Ti3SnâNiTi Syntactic Foams with Extremely High Specific Strength

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Functional Inorganic Materials and Devices 3

TiSn-NiTi syntactic foams with extremely high specific strength and damping capacity fabricated by pressure melt infiltration Changchun Xie, Hua Li, Bin Yuan, Yan Gao, Zhengtang Luo, and Min Zhu ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b08145 • Publication Date (Web): 16 Jul 2019 Downloaded from pubs.acs.org on July 17, 2019

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ACS Applied Materials & Interfaces

Ti3Sn-NiTi Syntactic Foams with Extremely High Specific Strength and Damping Capacity Fabricated by Pressure Melt Infiltration Changchun Xie1, Hua Li1, Bin Yuan1,2*, Yan Gao1,2, Zhengtang Luo3*, Min Zhu1,2

1School

of Materials Science and Engineering, South China University of Technology, Guangzhou 510640, China

2Key

Laboratory of Advanced Energy Storage Materials of Guangdong Province, Guangzhou, 51640, China

3Department

of Chemical and Biological Engineering, Hong Kong University of Science and Technology, Clear Water Bay, Kowloon, Hong Kong, 999077, China Correspondence to this article please direct to Dr. B. Yuan or Dr. Z. Luo Email: [email protected] or [email protected]

1

ACS Paragon Plus Environment

ACS Applied Materials & Interfaces 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Abstract: NiTi shape memory alloy foams have attracted much attention due to their unique superelasticity, excellent mechanical properties and damping capacities, but their high-temperature damping capacity and compressive strength remain to be a challenge. Herein, we demonstrate the preparation of Ti3Sn-NiTi syntactic foams using Ti58Ni34Sn8 alloy and alumina microspheres by a novel pressure melt infiltration and air-cooling strategy. The syntactic foams with 45% porosity, contain spherical and welldistributed pores of average size 500-600 μm. A fine lamellar Ti3Sn/NiTi eutectic with an inter-spacing distance of 600-900 nm and a Ti2Ni interfacial layer of 10 μm thickness were formed between the alumina microspheres and the matrix. The syntactic foams achieved a high specific compressive strength (110.2-110.8 MPa ∙ cm3/g) at a wide temperature range because of the large interfacial area and the good lattice strain matching in the lamellar Ti3Sn/NiTi. They also exhibited 2% recoverable strain and high specific energy absorption capacity (31.5 kJ/kg). Moreover, the foams showed ultrahigh damping capacity (0.066) at a temperature range of -150℃ to 200℃. Most interestingly, the Ti3Sn-NiTi syntactic foams showed the highest comprehensive coefficient (σ/ρ) ∙ Tanδ of 5.07 to date. Because of these impressive features, Ti3SnNiTi syntactic foams become a promising material for energy absorption and damping applications.

Key words: Ti3Sn/NiTi; syntactic foam; strength; damping capacity; melt infiltration

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1. Introduction Lightweight materials with high strength and damping capacity are highly attractive in engineering applications1-3. In particular, they are often used as energy absorbing materials to reduce vibration and noise in the fields of alternative fuel vehicles, high-speed rails, and aerospace. Metal and polymer foams, such as Al and Mg foams and their composites, are known to possess high damping capacity and low density, but their low strength and poor environmental adaptability resistance have limited their extensive applications4-9. It is of importance and interest, therefore, to design a novel low-density material that possesses both high strength and damping capacity within a wide range of service temperature. Recently, NiTi shape-memory alloys (SMAs) have attracted much attention as biomedical and energy-absorbing materials for their unique shape-memory effect, superelasticity, high damping capacity, good mechanical properties, and corrosion resistance10-11. NiTi SMAs can exhibit high damping capacity at martensite state because of their numerous movable interfaces, including those between martensite variants and twin boundaries12-13. However, dense NiTi shows some fatal drawbacks; for example, the damping capacity of B2-NiTi at high temperature (over 100℃) drops to a very low level (0.005) as the result of the loss of massive movable interfaces. In addition, the density (6.45g/cm3) and cost (100 US$/kg) of dense NiTi are relatively high10-13, thereby limiting its extensive applications as an energy-absorbing material. Fortunately, it has been reported that NiTi SMA foams can exhibit a lower density and high damping capacity at a wide temperature range from -100 to 200℃ by utilizing the bending and kinking of numerous pore walls and severe stress concentration around them1-3. But the strength and superelasticity of the NiTi foams would deteriorate poorly as a consequence of the terrible physical discontinuity of the foams and the non-uniform distribution, irregularity of the pores, which easily cause stress concentration and failure of foams2-3, 14-17. To date, there are two major approaches to improve the mechanical properties superelasticity and damping capacities of metal foams; the first of which is by 3



modification of their pore characteristics. At present, NiTi foams are mostly prepared by powder metallurgy (PM)18-20 and additive manufacturing (like selective laser melting)21-24 because of their high melting point (1310 ℃ ) . NiTi foams fabricated through PM exhibit pores that are large, irregularly shaped, and non-uniform in size, leading to poor mechanical properties. In addition, porous NiTi parts with any porous architectures can also be formed by additive manufacturing using the pre-alloyed NiTi powders. However, their properties are sensitive to fabrication orientation, post-process heat treatment, or laser power, and the defects (e.g. cracks, residual stresses) from the fabricating method may lead to deterioration during mechanical testing21. Guo et al.1 attempted to acquire spherical pores of homogeneous pore size distribution in NiTiNb by solidified phase controlling and etching process. The method could improve its strength to 270 MPa and damping capacity to about 0.04 at B2 phase, but the foam was only 31% porosity and thus becomes a limitation of this technique. Moreover, it is known that the addition of ceramic microspheres into the Al, Mg, Zn, steel, and metallic glass matrix can produce various syntactic foams25-30. As compared to pure foams, syntactic foams exhibit higher strength and better damping capacity because of the existence of hard microspheres, their spherical pores, and the large interfacial area between the matrix and the microspheres. The second method is the introduction of strengthening phases into NiTi SMAs matrix composites (SMAC). Soft/hard dualphase composites have received a lot of attention recently because of their excellent mechanical properties and high damping capacities31-35. For instance, it was reported that Ti3Sn showed excellent damping performance and desirable mechanical properties at a wide temperature range36-38. Zhang et al.39-41 observed that dense NiTi/Ti3Sn composites exhibited a high compressive strength of 3.0 GPa and a fracture strain of 33%. This excellent performance was attributed to lattice strain matching and the large interfacial area between brittle Ti3Sn phase and ductile NiTi phase in the composite. The deformation mechanism of Ti3Sn/NiTi differs from the dislocation slip in common metal matrix composites as the former can reduce stress concentration and restrain crack propagation. However, the problems with enhancing the damping capacity of the B2 phase and reducing its density are yet to be overcome. 4


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In order to combine the advantages of syntactic foams and dual phase composites, the preparation of Ti3Sn-NiTi syntactic foams with low density (i.e. high porosity), high strength and damping capacity was investigated. It was reported that the melt infiltration which adapts space holders is a valid method to fabricate metallic syntactic foams with high porosity of about 40% to 80%42-43. This method, however, is seldom used for alloys with high melting point. To the best of our knowledge, no NiTi-based foams have been reported to have used melt infiltration in order to obtain high porosity and spherical pores. Herein, Ti3Sn-NiTi syntactic foams were prepared by the pressure melt infiltration method in a vertical tube furnace designed by ourselves, where alumina microspheres were selected as space holders. The syntactic foams were then subjected to fast cooling, after which the fine matrix phase, compressive properties, and damping capacities were controlled.

2. Experimental Alloy ingots with nominal composition of Ti58Ni34Sn8 (at. %) were prepared from high-purity (>99.5 wt.%) elemental metals by vacuum arc melting in a water-cooled copper hearth under an argon atmosphere. Alumina microspheres with size of 500 μm to 700 μm and wall thickness of 10 μm to 30 μm were selected as space holder particles. The characteristics of the alumina microspheres used are shown in Figure S1 and Table S1. The pressure melt infiltration was carried out in an SK-1600 vertical vacuum-gas tube furnace. An alumina crucible of 28-mm diameter and 160-mm height was coated with BN before being used as a mold for the syntactic foams to avoid unwanted reactions and the adherence of the samples to the crucible surface. The crucible was then filled with alumina microspheres up to a height of about 40 mm before a 1-mm thick layer porous plate and 90 g of the Ti58Ni34Sn8 alloy were added in a manner according to the illustration in Fig. 1a. For the infiltration process (Fig. 1b) the crucible was heated for 20 minutes in a vacuum furnace at a constant temperature of 1200℃ after a ramp time of about 3 hours, melting the alloy completely. The crucible was then pressurized to 0.1 MPa using argon gas and was maintained at this pressure for 10 5



minutes to force the molten metal into the voids between the space holders. Afterwards, the samples were furnace-cooled and air-cooled, respectively, labelled as sample A and sample B. The expected result after infiltration is shown in Fig. 1c.

Figure 1. A schematic representation of the pressure infiltration: (a) the setup before infiltration, (b) the infiltration process under vacuum heating and the subsequent pressurization, (c) the product of pressure infiltration.

The microstructures of the Ti3Sn/NiTi syntactic foams were characterized by Xray diffraction (Philips X’pert XRD) and scanning electron microscope (Zeiss Supra 40 SEM) equipped with an energy-dispersive X-ray spectrometer (EDX). The phase transformation behaviors, on the other hand, were tested using a differential scanning calorimeter (DSC Q200, TA) at temperatures from -80℃ to 200℃ and a constant heating or cooling rate of 10℃/min. Mechanical properties and superelastic behavior were also measured from specimens of size 6 × 6 × 9 mm (including about 10 pores in each cross-section side of the samples) using an Instron 5984 universal testing machine at a strain rate of 3.33 × 10 ―4/s at temperature of 25℃ and -50℃. Damping tests for 25 × 6 × 1.8 mm samples were then conducted using a dynamic mechanical analyzer

(DMA Q800, TA) in a single cantilever mode (bending through loading at one end of the sample). Damping capacity (indexed by tan δ) as a function of temperature was inducted at a constant strain amplitude of 0.2%, a frequency of 0.1Hz, and a constant heating or cooling rate of 5℃/min from -120℃ to 200℃. Damping capacity as a 6


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function of strain was inducted at a variable strain amplitude of 0.005% to 0.2%, a frequency of 1Hz, at 150℃ (austenite state) and -100℃ (martensite state), respectively. All the mechanical or damping parameters were measured through at least two samples from different batches, and the obtained value were given in form of average value± error.

3. Results and discussion

Figure 2. Optical micrographs of sample A (a) and sample B (b), where insets show the macrographs (the areas filled with molten metal are marked with yellow arrows); The quantity and area in percent of total for different pore sizes: sample A (c), and sample B (d).

The macrographs of the samples (insets of Fig. 2a and 2b) prove that large and well-infiltrated TiNiSn foams could be obtained by the pressure melt infiltration process. Optical micrographs (as shown above) suggest that their pores completely replicated the shape of alumina microspheres, forming nearly spherical and uniformly distributed pores. Specifically, the distribution (Fig. 2c and 2d) indicates that mediumsized pores with diameter of 500-700 μm account for most of both the total surface area 7



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and total quantity. Further data analysis showed the average pore size for samples A and B were 552±18 μm and 537±27 μm, respectively. A low-density material is generally desired for engineering applications as previously mentioned. According to the Archimedes’ law and by using the measured average pore size, the calculated densities were about 3.75g/cm3 and 3.65g/cm3 for sample A and B, respectively. Their porosities, on the other hand, were also calculated as 43% and 45%, respectively, using the following equation. 𝑟i

𝑃 = 𝑉mb(𝑟o)3 +

𝜌t ― 𝜌m 𝜌t

(1)

The above formula25 is the sum of the porosity of alumina microspheres and the porosity due to either insufficient or excessive infiltration. Factors, such as the infiltration temperature and pressure, holding time and microsphere characteristics influence the degree of infiltration. As seen in the areas marked with yellow arrows (Figs. 2a and 2b), some pores were filled with molten metal, which may be caused by either excessive infiltration or the initial rupture of microspheres. As confirmed by XRD (Figure S2), the matrix phase in the syntactic foams were the B2-NiTi and Ti3Sn phases because of they had much higher concentration than the other phases, Al2O3, Ti2Ni, and B19’-NiTi. The Al2O3 phase detected may have originated from the alumina microspheres used. Subsequently, backscattered SEM images (Figs. 3a and 3b) show the microstructure of the matrix phase far away from the pore. EDX results (Figure S3-S6) identified that the dark gray area was the NiTi phase, the light gray area the Ti3Sn phase, and the pattern of dark and light areas the Ti3Sn/NiTi eutectic mixture. According to the pseudobinary phase diagram of NiTi-Ti3Sn37 (Figure S7), the Ti58Ni34Sn8 alloy should form a eutectic mixture after solidification. However, it was determined that furnace-cooled sample A was mainly composed of proeutectic NiTi (20-30 μm) and Ti3Sn (15-20 μm) phases, as well as some lamellar Ti3Sn/NiTi eutectics. In contrast, for the air-cooled sample B, the content of proeutectic phases decreased notably, while the proportion of lamellar eutectic region greatly increased. Upon higher magnification, the lamellar Ti3Sn/NiTi eutectic could be clearly observed (as shown in the inset of Fig. 8


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3b). Furthermore, quantitative image analysis determined that the volume fractions of lamellar eutectics were about 20% and 80% for sample A and B, respectively. Close investigation of the microstructure near the pore (Figs. 3c and 3d) indicated that the Ti2Ni interfacial layer, which was formed between alumina microspheres and matrix were about 20 μm and 10 μm thick for sample A and B, respectively. Additionally, EDX elemental mapping of lamellar eutectics (Figs. 3e and 3f) clearly confirmed that the eutectic structure was composed of alternating layers of NiTi and Ti3Sn phases, and the lamellar spacing of NiTi and Ti3Sn are about 3 μm and 2 μm for the sample A, while 0.9 μm and 0.6 μm for the sample B.

Figure 3. SEM images of the matrix phase far away from the pore: (a) sample A, (b) sample B, inset image is the magnification of lamellar Ti3Sn/NiTi eutectic structure; SEM images of the matrix phase near the pore: (c) sample A, (d) sample B; EDX elemental mapping of lamellar eutectics: (e) sample A, (f) sample B. 9



Significantly, the cooling rate was identified as the main reason for causing the difference in microstructures of these two samples. When the cooling rate was slow, the grains could easily grow as there was enough time to precipitate more proeutectic phases, however, the lamellar Ti3Sn/NiTi eutectic tends to be impeded, which leads to a low volume fraction of the eutectic structure. To add to this, the lamellar spacing of Ti3Sn for the furnace-cooled sample A was also much larger than that of the air-cooled sample B as previously mentioned. Another effect of cooling rate is in the formation of Ti2Ni interfacial layer, as there was a difference of thermal conductivities between the alumina microspheres and the matrix. The lower thermal conductivity of alumina reduced the rate of heat dissipation at the interface and caused overheating therein. This resulted to the formation of larger grains at the interface and the movement of more active Ti atoms towards the interface, forming the Ti2Ni interface layer. In conclusion, lower infiltration temperature, shorter holding time, and faster cooling rate could lead to a thinner Ti2Ni interface layer. The martensitic transformation behaviors of these two TiNi-based syntactic foams and the TiNiSn ingot were characterized by measuring their DSC curves (Figure S8) and determining the phase transformation temperature using the tangential method, the transformation temperature and transformation enthalpy are listed in the Table S2. It is well known that no solid phase transformation occurs in Ti3Sn phase from -50℃ to 100℃ 40. Therefore, the two overlapping exothermic (and endothermic) peaks of the TiNiSn ingot in the cooling (and heating) curve corresponded to the B2R and RB19’ (and B19’R and RB2) phase transformation38-40. However, for the aircooled sample B, the exothermic (and endothermic) peak in the cooling (and heating) curve indicated the B2B19’ (and B19’B2). Further measurements showed that the austenite transformation finish temperatures (Af) of sample A, sample B, and the TiNiSn ingot were 18℃, 19℃ and 83℃. These data suggested that the samples were mainly composed of B2-NiTi phases at room temperature (25℃), verifying the XRD analysis presented. Moreover, as the phase transformation temperature of the syntactic foams shifted to a lower temperature, the transformation enthalpy decreased and thus widened the temperature range of martensitic transformation. One of the causes of this 10


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observation was the mismatch between alumina microspheres and the matrix in both thermal expansion coefficient and elastic modulus, which could have caused the thermal mismatch strain near the interface. This delayed martensitic transformation, leading to a lower and wider range of phase transformation temperature44-45. The second possible reason was the formation of the Ti2Ni interfacial layer, which could have increased the nickel content in the NiTi matrix, greatly decreasing the martensitic transformation temperature46-47. 3.1 Compressive mechanical properties

Furnace-cooled sample A Air-cooled sample B

Compressive Stress (MPa)

450

o

(a) 25 C

120

400

100

350

80

300 250

60

200

40

150 100

20

50 0

0

1

2

3

4

5

6

7

8

0

Specific strength (MPacm 3/g )

500

Compressive Strain (%)

o

(b) -50 C

110 100

400

90

350

80

300

70

250

60

200

50 40

150

30

100

20

50 0

10 0

1

2

3

4

5

3

Furnace-cooled sample A Air-cooled sample B

450

Specific strength (MPacm /g )

500

Compressive Stress (MPa)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


0

Compressive Strain (%)

Figure 4. The compressive mechanical properties of sample A and sample B described by the compressive stress-strain curve: (a) at 25℃ (B2 phase) and (b) at -50℃ (B19’ phase). 11



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The compressive mechanical properties of the Ti3Sn-NiTi syntactic foams at 25℃ and at -50℃ (Fig. 4) provide a good measure of excellent mechanical properties. As mentioned, the matrix alloy, which is mainly composed of B2-NiTi and Ti3Sn phases at 25℃, transformed into B19’-NiTi and Ti3Sn phases at -50℃ for the both samples. At 25℃, samples A and B exhibited ultimate compressive strengths of about 335.0±6.0 MPa and 402.0±1.0 MPa and specific strengths (ratio of strength to density) of 89.3±1.6 and 110.2±0.3 MPa ∙ cm3/g, respectively (Table S3). Lowering the temperature to -50 ℃

slightly improved their ultimate compressive strengths to

355.8±5.1 MPa and 404.5±11.5 MPa, respectively. Table 1 compares compressive strengths of sample A and B with the results reported for other metal matrix syntactic foams (MMSF). Clearly, the air-cooled Ti3Sn-NiTi syntactic foams exhibited a remarkable advantage in specific compressive strength. Factors that enhanced the mechanical properties of Ti3Sn-NiTi matrix syntactic foams were proposed as follows. Firstly, the formation of some submicron lamellar Ti3Sn/NiTi eutectic structure in the matrix alloy improved the strength of the foams, resembling the strengthening mechanism of pearlite. In addition, a good lattice strain matching between NiTi and Ti3Sn could have occurred under external loading, the composite with this fine lamellar structure were proven to exhibit extremely high compressive strength before39-41. Secondly, the uniformly distributed, spherical pores reduced the local stress concentration and thus helped the sample to bear higher loads, moreover the alumina microspheres could bear the partial loads and thus prevent the propagation of cracks in the matrix30.

Table 1. The comparison of the compressive strengths of Ti3Sn-NiTi syntactic foam with other metal matrix syntactic foams at room temperature. Foams

Compressive strength (MPa)

Specific strength (MPa ∙ cm3/g)

Specific energy absorption (kJ/kg)

Ref.

Ti3Sn-NiTi

402.0±1.0

110.2±0.3

31.5±0.8

This work

NiTiNb

270

66

-

1

Zn matrix

138-217

48-75

25.3-34.7

28

12


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Al matrix

36-80

29-65

14.8-24.8

48-49

Mg matrix

118-168

65-93

42.2-53.7

50-51

Steel matrix

150-300

36-72

27.5-39.2

29

Compressive stress (MPa)

350

st

(a)

1

nd

rd

2

3

th

5

300

th

10

250 200 150 100 50

1%

0

Compressive strain (%)

105

(b)

100 Recovery ratio (%)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


95 90 85 80

1

2

3

4

5 6 7 Cycle number

8

9

10

11

Figure 5. (a) The superelastic stress-strain curve and (b) recovery ratio under compressive cycling for the air-cooled sample B.

The mechanical fatigue of the air-cooled Ti3Sn-NiTi syntactic foam (sample B) was tested by subjecting the sample to strain cycles ten times at 3% pre-strain. The stress-strain curve (Fig. 5a) and recovery ratio (Fig. 5b) are shown. The Ti3Sn-NiTi syntactic foam demonstrated 2.4% recoverable strain and partial superelasticity of 0.5% 13



after the first cycle, as the recovery ratio for the first four cycles was 81%, 97.9%, 98.3%, and 99.1%, respectively, even reaching more than 99.5% afterwards. Moreover, the sample can almost completely recover under 2.2% prestrain and behave 0.5% superelasticity for the 10th cycle. In contrast with other MMSFs, the unique superelasticity and high recoverable strain would allow the Ti3Sn-NiTi syntactic foam to be reused repeatedly under strains of less than 2%, making it a potential material for some engineering applications. 3.2 Energy absorbing and damping capacity

3

Energy absorption capacity (MJ/m )

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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Air-cooled sample B

120 100 80 60 40 20 0

0

5

10 15 20 25 30 35 40 45 50 55 60 Strain (%)

Figure 6. The energy absorption capacity of the air-cooled Ti3Sn-NiTi syntactic foam (sample B) as a function of strain.

The energy absorption capacity, defined as the area under the compressive stressstrain curve from zero up to the densification strain (𝜀𝑑)52, is calculated as follows: 𝜀

W=∫0𝑑𝜎𝜀𝑑𝜀

(2)

where W means the energy per unit volume of the material (MJ/m3). The energy absorption capacity of the air-cooled Ti3Sn-NiTi syntactic foam (sample B) was measured as a function of strain (Fig. 6). Because of the gradual fracture and collapse of the alumina microspheres for the TiNiSn syntactic foam, the compression stress14


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strain curve fluctuated, showing a zigzag shape in its graph (Figure S9). Listed in Table 1 are the calculated energy absorption capacity of the air-cooled sample B, as well as the results from other MMSFs. The specific energy absorption of TiNiSn syntactic foam was well above those of most MMSFs. Moreover, it could exhibit better corrosion resistance and higher strength than the Mg-based and steel syntactic foams, indicating that the TiNiSn syntactic foam could also be used in the energy absorption fields.

(a)

0.30

Cooling damping peak

0.25

0.30 0.25

Heating damping peak

0.20

Tan 

0.35

NiTi ingot TiNiSn ingot Sample A Sample B

0.20

0.15

0.15

0.10

0.10

0.05

0.05

0.00 Intrinsic damping at M -100

-50

Intrinsic damping at A

0

50

100

150

Tan 

0.35

0.00

200

o

Temperature ( C)

0.06

(b)

A M

0.05 0.04

Tan 

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


0.03 0.02 0.01

The air-cooled sample B 0.00 0.00

0.05

0.10

0.15

0.20

Strain (%)

Figure 7. (a) The damping capacity (indexed by tan δ) for different samples as a function of temperature during the cooling and heating process at frequency of 0.1 Hz, strain amplitude of 0.2%, and cooling/heating rate of 5℃/min; (b) the damping capacity as a function of strain for sample B at 150℃ (austenite state) and at -100℃ (martensite state) under a frequency of 1 Hz, a temperature varying rate of 5℃/min, and a strain amplitude of 0.005%-0.2%. 15



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Table 2. The intrinsic damping capacities of sample A, sample B, TiNiSn and NiTi ingot at the austenite and martensite states Intrinsic damping at

Intrinsic damping at

austenite state

martensite state

Syntactic foam A

0.052±0.001

0.053±0.001

Syntactic foam B

0.071±0.001

0.064±0.001

TiNiSn ingot

0.014±0.001

0.040±0.002

NiTi ingot

0.009±0.001

0.036±0.002

Alloys

In order to characterize the energy dissipation capacity at a low strain level, the damping capacity was measured as a function of temperature, showed in Fig. 7a, during the cooling and heating of Ti3Sn-NiTi syntactic foams, TiNiSn ingot, and NiTi ingot, and their intrinsic damping capacities are also presented (Table 2). Two damping capacity peaks appear for all alloys, which were caused by the martensitic and the reverse martensitic transformation. As compared with the TiNiSn ingot, the damping capacity peaks of syntactic foams shifted to a lower temperature, which was consistent with the DSC measurement. Although the damping capacity peak of NiTi was high and close to 0.3, its intrinsic damping capacity independent of temperature variation at the austenite state and the martensite state are relatively low. However, much more consideration was taken with the intrinsic damping capacity, as the damping capacity peak has generally no practical application. Interestingly, the intrinsic damping capacities of the TiNiSn alloy ingot at austenite and martensite states were 56% and 11% higher than those of the dense NiTi alloy ingot. Moreover, for sample A (or sample B), the intrinsic damping capacities at austenite and martensite states improved significantly by 478% (689%) and 47% (78%), respectively, in comparison with that of the NiTi alloy ingot. This huge improvement in intrinsic damping capacities at austenite state was caused by the introduction of spherical alumina microspheres and fine lamellar Ti3Sn/NiTi eutectic. Fig. 7b presents the damping capacity as a function of strain for the air-cooled sample B at 150℃ (austenite state) and -100℃ (martensite state). With the gradually 16


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increase in strain, the damping capacity at either state quickly reached a maximum value (0.058 for both) at a small strain amplitude (

Ti3SnâNiTi Syntactic Foams with Extremely High Specific Strength

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Ti3SnâNiTi Syntactic Foams with Extremely High Specific Strength