- Email: [email protected]
Synthesis and mechanical properties of bulk Al76Ni8Ti8Zr4Y4 alloy fabricated by consolidation of mechanically alloyed amorphous powders
Journal of Alloys and Compounds 632 (2015) 617–622
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Synthesis and mechanical properties of bulk Al76Ni8Ti8Zr4Y4 alloy fabricated by consolidation of mechanically alloyed amorphous powders Xinfu Wang, Kun Wang, Zhendong Li, Xingfu Wang, Dan Wang, Fusheng Han ⇑ Key Laboratory of Materials Physics, Institute of Solid State Physics, Chinese Academy of Sciences, Hefei 230031, China
a r t i c l e
i n f o
Article history: Received 26 December 2014 Received in revised form 26 January 2015 Accepted 29 January 2015 Available online 4 February 2015 Keywords: Amorphous alloy High-pressure sintering Vickers microhardness Nanohardness
a b s t r a c t Mechanically alloyed amorphous Al76Ni8Ti8Zr4Y4 (at.%) alloy powder was consolidated by high-pressure sintering process. The influence of the consolidation temperature on the structure and mechanical properties of the consolidated bulk alloys was examined by X-ray diffraction (XRD), Optical microscopy (OM), Scanning electron microscopy (SEM), Vickers Hardness Tester and Nano Indenter. Structural investigations of the bulk materials revealed that most of the amorphous structure was retained after consolidation at 623 K, however, compaction at 723 K and 823 K caused crystallization of the amorphous phase with the appearance of white regions. The results also indicate that application of high pressure affected the crystallization products of the present alloy. Micro mechanical analysis showed that the microhardness of the bulk composites increased from 945HV0.1 to 1177HV0.1 with the consolidation temperature increasing. The evolution trends of nanohardness and Young’s modulus agree well with the Vickers hardness testing results, indentation size effect (ISE) is observed in the multicomponent alloy systems. The yield strength, Meyer exponent and the indentation impressions of these bulk alloys have also been compared. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Bulk amorphous–nanocrystalline alloys have breathed new life both into materials science and application areas, due to their unique combination of properties, such as extraordinary strength, high fracture toughness and good corrosion resistance [1–3]. Among them, Al-based amorphous–nanocrystalline alloys have attracted more and more attention owe to their advantageous mechanical properties in particular a high specific strength to weight ratio and a high hardness [4,5], which are not reached by conventional Al-based alloys. Unfortunately, the Al-rich glasses are difficult to manufacture on account of the low glass-forming ability (GFA) with extremely high cooling rate (106 K/s) [4,6], the largest critical thickness obtained so far is about 1 mm for Al86Ni6Y4.5Co2La1.5 alloy [7] fabricated by rapid cooling method, which means that the application of Al-based metallic glasses as a structural material has been restricted. Given the challenges for the fabrication of Al-riched bulk metallic glasses (BMGs) from liquid, an alternative way of producing bulk amorphous or nanocrystalline materials is powder metallurgy route. Since it is a solid state alloying process, a high cooling rate is
⇑ Corresponding author. Tel.: +86 551 65591435; fax: +86 551 65591434. E-mail address: [email protected] (F. Han). http://dx.doi.org/10.1016/j.jallcom.2015.01.269 0925-8388/Ó 2015 Elsevier B.V. All rights reserved.
not necessary and the alloy composition can be more flexible. Based on the glassy powders produced by mechanical alloying (MA), the bulk glasses can be produced via powder metallurgy route, which makes it possible to produce bulk metallic glasses in more extensive ranges of size and shape [8]. However, it faces a great challenge to consolidate the amorphous powders into dense bulk material maintaining amorphous structure or creating nanocrystalline, because the amorphous phase will crystallize and grain coarsening will occur during conventional consolidation at elevated temperature. Nevertheless, high temperature is indispensable to obtain good interparticle bonding. To achieve this, considerable efforts have been made to obtain bulk Al-based amorphous–nanocrystalline alloys by applying of a high pressure as well as limiting of the high temperature exposure time during the consolidation. Based on such consideration, Al-based bulk amorphous–nanocrystalline alloys have been successfully fabricated by high-pressure sintering (HPS) method, spark plasma sintering (SPS) process, and so on [4,9–11]. Among the different consolidation techniques, HPS consolidation, which involves the simultaneous application of pressure and temperature, shows a large potential for achieving fast and highly dense samples. Meanwhile, HPS route can also break the oxide layers coating on the powder particles, and hence gives improved bonding between the particles.
618
X. Wang et al. / Journal of Alloys and Compounds 632 (2015) 617–622
In the present work, the synthesis of the bulk Al76Ni8Ti8Zr4Y4 amorphous and nanocrystalline alloy by HPS process was studied, the effects of sinter temperatures on the structure evolution and micro-mechanical properties of the bulk alloys were investigated using XRD, OM, SEM, Vickers Hardness Tester and Nano Indenter. The hardness, yield strength and Meyer exponent of these alloys have been compared, the related deformation behavior and mechanism were also discussed. 2. Experimental methods The mechanically alloyed Al76Ni8Ti8Zr4Y4 powders milled for 120 h were used for consolidation, the structure and thermal properties have been investigated elsewhere [12]. Disk-typed Al76Ni8Ti8Zr4Y4 bulk samples with a size of £10 ⁄ 2 mm were prepared by HPS method, HPS was conducted using a cubic hinge press machine, which is widely used in the ultra-hard materials field. The consolidation processes were conducted for 5 min under high vacuum conditions (102 Pa), the applied uniaxial pressure was 4 GPa at three different temperatures. The density of the sintered samples was evaluated using the Archimedes method. The specimens were mechanically polished to a mirror finish before structural analysis and micro-mechanical experiments. The structure of the sintered samples was studied by XRD analysis using Cu Ka radiation (X’ Pert Pro MPD). For comparison, the structure of as-milled powders isothermally annealed at the same temperature for the same time was also studied by XRD. Microhardness experiments were performed by MH-3 Vickers microhardness, the tests were carried out at five different loads of 25, 50, 100, 200 and 500 g and at a constant dwell time of 15 s. At least six points on each disk were selected for Vickers hardness testing. Nanohardness of the bulk samples was evaluated using Nano Indenter XP system (MTS NANO Indenter) with a Berkovich diamond tip at room temperature. The experiments were conducted in the displacement-control mode with a maximum indentation depth of 1700 nm. The strain rate was controlled at 0.05 s1. The hardness and elastic modulus were averaged between the indentation depths of 200– 1500 nm, six indentations were performed to verify the accuracy and scatter of the indentation data. The indented impression was further observed under an optical microscope (OM Zeiss Imager. A1 m) and scanning electron microscope (SEM FEI Sirion 200) equipped with an X-ray energy dispersive spectroscopy (EDS) system for structural characterization and for determination of chemical composition.
3. Results and discussion According to the DSC results of the Al76Ni8Ti8Zr4Y4 powders milled for 120 h, 623 K, 723 K and 823 K were selected to consolidate the powders to a disk shape with a diameter of 10 mm and a thickness of 2 mm. Detailed XRD patterns of the samples consolidated under different temperatures are summarized in Fig. 1a. It can be seen that the bulk material reveals two weak crystallization peaks of fcc Al and an ambiguous amorphous background after sintering at 623 K, which suggesting that the most of amorphous state was retained. With the sintering temperature increased to 723 K, more crystalline peaks appear and the intensity of the peaks increases, meaning that the crystalline fraction in the BMG increases as the sinter temperature increased. The average grain size is about 90 nm and the crystallization products are fcc Al, Al3Zr, Al3Ti, AlY, Al3Ni and other crystalline phases signed in Fig. 1a. There are some crystallized phases cannot be identified at present, because of the complexity of the multicomponent alloy. As the consolidated temperature increased to 823 K, the crystalline phase and grain size have not noticeable changes compared to the one consolidated at 723 K. The crystallinity of the samples consolidated at 4 GPa calculated from the software of X’ Pert HighScore Plus was 3.3%, 24.0%, 29.5% at the sinter temperature of 623 K, 723 K, and 823 K, respectively, indicating the BMGs crystallize fraction increased with the enhanced sinter temperature. For comparison, the as-milled powders were isothermally annealed in a tube resistance electric furnace filled with flowing high purity argon under ambient pressure, the annealing temperature and time are the same as those of the HPS bulk samples, the XRD patterns of the isothermally annealed powders are shown in Fig. 1b. It is seen that, under ambient pressure, the amorphous
Fig. 1. XRD patterns of (a) the Al76Ni8Ti8Zr4Y4 bulk samples consolidated at different temperatures; (b) the as-milled Al76Ni8Ti8Zr4Y4 powders annealed at the same temperature for the same time with the bulk samples.
phase is retained when the annealing temperature is below 723 K. Further increase in annealing temperature results in the coexistence of amorphous and crystalline phases (fcc Al, Al3Zr, AlTi and AlTi3) with the extremely low intensity. It is clear that the crystallization products are not consistent with the consolidated bulk samples, which suggesting that the amorphous state of the fabricated samples are sensitive to the pressure applied. In the present study, the major difference between the annealed and HPS specimens is the employment of high pressure. Thus, high pressure involved during consolidation can cut down on the existence of amorphous phase within the Al76Ni8Ti8Zr4Y4 alloy. The fcc A1 precipitation temperature decreases dramatically with an increase of the applied pressure, from about 792 K to 623 K, which implies that the crystallization kinetic of Al–Ni–Ti–Zr–Y amorphous powders is altered due to the applied pressure. Pressure effect on crystallization kinetics of amorphous alloys has been investigated in several systems [13–18]. It was found that the crystallization kinetics is associated with the atomic diffusion process and the volume change effect during the initial stage of nucleation of crystals in the amorphous phase. In other words, the applied pressure effect may be twofold. On the one hand, amorphous state is thermodynamically unstable and has a larger excessive volume than the counterpart of its crystalline, crystallization is a process of increasing density, so pressure may enhance the crystallization process from a thermodynamic point of view, some researchers [14,16,18] have confirmed the conclusion. On the other hand, an
X. Wang et al. / Journal of Alloys and Compounds 632 (2015) 617–622
applied pressure may suppress the atom diffusion, therefore, the crystallization process is retarded and then stabilized the amorphous phase. In this research, the observed decreasing crystallization temperature implied that the first effect (negative volume change) may play a dominant role. In addition, the composition of amorphous alloy will result in different crystallization modes and products, thus the pressure effect on crystallization will be varied. So, the intrinsic mechanism for the pressure effect on intermetallic formation from the remaining amorphous phase needs further studies in detail. To examine the microstructure and to provide information about the quality of consolidation, all the compacted samples were investigated using OM. Fig. 2a–c shows the polished cross-sectional plane of Al76Ni8Ti8Zr4Y4 bulk samples consolidated at 623 K, 723 K, 823 K, respectively. It can be seen that there are white and black regions except for the gray matrices, the amount and size of the white regions increase with the consolidated temperature rising, their sizes are ranged from 1 to 20 lm, the similar phenomenon is also arising in Ref. [19]. The chemical composition of these white regions and of the matrices was checked using EDS system in SEM, which proved that the white regions contain only Al, Ni and Ti elements, while Al, Ni, Ti, Zr and Y are present in the matrices. The black regions were conformed as pores by SEM (inset in Fig. 2a). There are almost no pores present as the consolidated temperature above 723 K, which evidences good quality of consolidation. The densities are 3.61 g/cm3, 3.80 g/cm3, 3.94 g/cm3 for the samples at the consolidation temperature of 623 K, 723 K and 823 K, respectively. This means that the density of the amorphous Al76Ni8Ti8Zr4Y4 alloy increases with crystallization. The microhardness of the bulk samples was investigated by Vickers microhardness measurements on the polished surfaces of the bulk samples, the white regions were also tested for comparison. The measured microhardness, HV, as a function of the applied test load, P, for each sample is listed in Table 1. As can be seen, the highest hardness is observed in the sample consolidated at 823 K followed by 723 K consolidated specimen, while, the almost amorphous one shows the lowest hardness. Similar results have also been reported [20,21], which was attributed to the presence of nanocrystals suppressing shear band multiplication and therefore increase in hardness. Moreover, annihilation of free volume takes place leading to restricted shear band movement and thus increases in hardness. It is clear that the microhardness in the matrix is a little higher than the white regions, which could be caused by the difference of the chemical composition and chemical bonding forces between them. The standard deviation of the measured microhardness values is small (±5%), which indicates uniformity of this property in each bulk sample. It can be also found that the microhardness for each sample decreases significantly with applied load. The behavior that the apparent hardness decreases with increasing the applied load is similar to that observed in brittle ceramics, glasses and the conventional crystalline materials [22]. Such phenomena usually referred to the indentation size effect (ISE). The ISE was traditionally described through the application of the Meyer’s law, in which the applied load, P (g), and indentation diagonal length, d (lm), is correlated as:
P ¼ kd
n
ð1Þ
where k is a material constant pertaining to the resistance against penetration by the indenter, and n is Meyer exponent also known as a material constant related to strain hardening of the material system, which can be derived from the curve fitting of the experimental results. Fig. 3 shows log P versus log d curves of the bulk samples consolidated at different temperatures. Table 2 presents
619
Fig. 2. OM images of the bulk Al76Ni8Ti8Zr4Y4 alloys consolidated at (a) 623 K; (b) 723 K; (c) 823 K. The inset in Fig. 2a shows the SEM image of black region of 623 K sintered sample.
the n and K values at different temperatures. It can be seen that both n and K are found to be maximum at 823 K sinter. The values of n are all less than 2. For materials exhibiting ISE, the Meyer exponent n was experimentally observed to be between 1 and 2. When n = 2, the hardness is independent of the applied test load. For the normal ISE behavior, the exponent 1 < n < 2, the measured hardness apparently decreases with increasing applied test load. As a result, a relatively high microhardness is obtained when the applied load is small.
620
X. Wang et al. / Journal of Alloys and Compounds 632 (2015) 617–622
Table 1 The values of microhardness (HV) of consolidated Al76Ni8Ti8Zr4Y4 bulk alloys. Consolidated temperature (K)
Region
623 723
Matrix Matrix White region Matrix White region
823
Microhardness (HV) 25 g
50 g
100 g
200 g
500 g
1092 1367 – 1397 –
982 1197 1011 1275 1080
945 1117 – 1177 –
804 1087 – 1094 –
753 893 – 999 –
Fig. 3. Log P versus log d plots for the as synthesized Al76Ni8Ti8Zr4Y4 alloys.
The Vickers hardness H (in GPa) was calculated by the following formula [23]
H ¼ 1:854 9:8
P
ð2Þ
2
d
The 0.2% offset yield strength can be computed by the microhardness values according to the following relationship [24];
r0 ¼ ðH=3Þð0:1Þn2
ð3Þ
where r0 is 0.2% offset yield strength and n is the Meyer exponent. The values of r0 and relevant parameters that could be inferred based on HV are also given in Table 2. Apparently the microhardness of the present Al76Ni8Ti8Zr4Y4 alloy is high compared to other Albased bulk amorphous [22,25,26]. The difference of the consolidating temperature and pressure may be responsible for the observed changes in the microhardness. Another important reason is that the chemical and physical properties of the various elements present in the alloy, atomic configuration and chemical bonding forces maybe account for the high microhardness of the metallic glass. OM was conducted around the indented areas with a load of 50 g, 100 g and 200 g, respectively. All the indents are of a regular pyramidal shape, as shown in Fig. 4. It should be noted that indentation-induced cracks and circumferential crushing are observed around the indented area in the bulk specimen consolidated at 623 K (Fig. 4a). As the sinter temperature increased, no circumferential crushing is observed around each indent of the specimens
(Fig. 4b and c) and the cracks become smaller, which suggested that the latter two bulk alloys have better fracture toughness than the former one. The crack paths surrounding the indents of all these samples emanate from the indent corners, indicating that the propagation of cracks occurred by radiation. For comparison, the indentation impressions of the white regions consolidated at 723 K and 823 K are shown in Fig. 5. Interestingly, the formation of cracks around the indentation periphery has been observed, which are not shown in the matrix. The possible reason was that brittle nano-crystalline intermetallic like NiAl and TiAl was formed during the consolidation process, resulting in the formation of cracks. Fig. 6 compares the representative load–displacement (P–h) responses of the bulk samples consolidated at different temperatures. It shows that the sample sintered at 823 K exhibits maximum load during loading, while the sample consolidated at 623 K shows the minimum load. Compared to higher temperature sintered ones, larger and distinct displacement bursts (pop-in events) are observed in the P–h curves of the lowest temperature sintered sample as shown in black arrows in Fig. 6. The pop-in events associated with the formation and propagation of shear bands, which suggests that large numbers of shear bands are activated during indentation in the amorphous region. Although a few serrations are observed in P–h curve of the bulk sample consolidated at 623 K, the indentation morphology does not reveal any shear band formation, as indicated in the SEM micrographs of the indentation imprints (insert in Fig. 6). This can be attributed to the fact that all the ‘pop-in’ events cannot be accounted for shear bands, since pop-ins can lead to nucleation and propagation of shear bands inside of the material as well [27]. The absence of pop-in events indicates that shear band activity during plastic deformation of this alloy is not very pronounced. Fig. 7 depicts the variation of nanohardness (Hn) and Young’s modulus (E) with the indentation depth for the different consolidation temperatures. As is seen, these evolution trends of Hn and E agree well with the Vickers hardness testing results, i.e., the Hn and E of the bulk sample increase with the consolidation temperature due to the free volume annihilation, similar to the Vickers hardness evolution during consolidation, as shown in Table 1. It can mainly be attributed to the partial nanocrystallization, since partial nanocrystallization increases interatomic distance and atomic bonding energy. It should be also noted that the nanohardness decreased with the indentation depth, which is similar to that reported for other metallic glasses during nanoindentation test. Similarly, Young’s modulus also decreased with the applied load. Such mechanical
Table 2 The values of microhardness (H), Meyer’s exponent (n), material constant (K) and yield strength (r0) of consolidated Al76Ni8Ti8Zr4Y4 bulk alloys. Consolidated temperature (K)
n
K
H at 50 g load (GPa)
r0 at 50 g load (GPa)
623 723 823
1.773 1.761 1.794
187.5 214.3 258.8
9.61 11.71 12.48
5.48 6.55 6.90
X. Wang et al. / Journal of Alloys and Compounds 632 (2015) 617–622
621
Fig. 5. Different regions indentation morphologies under 50 g load consolidated at (a) 723 K; (b) 823 K.
Fig. 4. OM indentation morphologies under various loads (50 g, 100 g, 200 g) at different consolidation temperatures (a) 623 K; (b) 723 K; (c) 823 K.
softening has been reported for other metallic glasses [28,29] and attributed to the net creation of free volume during plastic deformation. The average values of nanohardness and young’s modulus are listed in Table 3. It is evident that nanohardness values are little higher in comparison with the microindentation results, while the maximum load (50 g 500 mN) and the size of indention is almost the same with the Vickers hardness test, therefore, the ISE effect should be unconspicuous. This may be caused by experimental errors including the measurement of indentation size, the applied load and the geometry of indenter [30]. In addition to Hn and E, Table 3 shows the values of Hn/E ratios, which are indicative of the wear resistance of the metallic glass. A
Fig. 6. Representative nanoindentation load–displacement curves of the bulk sample consolidated at three different temperatures, where the black arrow indicates a displacement burst. The inset shows the SEM nanoindentation impression of 623 K sintered sample.
high Hn/E ratio is often a reliable indicator of good wear resistance in a coating or layers, therefore, the almost amorphous bulk sample seems to have slightly higher wear resistance than the bulk alloy because of its lower E.
622
X. Wang et al. / Journal of Alloys and Compounds 632 (2015) 617–622
tested by Vickers hardness technique. The SBs corresponding to pop-ins in P–h curve were not observed around the indented area in the bulk specimen consolidated at 623 K, due to the unconspicuous shear band activity during plastic deformation of the present alloy. The microhardness in the matrix is a little higher than the white regions, because brittle nano-crystalline intermetallic like NiAl and TiAl was formed in the white regions, which caused the cracks formed around the indentation periphery. The indentation-induced microcracking become smaller with increasing consolidated temperature, which suggesting that the bulk alloys consolidated at elevated temperature have better fracture toughness than the low temperature sintering one. The crack paths around the indents of the sample emanate from the indent corners, indicating that the propagation of cracks occurred by radiation. Acknowledgments This work was jointly supported by the National Science Foundation of China (51301168 and 513711), the National Basic Research Program of China (2011CB610300) and the Instrument Developing Project of Chinese Academy of Sciences (YG2012065) References
Fig. 7. Nanohardness (a) and Young’s modulus (b) of the Al76Ni8Ti8Zr4Y4 alloy with the indentation depth for the different consolidation temperatures.
Table 3 The nanohardness, Young’s modulus and Hn/E for the Al76Ni8Ti8Zr4Y4 bulk alloys. Consolidated temperature (K)
Nanohardness, Hn (GPa)
Young’s modulus, E (GPa)
Hn/E ratio
623 723 823
11.16 12.74 13.27
173.3 213.9 223.1
0.064 0.060 0.059
4. Conclusions Bulk amorphous and nanocrystalline Al76Ni8Ti8Zr4Y4 alloy has been successfully prepared by mechanical alloying and high pressure sintering process. Structural examinations of the bulk materials reveal that most of the amorphous structure was retained after consolidation at 623 K, the crystallinity increased with the increasing temperature, the results also indicate that application of high pressure enhanced the crystallization process due to the negative volume change and affected the crystalline products of this process. Micro mechanical analysis showed that the microhardness of the bulk composites increased from 945HV0.1 to 1177HV0.1 with the consolidation temperature increasing. The evolution trends of the nanohardness and young’s modulus are similar with the values
[1] M.M. Trexler, N.N. Thadhani, Prog. Mater. Sci. 55 (2010) 759–839. [2] H.W. Yang, M.J. Tan, R.V. Ramanujan, Scripta Mater. 66 (2012) 382–385. [3] B.J. Yang, J.H. Yao, Y.S. Chao, J.Q. Wang, E. Ma, Philos. Mag. 90 (2010) 3215– 3231. [4] J.Q. Wang, P. Dong, W.L. Hou, X.C. Chang, M.X. Quan, J. Alloys Comp. 554 (2013) 419–425. [5] X.P. Li, M. Yan, H. Imai, K. Kondoh, J.Q. Wang, G.B. Schaffer, M. Qian, Mater. Sci. Eng. A 568 (2013) 155–159. [6] A.H. Cai, D.W. Ding, X. Xiong, Y. Liu, W.K. An, G.J. Zhou, Y. Luo, T.L. Li, X.S. Li, Mater. Des. 63 (2014) 233–237. [7] B.J. Yang, J.H. Yao, J. Zhang, H.W. Yang, J.Q. Wang, E. Ma, Scripta Mater. 61 (2009) 423–426. [8] M.H. Lee, D.H. Bae, D.H. Kim, D.J. Sordelet, J. Mater. Res. 18 (2003) 2101–2108. [9] T.T. Sasaki, K. Hono, J. Vierke, M. Wollgarten, J. Banhart, Mater. Sci. Eng. A 490 (2008) 343–350. [10] S. Mula, K. Mondal, S. Ghosh, S.K. Pabi, Mater. Sci. Eng. A 527 (2010) 3757– 3763. [11] M. Krasnowski, T. Kulik, J. Alloys Comp. 495 (2010) 382–385. [12] X. Wang, D. Wang, B. Zhu, Y. Li, F. Han, J. Non-Cryst. Solids 385 (2014) 111– 116. [13] W.H. Wang, D.W. He, D.Q. Zhao, Y.S. Yao, M. He, Appl. Phys. Lett. 75 (1999) 2770–2772. [14] F. Ye, K. Lu, Acta Mater. 47 (1999) 2449–2454. [15] Y.X. Zhuang, J.Z. Jiang, T.J. Zhou, H. Rasmussen, L. Gerward, M. Mezouar, W. Crichton, A. Inoue, Appl. Phys. Lett. 77 (2000) 4133–4135. [16] X.J. Gu, J.Q. Wang, F. Ye, K. Lu, J. Non-Cryst. Solids 296 (2001) 74–80. [17] P.F. Xing, Y.X. Zhuang, W.H. Wang, L. Gerward, J.Z. Jiang, J. Appl. Phys. 91 (2002) 4956–4960. [18] Z.X. Wang, F.Y. Li, M.X. Pan, D.Q. Zhao, W.H. Wang, J. Alloys Comp. 388 (2005) 262–265. [19] M. Krasnowski, A. Antolak-Dudka, T. Kulik, Intermetallics 19 (2011) 1243– 1249. [20] S. Vincent, J. Basu, B.S. Murty, J. Bhatt, Mater. Sci. Eng. A 550 (2012) 160–166. [21] Y.H. Kim, K. Hiraga, A. Inoue, T. Masumoto, H.H. Jo, Mater. Trans. JIM 35 (1994) 293–302. [22] X. Wei, F.S. Han, X.F. Wang, C.E. Wen, J. Alloys Comp. 501 (2010) 164–167. [23] N.K. Mukhopadhyay, G.C. Weatherly, J.D. Embury, Mater. Sci. Eng. A 315 (2001) 202–210. [24] D. Singh, D. Singh, R.K. Mandal, O.N. Srivastava, R.S. Tiwari, J. Alloys Comp. 590 (2014) 15–20. [25] B. Zhu, X.F. Wang, X.Y. Li, D. Wang, Y. Qin, F.S. Han, J. Alloys Comp. 586 (2014) 645–649. [26] I. Borner, J. Eckert, Scripta Mater. 45 (2001) 237–244. [27] S. Vincent, B.S. Murty, M.J. Kramer, J. Bhatt, Mater. Des. 65 (2015) 98–103. [28] N. Van Steenberge, J. Sort, A. Concustell, J. Das, S. Scudino, S. Surinach, J. Eckert, M.D. Baro, Scripta Mater. 56 (2007) 605–608. [29] J.I. Jang, B.G. Yoo, Y.J. Kim, J.H. Oh, I.C. Choi, H.B. Bei, Scripta Mater. 64 (2011) 753–756. [30] K. Sangwal, B. Surowska, P. Blaziak, Mater. Chem. Phys. 77 (2003) 511–520.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Synthesis and mechanical properties of bulk Al76Ni8Ti8Zr4Y4 alloy fabricated by consolidation of mechanically alloyed amorphous powders Xinfu Wang, Kun Wang, Zhendong Li, Xingfu Wang, Dan Wang, Fusheng Han ⇑ Key Laboratory of Materials Physics, Institute of Solid State Physics, Chinese Academy of Sciences, Hefei 230031, China
a r t i c l e
i n f o
Article history: Received 26 December 2014 Received in revised form 26 January 2015 Accepted 29 January 2015 Available online 4 February 2015 Keywords: Amorphous alloy High-pressure sintering Vickers microhardness Nanohardness
a b s t r a c t Mechanically alloyed amorphous Al76Ni8Ti8Zr4Y4 (at.%) alloy powder was consolidated by high-pressure sintering process. The influence of the consolidation temperature on the structure and mechanical properties of the consolidated bulk alloys was examined by X-ray diffraction (XRD), Optical microscopy (OM), Scanning electron microscopy (SEM), Vickers Hardness Tester and Nano Indenter. Structural investigations of the bulk materials revealed that most of the amorphous structure was retained after consolidation at 623 K, however, compaction at 723 K and 823 K caused crystallization of the amorphous phase with the appearance of white regions. The results also indicate that application of high pressure affected the crystallization products of the present alloy. Micro mechanical analysis showed that the microhardness of the bulk composites increased from 945HV0.1 to 1177HV0.1 with the consolidation temperature increasing. The evolution trends of nanohardness and Young’s modulus agree well with the Vickers hardness testing results, indentation size effect (ISE) is observed in the multicomponent alloy systems. The yield strength, Meyer exponent and the indentation impressions of these bulk alloys have also been compared. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Bulk amorphous–nanocrystalline alloys have breathed new life both into materials science and application areas, due to their unique combination of properties, such as extraordinary strength, high fracture toughness and good corrosion resistance [1–3]. Among them, Al-based amorphous–nanocrystalline alloys have attracted more and more attention owe to their advantageous mechanical properties in particular a high specific strength to weight ratio and a high hardness [4,5], which are not reached by conventional Al-based alloys. Unfortunately, the Al-rich glasses are difficult to manufacture on account of the low glass-forming ability (GFA) with extremely high cooling rate (106 K/s) [4,6], the largest critical thickness obtained so far is about 1 mm for Al86Ni6Y4.5Co2La1.5 alloy [7] fabricated by rapid cooling method, which means that the application of Al-based metallic glasses as a structural material has been restricted. Given the challenges for the fabrication of Al-riched bulk metallic glasses (BMGs) from liquid, an alternative way of producing bulk amorphous or nanocrystalline materials is powder metallurgy route. Since it is a solid state alloying process, a high cooling rate is
⇑ Corresponding author. Tel.: +86 551 65591435; fax: +86 551 65591434. E-mail address: [email protected] (F. Han). http://dx.doi.org/10.1016/j.jallcom.2015.01.269 0925-8388/Ó 2015 Elsevier B.V. All rights reserved.
not necessary and the alloy composition can be more flexible. Based on the glassy powders produced by mechanical alloying (MA), the bulk glasses can be produced via powder metallurgy route, which makes it possible to produce bulk metallic glasses in more extensive ranges of size and shape [8]. However, it faces a great challenge to consolidate the amorphous powders into dense bulk material maintaining amorphous structure or creating nanocrystalline, because the amorphous phase will crystallize and grain coarsening will occur during conventional consolidation at elevated temperature. Nevertheless, high temperature is indispensable to obtain good interparticle bonding. To achieve this, considerable efforts have been made to obtain bulk Al-based amorphous–nanocrystalline alloys by applying of a high pressure as well as limiting of the high temperature exposure time during the consolidation. Based on such consideration, Al-based bulk amorphous–nanocrystalline alloys have been successfully fabricated by high-pressure sintering (HPS) method, spark plasma sintering (SPS) process, and so on [4,9–11]. Among the different consolidation techniques, HPS consolidation, which involves the simultaneous application of pressure and temperature, shows a large potential for achieving fast and highly dense samples. Meanwhile, HPS route can also break the oxide layers coating on the powder particles, and hence gives improved bonding between the particles.
618
X. Wang et al. / Journal of Alloys and Compounds 632 (2015) 617–622
In the present work, the synthesis of the bulk Al76Ni8Ti8Zr4Y4 amorphous and nanocrystalline alloy by HPS process was studied, the effects of sinter temperatures on the structure evolution and micro-mechanical properties of the bulk alloys were investigated using XRD, OM, SEM, Vickers Hardness Tester and Nano Indenter. The hardness, yield strength and Meyer exponent of these alloys have been compared, the related deformation behavior and mechanism were also discussed. 2. Experimental methods The mechanically alloyed Al76Ni8Ti8Zr4Y4 powders milled for 120 h were used for consolidation, the structure and thermal properties have been investigated elsewhere [12]. Disk-typed Al76Ni8Ti8Zr4Y4 bulk samples with a size of £10 ⁄ 2 mm were prepared by HPS method, HPS was conducted using a cubic hinge press machine, which is widely used in the ultra-hard materials field. The consolidation processes were conducted for 5 min under high vacuum conditions (102 Pa), the applied uniaxial pressure was 4 GPa at three different temperatures. The density of the sintered samples was evaluated using the Archimedes method. The specimens were mechanically polished to a mirror finish before structural analysis and micro-mechanical experiments. The structure of the sintered samples was studied by XRD analysis using Cu Ka radiation (X’ Pert Pro MPD). For comparison, the structure of as-milled powders isothermally annealed at the same temperature for the same time was also studied by XRD. Microhardness experiments were performed by MH-3 Vickers microhardness, the tests were carried out at five different loads of 25, 50, 100, 200 and 500 g and at a constant dwell time of 15 s. At least six points on each disk were selected for Vickers hardness testing. Nanohardness of the bulk samples was evaluated using Nano Indenter XP system (MTS NANO Indenter) with a Berkovich diamond tip at room temperature. The experiments were conducted in the displacement-control mode with a maximum indentation depth of 1700 nm. The strain rate was controlled at 0.05 s1. The hardness and elastic modulus were averaged between the indentation depths of 200– 1500 nm, six indentations were performed to verify the accuracy and scatter of the indentation data. The indented impression was further observed under an optical microscope (OM Zeiss Imager. A1 m) and scanning electron microscope (SEM FEI Sirion 200) equipped with an X-ray energy dispersive spectroscopy (EDS) system for structural characterization and for determination of chemical composition.
3. Results and discussion According to the DSC results of the Al76Ni8Ti8Zr4Y4 powders milled for 120 h, 623 K, 723 K and 823 K were selected to consolidate the powders to a disk shape with a diameter of 10 mm and a thickness of 2 mm. Detailed XRD patterns of the samples consolidated under different temperatures are summarized in Fig. 1a. It can be seen that the bulk material reveals two weak crystallization peaks of fcc Al and an ambiguous amorphous background after sintering at 623 K, which suggesting that the most of amorphous state was retained. With the sintering temperature increased to 723 K, more crystalline peaks appear and the intensity of the peaks increases, meaning that the crystalline fraction in the BMG increases as the sinter temperature increased. The average grain size is about 90 nm and the crystallization products are fcc Al, Al3Zr, Al3Ti, AlY, Al3Ni and other crystalline phases signed in Fig. 1a. There are some crystallized phases cannot be identified at present, because of the complexity of the multicomponent alloy. As the consolidated temperature increased to 823 K, the crystalline phase and grain size have not noticeable changes compared to the one consolidated at 723 K. The crystallinity of the samples consolidated at 4 GPa calculated from the software of X’ Pert HighScore Plus was 3.3%, 24.0%, 29.5% at the sinter temperature of 623 K, 723 K, and 823 K, respectively, indicating the BMGs crystallize fraction increased with the enhanced sinter temperature. For comparison, the as-milled powders were isothermally annealed in a tube resistance electric furnace filled with flowing high purity argon under ambient pressure, the annealing temperature and time are the same as those of the HPS bulk samples, the XRD patterns of the isothermally annealed powders are shown in Fig. 1b. It is seen that, under ambient pressure, the amorphous
Fig. 1. XRD patterns of (a) the Al76Ni8Ti8Zr4Y4 bulk samples consolidated at different temperatures; (b) the as-milled Al76Ni8Ti8Zr4Y4 powders annealed at the same temperature for the same time with the bulk samples.
phase is retained when the annealing temperature is below 723 K. Further increase in annealing temperature results in the coexistence of amorphous and crystalline phases (fcc Al, Al3Zr, AlTi and AlTi3) with the extremely low intensity. It is clear that the crystallization products are not consistent with the consolidated bulk samples, which suggesting that the amorphous state of the fabricated samples are sensitive to the pressure applied. In the present study, the major difference between the annealed and HPS specimens is the employment of high pressure. Thus, high pressure involved during consolidation can cut down on the existence of amorphous phase within the Al76Ni8Ti8Zr4Y4 alloy. The fcc A1 precipitation temperature decreases dramatically with an increase of the applied pressure, from about 792 K to 623 K, which implies that the crystallization kinetic of Al–Ni–Ti–Zr–Y amorphous powders is altered due to the applied pressure. Pressure effect on crystallization kinetics of amorphous alloys has been investigated in several systems [13–18]. It was found that the crystallization kinetics is associated with the atomic diffusion process and the volume change effect during the initial stage of nucleation of crystals in the amorphous phase. In other words, the applied pressure effect may be twofold. On the one hand, amorphous state is thermodynamically unstable and has a larger excessive volume than the counterpart of its crystalline, crystallization is a process of increasing density, so pressure may enhance the crystallization process from a thermodynamic point of view, some researchers [14,16,18] have confirmed the conclusion. On the other hand, an
X. Wang et al. / Journal of Alloys and Compounds 632 (2015) 617–622
applied pressure may suppress the atom diffusion, therefore, the crystallization process is retarded and then stabilized the amorphous phase. In this research, the observed decreasing crystallization temperature implied that the first effect (negative volume change) may play a dominant role. In addition, the composition of amorphous alloy will result in different crystallization modes and products, thus the pressure effect on crystallization will be varied. So, the intrinsic mechanism for the pressure effect on intermetallic formation from the remaining amorphous phase needs further studies in detail. To examine the microstructure and to provide information about the quality of consolidation, all the compacted samples were investigated using OM. Fig. 2a–c shows the polished cross-sectional plane of Al76Ni8Ti8Zr4Y4 bulk samples consolidated at 623 K, 723 K, 823 K, respectively. It can be seen that there are white and black regions except for the gray matrices, the amount and size of the white regions increase with the consolidated temperature rising, their sizes are ranged from 1 to 20 lm, the similar phenomenon is also arising in Ref. [19]. The chemical composition of these white regions and of the matrices was checked using EDS system in SEM, which proved that the white regions contain only Al, Ni and Ti elements, while Al, Ni, Ti, Zr and Y are present in the matrices. The black regions were conformed as pores by SEM (inset in Fig. 2a). There are almost no pores present as the consolidated temperature above 723 K, which evidences good quality of consolidation. The densities are 3.61 g/cm3, 3.80 g/cm3, 3.94 g/cm3 for the samples at the consolidation temperature of 623 K, 723 K and 823 K, respectively. This means that the density of the amorphous Al76Ni8Ti8Zr4Y4 alloy increases with crystallization. The microhardness of the bulk samples was investigated by Vickers microhardness measurements on the polished surfaces of the bulk samples, the white regions were also tested for comparison. The measured microhardness, HV, as a function of the applied test load, P, for each sample is listed in Table 1. As can be seen, the highest hardness is observed in the sample consolidated at 823 K followed by 723 K consolidated specimen, while, the almost amorphous one shows the lowest hardness. Similar results have also been reported [20,21], which was attributed to the presence of nanocrystals suppressing shear band multiplication and therefore increase in hardness. Moreover, annihilation of free volume takes place leading to restricted shear band movement and thus increases in hardness. It is clear that the microhardness in the matrix is a little higher than the white regions, which could be caused by the difference of the chemical composition and chemical bonding forces between them. The standard deviation of the measured microhardness values is small (±5%), which indicates uniformity of this property in each bulk sample. It can be also found that the microhardness for each sample decreases significantly with applied load. The behavior that the apparent hardness decreases with increasing the applied load is similar to that observed in brittle ceramics, glasses and the conventional crystalline materials [22]. Such phenomena usually referred to the indentation size effect (ISE). The ISE was traditionally described through the application of the Meyer’s law, in which the applied load, P (g), and indentation diagonal length, d (lm), is correlated as:
P ¼ kd
n
ð1Þ
where k is a material constant pertaining to the resistance against penetration by the indenter, and n is Meyer exponent also known as a material constant related to strain hardening of the material system, which can be derived from the curve fitting of the experimental results. Fig. 3 shows log P versus log d curves of the bulk samples consolidated at different temperatures. Table 2 presents
619
Fig. 2. OM images of the bulk Al76Ni8Ti8Zr4Y4 alloys consolidated at (a) 623 K; (b) 723 K; (c) 823 K. The inset in Fig. 2a shows the SEM image of black region of 623 K sintered sample.
the n and K values at different temperatures. It can be seen that both n and K are found to be maximum at 823 K sinter. The values of n are all less than 2. For materials exhibiting ISE, the Meyer exponent n was experimentally observed to be between 1 and 2. When n = 2, the hardness is independent of the applied test load. For the normal ISE behavior, the exponent 1 < n < 2, the measured hardness apparently decreases with increasing applied test load. As a result, a relatively high microhardness is obtained when the applied load is small.
620
X. Wang et al. / Journal of Alloys and Compounds 632 (2015) 617–622
Table 1 The values of microhardness (HV) of consolidated Al76Ni8Ti8Zr4Y4 bulk alloys. Consolidated temperature (K)
Region
623 723
Matrix Matrix White region Matrix White region
823
Microhardness (HV) 25 g
50 g
100 g
200 g
500 g
1092 1367 – 1397 –
982 1197 1011 1275 1080
945 1117 – 1177 –
804 1087 – 1094 –
753 893 – 999 –
Fig. 3. Log P versus log d plots for the as synthesized Al76Ni8Ti8Zr4Y4 alloys.
The Vickers hardness H (in GPa) was calculated by the following formula [23]
H ¼ 1:854 9:8
P
ð2Þ
2
d
The 0.2% offset yield strength can be computed by the microhardness values according to the following relationship [24];
r0 ¼ ðH=3Þð0:1Þn2
ð3Þ
where r0 is 0.2% offset yield strength and n is the Meyer exponent. The values of r0 and relevant parameters that could be inferred based on HV are also given in Table 2. Apparently the microhardness of the present Al76Ni8Ti8Zr4Y4 alloy is high compared to other Albased bulk amorphous [22,25,26]. The difference of the consolidating temperature and pressure may be responsible for the observed changes in the microhardness. Another important reason is that the chemical and physical properties of the various elements present in the alloy, atomic configuration and chemical bonding forces maybe account for the high microhardness of the metallic glass. OM was conducted around the indented areas with a load of 50 g, 100 g and 200 g, respectively. All the indents are of a regular pyramidal shape, as shown in Fig. 4. It should be noted that indentation-induced cracks and circumferential crushing are observed around the indented area in the bulk specimen consolidated at 623 K (Fig. 4a). As the sinter temperature increased, no circumferential crushing is observed around each indent of the specimens
(Fig. 4b and c) and the cracks become smaller, which suggested that the latter two bulk alloys have better fracture toughness than the former one. The crack paths surrounding the indents of all these samples emanate from the indent corners, indicating that the propagation of cracks occurred by radiation. For comparison, the indentation impressions of the white regions consolidated at 723 K and 823 K are shown in Fig. 5. Interestingly, the formation of cracks around the indentation periphery has been observed, which are not shown in the matrix. The possible reason was that brittle nano-crystalline intermetallic like NiAl and TiAl was formed during the consolidation process, resulting in the formation of cracks. Fig. 6 compares the representative load–displacement (P–h) responses of the bulk samples consolidated at different temperatures. It shows that the sample sintered at 823 K exhibits maximum load during loading, while the sample consolidated at 623 K shows the minimum load. Compared to higher temperature sintered ones, larger and distinct displacement bursts (pop-in events) are observed in the P–h curves of the lowest temperature sintered sample as shown in black arrows in Fig. 6. The pop-in events associated with the formation and propagation of shear bands, which suggests that large numbers of shear bands are activated during indentation in the amorphous region. Although a few serrations are observed in P–h curve of the bulk sample consolidated at 623 K, the indentation morphology does not reveal any shear band formation, as indicated in the SEM micrographs of the indentation imprints (insert in Fig. 6). This can be attributed to the fact that all the ‘pop-in’ events cannot be accounted for shear bands, since pop-ins can lead to nucleation and propagation of shear bands inside of the material as well [27]. The absence of pop-in events indicates that shear band activity during plastic deformation of this alloy is not very pronounced. Fig. 7 depicts the variation of nanohardness (Hn) and Young’s modulus (E) with the indentation depth for the different consolidation temperatures. As is seen, these evolution trends of Hn and E agree well with the Vickers hardness testing results, i.e., the Hn and E of the bulk sample increase with the consolidation temperature due to the free volume annihilation, similar to the Vickers hardness evolution during consolidation, as shown in Table 1. It can mainly be attributed to the partial nanocrystallization, since partial nanocrystallization increases interatomic distance and atomic bonding energy. It should be also noted that the nanohardness decreased with the indentation depth, which is similar to that reported for other metallic glasses during nanoindentation test. Similarly, Young’s modulus also decreased with the applied load. Such mechanical
Table 2 The values of microhardness (H), Meyer’s exponent (n), material constant (K) and yield strength (r0) of consolidated Al76Ni8Ti8Zr4Y4 bulk alloys. Consolidated temperature (K)
n
K
H at 50 g load (GPa)
r0 at 50 g load (GPa)
623 723 823
1.773 1.761 1.794
187.5 214.3 258.8
9.61 11.71 12.48
5.48 6.55 6.90
X. Wang et al. / Journal of Alloys and Compounds 632 (2015) 617–622
621
Fig. 5. Different regions indentation morphologies under 50 g load consolidated at (a) 723 K; (b) 823 K.
Fig. 4. OM indentation morphologies under various loads (50 g, 100 g, 200 g) at different consolidation temperatures (a) 623 K; (b) 723 K; (c) 823 K.
softening has been reported for other metallic glasses [28,29] and attributed to the net creation of free volume during plastic deformation. The average values of nanohardness and young’s modulus are listed in Table 3. It is evident that nanohardness values are little higher in comparison with the microindentation results, while the maximum load (50 g 500 mN) and the size of indention is almost the same with the Vickers hardness test, therefore, the ISE effect should be unconspicuous. This may be caused by experimental errors including the measurement of indentation size, the applied load and the geometry of indenter [30]. In addition to Hn and E, Table 3 shows the values of Hn/E ratios, which are indicative of the wear resistance of the metallic glass. A
Fig. 6. Representative nanoindentation load–displacement curves of the bulk sample consolidated at three different temperatures, where the black arrow indicates a displacement burst. The inset shows the SEM nanoindentation impression of 623 K sintered sample.
high Hn/E ratio is often a reliable indicator of good wear resistance in a coating or layers, therefore, the almost amorphous bulk sample seems to have slightly higher wear resistance than the bulk alloy because of its lower E.
622
X. Wang et al. / Journal of Alloys and Compounds 632 (2015) 617–622
tested by Vickers hardness technique. The SBs corresponding to pop-ins in P–h curve were not observed around the indented area in the bulk specimen consolidated at 623 K, due to the unconspicuous shear band activity during plastic deformation of the present alloy. The microhardness in the matrix is a little higher than the white regions, because brittle nano-crystalline intermetallic like NiAl and TiAl was formed in the white regions, which caused the cracks formed around the indentation periphery. The indentation-induced microcracking become smaller with increasing consolidated temperature, which suggesting that the bulk alloys consolidated at elevated temperature have better fracture toughness than the low temperature sintering one. The crack paths around the indents of the sample emanate from the indent corners, indicating that the propagation of cracks occurred by radiation. Acknowledgments This work was jointly supported by the National Science Foundation of China (51301168 and 513711), the National Basic Research Program of China (2011CB610300) and the Instrument Developing Project of Chinese Academy of Sciences (YG2012065) References
Fig. 7. Nanohardness (a) and Young’s modulus (b) of the Al76Ni8Ti8Zr4Y4 alloy with the indentation depth for the different consolidation temperatures.
Table 3 The nanohardness, Young’s modulus and Hn/E for the Al76Ni8Ti8Zr4Y4 bulk alloys. Consolidated temperature (K)
Nanohardness, Hn (GPa)
Young’s modulus, E (GPa)
Hn/E ratio
623 723 823
11.16 12.74 13.27
173.3 213.9 223.1
0.064 0.060 0.059
4. Conclusions Bulk amorphous and nanocrystalline Al76Ni8Ti8Zr4Y4 alloy has been successfully prepared by mechanical alloying and high pressure sintering process. Structural examinations of the bulk materials reveal that most of the amorphous structure was retained after consolidation at 623 K, the crystallinity increased with the increasing temperature, the results also indicate that application of high pressure enhanced the crystallization process due to the negative volume change and affected the crystalline products of this process. Micro mechanical analysis showed that the microhardness of the bulk composites increased from 945HV0.1 to 1177HV0.1 with the consolidation temperature increasing. The evolution trends of the nanohardness and young’s modulus are similar with the values
[1] M.M. Trexler, N.N. Thadhani, Prog. Mater. Sci. 55 (2010) 759–839. [2] H.W. Yang, M.J. Tan, R.V. Ramanujan, Scripta Mater. 66 (2012) 382–385. [3] B.J. Yang, J.H. Yao, Y.S. Chao, J.Q. Wang, E. Ma, Philos. Mag. 90 (2010) 3215– 3231. [4] J.Q. Wang, P. Dong, W.L. Hou, X.C. Chang, M.X. Quan, J. Alloys Comp. 554 (2013) 419–425. [5] X.P. Li, M. Yan, H. Imai, K. Kondoh, J.Q. Wang, G.B. Schaffer, M. Qian, Mater. Sci. Eng. A 568 (2013) 155–159. [6] A.H. Cai, D.W. Ding, X. Xiong, Y. Liu, W.K. An, G.J. Zhou, Y. Luo, T.L. Li, X.S. Li, Mater. Des. 63 (2014) 233–237. [7] B.J. Yang, J.H. Yao, J. Zhang, H.W. Yang, J.Q. Wang, E. Ma, Scripta Mater. 61 (2009) 423–426. [8] M.H. Lee, D.H. Bae, D.H. Kim, D.J. Sordelet, J. Mater. Res. 18 (2003) 2101–2108. [9] T.T. Sasaki, K. Hono, J. Vierke, M. Wollgarten, J. Banhart, Mater. Sci. Eng. A 490 (2008) 343–350. [10] S. Mula, K. Mondal, S. Ghosh, S.K. Pabi, Mater. Sci. Eng. A 527 (2010) 3757– 3763. [11] M. Krasnowski, T. Kulik, J. Alloys Comp. 495 (2010) 382–385. [12] X. Wang, D. Wang, B. Zhu, Y. Li, F. Han, J. Non-Cryst. Solids 385 (2014) 111– 116. [13] W.H. Wang, D.W. He, D.Q. Zhao, Y.S. Yao, M. He, Appl. Phys. Lett. 75 (1999) 2770–2772. [14] F. Ye, K. Lu, Acta Mater. 47 (1999) 2449–2454. [15] Y.X. Zhuang, J.Z. Jiang, T.J. Zhou, H. Rasmussen, L. Gerward, M. Mezouar, W. Crichton, A. Inoue, Appl. Phys. Lett. 77 (2000) 4133–4135. [16] X.J. Gu, J.Q. Wang, F. Ye, K. Lu, J. Non-Cryst. Solids 296 (2001) 74–80. [17] P.F. Xing, Y.X. Zhuang, W.H. Wang, L. Gerward, J.Z. Jiang, J. Appl. Phys. 91 (2002) 4956–4960. [18] Z.X. Wang, F.Y. Li, M.X. Pan, D.Q. Zhao, W.H. Wang, J. Alloys Comp. 388 (2005) 262–265. [19] M. Krasnowski, A. Antolak-Dudka, T. Kulik, Intermetallics 19 (2011) 1243– 1249. [20] S. Vincent, J. Basu, B.S. Murty, J. Bhatt, Mater. Sci. Eng. A 550 (2012) 160–166. [21] Y.H. Kim, K. Hiraga, A. Inoue, T. Masumoto, H.H. Jo, Mater. Trans. JIM 35 (1994) 293–302. [22] X. Wei, F.S. Han, X.F. Wang, C.E. Wen, J. Alloys Comp. 501 (2010) 164–167. [23] N.K. Mukhopadhyay, G.C. Weatherly, J.D. Embury, Mater. Sci. Eng. A 315 (2001) 202–210. [24] D. Singh, D. Singh, R.K. Mandal, O.N. Srivastava, R.S. Tiwari, J. Alloys Comp. 590 (2014) 15–20. [25] B. Zhu, X.F. Wang, X.Y. Li, D. Wang, Y. Qin, F.S. Han, J. Alloys Comp. 586 (2014) 645–649. [26] I. Borner, J. Eckert, Scripta Mater. 45 (2001) 237–244. [27] S. Vincent, B.S. Murty, M.J. Kramer, J. Bhatt, Mater. Des. 65 (2015) 98–103. [28] N. Van Steenberge, J. Sort, A. Concustell, J. Das, S. Scudino, S. Surinach, J. Eckert, M.D. Baro, Scripta Mater. 56 (2007) 605–608. [29] J.I. Jang, B.G. Yoo, Y.J. Kim, J.H. Oh, I.C. Choi, H.B. Bei, Scripta Mater. 64 (2011) 753–756. [30] K. Sangwal, B. Surowska, P. Blaziak, Mater. Chem. Phys. 77 (2003) 511–520.