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Article

Transformation Behavior and Shape Memory Effect of Ni47Ti44Nb9 Alloy Synthesized by Laser Powder Bed Fusion and Heat Treating

by
Mingyan Sun
,
Jie Chen
,
Qichao Fan
,
Chuan Yang
,
Guowei Wang
,
Xianfeng Shen
,
Yangyang Wang
,
Yonghao Zhang
* and
Shuke Huang
*
Institute of Machinery Manufacturing Technology, China Academy of Engineering Physics, Mianyang 621900, China
*
Authors to whom correspondence should be addressed.
Metals 2022, 12(9), 1438; https://doi.org/10.3390/met12091438
Submission received: 26 July 2022 / Revised: 22 August 2022 / Accepted: 26 August 2022 / Published: 29 August 2022
(This article belongs to the Section Additive Manufacturing)
Browse Figures
Review Reports Versions Notes

Abstract

:
Ni47Ti44Nb9 alloys were successfully fabricated by laser powder bed fusion (LPBF) technique with different laser powers. The phase transformation behavior, tensile properties and shape memory response before and after heat treating were also investigated. The Ni47Ti44Nb9 LPBF alloys have good shaping properties, though a few defects were discovered. Phase transformation peaks did not appear in the as-built samples, but were observed in the heat-treated samples. The phase transformation temperatures of the heat-treated samples increase with the increase in laser power. The tension test at room temperature indicates that the LPBF samples exhibit poor tensile ductility, which may be related to the existence of pores and Ti2Ni or Ti4Ni2Ox phase during the LPBF process. However, the LPBF samples after heat treating still possess good shape memory effect (with recovery strain about 7.82–8%) and relatively high reverse transformation temperature (about 36–52.6 °C) when deformed to 8%.

1. Introduction

NiTiNb shape memory alloys (SMAs) have attracted considerable attention in aeronautics and astronautics and civil industrial, medical and other applications. One common application is coupling and sealing, since the materials not only demonstrate good shape memory effect (SME), but also extraordinarily wide hysteresis [1,2,3,4,5]. Take the representative Ni47Ti44Nb9 (at.%) alloy, for example; the products made of this material can be stored and shipped as martensite at room temperature after predeformation [5]. Another application for NiTiNb SMAs is used as energy absorption or damping materials, due to their high damping capacity [6,7,8,9]. In addition, NiTiNb SMAs also can be applied in biomedical implants, as Nb exhibits superior corrosion resistance and outstanding biocompatibility [8,10,11].
To date, different techniques have been used to manufacture NiTiNb alloys. One common conventional method is casting followed by thermomechanical deformation, heat treating and machining [2,12]. Another is powder metallurgy (PM) routes [8,13,14,15,16]. The former is very successful in achieving simple structures (e.g., rods, tubes, wires and sheets), which can satisfy the requirement for most applications [2]. However, it is not suitable for complex structures such as porous structures, as it is known that the damping capacity of NiTiNb could be improved by making them porous and most biomedical implants are also porous [7,8,11,17]. PM is used for producing near-net-shape devices, but is still limited in the complexity of the resulting parts and in controlling the size and shape of pores [17,18]. Recently, laser powder bed fusion (LPBF), as a popular form of additive manufacturing (AM), is attracting increasing attention due to the ability to fabricate free-form geometries and complex near-net-shape parts without any extensive machining or complex dies [19,20]. This technique has been used in many complex metal components such as steel, aluminum, titanium and nickel alloys [21,22,23,24,25,26]. Thus, it will also offer a promising solution to fabricate NiTiNb complex components. Moreover, the materials will experience a rapid melting–solidification process and a complex aging process during LPBF, which is very different from traditional methods [27]. These excessive thermal conditions and specific thermal histories may result in complex microstructures, thus remarkably affecting the properties of NiTiNb SMAs.
Over the past decade, many studies have been conducted in manufacturing NiTi binary alloys by using LPBF. These studies have focused on the influence of process parameters on microstructure, phase transformation behavior and mechanical and functional properties of the LPBF NiTi alloys [27,28,29,30,31,32,33]. In contrast, there are only a few studies about ternary NiTi-based SMAs. Mohammad Elahinia et al. investigated the thermomechanical and shape memory response of NiTiHf high-temperature SMAs fabricated via LPBF process [34]. Shiva et al. investigated the mechanical properties and phase transformation behavior of TiNiCu SMAs fabricated via LPBF process [35]. However, little work has been performed on the LPBF of NiTiNb alloys. In this paper, Ni47Ti44Nb9 SMAs were fabricated for the first time via LPBF. In addition, the phase transformation behavior and SME of LPBF Ni47Ti44Nb9 alloys were also investigated.

2. Experimental Procedures

2.1. LPBF Process

The Ni47Ti44Nb9 (at.%) rod, produced by vacuum induction melting followed by hot-forging, was atomized to powder through an electrode induction-melting gas atomization (EIGA) process by Avimetal Powder Metallurgy Technology Co., Ltd. (Beijing, China) and then the powder was sieved with 15–53 micron mesh. The chemical composition of the Ni47Ti44Nb9 pre-alloyed powder and rod used for producing the powder is given in Table 1. Compared with the rod, there is a slight decrease in Ni content for the powder. The Ni loss associated with powder preparation can mostly be attributed to the fact that the Ni element tends to evaporate sooner than the Ti and Nb element due to its lower boiling point. The secondary electron image characterized via scanning electron microscope (SEM) shows the morphology of the powder (Figure 1a). It is clear that most powder particles exhibit a spherical shape. Figure 1b shows the corresponding particle size distribution of the powder. The particle size ranges from d10 = 20 μm to d90 = 56 μm.
SLM 100B AM machine (IPG fiber laser with maximum power of 200 W and beam size of 50 μm) developed by Nanjing University of Aeronautics and Astronautics (Nanjing, China) was used to fabricate NiTiNb samples under argon atmosphere. NiTi build platform was utilized without pre-heating. Bidirectional scanning strategy with a hatching rotation of 67° between the layers was employed during the LPBF process. Four groups of parameters with different laser powers (130–190 W) were used, which are listed in Table 2. The LPBF energy density is calculated from laser power (P), scanning velocity (v), hatch spacing (h) and layer thickness (t) as the following formula [18,19]:
E = P v × h × t
The corresponding LPBF-fabricated NiTiNb samples are illustrated in Figure 1c. Rectangular samples with size of 10 mm (length) × 5 mm (width) × 5 mm (height) were produced for characterizing microstructure, phase composition and phase transformation. Strip-shaped samples with size of 35 mm (length) × 3 mm (width) × 5 mm (height) were produced for characterizing tension property, transformation hysteresis and shape memory effect. X-axis runs parallel to the front of the machine, Z-axis runs normal to the layers and Y-axis runs perpendicular to the X-axis and Z-axis. The corresponding planes are noted as XZ, XY and YZ planes in this paper.

2.2. Characterization

Particle size distribution was measured by laser diffraction method on S3500-SI particle size analyzer (Microtrac Inc., Largo, FL, USA). The microstructures and microarea compositions were characterized by GX53 optical microscopy (OM, Olympus, Tokyo, Japan), Talos F200X transmission electron microscope (TEM, FEI, Hillsboro, OR, USA) and Auriga SEM (Zeiss, Oberkochen, Germany) with backscattered electron (BSE) detector and energy-dispersive spectroscopy (EDS, Oxford Instruments, Oxford, UK). Phase identification was carried out by Cu Kα radiation using X-ray diffractometer (XRD, Empyrean, Panalytical, Almelo, The Netherlands). Phase-transformation behavior was determined by differential scanning calorimeter (DSC, DSC 214 Polyma, Netzsch, Selb, Germany) at a 10 °C/min heating/cooling rate. The tensile stress–strain response was tested in 5985 electronic universal testing system (Instron, Boston, MA, USA) at strain rate of 2.5 × 10−4 s−1. The SME was explored under different tensile deformations of 2%, 4%, 6% and 8% and conducted on cryogenic tensile loading–unloading and subsequent heating. Loading, unloading and heating operations were also carried out in Instron 5985 electronic universal testing system accompanied with a high and low temperature control system at a loading–unloading strain rate of 2.5 × 10−4 s−1. Extensometer with gauge length of 8 mm and temperature range from −100 °C to 200 °C was used in both tension test and SME test to measure the strain variation. Dog-bone-shaped tensile specimens with gauge length of 15 mm, gauge width of 1.5 mm and thickness of 1.2 mm were electrical-discharge-machined from the strip-shaped samples. The thickness direction of the tensile specimens is perpendicular to Z-axis. In addition to the as-built samples such as AB-130 W, AB-150 W, AB-170 W and AB-190 W (AB represents the as-built samples), analysis was also carried out in the heat-treated samples such as HT-130 W, HT-150 W, HT-170 W and HT-190 W (HT represents the heat-treated samples), which were heated at 900 °C for 2 h in an evacuated quartz tube followed by water quenching.

3. Results

3.1. Microstructure and Phase Composition

Figure 2 depicts the OM images from XZ planes of the as-built samples and the heat-treated samples. As seen in Figure 2a–d, the melt pool boundaries are clearly observed in the as-built samples, which reflect sectorial melt pools, while the grain boundaries are actually not very clear. The depth of the melt pools increases with the increase in the laser power. Previous research has reported that when increasing the energy density, more heat is induced in the surface of the melt pools; therefore, the temperature increases and the reflectivity ratio reduces, which leads to an increase in the absorption ratio, and then the deeper melt pools are obtained [25]. After heat treating, the grain boundaries are clearer and numerous grains along the building direction are observed (Figure 2e–h). In addition, defects are seen in as-built and heat-treated samples. Unmelted powders are observed in the sample AB-130 W and sample HT-130 W, due to relatively low laser power and input energy density, as shown in Figure 2a,e. However, when the laser power increases, a few micropores are observed (Figure 2b,f). The micropores are mostly attributed to gas entrapment [36,37]. With the further increase in laser power, larger pores appear and the number of the pores increases, as seen in Figure 2c,d,g,h. The above defects fabricated by LPBF with different laser powers generally show similar trends as previously reported in the other metallic materials [18,36].
To figure out the microstructure in detail, especially the distribution of the Nb element, SEM was characterized. Figure 3a,d show the BSE image of the as-built sample and heat-treated sample, respectively. The melt pools and grains are clearly observed. Figure 3b,c show a magnified image of the inner melt pool (rectangle A) and melt pool boundary (rectangle B) in Figure 3a. Figure 3e,f show magnified image of inner melt pool (rectangle A) and melt pool boundary (rectangle B) in Figure 3d. The dark-gray areas represent the B2-NiTi matrix phase, the light-white areas represent the Nb-rich phase and the black particles are considered to be the Ti2Ni or Ti4Ni2Ox phase. The above phases have also been identified by XRD (Figure 4). As shown in Figure 3b,c, the shape, size and distribution of the Nb-rich phase are different from traditional Ni47Ti44Nb9 cast alloy [2,38]. Most Nb-rich phases in the as-built sample are arranged in a fine cellular configuration. The size of the cells is hundreds of nanometers, which is far smaller than that in traditional cast alloy. Moreover, there are also some Nb-rich nanoparticles distributed inside the cells. As shown in Figure 3b,c, the size of cells and Nb-rich particles is smaller inside the melt pools than on melt pool boundaries. This may be related to the more rapid cooling rate inside the melt pools, which inhibits grain growth. After heat treating, the cellular structures become discrete, and the number of Nb-rich nanoparticles increases significantly due to the diffusion of Nb atoms.
Figure 4 shows the XRD results of the as-built samples and heat-treated samples that are produced by two representative laser powers (130 W and 190 W); the XRD pattern of the pre-alloyed powders is also given for contrast. The main reflections of all samples match with B2-structured NiTi austenite and Nb-rich phases. However, the as-built samples and heat-treated samples also contain some reflections (marked by arrows and circles) belonging to the fcc-structured Ti2Ni or Ti4Ni2Ox phase, as shown in Figure 4b,c. It has been reported that the Ti2Ni phase is formed directly from liquid, and the formation of Ti4Ni2Ox phase may be related to oxygen pick-up from the build chamber, though the LPBF process is implemented in argon atmosphere [16,39,40]. In addition, as seen in Figure 4b, the XRD patterns of the as-built samples with different laser powers are very similar, indicating that the laser power has tiny effect on the phase composition.

3.2. Phase Transformation and Tensile Properties

Figure 5 shows the transformation behavior of the powder, the as-built samples and the heat-treated samples characterized by differential scanning calorimeter (DSC). As seen in Figure 5a, DSC curves are almost flat for the powder and the as-built samples, indicating that no obvious phase transformation peak was detected. This is very different from the results in traditional Ni47Ti44Nb9 cast alloys, as the traditional cast alloys own four phase transformation peaks [2]. After heat treating, the phase transformation peaks were observed, as shown in Figure 5b, suggesting that martensitic transformation occurs during the cooling process and reverse transformation occurs during the heating process. The four transformation temperatures (martensitic start temperature Ms, martensitic finish temperature Mf, austenitic start temperature As and austenitic finish temperature Af) of the heat-treated samples under different laser powers are shown in Figure 5b. The phase transformation temperatures increase with the increase in laser power.
Figure 6 shows the stress–strain response of the as-built samples and heat-treated samples fabricated under different laser powers. The results are achieved under tension mode at room temperature. As shown in Figure 6a, the as-built samples possess high tensile strength of about 800 MPa and very low total elongation (<1.5%). The tensile strength decreases and the total elongation is not significantly enhanced after heat treating (Figure 6b), and is still far smaller than conventionally fabricated Ni47Ti44Nb9 alloys [3]. One reason for the poor ductility may be related to the existence of Ti2Ni or Ti4Ni2Ox phase, which is very hard and brittle, thus resulting in deterioration of the ductility, as reported in previous studies [27,41]. Another reason may be attributed to the presence of defects (e.g., pores) in the LPBF-fabricated samples [32,42].

3.3. Shape Memory Response

The recovery behavior of the Ni47Ti44Nb9 heat-treated samples was characterized by cryogenic loading–unloading and subsequent heating. Each sample was explored under different tensile strains of 2%, 4%, 6% and 8% until fracture. The deformation temperature (Td) was set as (Ms﹢8 °C). Figure 7 shows the stress–strain–temperature curves of the heat-treated sample HT-130 W. The shape memory rates are 100% when deformed from 2% to 8%. The reverse transformation temperature As’ increases from −10.0 °C to 36.0 °C with the increase in the predeformation strain from 2% to 8%. The increase in the As’ comes from the fact that larger tensile strain tends to bring about more dislocations, thus leading to a higher extent of martensitic stabilization [38,43]. In addition, the total elongation reaches to 8.9% until fracture, which is larger than the total elongation tested at room temperature. This may be attributed to the yielding by stress-induced martensitic transformation (SIM), martensitic reorientation and deformation of oriented martensite when loaded in lower temperature, as reported in previous studies [3,42].
Figure 8a shows the stress–strain–temperature curves (under deformation of 8%) of the heat-treated samples which were fabricated by different laser powers. The cryogenic stress–strain curves are very similar. All the samples exhibit nearly the same phase-transformation plateau, which means the same stress-induced martensitic critical stress σsim. According to Fan’s results, when other conditions are the same, the temperature of ΔT = Td − Ms determines σsim [44]. Therefore, when Td is set as (Ms + 8 °C), the temperature of ΔT is constant, thus resulting in the same σsim, which means the same phase-transformation plateau. On the other hand, the samples exhibit different recovery behavior during the heating. Figure 8b depicts the recovery strain εr and As’ variation with the laser power. The εr slightly decreases from 8% to 7.82% and the As’ increases from 36.0 °C to 52.6 °C with the increase in laser power.

4. Discussions

4.1. Influence of Laser Power and Heat Treatment on Phase-Transformation Behavior

In order to figure out the reason for the change of the phase-transformation temperatures, EDS quantitative analysis was used to characterize the as-built samples and heat-treated samples that were produced by two representative laser powers (130 W and 190 W). The analyzed chemical composition and corresponding Ni/Ti ratio are listed in Table 3. The atomic fraction of Ni decreases and Ti increases with the increase in laser power, leading to a decrease in the Ni/Ti atomic ratio. This may be the main reason for the change in transformation temperatures in the heat-treated samples fabricated under different laser powers. It is well-known that the alloying elements will evaporate from the melt pool during the LPBF process [27]. It is hard for Nb to evaporate, owing to its high boiling point (4744 °C), while Ni has a lower boiling point (2913 °C) compared with Ti (3287 °C) and Ni has the highest equilibrium vapor pressure at elevated temperatures [18,27]. As a result, more Ni loss will occur during the LPBF process and the increase in laser power will lead to a further increase in Ni loss, leading to the decrease in Ni/Ti ratio and thus the increase in transformation temperatures.
It is very interesting in our work that phase transformation did not occur in the as-built samples upon heating and cooling, but happened in the heat-treated samples. It is well-known that the materials undergo a rapid solidification during the LPBF process [28,45]. The nonequilibrium solidification process may result in unique chemical composition, precipitate, high-density dislocation and internal stress, which will influence the phase-transformation behavior markedly. As shown in Table 3 and Figure 5b, the Ms of the heat-treated samples increases by only 26 °C with the decrease of ~0.5 at.% Ni. This means that the change in phase transformation caused by chemical composition is relatively small. Furthermore, the total chemical composition of the heat-treated samples is very close to the as-built samples. Therefore, the change in chemical composition may not be the main reason for the extraordinary change in the phase-transformation temperatures before and after heat treating. On the other hand, previous studies have reported that Nb-rich precipitates suppress martensitic transformation by hampering the growth of martensitic plates, thus resulting in the decrease in Ms [46]. As shown in Figure 3, compared with the the heat-treated samples, the as-built samples own fewer Nb-rich precipitates, which should exhibit higher transformation temperatures. However, no obvious phase-transformation peak was detected in the as-built samples. This indicates that the precipitate may not be the main reason either. To find out the reason, TEM analysis was employed to further characterize the microstructure. As shown in Figure 9, many dislocations were observed in the as-built sample, and the dislocation density decreases a lot in the heat-treated sample. The high-density dislocations can be attributed to the large internal stress during the nonequilibrium solidification process. Previous studies have reported that high-density dislocations and large internal stress can prevent martensitic transformation [27,47]. Therefore, we suggest that the dislocations and internal stress may be the main reason for the disappearance of the phase transformation peak in the as-built samples. Naturally, when the as-built samples are heat-treated, the dislocation densities and internal stress will decrease largely, resulting in the appearance of the phase-transformation peak.

4.2. Influence of Laser Power on εr and As’

As shown in Figure 8, the heat-treated samples fabricated under different laser powers exhibit different SME. The εr decreases with the increase in laser power. It is well-known that the decrease in shape memory effect mainly originates from irreversible slip deformation. Samples with lower dislocation-slip critical stress tend to produce more dislocations or other defects during the predeformation, which can prevent reverse martensitic transformation, resulting in poor recoverability. Figure 10a,b give the magnified OM image of the heat-treated sample fabricated by laser power of 130 W and 190 W, respectively. The sample fabricated with higher laser power tends to have larger average grain size. According to grain-boundary-strengthening theory, a sample with fewer grain boundaries tends to perform lower dislocation-slip critical stress. Thus, the sample fabricated with higher laser power performs lower dislocation-slip critical stress, resulting in lower recoverability. In addition, there are more micropores in the sample fabricated with higher laser power (Figure 10b), which may also contribute to lower dislocation-slip critical stress and lower recoverability.
As shown in Figure 8, the As’ increases with the increase in laser power. It is well-known that the As’ is determined by Ms and transformation hysteresis. On the one hand, the Ms largely increases from −72.2 °C to −45.6 °C with the increase in laser power (Figure 11), and the reason has been discussed in Section 4.1. On the other hand, the transformation hysteresis slightly decreases from 108.2 °C to 98.2 °C with the increase in laser power (Figure 11). A new theory considers that the fundamental cause of transformation hysteresis is owing to crystal symmetry and geometric compatibility of the martensite and the austenite [48]. It predicts that the hysteresis can be minimized by improving the geometric compatibility between the two phases, which has already been confirmed in NiTiCu alloys with narrow hysteresis. The theory specifies two conditions in order for an extremely small hysteresis. The first one is detU = λ1λ2λ3 = 1, which means that there is no volume change due to transformation. The second one is λ2 = 1, which means that the austenite is directly compatible with a single variant of martensite. According to the research, the only inputs to the theory are lattice parameters which can be tuned by compositions. Therefore, we can speculate that the decrease in the transformation hysteresis with the increase in laser power is related to the variation of chemical composition—maybe the decrease in Ni/Ti ratio. Of course, detailed studies about the relation between the chemical composition and lattice parameters of both phases are still required to figure out the mechanism in future. In general, it is obvious that the increase in Ms overcomes the decrease in transformation hysteresis, thus resulting in the increase in As’ with the increase in laser power.

5. Conclusions

(1)
Ni47Ti44Nb9 alloys were successfully fabricated via LPBF technique with different laser powers. The as-built alloys have good shaping properties, though a few defects were discovered. The unmelted powders are owing to relatively low laser power and input energy density, while the micropores are owing to gas entrapment under high laser power.
(2)
There is no phase-transformation peak observed in the as-built samples, which may be due to the high dislocation density and internal stress caused by the nonequilibrium solidification process in LPBF. After heat treating, the phase-transformation peaks appear, and the phase-transformation temperatures increase with the increase in laser power, which may be due to the decrease in Ni/Ti ratio.
(3)
The tension test at room temperature indicates that the LPBF samples exhibit poor tensile ductility with total elongation below 1.5%, which may be related to the existence of the Ti2Ni or Ti4Ni2Ox phase or LPBF pores. However, the LPBF samples after heat treating still possess good shape memory effect (εr = 7.82–8%) and relatively high reverse transformation temperature (As’ = 36.0–52.6 °C) when deformed to 8%. In addition, the εr decreases and As’ increases with the increase in laser power.

Author Contributions

Conceptualization, M.S.; methodology, M.S., J.C., Q.F., C.Y., G.W. and X.S.; formal analysis, M.S., J.C., Q.F. and Y.W.; investigation, M.S. and Y.W.; resources, G.W. and X.S.; data curation, M.S., J.C., C.Y. and Y.W.; writing—original draft preparation, M.S.; writing—review and editing, Y.Z. and S.H.; supervision, Y.Z. and S.H.; project administration, Y.Z. and S.H.; funding acquisition, J.C., Q.F., C.Y., G.W. and X.S. All authors have read and agreed to the published version of the manuscript.

Funding

This work was financially supported by National Safety Academic Fund (Grant No. U1930207) and National Natural Science Foundation of China (Grants No. 51901214, No. 52001289, No. 52001290, No. 52105410, No. 52101057 and No. 52101158).

Data Availability Statement

Not applicable.

Conflicts of Interest

The authors declare no conflict of interest.

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Figure 1. (a) Secondary electron image of NiTiNb pre-alloyed powder. (b) Particle size distribution of NiTiNb pre-alloyed powder. (c) The as-built NiTiNb samples fabricated by laser powder bed fusion (LPBF).
Figure 1. (a) Secondary electron image of NiTiNb pre-alloyed powder. (b) Particle size distribution of NiTiNb pre-alloyed powder. (c) The as-built NiTiNb samples fabricated by laser powder bed fusion (LPBF).
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Figure 2. Optical microscopy (OM) images of (a–d) the as-built samples fabricated by LPBF with laser power of 130W-190W and (e,f) heat-treated samples of (a–d); the rectangles and circles indicate the unmelted powders and pores, respectively.
Figure 2. Optical microscopy (OM) images of (a–d) the as-built samples fabricated by LPBF with laser power of 130W-190W and (e,f) heat-treated samples of (a–d); the rectangles and circles indicate the unmelted powders and pores, respectively.
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Figure 3. Backscattered electron (BSE) images of (a) the as-built sample fabricated under laser power of 130 W; (b) magnified image of rectangle A in (a); (c) magnified image of rectangle B in (a); (d) the heat-treated sample fabricated under laser power of 130 W; (e) magnified image of rectangle A in (d); (f) magnified image of rectangle B in (d). Rectangle A represents inner melt pool, rectangle B represents melt pool boundary.
Figure 3. Backscattered electron (BSE) images of (a) the as-built sample fabricated under laser power of 130 W; (b) magnified image of rectangle A in (a); (c) magnified image of rectangle B in (a); (d) the heat-treated sample fabricated under laser power of 130 W; (e) magnified image of rectangle A in (d); (f) magnified image of rectangle B in (d). Rectangle A represents inner melt pool, rectangle B represents melt pool boundary.
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Figure 4. X-ray diffractometer (XRD) patterns for (a) Ni47Ti44Nb9 pre-alloyed powder, (b) as-built samples and (c) heat-treated samples.
Figure 4. X-ray diffractometer (XRD) patterns for (a) Ni47Ti44Nb9 pre-alloyed powder, (b) as-built samples and (c) heat-treated samples.
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Figure 5. Differential scanning calorimeter (DSC) curves of the Ni47Ti44Nb9 (a) powder and the as-built samples; (b) the heat-treated samples.
Figure 5. Differential scanning calorimeter (DSC) curves of the Ni47Ti44Nb9 (a) powder and the as-built samples; (b) the heat-treated samples.
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Figure 6. Stress–strain curves of (a) the as-built samples and (b) the heat-treated samples. The arrow represents the moment of removing the extensometer.
Figure 6. Stress–strain curves of (a) the as-built samples and (b) the heat-treated samples. The arrow represents the moment of removing the extensometer.
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Figure 7. Stress–strain–temperature curves obtained by tensile loading–unloading–heating process of the heat-treated samples fabricated by laser power of 130 W.
Figure 7. Stress–strain–temperature curves obtained by tensile loading–unloading–heating process of the heat-treated samples fabricated by laser power of 130 W.
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Figure 8. (a) Stress–strain–temperature curves of the heat-treated samples fabricated by different laser powers; (b) the εr and As’ variation with the laser power.
Figure 8. (a) Stress–strain–temperature curves of the heat-treated samples fabricated by different laser powers; (b) the εr and As’ variation with the laser power.
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Figure 9. Transmission electron microscope (TEM) images of (a) as-built sample and (b) heat-treated sample. The arrows represent the dislocations.
Figure 9. Transmission electron microscope (TEM) images of (a) as-built sample and (b) heat-treated sample. The arrows represent the dislocations.
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Figure 10. OM images of the heat-treated sample fabricated by LPBF with laser power of (a) 130 W, (b) 190 W.
Figure 10. OM images of the heat-treated sample fabricated by LPBF with laser power of (a) 130 W, (b) 190 W.
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Figure 11. Ms and transformation hysteresis width variation with the laser power.
Figure 11. Ms and transformation hysteresis width variation with the laser power.
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Table
Table 1. Chemical compositions of the NiTiNb rod and pre-alloyed powder.
Table 1. Chemical compositions of the NiTiNb rod and pre-alloyed powder.
SampleNi (wt.%)Ti (wt.%)Nb (wt.%)C (wt.%)H (wt.%)O (wt.%)
rod48.32Bal.14.690.0030<0.00100.017
powder48.03Bal.14.910.0036<0.00100.028
Table
Table 2. The LPBF process parameters and calculated energy density; AB represents the as-built samples.
Table 2. The LPBF process parameters and calculated energy density; AB represents the as-built samples.
SampleLaser Power
P (W)		Scanning Velocity
v (mm/s)		Hatch Spacing
h (μm)		Layer Thickness
t (μm)		Energy Density
Eν (J/mm3)
AB-130 W1307005030123.8
AB-150 W1507005030142.9
AB-170 W1707005030161.9
AB-190 W1907005030181.0
Table
Table 3. The analyzed chemical composition and corresponding Ni/Ti ratio of the as-built samples and heat-treated samples.
Table 3. The analyzed chemical composition and corresponding Ni/Ti ratio of the as-built samples and heat-treated samples.
SampleNi (wt.%)Ti (wt.%)Nb (wt.%)Ni (at.%)Ti (at.%)Nb (at.%)Ni/Ti
AB-130 W48.9836.5514.4747.5743.538.911.094
AB-190 W48.3536.8114.8446.9943.879.141.071
HT-130 W48.9136.6114.4847.5043.598.911.089
HT-190 W48.4937.0514.4647.0444.078.891.068
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Sun, M.; Chen, J.; Fan, Q.; Yang, C.; Wang, G.; Shen, X.; Wang, Y.; Zhang, Y.; Huang, S. Transformation Behavior and Shape Memory Effect of Ni47Ti44Nb9 Alloy Synthesized by Laser Powder Bed Fusion and Heat Treating. Metals 2022, 12, 1438. https://doi.org/10.3390/met12091438

AMA Style

Sun M, Chen J, Fan Q, Yang C, Wang G, Shen X, Wang Y, Zhang Y, Huang S. Transformation Behavior and Shape Memory Effect of Ni47Ti44Nb9 Alloy Synthesized by Laser Powder Bed Fusion and Heat Treating. Metals. 2022; 12(9):1438. https://doi.org/10.3390/met12091438

Chicago/Turabian Style

Sun, Mingyan, Jie Chen, Qichao Fan, Chuan Yang, Guowei Wang, Xianfeng Shen, Yangyang Wang, Yonghao Zhang, and Shuke Huang. 2022. "Transformation Behavior and Shape Memory Effect of Ni47Ti44Nb9 Alloy Synthesized by Laser Powder Bed Fusion and Heat Treating" Metals 12, no. 9: 1438. https://doi.org/10.3390/met12091438

APA Style

Sun, M., Chen, J., Fan, Q., Yang, C., Wang, G., Shen, X., Wang, Y., Zhang, Y., & Huang, S. (2022). Transformation Behavior and Shape Memory Effect of Ni47Ti44Nb9 Alloy Synthesized by Laser Powder Bed Fusion and Heat Treating. Metals, 12(9), 1438. https://doi.org/10.3390/met12091438

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