The peculiarities of the nanocrystallization process and copper addition effects for an amorphous Fe75.5Si13.5B9Cu2 alloy

Journal of Alloys and Compounds 475 (2009) 706–711

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The peculiarities of the nanocrystallization process and copper addition effects for an amorphous Fe75.5 Si13.5 B9 Cu2 alloy H.T. Zhou a,∗ , Z.K. Zhao a , X. Zhou b , B. Yan c , J.W. Zhong a , Q.B. Li a a

School of Materials Science and Engineering, Central South University, Changsha 410083, PR China School of Materials and Chemical Engineering, Jiangxi University of Science and Technology, Ganzhou 341000, PR China c Shanghai Key Lab of D&A for Metal-Functional Materials, Tongji University, Shanghai 200092, PR China b

a r t i c l e

i n f o

Article history: Received 2 December 2007 Received in revised form 28 July 2008 Accepted 30 July 2008 Available online 18 September 2008 Keywords: Amorphous alloy Cu addition Phase transformation Thermal property Grain evolution

a b s t r a c t The crystallization process and copper addition effects of an amorphous Fe75.5 Si13.5 B9 Cu2 alloy were investigated by DSC, DIL, XRD and TEM. The experimental results show that copper addition decreases the activation energy of primary crystallization as well as the structure relaxation temperature, and thus causes Fe–Si–B metallic glasses less stable. During annealing process of Fe75.5 Si13.5 B9 Cu2 alloy, it is found that metastable Fe3 B phase precipitates which is confirmed by both XRD results directly and thermal expansion coefficient curves at the heating rates of higher than 10 K/min indirectly. Moreover, equiaxed nanometer-sized grains of Fe–Si are promoted without forming coarse dentritic morphology due to copper addition of 2 at.%. By annealing isothermally at 773 K for 60 min, homogeneous Fe–Si nanograins (about 25 nm) with more than 65% volume fraction are obtained in favor of good soft magnetic properties. © 2008 Elsevier B.V. All rights reserved.

1. Introduction There are considerable interests in amorphous Fe–Si–B-based alloys for many years due to their excellent soft magnetic properties and commercial application prospects. Some investigations have been carried out on crystallization kinetics, magnetic properties, electrochemical properties and microstructure [1–5]. However, most researches are concentrated on the composition of 5–10 at.% Si. When Fe–Si–B-based alloys are with high Si content and added with some elements (e.g. Cu, Nb, Ta, Cr, V and Ni), metallic glasses can be formed easily. For instance, Yoshizawa developed commercial metallic glass Finemet with the nominal composition Fe73.5 Si13.5 B9 Cu1 Nb3 with little nanocrystallites exhibiting good soft magnetic properties [6]. In this alloy system, the microstructure and soft magnetic properties are dependent much on the composition and annealing processing. Note that copper plays a key role in forming ultrafine grain structure for excellent soft magnetic properties, though there are still conflicting elucidations regarding the crucial role of copper. Ayers suggested that Cu clusters were enveloped by Fe–Si nanocrystals [7]. In contrast, Hono assumed that Cu clusters were formed

∗ Corresponding author. Tel.: +86 731 8830257. E-mail address: [email protected] (H.T. Zhou). 0925-8388/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2008.07.125

in the as-quenched state and served as nucleation sites [8]. During annealing processing, a series of changes in these high silicon content Fe–Si–B alloys would occur, especially the crystallization behavior. For example, the Fe75−x Si9 B16 Cux alloy undergoes three stages during annealing [9]: (a) dendritic growth of bcc ␣-Fe(Si, Cu), (b) eutectic crystallization of bcc ␣-Fe(Si, Cu) and bct Fe3 B, (c) precipitation of bct Fe3 B phase. Afterwards, the metastable compound Fe3 B dissolves into ␣-Fe and Fe2 B during the second and third stages. Besides, Cu addition can accelerate the crystallization and the constituent of 2 at.% Cu results in the minimum annealing time (less than 1 h) to obtain the maximum saturation magnetization [1]. Moreover, the composite microstructure of primary crystals Fe–Si with a rather small size homogeneously distributed in the amorphous matrix can be obtained by isothermally annealing, which makes great contribution to good soft magnetic properties [10]. In this study, we try to focus on the Fe–Si–B alloy with high Si content as well as Cu addition and analyze the thermodynamics and thermal expansion changes. Because metallic glasses can relax always companying with a characteristic of volume contraction by annealing at a high temperature due to the existence of large free volume during rapid solidification [11], this provides a new method to monitor the crystallization behavior of amorphous Fe–Si–Bbased alloys. And this research will provide a valuable reference to understand copper addition effects on amorphous Fe–Si–B-based alloys.

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2. Experimental procedure The amorphous Fe75.5 Si13.5 B9 Cu2 (at.%) ribbons received were 25 ± 1 ␮m thick and 55.0 ± 0.1 mm wide prepared by single-wheel melt spinning processing in Shanghai key lab of D&A for Metal-Functional materials. As-quenched samples were analyzed by Simultaneous DSC–TGA Q600 in argon atmosphere with different heating rates of 5, 10, 15, 20 and 25 K/min, respectively. The dilatometric (DIL) measurements of as-quenched ribbons were performed using Netzsch DIL 402C dilatometer in argon atmosphere with heating rates of 5, 10 and 20 K/min, respectively. Before these measurements, the 200-mm long ribbons were coiled into rolls of Ø6.70 ± 0.01 mm × 22.0 ± 0.1 mm. Rigaku D/Max 2550 with Cu K␣ radiation was used for X-ray diffraction (XRD) analysis. According to XRD patterns, phase identification and grain-size calculation by Scherrer’s formula were performed. Annealing processing was in vacuum (0.1 Pa) using a CVD (G)-06/50/2 electronic high-temperature tube furnace (argon gas protection) with controllable heating rates. Annealing temperatures were selected from 773 to 953 K according to DSC results. The samples were heated rapidly at 10 K/min and then annealed isothermally for 5–60 min at the selected temperatures. When annealing was completed, the samples were taken out and cooled quickly in the air. Hitachi H-800 transmission electron microscope (TEM) with an accelerating voltage of 150 kV was utilized for microstructure investigations.

3. Results and discussion 3.1. Phase transformation during crystallization Fig. 1 is the XRD results of samples after isothermal annealing for 1 h from 773 to 953 K and shows the annealing temperature dependence of phase transformation. In the as-quenched samples, only a broad diffraction peak at the position of 2 ≈ 45◦ is observed in Fig. 1(a) and bcc ␣-Fe(Si) precipitation occurs with increasing the temperature below 773 K. The growth of Fe2 B peaks and the decay of Fe3 B peaks in Fig. 1(c), (d) and (f) suggest that Fe2 B forms gradually while metastable Fe3 B phase forms at the beginning of the secondary crystallization and then decomposes into Fe2 B phase consequently, which is similar to XRD analysis of the literatures

Fig. 2. DSC curves at different heating rates.

Table 1 Characteristic temperatures of an amorphous Fe75.5 Si13.5 B9 Cu2 alloy at different heating rates ˇ (K/min)

Tc (K)

Tg (K)

Tx1 (K)

Tx2 (K)

Tx (K)

Tp1 (K)

Tp2 (K)

5 10 15 20 25

688.2 686.0 685.4 684.8 684.2

693.7 713.3 715.9 720.5 742.1

791.5 798.9 804.0 810.0 810.3

812.4 821.8 828.5 832.8 835.1

28.3 30.8 34.9 35.7 39.2

799.4 808.4 814.6 818.0 822.2

818.0 825.4 832.1 836.0 837.4

ˇ: heating rate; Tc : Curie temperature; Tg : glass transition temperature; Tx : the temperature of the whole crystallization process; Tx1 , Tx2 : the onset temperature of the first and second crystallization; Tp1 , Tp2 : peak temperature.

[12,13]. When crystallization finishes, amorphous phase vanishes and residual phases are ␣-Fe(Si) and Fe2 B. 3.2. Thermal properties during crystallization Fig. 2 shows the DSC curves of amorphous Fe75.5 Si13.5 B9 Cu2 alloy at different heating rates, and Table 1 presents the characteristic temperatures. It is found that there are two exothermic peaks indicating two main crystallization stages. When the heating rates are 5, 10 and 20 K/min, the onset temperature of crystallization Tx are 791.5, 798.9 and 810.0 K, respectively. When the heating rate increases, Curie temperature Tc declines due to the decease of structural relaxation temperature. Notably, Tc has a shift tendency which is different from other characteristic temperatures such as Tg and Tp . According to the melting curves as shown in Fig. 3, endothermic peaks exist at 1442.2 and 1451.5 K, respec-

Fig. 1. XRD patterns of an amorphous Fe75.5 Si13.5 B9 Cu2 alloy annealed at different temperatures for 1 h: (a) as-cast; (b) 773 K; (c) 833 K; (d) 873 K; (e) 953 K; (f) 883 K. ␣-Fe(Si) (), Fe3 B () and Fe2 B (↓).

Fig. 3. The melting and cooling curves in DSC measurements.

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Table 2 Characteristic temperatures at the rate of 20 K/min and some criteria for evaluating GFA Fe75.5 Si13.5 B9 Cu2 Tg (K) Tx (K) Te (K) Tm (K) Tl (K) Tg /Tm Trg Txg (K) 

720.5 810.0 1394.1 1419.6 1448.8 0.507 0.497 89.5 0.373

Tg : glass transition temperature; Tx : the onset temperature of the crystallization; Te : eutectic transformation temperature; Tl : liquidus temperature; Trg = Tg /Tl ; Txg = Tx − Tg ;  = Tx /(Tg + Tl ).

tively. Combining with XRD results, these two peaks corresponds to Fe–Si solid solution and Fe2 B phase during melting. Table 2 shows some values which are criteria for evaluating the glass-forming ability (GFA). Compared with the Fe77.5 Si13.5 B9 alloy [14], copper addition weakens the GFA of Fe–Si–B-based alloys. This results from the mixing enthalpies with unwanted large positive values. The mixing enthalpies between Cu and other constituent elements are 13 kJ/mol for the Cu–Fe pair, 0 kJ/mol for the Cu–B pair and −19 kJ/mol for the Cu–Si pair, respectively [15]. For this reason, with Cu addition, the negative value for the overall mixing enthalpy becomes smaller and finally leads to the weakness of GFA. Usually, the crystallization kinetics of amorphous Fe–Si–Bbased alloys is described by JMA equation as follows [16]: n

˛ = 1 − exp[−(kt) ]

(1)

where k is a rate constant as a function of the temperature and depends on the nucleation rate and on the growth speed of the crystallites; n is the Avrami exponent which is related to the growth mechanism; ˛ is the crystallization volume fraction. Taking account of the case of the partial overlap for two exothermic peaks and the precipitation of two different phases shown in the XRD results (Fig. 1), it is difficult to arrest the crystallization process precisely at a predetermined crystallization fraction, or in other words, the two peaks cannot be divided accurately. So the models based on crystallization fraction are unsuitable to characterize the crystallization kinetics of this alloy. As a result, two nonisothermal kinetics models are introduced to calculate the crystallization activation energy. Firstly, Kissinger’s method [17] is used as follows:



ln

Tp2



ˇ

= ln

E  a

R

− ln  +

Ea RTp

(2)

where ˇ is the heating rate; Ea is the activation energy; Tp is the peak temperature and R is gas constant: 8.314 J/mol. Plotting ln(Tp2 /ˇ) as y-axis and 1/Tp as x-axis, a straight fit line can be obtained with the slope of Ea /R. Therefore, the activation energy can be calculated as shown in Fig. 4. Secondly, the crystallization activation energies are also calculated by Doyle–Ozawa’s means [18] with the following equation: lg ˇ = lg

 AE  a RF(˛)

− 2.315 −

0.4567Ea RT

(3)

where ˇ is the heating rate; A is a constant; Ea is the activation energy; F(˛) is the crystallization fraction and R is gas constant: 8.314 J/mol. A straight line can also be obtained by plotting lg ˇ as yaxis and 1/T as x-axis with the slope of −0.4567Ea /R. The activation energy can then be calculated as seen in Fig. 5. Combining with the XRD results, the first and second exothermic peaks in Fig. 2 correspond with the formation of bcc ␣-Fe(Si) crystallites and Fe–B intermetallic compounds. Table 3 shows the

Fig. 4. The relationship between ln(Tp2 /ˇ) and 1/Tp of crystallization for an amorphous Fe75.5 Si13.5 B9 Cu2 alloy: (a) peak 1; (b) peak 2.

activation energies of two crystallization stages. On the basis of calculation with Doyle–Ozawa’s method, the activation energy for the first peak is 366 kJ/mol while that the second one is 426 kJ/mol. Compared with the activation energies of Fe77.5 Si13.5 B9 to be 536 and 430 kJ/mol [12], it is concluded that the crystallization process can start more easily with Cu addition or the metallic glass becomes less stable [19]. At the same time, in comparison with the first crystallization activation energy (225 kJ/mol) of Fe73 Si9 B16 Cu2 alloy [1], the amorphous matrix is stabilized by replacing B with Si in the Fe–Si–B–Cu alloy, which is consistent with the early reports [20,21]. Obviously, the primary crystallization has a lower value of activation energy than that of the secondary crystallization. The reason is that Cu atoms are slightly soluble in ␣-Fe and cannot form borides [9], they favor microsegregation and concentrate into clusters with few nanometers in size, and then act as heterogeneous nucleation sites for primary crystallization product ␣-Fe(Si) in the early stage [8]. As a result, the activation energy of primary crystallization decreases dramatically while the secondary crystallization, Table 3 Crystallization activation energy Ea of an amorphous Fe75.5 Si13.5 B9 Cu2 alloy Crystallization process

Primary Secondary

Crystallization activation energy Ea (kJ/mol) Kissinger’s method

Doyle–Ozawa’s method

371.9 434.2

366.4 426.2

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Fig. 7. Thermal expansion coefficient curves of an amorphous Fe75.5 Si13.5 B9 Cu2 alloy at different heating rates: (a) low temperature range; (b) high temperature range.

Fig. 5. The relationship between lg ˇ and 1/T of crystallization for an amorphous Fe75.5 Si13.5 B9 Cu2 alloy: (a) peak 1; (b) peak 2.

which is the formation of Fe–B intermetallics, is not affected nearly by Cu addition effect. Fig. 6 shows the thermal expansion curves of samples at different heating rates. The curves exhibit almost the same complex shape and shift towards higher temperatures with heating rates

Fig. 6. Thermal expansion curves of an amorphous Fe75.5 Si13.5 B9 Cu2 alloy at different heating rates.

increasing. Each curve drops drastically during the crystallization at around 770 K and displays a positive linear expansion tendency after complete crystallization. The value of contraction is around 0.0033–0.0038, a little lower than that of Finemet 0.0042 [22]. Referring to the contraction value of Fe–Si–B about 0.003–0.005 [23], there is no much difference as to the contraction of Fe–Si–Bbased metallic glasses during crystallization. Furthermore, the anomalous change of thermal expansion occurs near Curie temperature Tc due to large spontaneous volume magnetostriction in Fe-based metallic glasses, in contrast, cobalt-based metallic glasses have nearly zero magnetostriction. And this large spontaneous volume magnetostriction is mainly ascribed to the minimization of the total value for thermal energy, elastic energy and magnetic energy [24]. Fig. 7 shows the thermal expansion coefficient variation of Fe75.5 Si13.5 B9 Cu2 alloy at different heating rates. In the low temperature range (Fig. 7(a)), the temperatures of initial structure relaxation are 385.5, 441.7 and 520.8 K, respectively. They are much lower than those of Fe–Si–B [23], which also manifests that the stability of amorphous Fe–Si–B-based alloy decreases due to Cu addition. Moreover, the alloy has a ferromagnetic state and a transition from ferromagnetic to paramagnetic state occurs at around 686–687 K close to the DSC results. Furthermore, Fig. 6(b) displays an inflexion when samples contract dramatically during crystallization. Little attention has been paid to this change in previous researches. This inflexion drops a hint of the occurrence of different crystallization stages and formation of new phases. According to the results of XRD, the two contraction peaks at 5 K/min shown in Fig. 7(b) are much similar to the two exothermic peaks of DSC curves owing to the sequential formation of bcc ␣-Fe(Si) crystallites and Fe–B intermetallic compounds. The inflexion in Fig. 6(b) can also be well explained that Fe–B compounds form when the second crystallization starts. Some researchers claimed that there were three stages of crystallization process for amorphous Fe–Si–B-based alloys [9,25]. In our experiment, the formation of Fe3 B and Fe2 B and the decomposition of Fe3 B occur almost simultaneously with the second

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contrast to the symmetrical shape of the second contraction peak at 5 K/min. And the second peak can be split into two at higher heating rates. This can be explained that metastable Fe3 B phase does not have enough time to complete the decomposition into Fe2 B at high heating rates (e.g. 10 and 20 K/min) and residual Fe3 B exists causing the occurrence of another peak. Since the magnetocrystalline anisotropy of bct Fe3 B is larger than bcc ␣Fe(Si), the appearance of Fe3 B markedly deteriorates soft magnetic properties [27]. So, the upper limited annealing temperature for primary crystallization is about 803 K according to the thermal expansion coefficient curve at 5 K/min. 3.3. Grain evolution during crystallization

Fig. 8. The relationship between the grain size of ␣-Fe(Si) and annealing temperatures for 1 h or annealing time at 773 K.

exothermic peak, so the bodies of peaks overlap and only one for the secondary crystallization can be distinguished in heat flow measurements [12,26]. Additionally, because of the same crystal structure (i.e. bct) for Fe3 B and Fe2 B phases, the Bragg intensive peaks of these two phases might not be appropriately identified and indexed sometimes (i.e. both Fe3 B and Fe2 B phases make contributions to the intensity of the peak (2 ≈ 43◦ ), this peak cannot be distinguished precisely). Nevertheless, in our DIL study in Fig. 7(b), a step appears at 10 K/min in the uphill part of the second main peak and one more peak comes up at 20 K/min in

Fig. 8 shows the grain size variation of ␣-Fe(Si) calculated by Scherer’s formula. From Fig. 8, it is found that grains grow in the sigmoid style and the growth rate increases sharply in the temperature range from 813 to 953 K. The minimum value of 13.7 nm appears at 773 K and the maximum is 41.2 nm at 953 K. According to Herzer’s theoretical model [10], excellent soft magnetic properties can be obtained by minimizing the grain size of primary crystalline phase ␣-Fe(Si). So, in order to refine grains further, isothermal annealing at 773 K was proceeded. The grain size is further reduced to 12.1 nm by shortening annealing time to 30 min. In order to further observe the grain evolution, transition electron microscopy was employed as shown in Fig. 9. Evidently, the as-quenched alloy has an entire amorphous state in Fig. 9(a). Quite a small volume fraction of visible nanograins of Fe–Si phase with less than 5 nm in size are randomly dispersed in the amorphous matrix at the beginning of crystallization as shown in Fig. 9(b). In this stage, nucleation occurs but is not saturated. Fig. 9(c) shows that Fe–Si

Fig. 9. TEM morphology and SAED patterns of an amorphous Fe75.5 Si13.5 B9 Cu2 alloy after isothermal annealing: (a) as-quenched; (b) 773 K, 5 min; (c) 773 K, 30 min; (d) 773 K, 60 min.

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nanograins continue precipitating from the amorphous matrix and crystallites with the size of less than 10 nm are distributed uniformly after annealing for 30 min at 773 K. Meanwhile, due to the increment of the crystallization volume fraction, the rings in the SAED pattern become sharper and consist with the Bragg peaks of ␣-Fe(Si) indexed by XRD results. For minimizing the magnetostriction anisotropy effect, an adequate volume fraction of Fe–Si, about 70–75%, and ultrafine soft magnetic grains (e.g. ␣-Fe(Si)) are required [8,28,29] through optimizing annealing processing. In our experiment, it is found that nanograins with a small size (about 25 nm) account for more than 65% of the total volume by annealing at 773 K for 60 min as shown in Fig. 9(d), and the saturated magnetization Ms and coercivity Hc by VSM measurements are 1.55 T and 37 A/m, respectively. Instead of the coarse dendritic morphology for Fe77.5 Si13.5 B9 [12] and Fe75 Si9 B16 Cux [9], equiaxed crystallites can be observed clearly for Fe77.5 Si13.5 B9 Cu2 with copper addition as shown in Fig. 9(c) and (d). In the absence of copper, more interfaces are created to reduce the overall interfacial energy of the system, or in other words the interfaces of dendrite branches come to be the critical factor for preferential nucleation sites of the precipitation of Fe–Si grains from the amorphous matrix. On the contrary, in the case of copper addition, because Cu atoms cannot dissolve into the matrix, they get together into nanometer-sized clusters in the amorphous state prior to the onset of primary crystallization according to Hono’s observation [30]. These clusters acting as heterogeneous nucleation sites can substantially reduce the free energy of nucleation and small crystals of Fe–Si are preferably nucleated at the Cu/amorphous interfaces. The nucleation rate is increased significantly or crystallization is accelerated, and thus the effect of refining grains can be achieved. So, to great extent, the average size and morphology of Fe–Si crystals depend on the number density of Cu clusters. Furthermore, according to a small-angle neutron scattering study, both of the primary crystallization temperature the precipitation temperature of Fe–B compounds are high in high-Si-content (11 and 13.5 at.%) Fe–Si–B–Cu–Nb alloy [31]. It is concluded that the kinetics of Cu clusters is affected by the thermal stability of the amorphous state. In other words, the opportunity for precipitation of Cu clusters is rare in the stable amorphous matrix [27]. As described in Section 3.2, the thermal stability of the remaining amorphous state is improved by the substitution of high Si content for B, suppressing the precipitation of Fe3 B and expanding the annealing temperature rage [32]. Therefore, a high Cu content (e.g. 2 at.% in our experimental alloy) is necessary for obtaining the optimum number density of Cu clusters in high-Si-content alloys, which leads to the microstructure of adequate ultrafine primary crystals for excellent soft magnetic properties. 4. Conclusions The crystallization activation energy of Fe75.5 Si13.5 B9 Cu2 for the first and second peaks calculated by Doyle–Ozawa’s method is 366 and 426 kJ/mol, respectively. Copper addition causes the Fe–Si–B metallic glass less stable detected by the decrease of the activa-

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tion energy of primary crystallization and the structure relaxation temperature due to copper’s microsegregation, concentration and heterogeneous nucleation effect. One more peak appears during the second crystallization in thermal expansion coefficient curves at higher than 10 K/min, which provides indirect evidence of the presence and transformation of metastable Fe3 B phase. Moreover, equiaxed nanometer-sized grains are promoted without forming coarse dentritic morphology by adding 2 at.% copper in Fe–Si–B. Both ultrafine-grain structure (about 25 nm) and the sufficient crystallization volume fraction (more than 65%) of Fe–Si phase are obtained by annealing at 773 K for 60 min and the corresponding alloy exhibits good soft magnetic properties with a highsaturated magnetization of 1.55 T and a relatively low coercivity of 37 A/m. Acknowledgements This work was partially supported by Shanghai Nanotechnology Special Program (Grant No. 065nm004). References [1] K.G. Efthimiadis, C.A. Achilleos, S.C. Chadjivasiliou, I.A. Tsoukalas, J. Magn. Magn. Mater. 171 (1997) 141–146. [2] M.A. Gibson, G.W. Delamore, J. Mater. Sci. 27 (1992) 3533–3538. [3] M. Pekala, M. Jachimowicz, V.I. Fadeeva, H. Matyja, A. Grabias, J. Non-Cryst. Solids 287 (2001) 380–384. [4] V. Cremaschi, I. Avram, T. Perez, H. Sirkin, Scripta Mater. 46 (2002) 95–100. [5] H. Chiriac, C.S. Marinescu, Sens. Actuators A 81 (2000) 174–175. [6] Y. Yoshizawa, S. Oguma, K. Yamauchi, J. Appl. Phys. 64 (1988) 6044–6046. [7] J.D. Ayers, V.G. Harris, J.A. Sprague, W.T. Elam, H.N. Jones, Acta Mater. 46 (1998) 1861–1874. [8] K. Hono, D.H. Ping, M. Ohnuma, H. Onodera, Acta Mater. 47 (1999) 997–1006. [9] K.G. Efthimiadis, E.K. Polychroniadis, S.C. Chadjivasiliou, I.A. Tsoukalas, Mater. Res. Bull. 35 (2000) 937–944. [10] G. Herzer, Scripta Metall. Mater. 33 (1995) 1741–1756. [11] F.E. Luborsky (Ed.), Translated by C. Ke, Y. Luo, K.Y. He, Metallurgical Industry, Beijing, 1989, p. 229. [12] Y.R. Zhang, R.V. Ramanujan, Mater. Sci. Eng. A 416 (2006) 161–168. [13] Y. Takahara, N. Narita, Mater. Trans. JIM 41 (2000) 1077–1081. [14] S.D. Kaloshkin, I.A. Tomilin, Thermochim. Acta 280–281 (1996) 303–317. [15] A. Takeuchi, A. Inoue, Mater. Trans. JIM 46 (2005) 2817–2829. [16] W. Lu, L. Yang, B. Yan, W.-H. Huang, J. Alloys Compd. 420 (2006) 186–192. [17] H.E. Kissinger, Anal. Chem. 29 (1957) 1170–1702. [18] T. Ozawa, J. Thermal Anal. 2 (1970) 301–324. [19] F.E. Luborsky (Ed.), Translated by C. Ke, Y. Luo, K.Y. He, Metallurgical Industry, Beijing, 1989, p. 217. [20] A.R. Bhatti, B. Cantor, J. Mater. Sci. 29 (1994) 816–823. [21] J.Y. Bang, R.Y. Lee, J. Mater. Sci. 26 (1991) 4961–4965. [22] Y.C. Niu, X.F. Bian, W.M. Wang, S.F. Jin, G.H. Li, F.M. Chu, W.G. Zhang, J. Alloys Compd. 433 (2007) 296–301. [23] Y.C. Niu, X.F. Bian, W.M. Wang, S.F. Jin, X.J. Liu, J.Y. Zhang, G.L. Qin, J. Non-Cryst. Solids 351 (2005) 3854–3860. [24] F.E. Luborsky (Ed.), Translated by C. Ke, Y. Luo, K.Y. He, Metallurgical Industry, Beijing, 1989, p. 428. [25] K. Chrissafis, M.I. Maragakis, K.G. Efthimiadis, E.K. Polychroniadis, J. Alloys Compd. 386 (2005) 165–173. [26] Y.R. Zhang, R.V. Ramanujan, Thin Solid Films 505 (2006) 97–102. [27] M. Ohta, Y. Yoshizawa, J. Appl. Phys. 103 (07) (2008) E722. [28] T. Kulik, J. Non-Cryst. Solids 287 (2001) 145–161. [29] T. Kulik, G. Vlasá, R. Zuberek, Mater. Sci. Eng. A 226-228 (1997) 701–705. [30] K. Hono, Acta Mater. 47 (1999) 3127–3145. [31] M. Ohnuma, K. Hono, S. Linderoth, J.S. Pedersen, Y. Yoshizawa, H. Onodera, Acta Mater. 48 (2000) 4783–4790. [32] M. Ohta, Y. Yoshizawa, Appl. Phys. Lett. 91 (2007) 062517.