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An Experimental and Modeling Investigation of Al-based Nanocomposites Manufactured via Ultrasonic Cavitation and Solidification Processing
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ScienceDirect Materials Today: Proceedings 5 (2018) 16022–16031
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NN17
An Experimental and Modeling Investigation of Al-based Nanocomposites Manufactured via Ultrasonic Cavitation and Solidification Processing Yang Xuana, Daojie Zhanga, Laurentiu Nastaca,* a
The University of Alabama, Department of Metallurgical and Materials Engineering, Box 870202, Tuscaloosa, AL, 35487, USA
Abstract In the present study, A356 based nano-composites are fabricated by using the ultrasonic stirring technology (UST) in a coreless induction furnace. SiC nanoparticles were used as the reinforcement. Nanoparticles were added into the molten metal and then dispersed by ultrasonic cavitation and acoustic streaming assisted by electromagnetic stirring. The UST was also applied during the solidification process. The microstructure of the nanocomposites has been investigated by optical microscopy and scanning electron microscopy (SEM). The distribution of SiC nanoparticles in the A356 alloy matrix has also been analyzed. The SEM and energy dispersive X-ray spectroscopy (EDS) analyses showed that the matrix microstructure of the A356 alloy and the dispersion of the SiC nanoparticles into the matrix can be significantly improved when ultrasonic cavitation, induction melting and stirring, and solidification processing techniques are used together. Molecular dynamics (MD) simulations were conducted to analyze the complex interactions between the nanoparticle and the liquid/solid interface. The assumption that nanoparticles will be engulfed by the solidification front instead of being pushed was proved through MD simulations. © 2018 Elsevier Ltd. All rights reserved. Selection and/or Peer-review under responsibility of 14th International Conference on Nanosciences & Nanotechnologies (NN17). Keywords: Al-based nanocomposites; Microstructure refinement; SiC nanoparticle dispersion, Molecular dynamics simulation, Particle pushing and engulfment.
* Laurentiu Nastac. Tel: +1-205-348-4844, Fax: +1-205-348-2164 E-mail address: [email protected] 2214-7853 © 2018 Elsevier Ltd. All rights reserved. Selection and/or Peer-review under responsibility of 14th International Conference on Nanosciences & Nanotechnologies (NN17).
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1. Introduction Al metal-matrix-composites (MMCs) have been widely used in aerospace, military application, automobile and so on, since their low cost and desirable properties, such as high wear resistance, and high specific strength, etc. [110]. Because normally the use of micro-size particles could jeopardize the ductility of composites, nano-size ceramic particles has been studied as reinforcement to fabricated metal-matrix-nano-composites (MMNCs). It has been proven that MMNCs could significantly improve matrix properties while maintaining good ductility and high temperature creep resistance at the same time [11-13]. Mechanical stirring is widely used to produce MMCs. However, it is not suitable to fabricate MMNCs. Nanoparticles have high specific surface area, which means it is difficult for mechanical stirring to distribute and disperse them uniformly in the melt. In addition, nanoparticles will agglomerate again and float on the liquid surface after stirring due to their poor wettability with the metallic alloy melts [13]. UST treatment is a promising way to fabricate MMNCs, since it will induce nonlinear effects into the melt, such as cavitation and acoustic streaming, which can spread and break the nanoparticle clusters in the melt. Also, extremely high temperature (possibly larger than 5,000K) could help to enhance the wettability between nanoparticles and metal melts, which means the nanoparticles would not gather together or float on liquid surface when the UST processing is stopped [6-8, 14, 15]. It has been proved that by applying UST during the melt solidification process would help to refine the matrix microstructure [16-21]. Previous works have indicated that by applying UST for A356-Al2O3 nanocomposites around liquidus temperature could change the matrix microstructure from dendritic grains to globular grains [6, 8, 22, 23]. In this study, the effect of the UST processing on A356-SiC nanocomposites has been studied in detail. Moreover, the effect of nanoparticles size on their distribution in the matrix has been discussed. During the solidification process, the particles will be pushed, engulfed or entrapped by the solidification front, among which particle pushing will always lead to particle clustering, which is undesirable as it results in nonhomogeneous and lower macroscopic mechanical properties [24]. The particle pushing and engulfment (PEP) is a complicated phenomenon and it is affected by many factors [25] including particle size and shape, interfacial energy between particle, liquid, and solid, as well as a temperature gradient in the melt ahead of the solidification interface. However, these models only predict the behavior in the coarse (>> 1 µm) and fine particle (~ 1 µm) systems, and they don’t accurately describe the ultrafine particle (<< 1 µm) system, presumably because the models rely on continuum mechanics. But these models cannot explain the evidence in MMNCs that nanoparticles can indeed be engulfed and distributed throughout the material and are not necessarily concentrated in grain boundary or interdendritic regions. As proposed by Ferguson [26, 27], for sufficiently small particles, Brownian Motion can partially or completely counteract forces such as viscous drag, gravity and thermal/concentration gradients, thus leading to engulfment rather than pushing. Chernov [28, 29] introduced the idea of a positive disjoining pressure in the liquid film separating the particle from the solidification interface which results in a repulsive interfacial force, while the viscous drag force is still the force responsible for particle engulfment. The presence of a positive disjoining pressure leads to particle repulsion while a negative pressure would lead to entrapment. However, it is extremely difficult to investigate the interfacial properties experimentally at the atomic level. Thus, molecular dynamics (MD) simulations (atomistic simulations) conducted at the atomic level offer a good alternative in studying the interface dynamics. In this study, the open source MD program LAMMPS (Large-scale Atomic/Molecular Massively Parallel Simulator) [30] was used to conduct MD simulations to analyze the complex interactions between the nanoparticle and the liquid/solid interface. 2. Experimental Approach Aluminum alloy A356 was used as the metallic alloy matrix and SiC nanoparticles (spherical shape, average diameter range is about 45-55 nm) were used as the reinforcement. Jia et al. [6] indicated that the optimum amount of SiC is about 1.0 wt. %. An Inductotherm induction furnace was used to melt the alloy. When the melt temperature reached about 750℃ ultrasound cavitation and stirring processing was performed onto the melt. 1.0 wt. % SiC was added into the cavitation area during a 15 min time-frame. During the UST processing, an Nb ultrasonic probe (40
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mm diameter) was inserted to about 50 mm beneath the melt surface. The sketch of the UST system is shown in Fig. 1. The system parameters are: probe amplitude is 20 microns, maximum applied power is 1.5kW and frequency is18 kHz. In order to investigate the effect of the UST and the added nanoparticles on the matrix microstructure, five types of experiments were designed as follows: (a) A356 alloy with non-treatment UST processing: When A356 melt temperature reached about 750℃, held 15min and then turned off the furnace, and let the melt cooling in the furnace to the room temperature. (b) A356 alloy with 15 min UST at 750℃: When A356 melt temperature reached about 750℃, applied UST at 750℃ about 15min, then turned off the furnace, and let the melt cooling in the furnace to the room temperature. (c) A356 alloy with 15 min UST at 750℃ and UST during solidification: When A356 melt temperature reached about 750℃, applied UST at 750℃ about 15min, then turned off the furnace, keep treated melt with UST for about 3min, then let the melt solidify in the furnace. (d) Nanocomposites with non-treatment UST processing during solidification: when A356 melt reached about 750℃, added 1 wt.% SiC nanoparticles into the melt and treated with UST about 15min at 750℃, then turned off the furnace, and let the melt cooling in the furnace to the room temperature. (e) Nanocomposites with UST processing during solidification: similar to experiment 4, but after the furnace was turned off, melt was continuously treated with UST for about 3 min, then let the melt solidify in the furnace. The liquidus temperature of A356 is about 614℃. It has been identified that the added nanoparticles will not change the melt liquidus temperature significantly [23]. Accordingly, the UST was applied to a melt temperature of about 136℃ higher than the melt liquidus temperature. The cast ingots were cylindrical in shape with a diameter of 85 mm and a height of 120 mm. To avoid the effects of different locations (cooling rates) on the microstructure [8], all samples were taken from the middle location of the ingots. During the UST processing, the ultrasonic probe was inserted to about 50 mm beneath the melt, thus all the samples were taken from a section that is below the probe. Optical microscopy (Nikon EPIPHOT 200) photos and SEM (JEOL 7000) photos were prepared by grinding with 240, 600, 800, 1200 sand papers followed by polishing with 9 µm, 3 µm, and 1µm water based diamond suspensions.
Fig. 1. Sketch of the UST system.
3. Molecular Dynamics Modeling Approach The dimensions of the MD geometry in the x, y, and z-direction are 40 nm × 40 nm × 40 nm. The diameter of the SiC nanoparticle is 20 nm. The simulated structure of pure aluminum is a face centered cubic (FCC) structure with a lattice parameter of 4.05 Å and with the <100> directions coincident with the Cartesian coordinates, since planes of looser packing, such as {100}, are better able to accommodate an atom that leaves the liquid to join the solid than a
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closer packed plane, such as {111} [31]. The nanoparticle consists of a cubic crystalline polytype of silicon carbide (3C-SiC) with a lattice parameter of 4.36 Å [32]. In total, 3,900,208 atoms are generated. The solidification process was simulated by using periodic boundary conditions that were applied in three coordinate directions. The simulation starts with a solid SiC region, a solid Al region (thickness 5 nm in z direction) and a liquid Al region separated by a nominally flat liquid-solid boundary. The two-phase Al system is established by holding fixed the atoms in solid Al region of the original crystalline lattice and melting the liquid Al region by raising the temperature to some temperature above the melting point. Upon solidification, both the solid Al atoms at the bottom and the liquid Al atoms at the top can move freely into each other because of the periodic boundary conditions. The velocity of the atoms is then rescaled to some temperature below the melting point. Due to the density difference between solid and liquid, constant volume simulations will lead to a gradual buildup of pressure during crystallization which eventually halts solidification [33]. Instead, constant pressure conditions are required with the pressure maintained at zero throughout the run with a time step of 1 fs. During solidification the solid atoms at the bottom which were initially fixed in position are allowed to move. The simulation is terminated when complete solidification is achieved (about 600 ps). The EAM potential model developed by Mishin et al. [34] was used for the Al matrix. The interatomic potentials for Al are available from [35]. The Tersoff many-body potential [36, 37] was used for the SiC nanoparticles. The required equations along with the energy constants and corresponding cut-off distances for silicon carbide are presented in [36, 37]. To describe the interface between the Al matrix and the SiC nanoparticle, a two-body pairwise Morse [38] potential (two-body interactions between Al–Si and Al–C) was used [39]. The required parameters can be found in [39]. More details regarding the MD simulation procedure are provided in [40, 41]. 4. Results and Discussion 4.1. Experiments Figure 2 shows a comparison of experimental microstructures taken from all five different experiments described in the previous section. The matrix microstructure shown in Figs. 2(a), 2(b) and 2(d) is dendritic grain morphology, which is significantly different from the microstructures shown in Figs. 2(c) and 2(e), which exhibit a globular grain morphology. As shown in Figs. 2(a) and 2(d), for non-UST treatment of the melt during solidification, the microstructure is a dendritic grain morphology, as expected. However, for 3 min UST treatment of the melt during solidification, as shown in Figs. 2(c) and 2(e), the microstructure was refined significantly and modified from a dendritic to globular grain morphology. Moreover, it can be seen that SiC nanoparticles are distributed quite well in the matrix, and there are no micro-size nanoparticle clusters in the matrix, which confirmed that 15min UST at 750℃ is a suitable method to fabricate Al based MMNCs. Additionally, Fig. 2(b) indicates that by applying UST at high melt superheats there are no significant effects of the UST on the matrix microstructure. The results shown in Figs. 2(b), 2(d) and 2(e) indicate that applying 3min UST during the solidification process but not adding SiC nanoparticles is the main reason for the microstructure improvement, since only Fig. 2(e) illustrates globular grains. Thus, by applying the UST during the solidification (Figs. 2(b) and 2(e)) the matrix microstructure is refined and modified from the dendritic to globular grain morphology. This can be explained as follows: First, ultrasound will increase the nucleation potential in the melt during solidification. Because UST will cause the temperature and pressure characteristics of the melt to change periodically at high frequencies. As a result, this will increase the local number of nuclei into the melt. Moreover, strong convection produced by UST will promote the diffusion of solute, which will also increase the number of nuclei into the melt. The second reason is because shock waves produced by ultrasonic cavitation could break the dendrite tips during solidification, which will contribute to the microstructure refinement and modification [8, 42-45]. When applying UST at a temperature much higher than the melt liquidus temperature (more than 130℃ in this study), the ultrasonic cavitation could increase the local number of nuclei momentarily, but since the melt temperature is much higher than the liquidus temperature, these nuclei will remelt. Thus, this explains why when only applying the UST for 15min at 750℃ will not change the microstructure significantly, as shown in Figs. 2(b) and 2(d).
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(a)
(b)
(c)
(d)
(e) Fig. 2. Optical micrographs of samples at 50X (a) A356 alloy with non-treatment UST processing (b) A356 alloy with 15 min UST at 750℃ (c) A356 alloy with 15 min UST at 750℃ and UST during solidification (d) Nanocomposites with non-treatment UST processing (e) Nanocomposites with UST processing during solidification.
Figure 3 shows the SEM analysis result of the sample e, which is A356- 1.0 wt. % SiC nanocomposite processed with UST during solidification. The SEM-EDS analyses of the experimental results indicate that SiC nanoparticles are well dispersed into the matrix during the solidification process. Some nanoparticles have been found in the grain boundary area (Figs. 3(a) and 3(b)) and some of them in the matrix (Figs. 3(a) and 3(c)). The nanoparticle shown in Figure 3 have spherical shape with a diameter range of about 40-75 nm, which is similar with the SiC nanoparticles added into the melt. In EDS mapping photo (Fig. 3(d)), the dark red color represents C (carbon), the green color
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represents Al (aluminum). The EDS results indicate that some C is presented at the grain boundary area. By combining the EDS results with the SEM photos, it can be concluded that some of the SiC nanoparticles are gathering around the grain boundary area. At the same time, some SiC nanoparticles are dispersed within the grains.
(a)
(b)
(c)
(d)
Fig. 3. SEM results and EDS mapping photo of the sample processed via 3 min UST+1.0 mass% SiC: (a) SEM picture at X600; (b) SEM picture of red rectangular area showing SiC particles shown along and near the grain boundary at X110,000; (c) SEM picture of red rectangular area in the grain area showing SiC particles in the grain at X110,000; and (d) EDS mapping results at X40.
4.2. Simulation The two-phase Al system is established by holding fixed the atoms in solid Al region of the original crystalline lattice and melting the liquid Al region by raising the temperature to some temperature above the melting point. The whole system is equilibrated after 40ps, then the initial solid Al atoms are fixed at some temperature (500 K) below the melting point, and Si, C, and liquid Al atoms with will release heat through the solid atoms. Figure 4 shows the solid/liquid status of the system with a 20 nm SiC nanoparticle after 100 ps, 200 ps, 300 ps, 400 ps, 500 ps, and 600 ps, respectively. As it can be seen from Fig. 4(a), when the liquid/solid interfaces are far away from the SiC nanoparticle, both of the interfaces at the top and bottom are flat. When the interface at the bottom approaches the SiC nanoparticle, a trough is formed below the particle (see Fig. 4(b)). As time goes on, the interface at the bottom passes through the particle and the interface at the top approaches the particle (See Figs. 4(c)-(d)). Finally, these two interfaces meet each other, become one, and disappear (see Fig. 4(f)). Figure 5 shows the position of the 20 nm SiC nanoparticle with time. As it can be seen, the horizontal movement (x and y) of the nanoparticle is insignificant. But in z direction (vertical), the particle moves down towards the
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liquid/solid interface at the bottom in the beginning, and when the interface from the top approaches it, the particle moves up towards that interface again. This MD simulation confirmed that during the solidification process, the SiC nanoparticle will be engulfed by the solidification front instead of being pushed.
(a)
(b)
(c)
(d)
(e)
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Fig. 4. Solid/liquid status of the system with a 20 nm SiC nanoparticle after (a) 100 ps, (b) 200 ps, (c) 300 ps, (d) 400 ps, (e) 500 ps, (f) 600 ps.
Fig. 5. Position of a 20 nm SiC nanoparticle during solidification.
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4.3. PEP modeling Chernov et al. [28, 29] describes the critical velocity (Vcr) for PEP when the size of the particle is smaller than 1 mm as [46]:
Vcr
1/ 3 0.14 B 2 / 3 SL rp4 / 3
(1)
where B is a constant (fitting parameter) suggested to be 10-21 J, SL is the surface, is the viscosity, and rp is the radius of the particle. For Al alloys, SL = 0.093 J/m2, = 0.00105 [46]. Figure 6 presents a comparison of the Chernov’s model (PEP-CM) (Eq. (1)) using B = 10-21 J and Chernov’s modified model (PEP-CMM) (Eq. (1)) using a B = 10-23 J, which was corrected based on both the experimental and the simulation results. The MD simulation result for 10 nm nanoparticles is provided in [40]. The MD simulation result for 20 nm nanoparticles is shown in Figs. 4 and 5. Both MD simulations predict engulfment of the nanoparticles. The experimental results for 40nm and 80 nm nanoparticles were taken from Figs. 3(b) and 3(c), which also show engulfment of nanoparticles. The PEP-CM model predicts pushing while the PEP-CMM predicts engulfment for all nanoparticle sizes (10nm, 20nm, 40 mm, and 80 nm) and for a critical velocity of about 10-4 m/s, which is a representative value for the current solidification conditions in the furnace.
Fig. 6. Comparison between Chernov’s model (PEP-CM) and Chernov’s modified model (PEP-CMM).
5. Conclusions This study determined that the UST treatment during solidification and not the addition of SiC nanoparticles into the matrix plays a key role to refine and modify the matrix microstructure from dendritic to globular grain morphology. Also, the application of UST at high superheats would not considerably refine the microstructure. The SEM and EDS analyses of the experimental results indicated that SiC nanoparticles are well dispersed into the matrix during the solidification process. Some SiC nanoparticles are gathering around the grain boundary area.
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The complex interactions between a SiC nanoparticle and the Al liquid/solid interface were analyzed with a 3D molecular dynamics model using LAMMPS on the High Performance Computing Cluster (HPCC) at the University of Alabama. Molecular dynamics simulation results showed that nanoparticles with diameters less than about 20 nm would be engulfed by the solidification front instead of being pushed, which is beneficial for nano-dispersion. Acknowledgements The authors wish to express thanks to the China Scholarship Council (CSC) for the financial support. References [1] B. Su, H.G.Yan, G. Chen, J.L. Shi, J.H. Chen, P.L. Zeng, Mat Sci Eng a-Struct. 527(24-25), (2010) 6660-6665. [2] D. Bozic, B. Dimcic, O. Dimcic, J. Stasic, V. Rajkovic, Mater Design. 31(1), (2010) 134-41. [3] A.B.Elshalakany, T.A. Osman, A. Khattab, B. Azzam, M. Zaki, J Nanomater. 2014 (2014) 1-14. [4] X.D. Liu, S.A. Jia, L. Nastac, Int J Metalca St. 8(3), (2014) 51-58. [5] Jia, Zhang, L. Nastac, J Mater Eng Perform. 24(6), (2015) 2225-2233. [6] S.A. Jia, D.J. Zhang, Y. Xuan, L. Nastac, Appl Acoust. 103 (2016) 226-231. [7] S.A. Jia, Y. Xuan, L. Nastac, P.G. Allison, T.W. Rushing, Int J Cast Metal Res. 29(5), (2016) 286-289. [8] Y. Xuan S.A. Jia, L. Nastac, High Temperature Materials and Processing. 36(4), (2017) 381-388. [9] C.S. Goh, J. Wei, L.C. Lee, A. Gupta, Mat Sci Eng a-Struct. 2006;423(1-2):153-156. [10] S.M.S. Reihani, Mater Design. 27(3), (2006) 216-222. [11] M.K. Akbari, O. Mirzaee, H.R. Baharvandi, Mater Design. 46 (2013) 199-205. [12] M.K. Akbari, H.R. Baharvandi, O. Mirzaee, Compos Part B-Eng. 55 (2013) 426-432. [13] Y. Yang, J. Lan, X.C. Li, Mat Sci Eng a-Struct. 380(1-2), (2004) 378-383. [14] X. D. Liu. Master degree thesis, The University of Alabama; Tuscaloosa, Al, USA, 2013. [15] J.C. Yan, Z.W. Xu, L. Shi, X. Ma, S.Q. Yang, Mater Design. 32(1), (2011) 343-347. [16] A. Ramirez, M. Qian, B. Davis, T. Wilks, Int J Cast Metal Res. 22(1-4), (2009) 260-263. [17] M. Qian, A.Ramirez. In: Czerwinski F., editor. Magnesium Alloys - Design, Processing and Properties: InTech; 2011. pp. 163-186. [18] H.K. Feng, S.R. Yu, Y.L. Li, L.Y. Gong, J Mater Process Tech. 208(1-3), (2008) 330-335. [19] J.I. Youn, Y.J. Kim, Jpn J Appl Phys. 48 (2009) 07GM14. [20] J.G. Jung, S.H. Lee, J.M. Lee, Y.H. Cho, S.H. Kim, W.H. Yoon, Mat Sci Eng a-Struct. 669 (2016) 187-195. [21] J.I. Youn, Y.K. Lee, Y.J. Kim, J.W. Park, Jpn J Appl Phys. 55 (2016) 07KE10). [22] Y. Xuan, L.Nastac, Ultrasonics. online 1 July 2017, https://doi.org/10.1016/j.ultras.2017.06.023. [23] Y. Xuan, L. Nastac, TMS 2017 146th Annual Meeting & Exhibition Supplemental Proceedings. San Deigo: 2017. Part of MMMS book series, pp. 297-303. [24] S. Naher, D. Brabazon, L. Looney, Composites Part A: Applied Science and Manufacturing. 38 (2007) 719-729. [25] D.M. Stefanescu, B.K. Dhindaw, S.A. Kacar, A. Moitra, Metallurgical Transactions A. 19 (1988) 2847-2855. [26] J.B. Ferguson, B.F. Schultz, P.K. Rohatgi, C.S. Kim, Light Metals 2014. John Wiley & Sons, Inc., 2014. pp. 1383-1388. [27] J.B. Ferguson, B. Schultz, P. Rohatgi, C.S. Kim, Met. Mater. Int. 20 (2014) 747-755. [28] A. Chernov, D. Temkin, A. Mel’nikova, Sov. Phys. Crystallogr. 21 (1976) 369-373. [29] A. Chernov, D. Temkin, A. Mel’nikova, Sov. Phys. Crystallogr. 22 (1977) 656-658. [30] S. Plimpton. J. Comput. Phys. 117 (1995) 1-19. [31] C.R. Dandekar, Y.C. Shin, Compos Part A: Appl Sci Manuf. 42 (4), (2011) 355-363. [32] B. Chalmers, Trans. AIME. 200 (1954), pp. 519. [33] J. Hoyt, M. Asta, A. Karma, Interface Science 10 (2002) 181-189. [34] Y. Mishin, D. Farkas, M.J. Mehl, D.A. Papaconstantopoulos, Physical Review B, 59 (1999) 3393-3407. [35] J. Winey, A. Kubota, Y. Gupta, Modelling Simul. Master. Sci. Eng. 17 (2009) 055004. [36] J. Tersoff, Phys. Rev. B. 39 (1989) 5566-5568. [37] J. Tersoff, Phys. Rev. Lett. 64 (1990) 1757-1760. [38] P.M. Morse, Phys. Rev. 34 (1929) 57-64. [39] H. Zhao, N. Chen, Y. Long, J. Phys. Condens. Matter. 21 (2009) 225002.
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ScienceDirect Materials Today: Proceedings 5 (2018) 16022–16031
www.materialstoday.com/proceedings
NN17
An Experimental and Modeling Investigation of Al-based Nanocomposites Manufactured via Ultrasonic Cavitation and Solidification Processing Yang Xuana, Daojie Zhanga, Laurentiu Nastaca,* a
The University of Alabama, Department of Metallurgical and Materials Engineering, Box 870202, Tuscaloosa, AL, 35487, USA
Abstract In the present study, A356 based nano-composites are fabricated by using the ultrasonic stirring technology (UST) in a coreless induction furnace. SiC nanoparticles were used as the reinforcement. Nanoparticles were added into the molten metal and then dispersed by ultrasonic cavitation and acoustic streaming assisted by electromagnetic stirring. The UST was also applied during the solidification process. The microstructure of the nanocomposites has been investigated by optical microscopy and scanning electron microscopy (SEM). The distribution of SiC nanoparticles in the A356 alloy matrix has also been analyzed. The SEM and energy dispersive X-ray spectroscopy (EDS) analyses showed that the matrix microstructure of the A356 alloy and the dispersion of the SiC nanoparticles into the matrix can be significantly improved when ultrasonic cavitation, induction melting and stirring, and solidification processing techniques are used together. Molecular dynamics (MD) simulations were conducted to analyze the complex interactions between the nanoparticle and the liquid/solid interface. The assumption that nanoparticles will be engulfed by the solidification front instead of being pushed was proved through MD simulations. © 2018 Elsevier Ltd. All rights reserved. Selection and/or Peer-review under responsibility of 14th International Conference on Nanosciences & Nanotechnologies (NN17). Keywords: Al-based nanocomposites; Microstructure refinement; SiC nanoparticle dispersion, Molecular dynamics simulation, Particle pushing and engulfment.
* Laurentiu Nastac. Tel: +1-205-348-4844, Fax: +1-205-348-2164 E-mail address: [email protected] 2214-7853 © 2018 Elsevier Ltd. All rights reserved. Selection and/or Peer-review under responsibility of 14th International Conference on Nanosciences & Nanotechnologies (NN17).
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1. Introduction Al metal-matrix-composites (MMCs) have been widely used in aerospace, military application, automobile and so on, since their low cost and desirable properties, such as high wear resistance, and high specific strength, etc. [110]. Because normally the use of micro-size particles could jeopardize the ductility of composites, nano-size ceramic particles has been studied as reinforcement to fabricated metal-matrix-nano-composites (MMNCs). It has been proven that MMNCs could significantly improve matrix properties while maintaining good ductility and high temperature creep resistance at the same time [11-13]. Mechanical stirring is widely used to produce MMCs. However, it is not suitable to fabricate MMNCs. Nanoparticles have high specific surface area, which means it is difficult for mechanical stirring to distribute and disperse them uniformly in the melt. In addition, nanoparticles will agglomerate again and float on the liquid surface after stirring due to their poor wettability with the metallic alloy melts [13]. UST treatment is a promising way to fabricate MMNCs, since it will induce nonlinear effects into the melt, such as cavitation and acoustic streaming, which can spread and break the nanoparticle clusters in the melt. Also, extremely high temperature (possibly larger than 5,000K) could help to enhance the wettability between nanoparticles and metal melts, which means the nanoparticles would not gather together or float on liquid surface when the UST processing is stopped [6-8, 14, 15]. It has been proved that by applying UST during the melt solidification process would help to refine the matrix microstructure [16-21]. Previous works have indicated that by applying UST for A356-Al2O3 nanocomposites around liquidus temperature could change the matrix microstructure from dendritic grains to globular grains [6, 8, 22, 23]. In this study, the effect of the UST processing on A356-SiC nanocomposites has been studied in detail. Moreover, the effect of nanoparticles size on their distribution in the matrix has been discussed. During the solidification process, the particles will be pushed, engulfed or entrapped by the solidification front, among which particle pushing will always lead to particle clustering, which is undesirable as it results in nonhomogeneous and lower macroscopic mechanical properties [24]. The particle pushing and engulfment (PEP) is a complicated phenomenon and it is affected by many factors [25] including particle size and shape, interfacial energy between particle, liquid, and solid, as well as a temperature gradient in the melt ahead of the solidification interface. However, these models only predict the behavior in the coarse (>> 1 µm) and fine particle (~ 1 µm) systems, and they don’t accurately describe the ultrafine particle (<< 1 µm) system, presumably because the models rely on continuum mechanics. But these models cannot explain the evidence in MMNCs that nanoparticles can indeed be engulfed and distributed throughout the material and are not necessarily concentrated in grain boundary or interdendritic regions. As proposed by Ferguson [26, 27], for sufficiently small particles, Brownian Motion can partially or completely counteract forces such as viscous drag, gravity and thermal/concentration gradients, thus leading to engulfment rather than pushing. Chernov [28, 29] introduced the idea of a positive disjoining pressure in the liquid film separating the particle from the solidification interface which results in a repulsive interfacial force, while the viscous drag force is still the force responsible for particle engulfment. The presence of a positive disjoining pressure leads to particle repulsion while a negative pressure would lead to entrapment. However, it is extremely difficult to investigate the interfacial properties experimentally at the atomic level. Thus, molecular dynamics (MD) simulations (atomistic simulations) conducted at the atomic level offer a good alternative in studying the interface dynamics. In this study, the open source MD program LAMMPS (Large-scale Atomic/Molecular Massively Parallel Simulator) [30] was used to conduct MD simulations to analyze the complex interactions between the nanoparticle and the liquid/solid interface. 2. Experimental Approach Aluminum alloy A356 was used as the metallic alloy matrix and SiC nanoparticles (spherical shape, average diameter range is about 45-55 nm) were used as the reinforcement. Jia et al. [6] indicated that the optimum amount of SiC is about 1.0 wt. %. An Inductotherm induction furnace was used to melt the alloy. When the melt temperature reached about 750℃ ultrasound cavitation and stirring processing was performed onto the melt. 1.0 wt. % SiC was added into the cavitation area during a 15 min time-frame. During the UST processing, an Nb ultrasonic probe (40
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mm diameter) was inserted to about 50 mm beneath the melt surface. The sketch of the UST system is shown in Fig. 1. The system parameters are: probe amplitude is 20 microns, maximum applied power is 1.5kW and frequency is18 kHz. In order to investigate the effect of the UST and the added nanoparticles on the matrix microstructure, five types of experiments were designed as follows: (a) A356 alloy with non-treatment UST processing: When A356 melt temperature reached about 750℃, held 15min and then turned off the furnace, and let the melt cooling in the furnace to the room temperature. (b) A356 alloy with 15 min UST at 750℃: When A356 melt temperature reached about 750℃, applied UST at 750℃ about 15min, then turned off the furnace, and let the melt cooling in the furnace to the room temperature. (c) A356 alloy with 15 min UST at 750℃ and UST during solidification: When A356 melt temperature reached about 750℃, applied UST at 750℃ about 15min, then turned off the furnace, keep treated melt with UST for about 3min, then let the melt solidify in the furnace. (d) Nanocomposites with non-treatment UST processing during solidification: when A356 melt reached about 750℃, added 1 wt.% SiC nanoparticles into the melt and treated with UST about 15min at 750℃, then turned off the furnace, and let the melt cooling in the furnace to the room temperature. (e) Nanocomposites with UST processing during solidification: similar to experiment 4, but after the furnace was turned off, melt was continuously treated with UST for about 3 min, then let the melt solidify in the furnace. The liquidus temperature of A356 is about 614℃. It has been identified that the added nanoparticles will not change the melt liquidus temperature significantly [23]. Accordingly, the UST was applied to a melt temperature of about 136℃ higher than the melt liquidus temperature. The cast ingots were cylindrical in shape with a diameter of 85 mm and a height of 120 mm. To avoid the effects of different locations (cooling rates) on the microstructure [8], all samples were taken from the middle location of the ingots. During the UST processing, the ultrasonic probe was inserted to about 50 mm beneath the melt, thus all the samples were taken from a section that is below the probe. Optical microscopy (Nikon EPIPHOT 200) photos and SEM (JEOL 7000) photos were prepared by grinding with 240, 600, 800, 1200 sand papers followed by polishing with 9 µm, 3 µm, and 1µm water based diamond suspensions.
Fig. 1. Sketch of the UST system.
3. Molecular Dynamics Modeling Approach The dimensions of the MD geometry in the x, y, and z-direction are 40 nm × 40 nm × 40 nm. The diameter of the SiC nanoparticle is 20 nm. The simulated structure of pure aluminum is a face centered cubic (FCC) structure with a lattice parameter of 4.05 Å and with the <100> directions coincident with the Cartesian coordinates, since planes of looser packing, such as {100}, are better able to accommodate an atom that leaves the liquid to join the solid than a
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closer packed plane, such as {111} [31]. The nanoparticle consists of a cubic crystalline polytype of silicon carbide (3C-SiC) with a lattice parameter of 4.36 Å [32]. In total, 3,900,208 atoms are generated. The solidification process was simulated by using periodic boundary conditions that were applied in three coordinate directions. The simulation starts with a solid SiC region, a solid Al region (thickness 5 nm in z direction) and a liquid Al region separated by a nominally flat liquid-solid boundary. The two-phase Al system is established by holding fixed the atoms in solid Al region of the original crystalline lattice and melting the liquid Al region by raising the temperature to some temperature above the melting point. Upon solidification, both the solid Al atoms at the bottom and the liquid Al atoms at the top can move freely into each other because of the periodic boundary conditions. The velocity of the atoms is then rescaled to some temperature below the melting point. Due to the density difference between solid and liquid, constant volume simulations will lead to a gradual buildup of pressure during crystallization which eventually halts solidification [33]. Instead, constant pressure conditions are required with the pressure maintained at zero throughout the run with a time step of 1 fs. During solidification the solid atoms at the bottom which were initially fixed in position are allowed to move. The simulation is terminated when complete solidification is achieved (about 600 ps). The EAM potential model developed by Mishin et al. [34] was used for the Al matrix. The interatomic potentials for Al are available from [35]. The Tersoff many-body potential [36, 37] was used for the SiC nanoparticles. The required equations along with the energy constants and corresponding cut-off distances for silicon carbide are presented in [36, 37]. To describe the interface between the Al matrix and the SiC nanoparticle, a two-body pairwise Morse [38] potential (two-body interactions between Al–Si and Al–C) was used [39]. The required parameters can be found in [39]. More details regarding the MD simulation procedure are provided in [40, 41]. 4. Results and Discussion 4.1. Experiments Figure 2 shows a comparison of experimental microstructures taken from all five different experiments described in the previous section. The matrix microstructure shown in Figs. 2(a), 2(b) and 2(d) is dendritic grain morphology, which is significantly different from the microstructures shown in Figs. 2(c) and 2(e), which exhibit a globular grain morphology. As shown in Figs. 2(a) and 2(d), for non-UST treatment of the melt during solidification, the microstructure is a dendritic grain morphology, as expected. However, for 3 min UST treatment of the melt during solidification, as shown in Figs. 2(c) and 2(e), the microstructure was refined significantly and modified from a dendritic to globular grain morphology. Moreover, it can be seen that SiC nanoparticles are distributed quite well in the matrix, and there are no micro-size nanoparticle clusters in the matrix, which confirmed that 15min UST at 750℃ is a suitable method to fabricate Al based MMNCs. Additionally, Fig. 2(b) indicates that by applying UST at high melt superheats there are no significant effects of the UST on the matrix microstructure. The results shown in Figs. 2(b), 2(d) and 2(e) indicate that applying 3min UST during the solidification process but not adding SiC nanoparticles is the main reason for the microstructure improvement, since only Fig. 2(e) illustrates globular grains. Thus, by applying the UST during the solidification (Figs. 2(b) and 2(e)) the matrix microstructure is refined and modified from the dendritic to globular grain morphology. This can be explained as follows: First, ultrasound will increase the nucleation potential in the melt during solidification. Because UST will cause the temperature and pressure characteristics of the melt to change periodically at high frequencies. As a result, this will increase the local number of nuclei into the melt. Moreover, strong convection produced by UST will promote the diffusion of solute, which will also increase the number of nuclei into the melt. The second reason is because shock waves produced by ultrasonic cavitation could break the dendrite tips during solidification, which will contribute to the microstructure refinement and modification [8, 42-45]. When applying UST at a temperature much higher than the melt liquidus temperature (more than 130℃ in this study), the ultrasonic cavitation could increase the local number of nuclei momentarily, but since the melt temperature is much higher than the liquidus temperature, these nuclei will remelt. Thus, this explains why when only applying the UST for 15min at 750℃ will not change the microstructure significantly, as shown in Figs. 2(b) and 2(d).
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(a)
(b)
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(e) Fig. 2. Optical micrographs of samples at 50X (a) A356 alloy with non-treatment UST processing (b) A356 alloy with 15 min UST at 750℃ (c) A356 alloy with 15 min UST at 750℃ and UST during solidification (d) Nanocomposites with non-treatment UST processing (e) Nanocomposites with UST processing during solidification.
Figure 3 shows the SEM analysis result of the sample e, which is A356- 1.0 wt. % SiC nanocomposite processed with UST during solidification. The SEM-EDS analyses of the experimental results indicate that SiC nanoparticles are well dispersed into the matrix during the solidification process. Some nanoparticles have been found in the grain boundary area (Figs. 3(a) and 3(b)) and some of them in the matrix (Figs. 3(a) and 3(c)). The nanoparticle shown in Figure 3 have spherical shape with a diameter range of about 40-75 nm, which is similar with the SiC nanoparticles added into the melt. In EDS mapping photo (Fig. 3(d)), the dark red color represents C (carbon), the green color
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represents Al (aluminum). The EDS results indicate that some C is presented at the grain boundary area. By combining the EDS results with the SEM photos, it can be concluded that some of the SiC nanoparticles are gathering around the grain boundary area. At the same time, some SiC nanoparticles are dispersed within the grains.
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Fig. 3. SEM results and EDS mapping photo of the sample processed via 3 min UST+1.0 mass% SiC: (a) SEM picture at X600; (b) SEM picture of red rectangular area showing SiC particles shown along and near the grain boundary at X110,000; (c) SEM picture of red rectangular area in the grain area showing SiC particles in the grain at X110,000; and (d) EDS mapping results at X40.
4.2. Simulation The two-phase Al system is established by holding fixed the atoms in solid Al region of the original crystalline lattice and melting the liquid Al region by raising the temperature to some temperature above the melting point. The whole system is equilibrated after 40ps, then the initial solid Al atoms are fixed at some temperature (500 K) below the melting point, and Si, C, and liquid Al atoms with will release heat through the solid atoms. Figure 4 shows the solid/liquid status of the system with a 20 nm SiC nanoparticle after 100 ps, 200 ps, 300 ps, 400 ps, 500 ps, and 600 ps, respectively. As it can be seen from Fig. 4(a), when the liquid/solid interfaces are far away from the SiC nanoparticle, both of the interfaces at the top and bottom are flat. When the interface at the bottom approaches the SiC nanoparticle, a trough is formed below the particle (see Fig. 4(b)). As time goes on, the interface at the bottom passes through the particle and the interface at the top approaches the particle (See Figs. 4(c)-(d)). Finally, these two interfaces meet each other, become one, and disappear (see Fig. 4(f)). Figure 5 shows the position of the 20 nm SiC nanoparticle with time. As it can be seen, the horizontal movement (x and y) of the nanoparticle is insignificant. But in z direction (vertical), the particle moves down towards the
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liquid/solid interface at the bottom in the beginning, and when the interface from the top approaches it, the particle moves up towards that interface again. This MD simulation confirmed that during the solidification process, the SiC nanoparticle will be engulfed by the solidification front instead of being pushed.
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Fig. 4. Solid/liquid status of the system with a 20 nm SiC nanoparticle after (a) 100 ps, (b) 200 ps, (c) 300 ps, (d) 400 ps, (e) 500 ps, (f) 600 ps.
Fig. 5. Position of a 20 nm SiC nanoparticle during solidification.
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4.3. PEP modeling Chernov et al. [28, 29] describes the critical velocity (Vcr) for PEP when the size of the particle is smaller than 1 mm as [46]:
Vcr
1/ 3 0.14 B 2 / 3 SL rp4 / 3
(1)
where B is a constant (fitting parameter) suggested to be 10-21 J, SL is the surface, is the viscosity, and rp is the radius of the particle. For Al alloys, SL = 0.093 J/m2, = 0.00105 [46]. Figure 6 presents a comparison of the Chernov’s model (PEP-CM) (Eq. (1)) using B = 10-21 J and Chernov’s modified model (PEP-CMM) (Eq. (1)) using a B = 10-23 J, which was corrected based on both the experimental and the simulation results. The MD simulation result for 10 nm nanoparticles is provided in [40]. The MD simulation result for 20 nm nanoparticles is shown in Figs. 4 and 5. Both MD simulations predict engulfment of the nanoparticles. The experimental results for 40nm and 80 nm nanoparticles were taken from Figs. 3(b) and 3(c), which also show engulfment of nanoparticles. The PEP-CM model predicts pushing while the PEP-CMM predicts engulfment for all nanoparticle sizes (10nm, 20nm, 40 mm, and 80 nm) and for a critical velocity of about 10-4 m/s, which is a representative value for the current solidification conditions in the furnace.
Fig. 6. Comparison between Chernov’s model (PEP-CM) and Chernov’s modified model (PEP-CMM).
5. Conclusions This study determined that the UST treatment during solidification and not the addition of SiC nanoparticles into the matrix plays a key role to refine and modify the matrix microstructure from dendritic to globular grain morphology. Also, the application of UST at high superheats would not considerably refine the microstructure. The SEM and EDS analyses of the experimental results indicated that SiC nanoparticles are well dispersed into the matrix during the solidification process. Some SiC nanoparticles are gathering around the grain boundary area.
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The complex interactions between a SiC nanoparticle and the Al liquid/solid interface were analyzed with a 3D molecular dynamics model using LAMMPS on the High Performance Computing Cluster (HPCC) at the University of Alabama. Molecular dynamics simulation results showed that nanoparticles with diameters less than about 20 nm would be engulfed by the solidification front instead of being pushed, which is beneficial for nano-dispersion. Acknowledgements The authors wish to express thanks to the China Scholarship Council (CSC) for the financial support. References [1] B. Su, H.G.Yan, G. Chen, J.L. Shi, J.H. Chen, P.L. Zeng, Mat Sci Eng a-Struct. 527(24-25), (2010) 6660-6665. [2] D. Bozic, B. Dimcic, O. Dimcic, J. Stasic, V. Rajkovic, Mater Design. 31(1), (2010) 134-41. [3] A.B.Elshalakany, T.A. Osman, A. Khattab, B. Azzam, M. Zaki, J Nanomater. 2014 (2014) 1-14. [4] X.D. Liu, S.A. Jia, L. Nastac, Int J Metalca St. 8(3), (2014) 51-58. [5] Jia, Zhang, L. Nastac, J Mater Eng Perform. 24(6), (2015) 2225-2233. [6] S.A. Jia, D.J. Zhang, Y. Xuan, L. Nastac, Appl Acoust. 103 (2016) 226-231. [7] S.A. Jia, Y. Xuan, L. Nastac, P.G. Allison, T.W. Rushing, Int J Cast Metal Res. 29(5), (2016) 286-289. [8] Y. Xuan S.A. Jia, L. Nastac, High Temperature Materials and Processing. 36(4), (2017) 381-388. [9] C.S. Goh, J. Wei, L.C. Lee, A. Gupta, Mat Sci Eng a-Struct. 2006;423(1-2):153-156. [10] S.M.S. Reihani, Mater Design. 27(3), (2006) 216-222. [11] M.K. Akbari, O. Mirzaee, H.R. Baharvandi, Mater Design. 46 (2013) 199-205. [12] M.K. Akbari, H.R. Baharvandi, O. Mirzaee, Compos Part B-Eng. 55 (2013) 426-432. [13] Y. Yang, J. Lan, X.C. Li, Mat Sci Eng a-Struct. 380(1-2), (2004) 378-383. [14] X. D. Liu. Master degree thesis, The University of Alabama; Tuscaloosa, Al, USA, 2013. [15] J.C. Yan, Z.W. Xu, L. Shi, X. Ma, S.Q. Yang, Mater Design. 32(1), (2011) 343-347. [16] A. Ramirez, M. Qian, B. Davis, T. Wilks, Int J Cast Metal Res. 22(1-4), (2009) 260-263. [17] M. Qian, A.Ramirez. In: Czerwinski F., editor. Magnesium Alloys - Design, Processing and Properties: InTech; 2011. pp. 163-186. [18] H.K. Feng, S.R. Yu, Y.L. Li, L.Y. Gong, J Mater Process Tech. 208(1-3), (2008) 330-335. [19] J.I. Youn, Y.J. Kim, Jpn J Appl Phys. 48 (2009) 07GM14. [20] J.G. Jung, S.H. Lee, J.M. Lee, Y.H. Cho, S.H. Kim, W.H. Yoon, Mat Sci Eng a-Struct. 669 (2016) 187-195. [21] J.I. Youn, Y.K. Lee, Y.J. Kim, J.W. Park, Jpn J Appl Phys. 55 (2016) 07KE10). [22] Y. Xuan, L.Nastac, Ultrasonics. online 1 July 2017, https://doi.org/10.1016/j.ultras.2017.06.023. [23] Y. Xuan, L. Nastac, TMS 2017 146th Annual Meeting & Exhibition Supplemental Proceedings. San Deigo: 2017. Part of MMMS book series, pp. 297-303. [24] S. Naher, D. Brabazon, L. Looney, Composites Part A: Applied Science and Manufacturing. 38 (2007) 719-729. [25] D.M. Stefanescu, B.K. Dhindaw, S.A. Kacar, A. Moitra, Metallurgical Transactions A. 19 (1988) 2847-2855. [26] J.B. Ferguson, B.F. Schultz, P.K. Rohatgi, C.S. Kim, Light Metals 2014. John Wiley & Sons, Inc., 2014. pp. 1383-1388. [27] J.B. Ferguson, B. Schultz, P. Rohatgi, C.S. Kim, Met. Mater. Int. 20 (2014) 747-755. [28] A. Chernov, D. Temkin, A. Mel’nikova, Sov. Phys. Crystallogr. 21 (1976) 369-373. [29] A. Chernov, D. Temkin, A. Mel’nikova, Sov. Phys. Crystallogr. 22 (1977) 656-658. [30] S. Plimpton. J. Comput. Phys. 117 (1995) 1-19. [31] C.R. Dandekar, Y.C. Shin, Compos Part A: Appl Sci Manuf. 42 (4), (2011) 355-363. [32] B. Chalmers, Trans. AIME. 200 (1954), pp. 519. [33] J. Hoyt, M. Asta, A. Karma, Interface Science 10 (2002) 181-189. [34] Y. Mishin, D. Farkas, M.J. Mehl, D.A. Papaconstantopoulos, Physical Review B, 59 (1999) 3393-3407. [35] J. Winey, A. Kubota, Y. Gupta, Modelling Simul. Master. Sci. Eng. 17 (2009) 055004. [36] J. Tersoff, Phys. Rev. B. 39 (1989) 5566-5568. [37] J. Tersoff, Phys. Rev. Lett. 64 (1990) 1757-1760. [38] P.M. Morse, Phys. Rev. 34 (1929) 57-64. [39] H. Zhao, N. Chen, Y. Long, J. Phys. Condens. Matter. 21 (2009) 225002.
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