Abstract
Imitating the structure of steel-reinforced concrete, a composite coupling good damping capacity and mechanical property was fabricated by pressure infiltration progress. The aluminum (Al) matrix composite was hybrid reinforced by 20% volume fraction of SiC particle (SiCp) and 20% volume fraction of TiNi fiber (TiNif). The damping capacity of the composite in the temperature range from 30°C to 290°C was studied using a dynamic mechanical analyzer (DMA). Due to the B19′→B2 reverse martensitic transformation in TiNif, a damping peak showed up in the heating process. Furthermore, both the hysteretic effect of the martensite/variants interfaces in TiNif and the weak bonding interface between SiCp and TiNif were attributed to the high damping capacity of the composite. After tension deformation, a compressive stress was formed in the composite in the heating process. With the help of compressive stress, the value of the damping peak was much higher than before, since the movement of dislocation in the Al matrix was easier.
1 Introduction
As subjecting to cyclic and dynamic stresses, the structural materials in engineering applications always need high damping capacity and excellent mechanical properties. Aluminum (Al) alloys that are exceptionally lightweight and have high specific stiffness are often used as structural materials, while the application on the precision components is limited due to the poor damping capacity of Al alloys. Hence, the Al matrix composites coupling high damping capacity and good mechanical properties were developed [1], such as the 6061Al/SiC particle (SiCp)/Gr hybrid metal-matrix composites [2, 3]. Based on the previous studies, SiCp was often used as a reinforcement to improve the mechanical properties and modulus of the composite [4, 5]. The strengthening mechanism had been discussed by Miller and Humphreys [6]. Moreover, the addition of SiCp to Al alloy matrix increased the wear resistance of the composite [7, 8]. In recent years, TiNi alloy, which was one of the most common shape memory alloys (SMA), was used to make composites with Al [9, 10]. By utilizing the excellent shape memory effect, compressive stress could be formed in the composite. The mechanical properties [11] of the composite, such as yield strength [10, 12] and fatigue life [13, 14], could be improved with the help of compressive stress. Furthermore, the damping capacity [12, 15] of the composites with TiNi alloys could also be improved due to the good damping capacity of martensite.
With the hybrid reinforced by SiCp and TiNi fiber (TiNif), both the mechanical properties and the damping capacity of the composite might be improved. Based on the studies about TiNi short fiber and SiCp-reinforced Al alloy composites by Chaudhury et al. [16] and Akalin et al. [17], the mechanical properties and wear resistance of the composite were improved. In this study, a new Al matrix composite coupling good damping capacity and mechanical properties was fabricated by pressure infiltration progress. The structure of the composite that was hybrid reinforced by SiCp and TiNif is similar to the reinforced concrete structures. As shown in the schematic diagram in Figure 1, SiCp is uniformly distributed around TiNif in the Al matrix. The damping capacity of the composite before and after tensile deformation was studied using a dynamic mechanical analyzer (DMA).
Schematic diagram of the structure in the SiCp+TiNif/Al composite.
2 Materials and methods
The SiCp+TiNif/Al composite was fabricated by pressure infiltration process in ambient air. Pure Al (Northeast Light Alloy Co., Ltd, China) and Ti-50 at.% Ni fiber (Northwest Institute For Non-ferrous Metal Research, China) with the diameter of 200 μm was used in this study. As shown in the phase transformation temperatures in Table 1, the martensitic transformation start (Ms) and finish (Mf) temperature of TiNif was 64.8°C and 35.1°C, respectively, whereas the reverse martensitic transformation start (As) and finish (Af) temperature was 114.3°C and 121.5°C, respectively. Because Mf is higher than room temperature, the TiNi alloy is in the martensite phase at room temperature. The volume fraction of TiNif was set at approximately 20%. In order to prevent contact, TiNif was arrayed straight at a constant spacing. Pure Al powders with the average size of 10 μm were mixed with 20 vol.% SiC powers (Harbin Xin Ceramics Co., Ltd, China) with an average size of 5 μm in a ball mill for about 10 h. The mixed powers containing pure Al and SiC were poured into the gap between TiNif. The fiber holder with mixed powder was then placed in a mold into which molten pure Al was poured, followed by pressurization at 35 MPa. Then, the as-fabricated composite was annealed at 200°C for 30 min, and 20 vol.% SiCp/Al composite without TiNif was also prepared in the same way as described above for comparison purposes.
Transformation points of TiNif.
| Transformation points | Temperature (°C) |
|---|---|
| Mf | 35.1 |
| Ms | 64.8 |
| As | 114.3 |
| Af | 121.5 |
The distribution of TiNif and SiCp in the pure Al matrix was characterized using a scanning electron microscope (Quanta 200FEG, FEI, USA). The uniaxial tensile properties of the composites were conducted on Instron 5569 (Instron, USA) with tensile velocity of 1×10-2 mm/s at room temperature. A rectangle specimen with a width of 10 mm and a length of 80 mm was used in the tensile test. DMA (Q800, TA, USA) was used to study the damping capacity of the composites. Rectangle specimens with a geometry of 1.5×8×35 mm (thickness×width×length) were measured in the single-cantilever mode. Furthermore, the damping tests were conducted at a constant strain amplitude of 5×10-5 in the temperature range from 30°C to 290°C. The constant heating rate was 10°C/min and the frequency was 1 Hz. The length variation of the SiCp+TiNif/Al composite before and after tensile deformation was detected with the NETZSCH DIL 402C (NETZSCH, Germany) dilatometer in the temperature range between 0 and 290°C.
3 Results
3.1 Microstructure of the SiCp+TiNif/Al composite
Figure 2 shows the typical micrographs of the SiCp+TiNif/Al composite. It is obvious that the composite fabricated by pressure infiltration process was in good macroscale homogeneity. The distribution of TiNif in the Al matrix was uniform as shown in Figure 2A and the dispersion of SiCp in the Al matrix around TiNif was well distributed as shown in Figure 2B.
Microstructure of the SiCp+TiNif/Al composite: (A) arrangement of TiNif in the Al matrix and (B) dispersion of SiCp in the Al matrix around TiNif.
In order to observe more detailed microstructural features in the interface of the composite, a higher-magnification micrograph of the interface is conducted. As shown in Figure 3A and B, a clean interface was formed not only between SiC and Al but also between TiNi and Al. Furthermore, no obvious porosity was observed in the composite. An electron probe microanalysis (EPMA) was carried out to investigate the interfacial reaction between TiNif and Al matrix in the composite. As shown in the chemical composition changes in Figure 3D, about 3-μm-thick interfacial reaction layer was formed in the vicinity of the boundary. The main contents of O, Al, Ti, and Ni elements were observed.
SEM micrograph of the interface in the SiCp+TiNif/Al composite: (A) interface between SiC and Al, (B) interface between TiNi and Al, (C) energy-dispersive X-ray spectroscopy (EDX) scanned region (shown as a white line), and (d) EPMA analysis of the interface.
3.2 Tensile deformation of the composite
Figure 4 shows the stress-strain curves of the pure Al, SiCp/Al, and SiCp+TiNif/Al composites. Both the yield and tensile strength of the composites were markedly higher than that of the pure Al matrix. The tensile strength of the SiCp+TiNif/Al composite was 218 MPa. It is almost 3.5 times higher than 60 MPa, which is the tensile strength of the Al matrix. However, the elongation of the composites was reduced to about 6%. The uniaxial tensile deformation of the SiCp+TiNif/Al composite was similar to the SiCp/Al composite. The addition of SiCp to the Al matrix increased the dislocation density, resulting in a high equivalent stress and a low equivalent strain for the dynamic recovery of the composites [18]. The enhancement mechanism of SiCp was mainly attributed to the high strength and high modulus. Moreover, due to the excellent plastic property, TiNif were elongated with the Al matrix in the tensile deformation. They were pulled out and no fracture of the composite was observed even when the Al matrix with SiCp fractured.
Stress-strain curves of the Al, SiCp/Al, and SiCp+TiNif/Al composites at room temperature.
TiNif in the composite are in the martensite phase at room temperature because the Mf is higher than room temperature. It has been confirmed by previous studies that the reorientation of the martensite was induced with the increase of the strain [19]. After tensile deformation, a recovery stress could be induced in the follow-up heating process due to the shape memory effect of TiNif. The recovery stress resulted in the formation of compressive stress in the composite [20]. In order to study the effect of compressive stress on the damping capacity of the composite, a prestrain treatment was performed on the SiCp+TiNif/Al composite and the amount of the tensile strain was 2%.
3.3 Damping capacity of the composite
The damping capacity and storage modulus of the SiCp+TiNif/Al composite at different temperatures are shown in Figure 5A. Obviously, the damping capacity of the composite went up, while the storage modulus decreased with the increase of temperature. Moreover, a damping peak was observed at the temperature of 140°C. The damping peak in the heating process was associated with the B19′→B2 reverse martensitic transformation of TiNif. Because the thermal mismatch stress was induced between the Al matrix and TiNi alloys in the heating progress, in addition to the hysteretic effect induced by the coupled operations of temperature and stress under the testing conditions of DMA [15], the start temperature of the reverse martensitic transformation of TiNi was deferred to 120°C. The evolution of the damping capacity and storage modulus with the increase in temperature was contrary, except at the temperature close to 140°C. Since the storage modulus of B2 austenite was higher than that of B19′ martensite, the storage modulus of the composite increased slightly with the phase transformation from B19′ martensite to B2 austenite.
Tan δ and storage modulus of the SiCp+TiNif/Al composite as a function of the temperature from 20°C to 290°C at the frequency of 1 Hz (A) before and (B) after tensile deformation.
After the prestrain treatment, the damping capacity and storage modulus of the composite are shown in Figure 5B. It is clear that there are some differences with the composite before tensile deformation, especially when the temperature was higher than 130°C. In the temperature range from 130°C to 230°C, the martensite in TiNif transformed into austenite and a compressive stress was formed in the composite. It resulted in some differences. (a) The storage modulus of the composite was increased after the prestrain treatment. (b) The value of the damping peak was 0.06. It was higher than 0.025, which was the value of the damping peak before the prestrain treatment. (c) The damping peak was broadened and the temperature of the damping peak was about 160°C. It was higher than 140°C, which was the temperature of the damping peak before the prestrain treatment. Therefore, it was clear that the damping capacity of the composite was increased by the compressive stress.
Figure 6A shows the evolution of the damping capacity as a function of temperature at different frequency. The damping capacity of the composite was the same at different frequency in the range temperature from 20°C to 50°C. However, when the temperature was higher than 50°C, the damping capacity of the composite at 1 Hz was much better than that at 10 and 20 Hz. Moreover, with the increase in frequency, the value of the damping peak decreased. Due to the various substructures such as twin dislocation in the martensite phase state and the reverse martensitic transformation from B19′ to B2, the TiNi alloys have an outstanding damping capacity at low frequency [21, 22]. It can be confirmed that the high damping capacity of the composite at 1 Hz was attributed to TiNif. After the prestrain treatment, the damping peak of the composite at the frequency of 10 and 20 Hz was not affected as shown in Figure 6B.
Damping capacity vs. temperature curves of the SiCp+TiNif/Al composite at different frequency (A) before and (B) after tensile deformation.
4 Discussion
4.1 Origin of the damping capacity of the SiCp+TiNif/Al composite
The damping capacity of pure Al, SiCp/Al, and SiCp+TiNif/Al as a function of strain at room temperature is shown in Figure 7. The significant increase in the damping capacity of Al alloy was related to the climbing of dislocation when the strain was more than 5×10-4. Based on the previous studies, the effect of SiCp in the Al matrix was not only to share responsibility for the external stress but also to prevent the movement of dislocation in the Al matrix. It is resulted in the increase of yield strength of the composite and the difficulty of the gliding of dislocation in the Al matrix [4]. Then, the damping capacities of the SiCp/Al and SiCp+TiNif/Al composites were worse than pure Al when the strain was more than 5×10-4.
Damping capacity vs. strain curves of the Al matrix, SiCp/Al, and SiCp+TiNif/Al composites.
When the strain was 5×10-5, both the addition of SiCp and TiNif could increase the damping capacity of the composite. As shown in Figure 8, the damping capacity of the SiCp/Al composite was better than the Al matrix but worse than the SiCp+TiNif/Al composite. When the temperature was higher than 250°C, the high damping capacity of pure Al was attributed to the slippage of grain boundaries. The damping mechanism had been confirmed by previous studies on the inelasticity behavior of grain boundaries in the polycrystalline Al [23]. However, the damping capacity of the SiCp/Al composite approximately went up linearly with the increase in temperature. The increase of damping capacity was attributed to the damping of the interface between SiCp and the Al matrix.
Damping capacity vs. temperature curves of the Al matrix, SiCp/Al, and SiCp+TiNif/Al composites.
In addition to the damping of grain boundary in the Al matrix, the main origins of the damping capacity of the SiCp+TiNif/Al composite can be identified by the individual constituents tan δNiTi and tan δint, denoting the damping capacity of TiNif and the interface, respectively. The high damping capacity of TiNif is attributed to the hysteretic effect of the martensite/variants interfaces [19] and the reverse martensitic transformation from B19′ to B2 when the temperature is close to 120°C. Furthermore, there are three other kinds of interfaces in the composites as shown in Figure 9: (a) the strong bonding interface between SiCp and Al, which was confirmed by Leng and Wu [24]; (b) the strong bonding interface between TiNif and Al, which has been discussed by previous study [25]; and (c) the weak bonding interface between SiCp and TiNif. The strong bonding interfaces are beneficial to improve the strength and modulus of the composite, but it has a little effect on the increase in the damping capacity of the composite since the relative friction and sliding in the interface is difficult, while the influence of the weak bonding interface on the properties of the composite was the opposite. Therefore, due to the high damping capacity of TiNif and the damping of the interface between SiCp and TiNif, the damping capacity of the SiCp+TiNif/Al composite was better than the TiNif/Al composite and pure Al.
Microstructure of the interface in the SiCp+TiNif/Al composite.
4.2 Effect of prestrain on the damping capacity of the SiCp+TiNif/Al composite
Due to the shape memory effect of TiNif, a compressive stress was formed in the SiCp+TiNif/Al composite after the prestrain treatment. When the temperature was higher than As, TiNif that was constrained by the Al matrix went back to its original shape. The recovery stress of TiNif rose with the increase in temperature [26]. Thus, the increase of compressive stress induced the change in the thermal expansion behavior of the composite. Figure 10 shows the strain of the composite as a function of temperature. Before the prestrain treatment, the length of the composite went up linearly with the increase in temperature. After the prestrain treatment, the change of the length with temperature could be divided into three parts: (a) when the temperature was lower than As, the length of the composite rose linearly with the increase in temperature; (b) when the temperature was higher than As but lower than Af, the length of the composite decreased with the increase in temperature; and (c) when the temperature was higher than Af, the length of the composite slowly rose with the increase in temperature.
Strain of the SiCp+TiNif/Al composite as a function of temperature in the heating process.
The decrease in the length of the composite was attributed to the compressive stress. Corresponding to the decrease of the length, the compressive stress in the composite increased. With the help of compressive stress, the movement of dislocation in the Al matrix was easier and the hysteretic effect of martensite/variants interfaces in TiNif was greater than before. It resulted in the increase of the damping capacity of the composite as shown in Figure 5. Moreover, the temperatures of the reverse martensitic transformation of TiNif were deferred by the prestrain treatment. It was confirmed by previous studies that the increase of the temperature was caused by the martensite stabilization in TiNif after the prestrain treatment [27]. From the results of the above research, the 20% SiCp+20% TiNif/Al composite fabricated by the pressure infiltration method had high strength and good damping capacity. However, much further work is needed to characterize in detail the microstructure-damping performance relationships in the composite and to optimize the design and processing of the composites with required properties.
5 Conclusions
In the study, an Al matrix composite hybrid reinforced by SiCp and TiNif was fabricated by pressure infiltration progress. Both the volume fraction of SiCp and TiNif was 20%. The damping capacity of the composite was studied in the temperature range from 30°C to 290°C. The main conclusions can be summarized as follows:
With the mix of TiNif and SiCp, the tensile strength of the composite could be significantly enhanced. Both the tension strength and the damping capacity of the composite were much better than the Al matrix.
The high damping capacity of the SiCp+TiNif/Al composite was due to the hysteretic effect of the martensite/variants interfaces in TiNif and the weak bonding interface between SiCp and TiNif.
The damping peak that was observed in the heating process was attributed to the B19′→B2 reverse martensitic transformation in TiNif.
Due to the shape memory effect of TiNif, a compressive stress was induced in the composite after tension deformation. The value of the damping peak in the heating process was much higher than before since the movement of dislocation in the Al matrix was easier.
References
[1] Wolfenden A, Wolla JM. J. Mater. Sci. 1989, 24, 3205–3212.Search in Google Scholar
[2] Zhang J, Perez RJ, Lavernia EJ. J. Mater. Sci. 1993, 28, 2395–2404.Search in Google Scholar
[3] Liu WH, Zhu JS. J. Funct. Mater. 2001, 32, 440–442.Search in Google Scholar
[4] Chawla N, Andres C, Jones JW, Allison JE. Metall. Mater. Trans. A 1998, 29, 2843–2854.10.1007/s11661-998-0325-5Search in Google Scholar
[5] Starink MJ, Wang P, Sinclair I, Gregson PJ. Acta Mater. 1999, 47, 3855–3868.Search in Google Scholar
[6] Miller WS, Humphreys FJ. Scripta Mater. 1991, 25, 33–38.Search in Google Scholar
[7] Ozyurek D, Yildirim M, Ciftci I. Sci. Eng. Compos. Mater. 2012, 19, 351–356.Search in Google Scholar
[8] Ozyurek D, Tekeli S. Sci. Eng. Compos. Mater. 2010, 17, 31–38.Search in Google Scholar
[9] Xie CL, Hailat M, Wu X, Newaz G, Taya M, Raju B. J. Eng. Mater. Technol. 2007, 129, 69–76.Search in Google Scholar
[10] Han C, Choi I, Cho K, Park I. Metall. Mater. 2000, 6, 169–175.Search in Google Scholar
[11] Lee JK, Tay M. Scripta Mater. 2004, 51, 443–447.Search in Google Scholar
[12] Raghavan J, Bartkiewicz T, Boyko S, Kupriyanov M, Rajapakse N, Yu B. Compos. Part B 2010, 41, 214–222.10.1016/j.compositesb.2009.10.009Search in Google Scholar
[13] Park YC, Kang JH, Lee JK, Lee GC, Furuya Y. Smart Mater. Struct. 2007, 16, 982–988.Search in Google Scholar
[14] Park YC, Jo YJ, Baek SH, Furuya Y. Smart Mater. Struct. 2009, 18, 055009 (9pp).10.1088/0964-1726/18/5/055009Search in Google Scholar
[15] Li DS, Zhang XP, Xiong ZP, Mai YW. J. Alloys Compd. 2010, 490, L15–L19.Search in Google Scholar
[16] Chaudhury Z, Hailat M, Liu Y, Newaz G. J. Mater. Sci. 2011, 46, 1945–1955.Search in Google Scholar
[17] Akalin O, Ezirmik KV, Urgen M, Newaz GM. J. Tribol-T ASME 2010, 132, 041603.10.1115/1.4002332Search in Google Scholar
[18] Totik Y, Sadeler R, Gavgali M, Kaymaz I. Sci. Eng. Compos. Mater. 2002, 10, 279–286.Search in Google Scholar
[19] Otsuka K, Ren X. Prog. Mater. Sci. 2005, 50, 511–678.Search in Google Scholar
[20] Porter GA, Liaw PK, Tiegs TN, Wu KH. JOM 2000, 52, 52–56.10.1007/s11837-000-0084-2Search in Google Scholar
[21] Chang SH, Wu SK. Mater. Sci. Eng. A 2007, 454–455, 379–383.10.1016/j.msea.2006.11.157Search in Google Scholar
[22] Yoshida I, Ono T, Asai M. J. Alloys Compd. 2000, 310, 339–343.Search in Google Scholar
[23] Ke TS. Phys. Rev. 1947, 71, 533.Search in Google Scholar
[24] Leng JF, Wu GH. Trans. Nonferrous Soc. China 2006, 16, 1640–1644.Search in Google Scholar
[25] Hu J, Wu GH, Zhang Q, Kang PC, Liu Y. Micron 2014, 64, 57–65.10.1016/j.micron.2014.02.014Search in Google Scholar PubMed
[26] Edwards GR, Perkins J, Johnson JM. Scripta Metall. 1975, 9, 1167–1171.Search in Google Scholar
[27] Zheng YJ, Cui LS. Intermetallics 2004, 12, 1305–1309.10.1016/j.intermet.2004.03.021Search in Google Scholar
©2016 by De Gruyter
This article is distributed under the terms of the Creative Commons Attribution Non-Commercial License, which permits unrestricted non-commercial use, distribution, and reproduction in any medium, provided the original work is properly cited.
