Crystallization of melt-spun Mg63Ni22Pr15 amorphous alloy ribbon

Journal of Alloys and Compounds 441 (2007) 76–80

Crystallization of melt-spun Mg63Ni22Pr15 amorphous alloy ribbon Feng Xu a,∗ , Yulei Du a , Ping Gao a , Zhida Han b , Guang Chen a , Shuqin Wang c , Jianzhong Jiang d a

Department of Materials Science and Engineering, Nanjing University of Science and Technology, Nanjing 210094, China b Physics Department, Nanjing University, Nanjing 210093, China c Analysis Center, School of Chemical Engineering, Nanjing University of Science and Technology, Nanjing 210094, China d Department of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, China Received 18 August 2006; received in revised form 18 September 2006; accepted 18 September 2006 Available online 17 October 2006

Abstract The crystallization of melt-spun Mg63 Ni22 Pr15 amorphous ribbons has been studied. The isochronal differential scanning calorimeter (DSC) measurement shows that the alloy has a certain glass-forming ability (GFA) according to different criterions. The analysis of with Kissinger equation indicates that the activation energy for the primary crystallization is relatively low. Nanocrystalline Mg12 Pr and Mg2 Ni phases appeared after the primary crystallization. Discussion based on Johnson–Mehl–Avrami (JMA) equation indicates that the isothermal process of primary crystallization starts from small dimensions with an increasing nucleation rate. The further fitting and discussion with some simplified models infers that the increasing nucleation rate is caused by the large difference between diffusivity of Mg atom and that of Ni and Pr atoms. © 2006 Elsevier B.V. All rights reserved. Keywords: Metallic glasses; Isothermal crystallization; Thermal analysis

1. Introduction Since Mg-based bulk metallic glass (BMG) was reported [1], it has attracted lots of attention due to its high specific strength. With the development of bulk metallic glasses (BMGs) research field, new kinds of Mg-based BMGs, Mg–Ln–TM (Ln: lanthanide; TM: transitional metal), have been studied [1–12]. Among them, most of studies were focused on Mg–Cu–RE (RE: rare earth metals) systems [2–9], such as Mg–Cu–Nd [2], Mg–Cu–Y [3–7], Mg–Cu–Gd [8] and Mg–Cu–Er [9]. Since Ni and Cu are close to each other in the element periodic table, Mg–Ni–RE has also been investigated, such as Mg–Ni–Ce [1], Mg–Ni–Nd [10–11] and Mg–Ni–Y [12]. Recently, a new composition of Mg-based BMG, Mg–Ni–Pr alloy, has been reported to have unusual oxygen resistance during the preparation process, and to be a very potential candidate ∗ Corresponding author. Present address: Department of Materials Science and Engineering, National Tsing Hua University, Hsinchu 30013, Taiwan. Tel.: +86 25 84315797/886 3 5715131x35343; fax: +86 25 84315797/886 3 5722366. E-mail addresses: xufeng [email protected], [email protected] (F. Xu).

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for a hydrogen-storage material [13]. Although some investigation on the formation of this alloy has been made, the crystallization behavior needs further study. 2. Experimental Ni22 Pr15 master alloy was prepared by induction melting Ni (99.8%) and Pr (99.9%) under an argon atmosphere in a quartz tube. The ingot was remelted at least three times to guarantee the chemical homogeneity. In order to prevent the evaporation of Mg, the master alloy was then alloyed with Mg (99.99%) step by step. After the completion of the master ingot preparation, the weight loss was found negligible. Mg63 Ni22 Pr15 alloy ribbons, with a cross section of 1.5 mm × 0.022 mm, were prepared by melt-spinning in an argon atmosphere onto a copper wheel with a tangential spinning speed of 40 m/s. The ribbon samples were annealed after sealed into a vacuum quartz tube. The sample structures were characterized by an X-ray diffractometer (XRD) using Cu K␣ radiation. Thermal analyses were performed in a Shimadzu 50 power compensation differential scanning calorimeter (DSC) under a flow of purified argon (99.998%).

3. Results and discussion The XRD patterns of the as-quenched and annealed samples are presented in Fig. 1. The results show that the structure of the as-quenched sample is typical amorphous. Fig. 2 presents

F. Xu et al. / Journal of Alloys and Compounds 441 (2007) 76–80

77

Fig. 3. Kissinger analysis with the calculated activation energies of three crystallization transformations.

Fig. 1. XRD patterns of Mg63 Ni22 Pr15 alloy.

Fig. 2. The isochronous DSC measured at different heating rates for Kissinger analysis.

the DSC traces measured at different heating rates from 10 to 30 K/min. Each trace exhibits an endothermic event characteristic of the glass transition and a distinct supercooled liquid region, followed by three exothermic events characteristic of crystallization processes. It shows that the Mg63 Ni22 Pr15 alloy prepared by melt-spinning method is a typical BMG candidate as reported [13]. Same as in the BMG cases, Mg63 Ni22 Pr15 alloy ribbon experiences two separate melting stages during the DSC measurements. It indicates that the composition is off-eutectic. All the transition temperatures obtained at different heating rates, including glass transition temperature (Tg ), onset of the first crystallization peak (Tx ), three crystallization peaks (Tp1 , Tp2 , T), onset of melting temperature (Tm ) and liquidus temperature (Tl ) were listed into Table 1. Several thermal parameters at different heating rates are also concluded. The supercooled liquid region T (defined as Tx − Tg ) and the reduced glass transi-

tion temperature Trg (defined as Tg /Tm ) and Trg (defined as Tg /Tl ) are the most common used parameters to reflect the relative glass-forming ability (GFA). Recently, a new criterion defined as γ = Tx /(Tg + Tl ) was raised as an important supplement of parameters [14]. With increasing heating rate, the transition temperatures are shifted to higher temperatures. This phenomenon is quite common, and the physical reason has been discussed in Ref. [14]. Although the transition temperatures change, Trg , Trg and γ keep quite stable and the increases of their values are very limited, within 2%. That shows here both Trg (or Trg ) and γ can be the criterions of the glass formation of the alloy. And the values indicate that the alloy has a certain glass forming ability according to Refs. [14,15], respectively. The relation between the peak temperatures and the heating rate can usually be analyzed by Kissinger equation [16]:   β −Ex (1) = νx exp RTp (Tp )2 where β is the heating rate, νx a pre-exponential factor, Ex the activation energy for crystallization, R the gas constant and Tp is the temperature of maximum energy release of each peak in the DSC scan. And this equation can be converted to the expression:   β Ex ln =C− (2) Tp2 RTp where C is a constant (=ln νx ). The activation energy of each peak obtained through ln(β/Tp2 ) 1/Tp relation is 1.78 ± 0.03, 1.53 ± 0.02, and 2.26 ± 0.02 eV, respectively, as shown in Fig. 3. The activation energy of the first crystallization is close to that in the (Mg1−x Alx )60 Cu30 Y10 (x = 0.02 and 0.04) [4]. After first and second crystallization, the structures of the samples were examined by XRD, as shown in Fig. 1. One sam-

Table 1 Thermal parameters obtained from Fig. 2 at different heating rates Heating rate (K/min)

Tg (K)

Tx (K)

Tp1 (K)

Tp2 (K)

Tp3 (K)

T (K)

Tm (K)

Tl (K)

Trg

Trg

γ

10 20 30

449 452 456

489 495 502

497 505 510

536 547 555

640 650 656

40 43 46

740 741 741

758 763 766

0.59 0.59 0.60

0.61 0.61 0.615

0.405 0.407 0.411

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F. Xu et al. / Journal of Alloys and Compounds 441 (2007) 76–80

Fig. 4. The isothermal DSC traces of the primary crystallization of Mg63 Ni22 Pr15 ribbons at different temperatures.

ple was first heated from room temperature to 491 K at 20 K/min, and then isothermally annealed for 30 min, enough for the primary crystallization. The other sample was first heated from room temperature to 543 K at 20 K/min, and then isothermally annealed for 30 min. After the annealing, Mg2 Ni and Mg12 Pr phases can be detected from both patterns, similar as the cases in other Mg–Ni–RE alloy systems [12,17]. Some unknown phases appear after the second crystallization. For the primary crystallization, the crystallite sizes can be estimated from the broadening of the peaks by Scherrer equation. The estimated sizes of Mg12 Pr and Mg2 Ni crystals for 30 min annealing are about 14 and 22 nm, respectively. For the shorter time of the following isothermal annealing processes, the crystal sizes could be even smaller. In order to know more about the primary crystallization of Mg63 Ni22 Pr15 alloy ribbon, isothermal DSC measurements were performed between 481 and 491 K in the supercooled liquid region. The DSC traces of the isothermal treatment shown in Fig. 4 exhibit a single exothermic peak after a certain incubation period. The traces of this alloy are in high symmetrical forms and have relatively long incubation times. Here, we make an assumption which is often used in the crystallization kinetics of melt-glass systems [18] that the heat release measured in Fig. 4 is proportional to the relative crystallized volume, and then the relative volume fraction of crystallized phase as a function of time could be deduced, as shown in Fig. 5. The shape of the curves is typical “S” type. In many cases, the time evolution of the crystallization transformation fraction is often described by a phenomenological model known as the Johnson–Mehl–Avrami (JMA) model [19,20]. The essence of the model can be written as a very simple formula commonly referred to as the JMA equation: x(t) = 1 − exp{−[k(t − τ)]n }

Fig. 5. Relative volume transformation fractions of the primary crystallization as a function of time at different isothermal annealing temperatures with JMA fitting curves. Table 2 Parameters obtained from the JMA model, and model 10 fitting at three isothermal temperatures for the Mg63 Ni22 Pr15 ribbon alloys Parameters

481 K

486 K

491 K

From Fig. 5 k (min−1 ) τ (min) n

0.3428 ± 0.0003 6.10 ± 0.02 2.611 ± 0.004

0.5622 ± 0.0003 4.20 ± 0.02 2.613 ± 0.005

0.8637 ± 0.0004 2.85 ± 0.01 2.614 ± 0.004

From Fig. 6 k (min−1 ) n

0.342 ± 0.002 2.622 ± 0.005

0.559 ± 0.002 2.626 ± 0.004

0.862 ± 0.003 2.615 ± 0.007

3470 ± 5 46.2 ± 0.1

7140 ± 7 32.1 ± 0.1

From model 10 Ist V0 (min−1 ) τ (min)

914 ± 2 61.9 ± 0.1

fitting are listed in Table 2. Both k and n are good experimental parameters for kinetic studies and can also be estimated from the intercept and slope of the following equation: ln[−ln(1 − x)] = n ln k + n ln(t − τ)

(4)

Fig. 6 shows the plots of ln[−ln(1 − x)] versus ln(t − τ) at different annealing temperatures for the data of x = 20–80%, in which the values of τ obtained from Eq. (3) are used. The Avrami exponent n and the reaction rate constants k deduced from the slope and the intercept of the lines and the results are

(3)

where x(t) is the crystallized volume fraction, t the annealing time, τ the incubation time, n the Avrami exponent related to the dimensionality of nucleation and growth, and k a kinetic constant of the process which depends on temperature and effective activation energy Ea by k = k0 exp(−Ea /RT), where k0 is a constant and R is the gas constant. Eq. (3) fits the experiment results well, as shown in Fig. 5 by solid curves, for annealing time longer than τ. The values of k, τ and n obtained from the

Fig. 6. JMA plots, ln[−ln(1 − x)] vs. ln(t − τ), in which the data for 0.2 < x < 0.8 are used.

F. Xu et al. / Journal of Alloys and Compounds 441 (2007) 76–80

also listed in Table 2. The effective activation energy for the primary crystallization is estimated to be 182 ± 5 kJ/mol (1.88 eV), deduced from the plots of ln k versus 1/T. The value is close to that obtained from Kissinger model of the first crystallization. Compared with the activation energy of crystallization in ZrTiCuNiBe alloy (409 kJ/mol) [21] or CuZrTi alloy (533 kJ/mol) [22], the activation energy of the primary crystallization of Mg63 Ni22 Pr15 alloy is lower and it is determined by the low thermodynamic potential energy barrier of nucleation and diffusion activation energy, which has been indicated by the low transformation temperatures. Usually in most metallic glass alloys, the Avrami exponent changes with annealing temperatures [22,23]. However, in this Mg63 Ni22 Pr15 alloy ribbon, the Avrami exponent is found to be almost constant, n ≈ 2.6, for all three isothermal annealing temperatures. This similar phenomenon was ever observed in Mg–Ni–Y BMG alloy [12]. It indicates that the crystallization mechanism in the alloy is not very sensitive to the annealing temperature within the temperature region. According to the JMA kinetic model, the Avrami exponent, n > 2.5, means that crystallization starts from small dimensions with a increasing nucleation rate. Crystallization of a metallic glass is normally regarded as a process proceeding by nucleation and subsequent growth of crystals. In order to make a clearer description of the isothermal crystallization process, several simple models which consider about a transient nucleation process (time-dependent nucleation) for primary crystallization within the incubation time were applied to fit the data. In the past study of the crystallization of metallic glasses, a time-dependent nucleation model which is based on the Zeldovich equation [24] was applied [21,22], where the time-dependent nucleation rate, I(t) is written as  I(t) = Ist exp

−τ t

 (5)

where τ is the transient nucleation time and Ist is the steadystate nucleation rate. The relative crystallized volume fraction x(t) during the crystallization process in metallic glasses as a function of time can be expressed as x(t) = 1 − exp(−Y (t))

(6)

where the expression of Y(t) depends on the nucleation and growth models used. For Mg63 Ni22 Pr15 metallic glasses, the crystal sizes of Mg2 Ni and Mg12 Pr phases, formed through isothermal annealing in primary crystallization, are very small, as confirmed by XRD. The nanostructure of this alloy is determined by the large difference between the diffusivities of different atoms (the diffusivity of Mg is much larger than that of Ni and Pr atoms in the alloy). The high diffusivity of Mg leads to very fast saturation of crystal growth from nuclei, while the further growth is confined by the low diffusivity of Pr atoms. For the primary crystallization of Mg–Ni–RE metallic glasses, the development of the nanocrystals’ size with the time of isothermal annealing has been observed to follow a parabolic relationship, which can be

79

Fig. 7. Relative volume transformation fractions of the primary crystallization as a function of the annealing time at different isothermal temperatures with model 10 fitting curves.

expressed as [25]: √ d = 2α Dt

(7)

where d is the grain size, D the chemical diffusion coefficient and α is a constant factor. Since the typical D value for amorphous alloys is relatively high [26], this relationship indicated a rapid crystal growth from zero to d. Therefore, it can be inferred that the crystallization is a nucleation-controlled process. In this case, we have made an assumption that grain size is constant. Among the models listed in Ref. [21], the following models are chosen to fit the relation curves of the relative volume fraction data versus time. They are [21]: • Model 2: Quenched-in nucleation with constant grain size V0 . • Model 4: Steady-state nucleation with constant grain size. • Model 6: Quenched-in and steady-state nucleation with constant grain size. • Model 10: Time-dependent Z-model with constant grain size:  t  t τ  I(x)V0 dx = Ist V0 exp dx (8) Y (t) = x 0 0 • Model 14: Quenched-in and time-dependent Z-model with constant grain size:  t I(x)V0 dx Y (t) = NV0 + 0



= NV 0 + Ist V0



t

exp 0

−τ x

 dx

(9)

We found that models 2, 4 and 6 cannot fit the data at all, while models 10 and 14 can fit the data quite well. The fitting results from model 14 are extremely close to those from model 10, except for the negligible contribution from the quenched-in nucleation. The initial crystallization stage of “S” curve can be well fitted, as illustrated in the inset of Fig. 7. The parameters from model 10 were listed in Table 2. The good fitting indicates that the crystallization process is dominated by time-dependent nucleation. There is no detectable quenched-in nucleation found through fitting. This could be different from the BMG case where the quenched-in nucleation

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F. Xu et al. / Journal of Alloys and Compounds 441 (2007) 76–80

is usually unavoidable. The incubation times obtained through model 10 are much longer than the actual ones, and even longer than the measurement times. It indicates that the nucleation rate increases during the whole the measurement protocol. And this finding is in good accordance with the Avrami exponent (n ≈ 2.6) obtained from JMA model. With the results deduced from Ist V0 of model 10, the nucleation rate of either crystalline phase for 50% crystallization is estimated to be in the order of 1 × 1021 m−3 s−1 . Although it is quite large compared with those calculated in other metallic glasses [21,22], it is still a reasonable value for primary crystallization processes [27]. However, although the exothermic behavior of primary crystallization has been satisfactorily fitted, the increasing nucleation mechanism needs further consideration. As discussed before, for the primary crystallization of both Mg12 Pr and Mg2 Ni, the crystal growth is confined within nanosize by the low diffusivity of Pr and Ni atoms. However the high diffusivity of Mg atom makes the nucleation easy. Furthermore, because the composition of the alloy is off-eutectic, once the primary nucleation starts, the composition of the amorphous residue will become farther away from the eutectic composition. In this case, the earlier nuclei or nanocrystals will induce further nucleation, and therefore increase the nucleation rate. High nucleation rate dominates the crystallization mechanism and also reduces the activation energy. Although the energy barrier for growth is still high, the activation energy contributed from both nucleation and growth is relatively low, because of the easy nucleation. Although model 10 is a good fitting of the “S” curve, there are definitely some limitations in model 10 because of the simplified precondition. For example, the fitting cannot well explain the nucleation and growth process when crystallization inclines to saturation. When the crystallization fraction is very high and gets close to saturation, the nucleation rate will get lower. As seen from the isochronal measurement, the exothermic enthalpy from the primary crystallization takes up 60% of the whole crystallization enthalpy. Such a high percentage indicates the overlapping of the diffusion fields of the growing nanocrystals, which will suppress further nucleation and growth [12]. 4. Conclusion Although the alloy is off-eutectic, it has a certain glassforming ability according to different criterions. However, the activation energy for the primary crystallization calculated from either isochronal DSC measurement or isothermal DSC measurement is relatively lower than that of some other BMG alloys. The Avrami exponent indicates that the isothermal process of

primary crystallization starts from small dimensions with an increasing nucleation rate, which is further supported by the fitting with model 10. Discussion infers that the increasing nucleation rate is caused by the large difference between diffusivity of Mg atom and that of Ni and Pr atoms. Acknowledgements We would like to acknowledge Prof. Youwei Du in Physics Department of Nanjing University for giving the convenience of sample preparation, and acknowledge Dr. Qingping Cao in Zhejiang University for discussion. References [1] A. Inoue, K. Ohtera, K. Kita, T. Masumoto, Jpn. J. Appl. Phys. 27 (1988) L2248. [2] L.J. Huang, G.Y. Liang, Z.B. Sun, Y.F. Zhou, J. Alloys Compd. 432 (2007) 172. [3] U. Wolff, N. Pryds, J.A. Wert, Scripta Mater. 50 (2004) 1385. [4] S. Linderoth, N.H. Pryds, M. Ohnuma, A.S. Pedersen, M. Eldrup, N. Nishiyama, A. Inoue, Mater. Sci. Eng. A 304–306 (2001) 656. [5] S.V. Madge, A.L. Greer, Mater. Sci. Eng. A 375–377 (2004) 759. [6] H. Men, Z.Q. Hu, J. Xu, Scripta Mater. 46 (2002) 699. [7] D. Kim, B.J. Lee, N.J. Kim, Scripta Mater. 52 (2005) 969. [8] G.Y. Yuan, A. Inoue, J. Alloys Compd. 387 (2005) 134. [9] W.Y. Liu, H.F. Zhang, Z.Q. Hu, H. Wang, J. Alloys Compd. 397 (2005) 202. [10] S.V. Madge, D.T.L. Alexander, A.L. Greer, J. Non-Cryst. Solids 317 (2003) 23. [11] Z.P. Lu, C.T. Liu, Y. Li, Intermetallics 12 (2004) 869. [12] V. Rangelova, T. Spassov, J. Alloys Compd. 345 (2002) 148. [13] Y.X. Wei, X.K. Xi, D.Q. Zhao, M.X. Pan, W.H. Pan, Mater. Lett. 59 (2005) 945. [14] Z.P. Lu, C.T. Liu, Phys. Rev. Lett. 91 (2003) 115505. [15] A. Inoue, Acta Mater. 48 (2000) 279. [16] H.E. Kissinger, J. Res. Nat. Bur. Stand. 57 (1956) 217. [17] T. Spassov, V. Rangelova, N. Neykov, J. Alloys Compd. 334 (2002) 219. [18] U. Wolff, N. Pryds, E. Johnson, J.A. Wert, Acta Mater. 52 (2004) 1989. [19] M.A. Johnson, R.F. Mehl, Trans. Am. Inst. Min. Metall. Pet. Eng. 135 (1939) 416. [20] M. Avrami, J. Chem. Phys. 9 (1941) 177. [21] J.Z. Jiang, Y.X. Zhuang, H. Rasmussen, Phys. Rev. B 64 (2001) 094208. [22] Q.P. Cao, Y.H. Zhou, A. Horsewell, J.Z. Jiang, J. Phys.: Condens. Matter 15 (2003) 8703. [23] F. Xu, J.Z. Jiang, Q.P. Cao, Y.W. Du, J. Alloys Compd. 392 (2005) 173. [24] J.B. Zeldovich, Acta Physicochim., URSS 18 (1943) 1. [25] Z.P. Lu, C.T. Liu, C.H. Kam, Y. Li, Appl. Phys. Lett. 82 (2003) 862. [26] A.L. Greer, in: H.H. Lieberman (Ed.), Rapidly Solidified Alloys: Processes, Structures, Properties and Applications, Marcel Dekker, New York, 1993, p. 269. [27] J.C. Foley, D.R. Allen, J.H. Perepezko, Scripta Mater. 35 (1996) 655.