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Solidification microstructures of the undercooled Co–24 at%Sn eutectic alloy containing 0.5 at%Mn
Materials & Design 83 (2015) 138–143
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Solidification microstructures of the undercooled Co–24 at%Sn eutectic alloy containing 0.5 at%Mn Xiaoli Ma ⇑, Li Liu School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China
a r t i c l e
i n f o
Article history: Received 20 January 2015 Revised 13 March 2015 Accepted 4 June 2015 Available online 19 June 2015 Keywords: Co–24 at%Sn eutectic alloy 0.5 at%Mn Undercooling Growth interface morphology Interface energy anisotropy Anomalous eutectic
a b s t r a c t Solidification samples of undercooled Co–24 at%Sn eutectic alloy containing a small amount of Mn (<1.0 at%) were prepared by the glass fluxing technique. The surface and internal solidification microstructures of the samples were observed by a scanning electron microscope (SEM) and an optical microscope (OM), respectively. The experiment results revealed that the addition of 0.5 at%Mn remarkably changed the solidification behaviors of the undercooled Co–24 at%Sn eutectic alloy. The addition of 0.5 at%Mn influenced the morphological selection of eutectic growth interface by increasing the interface energy anisotropy during the solidification of the undercooled Co–24 at%Sn eutectic melt. As undercooling increases, the coupled eutectic growth interface morphology successively experienced dendritic pattern, factual seaweed pattern and compact seaweed pattern. Besides, the addition of 0.5 at%Mn decreased the critical undercooling for the formation of anomalous eutectic by introducing a new formation mechanism of anomalous eutectic, i.e. divorce eutectic mechanism. Ó 2015 Elsevier Ltd. All rights reserved.
1. Introduction The solidification behaviors of undercooled eutectic alloys, especially binary eutectic alloys, have been widely studied through experimental methods [1–9] and theoretical models, such as the JH model [1], the TMK model [2], and the LZ model [3]. Various solidification behaviors, including anomalous eutectic, eutectic dendrite, and growth velocity, were focused in these investigations. However, other elements were added into binary eutectic alloys to improve their properties in practices. Eutectic solidification behaviors can be influenced by the third element addition. Yamauchi et al. [10] reported that the critical growth rate for the transition from lamellar to rod eutectics decreased as the content of the third element increased. Li et al. [11] found that even low content of Dy in Ni–Al eutectic alloys could break the co-operative growth of two eutectic phases. It was well known that the additions of P into the Al–Si alloy can increase the nucleation frequency of eutectic grains due to the formation of AlP particles acting as nucleation sites for eutectic silicon [12]. When a third element is added to a binary eutectic alloy melt, an additional solute diffusion field forms ahead of the solidification interface, and thus imposes obvious effects on the solidification behaviors. In the previous work [13], a theoretical model for lamellar eutectic growth in
⇑ Corresponding author. E-mail address: [email protected] (X. Ma). http://dx.doi.org/10.1016/j.matdes.2015.06.010 0264-1275/Ó 2015 Elsevier Ltd. All rights reserved.
an undercooled eutectic alloy melt containing a third element had been developed and analyzed the effects of a third element on the solidification behaviors in detail, such as eutectic lamellar spacing, growth velocity and eutectic dendrite tip radius. However, precise experimental evidences were still insufficient to verify these solidification behaviors deduced by the theoretical analysis up to now. The solidification behaviors of undercooled Co–24 at%Sn eutectic alloys were investigated in another related publication [14], in which both fractal and compact eutectic seaweed microstructures were observed. To study the effects of a third element on the solidification behaviors of undercooled Co–24 at%Sn eutectic alloys various elements, such as Mn, Ni, Cu, Al, Si, Sb, were added into the Co–24 at%Sn eutectic alloy in the early works. It had been found that the solidification behavior changes of undercooled Co–24 at%Sn eutectic alloy depend on the element added. In this paper, the morphology changes of coupled eutectic growth interface were the mainly studied solidification phenomenon. Therefore, the element added should not break the coupled eutectic growth of a-Co phase and b-Co3Sn2 phase. Besides, the interface morphology transformation between dendritic pattern and seaweed pattern aroused extreme interest. Among the various elements added, only a small addition of Mn could satisfy these requirements. Therefore, Mn was selected to be the third element added into Co–24 at%Sn eutectic alloys. In the present work, the solidification behaviors of undercooled Co–24 at%Sn eutectic alloys containing a small amount of Mn were studied to explore the
X. Ma, L. Liu / Materials & Design 83 (2015) 138–143
influence of Mn on the eutectic growth interface morphology, the growth velocity and the anomalous eutectic microstructure. 2. Materials and experimental procedures The Co–24.0 at%Sn eutectic alloys containing a small amount of Mn (less than 1.0 at%) were prepared by the glass fluxing technique. The Co–24.0 at%Sn eutectic alloy samples containing 0.5 at%Mn were cooled to different undercooling before solidification. In this experiment high purity components are used, i.e. 99.99%Co, 99.999%Sn and 99.99Mn%. The surface and internal solidification microstructures of the samples were observed by a scanning electron microscope (SEM) and an optical microscope (OM), respectively. In the solidification of undercooled Co–24 at%Sn eutectic alloys fractal eutectic seaweed microstructures were observed. In this work, the fractal dimension of the solidification microstructure needed to be measured to confirm whether the eutectic growth interface morphology is fractal or not. A method called box-dimension calculation was used. When the method is used, the smallest number of boxes of side length l needed to cover the shape is counted for different side length l. The expression is,
dbox ¼ lim l!0
log NðlÞ logð1=lÞ
ð1Þ
where dbox is the box-dimension; N(l) is the number of boxes of side length e to cover the shape. If the limit of Eq. (1) exists, and it is not an integer, the shape is fractal. The box-dimension calculation can be divided into two steps. First, we depicted the contour profile of the coupled eutectic growth interface in the metallograph. Then the fractal dimension of the contour profile was measured by box-dimension calculation software. The crystal growth velocities were measured by the dual-probe method. The schematic diagram of the dual-probe method is shown in Fig. 1. A long sample of /8 30 mm was solidified in a cylindrical quartz glass crucible, followed by the same solidification procedure as described above. When the melt was undercooled to the predetermined degree, solidification was triggered using a thin Al2O3 needle at the center of the top surface of the
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melt column. Two identical infrared pyrometers were placed along the axial direction of the sample with a distance of 20 mm (as so-called double probe method). The growth velocity was evaluated by dividing the distance difference from the triggering point to the respective view-points of the two pyrometers at the sample surface by the time interval between the moments at which the recalescence front reaches the respective viewpoints. The solute partition coefficients of Mn in Co–Sn alloys were determined by the rapid-quenching experiment. Samples of Co–18 at%Sn hypo-eutectic and Co–30 at%Sn hyper-eutectic alloy containing 1.0 at%Mn addition were rapidly quenched from a coexisting solid–liquid state after 1 h holding. An area close to the surface of the quenched sample was selected, in which the distributions of Mn in the two eutectic phases were measured by the energy dispersive X-ray spectroscopy (EDX). The measured results were used to estimate the solute partitioning coefficients of Mn. The phase transition temperatures of the Co–24.0 at%Sn eutectic and Co–10 at%Sn hypo-eutectic alloys with and without 0.5 at%Mn addition were measured by a differential scanning calorimetry (DSC) instrument. The specimens used for this measurement were cut from the solidified samples, and were small round sheets with a diameter of 5 mm and height of 1 mm. The heating rate and cooling rate were set as 30 K/min. The DSC results could be used to evaluate the fusion enthalpy. The fusion heat produced by the phase transition can be evaluated by calculating the area of the phase transition peak on the DSC curve. The fusion enthalpy of the Co–24 at%Sn eutectic alloy equals to the fusion enthalpy summation of the a-Co phase and b-Co3Sn2 phase. Based on the Co–Sn binary phase diagram, the proportions of the a-Co phase and b-Co3Sn2 phase can be obtained. The corporeal quantity of primary a-Co phase is about 63.6% of the whole Co–10 at%Sn alloy. The corporeal quantity ratio of the a-Co phase to the b-Co3Sn2 phase is about 0.55 in the Co–24.0 at%Sn eutectic alloy. Combining with the quantity of the sample, the corporeal quantity of different phases can be obtained. The fusion enthalpy equaled to the ratio of fusion heat to the corporeal quantity of the sample. In the present work, the fusion enthalpy of the a-Co phase, b-Co3Sn2 phase and Co–24.0 at%Sn eutectic alloy with and without Mn addition were evaluated. 3. Results 3.1. As-cast solidification microstructures Fig. 2 shows the cross-sectional as-cast microstructures of Co–24 at%Sn eutectic alloys containing Mn from 0 to 1.0 at%. With different Mn contents, the solidification microstructures show various features. The pure Co–24 at%Sn eutectic alloy exhibits eutectic seaweed microstructures (Fig. 2a). After adding 0.5 at%Mn, the as-cast microstructure transforms from eutectic seaweed to eutectic dendrite (Fig. 2b). When the Mn content increases up to 0.7 at%, the eutectic dendrites become very irregular (Fig. 2c). As shown in Fig. 2c, the primary arms look like a series of continuous bamboo joints, with only one a-Co phase central lamellar and two b-Co3Sn2 lamellae on its two sides. Some eutectic cells distribute around primary eutectic dendrite arms. Once the Mn addition increases up to 1.0 at%, primary b-Co3Sn2 phase forms in the solidification microstructures. Obviously, the as-cast solidification microstructures of Co–24 at%Sn eutectic alloy are very sensitive to the addition amount of Mn. 3.2. Growth interface morphology
Fig. 1. Schematic diagram of the dual-probe method for growth velocity measurement. (a) Measurement setting and (b) temperature recalescence curves.
The SEM surface images of the Co–24 at%Sn samples containing 0.5 at%Mn solidified at different undercoolings are presented in
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Fig. 2. Cross-sectional as-cast microstructures of Co–24 at%Sn eutectic alloy with different Mn additions. (a) 0 at%Mn; (b) 0.5 at%Mn; (c) 0.7 at%Mn; (d) 1.0 at%Mn.
Fig. 3, in which the crystal growth direction is from right to left. As shown in Fig. 3a, the solidification microstructure exhibits eutectic seaweed morphology at the undercooling of 55 K, which is similar with that in the solidification of pure Co–24 at%Sn undercooled melt at low undercooling [14]. The result of box-dimension calculation shows that the average dimension of eutectic seaweed microstructure is about 1.7. Therefore, the eutectic seaweed microstructure shown in Fig. 3a is fractal. Obviously, the solidification microstructures transform from eutectic dendrite to fractal eutectic seaweed as the undercooling increases up to 55 K. The eutectic growth interface keeps fractal seaweed morphology up to the undercooling of 182 K (Fig. 3b). As the undercooling rises to 185 K, the coupled eutectic growth interface morphology changes greatly, i.e. the two branches grow relatively independent after tip-splitting and symmetrically distributed short side-branches form behind the their tips (Fig. 3c). Such a solidification microstructure feature is very similar to that in the solidification of pure Co–24 at%Sn undercooled melt with a undercooling of 175 K [14]. According to the box-dimension calculation result, the solidification microstructures shown in Fig. 3c exhibit the morphology of compact eutectic seaweed pattern. This indicates that the addition of 0.5 at%Mn into the Co–24 at%Sn eutectic alloy enlarges the critical undercooling to 185 K for the growth interface morphology transformation from fractal seaweed to compacted seaweed. 3.3. Anomalous eutectics Anomalous eutectic microstructures of the Co–24 at%Sn eutectic alloy containing 0.5 at%Mn undercooled at different undercoolings are shown in Fig. 4. In the previous works, it had already been found that the anomalous eutectic firstly forms in the solidification of Co–24 at%Sn alloys undercooled by 21 K [14]. The anomalous eutectics can be observed in the sample undercooled by only 10 K (Fig. 4a). Obviously, the addition of 0.5 at%Mn decreases the critical undercooling for the formation of anomalous eutectics in the Co–24 at%Sn eutectic alloys. Meanwhile, it should be noted that the first formed anomalous eutectics mainly exist in the
region between eutectic branches, rather than the interior of eutectic branches in the Co–24 at%Sn eutectic alloy. When undercooling increases to about 20 K, another type of anomalous eutectic forms inside the eutectic dendrites coexisting with the first type of anomalous eutectic. The dimensions of a-Co particles in the two types of anomalous eutectics are remarkably different (Fig. 4b). As undercooling increases further, e.g. at 120 K, only one type of anomalous eutectic is observed (Fig. 4c), which has the same behavior with that in the pure Co–24 at%Sn eutectic alloy. 3.4. Growth velocity The measured solidification velocities of Co–24 at%Sn eutectic alloys with and without 0.5 at%Mn at different undercoolings are listed in Fig. 5. The addition of 0.5 at%Mn does not change the overall variation trend of the solidification velocity as undercooling increases, but increases the critical undercooling for the sudden rapid rise of solidification velocity from 175 K to 185 K. Below 175 K, the addition of 0.5 at%Mn hardly change the solidification velocity. Above 175 K, however, 0.5 at%Mn greatly decreases the growth velocity of Co–24 at%Sn alloys. 3.5. Solute partitioning coefficient In the present work, the solute partitioning coefficient of Mn in the a-Co, ka, and b-Co3Sn2 phases, kb, were measured through the rapid-quenching experiments of Co–10 at%Sn hypoeutectic and Co–30 at%Sn hypereutectic alloys containing 1.0 at%Mn. The experimental results reveal that ka and kb are 0.57 and 1.2, respectively (Table 1). Calculated according to the results, the effective solute partitioning of Mn in the solidification of Co–24 at%Sn eutectic alloy, keff, is 0.94, which is slightly less than 1. 3.6. DSC curves and fusion enthalpy Fig. 6 shows the DSC heating curves of Co–Sn alloys with different compositions, including Co–24 at%Sn, Co–24 at%–0.5 at%Mn, Co–10 at%Sn and Co–10 at%Sn–0.5 at%Mn. For Co–24 at%Sn and
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Fig. 3. SEM surface images of the Co–24 at%Sn samples with 0.5 at%Mn addition solidified at undercooling. (a) 55 K; (b) 182 K and (c) 185 K, respectively.
Fig. 4. Anomalous eutectics of the Co–24 at%Sn alloy with 0.5 at%Mn addition undercooled by (a) 10 K; (b) 20 K; (c) 120 K.
Fig. 5. The measured solidification velocities of Co–24 at%Sn eutectic alloys with and without 0.5 at%Mn alloys at different undercoolings.
Co–24 at%–0.5 at%Mn alloys, only one endothermic peak exists on the heating curve, corresponding to the fusion of the eutectic microstructure. There are two endothermic peaks on the heating curve of Co–10 at%Sn alloy with and without 0.5 at%Mn addition. The two peaks correspond to the fusion of eutectic microstructure and the single a-Co phase, respectively. It can be obtained from the DSC curves that the eutectic temperature Teutectic of Co–24 at%Sn alloy containing 0.5 at%Mn is 1387 K, being 2 K larger than that of pure Co–24 at%Sn alloy, i.e. 1385 K. The fusion enthalpies of the eutectic alloy DH, a-Co phase DHa and b-Co3Sn2 phase DHb were evaluated through the DSC results. The fusion enthalpies of Co–24 at%Sn eutectic alloys with and without Mn addition were
Table 1 The experimentally measured solute partitioning coefficient of Mn in the solidification of Co–Sn alloy. The third element
ka
kb
keff
Mn
0.57
1.2
0.94
Fig. 6. DSC heating curves of Co–Sn alloys with different compositions.
listed in Table 2. The addition of 0.5 at%Mn increases the fusion enthalpy of Co–24 at%Sn eutectic alloy and the b-Co3Sn2 phase. Conversely, the fusion enthalpy of the a-Co phase is reduced. 4. Discussions 4.1. Influence of Mn on interface energy anisotropy The addition of a new element will change the interatomic action potential energy of growth interface in the solidification of
Table 2 The fusion enthalpy of Co–24 at%Sn eutectic alloys with and without 0.5 at% Mn addition. The third element
DH (J/mol)
DHa (J/mol)
DHb (J/mol)
No addition 0.5 at%Mn
10,900 12,200
7300 6500
12,400 15,690
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alloy melt, and thus cause the interface energy variation of growth interface [15,16]. In the present paper, alloy thermodynamics is used to evaluate the effects of Mn addition on the interface energy in the solidification of Co–24 at%Sn eutectic alloy. From crystal structure, Spaepen [17] advanced a calculation model of solid/liquid (S/L) interface energy c,
amplitude closely relates with the interface energy anisotropy, and the relation can be expressed as,
F P exp
e7=8
3
P ¼ ðT=T 0 Þ
c ¼ am
DH NA V m 2
1=3
ð2Þ
where NA is Avogadro constant, V m is the average atomic volume, am is a factor related with crystal structure. Since the atomic volume of Mn is close to that of Co, and its addition amount, 0.5 at%, is very small, the influence of Mn addition on the V m is very limited. Besides, the addition of 0.5 at%Mn does not change the crystal structure of a-Co and b-Co3Sn2 phases. According to Eq. (1) the effect of Mn on the interface energy can be evaluated through determining the fusion enthalpy variation of Co–24 at%Sn eutectic alloy caused by the addition of 0.5 at%Mn. As shown in Table 2, the addition of 0.5 at%Mn increases the fusion enthalpy of Co–24 at%Sn eutectic alloy. Based on the Eq. (2), the addition of 0.5 at%Mn enlarges the S/L interface energy in the solidification of Co–24 at%Sn eutectic alloy. Meantime, the addition of Mn reduces a/L interface energy and increases b/L interface energy. Since the fusion entropy of the intermetallic compound b-Co3Sn2 is much higher than that of the a-Co phase and the solid solubility of Mn is larger in the b-Co3Sn2 phase, the addition of 0.5 at%Mn imposes more effects on b/L interface energy than those on a/L interface energy, leading to the enlargement of the coupled eutectic growth interface energy. Since the atomic structures of different crystal planes in one crystal are different, for a curved growth interface, different interface positions have various atomic structures. The addition of a third element will impose different effects on the lattice structures and the interatomic potentials of different crystal planes. Therefore, the variations of interface energy caused by the third element along the curved growth interface tip are different. In the solidification of eutectic alloy with small addition of a third element, the concentration of the third element in the liquid ahead of the growth interface relates with the growth velocity and eutectic lamellar spacing [13]. The normal growth velocities and eutectic lamellar spacing are affected by the position change, resulting in the concentration variation of the third element along a curved growth interface, which also affect the interface energy and its anisotropy. It has already been confirmed that the interface energy anisotropy of eutectic growth interface is very weak in the solidification of undercooled Co–24 at%Sn eutectic alloy [14]. According to the analyses above, the addition of 0.5 at%Mn not only increases the interface energy, but also enlarges the interface energy anisotropy of the coupled eutectic growth interface. 4.2. Influence of Mn on growth morphology pattern The effective partitioning coefficient is slightly less than 1, implying that only small amounts of Mn atoms enrich ahead of the growth interface. Therefore, the solute diffusion field of Mn has little effect on the growth interface morphology. Then, the small addition of Mn mainly influences the growth interface morphology by changing the interface energy and its anisotropy in the solidification of Co–24 at%Sn eutectic. The addition of 0.5 at%Mn strengthens the interface energy anisotropy of the eutectic growth interface. Once the interface energy anisotropy exceeds a critical value, the interface tip is no longer the coldest position along the whole interface, improving the formation of side-branches. Brener et al. [18] proposed that the interference
2ld0
ð3Þ !1=2
R4
ð4Þ
where F is the interference amplitude, P is the relative noise strength, e is the interface energy anisotropy strength, T is the factual temperature, d0 is the capillary length, l is the solute diffusion length, R is the interface tip radius. According to the Eq. (3), the increasing of interface energy anisotropy will decrease the interface interference amplitude when the undercooling keeps constant. This means that the addition of 0.5 at%Mn increases the interface tip stability in the solidification of undercooled Co–24 at%Sn eutectic alloy. Therefore, the addition of 0.5 at%Mn leads to the formation of eutectic dendrite rather than the eutectic seaweed in the solidification of Co–24 at%Sn eutectic alloy at low undercooling. As undercooling increases, the noise strength rises, and the interference amplitude on the growth interface is enlarged correspondingly. When the interference amplitude exceeds a critical value, the stability effect caused by Mn addition is not effective enough to sustain a stable interface tip. So there is a critical undercooling above which the noise interference plays the main role, resulting in the splitting of interface tip. As a result, the eutectic growth interface transform from dendrite to seaweed morphology. When undercooling exceeds a certain value, the crystal grows so fast that the interference on the interface tip rapidly leaves the tip region through tip convection and thus keeps the tip stable in the solidification of undercooled Co–24 at%Sn eutectic melt. In this case, the influence of the third element can be ignored. Those interferences leaving the interface tip lead to the formation of side-branches at the positions behind the tip with a certain distance. Therefore, the eutectic growth interface morphology transforms from factual seaweed to compact seaweed as the undercooling beyond a critical value in the solidification of Co–24 at%Sn eutectic alloy. Brener et al. [18] advanced that the critical undercooling D⁄ for the transition from factual seaweed to compact seaweed relates with the relative noise strength, P, and corresponding law is
D jPj2=5
ð5Þ
Since the value of P is far less than 1 in the undercooled eutectic melt, D⁄ is in direct proportion to P, meaning that the enlargement of P will increase D⁄. Small Mn addition increases the eutectic growth interface energy in the solidification of Co–24 at%Sn eutectic melt, which enlarges the relative noise strength according to Eq. (4), and thus increases the interface interference amplitude. The influence of Mn on the eutectic growth velocity is small, so the growth velocity is not enough to sustain the interface tip stability until the undercooling up to 185 K, at which the seaweed shaped eutectic growth interface transform from factual to compact mode. 4.3. Influence of Mn on the anomalous eutectic In the rapid solidification of eutectic alloy containing a third element, the eutectic lamellar spacing is very small, and thus the enrichment of the third element atom is not enough to break the stability of single eutectic lamellar. As the rapid solidification is to the end, the eutectic lamellar spacing increases gradually, the single eutectic lamellar may split due to the instability effect produced by the enrichment of the third element atoms. These
X. Ma, L. Liu / Materials & Design 83 (2015) 138–143
splitting eutectic lamellae grow into the remaining liquid as branches and thus break the coupled eutectic growth interface. Since ka of Mn is less than 1 and kb of Mn is larger than 1, more Mn atoms enrich ahead of the a/L interface than that of the b/L interface, meaning that the solute undercooling caused by Mn addition in front of the a-Co phase is larger. Therefore, at the end of rapid solidification the a-Co phase first deviates from the eutectic coupled growth interface and grows into the remaining liquid between adjacent eutectic dendrite arms as a single dendrite. This process is so-called the divorced eutectic growth. During the divorced eutectic growth the tip radius of these a-Co dendrites is much larger than the eutectic lamellar spacing. These a-Co dendrites stay in the liquid phase for long time, subjecting to the re-melting and ripening due to the S/L interface tension, and form another type anomalous eutectic finally. Such a formation mechanism of anomalous eutectic can be named divorced eutectic mechanism. Since the undercooling is very low, at 11 K, the primary eutectic lamellar cannot be re-melted into anomalous eutectic. So the anomalous eutectic lies between eutectic dendrite arms (Fig. 4a). When the undercooling is large enough, the primary formed eutectic lamellae during the rapid solidification are re-melted due to overheating and then transform into anomalous eutectic microstructures. As a result, the anomalous eutectics are formed through two mechanisms in the solidification of Co– 24 at%Sn eutectic alloy with 0.5 at%Mn addition. The anomalous eutectic microstructure with large-size a-Co particles, lying between the eutectic seaweed arms resulted from divorced eutectic growth. While the one with small-size a-Co particles, lying in the internal eutectic seaweed arms, are produced by the re-melting and ripening of primary eutectic lamellae (Fig. 4b). As the undercooling increases further, the primary eutectic lamellae grow faster and the regions between eutectic seaweed arms shrink greatly so that the divorced eutectic microstructures disappear. As a result, only one type of anomalous eutectic forms, which are the productions of rapidly solidified primary eutectic lamellae (Fig. 4c). 5. Conclusions 1. The addition of 0.5 at%Mn effectively changes the interface energy and its anisotropy of the coupled eutectic growth interface during the solidification of undercooled Co–24 at%Sn eutectic alloy. 0.5 at%Mn addition improves the eutectic dendrite growth at a low undercooling. As undercooling increases the morphology of eutectic growth interface transforms from dendrite pattern to fractal seaweed pattern. When undercooling exceeds a critical value the fractal seaweed interface morphology change to compact seaweed morphology, just like that happens in the pure Co–24 at%Sn eutectic alloy. The only difference is that the addition of 0.5 at%Mn increases critical transition undercooling from 175 K to 185 K. 2. Below undercooling of about 175 K, the addition of 0.5 at%Mn has little effect on the solidification velocity of Co–24 at%Sn eutectic alloy. Above the undercooling of 175 K, the addition of 0.5 at%Mn obviously decreases the solidification velocity.
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3. The undercooling for the primary formation of anomalous eutectic is 10 K in the solidification of Co–24 at%Sn eutectic alloy containing 0.5 at%Mn, which is 11 K less than the 21 K for that without Mn addition. The addition of 0.5 at%Mn introduces a new formation mechanism of anomalous eutectic, i.e. divorce eutectic mechanism, which makes anomalous eutectic first form at the inter-dendrite region at a very low undercooling. In the intermediate undercooling range, two mechanisms work together for the formation of anomalous eutectic, and only one type of anomalous eutectic microstructure forms at large undercooling.
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Solidification microstructures of the undercooled Co–24 at%Sn eutectic alloy containing 0.5 at%Mn Xiaoli Ma ⇑, Li Liu School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China
a r t i c l e
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Article history: Received 20 January 2015 Revised 13 March 2015 Accepted 4 June 2015 Available online 19 June 2015 Keywords: Co–24 at%Sn eutectic alloy 0.5 at%Mn Undercooling Growth interface morphology Interface energy anisotropy Anomalous eutectic
a b s t r a c t Solidification samples of undercooled Co–24 at%Sn eutectic alloy containing a small amount of Mn (<1.0 at%) were prepared by the glass fluxing technique. The surface and internal solidification microstructures of the samples were observed by a scanning electron microscope (SEM) and an optical microscope (OM), respectively. The experiment results revealed that the addition of 0.5 at%Mn remarkably changed the solidification behaviors of the undercooled Co–24 at%Sn eutectic alloy. The addition of 0.5 at%Mn influenced the morphological selection of eutectic growth interface by increasing the interface energy anisotropy during the solidification of the undercooled Co–24 at%Sn eutectic melt. As undercooling increases, the coupled eutectic growth interface morphology successively experienced dendritic pattern, factual seaweed pattern and compact seaweed pattern. Besides, the addition of 0.5 at%Mn decreased the critical undercooling for the formation of anomalous eutectic by introducing a new formation mechanism of anomalous eutectic, i.e. divorce eutectic mechanism. Ó 2015 Elsevier Ltd. All rights reserved.
1. Introduction The solidification behaviors of undercooled eutectic alloys, especially binary eutectic alloys, have been widely studied through experimental methods [1–9] and theoretical models, such as the JH model [1], the TMK model [2], and the LZ model [3]. Various solidification behaviors, including anomalous eutectic, eutectic dendrite, and growth velocity, were focused in these investigations. However, other elements were added into binary eutectic alloys to improve their properties in practices. Eutectic solidification behaviors can be influenced by the third element addition. Yamauchi et al. [10] reported that the critical growth rate for the transition from lamellar to rod eutectics decreased as the content of the third element increased. Li et al. [11] found that even low content of Dy in Ni–Al eutectic alloys could break the co-operative growth of two eutectic phases. It was well known that the additions of P into the Al–Si alloy can increase the nucleation frequency of eutectic grains due to the formation of AlP particles acting as nucleation sites for eutectic silicon [12]. When a third element is added to a binary eutectic alloy melt, an additional solute diffusion field forms ahead of the solidification interface, and thus imposes obvious effects on the solidification behaviors. In the previous work [13], a theoretical model for lamellar eutectic growth in
⇑ Corresponding author. E-mail address: [email protected] (X. Ma). http://dx.doi.org/10.1016/j.matdes.2015.06.010 0264-1275/Ó 2015 Elsevier Ltd. All rights reserved.
an undercooled eutectic alloy melt containing a third element had been developed and analyzed the effects of a third element on the solidification behaviors in detail, such as eutectic lamellar spacing, growth velocity and eutectic dendrite tip radius. However, precise experimental evidences were still insufficient to verify these solidification behaviors deduced by the theoretical analysis up to now. The solidification behaviors of undercooled Co–24 at%Sn eutectic alloys were investigated in another related publication [14], in which both fractal and compact eutectic seaweed microstructures were observed. To study the effects of a third element on the solidification behaviors of undercooled Co–24 at%Sn eutectic alloys various elements, such as Mn, Ni, Cu, Al, Si, Sb, were added into the Co–24 at%Sn eutectic alloy in the early works. It had been found that the solidification behavior changes of undercooled Co–24 at%Sn eutectic alloy depend on the element added. In this paper, the morphology changes of coupled eutectic growth interface were the mainly studied solidification phenomenon. Therefore, the element added should not break the coupled eutectic growth of a-Co phase and b-Co3Sn2 phase. Besides, the interface morphology transformation between dendritic pattern and seaweed pattern aroused extreme interest. Among the various elements added, only a small addition of Mn could satisfy these requirements. Therefore, Mn was selected to be the third element added into Co–24 at%Sn eutectic alloys. In the present work, the solidification behaviors of undercooled Co–24 at%Sn eutectic alloys containing a small amount of Mn were studied to explore the
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influence of Mn on the eutectic growth interface morphology, the growth velocity and the anomalous eutectic microstructure. 2. Materials and experimental procedures The Co–24.0 at%Sn eutectic alloys containing a small amount of Mn (less than 1.0 at%) were prepared by the glass fluxing technique. The Co–24.0 at%Sn eutectic alloy samples containing 0.5 at%Mn were cooled to different undercooling before solidification. In this experiment high purity components are used, i.e. 99.99%Co, 99.999%Sn and 99.99Mn%. The surface and internal solidification microstructures of the samples were observed by a scanning electron microscope (SEM) and an optical microscope (OM), respectively. In the solidification of undercooled Co–24 at%Sn eutectic alloys fractal eutectic seaweed microstructures were observed. In this work, the fractal dimension of the solidification microstructure needed to be measured to confirm whether the eutectic growth interface morphology is fractal or not. A method called box-dimension calculation was used. When the method is used, the smallest number of boxes of side length l needed to cover the shape is counted for different side length l. The expression is,
dbox ¼ lim l!0
log NðlÞ logð1=lÞ
ð1Þ
where dbox is the box-dimension; N(l) is the number of boxes of side length e to cover the shape. If the limit of Eq. (1) exists, and it is not an integer, the shape is fractal. The box-dimension calculation can be divided into two steps. First, we depicted the contour profile of the coupled eutectic growth interface in the metallograph. Then the fractal dimension of the contour profile was measured by box-dimension calculation software. The crystal growth velocities were measured by the dual-probe method. The schematic diagram of the dual-probe method is shown in Fig. 1. A long sample of /8 30 mm was solidified in a cylindrical quartz glass crucible, followed by the same solidification procedure as described above. When the melt was undercooled to the predetermined degree, solidification was triggered using a thin Al2O3 needle at the center of the top surface of the
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melt column. Two identical infrared pyrometers were placed along the axial direction of the sample with a distance of 20 mm (as so-called double probe method). The growth velocity was evaluated by dividing the distance difference from the triggering point to the respective view-points of the two pyrometers at the sample surface by the time interval between the moments at which the recalescence front reaches the respective viewpoints. The solute partition coefficients of Mn in Co–Sn alloys were determined by the rapid-quenching experiment. Samples of Co–18 at%Sn hypo-eutectic and Co–30 at%Sn hyper-eutectic alloy containing 1.0 at%Mn addition were rapidly quenched from a coexisting solid–liquid state after 1 h holding. An area close to the surface of the quenched sample was selected, in which the distributions of Mn in the two eutectic phases were measured by the energy dispersive X-ray spectroscopy (EDX). The measured results were used to estimate the solute partitioning coefficients of Mn. The phase transition temperatures of the Co–24.0 at%Sn eutectic and Co–10 at%Sn hypo-eutectic alloys with and without 0.5 at%Mn addition were measured by a differential scanning calorimetry (DSC) instrument. The specimens used for this measurement were cut from the solidified samples, and were small round sheets with a diameter of 5 mm and height of 1 mm. The heating rate and cooling rate were set as 30 K/min. The DSC results could be used to evaluate the fusion enthalpy. The fusion heat produced by the phase transition can be evaluated by calculating the area of the phase transition peak on the DSC curve. The fusion enthalpy of the Co–24 at%Sn eutectic alloy equals to the fusion enthalpy summation of the a-Co phase and b-Co3Sn2 phase. Based on the Co–Sn binary phase diagram, the proportions of the a-Co phase and b-Co3Sn2 phase can be obtained. The corporeal quantity of primary a-Co phase is about 63.6% of the whole Co–10 at%Sn alloy. The corporeal quantity ratio of the a-Co phase to the b-Co3Sn2 phase is about 0.55 in the Co–24.0 at%Sn eutectic alloy. Combining with the quantity of the sample, the corporeal quantity of different phases can be obtained. The fusion enthalpy equaled to the ratio of fusion heat to the corporeal quantity of the sample. In the present work, the fusion enthalpy of the a-Co phase, b-Co3Sn2 phase and Co–24.0 at%Sn eutectic alloy with and without Mn addition were evaluated. 3. Results 3.1. As-cast solidification microstructures Fig. 2 shows the cross-sectional as-cast microstructures of Co–24 at%Sn eutectic alloys containing Mn from 0 to 1.0 at%. With different Mn contents, the solidification microstructures show various features. The pure Co–24 at%Sn eutectic alloy exhibits eutectic seaweed microstructures (Fig. 2a). After adding 0.5 at%Mn, the as-cast microstructure transforms from eutectic seaweed to eutectic dendrite (Fig. 2b). When the Mn content increases up to 0.7 at%, the eutectic dendrites become very irregular (Fig. 2c). As shown in Fig. 2c, the primary arms look like a series of continuous bamboo joints, with only one a-Co phase central lamellar and two b-Co3Sn2 lamellae on its two sides. Some eutectic cells distribute around primary eutectic dendrite arms. Once the Mn addition increases up to 1.0 at%, primary b-Co3Sn2 phase forms in the solidification microstructures. Obviously, the as-cast solidification microstructures of Co–24 at%Sn eutectic alloy are very sensitive to the addition amount of Mn. 3.2. Growth interface morphology
Fig. 1. Schematic diagram of the dual-probe method for growth velocity measurement. (a) Measurement setting and (b) temperature recalescence curves.
The SEM surface images of the Co–24 at%Sn samples containing 0.5 at%Mn solidified at different undercoolings are presented in
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Fig. 2. Cross-sectional as-cast microstructures of Co–24 at%Sn eutectic alloy with different Mn additions. (a) 0 at%Mn; (b) 0.5 at%Mn; (c) 0.7 at%Mn; (d) 1.0 at%Mn.
Fig. 3, in which the crystal growth direction is from right to left. As shown in Fig. 3a, the solidification microstructure exhibits eutectic seaweed morphology at the undercooling of 55 K, which is similar with that in the solidification of pure Co–24 at%Sn undercooled melt at low undercooling [14]. The result of box-dimension calculation shows that the average dimension of eutectic seaweed microstructure is about 1.7. Therefore, the eutectic seaweed microstructure shown in Fig. 3a is fractal. Obviously, the solidification microstructures transform from eutectic dendrite to fractal eutectic seaweed as the undercooling increases up to 55 K. The eutectic growth interface keeps fractal seaweed morphology up to the undercooling of 182 K (Fig. 3b). As the undercooling rises to 185 K, the coupled eutectic growth interface morphology changes greatly, i.e. the two branches grow relatively independent after tip-splitting and symmetrically distributed short side-branches form behind the their tips (Fig. 3c). Such a solidification microstructure feature is very similar to that in the solidification of pure Co–24 at%Sn undercooled melt with a undercooling of 175 K [14]. According to the box-dimension calculation result, the solidification microstructures shown in Fig. 3c exhibit the morphology of compact eutectic seaweed pattern. This indicates that the addition of 0.5 at%Mn into the Co–24 at%Sn eutectic alloy enlarges the critical undercooling to 185 K for the growth interface morphology transformation from fractal seaweed to compacted seaweed. 3.3. Anomalous eutectics Anomalous eutectic microstructures of the Co–24 at%Sn eutectic alloy containing 0.5 at%Mn undercooled at different undercoolings are shown in Fig. 4. In the previous works, it had already been found that the anomalous eutectic firstly forms in the solidification of Co–24 at%Sn alloys undercooled by 21 K [14]. The anomalous eutectics can be observed in the sample undercooled by only 10 K (Fig. 4a). Obviously, the addition of 0.5 at%Mn decreases the critical undercooling for the formation of anomalous eutectics in the Co–24 at%Sn eutectic alloys. Meanwhile, it should be noted that the first formed anomalous eutectics mainly exist in the
region between eutectic branches, rather than the interior of eutectic branches in the Co–24 at%Sn eutectic alloy. When undercooling increases to about 20 K, another type of anomalous eutectic forms inside the eutectic dendrites coexisting with the first type of anomalous eutectic. The dimensions of a-Co particles in the two types of anomalous eutectics are remarkably different (Fig. 4b). As undercooling increases further, e.g. at 120 K, only one type of anomalous eutectic is observed (Fig. 4c), which has the same behavior with that in the pure Co–24 at%Sn eutectic alloy. 3.4. Growth velocity The measured solidification velocities of Co–24 at%Sn eutectic alloys with and without 0.5 at%Mn at different undercoolings are listed in Fig. 5. The addition of 0.5 at%Mn does not change the overall variation trend of the solidification velocity as undercooling increases, but increases the critical undercooling for the sudden rapid rise of solidification velocity from 175 K to 185 K. Below 175 K, the addition of 0.5 at%Mn hardly change the solidification velocity. Above 175 K, however, 0.5 at%Mn greatly decreases the growth velocity of Co–24 at%Sn alloys. 3.5. Solute partitioning coefficient In the present work, the solute partitioning coefficient of Mn in the a-Co, ka, and b-Co3Sn2 phases, kb, were measured through the rapid-quenching experiments of Co–10 at%Sn hypoeutectic and Co–30 at%Sn hypereutectic alloys containing 1.0 at%Mn. The experimental results reveal that ka and kb are 0.57 and 1.2, respectively (Table 1). Calculated according to the results, the effective solute partitioning of Mn in the solidification of Co–24 at%Sn eutectic alloy, keff, is 0.94, which is slightly less than 1. 3.6. DSC curves and fusion enthalpy Fig. 6 shows the DSC heating curves of Co–Sn alloys with different compositions, including Co–24 at%Sn, Co–24 at%–0.5 at%Mn, Co–10 at%Sn and Co–10 at%Sn–0.5 at%Mn. For Co–24 at%Sn and
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Fig. 3. SEM surface images of the Co–24 at%Sn samples with 0.5 at%Mn addition solidified at undercooling. (a) 55 K; (b) 182 K and (c) 185 K, respectively.
Fig. 4. Anomalous eutectics of the Co–24 at%Sn alloy with 0.5 at%Mn addition undercooled by (a) 10 K; (b) 20 K; (c) 120 K.
Fig. 5. The measured solidification velocities of Co–24 at%Sn eutectic alloys with and without 0.5 at%Mn alloys at different undercoolings.
Co–24 at%–0.5 at%Mn alloys, only one endothermic peak exists on the heating curve, corresponding to the fusion of the eutectic microstructure. There are two endothermic peaks on the heating curve of Co–10 at%Sn alloy with and without 0.5 at%Mn addition. The two peaks correspond to the fusion of eutectic microstructure and the single a-Co phase, respectively. It can be obtained from the DSC curves that the eutectic temperature Teutectic of Co–24 at%Sn alloy containing 0.5 at%Mn is 1387 K, being 2 K larger than that of pure Co–24 at%Sn alloy, i.e. 1385 K. The fusion enthalpies of the eutectic alloy DH, a-Co phase DHa and b-Co3Sn2 phase DHb were evaluated through the DSC results. The fusion enthalpies of Co–24 at%Sn eutectic alloys with and without Mn addition were
Table 1 The experimentally measured solute partitioning coefficient of Mn in the solidification of Co–Sn alloy. The third element
ka
kb
keff
Mn
0.57
1.2
0.94
Fig. 6. DSC heating curves of Co–Sn alloys with different compositions.
listed in Table 2. The addition of 0.5 at%Mn increases the fusion enthalpy of Co–24 at%Sn eutectic alloy and the b-Co3Sn2 phase. Conversely, the fusion enthalpy of the a-Co phase is reduced. 4. Discussions 4.1. Influence of Mn on interface energy anisotropy The addition of a new element will change the interatomic action potential energy of growth interface in the solidification of
Table 2 The fusion enthalpy of Co–24 at%Sn eutectic alloys with and without 0.5 at% Mn addition. The third element
DH (J/mol)
DHa (J/mol)
DHb (J/mol)
No addition 0.5 at%Mn
10,900 12,200
7300 6500
12,400 15,690
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alloy melt, and thus cause the interface energy variation of growth interface [15,16]. In the present paper, alloy thermodynamics is used to evaluate the effects of Mn addition on the interface energy in the solidification of Co–24 at%Sn eutectic alloy. From crystal structure, Spaepen [17] advanced a calculation model of solid/liquid (S/L) interface energy c,
amplitude closely relates with the interface energy anisotropy, and the relation can be expressed as,
F P exp
e7=8
3
P ¼ ðT=T 0 Þ
c ¼ am
DH NA V m 2
1=3
ð2Þ
where NA is Avogadro constant, V m is the average atomic volume, am is a factor related with crystal structure. Since the atomic volume of Mn is close to that of Co, and its addition amount, 0.5 at%, is very small, the influence of Mn addition on the V m is very limited. Besides, the addition of 0.5 at%Mn does not change the crystal structure of a-Co and b-Co3Sn2 phases. According to Eq. (1) the effect of Mn on the interface energy can be evaluated through determining the fusion enthalpy variation of Co–24 at%Sn eutectic alloy caused by the addition of 0.5 at%Mn. As shown in Table 2, the addition of 0.5 at%Mn increases the fusion enthalpy of Co–24 at%Sn eutectic alloy. Based on the Eq. (2), the addition of 0.5 at%Mn enlarges the S/L interface energy in the solidification of Co–24 at%Sn eutectic alloy. Meantime, the addition of Mn reduces a/L interface energy and increases b/L interface energy. Since the fusion entropy of the intermetallic compound b-Co3Sn2 is much higher than that of the a-Co phase and the solid solubility of Mn is larger in the b-Co3Sn2 phase, the addition of 0.5 at%Mn imposes more effects on b/L interface energy than those on a/L interface energy, leading to the enlargement of the coupled eutectic growth interface energy. Since the atomic structures of different crystal planes in one crystal are different, for a curved growth interface, different interface positions have various atomic structures. The addition of a third element will impose different effects on the lattice structures and the interatomic potentials of different crystal planes. Therefore, the variations of interface energy caused by the third element along the curved growth interface tip are different. In the solidification of eutectic alloy with small addition of a third element, the concentration of the third element in the liquid ahead of the growth interface relates with the growth velocity and eutectic lamellar spacing [13]. The normal growth velocities and eutectic lamellar spacing are affected by the position change, resulting in the concentration variation of the third element along a curved growth interface, which also affect the interface energy and its anisotropy. It has already been confirmed that the interface energy anisotropy of eutectic growth interface is very weak in the solidification of undercooled Co–24 at%Sn eutectic alloy [14]. According to the analyses above, the addition of 0.5 at%Mn not only increases the interface energy, but also enlarges the interface energy anisotropy of the coupled eutectic growth interface. 4.2. Influence of Mn on growth morphology pattern The effective partitioning coefficient is slightly less than 1, implying that only small amounts of Mn atoms enrich ahead of the growth interface. Therefore, the solute diffusion field of Mn has little effect on the growth interface morphology. Then, the small addition of Mn mainly influences the growth interface morphology by changing the interface energy and its anisotropy in the solidification of Co–24 at%Sn eutectic. The addition of 0.5 at%Mn strengthens the interface energy anisotropy of the eutectic growth interface. Once the interface energy anisotropy exceeds a critical value, the interface tip is no longer the coldest position along the whole interface, improving the formation of side-branches. Brener et al. [18] proposed that the interference
2ld0
ð3Þ !1=2
R4
ð4Þ
where F is the interference amplitude, P is the relative noise strength, e is the interface energy anisotropy strength, T is the factual temperature, d0 is the capillary length, l is the solute diffusion length, R is the interface tip radius. According to the Eq. (3), the increasing of interface energy anisotropy will decrease the interface interference amplitude when the undercooling keeps constant. This means that the addition of 0.5 at%Mn increases the interface tip stability in the solidification of undercooled Co–24 at%Sn eutectic alloy. Therefore, the addition of 0.5 at%Mn leads to the formation of eutectic dendrite rather than the eutectic seaweed in the solidification of Co–24 at%Sn eutectic alloy at low undercooling. As undercooling increases, the noise strength rises, and the interference amplitude on the growth interface is enlarged correspondingly. When the interference amplitude exceeds a critical value, the stability effect caused by Mn addition is not effective enough to sustain a stable interface tip. So there is a critical undercooling above which the noise interference plays the main role, resulting in the splitting of interface tip. As a result, the eutectic growth interface transform from dendrite to seaweed morphology. When undercooling exceeds a certain value, the crystal grows so fast that the interference on the interface tip rapidly leaves the tip region through tip convection and thus keeps the tip stable in the solidification of undercooled Co–24 at%Sn eutectic melt. In this case, the influence of the third element can be ignored. Those interferences leaving the interface tip lead to the formation of side-branches at the positions behind the tip with a certain distance. Therefore, the eutectic growth interface morphology transforms from factual seaweed to compact seaweed as the undercooling beyond a critical value in the solidification of Co–24 at%Sn eutectic alloy. Brener et al. [18] advanced that the critical undercooling D⁄ for the transition from factual seaweed to compact seaweed relates with the relative noise strength, P, and corresponding law is
D jPj2=5
ð5Þ
Since the value of P is far less than 1 in the undercooled eutectic melt, D⁄ is in direct proportion to P, meaning that the enlargement of P will increase D⁄. Small Mn addition increases the eutectic growth interface energy in the solidification of Co–24 at%Sn eutectic melt, which enlarges the relative noise strength according to Eq. (4), and thus increases the interface interference amplitude. The influence of Mn on the eutectic growth velocity is small, so the growth velocity is not enough to sustain the interface tip stability until the undercooling up to 185 K, at which the seaweed shaped eutectic growth interface transform from factual to compact mode. 4.3. Influence of Mn on the anomalous eutectic In the rapid solidification of eutectic alloy containing a third element, the eutectic lamellar spacing is very small, and thus the enrichment of the third element atom is not enough to break the stability of single eutectic lamellar. As the rapid solidification is to the end, the eutectic lamellar spacing increases gradually, the single eutectic lamellar may split due to the instability effect produced by the enrichment of the third element atoms. These
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splitting eutectic lamellae grow into the remaining liquid as branches and thus break the coupled eutectic growth interface. Since ka of Mn is less than 1 and kb of Mn is larger than 1, more Mn atoms enrich ahead of the a/L interface than that of the b/L interface, meaning that the solute undercooling caused by Mn addition in front of the a-Co phase is larger. Therefore, at the end of rapid solidification the a-Co phase first deviates from the eutectic coupled growth interface and grows into the remaining liquid between adjacent eutectic dendrite arms as a single dendrite. This process is so-called the divorced eutectic growth. During the divorced eutectic growth the tip radius of these a-Co dendrites is much larger than the eutectic lamellar spacing. These a-Co dendrites stay in the liquid phase for long time, subjecting to the re-melting and ripening due to the S/L interface tension, and form another type anomalous eutectic finally. Such a formation mechanism of anomalous eutectic can be named divorced eutectic mechanism. Since the undercooling is very low, at 11 K, the primary eutectic lamellar cannot be re-melted into anomalous eutectic. So the anomalous eutectic lies between eutectic dendrite arms (Fig. 4a). When the undercooling is large enough, the primary formed eutectic lamellae during the rapid solidification are re-melted due to overheating and then transform into anomalous eutectic microstructures. As a result, the anomalous eutectics are formed through two mechanisms in the solidification of Co– 24 at%Sn eutectic alloy with 0.5 at%Mn addition. The anomalous eutectic microstructure with large-size a-Co particles, lying between the eutectic seaweed arms resulted from divorced eutectic growth. While the one with small-size a-Co particles, lying in the internal eutectic seaweed arms, are produced by the re-melting and ripening of primary eutectic lamellae (Fig. 4b). As the undercooling increases further, the primary eutectic lamellae grow faster and the regions between eutectic seaweed arms shrink greatly so that the divorced eutectic microstructures disappear. As a result, only one type of anomalous eutectic forms, which are the productions of rapidly solidified primary eutectic lamellae (Fig. 4c). 5. Conclusions 1. The addition of 0.5 at%Mn effectively changes the interface energy and its anisotropy of the coupled eutectic growth interface during the solidification of undercooled Co–24 at%Sn eutectic alloy. 0.5 at%Mn addition improves the eutectic dendrite growth at a low undercooling. As undercooling increases the morphology of eutectic growth interface transforms from dendrite pattern to fractal seaweed pattern. When undercooling exceeds a critical value the fractal seaweed interface morphology change to compact seaweed morphology, just like that happens in the pure Co–24 at%Sn eutectic alloy. The only difference is that the addition of 0.5 at%Mn increases critical transition undercooling from 175 K to 185 K. 2. Below undercooling of about 175 K, the addition of 0.5 at%Mn has little effect on the solidification velocity of Co–24 at%Sn eutectic alloy. Above the undercooling of 175 K, the addition of 0.5 at%Mn obviously decreases the solidification velocity.
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3. The undercooling for the primary formation of anomalous eutectic is 10 K in the solidification of Co–24 at%Sn eutectic alloy containing 0.5 at%Mn, which is 11 K less than the 21 K for that without Mn addition. The addition of 0.5 at%Mn introduces a new formation mechanism of anomalous eutectic, i.e. divorce eutectic mechanism, which makes anomalous eutectic first form at the inter-dendrite region at a very low undercooling. In the intermediate undercooling range, two mechanisms work together for the formation of anomalous eutectic, and only one type of anomalous eutectic microstructure forms at large undercooling.
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