. Since the transformation temperature of Cu-based SMAs is sensitive to the variation of composition [3], it can be tailored by adjusting the ratios of elements in the alloy to satisfy various temperature requirements in different application environments. More significantly, Cu-Al-Ni exhibits a higher martensite transformation temperature than do Cu-Zn-Al alloys [13]. For example, Sugimoto et al. [14] developed a Cu-Al-Ni-Mn alloy with good thermodynamic stability and anti-ageing behaviour at temperatures above 100 C, and this alloy showed the capability for high-temperature applications. However, Cu-Al-Ni SMAs are brittle in nature due to the intergranular cracking, which is caused by the high elastic anisotropy
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[15]. For example, the elastic anisotropy factors of the Cu-Al-Ni, CuZn-Al, and Ni-Ti alloys are 13 [16], 15 [17], and 2 [18], respectively. The large elastic anisotropy makes the intergranular elasticity and plastic deformation of the copper-based SMAs extremely uncoordinated, which causes the grain boundary stress concentration and intergranular cracking. Currently, grain refinement and plasticity enhancement are two main methods to improve the mechanical properties of Cu-based SMAs. Researchers found that the addition of B, Ce and V elements could result in the grain refinement of CuZn-Al alloys [19], whereas the addition of Ti and Zr elements similarly affected Cu-Al-Ni alloys [20,21]. Alternative routes to increase the performance of Cu-Al-Ni alloys are rapid solidification methods such as melt-spinning [22], melt extraction [23], spray forming [24] and selective laser melting [25]. As an additive manufacturing (AM) technique, selective laser melting (SLM) can create bulk parts by melting specific predefined areas of a powder bed layer by layer [26]. The processing of a thin powder layer (20e50 mm) on massive substrate plates in combination with small laser spot diameters (approximately 100 mm) results in high intrinsic cooling rates of 2.13e2.97 106 C/s [27]. Fine grain sizes within approximately one micron can be obtained because of the high cooling rates during SLM. If the parts prepared by SLM have a high relative density, the yield and ultimate tensile strength are comparable to (even better than) those of conventional manufactured samples, whereas the fracture strain remains lower [26]. Recently, several studies investigated the fabrication and characterization of SLM-fabricated Ni-Ti alloys. Fabricated parts using selective laser melting (SLM) technology show low impurity [28], microstructure control [29], high temperature properties [30] and satisfy the requirements in the ASTM F2063-05 for medical NiTi devices [31]. However, those of copper-based SMAs are notably limited. Gustmann et al. [32,33] studied the effect of SLM processing parameters on the microstructure and mechanical properties of Cu-11.85Al-3.2Ni-3Mn SMAs and obtained nearly dense samples with higher mechanical properties than those of the casting counterparts. The SLM-fabricated Cu-11.85Al-3.2Ni-3Mn alloys have an average ultimate tensile strength of 620 ± 50 MPa and elongation of 8.2 ± 0.9% at room temperature (RT). The average hardness is 245 ± 20 HV. However, the mechanical properties of Cubased SMAs must be improved, and there is a lack of hightemperature performance studies. This study examines the SLM of a new Cu-based SMA, Cu13.5Al-4Ni-0.5Ti. According to the phase diagram of Cu-Al-Ni ternary alloy, the alloy with 4 wt % Ni and 13.5 wt % Al exhibits a good shape memory effect. The addition of Ti element is expected to improve the material strength because of the grain refinement effect [20]. The objectives are to obtain the optimal SLM processing parameters to fabricate fully dense samples (r 99%), clarify the phase formation, microstructures and tensile properties at RT and 200 C, and understand their performance characteristic and difference from casting alloys. The findings provide a basic insight on SLM-fabricating high-performance Cu-based SMAs with tailored microstructure and elevated mechanical properties. 2. Experimental 2.1. Materials Gas-atomized Cue13.5Ale4Nie0.5Ti powder was produced by using high-purity elements (>99.9 wt %) in an induction furnace in argon atmosphere. The density (7.11 g/cm3) of the produced alloy was determined by a fully automatic true density analyser (AccuPyc 1330, USA) as the reference value to calculate the relative densities of the SLM-fabricated samples. The morphology of the powder was observed by scanning electron microscopy (SEM JSM-7600F, Japan),
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as shown in Fig. 1(a). The powder particle shows a nearly spherical shape. Small satellite particles were observed to adhere to the surface of large particles. The powder size distribution was measured by a laser particle size analyser (Mastersizer 3000, UK), as illustrated in Fig. 1(b). The powder size was 19.8e46.7 mm with an average value of 30.5 mm. 2.2. Selective laser melting and samples preparation The samples were fabricated by an SLM 125HL machine (Solutions GmbH, Germany). The SLM machine was equipped with a 400-W Gaussian beam fibre laser. The laser beam diameter (F90%) is approximately 80 mm. The process was performed in an argon atmosphere (approximately 500 ppm oxygen). Based on previous extensive SLM forming experiments, the laser power and scanning speed have greater effect on the formed parts than the layer thickness and scan spacing. Therefore, the laser power and scanning speed were chosen as the experimental variables of SLM process optimization experiments. Some cube samples (in 8 8 5 mm3) for process optimization were fabricated by SLM. The processing parameters with fixed layer thickness (0.04 mm), fixed scanning spacing (0.09 mm), varying laser power (250e340 W) and varying scanning speed (600e900 mm/s) were chosen. The long bidirectional scanning strategy was used, as shown in Fig. 2. To reduce the residual thermal stress during SLM, a 67 rotation of the scanning direction between two consecutive layers N and N þ 1 was set [34]. The substrate was preheated to 200 C before processing. Fig. 3(a) illustrates the Cu-13.5Al-4Ni0.5Ti alloy cubes (8 8 5 mm3) fabricated by SLM using various processing parameters. In addition, the hardness testing samples (cubes of 8 8 5 mm3) and tensile samples were fabricated with the optimized processing parameter (maximum relative density), as shown in Fig. 3. 2.3. Materials characterization The density of the cube samples was measured using a highprecision balance (Sartorius MC210P, Germany), whose measurement principle is Archimedes method. The phase identification was performed by XRD (XRD-7000S, Shimadzu, Japan) with a Cu target in a conventional X-ray tube, which operated at an accelerating voltage of 40 kV and an electron beam current of 30 mA. The SLMfabricated samples for microstructure observations were cut from the substrate by wire-electrode cutting, ground and polished with automatic grinding and polishing machine (Ecomet300/Automet300, Buehler, USA), and chemically etched with a solution (vol. 50% H2O and vol. 50% HNO3) at room temperature for approximately 10 seconds. The microstructures were investigated by optical microscopy (OM VHX-1000C, Keyence, Japan), scanning electron microscopy (SEM JSM-7600F, Japan) and transmission electron microscopy (TEM JEM-2100, Japan). The Vickers microhardness of the samples was determined by a hardness test machine (430-SVD, Wilson Hardness, USA) with a 5kg load and 30-s dwell time. Five points were chosen and measured at different positions of each surface to obtain an average value. The tensile samples were cut to the required dimension by a CNC machine (WDW3200, China), as shown in Fig. 3(b). Because the working environment of Cu-Al-Ni alloys at high temperature is mainly at 100e200 C [35], the tensile tests were performed on a mechanical testing machine (Zwick/Roell Z020, Zwick, Germany) at RT and 200 C. The tests were performed using a cross speed of 0.5 mm/min and a constant tension load. The strain was recorded by a laser extensometer. The fracture morphology of the tensile samples was observed by scanning electron microscopy (SEM JSM7600F, Japan).
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Fig. 1. Cu-13.5Al-4Ni-0.5Ti powder: (a) The micromorphology under SEM, (b) Powder particle size distribution.
A), the density substantially decreases with the increase in laser power but increases with the increase in scanning speed. The combination of the processing parameter, a high laser power and a low scanning speed easily makes the metal solution boil in molten pools due to an extremely high temperature, which causes the splashing of small metal droplets and finally the formation of small pores in the samples and a lower density [36]. When the scanning speed is improved to 800e900 mm/s (Fig. 4(b), part B), the relative density exceeds 99%. In this case, the density only slightly varies with the laser power and scanning speed. Thus, the laser energy is sufficient and appropriate to completely melt powder to form a continuous molten pool. The input energy density is usually used to evaluate the effects of the processing parameters on the density of SLM-fabricated samples, which can be described as [37]:
E¼
Fig. 2. Schematic of laser scanning strategies during SLM.
3. Results and discussion 3.1. Effects of the SLM processing parameters on the relative density Fig. 4 illustrates the relative densities of Cue13.5Ale4Nie0.5Ti SMAs fabricated by SLM with various laser powers and scanning speeds. When the scanning speed is 600e700 mm/s (Fig. 4(a), part
P HDV
(1)
where E is the input energy density (in J/mm3), P is the laser power (in W), V is the scanning speed (in mm/s), H is the scanning space (in mm), and D is the layer thickness (in mm). Fig. 5 illustrates the relationship between the input energy density and the relative density of SLM-fabricated samples. When the input energy density is 77e110 J/mm3, the relative density exceeds 99% and slightly increases with the increase in energy density. However, the relative density significantly decreases with the increase in energy density when the input energy density is over 110 J/mm3. Fig. 6 displays the morphology in SLM-fabricated samples at three input laser energy densities values of points A, B and C in
Fig. 3. The Cu-13.5Al-4Ni-0.5Ti samples fabricated by SLM: (a) Cubes (8 8 5 mm3) fabricated by various processing parameters (layer thickness: 0.04 mm, scanning space: 0.09 mm); (b) Tensile samples (laser power: 310 W, scanning speed: 800 mm/s, layer thickness: 0.04 mm, scanning space: 0.09 mm).
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Fig. 4. (a) The relationship between the laser power and the scanning speed and the sample density; (b) Top view of Fig. 4a: Part A is the density value of the sample at the lower scanning speed (600e700 mm/s); Part B is the density value of the sample at the higher scanning speed (800e900 mm/s).
many pores to form in the SLM-fabricated sample, which decreases the density of the sample. Simultaneously, spheroidization will increase the molten pool surface roughness, affect the powder spreading and follow-up melt spread and accumulation, and reduce the relative density. Based on the density analysis of the test samples fabricated with different parameters, when the laser energy density is approximately 110 J/mm3 (the power is 310 W, the scanning speed is 800 mm/s, the layer thickness is 0.04 mm, and the scanning space is 0.09 mm), the sample has the highest density and best formability. 3.2. Phases and microstructures
Fig. 5. Relationship between input energy density and relative density.
Fig. 5. When the input energy density is extremely low, a discontinuous spread of melting tracks is formed because some of the powder has not been completely melted [36]. Meanwhile, the irregular boundaries of molten pools are produced even without fusion overlapping, as shown in Fig. 6(a). Consequently, irregular micro pores result from the insufficient melt (Fig. 6(a)), and a small amount of round gas pores (Fig. 6(b)) is found. With the increase in input energy density, the melting tracks become more continuous and fully overlap one another, as shown in Fig. 6(c). No pores are found on the overlapping boundaries (Fig. 6(d)). Fig. 6(e) shows that the samples that were fabricated with a very high energy density have nearly linear overlapping boundaries, and macro cracks are generated because of excessive residual stress. Moreover, many micro metal balls can be found, as shown in Fig. 6(f). The extremely high energy density can result in molten pool boiling due to the high metal temperatures and splashing of small powder particles [38]. In addition, the large laser energy density input produces large spheroidal particles that separate from one another, as shown in Fig. 7, which will create more void (Fig. 7(c)) formation during the powder-spreading process. The voids will finally cause
Fig. 8 shows the phase composition of the Cu-Al-Ni alloy from the phase diagram and XRD measurements. The theoretical solidification pathway of Cu-13.5Al-4Ni alloy is: L/Lþb/b/aþNiAlþg2. Thus, during the slow cooling process, the Cu-13.5Al-4Ni alloy is composed of a, NiAl and g2 phases at room temperature (RT), where a is a solid solution with the FCC structure of the Al element incorporated into Cu element, NiAl is an intermetallic compound, and g2 is a solid solution based on an electron compound [5]. However, the XRD results in Fig. 8(b) indicate only the b10 phase (the martensite phases that inherit the composition and order of austenite) in the SLM-fabricated Cu13.5Al-4Ni-0.5Ti samples, which is attributed to the rapid cooling and solidification process of SLM. Because the cooling time is insufficient, the transition of b1 phase into a phase and g2 phase can be suppressed. When the cooling rate is greater than 5e6 C/min, the eutectoid transition is suppressed [5]. The cooling rate of SLM is up to 2.13 106e2.97 106 C/s [27]. Therefore, the solidification route at a high cooling rate is: L/Lþb/b/b1/b1'. Commonly, the direct casting Cu-Al-Ni SMAs must be quenched to complete the thermoelastic martensitic transformation. Therefore, more thermo-elastic martensite (b10 ) remain at RT, which improves the shape memory properties of SLM-fabricated Cu-13.5Al-4Ni-0.5Ti SMAs. Fig. 9 illustrates the SEM micro morphologies of the SLMfabricated samples (laser power: 310 W; scanning speed: 800 mm/s; layer thickness: 0.04 mm; scanning space: 0.09 mm; laser energy density: 107 J/mm3). Fig. 9(a) and (b) show the grain morphology on the X-Y plane. The SLM-fabricated
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Fig. 6. Micro morphologies by OM: (a), (b) Point A (77 J/mm3, 250W and 900 mm/s); (c), (d) Point B (107 J/mm3, 310 W and 800 mm/s); (e), (f) Point C (147 J/mm3, 310 W and 600 mm/s).
Fig. 7. After spheroidization the schematic diagram of the spreading process of powder during SLM [39].
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Fig. 8. Phase composition of Cu-Al-Ni alloys: (a) Cu-Al-Ni (4% Ni fraction) ternary phase diagram (The red arrow shows the solidification path and phase transition of Cu-13.5Al-4Ni) [40]; (b) X-ray diffraction patterns of SLM-fabricated Cu-13.5Al-4Ni-0.5Ti samples (laser power: 310 W; scanning speed: 800 mm/s; layer thickness: 0.04 mm; scanning space: 0.09 mm; laser energy density: 107 J/mm3). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
Cue13.5Ale4Nie0.5Ti sample exhibits a ‘‘bimodal’’ grain size distribution [41] with equiaxed grains having diameters in the range of 10e30 mm in the regions where the melt tracks overlap. In the centre of the melting tracks, the grains are elongated perpendicular to the scanning direction and have typical sizes of approximately 30e80 mm. Fig. 9(c) with a high magnification in Fig. 9(b) presents a typical lath martensite and in the internal cross grain growth, which is consistent with the XRD test. The average grain size can be estimated as follows [42],
d¼2
rffiffiffiffiffiffiffi s pN
(2)
where s is the area of a certain circle, N is the effective number of grains inside this circle, and d is the average diameter of grains. According to the high magnification morphology of grains, as shown in Fig. 9(d), the average diameter of grains calculated by Eq. (2) is approximately 43 mm. Saud et al. [43] fabricated a Cu-11.9Al4Ni-0.7Ti SMA by a conventional casting method with an average grain size of 400 mm. The grain size of SLM-fabricated samples is only 1/10 of that of the casting samples, which is the combined effect of the high cooling rate during SLM and the grain refining of the Ti element. Fig. 9(d) shows a micro pore with the diameter of approximately 10 mm. The reason is the fast cooling and solidification speed during SLM processing, so the gas carried in molten pools is too late to overflow [36]. The presence of these pores may
Fig. 9. SEM microstructure of the Cu-13.5Al-4Ni-0.5Ti alloy fabricated by SLM: (a) Microstructure on X-Y plane; (b) High magnification microstructure of the circle area marked in (a); (c) Martensite morphology; (d) Micro pore.
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prevent the SLM-fabricated samples from achieving full density. However, the so-called X-phase (Cu2TiAl) [44] described in the reference is not observed by SEM, which we will discuss next. The result of TEM will prove the existence of the X-phase. To prove the existence of the X-phase (Cu2TiAl), transmission electron microscopy (TEM) is applied. Fig. 10(a) shows the brightfield image of the SLM-fabricated sample (laser power: 310 W; scanning speed: 800 mm/s; layer thickness: 0.04 mm; scanning space: 0.09 mm; laser energy density: 107 J/mm3), which indicates that the X-phase (Cu2TiAl) is randomly distributed on the surface of the matrix. Fig. 10(b) and (c) present the selected area diffraction pattern (SADP) of Fig. 10(a). The X-phase (Cu2TiAl) and twin martensite can be identified from SADP. The X-phase is granular with a size of 20e50 nm, which indicates that the selective laser melting does not affect the structure of the precipitates but merely causes size refinement. The X-phase is dispersed on the substrate and hinders the growth of the grains, resulting in grain refinement [44]. Thus, the Ti element in the alloy can refine the grains. 3.3. Mechanical properties 3.3.1. Hardness Fig. 11 shows the hardness values of the Cu-13.5Al-4Ni-0.5Ti alloy fabricated by SLM (laser power: 310 W; scanning speed: 800 mm/s; layer thickness: 0.04 mm; scanning space: 0.09 mm; laser energy density: 107 J/mm3). The hardness values at different measurement points are similar, which indicates the uniform composition distribution in the sample. The average hardness on the X-Z plane (289.1 ± 16.9 HV, Fig. 11(b)) is slightly higher than that on the X-Y plane (267.1 ± 24.2 HV, Fig. 11(b)) because the X-Z plane has a slightly higher molten pool boundary density than the X-Y plane. However, there is no obvious difference in phases between X-Y and X-Z planes. Silva et al. [45] fabricated a Cu-11.85Al3.2Ni-3Mn SMA alloy by SLM (SLM Solutions GmbH) with an average hardness of 249.3 HV, which exceeds the casting standard by approximately 20.9 HV because of grain refinement. The average grain size of the Cu-11.85Al-3.2Ni-3Mn alloy exceeds 123 mm, which is 3 times higher than that of the alloy fabricated in this study. Therefore, the grain refinement effect of the Ti element
contributes to the high hardness of the SLM-fabricated Cu-13.5Al4Ni-0.5Ti alloy. The excessive cooling rate of SLM also inhibits the precipitation of the brittle g2 phase and increases the hardness of the alloy. 3.3.2. Tensile properties at RT Fig. 12 illustrates the room temperature stress-strain curves of the SLM-fabricated Cue13.5Ale4Nie0.5Ti SMAs for three process parameters: (a) 77 J/mm3 (250 W and 900 mm/s); (b) 107 J/mm3 (310 W and 800 mm/s); (c) 147 J/mm3 (310 W and 600 mm/s). The average ultimate tensile strength of the SLM-fabricated samples for the three process parameters are 420 ± 16 MPa, 541 ± 26 MPa and 476 ± 14 MPa, respectively. The average ultimate elongation of the SLM-fabricated samples for the three process parameters are 5.72 ± 0.31%, 7.63 ± 0.39% and 6.32 ± 0.23%, respectively. The results show that the SLM-fabricated sample has the best tensile properties at moderate energy densities (107 J/mm3, 310 W and 800 mm/s), which is consistent with the microscopic analysis results of the forming quality in Fig. 6. At low energy density, the powder are not fully melted to form continuous melting tracks and overlap zones, which results in the formation of voids and reduces the mechanical properties of the alloy. When the laser energy is too high, the temperature of the molten pool is too high to cause the spatter of the molten pool to form pores, which reduces the mechanical properties of the alloy. Compared with the as-cast Cu11.9Al-4Ni-0.7Ti SMAs [43], the ultimate tensile strength decreases by 23%, and the elongation increases by 163%. The high elongation is related to the grain refinement and phase composition in the alloy. The grain size of the as-cast Cu-11.9Al-4Ni-0.7Ti by Saud et al. [43] is approximately 400 mm, which is approximately 9e10 times larger than that of the alloy fabricated in this study. Additionally, the absence of the g2 brittle phase in samples should benefit the ductility improvement of SLM-fabricated Cue13.5Ale4Nie0.5Ti alloys. However, the ultimate tensile strength is 79 MPa lower than that of the as-cast Cu-11.9Al-4Ni-0.7Ti SMAs, which may be attributed to the presence of pores in Fig. 9(d). The reason is that the stress easily concentrates where pores are under tensile stress or compressive stress, so cracks are first formed from the pores and finally decreases the ultimate strength value [46].
Fig. 10. (a) TEM microstructure of the Cu-13.5Al-4Ni-0.5Ti alloy fabricated by SLM; (b), (c) Corresponding selected area diffraction pattern (SADP) data.
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Fig. 11. (a) Microhardness test specimen of the Cu-13.5Al-4Ni-0.5Ti alloy fabricated by SLM: (b) Average hardness of X-Y and X-Z planes.
Fig. 12. Room temperature stressestrain curves of the SLM-fabricated Cue13.5Ale4Nie0.5Ti SMAs under three process parameters: (a) 77 J/mm3 (250W and 900 mm/s); (b) 107 J/ mm3 (310 W and 800 mm/s); (c) 147 J/mm3 (310 W and 600 mm/s).
Fig. 13 illustrates the room temperature tensile fracture morphologies of the SLM-fabricated Cue13.5Ale4Nie0.5Ti alloys. The macroscopic fracture (Fig. 13(a)) is composed of many facets with irregular orientations, from which radial stripes (Fig. 13(b)) and bright facets (Fig. 13(b)) can be observed. Thus, the macroscopic fracture features show that the tensile samples have no obvious plastic deformation and exhibit a brittle fracture. Moreover, many cleavage steps, cleavage planes and river-like patterns can be found
on the microscopic fracture, as shown in Fig. 13(c). This appearance also indicates that the tensile samples are a cleavage fracture. Remarkably, gas pores (Fig. 13(d)) are observed in the fracture, around which there is a crack distribution. Thus, the cracks first appear at the pores due to stress concentration in the tension process. Then, the crack propagations interconnect to form steps, which will merge or disappear during propagation and eventually form a river-like pattern.
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Fig. 13. Room temperature tensile fracture morphologies of the SLM-fabricated Cue13.5Ale4Nie0.5Ti alloys (laser power: 310 W; scanning speed: 800 mm/s; layer thickness: 0.04 mm; scanning space: 0.09 mm): (a) SEM at low magnification; (b), (c) and (d) SEM at high magnification.
3.3.3. Tensile properties at 200 C Fig. 14 illustrates the DSC curve and tensile curves at 200 C of the SLM-fabricated Cue13.5Ale4Nie0.5Ti alloys (laser power: 310 W; scanning speed: 800 mm/s, layer thickness: 0.04 mm; scanning space: 0.09 mm; laser energy density: 107 J/mm3). Fig. 14(a) shows the martensitic transformation starting temperature of the alloy is approximately 83 C, the end temperature of the martensite transformation is approximately 40 C, the starting temperature of the austenite transformation is approximately 63 C, and the end temperature transition temperature of the austenite is approximately 117 C, which indicates that the alloy is in a martensite state at room temperature and has high-
temperature application potential. Because the working environment of Cu-Al-Ni alloys at high temperature is mainly at 100e200 C [35], the tensile tests were performed at 200 C, as shown in Fig. 14(b). The average ultimate tensile strength is 611 ± 9 MPa, which is nearly 70 MPa higher than that obtained at RT, whereas the average ultimate elongation is 10.78 ± 1.87%, which is nearly 3.15% higher than that obtained at RT. The strength is improved because the tensile samples are in the austenitic state at 200 C. The austenitic parent phase is a bcc structure, which is not prone to plastic slip under external force, whereas martensite is prone to twin deformation and martensite reorientation, so the strength of the parent phase is higher than that of the martensite
Fig. 14. (a) DSC curve of the SLM-fabricated Cu-13.5Al-4Ni-0.5Ti; (b) Tensile curves of the SLM-fabricated Cu-13.5Al-4Ni-0.5Ti alloys at 200 C.
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Fig. 15. High temperature (200 C) tensile fracture morphologies of the SLM-fabricated Cue13.5Ale4Nie0.5Ti alloys: (a) SEM at low magnification; (b), (c) and (d) SEM at high magnification.
phase. Fig. 14(b) shows that the samples undergo two stages of deformation during the stretching process at 200 C. Stage I is the elastic deformation stage, where the sample is in the austenite state. In stage Ⅱ, a stress-induced martensitic transformation occurs after the stress reaches a certain level [47], and a strain platform appears in the curve. Then, martensite deformation and brittle fracture occur after a certain degree of deformation. Thus, the plasticity of the alloy increases at 200 C. However, the stressinduced martensite is not stable. When the external force disappears, the martensite transforms into the austenite parent phase. Fig. 15 shows the high-temperature (200 C) tensile fracture morphologies of the SLM-fabricated Cue13.5Ale4Nie0.5Ti alloys. Apparent cleavage characteristics are observed on the fracture surface, as shown in Fig. 15(a). However, unlike the room temperature tensile, there are obvious fibrous regions and dimples in the high-magnification SEM morphologies, as shown in Fig. 15(b), (c) and (d). As a result, the fracture is characterized by the combination of brittle fracture and ductile fracture, so the strength and plasticity of the alloy are enhanced at 200 C, and there is obvious yielding. The study shows that the SLM-fabricated Cue13.5Ale4Nie0.5Ti alloys have good mechanical properties and application potential at high temperatures. 4. Conclusions The Cu-13.5Al-4Ni-0.5Ti copper-based SMA was fabricated by SLM from pre-alloy powder. The effect of the SLM processing parameters on the relative density was investigated, and the optimized parameters were determined. In addition, the microstructure, phase composition, and tensile properties at RT and 200 C of the SLM-fabricated Cu-13.5Al-4Ni-0.5Ti alloy were studied by OM, SEM, XRD, TEM, and tensile measurements. The main findings and conclusions can be summarized as follows. (1) When the scanning speed is 800e900 mm/s, the relative density is more than 99%, and the density slightly varies with the laser power and scanning speed. The alloy density is relatively high when the input energy density is approximately 110 J/mm3. When the energy density is too low, a
discontinuous spread of melting tracks will be formed because the powder is not completely melted. Meanwhile, the irregular boundaries of molten pools are produced; there may even be no fusion overlapping. Consequently, there are irregular micro pores as a result of insufficient melt and a small amount of round gas pores. A high input laser energy density will produce large, separate spheroidal particles, which form more voids during the powder-spreading process. Finally, the voids cause many pores to form in the SLMfabricated sample and decrease the density of the sample. (2) Due to the rapid cooling rate during the SLM process, the aphase and g2-phase are suppressed, and the b10 -phase and Xphase (Cu2TiAl) are generated in the SLM-fabricated Cu13.5Al-4Ni-0.5Ti alloy. Elongated strip-like grains traverse the melting tracks. Large-size grains are found at the centre of the melting tracks, whereas the grain size in the overlap zone is obviously refined. The average grain size is approximately 43 mm. (3) The hardness values of the SLM-fabricated Cu-13.5Al-4Ni0.5Ti alloy are 267.1e289.1 HV. The SLM-fabricated sample has the best tensile properties at moderate energy densities (107 J/mm3, 310 W and 800 mm/s). The average ultimate tensile strength and elongation are 541 ± 26 MPa and 7.63 ± 0.39% at RT, respectively. The elongation is approximately 5.41% higher than that of as-cast Cu-11.9Al-4Ni-0.7Ti SMAs. It should be related to the grain refinement, which is the combined effect of the high cooling rate during SLM and grain refining of the Ti element. The average ultimate tensile strength is 611 ± 9 MPa at 200 C, which increases by nearly 70 MPa compared to that obtained at RT. Accordingly, the elongation is 10.78 ± 1.87%, which increases by nearly 3.15% compared to that obtained at RT. The parent phase with the bcc structure is higher strength than martensite, and the stress-induced martensite transformation at 200 C increases the plasticity of the alloy. This study preliminary verifies the feasibility of SLM to prepare high-performance copper-based SMAs, but the effects of the element ratios, thermal and shape memory properties have not been considered. These topics may be the focus of our future study.
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Acknowledgments This work was funded by the academic frontier youth team at Huazhong University of Science and Technology (HUST), Hubei Science and Technology Support Program (2014AAA020), National Natural Science Fund Youth Fund (51701078) and Wuhan Key Technology Research Project (201501020201008). Meantime, thanks for the supporting tests of State Key Laboratory of Material Processing and Die & Mold Technology and Analysis and Test Centre at HUST. References [1] J.M. Jani, M. Leary, A. Subic, M.A. Gibson, Review of shape memory alloy research, applications and opportunities, Mater. Des. 56 (2014) 1078e1113. [2] N.B. Morgan, Medical shape memory alloy applicationsdthe market and its products, Mater. Sci. Eng. A 378 (2004) 16e23. [3] D.C. Lagoudas, Shape memory alloys, J. Model. Eng. Appl. (2008). [4] U. Sari, Influences of 2.5wt.% Mn addition on the microstructure and mechanical properties of Cu-Al-Ni shape memory alloys, Int. J. Miner. Metall. Mater. 17 (2010) 192e198. [5] X.J. Liu, R. Kaimuma, I. Qnuma, K. Ishida, Phase equilibria in the Cu-rich portion of the Cu-Al binary system, J. Alloys Compd. 264 (2000) 201e208. [6] L. Kovarik, F. Yang, A. Garg, D. Diercks, M. Kaufman, R.D. Noebe, M.J. Mills, Structural analysis of a new precipitate phase in high-temperature TiNiPt shape memory alloys, Acta Mater. 58 (2010) 4660e4673. [7] B. Kockar, K.C. Atli, J. Ma, M. Haouaoui, I. Karaman, M. Nagasako, R. Kainuma, Role of severe plastic deformation on the cyclic reversibility of a Ti33.716 high temperature shape memory alloy, Acta Mater. 58 (2010) 6411e6420. [8] T. Sawaguchi, M. Sato, A. Ishida, Grain-size effect on shape-memory behavior of Ti 35.0 Ni 49.7 Zr 15.4, thin films, Metall. Mater. Trans. A 35 (2004) 111e119. [9] F. Dalle, E. Perrin, P. Vermaut, M. Masse, R. Portier, Interface mobility in Ni 49.8 Ti 42.2 Hf 8, shape memory alloy, Acta Mater. 50 (2002) 3557e3565. rez-Landaza bal, V. Recarte, V. Sa nchez-Alarcos, M.L. No , Study of the [10] J.I. Pe stability and decomposition process of the b phase in CueAleNi shape memory alloys, Mater. Sci. Eng. A 438 (2006) 734e737. rez-Landaza bal, M.L. No , J.S. Juan, Study by resonant ultra[11] V. Recarte, J.I. Pe sound spectroscopy of the elastic constants of the b phase in Cu-Al-Ni shape memory alloys, Mater. Sci. Eng. A 370 (2004) 488e491. [12] S.N. Saud, T.A.A. Bakar, E. Hamzah, M.K. Ibrahim, A. Bahador, Effect of quarterly element addition of cobalt on phase transformation characteristics of CuAl-Ni shape memory alloys, Metall. Mater. Trans. A 46 (2015) 1e15. [13] J. Font, E. Cesari, J. Muntasell, J. Pons, Thermomechanical cycling in Cu-Al-Nibased melt-spun shape-memory ribbons, Mater. Sci. Eng. A 354 (2003) 207e211. [14] J.V. Humbeeck, R. Stalmans, M. Chandrasekaran, L. Delaey, On the stability of shape memory alloys, Eng. Asp. S. M. A. 35 (1990) 96e105. [15] K. Sugimoto, K. Kamei, M. Nakaniwa, Cu–Al–Ni–Mn: a new shape memory alloy for high temperature applications, Asp. S. M. A. 94 (1990) 89e95. [16] G. Motoyasu, M. Kaneko, H. Soda, A. Mclean, Continuously cast Cu-Al-Ni shape memory wires with a unidirectional morphology, Metall. Mater. Trans. A 32 (2001) 585e593. [17] M. Suezawa, K. Sumino, Behaviour of elastic constants in Cu-Al-Ni alloy in the close vicinity of Ms-point, Scr. Metall. 10 (1976) 789e792. [18] G. Guenin, M. Morin, P.F. Gobin, W. Dejonghe, L. Delaey, Elastic constant measurements in beta Cu-Zn-Al near the martensitic transformation temperature, Scr. Metall. 11 (1977) 1071e1075. [19] O. Mercier, K.N. Melton, G. Gremaud, et al., Single-crystal elastic constants of the equiatomic NiTi alloy near the martensitic transformation, J. Appl. Phys. 51 (1980) 1833e1834. [20] J.M. Long, S.Y. Zhou, Effect of heat treatment on Ms point of Cu-26.23Zn3.92Al-0.033B shape memory alloy, Nonferrous Met. 10 (1989) 1e7. [21] R. Elst, J.V. Humbeeck, L. Delaey, Grain refinement of CueZneAl and CueAleNi by Ti addition, J. Mater. Sci. Technol. 4 (2013) 644e648. [22] J. Dutkiewicz, T. Czeppe, J. Morgiel, Effect of titanium on structure and martensic transformation in rapidly solidified CueAleNieMneTi alloys, Mater. Sci. Eng. A s 273e275 (1999) 703e707. [23] G. Lojen, M. Goji c, I. An zel, Continuously cast CueAleNi shape memory alloy e properties in as-cast condition, J. Alloys Compd. 580 (2013) 497e505. [24] R.D. Cava, C. Bolfarini, C.S. Kiminami, E.M. Mazzer, W.J.B. Filho, P. Gargarella, J. Eckert, Spray forming of Cue11.85Ale3.2Nie3Mn (wt%) shape memory
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