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Thermal behaviors of Al-based amorphous alloys bearing nanocrystalline In particles
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Acta Materialia 58 (2010) 6267–6275 www.elsevier.com/locate/actamat
Thermal behaviors of Al-based amorphous alloys bearing nanocrystalline In particles J. Mu a,b, Z.W. Zhu a, H.F. Zhang a,*, Z.Q. Hu a, Y.D. Wang c, Y. Ren d a
Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China b Graduate School of the Chinese Academy of Sciences, Beijing 100039, China c School of Materials Science and Engineering, Beijing Institute of Technology, Beijing 100081, China d X-ray Science Division, Argonne National Laboratory, Argonne, IL 60439, USA Received 28 December 2009; received in revised form 19 July 2010; accepted 28 July 2010 Available online 20 August 2010
Abstract We successfully fabricated nanocrystalline (NC) indium (In) particles embedded in Al-based amorphous matrix. Systematic investigations indicate that thermal interaction between the NC In and the amorphous matrix significantly influences their respective thermal behaviors. The melting temperature of NC In was found to be depressed by 10–30 K, owing to the specific interfacial structure of NC In/amorphous system. The simultaneous appearance of the liquid/amorphous interface destabilizes the amorphous matrix, leading to face-centered cubic-Al precipitation at the interface of In sphere/amorphous matrix at a relatively lower temperature. This effect is attributed to the diffusion of La from the matrix to the liquid In particles. Ó 2010 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Melt-spinning; Nanocrystalline materials; Interface structure; Melting behavior; Crystallization
1. Introduction Compared to conventional crystalline alloys, the amorphous alloys have great potential applications as structural materials because of the superior properties [1–3]. However, lack of plasticity at room temperature is a fatal disadvantage [1–4]. Many works have suggested that the structural heterogeneities with different length scales, such as nano-particles and fiber, effectively improve the plasticity of bulk metallic glasses (BMGs) by inducing the formation of multiple shear bands [5–11]. Hence BMG matrix composite (BMGC) has been widely studied and exhibits outstanding properties [5–13]. For BMGC, the interface occurring due to the introduction of structural inhomogeneities plays an important role
*
Corresponding author. E-mail address: [email protected] (H.F. Zhang).
in their properties [5–7,12–15]. Up to now, abundant studies have been concentrated on the influence of the interface on mechanical performance [5–7,12–15]. As we know, mechanical properties of BMGs and BMGCs are closely correlated to their thermal properties [16–19]. Temperature rising, thermal softening and deformationinduced nanocrystallization are believed to occur during shear banding events [4,20,21]. Moreover, since the amorphous phase is thermodynamically metastable, it tends to crystallize on heating. Resultant partial crystallization may lead to embrittlement [22–25]. Thus for BMGC alloys, besides the mechanical property, investigations on thermal interaction between the second phase and the amorphous matrix are of scientific and technological importance. As a saltation of the structure, energy, and even the composition, the interface may play an important role in the nucleation and growth in the thermal behaviors. Elemental diffusion and interface reaction occur through it [5,26,27], and phase transition also happens under the influence of it
1359-6454/$36.00 Ó 2010 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2010.07.048
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[28–31]. However, how does the interface influence the thermodynamic properties of the BMGC alloy in detail? Few works focus on this point. In this work, the composite of nanocrystalline (NC) indium (In) particles dispersing on the amorphous matrix was synthesized. The detailed works were performed on the thermal interaction of NC In and the amorphous matrix. Melting temperature depression has been observed for In particles, and distinct crystallization, especially the multistage precipitation of face centered cubic (fcc)-Al, occurs for the amorphous matrix. 2. Experimental Ingots of (Al0.86Ni0.09La0.05)(100x)Inx (x = 0, 0.5, 3, 5) (in at.%) were obtained by melting the elements of spectroscopic grades in a graphite crucible under high purity argon atmosphere. The ribbons were prepared using melt-spinning method. The linear velocity of the copper wheel was 40 m s–1. The thickness of the ribbons is 50 lm. The structure of the ribbons was examined by X-ray diffractometry (XRD; Rigaku, D/Max-2500PC, Tokyo, Japan, Cu Ka) and high-resolution transmission electron microscopy (HRTEM, Tecnai G2 F30, 200 kV). Thermal behaviors were analyzed by differential scanning calorimetry (DSC; Netzsch DSC 204F, Germany) at the heating rates of 10, 20, 40 and 80 K min1 in a flowing argon. The melting behavior of the nano-particles was studied by in situ synchrotron high energy X-ray diffraction at the 11-ID-C beam line of the Advanced Photon Source, Argonne National Laboratory, USA. In situ observations were performed using TEM (JEOL 2010, 200 kV) coupled with a heating apparatus. In order to identify the crystallization products of the systems, the ex situ experiments were performed such that the samples were heated to every peak temperature (Pi) at 10 K min1 and subsequently cooled to the room temperature, examined by XRD and TEM.
The detailed microstructural morphology of the composites was examined by TEM observations. Fig. 2a shows the bright-field (BF) TEM micrograph and the corresponding selected-area electron diffraction (SAED) patterns (inset in Fig. 2a) of the melt-spun (Al0.86Ni0.09La0.05)95In5 alloy. The isolated spherical particles with the size of below 40 nm in diameter are embedded in the featureless contrast matrix. The particles and matrix are identified as In (tetragonal, body-centered, a = b = 3.215, c = 4.932 [32]) and amorphous phase, respectively, from their corresponding SAED patterns. The EDS results support these identifications. These results confirm the XRD analysis as shown in Fig. 1, suggesting that In NC particles are embedded in the amorphous matrix. The interface between In particle and the amorphous matrix is shown in the high-resolution TEM observations (Fig. 2b), which further confirmed that there is no reaction product in the interface between the two phases. Diffraction patterns are the same as the inset in Fig. 2a.
3. Results
3.2. Thermal behaviors
3.1. Microstructures
Fig. 3 shows the DSC curves of the melt-spun ribbons at a heating rate of 10 K min1. For Al–La–Ni alloy, the x = 0 curve in Fig. 3a shows two exothermic peaks at the experimental temperature range. The location of peaks change gradually as adding In to the system. As enlarged in Fig. 3b, for the x = 0.5 alloy, the peaks shift to lower temperature. When the content of In increases to 5 at.%, with NC In particles forming in the matrix as shown in Figs. 1 and 2, the onset crystallization temperatures (Tx) of the composites become much lower and the four exothermic peaks are shown. The first two peaks overlap each other. A similar condition is shown for the x = 3 alloy. These results apparently show that the occurrence of NC In particles dramatically alter the crystallization mechanism of the amorphous matrix during continuous heating. The temperatures obtained from the DSC curve are summarized in Table 1. (Since the first three peaks of the composites and the first peak of the x = 0 and 0.5 alloys are all
Fig. 1 shows the XRD patterns of the melt-spun (Al0.86Ni0.09La0.05)(100x)Inx (x = 0, 0.5, 3 and 5) alloys. For (Al0.86Ni0.09La0.05)95In5 (x = 5) alloy, it is seen that sharp peaks superimposed on a broad scattering hump typical of amorphous phase can be well indexed as pure In crystalline phase [32]. This suggests that a mixing structure of In crystal and amorphous matrix is formed in the studied alloy system. The similar case is also found for the x = 3 alloy, except that the intensity of the In peaks is much lower. However, as the content of In decreases to 0.5 at.%, the patterns show only one broad peak, indicating the formation of a single amorphous phase. For comparison, the XRD patterns of the melt-spun Al86Ni9La5 (x = 0) amorphous alloy are also displayed in Fig. 1. These results show that the composites are indeed formed when the proper content of In is added.
Fig. 1. The XRD pattern of the melt-spun (Al0.86Ni0.09La0.05)(100x)Inx (x = 0, 0.5, 3 and 5) alloys.
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Fig. 3. DSC curves of the melt-spun (Al0.86Ni0.09La0.05)(100x)Inx (x = 0, 0.5, 3 and 5) ribbons. (a) Continuous heating DSC curves of the melt-spun ribbons. The low temperature part is enlarged in (b).
Fig. 2. TEM image of melt-spun (Al0.86Ni0.09La0.05)95In5 alloy. (a) Nanoparticles are embedded in featureless contrast matrix in the bright-field (BF) TEM image and the inset shows the corresponding SAED pattern. (b) High-resolution transmission electron microscopy (HRTEM) image is verified the structure of the crystal particles and the amorphous matrix and shows the detail structure of the interface.
corresponding to the precipitation of fcc-Al, which will be discussed in detail in Section 3.4, their peak temperatures (Tpi) are summarized together in one row). 3.3. Melting behaviors of NC In The inset in Fig. 3a enlarges the DSC curve of the meltspun (Al0.86Ni0.09La0.05)95In5, where an endothermic peak can be found. In a heating and cooling circle, an endothermic peak appeared during heating process and an exothermic
peak occurred during cooling process in the temperature ranges of 400–419 K and 387–412 K, respectively. Similar phenomena cannot be observed on the scanning of Al86Ni9La5 system (Fig. 3). The detailed microstructure investigations as described in Fig. 2 reveal that the melting of NC In spheres could be responsible for the peaks of 400–419 K displayed in the DSC curve of (Al0.86Ni0.09La0.05)95In5 alloy. The melting of In particles was further verified using in situ high-energy synchrotron X-ray diffraction and in situ TEM investigations at a heating rate of 10 K min1. The XRD patterns of (Al0.86Ni0.09La0.05)95In5 ribbons collected at different temperatures are displayed in Fig. 4. There is one point needing to be noticed. The crystallization of the matrix around the particle took place as soon as the NC In particles melt, which will be discussed in detail in Section 4.2. Although the positions of some NC In peaks are coincided with those of the crystallized products, some peaks reflecting from NC In can be selected to characterize its phase transformation during the heating process. For example, the diffraction peaks at 2.48°, 4.17° and 4.43° of In, marked as 1, 2 and 3 in Fig. 4, are chosen to avoid the influence of the other phases. At 300 K, XRD patterns (the bottom one in Fig. 4a) of the melt-spun sample show that the studied alloy contains the crystalline In
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Table 1 The onset crystallization (Tx) and peak (Tpi) temperatures gained from the DSC experiment. Composition
Tx (K)
Al0.86Ni0.09La0.05 (Al0.86Ni0.09La0.05)99.5In0.5 (Al0.86Ni0.09La0.05)97In3 (Al0.86Ni0.09La0.05)95In5
512.4 475.7 Not clear 435.9
Tp1/Tp1, Tp2, Tp3 (K)
446.4 447.1
Fig. 4. In situ studies of the melt-spun (Al0.86Ni0.09La0.05)95In5 ribbon at different temperature by high-energy XRD. The small angle area is enlarged in (b). With rising temperature, the diffraction peaks, which belong to crystal In, become wide, even disappearing, indicating the melting behavior of the In nano-particles.
along with a tiny peak corresponding to small amount of crystalline Al. The occurrence of Al-ordering structure, such as pre-existing nuclei [33] or medium-range ordering
Tp2/Tp4 (K) 518.6 504.5 465.7 466.0
493.5 496.2
591.4 588.5 572.4 561.4
[34,35], were previously reported in Al-based amorphous alloys due to their low glass forming ability and high Al content. This kind of tiny and trace Al-ordering structure is rarely detected using commonly used XRD and TEM, as shown in Figs. 1 and 2. With increasing temperature, the diffraction peaks corresponding to crystalline In become wider and weaker. When the temperature reached 410 K, peak 2 of In began to disappear. At 425 K, the 1–3 peaks all disappeared. After heating to 425 K, the sample was cooled down to 300 K, and all peaks that disappeared after heating to 425 K come out again (the top curve shown in Fig. 4). The dependence of the intensity of the diffraction peaks for In with the temperature described above reveals that the melting took place for In particles [36]. Fig. 5 shows the dark field TEM micrographs of the same area of melt-spun (Al0.86Ni0.09La0.05)95In5 specimen in situ taken at room temperature, 393 K, 403 K and 413 K, respectively. The particles are marked 1 to 6. The small particles 5 and 6 melt at the lower temperature, indicated by the disappearance of the corresponding contrast as shown in Fig. 5b [37]. However, the big particles 1 and 2 do not melt completely until 413 K, although they present some melting features in Fig. 5d. The in situ investigations
Fig. 5. Dark field TEM micrographs of the same area of melt-spun (Al0.86Ni0.09La0.05)95In5 specimen in situ taken at (a) room temperature, (b) 393 K, (c) 403 K and (d) 413 K, respectively.
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are completely in a good agreement with the endothermic event in the range of 400–419 K, indicated by the DSC measurement (Fig. 3). These results surely indicate that the melting took place for In particles in the temperature range of 400–419 K, which is 10–30 K lower than the melting temperature T0 (429 K) [38] of the In bulk sample. The melting behaviors of NC In particles embedded in the amorphous matrix as shown in the present study are very different from the superheating of NC In particles embedded in the crystalline matrix [30,37,39]. The different interfacial structures between NC In/amorphous and NC In/ crystalline are thought to be responsible for these different melting behaviors of NC In. The mechanism underlying these unique phenomena will be discussed in Section 4.1 in detail.
Fig. 6a, the amorphous primarily precipitated fcc-Al and then decomposed into fcc-Al, Al3Ni and Al11La3. With the content of In higher than 3%, which leads to the formation of NC In particles, the first three peaks correspond to the formation of fcc-Al, as shown in Fig. 6b, and the last peak reaction products are the same as those of the x = 0 and 0.5 alloys. That is to say, the formation of NC In particles changes the mechanism of fcc-Al precipitation and leads to the loss of the thermal stability of the amorphous matrix. Fig. 7 shows Kissinger plots of the peak temperatures for the present alloys. The apparent activation energies Epi for crystallization reaction could be determined by Kissinger’s equation [40]: ln
3.4. Crystallization of amorphous matrix As stated in Section 3.2, the thermal behaviors were distinctly altered by the formation of NC In particles. In this part, the crystallization behaviors were studied in detail. Fig. 6 shows XRD patterns of the samples annealed to each peak temperature at a heating rate of 10 K min1. With minor addition of In (x = 0.5), which does not result in the formation of NC In particle, the crystallization behaviors were not tuned, as shown in Fig. 3a. In
Fig. 6. XRD patterns of the melt-spun (Al0.86Ni0.09La0.05)(100x)Inx (x = 0, 0.5, 3 and 5) ribbons after heating up to every crystallization peak temperature with 10 K min1.
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T 2pi Epi ¼ þC / k B T pi
ð1Þ
where / is the heating rate, kB is Boltzmann’s constant, Tpi is the ith peak temperature and C is a constant. The measured values can be linear-fitted well and accordingly Epi is calculated, as shown in Fig. 7. From Fig. 7a, it is seen that In addition decreases the activation energy of the fccAl precipitation, especially that for the x = 5 alloy, the activation energy of the first peak reaction is much lowered. The advanced melting of NC In particles as described in
Fig. 7. The Kissinger plots of crystallization peaks’ temperatures for the (Al0.86Ni0.09La0.05)(100x)Inx (x = 0, 0.5 and 5) ribbons.
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Section 3.3 may be responsible for this behavior. It suggests that the interaction of NC In and the amorphous matrix strongly decrease the thermal stability of the amorphous matrix. The primary precipitation of fcc-Al destabilizes the amorphous matrix. As displayed in Fig. 7b, the activation energy of the reaction of fcc-Al, Al3Ni and Al11La3, likewise reduces with increasing the In content. Multistage formation of fcc-Al in (Al0.86Ni0.09La0.05)95In5 system has attracted our interest. The sample which was heated to the first crystallization peak temperature and cooled down was examined by TEM. Fig. 8a shows the shape of the NC In particles is not a regular sphere. The interface between the In particle and the amorphous matrix is misty. Some new phase appears around the In spheres. However, the matrix far away from the In particles is still in featureless amorphous contrast. A high-resolution TEM image exhibits the interfacial structure of the system. As shown in Fig. 8b, there is indeed some new phase between the In particles and the amorphous matrix. It is confirmed to be crystalline Al through the Fourier transform. These results illustrate that the first stage of crystallization is precipitation of fcc-Al around the In particles. 4. Discussion Melting and crystallization are generally both nucleation and growth processes, which are largely influenced by the fluctuation of structure, energy and even composition, etc., in the materials. For BMGCs, the interface is very crucial to the distribution of such factors. Undoubtedly, the interface plays an important role in the thermal properties of the materials. In the following section, the influence of crystal/amorphous interface on the melting of the In particles is firstly discussed, and the influence of the resultant liquid/amorphous interface on the crystallization of the amorphous matrix is be discussed in detail. 4.1. Melting behaviors of NC In particles The melting of embedded NC particles has been widely studied. The superheating [30,31,37,39,41–43] and depression [28,29,42,44–47] of melting temperature have been both reported. Different theories were employed to interpret these phenomena [28,30,31,37,39,42,47]. Literatures on studying the melting of NC particle reported that for most of the rapid solidification systems, the NC particles embedded in crystalline matrix can be easily superheated [30,37,41,43,48,49]. However, when the matrix becomes amorphous, melting temperature depression is observed for NC In particles in this work. These results imply that the matrix state will influence the melting behaviors of the embedded NC particles. Herein, it is thought that the interfacial structure determines the melting behaviors of NC particle. According to the thermodynamic model [50], the melting temperature Tme of embedded particles can be simply expressed as
Fig. 8. (a) TEM image and (b) HRTEM image of (Al86Ni9La5)0.95In5 melt-spun ribbons annealed at the first crystalline peak temperature.
T me ðrÞ ¼ T 0
i 3 h 2=3 1 csm clm ðqs =ql Þ qs rLm
ð2Þ
where T0 is the melting temperature of bulk crystalline phase, r is the radius of the particle, Lm is the latent heat of melting, csm and clm are the interfacial energies between the solid/matrix and liquid/matrix, respectively, qs and ql are the densities of solid and liquid, respectively. From Eq. (2) it can be inferred that with supposing qs ql, the change in Tme depends on csm and clm. When the value of csm is larger than that of clm, the melting temperature of the embedded NC particles will be depressed. Contrarily, the embedded NC particles are superheated with csm smaller than clm, which is deducted in the most of
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melt-spinning NC-particle/crystalline-matrix systems [30,41]. In our work, considering the atomic structure of metallic glass is similar to that of liquid [51], it is expected that csm > clm for the In spheres embedded in the amorphous matrix system. To calculate the interfacial energies between solid/ matrix and liquid/matrix, a simple configuration model of the interfaces between both single component systems, which is modified based on Spaepon’s model, is proposed (the details on the model and the calculation are described in Supplementary material). Since the model is ideal and does not contain mismatch and disordered area, it only fits the well-matched interface in the rapid solidification system. The calculation is under the viewpoint that the interfacial energy is the summation of the binding energy [52]. The energy for different interfaces is estimated with the energy caused by phase transition and the mixing of the two elements. Thus the interfacial energy is calculated as follows:
1=3
3 1 16 N 0 V 2 5 5 DH Am þ DH Bm þ 2DH mix AB 2 2
cA=B cc ¼
ð3Þ
(ii) Interfacial energy of crystalline A/liquid B
A=B ccl
1=3 3 1 ¼ ð5DH Am þ 2DH mix AB Þ 16 N 0 V 2
ð4Þ
(iii) Interfacial energy of amorphous A/crystalline B
cA=B ac ¼
1=3 3 1 5 B mix DH þ 2DH m AB 16 N 0 V 2 2
ð5Þ
(iv) Interfacial energy of amorphous A/liquid B
A=B cal
1=3 3 1 ¼ 2DH mix AB 16 N 0 V 2
In/Al matrix [30], NC–Pb/Al matrix [43], NC Pb/Cu matrix [43] and NC Ag/Ni matrix [41] systems, which coincide with the qualitative deduction in previous literatures. Contrarily, ccl < ccc occurs in the melt-spun NC Bi/ Zn matrix system, where melting temperature depression is investigated for NC Bi particles [47]. However, since the model is ideal, and ignores many factors, there are still some cases that could not be explained, such as no superheating of the NC Pb embedded in Ni matrix [43] and premelting of the NC particles dispersed at the grain boundary [28], which are ascribed to the disordered structure of the interface. In our work, for the In NC particles and Al-based amorphous matrix system, Al is A and In is B. Because of the high content of Al in Al-based amorphous alloy, we consider the amorphous matrix contains only Al element for simplicity. According to the Eqs. (5) and (6), the calculated results are cac ¼ 0:098 J=m2 cal ¼ 0:067 J=m2
(i) Interfacial energy of crystalline A/crystalline B
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ð6Þ
where V is the molar volume of element, DH X m is the enthalpy of the melting process, DH mix is the enthalpy of AB mixing A and B, N0 is the Avogadro constant. Eqs. (3) and (4) are applied to the NC particles embedded in crystallization matrix systems. The calculated results (shown in Supplementary material) indicate that the value of ccl is indeed larger than that of ccc in the melt-spun NC
Then the calculated interfacial energies are put into Eq. (2), and it indicates that the melting temperature of the particles embedded in the amorphous matrix is depressed. The calculation results coincide with the expectation and phenomena stated above. Thus the interfacial energy induces the melting temperature depression. The size of NC particles is another important factor influencing the melting behaviors. From Eq. (2), it is inferred that the smaller the size of NC particles, the higher the deviation of the melting temperature of NC particles from that of bulk samples. It is evidenced in the works of the melt-spun NC In/Al matrix [30], NC–Pb/Al matrix [37] and NC–Ag/Ni matrix [41] systems. In the present work, it works very well. As observed in the in situ TEM observations in Fig. 5, it clearly shows that NC In with the smaller size melts at the lower temperature. However, the quantitative relationship between the size and the melting temperature cannot be presented in this work, which is ongoing. In addition, other factors influencing the melting behaviors of NC particles embedded in amorphous matrix are argued. The first one is the alloying effect. Alloying significantly influences the melting temperature of the system. In the present work, Al–In is an immiscibility system [38]. EDS results show the NC In is pure, indicating extremely limited content of La or Ni might dissolve into In within an error. Even if the elements dissolve into the In spheres, it only can cause a 1 K change in melting temperature for the bulk materials [38]. This cannot explain our experimental observation on the change in melting temperature. So solution of other elements cannot be responsible for the depression of In NC particles as evidenced in our work. The second one is the effect of the disordered region of the interface. Zhang and Cantor [37] considered that melting behaviors of the embedded NC particles are associated
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with the density of defects at the interface, which is the preferential place for melting nucleation. According to the study of Chen and Chuang [53], when the only existing phases boundaries are those between the crystal and the glass, hardly any defects are detected at the interface. Therefore the nucleation of melting may be more difficult in the spherical In phase embedded in amorphous matrix. Thus the disordered region of the interface from the reasons for the depression of the melting temperature in this work can be excluded. The third one is the strain energy effect. The strain energy usually arises from the mismatch in coefficients of thermal expansion or volume expansion during melting. Such a strain energy can elevate the melting temperatures of the embedded particles [41,54]. It should be considered in the NC-particle/crystalline-matrix system rather than our present system. 4.2. Crystallization of the amorphous matrix For the Al–La–Ni amorphous system [55], the core–shell structure occurs during primary crystallization. La-rich shells around the Al core build up when La atoms are rejected from the growing crystals of almost pure Al. These shells hinder further growth and, thus, limit the size of the Al nanocrystals, stabilizing the amorphous matrix. In the present study, as stated in Section 3.4, the addition of In damages the thermal stability of the amorphous matrix, leading to the change of the crystallization behaviors. From the structural analysis (Section 3.1), such behaviors may result from two contributions, one of which is the interface between In and amorphous matrix, the other is the dissolved In. As discussed in Section 3.3, the NC In particles melt prior to the crystallization of the amorphous matrix during continuous heating process. Apparently, the resultant liquid In/amorphous has a more significant effect than the solid/solid interface [5,26,27] on the crystallization behaviors. Because the value of the diffusion coefficient in liquid is several magnitudes higher than that in solid, the element is more easily diffused from the amorphous matrix to the In through this interface. According to Fig. 6b, the peaks of In, especially the main peak at 33°, shift to the small angle gradually. It means there are some La dissolved into In, since in Al–Ni–La–In system, only the radius of La atom is larger than In (The radius of Al, Ni, La and In are 0.14317 [56], 0.12459 [56], 0.18790 [56] and 0.16590 nm [57], respectively), and La is the only one that has negative mixing enthalpy with In in this system [38]. When La is dissolved into In liquid, the matrix around the In nano-particles is poor in La, which will delay the formation of La-rich shells and make the primary precipitation of fcc-Al in the zone proceed more easily. Correspondingly, the activation energy is largely reduced, as shown in Fig. 7. This conclusion is also consistent with the TEM observations in Fig. 8, in which the fcc-Al was indeed observed near the interface. Furthermore, since the melting of the particles occurs in a
temperature range of 20 K, the crystallization of the matrix around the smaller particles occurs at the same time with the melting process of the larger particles. Thus, the melting process of the particles is not an absolutely separated step from the crystallization of the matrix. Different from the diffusion induced crystallization around the NC particles, the crystallization of the matrix far away from the NC particles should be mainly under the influence of the dissolved In. For the x = 0.5 alloy, the melt-spun ribbon is in amorphous phase (as shown in Fig. 1) and the crystallization products are the same as the alloy without In (see Fig. 6). Thus the affect of dissolved In can be studied in this alloy. In Figs. 3 and 7, the dissolved In slightly decreases the temperature of fccAl precipitation. It is probably owing to the positive enthalpy of Al–In, which can reduce the pack density of the amorphous phase and make the atomic diffusion undertake more easily [Q1]. The effect of the so-called core–shell structure undermined [Q2]. Similar cases occur in the Zr– Ni–Cu–Al system with Ti partially replacing Zr [58] and Fe–P–C or Fe–Si–B systems with Cu partially replacing Fe [59,60]. There is still another factor that cannot be ignored. Under the free condition, the melting of In sphere must cause the increase of volume; however, in this system the amorphous matrix prevents this change. It brings the compressing stress field to both the liquid and the matrix around the sphere. According to Ref. [61], the applied pressure enhances the precipitation of the fcc-Al from the amorphous phase. Once the NC In particles are formed, these effects destabilize the amorphous matrix and obviously result in the change of the crystallization behaviors. 5. Conclusions A homogeneous dispersion of small spherical In particles within an amorphous Al–Ni–La matrix has been produced by melt-spinning method. The interaction between the NC particles and the amorphous matrix through the interface is studied. Melting temperature depression of the In NC particles and multistage crystallization of the amorphous matrix have been observed. Melting temperature depression is attributed to the specified structure at interface between the NC metal and amorphous matrix. A simple physical model considering the detailed atomic configuration at interface between NC metal and amorphous phase is used to elucidate the interface-mediated melting phenomenon. The interfacial energy of the crystal/glass state, which is higher than that of liquid/amorphous matrix, provides a driving force to promote the melting of NC metallic particles. The crystallization of the (Al0.86Ni0.09La0.05)95In5 composite starts soon after the melting of the In particles. Under the influence of the diffusion and the compressing stress, fcc-Al firstly precipitates from the matrix around the In particles. Then the primary crystallization of the matrix far away from the NC particles occurs. Finally, as
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the temperature goes up, the rest of the matrix turns into crystalline phase totally. Acknowledgments PJ Shang and ZQ Liu in Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences are appreciated for performing the in situ TEM experiments. The authors gratefully acknowledge the financial support from the Ministry of Science and Technology of China (Grants No. 2006CB605201) and the National Natural Science Foundation of China (Grant Nos. 50825402 and 50725102). Use of the Advanced Photon Source was supported by the US Department of Energy, Office of Science, Office of Basic Energy Science, under Contract No. DE-AC02-06CH11357. Appendix A. Supplementary material Supplementary data associated with this article can be found, in the online version, at doi:10.1016/j.actamat. 2010.07.048. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19]
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Acta Materialia 58 (2010) 6267–6275 www.elsevier.com/locate/actamat
Thermal behaviors of Al-based amorphous alloys bearing nanocrystalline In particles J. Mu a,b, Z.W. Zhu a, H.F. Zhang a,*, Z.Q. Hu a, Y.D. Wang c, Y. Ren d a
Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China b Graduate School of the Chinese Academy of Sciences, Beijing 100039, China c School of Materials Science and Engineering, Beijing Institute of Technology, Beijing 100081, China d X-ray Science Division, Argonne National Laboratory, Argonne, IL 60439, USA Received 28 December 2009; received in revised form 19 July 2010; accepted 28 July 2010 Available online 20 August 2010
Abstract We successfully fabricated nanocrystalline (NC) indium (In) particles embedded in Al-based amorphous matrix. Systematic investigations indicate that thermal interaction between the NC In and the amorphous matrix significantly influences their respective thermal behaviors. The melting temperature of NC In was found to be depressed by 10–30 K, owing to the specific interfacial structure of NC In/amorphous system. The simultaneous appearance of the liquid/amorphous interface destabilizes the amorphous matrix, leading to face-centered cubic-Al precipitation at the interface of In sphere/amorphous matrix at a relatively lower temperature. This effect is attributed to the diffusion of La from the matrix to the liquid In particles. Ó 2010 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Melt-spinning; Nanocrystalline materials; Interface structure; Melting behavior; Crystallization
1. Introduction Compared to conventional crystalline alloys, the amorphous alloys have great potential applications as structural materials because of the superior properties [1–3]. However, lack of plasticity at room temperature is a fatal disadvantage [1–4]. Many works have suggested that the structural heterogeneities with different length scales, such as nano-particles and fiber, effectively improve the plasticity of bulk metallic glasses (BMGs) by inducing the formation of multiple shear bands [5–11]. Hence BMG matrix composite (BMGC) has been widely studied and exhibits outstanding properties [5–13]. For BMGC, the interface occurring due to the introduction of structural inhomogeneities plays an important role
*
Corresponding author. E-mail address: [email protected] (H.F. Zhang).
in their properties [5–7,12–15]. Up to now, abundant studies have been concentrated on the influence of the interface on mechanical performance [5–7,12–15]. As we know, mechanical properties of BMGs and BMGCs are closely correlated to their thermal properties [16–19]. Temperature rising, thermal softening and deformationinduced nanocrystallization are believed to occur during shear banding events [4,20,21]. Moreover, since the amorphous phase is thermodynamically metastable, it tends to crystallize on heating. Resultant partial crystallization may lead to embrittlement [22–25]. Thus for BMGC alloys, besides the mechanical property, investigations on thermal interaction between the second phase and the amorphous matrix are of scientific and technological importance. As a saltation of the structure, energy, and even the composition, the interface may play an important role in the nucleation and growth in the thermal behaviors. Elemental diffusion and interface reaction occur through it [5,26,27], and phase transition also happens under the influence of it
1359-6454/$36.00 Ó 2010 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2010.07.048
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[28–31]. However, how does the interface influence the thermodynamic properties of the BMGC alloy in detail? Few works focus on this point. In this work, the composite of nanocrystalline (NC) indium (In) particles dispersing on the amorphous matrix was synthesized. The detailed works were performed on the thermal interaction of NC In and the amorphous matrix. Melting temperature depression has been observed for In particles, and distinct crystallization, especially the multistage precipitation of face centered cubic (fcc)-Al, occurs for the amorphous matrix. 2. Experimental Ingots of (Al0.86Ni0.09La0.05)(100x)Inx (x = 0, 0.5, 3, 5) (in at.%) were obtained by melting the elements of spectroscopic grades in a graphite crucible under high purity argon atmosphere. The ribbons were prepared using melt-spinning method. The linear velocity of the copper wheel was 40 m s–1. The thickness of the ribbons is 50 lm. The structure of the ribbons was examined by X-ray diffractometry (XRD; Rigaku, D/Max-2500PC, Tokyo, Japan, Cu Ka) and high-resolution transmission electron microscopy (HRTEM, Tecnai G2 F30, 200 kV). Thermal behaviors were analyzed by differential scanning calorimetry (DSC; Netzsch DSC 204F, Germany) at the heating rates of 10, 20, 40 and 80 K min1 in a flowing argon. The melting behavior of the nano-particles was studied by in situ synchrotron high energy X-ray diffraction at the 11-ID-C beam line of the Advanced Photon Source, Argonne National Laboratory, USA. In situ observations were performed using TEM (JEOL 2010, 200 kV) coupled with a heating apparatus. In order to identify the crystallization products of the systems, the ex situ experiments were performed such that the samples were heated to every peak temperature (Pi) at 10 K min1 and subsequently cooled to the room temperature, examined by XRD and TEM.
The detailed microstructural morphology of the composites was examined by TEM observations. Fig. 2a shows the bright-field (BF) TEM micrograph and the corresponding selected-area electron diffraction (SAED) patterns (inset in Fig. 2a) of the melt-spun (Al0.86Ni0.09La0.05)95In5 alloy. The isolated spherical particles with the size of below 40 nm in diameter are embedded in the featureless contrast matrix. The particles and matrix are identified as In (tetragonal, body-centered, a = b = 3.215, c = 4.932 [32]) and amorphous phase, respectively, from their corresponding SAED patterns. The EDS results support these identifications. These results confirm the XRD analysis as shown in Fig. 1, suggesting that In NC particles are embedded in the amorphous matrix. The interface between In particle and the amorphous matrix is shown in the high-resolution TEM observations (Fig. 2b), which further confirmed that there is no reaction product in the interface between the two phases. Diffraction patterns are the same as the inset in Fig. 2a.
3. Results
3.2. Thermal behaviors
3.1. Microstructures
Fig. 3 shows the DSC curves of the melt-spun ribbons at a heating rate of 10 K min1. For Al–La–Ni alloy, the x = 0 curve in Fig. 3a shows two exothermic peaks at the experimental temperature range. The location of peaks change gradually as adding In to the system. As enlarged in Fig. 3b, for the x = 0.5 alloy, the peaks shift to lower temperature. When the content of In increases to 5 at.%, with NC In particles forming in the matrix as shown in Figs. 1 and 2, the onset crystallization temperatures (Tx) of the composites become much lower and the four exothermic peaks are shown. The first two peaks overlap each other. A similar condition is shown for the x = 3 alloy. These results apparently show that the occurrence of NC In particles dramatically alter the crystallization mechanism of the amorphous matrix during continuous heating. The temperatures obtained from the DSC curve are summarized in Table 1. (Since the first three peaks of the composites and the first peak of the x = 0 and 0.5 alloys are all
Fig. 1 shows the XRD patterns of the melt-spun (Al0.86Ni0.09La0.05)(100x)Inx (x = 0, 0.5, 3 and 5) alloys. For (Al0.86Ni0.09La0.05)95In5 (x = 5) alloy, it is seen that sharp peaks superimposed on a broad scattering hump typical of amorphous phase can be well indexed as pure In crystalline phase [32]. This suggests that a mixing structure of In crystal and amorphous matrix is formed in the studied alloy system. The similar case is also found for the x = 3 alloy, except that the intensity of the In peaks is much lower. However, as the content of In decreases to 0.5 at.%, the patterns show only one broad peak, indicating the formation of a single amorphous phase. For comparison, the XRD patterns of the melt-spun Al86Ni9La5 (x = 0) amorphous alloy are also displayed in Fig. 1. These results show that the composites are indeed formed when the proper content of In is added.
Fig. 1. The XRD pattern of the melt-spun (Al0.86Ni0.09La0.05)(100x)Inx (x = 0, 0.5, 3 and 5) alloys.
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Fig. 3. DSC curves of the melt-spun (Al0.86Ni0.09La0.05)(100x)Inx (x = 0, 0.5, 3 and 5) ribbons. (a) Continuous heating DSC curves of the melt-spun ribbons. The low temperature part is enlarged in (b).
Fig. 2. TEM image of melt-spun (Al0.86Ni0.09La0.05)95In5 alloy. (a) Nanoparticles are embedded in featureless contrast matrix in the bright-field (BF) TEM image and the inset shows the corresponding SAED pattern. (b) High-resolution transmission electron microscopy (HRTEM) image is verified the structure of the crystal particles and the amorphous matrix and shows the detail structure of the interface.
corresponding to the precipitation of fcc-Al, which will be discussed in detail in Section 3.4, their peak temperatures (Tpi) are summarized together in one row). 3.3. Melting behaviors of NC In The inset in Fig. 3a enlarges the DSC curve of the meltspun (Al0.86Ni0.09La0.05)95In5, where an endothermic peak can be found. In a heating and cooling circle, an endothermic peak appeared during heating process and an exothermic
peak occurred during cooling process in the temperature ranges of 400–419 K and 387–412 K, respectively. Similar phenomena cannot be observed on the scanning of Al86Ni9La5 system (Fig. 3). The detailed microstructure investigations as described in Fig. 2 reveal that the melting of NC In spheres could be responsible for the peaks of 400–419 K displayed in the DSC curve of (Al0.86Ni0.09La0.05)95In5 alloy. The melting of In particles was further verified using in situ high-energy synchrotron X-ray diffraction and in situ TEM investigations at a heating rate of 10 K min1. The XRD patterns of (Al0.86Ni0.09La0.05)95In5 ribbons collected at different temperatures are displayed in Fig. 4. There is one point needing to be noticed. The crystallization of the matrix around the particle took place as soon as the NC In particles melt, which will be discussed in detail in Section 4.2. Although the positions of some NC In peaks are coincided with those of the crystallized products, some peaks reflecting from NC In can be selected to characterize its phase transformation during the heating process. For example, the diffraction peaks at 2.48°, 4.17° and 4.43° of In, marked as 1, 2 and 3 in Fig. 4, are chosen to avoid the influence of the other phases. At 300 K, XRD patterns (the bottom one in Fig. 4a) of the melt-spun sample show that the studied alloy contains the crystalline In
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Table 1 The onset crystallization (Tx) and peak (Tpi) temperatures gained from the DSC experiment. Composition
Tx (K)
Al0.86Ni0.09La0.05 (Al0.86Ni0.09La0.05)99.5In0.5 (Al0.86Ni0.09La0.05)97In3 (Al0.86Ni0.09La0.05)95In5
512.4 475.7 Not clear 435.9
Tp1/Tp1, Tp2, Tp3 (K)
446.4 447.1
Fig. 4. In situ studies of the melt-spun (Al0.86Ni0.09La0.05)95In5 ribbon at different temperature by high-energy XRD. The small angle area is enlarged in (b). With rising temperature, the diffraction peaks, which belong to crystal In, become wide, even disappearing, indicating the melting behavior of the In nano-particles.
along with a tiny peak corresponding to small amount of crystalline Al. The occurrence of Al-ordering structure, such as pre-existing nuclei [33] or medium-range ordering
Tp2/Tp4 (K) 518.6 504.5 465.7 466.0
493.5 496.2
591.4 588.5 572.4 561.4
[34,35], were previously reported in Al-based amorphous alloys due to their low glass forming ability and high Al content. This kind of tiny and trace Al-ordering structure is rarely detected using commonly used XRD and TEM, as shown in Figs. 1 and 2. With increasing temperature, the diffraction peaks corresponding to crystalline In become wider and weaker. When the temperature reached 410 K, peak 2 of In began to disappear. At 425 K, the 1–3 peaks all disappeared. After heating to 425 K, the sample was cooled down to 300 K, and all peaks that disappeared after heating to 425 K come out again (the top curve shown in Fig. 4). The dependence of the intensity of the diffraction peaks for In with the temperature described above reveals that the melting took place for In particles [36]. Fig. 5 shows the dark field TEM micrographs of the same area of melt-spun (Al0.86Ni0.09La0.05)95In5 specimen in situ taken at room temperature, 393 K, 403 K and 413 K, respectively. The particles are marked 1 to 6. The small particles 5 and 6 melt at the lower temperature, indicated by the disappearance of the corresponding contrast as shown in Fig. 5b [37]. However, the big particles 1 and 2 do not melt completely until 413 K, although they present some melting features in Fig. 5d. The in situ investigations
Fig. 5. Dark field TEM micrographs of the same area of melt-spun (Al0.86Ni0.09La0.05)95In5 specimen in situ taken at (a) room temperature, (b) 393 K, (c) 403 K and (d) 413 K, respectively.
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are completely in a good agreement with the endothermic event in the range of 400–419 K, indicated by the DSC measurement (Fig. 3). These results surely indicate that the melting took place for In particles in the temperature range of 400–419 K, which is 10–30 K lower than the melting temperature T0 (429 K) [38] of the In bulk sample. The melting behaviors of NC In particles embedded in the amorphous matrix as shown in the present study are very different from the superheating of NC In particles embedded in the crystalline matrix [30,37,39]. The different interfacial structures between NC In/amorphous and NC In/ crystalline are thought to be responsible for these different melting behaviors of NC In. The mechanism underlying these unique phenomena will be discussed in Section 4.1 in detail.
Fig. 6a, the amorphous primarily precipitated fcc-Al and then decomposed into fcc-Al, Al3Ni and Al11La3. With the content of In higher than 3%, which leads to the formation of NC In particles, the first three peaks correspond to the formation of fcc-Al, as shown in Fig. 6b, and the last peak reaction products are the same as those of the x = 0 and 0.5 alloys. That is to say, the formation of NC In particles changes the mechanism of fcc-Al precipitation and leads to the loss of the thermal stability of the amorphous matrix. Fig. 7 shows Kissinger plots of the peak temperatures for the present alloys. The apparent activation energies Epi for crystallization reaction could be determined by Kissinger’s equation [40]: ln
3.4. Crystallization of amorphous matrix As stated in Section 3.2, the thermal behaviors were distinctly altered by the formation of NC In particles. In this part, the crystallization behaviors were studied in detail. Fig. 6 shows XRD patterns of the samples annealed to each peak temperature at a heating rate of 10 K min1. With minor addition of In (x = 0.5), which does not result in the formation of NC In particle, the crystallization behaviors were not tuned, as shown in Fig. 3a. In
Fig. 6. XRD patterns of the melt-spun (Al0.86Ni0.09La0.05)(100x)Inx (x = 0, 0.5, 3 and 5) ribbons after heating up to every crystallization peak temperature with 10 K min1.
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T 2pi Epi ¼ þC / k B T pi
ð1Þ
where / is the heating rate, kB is Boltzmann’s constant, Tpi is the ith peak temperature and C is a constant. The measured values can be linear-fitted well and accordingly Epi is calculated, as shown in Fig. 7. From Fig. 7a, it is seen that In addition decreases the activation energy of the fccAl precipitation, especially that for the x = 5 alloy, the activation energy of the first peak reaction is much lowered. The advanced melting of NC In particles as described in
Fig. 7. The Kissinger plots of crystallization peaks’ temperatures for the (Al0.86Ni0.09La0.05)(100x)Inx (x = 0, 0.5 and 5) ribbons.
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Section 3.3 may be responsible for this behavior. It suggests that the interaction of NC In and the amorphous matrix strongly decrease the thermal stability of the amorphous matrix. The primary precipitation of fcc-Al destabilizes the amorphous matrix. As displayed in Fig. 7b, the activation energy of the reaction of fcc-Al, Al3Ni and Al11La3, likewise reduces with increasing the In content. Multistage formation of fcc-Al in (Al0.86Ni0.09La0.05)95In5 system has attracted our interest. The sample which was heated to the first crystallization peak temperature and cooled down was examined by TEM. Fig. 8a shows the shape of the NC In particles is not a regular sphere. The interface between the In particle and the amorphous matrix is misty. Some new phase appears around the In spheres. However, the matrix far away from the In particles is still in featureless amorphous contrast. A high-resolution TEM image exhibits the interfacial structure of the system. As shown in Fig. 8b, there is indeed some new phase between the In particles and the amorphous matrix. It is confirmed to be crystalline Al through the Fourier transform. These results illustrate that the first stage of crystallization is precipitation of fcc-Al around the In particles. 4. Discussion Melting and crystallization are generally both nucleation and growth processes, which are largely influenced by the fluctuation of structure, energy and even composition, etc., in the materials. For BMGCs, the interface is very crucial to the distribution of such factors. Undoubtedly, the interface plays an important role in the thermal properties of the materials. In the following section, the influence of crystal/amorphous interface on the melting of the In particles is firstly discussed, and the influence of the resultant liquid/amorphous interface on the crystallization of the amorphous matrix is be discussed in detail. 4.1. Melting behaviors of NC In particles The melting of embedded NC particles has been widely studied. The superheating [30,31,37,39,41–43] and depression [28,29,42,44–47] of melting temperature have been both reported. Different theories were employed to interpret these phenomena [28,30,31,37,39,42,47]. Literatures on studying the melting of NC particle reported that for most of the rapid solidification systems, the NC particles embedded in crystalline matrix can be easily superheated [30,37,41,43,48,49]. However, when the matrix becomes amorphous, melting temperature depression is observed for NC In particles in this work. These results imply that the matrix state will influence the melting behaviors of the embedded NC particles. Herein, it is thought that the interfacial structure determines the melting behaviors of NC particle. According to the thermodynamic model [50], the melting temperature Tme of embedded particles can be simply expressed as
Fig. 8. (a) TEM image and (b) HRTEM image of (Al86Ni9La5)0.95In5 melt-spun ribbons annealed at the first crystalline peak temperature.
T me ðrÞ ¼ T 0
i 3 h 2=3 1 csm clm ðqs =ql Þ qs rLm
ð2Þ
where T0 is the melting temperature of bulk crystalline phase, r is the radius of the particle, Lm is the latent heat of melting, csm and clm are the interfacial energies between the solid/matrix and liquid/matrix, respectively, qs and ql are the densities of solid and liquid, respectively. From Eq. (2) it can be inferred that with supposing qs ql, the change in Tme depends on csm and clm. When the value of csm is larger than that of clm, the melting temperature of the embedded NC particles will be depressed. Contrarily, the embedded NC particles are superheated with csm smaller than clm, which is deducted in the most of
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melt-spinning NC-particle/crystalline-matrix systems [30,41]. In our work, considering the atomic structure of metallic glass is similar to that of liquid [51], it is expected that csm > clm for the In spheres embedded in the amorphous matrix system. To calculate the interfacial energies between solid/ matrix and liquid/matrix, a simple configuration model of the interfaces between both single component systems, which is modified based on Spaepon’s model, is proposed (the details on the model and the calculation are described in Supplementary material). Since the model is ideal and does not contain mismatch and disordered area, it only fits the well-matched interface in the rapid solidification system. The calculation is under the viewpoint that the interfacial energy is the summation of the binding energy [52]. The energy for different interfaces is estimated with the energy caused by phase transition and the mixing of the two elements. Thus the interfacial energy is calculated as follows:
1=3
3 1 16 N 0 V 2 5 5 DH Am þ DH Bm þ 2DH mix AB 2 2
cA=B cc ¼
ð3Þ
(ii) Interfacial energy of crystalline A/liquid B
A=B ccl
1=3 3 1 ¼ ð5DH Am þ 2DH mix AB Þ 16 N 0 V 2
ð4Þ
(iii) Interfacial energy of amorphous A/crystalline B
cA=B ac ¼
1=3 3 1 5 B mix DH þ 2DH m AB 16 N 0 V 2 2
ð5Þ
(iv) Interfacial energy of amorphous A/liquid B
A=B cal
1=3 3 1 ¼ 2DH mix AB 16 N 0 V 2
In/Al matrix [30], NC–Pb/Al matrix [43], NC Pb/Cu matrix [43] and NC Ag/Ni matrix [41] systems, which coincide with the qualitative deduction in previous literatures. Contrarily, ccl < ccc occurs in the melt-spun NC Bi/ Zn matrix system, where melting temperature depression is investigated for NC Bi particles [47]. However, since the model is ideal, and ignores many factors, there are still some cases that could not be explained, such as no superheating of the NC Pb embedded in Ni matrix [43] and premelting of the NC particles dispersed at the grain boundary [28], which are ascribed to the disordered structure of the interface. In our work, for the In NC particles and Al-based amorphous matrix system, Al is A and In is B. Because of the high content of Al in Al-based amorphous alloy, we consider the amorphous matrix contains only Al element for simplicity. According to the Eqs. (5) and (6), the calculated results are cac ¼ 0:098 J=m2 cal ¼ 0:067 J=m2
(i) Interfacial energy of crystalline A/crystalline B
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ð6Þ
where V is the molar volume of element, DH X m is the enthalpy of the melting process, DH mix is the enthalpy of AB mixing A and B, N0 is the Avogadro constant. Eqs. (3) and (4) are applied to the NC particles embedded in crystallization matrix systems. The calculated results (shown in Supplementary material) indicate that the value of ccl is indeed larger than that of ccc in the melt-spun NC
Then the calculated interfacial energies are put into Eq. (2), and it indicates that the melting temperature of the particles embedded in the amorphous matrix is depressed. The calculation results coincide with the expectation and phenomena stated above. Thus the interfacial energy induces the melting temperature depression. The size of NC particles is another important factor influencing the melting behaviors. From Eq. (2), it is inferred that the smaller the size of NC particles, the higher the deviation of the melting temperature of NC particles from that of bulk samples. It is evidenced in the works of the melt-spun NC In/Al matrix [30], NC–Pb/Al matrix [37] and NC–Ag/Ni matrix [41] systems. In the present work, it works very well. As observed in the in situ TEM observations in Fig. 5, it clearly shows that NC In with the smaller size melts at the lower temperature. However, the quantitative relationship between the size and the melting temperature cannot be presented in this work, which is ongoing. In addition, other factors influencing the melting behaviors of NC particles embedded in amorphous matrix are argued. The first one is the alloying effect. Alloying significantly influences the melting temperature of the system. In the present work, Al–In is an immiscibility system [38]. EDS results show the NC In is pure, indicating extremely limited content of La or Ni might dissolve into In within an error. Even if the elements dissolve into the In spheres, it only can cause a 1 K change in melting temperature for the bulk materials [38]. This cannot explain our experimental observation on the change in melting temperature. So solution of other elements cannot be responsible for the depression of In NC particles as evidenced in our work. The second one is the effect of the disordered region of the interface. Zhang and Cantor [37] considered that melting behaviors of the embedded NC particles are associated
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with the density of defects at the interface, which is the preferential place for melting nucleation. According to the study of Chen and Chuang [53], when the only existing phases boundaries are those between the crystal and the glass, hardly any defects are detected at the interface. Therefore the nucleation of melting may be more difficult in the spherical In phase embedded in amorphous matrix. Thus the disordered region of the interface from the reasons for the depression of the melting temperature in this work can be excluded. The third one is the strain energy effect. The strain energy usually arises from the mismatch in coefficients of thermal expansion or volume expansion during melting. Such a strain energy can elevate the melting temperatures of the embedded particles [41,54]. It should be considered in the NC-particle/crystalline-matrix system rather than our present system. 4.2. Crystallization of the amorphous matrix For the Al–La–Ni amorphous system [55], the core–shell structure occurs during primary crystallization. La-rich shells around the Al core build up when La atoms are rejected from the growing crystals of almost pure Al. These shells hinder further growth and, thus, limit the size of the Al nanocrystals, stabilizing the amorphous matrix. In the present study, as stated in Section 3.4, the addition of In damages the thermal stability of the amorphous matrix, leading to the change of the crystallization behaviors. From the structural analysis (Section 3.1), such behaviors may result from two contributions, one of which is the interface between In and amorphous matrix, the other is the dissolved In. As discussed in Section 3.3, the NC In particles melt prior to the crystallization of the amorphous matrix during continuous heating process. Apparently, the resultant liquid In/amorphous has a more significant effect than the solid/solid interface [5,26,27] on the crystallization behaviors. Because the value of the diffusion coefficient in liquid is several magnitudes higher than that in solid, the element is more easily diffused from the amorphous matrix to the In through this interface. According to Fig. 6b, the peaks of In, especially the main peak at 33°, shift to the small angle gradually. It means there are some La dissolved into In, since in Al–Ni–La–In system, only the radius of La atom is larger than In (The radius of Al, Ni, La and In are 0.14317 [56], 0.12459 [56], 0.18790 [56] and 0.16590 nm [57], respectively), and La is the only one that has negative mixing enthalpy with In in this system [38]. When La is dissolved into In liquid, the matrix around the In nano-particles is poor in La, which will delay the formation of La-rich shells and make the primary precipitation of fcc-Al in the zone proceed more easily. Correspondingly, the activation energy is largely reduced, as shown in Fig. 7. This conclusion is also consistent with the TEM observations in Fig. 8, in which the fcc-Al was indeed observed near the interface. Furthermore, since the melting of the particles occurs in a
temperature range of 20 K, the crystallization of the matrix around the smaller particles occurs at the same time with the melting process of the larger particles. Thus, the melting process of the particles is not an absolutely separated step from the crystallization of the matrix. Different from the diffusion induced crystallization around the NC particles, the crystallization of the matrix far away from the NC particles should be mainly under the influence of the dissolved In. For the x = 0.5 alloy, the melt-spun ribbon is in amorphous phase (as shown in Fig. 1) and the crystallization products are the same as the alloy without In (see Fig. 6). Thus the affect of dissolved In can be studied in this alloy. In Figs. 3 and 7, the dissolved In slightly decreases the temperature of fccAl precipitation. It is probably owing to the positive enthalpy of Al–In, which can reduce the pack density of the amorphous phase and make the atomic diffusion undertake more easily [Q1]. The effect of the so-called core–shell structure undermined [Q2]. Similar cases occur in the Zr– Ni–Cu–Al system with Ti partially replacing Zr [58] and Fe–P–C or Fe–Si–B systems with Cu partially replacing Fe [59,60]. There is still another factor that cannot be ignored. Under the free condition, the melting of In sphere must cause the increase of volume; however, in this system the amorphous matrix prevents this change. It brings the compressing stress field to both the liquid and the matrix around the sphere. According to Ref. [61], the applied pressure enhances the precipitation of the fcc-Al from the amorphous phase. Once the NC In particles are formed, these effects destabilize the amorphous matrix and obviously result in the change of the crystallization behaviors. 5. Conclusions A homogeneous dispersion of small spherical In particles within an amorphous Al–Ni–La matrix has been produced by melt-spinning method. The interaction between the NC particles and the amorphous matrix through the interface is studied. Melting temperature depression of the In NC particles and multistage crystallization of the amorphous matrix have been observed. Melting temperature depression is attributed to the specified structure at interface between the NC metal and amorphous matrix. A simple physical model considering the detailed atomic configuration at interface between NC metal and amorphous phase is used to elucidate the interface-mediated melting phenomenon. The interfacial energy of the crystal/glass state, which is higher than that of liquid/amorphous matrix, provides a driving force to promote the melting of NC metallic particles. The crystallization of the (Al0.86Ni0.09La0.05)95In5 composite starts soon after the melting of the In particles. Under the influence of the diffusion and the compressing stress, fcc-Al firstly precipitates from the matrix around the In particles. Then the primary crystallization of the matrix far away from the NC particles occurs. Finally, as
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the temperature goes up, the rest of the matrix turns into crystalline phase totally. Acknowledgments PJ Shang and ZQ Liu in Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences are appreciated for performing the in situ TEM experiments. The authors gratefully acknowledge the financial support from the Ministry of Science and Technology of China (Grants No. 2006CB605201) and the National Natural Science Foundation of China (Grant Nos. 50825402 and 50725102). Use of the Advanced Photon Source was supported by the US Department of Energy, Office of Science, Office of Basic Energy Science, under Contract No. DE-AC02-06CH11357. Appendix A. Supplementary material Supplementary data associated with this article can be found, in the online version, at doi:10.1016/j.actamat. 2010.07.048. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19]
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