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Submicron-porous NiTi and NiTiNb shape memory alloys with high damping capacity fabricated by a new top-down process
Materials and Design 78 (2015) 74–79
Contents lists available at ScienceDirect
Materials and Design journal homepage: www.elsevier.com/locate/matdes
Submicron-porous NiTi and NiTiNb shape memory alloys with high damping capacity fabricated by a new top-down process Wei Guo a, Hidemi Kato b,⇑ a b
Department of Materials Science, Graduate School of Engineering, Tohoku University, Sendai 980-8579, Japan Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan
a r t i c l e
i n f o
Article history: Received 11 December 2014 Revised 14 April 2015 Accepted 18 April 2015 Available online 18 April 2015 Keywords: Porous shape memory alloys Fine pore structure High damping capacity Top-down processing
a b s t r a c t In this paper, submicron-porous NiTi and NiTiNb shape memory alloys were firstly fabricated by a new top-down process, including solidified phase controlling and etching process. These porous alloys possessed fine pore size (0.4 lm) and homogeneous pore distribution. The mechanical properties of these porous alloys were similar to those prepared by conventional powder metallurgy techniques or bottomup process. However, the damping capacity of the submicron-porous NiTi was 0.07 and 40% higher than the counterparts produced by the conventional powder metallurgy techniques. The damping capacity of the submicron-porous NiTiNb was 0.18 and 50% higher than its dense counterpart. The internal friction value at austenite state for submicron-porous NiTiNb was 0.04 and 200% higher than its dense counterpart, 300% higher than submicron-porous NiTi. The high damping capacity of these submicron-porous shape memory alloys was considered to be caused partly by their fine pore structures, which enhanced the stress concentration and the microplastic deformation or collapse of thin nodes. These porous shape memory alloys exhibited moderate stiffness as well as high damping capacity and high specific strength, making them good candidate as practical engineering damping materials. Ó 2015 Elsevier Ltd. All rights reserved.
1. Introduction For applications including energy absorption devices or structural materials used under cyclic and dynamic stress require both high damping capacity and good mechanical properties. However, traditional damping materials such as rubbers and polymers with high damping capacity have low mechanical strength, and those with high mechanical strength have low damping capacity [1–5]. NiTi shape memory alloys (SMAs) have attracted much attention as damping materials in buildings, bridges, and high-speed trains. Compared with other damping materials, NiTi SMAs have superior mechanical properties—including their specific strength (300 kN m/kg), fracture toughness (26 MPa m1/2), corrosion resistance, as well as high damping capacity [6–10]. The damping capacity of NiTi alloys can be further improved by making them porous. However, for the high melting point of NiTi (1310 °C), to date, porous NiTi can be only produced by powder metallurgy techniques (bottom-up process), such as self-propagating high-temperature synthesis (SHS), spark plasma sintering (SPS),
⇑ Corresponding author. E-mail address: [email protected] (H. Kato). http://dx.doi.org/10.1016/j.matdes.2015.04.028 0261-3069/Ó 2015 Elsevier Ltd. All rights reserved.
hot isostatic pressing (HIP) with argon expansion, capsule-free HIP (CF-HIP), and conventional sintering (CS) [11–18]. The porous NiTi produced by these methods possess large pore sizes from several tens to several hundreds of micrometers, and their pore distributions are not homogenous. However, finer pore sizes and more homogenous pore distribution would be beneficial. Liu et al. found that the damping capacity of foamed Al increased with decreasing pore size [19]. Yuan et al. used three different methods to produce porous NiTi with various porosities and pore sizes. They found that decreasing the pore size could improve the damping capacity and concluded that improvements in damping capacity depended more on pore size than porosity [20]. These results suggest that, to further improve the damping capacity of porous NiTi, fine and homogenously distributed pore structures are desired. Beyond NiTi SMAs, NiTiNb ternary SMAs have shown good damping properties [21,22]. Ni47Ti44Nb9 has exhibited a very high damping peak value of 0.12 during cooling [23], possibly caused by the creation or motion of interphase boundaries. The NiTiNb alloys are characterized by an ordered NiTi(Nb) matrix containing a fine dispersion of b-Nb particles [24]. The interfaces between the b-Nb particles and the matrix provide nucleus sites for the martensite, generating increased amount of boundaries between the martensite and parent phase, which might greatly benefit the
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damping capacity during the phase transformation [22]. However, the damping capacity of NiTiNb could be further improved by foaming, based on the works on porous NiTi, but porous NiTiNb has not been successfully produced yet. Generally, porous materials can be fabricated not only by powder metallurgy techniques (bottom-up process), but also by etching techniques (top-down process). In etching techniques, a precursor containing two or more phases is fabricated and etched, preferentially removing certain phases, leaving a porous structure [25,26]. In fact, etching techniques are also applied in fabricating porous NiTi by powder metallurgy process. The space-holder materials such as NaCl are used to sinter the alloys and finally etched out to obtain porous structures. But the space-holder materials are limited to a large particle size, generating large pores. In the present research, a novel top-down process was designed to produce submicron-porous NiTi and submicron-porous NiTiNb. At a proper composition, a multiphase Ni–Ti–Gd precursor containing both NiTi and Ni–Gd phases can be prepared according to their phase diagrams [27]. Then, by etching out the Ni–Gd phase with an acid solution, porous NiTi may be produced. For Ni–Gd phase is in situ formed during Ni–Ti–Gd preparation, the final porous specimens may possess fine pore size and homogenous pore distribution (the size and distribution of dissolved Ni–Gd phase). Furthermore, Nb and Gd are immiscible [27]; thus, by adding Nb to Ni–Ti–Gd alloy, it may be possible to fabricate Ni–Ti–Nb–Gd precursor consisting of NiTiNb and Ni–Gd phases, then similarly by etching out Ni–Gd phase, porous NiTiNb may be produced. Ti-rich NiTi phase rather than Ni-rich NiTi was selected because the damping peaks for Ni-rich NiTi show at rather low temperatures (i.e., below room temperature), making it not useful for increasing the transformation temperature. Additionally, few studies on the damping properties of Ti-rich porous NiTi alloys have been reported [28]. The composition of NiTiNb phase is determined to be Ni47Ti44Nb9 for its good damping properties. After careful adjustment of composition, the precursors are determined to be Ni48Ti44Gd8 and Ni47.5Ti26Nb6.5Gd20, respectively, to obtain the expected Ti-rich NiTi and NiTiNb phases. After preparing the porous specimens, the microstructures, mechanical properties, and damping capacities were investigated in details.
2. Experimental The Ni–Ti–Gd and Ni–Ti–Nb–Gd precursors were prepared by arc melting the constituent elements in a Ti-gettered Ar atmosphere. Then, both rod-shaped precursors (3 mm in diameter) and plate-shaped precursors (35 5 1 mm3) were fabricated by copper mold casting because our previous work showed that higher cooling rate can generate finer and more homogenous microstructures [29]. For Ni–Ti–Nb–Gd rod and plate precursors, they were annealed at 850 °C for 1 h in quartz capsules, evacuated to 0.04 MPa, and then cooled in the furnace. Finally, all precursors were immersed in nitric acid solution (HNO3:H2O = 2:1) for 12 h. The structures were examined by X-ray diffraction (XRD; Bruker D8 Advance) with Cu Ka radiation and a scanning electron microscope with an attached energy-dispersive X-ray spectrometer (SEM-EDX; Carl Zeiss Ultra 55 with Bruker AXS). The pore characteristics including average pore size and porosity were evaluated with a mercury porosimeter (Micromeritics, AutoPore IV). Uniaxial compression tests (Instron 4204) were performed at a strain rate of 1 104 s1 at room temperature. The damping tests (RSA-G2, TA Instruments) were conducted in single-cantilever mode at a constant strain amplitude of 3 104 and a frequency of 1 Hz, from 40 °C to 100 °C for porous NiTi and 120 °C to 20 °C for porous NiTiNb, with constant heating and cooling rates of 5 °C/min.
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3. Results and discussion 3.1. Microstructures and properties of submicron-porous NiTi Fig. 1 shows the plate-shaped Ni–Ti–Gd precursor before etching and porous NiTi after etching, indicating that bulk specimens were successfully produced by the new top-down process. Fig. 2(a) shows the XRD patterns of Ni–Ti–Gd precursor, revealing B2-NiTi, B190 -NiTi, and NiGd phases, without other Ni–Ti compounds such as Ti2Ni, which was usually formed during sintering in powder metallurgy process. Fig. 2(b) and (c) shows SEM images of the precursor at different magnifications, revealing two different phases: the darker phase was Ni49Ti51 and the brighter phase was NiGd according to EDX results (not shown here). The volume fraction of NiGd was 27% according to the image analysis of Fig. 2(b), indicating that the porosity of porous NiTi after etching out NiGd would be 27%. Fig. 2(d)–(f) shows EDX elemental mapping taken from Fig. 2(c), also proving that NiTi and NiGd phases were formed in the Ni–Ti–Gd precursors. Fig. 3(a) shows XRD patterns of porous NiTi after etching in the nitric acid solution, revealing no NiGd peaks because of its dissolution and the main phase was B2-NiTi rather than B190 -NiTi for its much higher peak intensity. Fig. 3(b) and (c) shows SEM images of the porous NiTi, exhibiting homogenous distributed fine pores. The detailed pore characteristics such as average pore size and porosity were evaluated by a mercury porosimeter. As shown in Fig. 3(d), the average pore size was 0.39 lm and the porosity was 25%. This submicron pore size was much finer than those produced by conventional powder metallurgy techniques [30–32]. The porosity agreed well with the image analysis results from Ni–Ti–Gd precursor shown before. Fig. 4(a) shows compressive stress–strain curve for the submicron-porous NiTi, exhibiting the Young’s modulus of 4.8 GPa, ultimate stress of 250 MPa, and fracture strain of 7.9%. The detailed mechanical behaviors were: at first, the elastic deformation from bending or compressing the cell walls elastically contributed to the linear part. Then, buckling and cell collapsing contributed to the serration part and each serration corresponded to the fracture of the foam cell walls. Finally, the material failed [33,34]. Gibson et al. proposed a formula to calculate the crushing strength (rcr) of porous materials, which is related to the stress during serrated behavior; this formula is shown as follows for elastic-brittle open-cell foams:
Fig. 1. Photograph of plate-shaped Ni–Ti–Gd precursor before etching (top) and porous specimens after etching (down).
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Fig. 2. (a) XRD patterns and (b and c) SEM images of Ni–Ti–Gd precursor; (d–f) EDX elemental mapping of (d) Ni, (e) Ti, and (f) Gd.
Fig. 3. (a) XRD patterns and (b and c) SEM images of submicron-porous NiTi and (d) pore-size distribution of submicron-porous NiTi by mercury porosimeter, the red dot line is fitting curves of the experimental data. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
rcr q 0:65 qs rfs
3=2 ;
where rcr is the crushing stress of the foam, rfs is the tensile fracture strength for the dense counterpart, q is the foam density, and qs is the density of the dense counterpart [35]. In the present study, the density of dense and porous NiTi is calculated by dividing the mass by volume, finding that qq is 0.65. And rfs is 800 MPa s
from Ref. [9], so rcr is calculated to be 270 MPa, similar to the experimental data, 250 MPa, suggesting that elastic-brittle foam model can be applied to the present submicron-porous NiTi alloy.
Fig. 5 shows the internal friction (IF) and storage modulus as functions of the temperature for submicron-porous NiTi, measured in a single-cantilever mode. These results reveal damping peaks during heating and cooling, correlated to the austenite–martensite phase transformations. The submicron-porous NiTi exhibited a very high damping peak (0.07), 40% higher than those produced by conventional powder metallurgy process [36,37]. The high damping capacity of porous NiTi usually comes from the following factors: (1) stress concentration causes elastic and plastic bending, buckling, and even fracture of thin nodes or walls [38,39]; (2) reversible dislocation motion in stress–strain localization areas [38,39] and stress-assisted twin-boundary motion in martensite
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Fig. 4. Compressive stress–strain curves of (a) submicron-porous NiTi and (b) submicron-porous NiTiNb.
Fig. 5. Damping capacity and storage modulus as functions of temperature for submicron-porous NiTi.
platelets [40,41]; and (3) stress-induced martensite formation when loaded in the austenite state. For the present submicron-porous NiTi, the fine pore structure increases the interface areas between the NiTi and pores, enhancing the stress concentration and stress–strain localization areas. This behavior enhances factors (1) and (2), generating improved damping capacity even though its porosity is not very high. Ashby et al. summarized the relations between the loss coefficient (damping property) and Young’s modulus (mechanical property) for the most common damping materials at 30 °C [42]. As shown in Fig. 6, polymers possess high damping capacity but low Young’s modulus, while metals such as Mg alloys or Al alloys possess high Young’s modulus but low damping capacity. However, for engineering application, both high damping capacity and high Young’s modulus are desired. The present submicron-porous NiTi was marked with a triangle symbol and the counterpart fabricated by powder metallurgy process [36] was marked with a cycle symbol in Fig. 6. Both the Young’s modulus and damping capacity were higher for the present submicron-porous NiTi. Lead alloys also exhibit superior damping properties, but the toxicity and unfavorable high specific gravity are their drawbacks [43]. 3.2. Microstructures and properties of submicron-porous NiTiNb Fig. 7(a) shows the XRD patterns of Ni–Ti–Nb–Gd precursor and the porous Ni–Ti–Nb after etching, exhibiting B2-NiTi(Nb) and b-Nb phases for both specimens. Fig. 7(b) shows the SEM images
Fig. 6. Relations between the loss coefficient and Young’s modulus for commonly used damping materials at 30 °C [42]. The submicron-porous NiTi by a new topdown process is marked with a triangle symbol, and the counterpart fabricated by a conventional powder metallurgy techniques (bottom-up process) [37] is marked with a cycle symbol.
of the Ni–Ti–Nb–Gd precursor, revealing a darker phase (Ni47Ti44Nb9) and a brighter phase (NiGd). Fig. 7(c)–(f) shows EDX elemental mapping taken from Fig. 7(b), revealing b-Nb precipitates in Ni–Ti–Nb phase. The volume fraction of NiGd phase was 31% based on the image analysis of Fig. 7(b), indicating the porosity for porous NiTiNb after etching would be 31%. Fig. 8(a) shows the SEM image of porous NiTiNb, revealing homogenous submicron pores, with a NiTi(Nb) matrix and b-Nb precipitates. These results are supported by EDX elemental mapping of Nb, as shown in Fig. 8(b). The average pore size and porosity of the submicron-porous NiTiNb were 0.42 lm and 33% by mercury porosimeter, respectively, as shown in Fig. 8(c). Fig. 4(b) shows the compressive stress–strain curve for porous NiTiNb, exhibiting the Young’s modulus of 5 GPa, ultimate stress of 285 MPa, and fracture stain of 10.7%. The curve has three distinct parts: a linear elastic part in which the cell walls bended or compressed elastically, an increasingly steep yield part in which the cells collapsed by buckling and plastic yielding, and then the maximum stress followed by fracture failure. This behavior is similar to cancellous bone in compression, where a plateau of almost constant stress appears after the linear elastic part [44]. It should be pointed out that unlike the elastic stress–strain curve of submicron-porous NiTi before fracture as shown in Fig. 4(a), the
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Fig. 7. (a) XRD patterns and (b) SEM images of the Ni–Ti–Nb–Gd precursor; (c–f) EDX mapping taken from (b) for (c) Ni, (d) Ti, (e) Nb, and (f) Gd.
Fig. 8. (a) SEM images of submicron-porous NiTiNb, (b) elemental mapping of Nb taken from (a), and (c) pore-size distribution by mercury porosimeter, the red dot line is fitting curves of the experimental data. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
submicron-porous NiTiNb shows yielding phenomenon at about 160 MPa before fracture. This different behavior may be generated from the different microstructures. As shown before in Fig. 8(a) and (b), very fine ductile b-Nb phases precipitate inside and between the cell walls, and these b-Nb phases are considered to prevent the brittle collapse of thin nodes during deformation and to generate a yielding behavior. Moreover, from Ref. [35], the yield strength for elastic–plastic foams can be calculated as follows:
r q 0:3 qs rys
3=2 ;
where r is the yield stress of the foam, rys is the yield stress of the dense counterpart, q is the density of the foam, and qs is the density of the dense counterpart. Dividing the mass by volume, qq is found s
to be 0.6. And rys is 1000 MPa [45], so the calculated yield stress of the present submicron-porous NiTiNb is 140 MPa, similar to the experimental data, 160 MPa, indicating that the elastic–plastic foam model is suitable for the present submicron-porous NiTiNb alloy. Fig. 9 shows the temperature dependence of the damping capacity and storage modulus for the submicron-porous NiTiNb alloy. One damping capacity peak appeared during cooling, caused by the martensitic transformation. But before that, the damping
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(4) These new submicron-porous shape memory alloys are considered to be good candidate as engineering damping materials and this new top-down process can be used to produce various porous alloys with fine pore structures.
References
Fig. 9. Damping capacity and storage modulus as functions of temperature for submicron-porous NiTiNb.
capacity remained nearly constant over a wide temperature range. The maximum internal friction value and that at the austenite state were 0.18 and 0.04, 50% and 200% higher than those of dense NiTiNb, respectively [22]. The high damping capacity of submicron-porous NiTiNb was caused by both the composite microstructure, containing the B2-NiTi(Nb) matrix and b-Nb particles, and the porous structure as well. The interfaces between the b-Nb particles and NiTi(Nb) matrix provided nucleus sites for the martensite variants, increasing the number of boundaries between the martensite and parent phase; this behavior increased the energy dissipation, caused by movement of the martensite/parent interfaces, similar to dense NiTiNb alloy. The pore structures improved the stress concentration and stress–strain localization areas, enhancing the microplastic deformation, crack initiation, and dislocation motion, which could consume more energy, similar to porous NiTi alloy. The submicron-porous NiTiNb also exhibited high and steady internal friction value at the austenite state because of strain localization in the fine nodes and struts.
4. Conclusion (1) The submicron-porous NiTi and NiTiNb with homogeneously distributed fine pore structures have been firstly fabricated by a new top-down process, including solidified phase controlling and etching process. (2) The submicron-porous NiTi possessed much finer pore size (0.39 lm) and more homogeneous pore distribution than the counterparts by conventional powder metallurgy techniques (bottom-up process). It also showed similar mechanical properties to the counterparts, but 40% higher damping capacity (0.07) than those counterparts. (3) The submicron-porous NiTiNb also possessed fine pore size (0.5 lm). Its damping capacity was 0.18 and 50% higher than its dense counterpart. The internal friction value at austenite state was 0.04 and 200% higher than its dense counterpart.
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Contents lists available at ScienceDirect
Materials and Design journal homepage: www.elsevier.com/locate/matdes
Submicron-porous NiTi and NiTiNb shape memory alloys with high damping capacity fabricated by a new top-down process Wei Guo a, Hidemi Kato b,⇑ a b
Department of Materials Science, Graduate School of Engineering, Tohoku University, Sendai 980-8579, Japan Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan
a r t i c l e
i n f o
Article history: Received 11 December 2014 Revised 14 April 2015 Accepted 18 April 2015 Available online 18 April 2015 Keywords: Porous shape memory alloys Fine pore structure High damping capacity Top-down processing
a b s t r a c t In this paper, submicron-porous NiTi and NiTiNb shape memory alloys were firstly fabricated by a new top-down process, including solidified phase controlling and etching process. These porous alloys possessed fine pore size (0.4 lm) and homogeneous pore distribution. The mechanical properties of these porous alloys were similar to those prepared by conventional powder metallurgy techniques or bottomup process. However, the damping capacity of the submicron-porous NiTi was 0.07 and 40% higher than the counterparts produced by the conventional powder metallurgy techniques. The damping capacity of the submicron-porous NiTiNb was 0.18 and 50% higher than its dense counterpart. The internal friction value at austenite state for submicron-porous NiTiNb was 0.04 and 200% higher than its dense counterpart, 300% higher than submicron-porous NiTi. The high damping capacity of these submicron-porous shape memory alloys was considered to be caused partly by their fine pore structures, which enhanced the stress concentration and the microplastic deformation or collapse of thin nodes. These porous shape memory alloys exhibited moderate stiffness as well as high damping capacity and high specific strength, making them good candidate as practical engineering damping materials. Ó 2015 Elsevier Ltd. All rights reserved.
1. Introduction For applications including energy absorption devices or structural materials used under cyclic and dynamic stress require both high damping capacity and good mechanical properties. However, traditional damping materials such as rubbers and polymers with high damping capacity have low mechanical strength, and those with high mechanical strength have low damping capacity [1–5]. NiTi shape memory alloys (SMAs) have attracted much attention as damping materials in buildings, bridges, and high-speed trains. Compared with other damping materials, NiTi SMAs have superior mechanical properties—including their specific strength (300 kN m/kg), fracture toughness (26 MPa m1/2), corrosion resistance, as well as high damping capacity [6–10]. The damping capacity of NiTi alloys can be further improved by making them porous. However, for the high melting point of NiTi (1310 °C), to date, porous NiTi can be only produced by powder metallurgy techniques (bottom-up process), such as self-propagating high-temperature synthesis (SHS), spark plasma sintering (SPS),
⇑ Corresponding author. E-mail address: [email protected] (H. Kato). http://dx.doi.org/10.1016/j.matdes.2015.04.028 0261-3069/Ó 2015 Elsevier Ltd. All rights reserved.
hot isostatic pressing (HIP) with argon expansion, capsule-free HIP (CF-HIP), and conventional sintering (CS) [11–18]. The porous NiTi produced by these methods possess large pore sizes from several tens to several hundreds of micrometers, and their pore distributions are not homogenous. However, finer pore sizes and more homogenous pore distribution would be beneficial. Liu et al. found that the damping capacity of foamed Al increased with decreasing pore size [19]. Yuan et al. used three different methods to produce porous NiTi with various porosities and pore sizes. They found that decreasing the pore size could improve the damping capacity and concluded that improvements in damping capacity depended more on pore size than porosity [20]. These results suggest that, to further improve the damping capacity of porous NiTi, fine and homogenously distributed pore structures are desired. Beyond NiTi SMAs, NiTiNb ternary SMAs have shown good damping properties [21,22]. Ni47Ti44Nb9 has exhibited a very high damping peak value of 0.12 during cooling [23], possibly caused by the creation or motion of interphase boundaries. The NiTiNb alloys are characterized by an ordered NiTi(Nb) matrix containing a fine dispersion of b-Nb particles [24]. The interfaces between the b-Nb particles and the matrix provide nucleus sites for the martensite, generating increased amount of boundaries between the martensite and parent phase, which might greatly benefit the
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damping capacity during the phase transformation [22]. However, the damping capacity of NiTiNb could be further improved by foaming, based on the works on porous NiTi, but porous NiTiNb has not been successfully produced yet. Generally, porous materials can be fabricated not only by powder metallurgy techniques (bottom-up process), but also by etching techniques (top-down process). In etching techniques, a precursor containing two or more phases is fabricated and etched, preferentially removing certain phases, leaving a porous structure [25,26]. In fact, etching techniques are also applied in fabricating porous NiTi by powder metallurgy process. The space-holder materials such as NaCl are used to sinter the alloys and finally etched out to obtain porous structures. But the space-holder materials are limited to a large particle size, generating large pores. In the present research, a novel top-down process was designed to produce submicron-porous NiTi and submicron-porous NiTiNb. At a proper composition, a multiphase Ni–Ti–Gd precursor containing both NiTi and Ni–Gd phases can be prepared according to their phase diagrams [27]. Then, by etching out the Ni–Gd phase with an acid solution, porous NiTi may be produced. For Ni–Gd phase is in situ formed during Ni–Ti–Gd preparation, the final porous specimens may possess fine pore size and homogenous pore distribution (the size and distribution of dissolved Ni–Gd phase). Furthermore, Nb and Gd are immiscible [27]; thus, by adding Nb to Ni–Ti–Gd alloy, it may be possible to fabricate Ni–Ti–Nb–Gd precursor consisting of NiTiNb and Ni–Gd phases, then similarly by etching out Ni–Gd phase, porous NiTiNb may be produced. Ti-rich NiTi phase rather than Ni-rich NiTi was selected because the damping peaks for Ni-rich NiTi show at rather low temperatures (i.e., below room temperature), making it not useful for increasing the transformation temperature. Additionally, few studies on the damping properties of Ti-rich porous NiTi alloys have been reported [28]. The composition of NiTiNb phase is determined to be Ni47Ti44Nb9 for its good damping properties. After careful adjustment of composition, the precursors are determined to be Ni48Ti44Gd8 and Ni47.5Ti26Nb6.5Gd20, respectively, to obtain the expected Ti-rich NiTi and NiTiNb phases. After preparing the porous specimens, the microstructures, mechanical properties, and damping capacities were investigated in details.
2. Experimental The Ni–Ti–Gd and Ni–Ti–Nb–Gd precursors were prepared by arc melting the constituent elements in a Ti-gettered Ar atmosphere. Then, both rod-shaped precursors (3 mm in diameter) and plate-shaped precursors (35 5 1 mm3) were fabricated by copper mold casting because our previous work showed that higher cooling rate can generate finer and more homogenous microstructures [29]. For Ni–Ti–Nb–Gd rod and plate precursors, they were annealed at 850 °C for 1 h in quartz capsules, evacuated to 0.04 MPa, and then cooled in the furnace. Finally, all precursors were immersed in nitric acid solution (HNO3:H2O = 2:1) for 12 h. The structures were examined by X-ray diffraction (XRD; Bruker D8 Advance) with Cu Ka radiation and a scanning electron microscope with an attached energy-dispersive X-ray spectrometer (SEM-EDX; Carl Zeiss Ultra 55 with Bruker AXS). The pore characteristics including average pore size and porosity were evaluated with a mercury porosimeter (Micromeritics, AutoPore IV). Uniaxial compression tests (Instron 4204) were performed at a strain rate of 1 104 s1 at room temperature. The damping tests (RSA-G2, TA Instruments) were conducted in single-cantilever mode at a constant strain amplitude of 3 104 and a frequency of 1 Hz, from 40 °C to 100 °C for porous NiTi and 120 °C to 20 °C for porous NiTiNb, with constant heating and cooling rates of 5 °C/min.
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3. Results and discussion 3.1. Microstructures and properties of submicron-porous NiTi Fig. 1 shows the plate-shaped Ni–Ti–Gd precursor before etching and porous NiTi after etching, indicating that bulk specimens were successfully produced by the new top-down process. Fig. 2(a) shows the XRD patterns of Ni–Ti–Gd precursor, revealing B2-NiTi, B190 -NiTi, and NiGd phases, without other Ni–Ti compounds such as Ti2Ni, which was usually formed during sintering in powder metallurgy process. Fig. 2(b) and (c) shows SEM images of the precursor at different magnifications, revealing two different phases: the darker phase was Ni49Ti51 and the brighter phase was NiGd according to EDX results (not shown here). The volume fraction of NiGd was 27% according to the image analysis of Fig. 2(b), indicating that the porosity of porous NiTi after etching out NiGd would be 27%. Fig. 2(d)–(f) shows EDX elemental mapping taken from Fig. 2(c), also proving that NiTi and NiGd phases were formed in the Ni–Ti–Gd precursors. Fig. 3(a) shows XRD patterns of porous NiTi after etching in the nitric acid solution, revealing no NiGd peaks because of its dissolution and the main phase was B2-NiTi rather than B190 -NiTi for its much higher peak intensity. Fig. 3(b) and (c) shows SEM images of the porous NiTi, exhibiting homogenous distributed fine pores. The detailed pore characteristics such as average pore size and porosity were evaluated by a mercury porosimeter. As shown in Fig. 3(d), the average pore size was 0.39 lm and the porosity was 25%. This submicron pore size was much finer than those produced by conventional powder metallurgy techniques [30–32]. The porosity agreed well with the image analysis results from Ni–Ti–Gd precursor shown before. Fig. 4(a) shows compressive stress–strain curve for the submicron-porous NiTi, exhibiting the Young’s modulus of 4.8 GPa, ultimate stress of 250 MPa, and fracture strain of 7.9%. The detailed mechanical behaviors were: at first, the elastic deformation from bending or compressing the cell walls elastically contributed to the linear part. Then, buckling and cell collapsing contributed to the serration part and each serration corresponded to the fracture of the foam cell walls. Finally, the material failed [33,34]. Gibson et al. proposed a formula to calculate the crushing strength (rcr) of porous materials, which is related to the stress during serrated behavior; this formula is shown as follows for elastic-brittle open-cell foams:
Fig. 1. Photograph of plate-shaped Ni–Ti–Gd precursor before etching (top) and porous specimens after etching (down).
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Fig. 2. (a) XRD patterns and (b and c) SEM images of Ni–Ti–Gd precursor; (d–f) EDX elemental mapping of (d) Ni, (e) Ti, and (f) Gd.
Fig. 3. (a) XRD patterns and (b and c) SEM images of submicron-porous NiTi and (d) pore-size distribution of submicron-porous NiTi by mercury porosimeter, the red dot line is fitting curves of the experimental data. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
rcr q 0:65 qs rfs
3=2 ;
where rcr is the crushing stress of the foam, rfs is the tensile fracture strength for the dense counterpart, q is the foam density, and qs is the density of the dense counterpart [35]. In the present study, the density of dense and porous NiTi is calculated by dividing the mass by volume, finding that qq is 0.65. And rfs is 800 MPa s
from Ref. [9], so rcr is calculated to be 270 MPa, similar to the experimental data, 250 MPa, suggesting that elastic-brittle foam model can be applied to the present submicron-porous NiTi alloy.
Fig. 5 shows the internal friction (IF) and storage modulus as functions of the temperature for submicron-porous NiTi, measured in a single-cantilever mode. These results reveal damping peaks during heating and cooling, correlated to the austenite–martensite phase transformations. The submicron-porous NiTi exhibited a very high damping peak (0.07), 40% higher than those produced by conventional powder metallurgy process [36,37]. The high damping capacity of porous NiTi usually comes from the following factors: (1) stress concentration causes elastic and plastic bending, buckling, and even fracture of thin nodes or walls [38,39]; (2) reversible dislocation motion in stress–strain localization areas [38,39] and stress-assisted twin-boundary motion in martensite
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Fig. 4. Compressive stress–strain curves of (a) submicron-porous NiTi and (b) submicron-porous NiTiNb.
Fig. 5. Damping capacity and storage modulus as functions of temperature for submicron-porous NiTi.
platelets [40,41]; and (3) stress-induced martensite formation when loaded in the austenite state. For the present submicron-porous NiTi, the fine pore structure increases the interface areas between the NiTi and pores, enhancing the stress concentration and stress–strain localization areas. This behavior enhances factors (1) and (2), generating improved damping capacity even though its porosity is not very high. Ashby et al. summarized the relations between the loss coefficient (damping property) and Young’s modulus (mechanical property) for the most common damping materials at 30 °C [42]. As shown in Fig. 6, polymers possess high damping capacity but low Young’s modulus, while metals such as Mg alloys or Al alloys possess high Young’s modulus but low damping capacity. However, for engineering application, both high damping capacity and high Young’s modulus are desired. The present submicron-porous NiTi was marked with a triangle symbol and the counterpart fabricated by powder metallurgy process [36] was marked with a cycle symbol in Fig. 6. Both the Young’s modulus and damping capacity were higher for the present submicron-porous NiTi. Lead alloys also exhibit superior damping properties, but the toxicity and unfavorable high specific gravity are their drawbacks [43]. 3.2. Microstructures and properties of submicron-porous NiTiNb Fig. 7(a) shows the XRD patterns of Ni–Ti–Nb–Gd precursor and the porous Ni–Ti–Nb after etching, exhibiting B2-NiTi(Nb) and b-Nb phases for both specimens. Fig. 7(b) shows the SEM images
Fig. 6. Relations between the loss coefficient and Young’s modulus for commonly used damping materials at 30 °C [42]. The submicron-porous NiTi by a new topdown process is marked with a triangle symbol, and the counterpart fabricated by a conventional powder metallurgy techniques (bottom-up process) [37] is marked with a cycle symbol.
of the Ni–Ti–Nb–Gd precursor, revealing a darker phase (Ni47Ti44Nb9) and a brighter phase (NiGd). Fig. 7(c)–(f) shows EDX elemental mapping taken from Fig. 7(b), revealing b-Nb precipitates in Ni–Ti–Nb phase. The volume fraction of NiGd phase was 31% based on the image analysis of Fig. 7(b), indicating the porosity for porous NiTiNb after etching would be 31%. Fig. 8(a) shows the SEM image of porous NiTiNb, revealing homogenous submicron pores, with a NiTi(Nb) matrix and b-Nb precipitates. These results are supported by EDX elemental mapping of Nb, as shown in Fig. 8(b). The average pore size and porosity of the submicron-porous NiTiNb were 0.42 lm and 33% by mercury porosimeter, respectively, as shown in Fig. 8(c). Fig. 4(b) shows the compressive stress–strain curve for porous NiTiNb, exhibiting the Young’s modulus of 5 GPa, ultimate stress of 285 MPa, and fracture stain of 10.7%. The curve has three distinct parts: a linear elastic part in which the cell walls bended or compressed elastically, an increasingly steep yield part in which the cells collapsed by buckling and plastic yielding, and then the maximum stress followed by fracture failure. This behavior is similar to cancellous bone in compression, where a plateau of almost constant stress appears after the linear elastic part [44]. It should be pointed out that unlike the elastic stress–strain curve of submicron-porous NiTi before fracture as shown in Fig. 4(a), the
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Fig. 7. (a) XRD patterns and (b) SEM images of the Ni–Ti–Nb–Gd precursor; (c–f) EDX mapping taken from (b) for (c) Ni, (d) Ti, (e) Nb, and (f) Gd.
Fig. 8. (a) SEM images of submicron-porous NiTiNb, (b) elemental mapping of Nb taken from (a), and (c) pore-size distribution by mercury porosimeter, the red dot line is fitting curves of the experimental data. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
submicron-porous NiTiNb shows yielding phenomenon at about 160 MPa before fracture. This different behavior may be generated from the different microstructures. As shown before in Fig. 8(a) and (b), very fine ductile b-Nb phases precipitate inside and between the cell walls, and these b-Nb phases are considered to prevent the brittle collapse of thin nodes during deformation and to generate a yielding behavior. Moreover, from Ref. [35], the yield strength for elastic–plastic foams can be calculated as follows:
r q 0:3 qs rys
3=2 ;
where r is the yield stress of the foam, rys is the yield stress of the dense counterpart, q is the density of the foam, and qs is the density of the dense counterpart. Dividing the mass by volume, qq is found s
to be 0.6. And rys is 1000 MPa [45], so the calculated yield stress of the present submicron-porous NiTiNb is 140 MPa, similar to the experimental data, 160 MPa, indicating that the elastic–plastic foam model is suitable for the present submicron-porous NiTiNb alloy. Fig. 9 shows the temperature dependence of the damping capacity and storage modulus for the submicron-porous NiTiNb alloy. One damping capacity peak appeared during cooling, caused by the martensitic transformation. But before that, the damping
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(4) These new submicron-porous shape memory alloys are considered to be good candidate as engineering damping materials and this new top-down process can be used to produce various porous alloys with fine pore structures.
References
Fig. 9. Damping capacity and storage modulus as functions of temperature for submicron-porous NiTiNb.
capacity remained nearly constant over a wide temperature range. The maximum internal friction value and that at the austenite state were 0.18 and 0.04, 50% and 200% higher than those of dense NiTiNb, respectively [22]. The high damping capacity of submicron-porous NiTiNb was caused by both the composite microstructure, containing the B2-NiTi(Nb) matrix and b-Nb particles, and the porous structure as well. The interfaces between the b-Nb particles and NiTi(Nb) matrix provided nucleus sites for the martensite variants, increasing the number of boundaries between the martensite and parent phase; this behavior increased the energy dissipation, caused by movement of the martensite/parent interfaces, similar to dense NiTiNb alloy. The pore structures improved the stress concentration and stress–strain localization areas, enhancing the microplastic deformation, crack initiation, and dislocation motion, which could consume more energy, similar to porous NiTi alloy. The submicron-porous NiTiNb also exhibited high and steady internal friction value at the austenite state because of strain localization in the fine nodes and struts.
4. Conclusion (1) The submicron-porous NiTi and NiTiNb with homogeneously distributed fine pore structures have been firstly fabricated by a new top-down process, including solidified phase controlling and etching process. (2) The submicron-porous NiTi possessed much finer pore size (0.39 lm) and more homogeneous pore distribution than the counterparts by conventional powder metallurgy techniques (bottom-up process). It also showed similar mechanical properties to the counterparts, but 40% higher damping capacity (0.07) than those counterparts. (3) The submicron-porous NiTiNb also possessed fine pore size (0.5 lm). Its damping capacity was 0.18 and 50% higher than its dense counterpart. The internal friction value at austenite state was 0.04 and 200% higher than its dense counterpart.
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