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Effects of substitution of La by other rare-earth elements on the glass forming ability of Al86Ni9La5 alloy
Journal of Non-Crystalline Solids 369 (2013) 1–4
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Effects of substitution of La by other rare-earth elements on the glass forming ability of Al86Ni9La5 alloy Zhang Zhang, Xian-zhong Xiong, Jiao-jiao Yi, Jin-fu Li ⁎ State Key Laboratory of Metal Matrix Composites, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China
a r t i c l e
i n f o
Article history: Received 15 January 2013 Received in revised form 14 March 2013 Available online xxxx Keywords: Metallic glass; Al alloy; Glass forming ability; Rare earth
a b s t r a c t The influence of La partially substituted by Sc, Dy, Gd or Ce on the glass forming ability (GFA) of Al86Ni9La5 was investigated by suction casting wedge-shaped samples with preparation conditions controlled. The results show that except Sc, such substitutions enhance the GFA of the alloy, and Al86Ni9La3.5Gd1.5 exhibits the best GFA among them. The GFA was analyzed from the thermodynamic and topological points of view. It was found that the GFA of Al86Ni9La3.5RE1.5 is strongly correlated to not only the thermodynamic driving force for the glass formation but also the atomic packing efficiency. Higher packing efficiency and stability of Al17RE correspond to better GFA. © 2013 Elsevier B.V. All rights reserved.
1. Introduction The Al-rich Al–TM–RE (TM = transition metal; RE = rare earth element) metallic glasses are highlighted by their high specific strength, good ductility and corrosion resistance, which makes them promising for potential structural applications [1–4]. However, most of the Al-based metallic glasses developed so far are limited in the form of ribbons or powders due to their relatively poor glass forming ability (GFA). To synthesize Al-based bulk metallic glasses (BMGs), many efforts have been made to improve the GFA. Recently, some Al-based BMG rods with diameter of 1 mm were obtained by simultaneously substituting Ni and Y by Co and La respectively in Al–Ni–Y system [5–7], indicating that mutual substitution of similar atoms is a promising way to improve the GFA. However, the reason for the GFA improvement has not been revealed sufficiently. High atomic packing efficiency is believed to be favorable for glass formation. Recently, some structural models, such as the efficient cluster packing (ECP) model [8–11] and the short-to-medium-range order model [12,13], were proposed to interpret the atomic-level structure of metallic glasses and the relationship between the local structure and the GFA. According to the ECP model, the most efficiently packed clusters in the RE–Al system and their coordination numbers vary with the type of RE due to the difference in atomic radii of the RE elements. Investigations on the local structure of Al-based metallic glasses using X-ray absorption fine structure (XAFS), extended X-ray absorption fine structure (EXAFS) or small-angle X-ray scattering (SAXS) methods [14–16] indicated that the coordination numbers of different RE elements in Al-based metallic glasses are also different. This suggests ⁎ Corresponding author. Tel./fax: +86 21 5474 8530. E-mail address: jfl[email protected] (J. Li). 0022-3093/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jnoncrysol.2013.03.018
that the atomic packing efficiency and correspondingly the GFA of Al-based metallic glasses strongly depend on the RE selected for substitution. Besides, the Pauli electronegativity of RE elements is different from each other, which will result in different thermodynamic behaviors upon alloying. Therefore, both the differences in atomic radius and electronegativity would result in different thermodynamic driving forces for glass formation when the RE in an Al-based alloy is partially substituted by other RE elements and, consequently, should lead to different GFAs. Although the influence of mutual substitution between similar atoms on the crystallization behavior has been widely studied [17–21], it remains unclear how the GFA will change when the RE in an Al-based alloy is partially substituted by other RE elements hitherto. The present work will therefore explore on this. Our previous work showed that the Al86Ni9La5 alloy, whose composition slightly deviates from the optimal glass former Al85.5Ni9.5La5 in the Al–Ni–La system, was an ideal base alloy. When substituting 1.5at.% La by Ce, the critical thickness of the amorphous part in a wedge-shaped sample was found to be much greater than that of the optimal glass former [22]. Therefore, the present work will focus on the influence of partial substitution of La by different RE elements on the GFA of the Al86Ni9La5 alloys. 2. Experimental Ingots of Al86Ni9La5 and Al86Ni9La3.5RE1.5 (RE = Sc, Ce, Gd, Dy) alloys were prepared by arc melting the mixtures of pure Al (99.999%), Ni (99.99%), La (99.9%), Sc (99.9%), Ce (99.9%), Gd (99.9%), and Dy (99.9%) metals in a titanium-gettered argon atmosphere under a water-cooled copper crucible. The alloy ingots were remelted six times to ensure compositional homogeneity. Before casting, the oxidized surface layer of the ingot was abraded away and the ingot was cut by a diamond saw into
2
Z. Zhang et al. / Journal of Non-Crystalline Solids 369 (2013) 1–4
pieces weighing about 2 g. The ingot pieces were then suction cast into a wedge-shaped copper mold with an included angle of 5° and a width of 10 mm. The arc current and melting duration were controlled to be the same in all operations. For each composition, at least two wedgeshaped samples were prepared. The sample was cross-sectioned using the diamond saw. One half of the sample was mounted, polished, and then etched using a modified Keller's reagent. Under an optical microscope, the maximum thickness tc of the amorphous part that characterizes the GFA was measured within an error of 10 μm. The structure of the amorphous part in the sample was examined by X-ray diffraction (XRD) with Cu Kα radiation. Thermal analysis of the amorphous alloy was performed using a Perkin-Elmer Pyres Diamond differential scanning calorimeter (DSC) at a heating rate of 20 K/min under a flow of high purity Ar atmosphere to obtain the glass transition temperature Tg and the onset temperature of crystallization Tx. In addition, the melting and solidification reactions of the alloys were measured using a Netzsch STA 449C differential thermal analysis (DTA) instrument at the same heating and cooling rates of 20 K/min. The liquidus temperature TL was determined from the cooling curve. 3. Results and discussion Fig. 1 shows the optical micrographs of the sections of the as-cast Al86Ni9La5 and Al86Ni9La3.5Gd1.5 wedge-shaped samples. Two distinct regions representing the fully amorphous and crystalline structure, respectively, and a transition region between can be clearly distinguished. The maximum thickness tc of the amorphous part of each sample, as well as the average value t(—)c, for an alloy are listed in Table 1. Apparently, the GFA shows a degrading tendency in the sequence of substituting elements Gd, Ce, Dy and Sc. In the other hand, it is noteworthy that the partial substitution of La by Dy, Gd and Ce enhances the GFA of Al86Ni9La5, while that by Sc does not. Besides, Al86Ni9La3.5Gd1.5 in particular has the best GFA among the alloys, and its critical thickness reaches 850 μm, larger than that of the optimal glass former Al85.5Ni9.5La5 (605 μm) in the ternary Al–Ni–La alloys [22]. Such results indicate that the selection of substituting RE elements is of great concern in improving the GFA of Al–Ni–RE alloys. Fig. 2 shows the XRD patterns taken from the amorphous parts of the samples. They consist only of one large broad peak, and no diffraction peaks corresponding to crystalline phases are seen, indicating that the tip part is amorphous. The TEM and HRTEM images and the selected area electron diffraction (SAED) pattern of the amorphous region in the wedge-shaped sample of Al86Ni9La3.5Gd1.5 are shown in Fig. 3. The featureless contrast in the TEM image, no visible crystals
Fig. 1. Optical micrograph of the cross-section of the wedge-shaped samples.
Table 1 Summary of the data of the Al-based metallic glasses investigated. Alloys
tc (μm)
tc (μm)
Tg (K)
Tx (K)
TL (K)
ΔTx (K)
Trg
γ
Al86Ni9La5 Al86Ni9La3.5Gd1.5 Al86Ni9La3.5Ce1.5 Al86Ni9La3.5Dy1.5 Al86Ni9La3.5Sc1.5
520/560 850/820 690/650 650/600 340/300
540 835 670 625 320
503 494 501 493 –
517 509 516 509 501
1097 1095 1092 1093 1110
14 15 15 16 –
0.458 0.452 0.459 0.451 0.451
0.323 0.320 0.324 0.321 –
in the HRTEM images and the broad halo in the SAED pattern further confirm the fully amorphous structure at t ≤ tc. Fig. 4 shows the DSC curves of the amorphous part in the wedge-shaped samples at a heating rate of 20 K/min. A glass transition and supercooled liquid region can be observed in all the samples except for Al86Ni9La3.5Sc1.5. The DTA curves of the alloys at a cooling rate of 20 K/min are shown in Fig. 5. The measured Tg, Tx and TL for each alloy are also listed in Table 1, along with the calculated GFA parameters such as ΔTx = Tx − Tg, Trg = Tg / TL and γ = Tx / (TL + Tg) [23,24]. No direct correlations are found between t(—)c and these GFA parameters. In order to understand the effect of similar atom substitution on the GFA of alloys, Li et al. [25] proposed an interpretation in the viewpoint of thermodynamics based on the ideal substitutional solid solution model. They argued that the GFA is correlated with the difference in energy between the solid glass and its liquid counterpart, G S − G L, which reflects the stability of the glass state and is proportional to the mixing free energy of ΔGmix of the liquid metals [26], which is given by: ΔGmix ¼ ΔH mix −TΔSmix
ð1Þ
where ΔHmix is the mixing enthalpy, ΔSmix the mixing entropy and T the temperature. ΔSmix is given by the sum of the configurational entropy ΔSconf and the mismatch entropy ΔSmis. Incidentally, Liu et al. [7] also suggested that the main thermodynamic driving force for the GFA improvement of Al-rich Al–Ni–Co–Y–La BMGs through similar atom substitution is the configurational entropy ΔSconf. The mixing enthalpies of Al–La, Al–Ce, Al–Gd, Al–Dy and Al–Sc are −38, −38, −39, −38 and −38 kJ/mol, and those of Ni–La, Ni–Ce, Ni–Gd, Ni–Dy and Ni–Sc are −27, −28, −31, −32 and −39 kJ/mol, respectively. Following the treatment of Li et al. [25], the effects of substitution content x of La on the ΔGmix for the Al86Ni9(La1 − xREx)5 (RE = Sc, Dy, Gd, Ce) alloys were calculated and shown in Fig. 6. The temperature was chosen to be 1200 K, slightly above the liquidus temperature of all the alloys. Obviously, a substitution of 30% La will always result in a reduction of ΔGmix, and consequently an enhanced driving force for glass formation. The decrement in ΔGmix for different RE substitutions is in the order of Gd, Ce, Dy and Sc. This would in turn leads to a different driving force for the glass formation and, therefore, the GFA of
Fig. 2. XRD patterns of the amorphous part in the wedge-shaped samples.
Z. Zhang et al. / Journal of Non-Crystalline Solids 369 (2013) 1–4
3
Fig. 3. TEM bright-field image and SAED pattern of the amorphous part of a wedge-shaped sample of Al86Ni9La3.5Gd1.5 (a) and the corresponding HRTEM image (b).
these alloys will be expected to be in the order of Al86Ni9La3.5Gd1.5 > Al86Ni9La3.5Ce1.5 > Al86Ni9La3.5Dy1.5 > Al86Ni9La3.5Sc1.5. However, the above interpretation can only partially explain the present observations, as the GFA of Al86Ni9La3.5Sc1.5 is seen to be lower than that of Al86Ni9La5, although it has a somehow larger driving force. The discrepancy seen above suggests that only considering the thermodynamics factor is not sufficient to interpret the GFA change. Other factors, such as the atomic packing structure, should be taken into account as well. In this regard, the ECP model has been proposed, according to which the composition facilitating the formation of solute centered clusters with high packing efficiency will have a high GFA [11]. Efficient packing clusters reduce the local volume, which in turn reduces the free energy of the system, increases the viscosity, and finally results in an enhanced GFA. The type of the most efficiently packed clusters occurs at the specific atomic ratios RN⁎ between solute and solvent elements, which can be calculated using Eq. (2) 8 4π > > . 9 for 0:225≤Rb0:414 8 > > 1 = > < . > > 1 > 2 > −π 6 arccos sinðπ=3Þ 1− > > : ; > ðRþ1Þ2 > > > > > 4π > > > . 9 for 0:414≤Rb0:902 8 < 1 = T < . N ¼ 1 2 > −2π > 8 arccos: sinðπ=4Þ 1− ðRþ1Þ2 > ; > > > > > > 4π > > . 9 for R≥0:902 8 > > > 1 = > < . > > 1 2 > 10 arccos sinðπ=5Þ 1− −3π > > : : ; ðRþ1Þ2
ð2Þ
Fig. 4. DSC curves of the amorphous part in the wedge-shaped samples. The glass transition on two curves is magnified in the inset as examples.
where R is the ratio of the solute to solvent atomic radius. The generated theoretical coordination number, NT, represents the most efficient or optimal configuration of a tightly packed cluster [11]. The calculated NTs of different RE elements are listed in Table 2, together with their atomic radii. In fact, solute-centered clusters in an Al-based metallic glass with a same coordination number are not the identical structure units. Computer simulations carried out on Al89Ni6La5 metallic glass revealed that La has an average coordination number of 17.5 with a rather narrow distribution. The dominant coordination numbers are 16, 17 and 18 with a percentage of about 13%, 46% and 30%, respectively. The La-centered cluster with coordination number N = 19 that is much larger than NT may also exist in the structure with a fraction smaller than 5% due to its large strain involved [13]. A simple representation of the local packing efficiency is obtained by a ratio of N/N T for the coordination number N b N T. When N > N T, the packing efficiency can be given by (rRE / (RN⁎ × rAl)) 3, where RN⁎ is the specific atomic ratio corresponding to N, and rRE and rAl the atomic radii of RE and Al (0.1432 nm), respectively [11]. The packing efficiencies of the RE-centered clusters with different coordination numbers are calculated accordingly and listed in Table 2. Obviously, the packing efficiencies of RE-centered clusters with the same coordination number vary with the type of RE. For La-centered clusters, the packing efficiency of the clusters decreases in an order of Al18La, Al17La and Al16La. Although the coordination numbers of La-centered clusters in Al86Ni9La5 are not known precisely, it can be speculated that most of them should also be the integers nearest to the NT since its GFA is much higher than that of Al89Ni6La5 [27]. That is to say, Al17La and Al18La are also the dominant clusters in Al89Ni9La5 metallic glass while Al16La is the third abundant cluster type. The maximal allowed number of the nearest Al atoms in the first coordination shell of an RE atom resulting from the pure topological
Fig. 5. DTA traces of the alloys at a cooling rate of 20 K/min.
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Z. Zhang et al. / Journal of Non-Crystalline Solids 369 (2013) 1–4
glasses. Therefore, it can be expected that the atomic packing density of Al86Ni9La3.5Gd1.5 is the highest among the quaternary Al–Ni–La–RE alloys and thus it would exhibit the best GFA. It is also noted that N T of Dy is 16.84 and that of Ce is 17.41, indicating that the cluster stability of Al17Dy is worse than that of Al17Ce. This may explain why the GFA of Al86Ni9La3.5Ce1.5 is better than Al86Ni9La3.5Dy1.5. 4. Conclusions
Fig. 6. Variation of the mixing Gibbs free energy ΔGmix with the substitution content x of RE in Al86Ni9(La1 − xREx)5.
In summary, the partial substitutions of La by Dy, Gd and Ce can enhance the GFA of Al86Ni9La5, and Al86Ni9La3.5Gd1.5 has the best GFA among these alloys. The substitution of La by Sc, however, deteriorates the GFA. The variation in GFA with substitution of La by other RE is dependent on not only the change in thermodynamic driving force for the glass formation but also the difference of Al17RE in packing efficiency and cluster stability. Acknowledgment
consideration should be the nearest integer to N T. If more than N T Al atoms are present in the first coordination shell of an RE atom, large stain will be induced and consequently worsen the cluster stability. Therefore, if the number of Al atoms in the first coordination shell of RE is much larger than N T, it will be adjusted to the maximal allowed number so as to reduce the free energy. When La is partially substituted by Sc, the original La-centered cluster structure may be destroyed seriously, since the N T of Sc is much smaller than that of La. As a result, some Al atoms will be free in the structure. These Al atoms would prefer to first rearrange into fcc-Al in the following crystallization. Therefore, no glass transition could be observed on the DSC curve of Al86Ni9La3.5Sc1.5 metallic glass. Correspondingly, it is not surprising that its average critical thickness is only 320 μm, much smaller than that of Al86Ni9La5. The N Ts of Dy, Gd and Ce are very close to 17. When La is partially substituted by one of these three REs, the La in Al16La and Al17La might be the site preferentially substituted. There are no free Al atoms to be released and no significant difference in local structure among the alloys due to such substitution. From Table 2, it is further known that the packing efficiency of Al16RE or Al17RE is increased when La is replaced by Dy, Gd or Ce, respectively. As a result, it may construct a denser structure and thus result in an improved GFA. For instance, it has been found experimentally that partial substitution of La by Ce decreases the amount of free volume and thus the packing density in Al–Ni–La system, resulting in the higher GFA [28]. However, the degree of packing density improvement is different as seen in Table 2 that the packing efficiencies are in the order of Gd, Ce(= Dy), La for Al17RE and Dy, Gd, Ce, La for Al16RE, respectively. As mentioned above, the percentage of Al17La is greater than that of Al16La in the structure of Al–Ni–La metallic
Table 2 Theoretical coordination number NT and packing efficiency of the Al–RE clusters with different coordination numbers. Element
Atom radius (nm)
NT
Cluster
Packing efficiency
La
0.1877
17.99
SC
0.1641
15.44
Dy
0.1773
16.84
Gd
0.1802
17.16
Ce
0.1825
17.41
Al16La Al17La Al18La Al14Sc Al15Sc Al16Sc Al16Dy Al17Dy Al18Dy Al16Gd Al17Gd Al18Gd Al16Ce Al17Ce Al18Ce
88.9% 94.4% 99.4% 90.7% 97.1% 90.9% 95.0% 97.6% 84.2% 93.2% 99.1% 88.4% 91.9% 97.6% 91.9%
The authors would like to acknowledge the Innovation Program of Shanghai Municipal Education Commission (grant no. 13ZZ016), the National Natural Science Foundation of China (grant no. 50831003) and the National Basic Research Program of China (grant no. 2011CB610405) for financial supports. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28]
Y. He, S.J. Poon, G.J. Shiflet, Science 241 (1988) 1640. A. Inoue, Prog. Mater. Sci. 43 (1998) 365. A.L. Greer, Science 267 (1995) 1947. J.H. Perepezko, Prog. Mater. Sci. 49 (2004) 263. B.J. Yang, J.H. Yao, J. Zhang, H.W. Yang, J.Q. Wang, E. Ma, Scr. Mater. 61 (2009) 423. B.J. Yang, J.H. Yao, Y.S. Chao, J.Q. Wang, E. Ma, Philos. Mag. 90 (2010) 3215. Y. Liu, G. Schumacher, I. Zizak, A. Erko, J. Banhart, Mater. Lett. 70 (2012) 171. D.B. Miracle, Nat. Mater. 3 (2004) 697. D.B. Miracle, Acta Mater. 54 (2006) 4317. D.B. Miracle, O.N. Senkov, J. Non-Cryst. Solids 319 (2006) 174. D.B. Miracle, W.S. Sanders, O.N. Senkov, Philos. Mag. 83 (2003) 2409. H.W. Sheng, W.K. Luo, F.M. Alamgir, J.M. Bai, E. Ma, Nature 439 (2006) 419. H.W. Sheng, Y.Q. Cheng, P.L. Lee, S.D. Shastri, E. Ma, Acta Mater. 56 (2008) 6264. R. Bacewicz, J. Antonowicz, Scr. Mater. 54 (2006) 1187. W. Zalewski, J. Antonowicz, R. Bacewicz, J. Latuch, J. Alloys Compd. 468 (2009) 40. J. Antonowicz, J. Mater. Sci. 45 (2010) 5040. R.D. Sá Lisboa, C. Bolfarini, W.J. Botta, C.S. Kiminami, Appl. Phys. Lett. 86 (2005) 211904. D.V. Louzguine, S. Sobu, A. Inoue, Appl. Phys. Lett. 85 (2004) 17. D.V. Louzguine, A. Inoue, J. Non-Cryst. Solids 311 (2002) 281. J. Basu, S. Ranganathan, Intermetallics 12 (2004) 1045. A. Inoue, S. Sobu, D.V. Louzguine, H. Kimura, K. Sasamori, J. Mater. Res. 19 (2004) 1539. Z. Zhang, W. Zhou, X. Lin, X.Z. Xiong, J.F. Li, Submitted to Metall. Mater. Trans. A. Z.P. Lu, C.T. Liu, Acta Mater. 50 (2002) 3501. Z.P. Lu, C.T. Liu, Phys. Rev. Lett. 91 (2003) 115505. R. Li, S.J. Pang, C.L. Ma, T. Zhang, Acta Mater. 55 (2007) 3719. A. Takeuchi, A. Inoue, Mater. Sci. Eng., A 304–306 (2001) 446. W.S. Sanders, J.S. Warner, D.B. Miracle, Intermetallics 14 (2006) 348. G.H. Li, W.M. Wang, X.F. Bian, J.T. Zhang, R. Li, L. Wang, J. Alloys Compd. 478 (2009) 745.
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Journal of Non-Crystalline Solids journal homepage: www.elsevier.com/ locate/ jnoncrysol
Effects of substitution of La by other rare-earth elements on the glass forming ability of Al86Ni9La5 alloy Zhang Zhang, Xian-zhong Xiong, Jiao-jiao Yi, Jin-fu Li ⁎ State Key Laboratory of Metal Matrix Composites, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China
a r t i c l e
i n f o
Article history: Received 15 January 2013 Received in revised form 14 March 2013 Available online xxxx Keywords: Metallic glass; Al alloy; Glass forming ability; Rare earth
a b s t r a c t The influence of La partially substituted by Sc, Dy, Gd or Ce on the glass forming ability (GFA) of Al86Ni9La5 was investigated by suction casting wedge-shaped samples with preparation conditions controlled. The results show that except Sc, such substitutions enhance the GFA of the alloy, and Al86Ni9La3.5Gd1.5 exhibits the best GFA among them. The GFA was analyzed from the thermodynamic and topological points of view. It was found that the GFA of Al86Ni9La3.5RE1.5 is strongly correlated to not only the thermodynamic driving force for the glass formation but also the atomic packing efficiency. Higher packing efficiency and stability of Al17RE correspond to better GFA. © 2013 Elsevier B.V. All rights reserved.
1. Introduction The Al-rich Al–TM–RE (TM = transition metal; RE = rare earth element) metallic glasses are highlighted by their high specific strength, good ductility and corrosion resistance, which makes them promising for potential structural applications [1–4]. However, most of the Al-based metallic glasses developed so far are limited in the form of ribbons or powders due to their relatively poor glass forming ability (GFA). To synthesize Al-based bulk metallic glasses (BMGs), many efforts have been made to improve the GFA. Recently, some Al-based BMG rods with diameter of 1 mm were obtained by simultaneously substituting Ni and Y by Co and La respectively in Al–Ni–Y system [5–7], indicating that mutual substitution of similar atoms is a promising way to improve the GFA. However, the reason for the GFA improvement has not been revealed sufficiently. High atomic packing efficiency is believed to be favorable for glass formation. Recently, some structural models, such as the efficient cluster packing (ECP) model [8–11] and the short-to-medium-range order model [12,13], were proposed to interpret the atomic-level structure of metallic glasses and the relationship between the local structure and the GFA. According to the ECP model, the most efficiently packed clusters in the RE–Al system and their coordination numbers vary with the type of RE due to the difference in atomic radii of the RE elements. Investigations on the local structure of Al-based metallic glasses using X-ray absorption fine structure (XAFS), extended X-ray absorption fine structure (EXAFS) or small-angle X-ray scattering (SAXS) methods [14–16] indicated that the coordination numbers of different RE elements in Al-based metallic glasses are also different. This suggests ⁎ Corresponding author. Tel./fax: +86 21 5474 8530. E-mail address: jfl[email protected] (J. Li). 0022-3093/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jnoncrysol.2013.03.018
that the atomic packing efficiency and correspondingly the GFA of Al-based metallic glasses strongly depend on the RE selected for substitution. Besides, the Pauli electronegativity of RE elements is different from each other, which will result in different thermodynamic behaviors upon alloying. Therefore, both the differences in atomic radius and electronegativity would result in different thermodynamic driving forces for glass formation when the RE in an Al-based alloy is partially substituted by other RE elements and, consequently, should lead to different GFAs. Although the influence of mutual substitution between similar atoms on the crystallization behavior has been widely studied [17–21], it remains unclear how the GFA will change when the RE in an Al-based alloy is partially substituted by other RE elements hitherto. The present work will therefore explore on this. Our previous work showed that the Al86Ni9La5 alloy, whose composition slightly deviates from the optimal glass former Al85.5Ni9.5La5 in the Al–Ni–La system, was an ideal base alloy. When substituting 1.5at.% La by Ce, the critical thickness of the amorphous part in a wedge-shaped sample was found to be much greater than that of the optimal glass former [22]. Therefore, the present work will focus on the influence of partial substitution of La by different RE elements on the GFA of the Al86Ni9La5 alloys. 2. Experimental Ingots of Al86Ni9La5 and Al86Ni9La3.5RE1.5 (RE = Sc, Ce, Gd, Dy) alloys were prepared by arc melting the mixtures of pure Al (99.999%), Ni (99.99%), La (99.9%), Sc (99.9%), Ce (99.9%), Gd (99.9%), and Dy (99.9%) metals in a titanium-gettered argon atmosphere under a water-cooled copper crucible. The alloy ingots were remelted six times to ensure compositional homogeneity. Before casting, the oxidized surface layer of the ingot was abraded away and the ingot was cut by a diamond saw into
2
Z. Zhang et al. / Journal of Non-Crystalline Solids 369 (2013) 1–4
pieces weighing about 2 g. The ingot pieces were then suction cast into a wedge-shaped copper mold with an included angle of 5° and a width of 10 mm. The arc current and melting duration were controlled to be the same in all operations. For each composition, at least two wedgeshaped samples were prepared. The sample was cross-sectioned using the diamond saw. One half of the sample was mounted, polished, and then etched using a modified Keller's reagent. Under an optical microscope, the maximum thickness tc of the amorphous part that characterizes the GFA was measured within an error of 10 μm. The structure of the amorphous part in the sample was examined by X-ray diffraction (XRD) with Cu Kα radiation. Thermal analysis of the amorphous alloy was performed using a Perkin-Elmer Pyres Diamond differential scanning calorimeter (DSC) at a heating rate of 20 K/min under a flow of high purity Ar atmosphere to obtain the glass transition temperature Tg and the onset temperature of crystallization Tx. In addition, the melting and solidification reactions of the alloys were measured using a Netzsch STA 449C differential thermal analysis (DTA) instrument at the same heating and cooling rates of 20 K/min. The liquidus temperature TL was determined from the cooling curve. 3. Results and discussion Fig. 1 shows the optical micrographs of the sections of the as-cast Al86Ni9La5 and Al86Ni9La3.5Gd1.5 wedge-shaped samples. Two distinct regions representing the fully amorphous and crystalline structure, respectively, and a transition region between can be clearly distinguished. The maximum thickness tc of the amorphous part of each sample, as well as the average value t(—)c, for an alloy are listed in Table 1. Apparently, the GFA shows a degrading tendency in the sequence of substituting elements Gd, Ce, Dy and Sc. In the other hand, it is noteworthy that the partial substitution of La by Dy, Gd and Ce enhances the GFA of Al86Ni9La5, while that by Sc does not. Besides, Al86Ni9La3.5Gd1.5 in particular has the best GFA among the alloys, and its critical thickness reaches 850 μm, larger than that of the optimal glass former Al85.5Ni9.5La5 (605 μm) in the ternary Al–Ni–La alloys [22]. Such results indicate that the selection of substituting RE elements is of great concern in improving the GFA of Al–Ni–RE alloys. Fig. 2 shows the XRD patterns taken from the amorphous parts of the samples. They consist only of one large broad peak, and no diffraction peaks corresponding to crystalline phases are seen, indicating that the tip part is amorphous. The TEM and HRTEM images and the selected area electron diffraction (SAED) pattern of the amorphous region in the wedge-shaped sample of Al86Ni9La3.5Gd1.5 are shown in Fig. 3. The featureless contrast in the TEM image, no visible crystals
Fig. 1. Optical micrograph of the cross-section of the wedge-shaped samples.
Table 1 Summary of the data of the Al-based metallic glasses investigated. Alloys
tc (μm)
tc (μm)
Tg (K)
Tx (K)
TL (K)
ΔTx (K)
Trg
γ
Al86Ni9La5 Al86Ni9La3.5Gd1.5 Al86Ni9La3.5Ce1.5 Al86Ni9La3.5Dy1.5 Al86Ni9La3.5Sc1.5
520/560 850/820 690/650 650/600 340/300
540 835 670 625 320
503 494 501 493 –
517 509 516 509 501
1097 1095 1092 1093 1110
14 15 15 16 –
0.458 0.452 0.459 0.451 0.451
0.323 0.320 0.324 0.321 –
in the HRTEM images and the broad halo in the SAED pattern further confirm the fully amorphous structure at t ≤ tc. Fig. 4 shows the DSC curves of the amorphous part in the wedge-shaped samples at a heating rate of 20 K/min. A glass transition and supercooled liquid region can be observed in all the samples except for Al86Ni9La3.5Sc1.5. The DTA curves of the alloys at a cooling rate of 20 K/min are shown in Fig. 5. The measured Tg, Tx and TL for each alloy are also listed in Table 1, along with the calculated GFA parameters such as ΔTx = Tx − Tg, Trg = Tg / TL and γ = Tx / (TL + Tg) [23,24]. No direct correlations are found between t(—)c and these GFA parameters. In order to understand the effect of similar atom substitution on the GFA of alloys, Li et al. [25] proposed an interpretation in the viewpoint of thermodynamics based on the ideal substitutional solid solution model. They argued that the GFA is correlated with the difference in energy between the solid glass and its liquid counterpart, G S − G L, which reflects the stability of the glass state and is proportional to the mixing free energy of ΔGmix of the liquid metals [26], which is given by: ΔGmix ¼ ΔH mix −TΔSmix
ð1Þ
where ΔHmix is the mixing enthalpy, ΔSmix the mixing entropy and T the temperature. ΔSmix is given by the sum of the configurational entropy ΔSconf and the mismatch entropy ΔSmis. Incidentally, Liu et al. [7] also suggested that the main thermodynamic driving force for the GFA improvement of Al-rich Al–Ni–Co–Y–La BMGs through similar atom substitution is the configurational entropy ΔSconf. The mixing enthalpies of Al–La, Al–Ce, Al–Gd, Al–Dy and Al–Sc are −38, −38, −39, −38 and −38 kJ/mol, and those of Ni–La, Ni–Ce, Ni–Gd, Ni–Dy and Ni–Sc are −27, −28, −31, −32 and −39 kJ/mol, respectively. Following the treatment of Li et al. [25], the effects of substitution content x of La on the ΔGmix for the Al86Ni9(La1 − xREx)5 (RE = Sc, Dy, Gd, Ce) alloys were calculated and shown in Fig. 6. The temperature was chosen to be 1200 K, slightly above the liquidus temperature of all the alloys. Obviously, a substitution of 30% La will always result in a reduction of ΔGmix, and consequently an enhanced driving force for glass formation. The decrement in ΔGmix for different RE substitutions is in the order of Gd, Ce, Dy and Sc. This would in turn leads to a different driving force for the glass formation and, therefore, the GFA of
Fig. 2. XRD patterns of the amorphous part in the wedge-shaped samples.
Z. Zhang et al. / Journal of Non-Crystalline Solids 369 (2013) 1–4
3
Fig. 3. TEM bright-field image and SAED pattern of the amorphous part of a wedge-shaped sample of Al86Ni9La3.5Gd1.5 (a) and the corresponding HRTEM image (b).
these alloys will be expected to be in the order of Al86Ni9La3.5Gd1.5 > Al86Ni9La3.5Ce1.5 > Al86Ni9La3.5Dy1.5 > Al86Ni9La3.5Sc1.5. However, the above interpretation can only partially explain the present observations, as the GFA of Al86Ni9La3.5Sc1.5 is seen to be lower than that of Al86Ni9La5, although it has a somehow larger driving force. The discrepancy seen above suggests that only considering the thermodynamics factor is not sufficient to interpret the GFA change. Other factors, such as the atomic packing structure, should be taken into account as well. In this regard, the ECP model has been proposed, according to which the composition facilitating the formation of solute centered clusters with high packing efficiency will have a high GFA [11]. Efficient packing clusters reduce the local volume, which in turn reduces the free energy of the system, increases the viscosity, and finally results in an enhanced GFA. The type of the most efficiently packed clusters occurs at the specific atomic ratios RN⁎ between solute and solvent elements, which can be calculated using Eq. (2) 8 4π > > . 9 for 0:225≤Rb0:414 8 > > 1 = > < . > > 1 > 2 > −π 6 arccos sinðπ=3Þ 1− > > : ; > ðRþ1Þ2 > > > > > 4π > > > . 9 for 0:414≤Rb0:902 8 < 1 = T < . N ¼ 1 2 > −2π > 8 arccos: sinðπ=4Þ 1− ðRþ1Þ2 > ; > > > > > > 4π > > . 9 for R≥0:902 8 > > > 1 = > < . > > 1 2 > 10 arccos sinðπ=5Þ 1− −3π > > : : ; ðRþ1Þ2
ð2Þ
Fig. 4. DSC curves of the amorphous part in the wedge-shaped samples. The glass transition on two curves is magnified in the inset as examples.
where R is the ratio of the solute to solvent atomic radius. The generated theoretical coordination number, NT, represents the most efficient or optimal configuration of a tightly packed cluster [11]. The calculated NTs of different RE elements are listed in Table 2, together with their atomic radii. In fact, solute-centered clusters in an Al-based metallic glass with a same coordination number are not the identical structure units. Computer simulations carried out on Al89Ni6La5 metallic glass revealed that La has an average coordination number of 17.5 with a rather narrow distribution. The dominant coordination numbers are 16, 17 and 18 with a percentage of about 13%, 46% and 30%, respectively. The La-centered cluster with coordination number N = 19 that is much larger than NT may also exist in the structure with a fraction smaller than 5% due to its large strain involved [13]. A simple representation of the local packing efficiency is obtained by a ratio of N/N T for the coordination number N b N T. When N > N T, the packing efficiency can be given by (rRE / (RN⁎ × rAl)) 3, where RN⁎ is the specific atomic ratio corresponding to N, and rRE and rAl the atomic radii of RE and Al (0.1432 nm), respectively [11]. The packing efficiencies of the RE-centered clusters with different coordination numbers are calculated accordingly and listed in Table 2. Obviously, the packing efficiencies of RE-centered clusters with the same coordination number vary with the type of RE. For La-centered clusters, the packing efficiency of the clusters decreases in an order of Al18La, Al17La and Al16La. Although the coordination numbers of La-centered clusters in Al86Ni9La5 are not known precisely, it can be speculated that most of them should also be the integers nearest to the NT since its GFA is much higher than that of Al89Ni6La5 [27]. That is to say, Al17La and Al18La are also the dominant clusters in Al89Ni9La5 metallic glass while Al16La is the third abundant cluster type. The maximal allowed number of the nearest Al atoms in the first coordination shell of an RE atom resulting from the pure topological
Fig. 5. DTA traces of the alloys at a cooling rate of 20 K/min.
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Z. Zhang et al. / Journal of Non-Crystalline Solids 369 (2013) 1–4
glasses. Therefore, it can be expected that the atomic packing density of Al86Ni9La3.5Gd1.5 is the highest among the quaternary Al–Ni–La–RE alloys and thus it would exhibit the best GFA. It is also noted that N T of Dy is 16.84 and that of Ce is 17.41, indicating that the cluster stability of Al17Dy is worse than that of Al17Ce. This may explain why the GFA of Al86Ni9La3.5Ce1.5 is better than Al86Ni9La3.5Dy1.5. 4. Conclusions
Fig. 6. Variation of the mixing Gibbs free energy ΔGmix with the substitution content x of RE in Al86Ni9(La1 − xREx)5.
In summary, the partial substitutions of La by Dy, Gd and Ce can enhance the GFA of Al86Ni9La5, and Al86Ni9La3.5Gd1.5 has the best GFA among these alloys. The substitution of La by Sc, however, deteriorates the GFA. The variation in GFA with substitution of La by other RE is dependent on not only the change in thermodynamic driving force for the glass formation but also the difference of Al17RE in packing efficiency and cluster stability. Acknowledgment
consideration should be the nearest integer to N T. If more than N T Al atoms are present in the first coordination shell of an RE atom, large stain will be induced and consequently worsen the cluster stability. Therefore, if the number of Al atoms in the first coordination shell of RE is much larger than N T, it will be adjusted to the maximal allowed number so as to reduce the free energy. When La is partially substituted by Sc, the original La-centered cluster structure may be destroyed seriously, since the N T of Sc is much smaller than that of La. As a result, some Al atoms will be free in the structure. These Al atoms would prefer to first rearrange into fcc-Al in the following crystallization. Therefore, no glass transition could be observed on the DSC curve of Al86Ni9La3.5Sc1.5 metallic glass. Correspondingly, it is not surprising that its average critical thickness is only 320 μm, much smaller than that of Al86Ni9La5. The N Ts of Dy, Gd and Ce are very close to 17. When La is partially substituted by one of these three REs, the La in Al16La and Al17La might be the site preferentially substituted. There are no free Al atoms to be released and no significant difference in local structure among the alloys due to such substitution. From Table 2, it is further known that the packing efficiency of Al16RE or Al17RE is increased when La is replaced by Dy, Gd or Ce, respectively. As a result, it may construct a denser structure and thus result in an improved GFA. For instance, it has been found experimentally that partial substitution of La by Ce decreases the amount of free volume and thus the packing density in Al–Ni–La system, resulting in the higher GFA [28]. However, the degree of packing density improvement is different as seen in Table 2 that the packing efficiencies are in the order of Gd, Ce(= Dy), La for Al17RE and Dy, Gd, Ce, La for Al16RE, respectively. As mentioned above, the percentage of Al17La is greater than that of Al16La in the structure of Al–Ni–La metallic
Table 2 Theoretical coordination number NT and packing efficiency of the Al–RE clusters with different coordination numbers. Element
Atom radius (nm)
NT
Cluster
Packing efficiency
La
0.1877
17.99
SC
0.1641
15.44
Dy
0.1773
16.84
Gd
0.1802
17.16
Ce
0.1825
17.41
Al16La Al17La Al18La Al14Sc Al15Sc Al16Sc Al16Dy Al17Dy Al18Dy Al16Gd Al17Gd Al18Gd Al16Ce Al17Ce Al18Ce
88.9% 94.4% 99.4% 90.7% 97.1% 90.9% 95.0% 97.6% 84.2% 93.2% 99.1% 88.4% 91.9% 97.6% 91.9%
The authors would like to acknowledge the Innovation Program of Shanghai Municipal Education Commission (grant no. 13ZZ016), the National Natural Science Foundation of China (grant no. 50831003) and the National Basic Research Program of China (grant no. 2011CB610405) for financial supports. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28]
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