Cu metallic glasses

Accepted Manuscript Composition dependent relaxation in La-Al-Ni/Cu metallic glasses J. Zhang, X.D. Wang, Y. Su, Y. Tang, T.D. Xu, Q.P. Cao, J.Z. Jiang PII:

S0925-8388(17)32693-2

DOI:

10.1016/j.jallcom.2017.07.307

Reference:

JALCOM 42723

To appear in:

Journal of Alloys and Compounds

Received Date: 23 April 2017 Revised Date:

12 July 2017

Accepted Date: 28 July 2017

Please cite this article as: J. Zhang, X.D. Wang, Y. Su, Y. Tang, T.D. Xu, Q.P. Cao, J.Z. Jiang, Composition dependent relaxation in La-Al-Ni/Cu metallic glasses, Journal of Alloys and Compounds (2017), doi: 10.1016/j.jallcom.2017.07.307. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

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Composition dependent relaxation in La-Al-Ni/Cu metallic glasses

J. Zhang, X.D. Wang,a Y. Su, Y. Tang, T. D. Xu, Q.P. Cao, J.Z. Jiangb

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International Center for New-Structured Materials (ICNSM), Laboratory of New-Structured Materials, State Key Laboratory of Silicon Materials, and School of Materials Science and

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Engineering, Zhejiang University, Hangzhou, 310027, People’s Republic of China

Author to whom correspondence should be addressed: (a) [email protected], (b)

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[email protected]

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ACCEPTED MANUSCRIPT We investigate structural relaxations in La-Al-Ni/Cu metallic glasses and find that the La-Al-Ni glasses tend to have a significant β -relaxation, which becomes more pronounced and shifts to high temperatures when increasing Ni content from 10 to 30 at.% in La85-xAl15Nix or Al content from 10 to 25 at.% in La70AlxNi30-x glasses. In

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contrast, the β -relaxation in La-Al-Cu glasses is weak and almost disappears after annealing. Compared with La-Al-Cu glasses, La-Al-Ni glasses usually have higher Tg temperatures and bigger changes in Tg by sub-Tg annealing. Different relaxation

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behaviors between La-Al-Ni and La-Al-Cu glasses changing with compositions and thermal history indicate that the free volume annihilation is not the key factor

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controlling the β -relaxation in La-Al-Ni glasses.

Keywords: β -structural relaxation; Dynamical mechanical analysis (DMA); Free volume

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annihilation; Annealing.

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ACCEPTED MANUSCRIPT 1. Introduction Due to the disordered atomic packing formed by rapid quenching, metallic glasses (MGs) are metastable and would finally evolve into their stable crystalline counterparts after suitable time and temporal conditions [1-5]. This physical aging process is of great importance for

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metastable materials since it largely affects their properties [6,7]. Although this issue has been studied for decades, it is still unclear how the atomic structure evolves during the ageing and how to control it precisely. During ageing, two kinds of structural relaxations, the β - and α-relaxations, were proposed to exist in a glass [8]. From the curve of loss modulus vs.

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temperature measured by dynamical mechanical analysis (DMA), the β -relaxation of MGs

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usually appears as an excess wing at the left side of the α-relaxation [9]. More recently, it has been reported that the different element addition can affect the β-relaxation behaviors. For example, La-Al-Ni glasses exhibit an isolated hump below the glass transition Tg measured by DMA, while the hump becomes unpronounced when Ni atoms are replaced by Cu atoms [10-12]. However, how the β-relaxation changes with different element substitutions and their

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thermal histories are still not completely understood. In this work, we carried out a systematical investigation for the influences of Ni and Cu contents and pre-annealing treatments at 0.9Tg, to the β-relaxation behavior in both La-Al-Ni/Cu MG systems with compositions of La85-xAl15Nix (x =10 to 30 at.%) and La70AlxNi30-x (x =10 to 25 at. %)

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together with three La85-xAl15Cux (x =10 to 20 at.%) for a comparison.The different thermal responses in La-Al-Ni and La-Al-Cu glasses are well observed and the mechanism at behind

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is also suggested.

2. Experimental methods

Alloy ingots were prepared by arc melting constituent elements with purity higher than 99.9 at.% in a Zr-gettered argon atmosphere. Each ingot was remelted four times to improve the chemical homogeneity. Ribbon samples of about 0.03 mm thick and 3 mm wide were prepared by single-roller melt spinning on a rotating Cu wheel with a surface velocity of 30 m/s in Ar atmosphere. The amorphous structure of samples was verified by x-ray diffraction (XRD) with Cu Kα radiation and differential scanning calorimetry (Perkin Elmer diamond 3

ACCEPTED MANUSCRIPT DSC). The dynamical mechanical analysis (DMA) was measured on a TA Q800 DMA using film tension mode in a nitrogen-flushed atmosphere. The storage modulus E′ and loss modulus E″ of ribbon samples were measured by temperature ramp mode with constant stress of about 7.5 MPa and 1 Hz at a heating rate of 3 K/min. Before starting the temperature

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scanning, the system was held at 303 K for 5 min to guarantee a constant initial stage. The as-spun ribbons were also vacuum-sealed in quartz tubes and annealed in a furnace for 4 days at about 0.9Tg of their own. Then the pre-annealed samples were isochronally examined by DSC at a heating rate of 20 K/min and by DMA under temperature ramp mode at a heating

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rate of 3 K/min.

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3. Results and discussion

3.1 DMA, XRD and DSC measurements for the as-spun samples We used DMA to measure the storage modulus (E’) and loss modulus (E”) of all the ribbon samples changing with temperatures. Figs. 1(a) and (b) illustrate the changes of E’ and E” of La85-xAl15Nix (x = 10 to 30 at.%) glasses normalized by the modulus at the initial stage.

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A steep decrease of E’ corresponds to the happening of the glass transition, which can be also seen from a rapid increase of E”. However, a gradual increase in E’ of La70Al15Ni15 is distinct above 1.15T0 (T0= 305 K), which likely originates from the annihilation of excess free volume.

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With reducing La content (partially replaced by Ni) of alloys, not only Tg but also Tβ temperature, at which the β-relaxation happens as a hump below Tg, increases. In addition, the double α-relaxation processes seem to appear in the La75Al15Ni10 glass upon heating, on

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which more details will be discussed in the last section. Fig. 1(c) displays the normalized loss modulus (E”/E0”) changing with normalized temperature (T/Tg). It shows that the normalized β-relaxation temperatures of all the glasses are almost the same, indicating that the variation of La and Ni contents affects the β-relaxation and the α-relaxation almost in a similar manner. For La70AlxNi30-x (x = 10 to 25 at.%) glasses, Fig. 2(a) shows that the value of E’ changes with T/T0, the steep decrease corresponding to the beginning of the glass transition. It indicates that the more Al addition to substitute for Ni, the higher is the Tg temperature, which can also be observed from the peak shift on E” curves in Fig. 2(b). Meanwhile, the hump position below Tg, corresponding to the β-relaxation shifts to high temperatures with more Al 4

ACCEPTED MANUSCRIPT addition. Thus, using Al to replace Ni atoms in some sense can further increase interatomic interactions and promote the happening of both relaxation processes to high temperatures. Figs. 2(c) and (d) displays the curves of E’/E0’ and E”/E0”changing with T/T0 for La85-xAl15Cux (x = 10, 15 and 20 at. %) glasses. All three glasses do not show a clear hump at

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the left side of the major α-relaxation, which are different from La85-xAl15Nix (x = 10 ~ 30 at. %) glasses.

Fig. 3 shows the XRD patterns for as-spun La-Al-Ni/Cu ribbon samples, which display a major broad diffraction peak at about 25 ~ 35°and another weak hump at about 50 ~ 60°

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with Cu Kα radiation, demonstrating that all the samples are fully amorphous. In Fig. 3(a), note that when holding Al at 15 at.% the major peak position shifts to low 2θ values with

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increasing La content. It means that with addition of more large La atoms the average atomic distances in samples increase. However, when keeping La content to be constant as shown in Fig. 3(b), the peak position of the first maximum moves to large 2θ values, corresponding to the decrease of the average atomic distance with increase of Al content. This phenomenon indicates that the interatomic interactions between La and Al atoms could be much stronger

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than those between La and Ni atoms although the atomic radius of Al (1.43Å) is larger than that of Ni (1.25Å). For La-Al-Cu glasses, Fig. 3(c) indicates that the average distance has a similar trend with La-Al-Ni glasses for which Al content being held as 15 at.% and only

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tuning the content ratio between La and Ni.

Fig. 4 shows DSC traces for La-Al-Ni/-Cu ribbon glasses. One can see that the Tg temperature almost increases with decrease of La content in La85-xAl15Nix (x = 10 to 30 at.%)

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series and with increase of Al content in La70AlxNi30-x (x = 10 to 25 at. %) glasses. The detailed characteristic temperature values are listed in Table 1. Since the Tg temperature is in some sense correlated with the strength of the glass [13], i.e., the lower Tg, the weak interatomic interactions between components, consistent with their average atomic distance in Table 1 estimated by XRD results. With increase of La content, the first crystallization event shifts to low temperatures and other events of multiple phase transformations occur over a large temperature range prior to the melting, indicating that the compositions are likely off-eutectic [14]. For La75Al15Ni10 glass, the glass transition becomes unpronounced on the DSC curve due to a weak crystallization event just above Tg, forming some ordering of 5

ACCEPTED MANUSCRIPT nano-clusters. The Tg temperature of this alloy can be detected by using DMA or DSC after its annealing. As a comparison, the Tg temperatures of La-Al-Cu glasses are generally lower than those of the same composition La-Al-Ni glasses. Also, three La-Al-Cu glasses show reduced

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Tg temperatures and multiple crystallization behaviors.

3.2 Annealing effect on thermal stability

As a metastable system, the stability of MGs is tightly related to their thermal history. It was reported that the annealing plays important roles in the stability of MGs, which not only

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affects the atomic dynamics in the glass state but also the following crystallization behavior in the supercooled liquid region [15,16]. Thus, we also studied how the thermal behavior

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changes after the sample was annealed at 0.9Tg for 4 days. The XRD patterns of all the annealed samples in Fig. 3 are quite similar to those of as-spun ones, indicating that they are still fully amorphous. Usually, the annealing makes the atomic packing denser, which are exactly the cases in these three series of glasses and can be reflected from the average atomic distance change, i.e., the values of ∆R/R in Table 1.

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The DSC data in Fig. 4 show that the annealing improves the stability of the glass, increasing the Tg temperature of about 3 ~ 10 K for all samples. The exact temperature values can be found in Table 1. In contrast, the Tx temperatures for the annealed La-Al-Ni glasses

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generally increase a little while decrease for some annealed La-Al-Cu glasses, indicating that La-Al-Ni glasses have relatively high thermal stability compared with La-Al-Cu glasses. In addition, we found that the trend of average atomic distance change is coincided with the

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reduced glass transition temperature Trg = Tg/Tl [17], i.e., the shorter the average atomic distance, the larger the Trg values, suggesting that the closed atomic packing generally favors the high glass forming ability. Figs. 5(a) (b) and (c) show the common feature of E”/E0”changing with T/T0 for three selected La-Al-Ni glasses (Others have almost the same variation trends). Note that the annealing does not make the β-relaxation disappear but shifts it to high temperatures as well as the α-relaxation. Also, the β-relaxation becomes more pronounced to separate from the α-relaxation. Besides, the excess wing between the β-relaxation and the α-relaxation gets steeper after sub-Tg annealing. Different from La-Al-Ni glasses, Figs. 5(d), (e) and (f) show 6

ACCEPTED MANUSCRIPT the weak hump (or excess wing) on the curves of E”/E0”changing with T/T0 for La-Al-Cu glasses almost disappears for the samples after annealing at 0.9Tg for 4 days, which means that the excess wing of La-Al-Cu glasses is likely determined by annihilation of free volume or internal stress. The same phenomenon was found in Mg-Cu-Y glass [18], i.e., the annealing

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removed the excess wing on DMA curve of E” ~ T before Tg. The authors suggested that it might be not a real β-relaxation but the broadening of the α-relaxation and could be described by the fictive temperature (Tf) change. It was also indicated that the annealing annihilates excess free volume and reduces spatial heterogeneity formed during the glass formation, thus

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suppressing the atomic motion in the glass and usually decreasing the strength of the β-relaxation [19, 20]. However, such scenarios are almost consistent with the case in

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La-Al-Cu glasses but cannot explain the enhanced β-relaxation in annealed La-Al-Ni glasses here. The results that the β-relaxation becomes more isolated in La-Al-Ni glasses but unpronounced in La-Al-Cu glasses after annealing indicate that the relaxation behaviors in both types of MGs are not similar and could not be the globe free volume control in La-Al-Ni

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glasses.

3.3 Whether “double” α-relaxations exist

In Figs. 5(c) and (f), note that there are two peaks on DMA curves of E”/E0” changing

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with T/T0 for the as-spun La75Al15Ni10 and La75Al15Cu10 glasses. It seems that there might be two α-relaxation processes in these two glasses upon heating [21]. To confirm the existence of double α-relaxations, a sample of La75Al15Ni10 glass was preheated to 450 K, the valley

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temperature after the first α-relaxation peak, and then cooled to room temperature. Although the as-spun sample exhibits two “α-relaxation” peaks at about 430 K and 485 K, for the pre-heated sample in Fig. 6(a), the first α-relaxation peak becomes weak and shifts to 456 K and the second peak slightly moves to 490 K. Fig. 6(b) shows the DSC curve of as-spun La75Al15Ni10 glass, displaying a weak but broad exothermic peak starting from Tx1 of about 438 K and another distinct exothermic peak from Tx2 of about 510 K using a heating rate of 20 K/min. When preheated the as-spun sample to 473 K, the DSC curve shows that the weak broad exothermic peak at the low temperature almost disappears. When enlarging the DSC curve of the preheated sample, the Tg temperature still can be seen to be about 452 K, and Tx1 7

ACCEPTED MANUSCRIPT and Tx2 being about 470 K and 516 K, respectively. Compared to the as-spun glass, the preheating treatment promotes the first exothermic event to happen and increases the stability of the residual amorphous phase (Tg ~ 456 K). This can be proved by XRD patterns in Fig. 6 (c). When heating to 443 K, some crystalline phase has already precipitated from the

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amorphous matrix. With temperature rising to 473 K, the content of crystalline phase La with space group P63/mcc significantly increases. Thus, the preheating treatment makes part of the amorphous crystallize, which weakens the first α-relaxation peak. When the temperature continues to rise, the residual amorphous begins to crystallize. In brief, our results confirm

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that the first peak on E”/E0”vs. T/T0 curve of as-spun La75Al15Ni10 glass depends on the glass transition and the formation of crystalline nano-clusters [22, 23]. The second peak is due to

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the crystallization of the rest amorphous. In addition, the La75Al15Ni10 glass shows the low Tg temperature and low glass forming ability, and weak β -relaxation.

4. Conclusions

By using DMA and DSC techniques, the thermal behaviors of as-spun and annealed

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La-Al-Ni/Cu glasses have been studied. It is found that the β -relaxation is more pronounced and shifts to high temperatures when increasing Ni content from 10 to 30 at.% in La85-xAl15Nix or Al content from 10 to 25 at.% in La70AlxNi30-x glasses. Also, the

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pre-annealing can stabilize the β -relaxation and separate it from the α-relaxation. In contrast,

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the β -relaxation is unpronounced in La-Al-Cu glasses and almost disappears in the annealed samples. The results demonstrate that the relaxation in La-Al-Ni glasses could not be the globe free volume control. Moreover, the nano-clusters formation near the Tg in the as-spun La75Al15Ni10 glass makes the glass transition unclear, which can stabilize the rest amorphous to high temperatures by pre-annealing.

Acknowledgements Financial supports from the National Natural Science Foundation of China (U1532115, 51671169, 51371157, U1432105, U1432110, and 51671170), the National Key Research and 8

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ductility at room temperature, Mater. Today 19 (2016) 568–579.

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[6] M. Calvo-Dahlborg, Structure and embrittlement of metallic glasses, Mater. Sci. Eng. A226-228 (1997) 833-845.

[7] C.A. Schuh, T.C. Hufnagel, U. Ramamurty, Mechanical behavior of amorphous alloys,

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Acta Mater. 55 (2007) 4067-4109.

[8] K.L. Ngai, M. Paluch, Classification of secondary relaxation in glass-formers based on dynamic properties, J. Chem. Phys. 120 (2004) 857.

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[9] W.H. Wang, P. Wen, X.F. Liu, The excess wing of bulk metallic glass forming liquids, J. Non-Cryst. Solids. 352 (2006) 5103-5109. [10] Z. Wang, H.B. Yu, P. Wen, H.Y. Bai, W.H. Wang, Pronounced slow β-relaxation in La-based bulk metallic glasses, J. Phys.: Condens. Matter.23 (2011) 142202. [11] D.D. Liang, X.D. Wang, Y. Ma, K. Ge, Q.P. Cao, J.Z. Jiang. Decoupling of pronounced beta and alpha relaxations and related mechanical property change, J. Alloy. Compd. 577 (2013) 257-260. [12] X.D. Wang, B. Ruta, L.H. Xiong, D.W. Zhang, Y. Chushkin, H.W. Sheng, H.B. Lou, Q.P. Cao, J.Z. Jiang,Free-volume dependent atomic dynamics in beta relaxation pronounced 9

ACCEPTED MANUSCRIPT La-based metallic glasses, Acta Mater. 99 (2015) 290-296. [13] B. Yang, C.T. Liu, T.G. Nieh, Unified equation for the strength of bulk metallic glasses, Appl. Phys. Lett. 88 (2006) 221911. [14] Y. Li,Bulk metallic glasses: Eutectic coupled zone and amorphous formation, JOM 57

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(2005) 60-63. [15] B. Ruta, Y. Chushkin, G. Monaco, L. Cipelletti, E. Pineda, P. Bruna, V.M. Giordano, M. Gonzalez-Silveira, Atomic-scale relaxation dynamics and aging in a metallic glass probed by x-ray photon correlation spectroscopy, Phys. Rev. Lett. 109 (2012) 165701.

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[16] H. B. Yu, W.H. Wang, K. Samwer, The β relaxation in metallic glasses: an overview, Mater. Today 16 (2013) 183-191.

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[17] Z. P. Lu, X. Hu, Y. Li, S.C. Ng, Glass forming ability of La–Al–Ni–Cu and Pd–Si–Cu bulk metallic glasses, Mater. Sci. Eng. A304 (2001) 679-682.

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[20] J.C. Qiao, J.M. Pelletier, C. Esnouf, Y. Liu, H. Kato, Impact of the structural state on the mechanical properties in a Zr–Co–Al bulk metallic glass, J. Alloy. Compd. 607 (2014) 139-149.

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[21] B.J. Park, H.J. Chang, D.H. Kim, W.T. Kim, In situ formation of two amorphous phases by liquid phase separation in Y–Ti–Al–Co alloy, Appl. Phys. Lett. 85 (2004) 6353. [22] G. Kumar, D. Nagahama, M. Ohnuma, T. Ohkubo, K. Hono, Structural evolution in the supercooled liquid of Zr36Ti24Be40 metallic glass, Scripta Mater. 54 (2006) 801-805. [23] X.M. Huang, X.D. Wang, Y. He, Q.P. Cao, J.Z. Jiang, Are there two glass transitions in Fe–M–Y–B (M = Mo, W, Nb) bulk metallic glasses? Scripta Mater. 60 (2009) 152-155.

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ACCEPTED MANUSCRIPT Table caption Tab. 1 Characteristic temperatures, including the glass transition temperature Tg, the onset crystallization temperature Tx, the melting point Tm, the liquidus temperature Tl, and reduced glass transition temperature Trg, for the as-spun La-Al-Ni/Cu glasses, together with Tga and Txa

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for the annealed glasses, ∆Tga the difference between Tga and Tg, ∆Txa the difference between Txa and Tx, the average atomic distance Rav and Rava for the as-spun and annealed glasses estimated from the principle peak positions of 2θ in XRD patterns and their difference ratio a

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∆R/R = (Rav -Rav)/Rav*100%.

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ACCEPTED MANUSCRIPT Figure captions Fig. 1(a) Normalized storage modulus E’/E0’ and (b) loss modulus E”/E0” vs. normalized temperature T/T0. (c) Normalized loss modulus E”/E0” changing with normalized temperature (T/Tg) for as-spun La85-xAl15Nix (x = 10 to 30 at.%) glasses, where E0’, E0”and T0 are the initial storage modulus, loss modulus and temperature at aheating rate of 3 K/min under a

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frequency of 1 Hz.

Fig. 2 Normalized storage modulus E’/E0’ and loss modulus E”/E0” vs. T/T0 for (a) (b)

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as-spun La70AlxNi30-x(x = 10 to 25 at.%) glasses and (c) (d) La85-xAl15Cux(x = 10 to 20 at.%)

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glasses, respectively.

Fig. 3 XRD patterns for as-spun and annealed (a) La85-xAl15Nix (x = 10 to 30 at.%), (b) La70AlxNi30-x (x = 10 to 25 at.%) and (c) La85-xAl15Cux (x = 10 to 20 at.%) glasses.

Fig. 4 DSC traces for as-spun and annealed (a) La85-xAl15Nix (x = 10 to 30 at.%), (b)

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La70AlxNi30-x (x = 10 to 25 at.%) and (c) La85-xAl15Cux (x = 10 to 20 at.%) glasses.

Fig. 5 Normalized storage modulus E’/E0’ and loss modulus E”/E0” vs. T/T0 for as-spun and annealed (a) La65Al15Ni20, (b) La70Al15Ni15, (c) La75Al15Ni10, (d) La65Al15Cu20, (e)

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La70Al15Cu15 and (f) La75Al15Cu10 glasses, respectively.

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Fig. 6 (a) Normalized loss modulus E”/E0” vs. T/T0 curves for as-spun and preheated-to-450 K La75Al15Ni10 glass, (b) DSC trances for as-spun and preheated-to-473 K La75Al15Ni10 glass, the inset is the local magnification of the DSC curve, showing the increased the Tg, Tx1 and Tx2 after preheating, (c) XRD patterns for as-spun, preheated-to-450 K and preheated-to-473 K La75Al15Ni10 glasses.

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Tg (K)

Tx (K)

Tm (K)

Tl (K)

Trg (Tg/Tl)

Rav (Å)

Tga (K)

Txa (K)

∆Tga (K)

∆Txa (K)

Rava (Å)

∆R/R (%)

La55Al15Ni30 La60Al15Ni25 La65Al15Ni20 La70Al15Ni15 La75Al15Ni10 La70Al10Ni20 La70Al20Ni10 La70Al25Ni5 La65Al15Cu20 La70Al15Cu15 La75Al15Cu10

475 451 437 429 422 416 439 465 400 395 399

504 466 459 444 435 436 480 492 436 429 419

739 692 706 700 702 701 703 693 679 679 677

752 750 721 724 725 745 725 724 692 691 688

0.632 0.601 0.606 0.593 0.582 0.558 0.606 0.642 0.578 0.572 0.579

2.838 2.888 2.920 2.944 2.979 2.953 2.944 2.936 2.896 2.916 2.929

485 461 447 440 432 419 445 470 401 397 403

505 466 460 445 438 432 483 492 436 417 416

10 10 9 11 10 3 4 5 1 2 4

1 0 1 1 3 -4 3 0 0 -12 -3

2.807 2.855 2.886 2.902 2.937 2.919 2.903 2.897 2.865 2.876 2.895

-1.071 -1.122 -1.166 -1.459 -1.413 -1.149 -1.364 -1.329 -1.064 -1.382 -1.17

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Alloys (at.%)

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(a) 1

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3 1 1.4

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T/T0

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2

5

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E'/E'0

5 1. La55Al15Ni30 2. La60Al15Ni25 3. La65Al15Ni20 4. La70Al15Ni15 5. La75Al15Ni10

E"/E"0

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1 1.0

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E'/E'0

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E'/E'0

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1. La65Al15Cu20 2. La70Al15Cu15 3. La75Al15Cu10

3. La70Al20Ni10 4. La70Al25Ni5 0.2

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1.3

3

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3

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Temperature (K)

(a) Intensity (a.u)

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La55Al15Ni30 La60Al15Ni25 La65Al15Ni20

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La70Al15Ni15

La75Al15Ni10 20

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(b) Intensity (a.u)

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La70Al10Ni20 La70Al15Ni15 La70Al20Ni10 La70Al25Ni5

20

30

40

50

60

70

2 (degree)

(c) Intensity (a.u)

As-spun Annealed

La65Al15Cu20 La70Al15Cu15 La75Al15Cu10 20

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70

(a)

As-spun Annealed

Heat flow (a.u.)

La55Al15Ni30 La60Al15Ni25 La65Al15Ni20

TE D

La70Al15Ni15 La75Al15Ni10

350

400

450

500

550

600

650

Temperature (K)

(b)

EP

Heat flow (a.u.)

As-spun Annealed

La70Al10Ni20 La70Al15Ni15

AC C

La70Al20Ni10

La70Al25Ni5

350

400

450

500

550

600

650

Temperature (K)

(c) Heat flow (a.u.)

As-spun Annealed

La65Al15Cu20

La70Al15Cu15

La75Al15Cu10

350

400

450

500

550

Temperature (K)

600

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

650

(a)

(d)

As-spun Annealed

10

10

TE D 1

1 1.1

1.2

1.3

1.4

1.5

1.0

(b)

(e)

AC C

As-spun Annealed

E''/E''0

La70Al15Ni15

1.1

1.2

1.3

1.4

1.5

1 1.0

1.6

1.1

1.2

La75Al15Cu10

E'/E'0

1 1.0

As-spun Annealed

10

E''/E''0

10

1.3

T/T0

(f)

As-spun Annealed La75Al15Ni10

1.3

As-spun Annealed La70Al15Cu15

T/T0

(c)

1.2

T/T0

10

1 1.0

1.1

EP

T/T0

E'/E'0

1.0

10

As-spun Annealed

La65Al15Cu20

E''/E''0

E''/E''0

La65Al15Ni20

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

1

1.1

1.2

1.3

1.4

T/T0

1.5

1.6

1.0

1.1

1.2

1.3

T/T0

1.4

1.5

(a) 10

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

485 K 430 K 456 K

E''/E''0

La75Al15Ni10

450 K As-spun

1

3 K/min

350

450

500

Temperature (K) La75Al15Ni10

As-spun

473K

Tx2

AC C

Tx1

EP

Heat flow (a.u.)

(b)

400

TE D

Preheated to 450 K

Preheated to 473K Tx1 Tg

Tx2

Tx2

20 K/min

420 440 460 480 500 520

400

420

440

460

480

500

520

540

Temperature (K)

(c)

La75Al15Ni10

La (P63/mcc)

Intensity (a.u)

As-spun preheated to 450 K preheated to 473K

20

30

40

50

2theta (degree)

60

ACCEPTED MANUSCRIPT Highlights:




La-Al-Ni metallic glasses have more pronounced beta-relaxation after annealing.




A weak beta-relaxation in La-Al-Cu metallic glasses almost disappears after




RI PT

annealing. La-Al-Ni glasses have higher Tg temperatures and bigger changes in Tg by sub-Tg annealing.

The pronounced beta-relaxation more depends on the composition than the free

EP

TE D

M AN U

SC

volume.

AC C