High Damping of Lightweight TiNi-Ti2Ni Shape Memory Composites for Wide Temperature Range Usage
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A bimodal porous TiNi-Ti2Ni shape memory alloy composite (SMAC) with 59% porosity was fabricated by sintering Ti-46at.%Ni elemental powders with pore-forming agent. The porous TiNi-Ti2Ni SMAC contains two irregular pores of about 400 and 120 μm. We investigated the microstructure and pore morphology correlated with the mechanical properties and damping capacities of the SMAC. Ti2Ni intermetallic phases with size of 1-3 μm were homogeneously distributed in the TiNi matrix. The porous TiNi-Ti2Ni SMAC exhibits exceptionally high inverse mechanical quality factor (Q −1) of ~0.25 at < 40 °C, which is among the highest value reported for porous/dense shape memory alloys or composites to best of our knowledge, and it shows very high compressive fracture strain of about 25%. Moreover, the fabricated porous SMAC at relatively low strain amplitude can exhibit considerable high Q −1 of 0.06 ~ 0.11 for a wide range of temperature between − 90 and 200 °C, which is attributed to the stress concentration distribution provided by the bimodal structure of pores and the massive interfaces between pore/matrix and TiNi/Ti2Ni. These porous SMACs can be an ideal candidate for using as a lightweight damping material in the energy-saving applications.
Keywordscomposite damping NiTi porous alloys shape memory alloys
The lightweight materials with properties such as high strength, ductility and high energy absorption capability for wide temperature range (like − 100 ~ 200 °C) are an ideal choice for the fabrication of impact protection and noise isolation system in the aircrafts, ground and amphibious vehicles, tanks and submarines, etc (Ref 1). Metal foams (MFs) are lightweight and porous metallic structures fabricated by deliberately introducing interconnected pores into a solid metal matrix (Ref 2), such as Al foams, which exhibit high energy-absorbing capacity by continuous bending and collapsing of pores. Unfortunately, the Al foams have two major drawbacks of low strength and non-reusable after deformation. Recently, porous shape memory alloys (SMAs), especially TiNi SMAs (Ref 3), attract intensive attentions due to the unique properties like shape memory effect (SME, originated from martensite variant detwinning) or superelasticity [SE, stemmed from stress-induced martensite (SIM) transformation, can be reproduced more than ten thousand times at a high strain of 8% (Ref 4)], in addition to high strength and ductility (Ref 4). An extra merit of SMAs is the presence of generous interfaces existing in the martensitic phases, including the interfaces between martensite variants and twin boundaries in one martensite variant (Ref 5); thus, porous TiNi SMAs can exhibit much higher damping properties (or energy-absorbing capacity) and strength than the common metal foam (Ref 5); the key point of porous TiNi SMAs is reusable after relatively high deformation. However, porous TiNi SMAs would lose their high damping capacity in the parent (or austenitic) phase, usually above 100 °C, due to the vanishing of massive interfaces when the martensite phase transforms back to the parent phase. Thus, many researchers attempted to improve the damping capacity of the parent phase in the porous TiNi SMAs by tailoring the pore characteristics. It had been reported that the damping performance of parent phase can be enhanced with increasing the porosity due to the bending and collapsing of a great deal of pores (Ref 6). Furthermore, it can reach up to the maximum about 0.03 in inverse mechanical quality factor (Q −1, indicator of damping capacity) by adopting bimodal pore size (small pores distributed in the matrix between large pores) at high porosity of 60% (Ref 7). However, the porous TiNi SMAs with high porosity drastically lose their ductility and strength (Ref 7).
Recently, the SMAs matrix composites (especially nanostructural soft/hard dual-phase composites) received a lot of attentions owing to their high strength and large ductility simultaneously at a relatively wide temperature range (Ref 8, 9), as well as high damping properties (Ref 10). Their exceptional mechanical properties come from the synergistic effect of the soft and hard phases during the deformation (Ref 8, 9), and their high damping capacity also stems from energy absorption within the massive interfaces between nano-size soft/hard phases (Ref 10). For example, Zhang et al. (Ref 11) reported that the TiNi-Ti3Sn SMAC can exhibit as high as 3GPa compressive strength and 33% fracture strain, and moreover they showed high Q −1 value of 0.02 ~ 0.075 at temperature from − 120 to 300 °C (Ref 12). In addition, Guo and Kato (Ref 13) found that the bulk TiNi-Mg SMAC can exhibit high mechanical strength and damping capacity of 0.07 even for high-temperature parent phase above 70 °C, but poor ductility.
Therefore, porous TiNi SMAC could be an effective approach to overcome the drawback of low strength and ductility in porous TiNi SMAs, offering some ideal properties required for energy-saving applications, including lightweight, high strength and ductility, as well as remarkable high energy-absorbing capacity during wide temperature range. However, the porous TiNi SMAC has not been investigated yet, and very little is known about the mechanical behavior and damping capacity of such porous composite structure. In this work, we fabricated porous TiNi-Ti2Ni SMAC with bimodal pore sizes utilizing facile powder metallurgy method integrated with pore-forming technique. The mechanical properties and damping capacity of the synthesized porous SMAC are systematically studied by compressive stress–strain curves in combination with internal friction spectrum at different temperatures and amplitudes.
Materials and Methods
The designed porous TiNi-Ti2Ni SMAC had a nominal composition of Ti-46at.%Ni. For preparing porous TiNi-Ti2Ni SMAC, firstly, nickel powders (size: 50-75 μm) and titanium powders (size: 50-75 μm), with an atomic ratio of 46 to 54 were blended for 4 h. Then, 40 wt.% pore-forming agent (NH4HCO3) with two particle sizes (large size: 300-450 μm, 29 wt.%; small size: 70-125 μm, 11 wt.%) was added into the premixed Ni-Ti powders and continued blending for another 0.5 h. After that, the mixed powder was cold-pressed into some cylindrical green samples under 200 MPa in a hydraulic press. The green samples were heated at 200 °C for 2 h to decompose the NH4HCO3 particles firstly and sintered at 1050 °C for 10 h under argon gas flow as an inert atmosphere using a tube furnace (CVD(G)-07/50/2, Risine Inc.). The porosity of the as-sintered sample was determined as 59% by the Archimedes’ principle. A scanning electronic microscopy (SEM, Super 40, Zeiss) and an optical microscopy (DMI3000M, Leica) were used to characterize the microstructure and pore morphology, respectively. An x-ray diffractometer (XRD, MiniFlex 600, Rigaku) was used to determine the phase constituents present at room temperature. An energy-dispersive x-ray spectroscope (EDS, Bruker Quantax 200) apparatus attached to the SEM was used to characterize the composition of every phase in the porous TiNi-Ti2Ni SMAC. A TA Instruments differential scanning calorimeter (DSC) Q20 was used to analyze the phase transformation temperatures of all the samples. The DSC specimens were cut into pieces of 20-30 mg. For DSC analysis, the specimens were scanned between − 50 and 200 °C at a heating/cooling rate of 10 °C/min. A dynamic mechanical analyzer (DMA, Q800, TA Instruments) was employed to characterize the inverse mechanical quality factor (Q −1) spectrum. The DMA equipment was set in either multi-frequency–strain (Test: custom) mode or multi-strain (Test: strain-sweep) mode with single cantilever (clamp section). The porous specimens were wire cut with geometry of 20 × 4×1 (length × width × thickness, mm). In DMA multi-frequency–strain testing mode, the specimens were measured between − 90 and 200 °C at a rate of 5 °C/min, a frequency of 1 Hz and strain amplitude of 1.0%. For the multi-strain testing mode, the specimens were measured at − 50 and 130 °C, respectively, at a frequency of 1 Hz. The compressive mechanical behaviors were characterized using a material testing system (5984, Instron Inc.) with environmental chamber. The compressive specimens were machined by the wire cutting into cylindrical shape of 5 × 10 (diameter × length, mm) and then tested at room temperature and 130 °C at a strain rate of 3.33 × 10−4/s.
Results and Discussion
EDS results of different regions from Fig. 3
Besides the deformation behavior at RT, the cyclic compressive stress–strain curves of the porous TiNi-Ti2Ni SMAC with an increment of 1% strain at 130 °C are given in Fig. 4(b), and the sample may have austenitic B2 phase according to the DSC result discussed above. It can been seen that the sample exhibits about 1% superelasticity at 130 °C, and partial superelastic recovery maintained even if the pre-strain exceeds 3%. Furthermore, the maximum compressive stress can reach 44 MPa, as the dotted arrow shown in Fig. 4(b), which is still larger than that of porous TiNi SMAs with the similar porous structure at RT (Ref 7), and the sample does not fracture even if the pre-strain is about 3%. It is suggested that the ductility of bimodal porous TiNi-Ti2Ni SMAC can be greatly enhanced by introducing brittle Ti2Ni phase in either B19′ martensite phase or B2 parent phase. These results reflect that the fabricated porous TiNi-Ti2Ni SMAC exhibits high strength and fracture strain irrespective of the phases, i.e., austenite or martensite phase.
In order to determine the damping performance and storage modulus of the porous TiNi-Ti2Ni SMAC during a wide temperature range, the Q −1 and storage modulus spectrums with temperature are given in Fig. 6(b), and the strain amplitude was set as 1% which will not cause any plastic deformation according to the results presented in Fig. 4(b). We observed the SMAC exhibits superior damping capacity for all the temperature range tested, and the lowest Q −1 value can reach 0.06 for high temperature range (> 120 °C), and the Q −1 peak can reach 0.11 due to reverse martensitic transformation, which is close to the reported value for TiNi/Mg composite (Ref 13). However, it is worth to note that the storage modulus of the sample is only 0.2 ~ 0.4 GPa which is 1% of TiNi/Mg composite, and the density of the sample is about 2.67 g/m3 which is half of TiNi/Mg composite (Ref 13), and thus the damping capacity per weight of the porous TiNi-Ti2Ni SMAC is two times of TiNi/Mg composite.
To summarize, a bimodal porous TiNi-Ti2Ni SMAC with 59% porosity and with uniformly dispersed particle-like Ti2Ni phases in the TiNi matrix was successfully fabricated by a simple powder sintering method. The porous TiNi-Ti2Ni SMAC at the temperature below 40 °C exhibits high damping capacity of ~ 0.25, and high fracture strain of 25%. Furthermore, the porous SMAC shows high damping capacity of 0.12 and superelasticity of 1% at 130 °C owing to the synergic effect of TiNi, Ti2Ni and bimodal pores. This porous SMAC exhibits high damping capacity of at least 0.06 at 1% strain amplitude during the temperature ranged from − 90 to 200 °C. The porous SMAC could be considered as an ideal candidate for potential use as a lightweight damping materials in a harsh temperature environment.
This work was supported by the Foundation for Innovative Research Groups of the National Natural Science Foundation of China (No. 51621001), National Natural Science Foundation of China (No. 51571090), Provincial Natural Science Foundation of Guangdong (2016A030311012) and the Fundamental Research Funds for the Central Universities (2017ZD009).
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