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Transient directional solidification of a eutectic Al–Si–Ni alloy: Macrostructure, microstructure, dendritic growth and hardness
Materialia 7 (2019) 100358
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Transient directional solidification of a eutectic Al–Si–Ni alloy: Macrostructure, microstructure, dendritic growth and hardness Rafael Kakitani a, Clarissa B. Cruz a, Thiago S. Lima a, Crystopher Brito b, Amauri Garcia a, Noé Cheung a,∗ a b
Department of Manufacturing and Materials Engineering, University of Campinas - UNICAMP, 13083-860 Campinas, SP, Brazil Campus of São João da Boa Vista, São Paulo State University - UNESP, 13876-750 São João da Boa Vista, SP, Brazil
a r t i c l e Keywords: Al–Si–Ni eutectic alloy Solidification Macrostructure Dendritic growth Hardness
i n f o
a b s t r a c t The Al-11 wt%Si-5 wt%Ni eutectic alloy was directionally solidified (DS) under transient heat flow conditions at cooling rates in a range of about 1–25 °C/s along the length of the casting. For comparison purposes binary Al-11 wt%Si and Al-5 wt%Ni alloys were also solidified under similar experimental conditions. Initially characterized by columnar grains, the macrostructure of the ternary DS alloy casting evolved during solidification to a columnar to equiaxed transition (CET) at critical cooling rates of about 1.3–1.6 °C/s, i.e. more than 7 times higher than the CET critical value determined in the literature for Al–Ni and Al–Si alloys. The microstructure of the ternary alloy casting is shown to be constituted by primary Si crystals and 𝛼-Al dendritic branches surrounded by the eutectic phase (𝛼-Al+Si+Al3 Ni), with the Al3 Ni phase having initially a plate-like morphology, which with the decrease in the solidification kinetics, evolved to a fishbone morphology. The pertinent scaling laws representing the dendritic growth of the eutectic ternary alloy, properly compared to those of the examined Al–Si and Al–Ni binary alloys are outlined. The scale of the dendritic 𝛼-Al phase is shown to affect the Vickers hardness (HV) along the length of the ternary eutectic alloy casting, and a Hall–Petch type equation is proposed relating HV to the secondary dendritic arm spacing.
1. Introduction Alloys of eutectic composition emerge as candidates to applications in phase change energy storage, electronic and electrical automatic control, welding etc. Over the last 70 years, this class of materials attracted the attention of several researchers looking for how to explain physicochemical features of the formation and growth of eutectic microstructures. The classical works carried out their approaches over binary systems [1], while the most current works have considered multicomponent systems and have tried to explain how the solute redistribution, associated with thermosolutal effects ahead of the solidification interface, govern the growth of the final microstructure [2,3]. During transient solidification conditions, the solute rejection at the solid/liquid interface is not as fast as the displacement of the solidification front, promoting instabilities that could lead to morphological transitions, precipitation of metastable phases and formation of supersaturated solutions. Some cases in which such transition was associated with the solidification cooling rate (Ṫ) have been reported in the literature. Brito et al. [4] reported a reverse cell to dendrite transition at cooling rates above 2.0 °C/s, that is, unexpected high cooling rate cells
∗
characterized the morphology of the 𝛼-Al matrix of an Al-3 wt%Mg1 wt%Si alloy directionally solidified under transient conditions, with the dendritic pattern being stable only for solidification cooling rates lower than 0.8 °C/s. In an Al–Cu–Mg multicomponent alloy with Fe and Ni additions, the increase in cooling rate suppressed the formation of flower-like and radiated morphologies of Al9 FeNi intermetallic particles, changing them to a morphology of small blocks [5]. Supersaturated solid solutions of Mn and Cr in Al of binary Al-based alloys can reach about 3 times the maximum equilibrium solid solubility when solidified at cooling rates of about 600 °C/s and for Zr, this value is higher than 4 times [6]. In the case of Al–Fe alloys, metastable Al6 Fe particles with fiber/needle-like morphologies were realized at Ṫ > 1 °C/s [7,8] in contrast to the phase diagram, which indicates the formation of the stable Al3 Fe intermetallics. As stated by the coupled zone theory [1], the final morphology of eutectic or near-eutectic ternary alloys is controlled by the growth velocity of each phase that form the eutectic mixture, i.e., depending on the solidification growth rate, the microstructure can be fully eutectic, single-phase dendritic with eutectic or two-phase dendritic with a eutectic component [9]. However, according to the Mullins–Sekerka
Corresponding author. E-mail address: [email protected] (N. Cheung).
https://doi.org/10.1016/j.mtla.2019.100358 Received 1 March 2019; Accepted 19 May 2019 Available online 22 May 2019 2589-1529/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
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Materialia 7 (2019) 100358
Table 1 Chemical composition of elements used to prepare the alloys (wt%). Element
Al
Si
Ni
Cu
Zn
Sn
Fe
Pb
C
S
Co
Al Si Ni
Bal. 0.084 –
0.055 Bal. –
0.006 0.016 Bal.
0.01 – –
0.05 – –
0.005 – –
0.073 0.21 0.004
0.006 – 0.002
– – 0.004
– – 0.002
– – 0.017
interface stability criterion [10], the microstructural arrangement of the eutectic mixture can be modified, induced by non-equilibrium conditions of solidification. The Al-rich eutectic of the Al–Ni phase diagram has a symmetric coupled zone, so a fully eutectic morphology is expected to occur, which was confirmed by unidirectional steady-state solidification experiments [11,12], where aligned Al3 Ni fibers in an 𝛼-Al phase were shown to constitute the eutectic mixture. The unsteady-state solidification regime is responsible for the change from a simple and regular eutectic, to a bimodal structure, composed by eutectic colonies with 𝛼-Al + Al3 Ni in the inner part, and a coarse lamellar morphology at the boundaries [13]. In these two cases, despite the control of the microstructure fineness being exerted by the solidification cooling rate, the Mullins–Sekerka stability criterion was the main factor determining the final arrangement, since for a same Ṫ different microstructures resulted for steady and unsteady conditions of solidification. With respect to the morphology of the resulting macrostructure, some studies in the literature report that a columnar to equiaxed transition (CET) can occur according to critical values either of growth rate (v) or temperature gradient (G) in the liquid ahead the solidification front [14,15]. However, several works on aluminum binary alloys (Al– Cu, Al–Si, Al–Ni and Al–Sn) solidified under unsteady-state solidification conditions, state that this transition is imposed by the simultaneous action of G and v, that is for a critical value of the solidification cooling rate (Ṫ =Gv) for each alloy system, within a range comprised between 0.16 and 0.29 °C/s [16–18]. Al–Si–Ni alloys are one of the preferred materials for pistons, being recognized by their good castability, high mechanical strength, good wear resistance and low thermal expansion [19–21], and in addition, when close to the eutectic composition they are prone to form metallic glasses [22–24]. It is the intention of this work to develop a comparative investigation on the evolution of macrostructure, microstructure and dendritic growth as a function of a wide range of solidification thermal parameters (growth rate and cooling rate) during transient directional solidification of a ternary eutectic Al–Ni–Si alloy (Al-11 wt%Si-5 wt%Ni) and of the corresponding Al-5 wt%Ni and Al-11 wt%Si binary alloys. Experimental growth laws are envisaged relating the scale of the dendritic arrangements to v and Ṫ, as well as to the hardness along the length of the directionally solidified (DS) castings with a view to deriving useful information for the processing of these alloys through solidification. 2. Materials and methods The unsteady-state solidification of Al-11 wt%Si, Al-5 wt%Ni and Al11 wt%Si-5 wt%Ni alloys were carried out in a directional solidification apparatus, where radial electrical wiring allows to control the temperature in the melt and heat is extracted only through a water-cooled bottom made of AISI 1020 steel. A stainless steel mold was used, which has an internal diameter of 55 mm, a height of 110 mm, and a wall thickness of 5 mm. Once the melt temperature reached 10% above liquidus (for binary alloys) or eutectic temperatures (for the ternary alloy), the electric heaters were disconnected, and the water-flow system began to cool down permitting the onset of solidification. More details about the solidification setup have been described in a previous work [25]. Commercial pure elements, having the chemical compositions shown in Table 1, were used to prepare these alloys. The temperature during solidification was acquired by fine K-type thermocouples (1.6 mm
Fig. 1. Al-Xwt%Si-5 wt%Ni partial phase diagram calculated by the ThermoCalc software (TTAL5 database).
of external diameter) placed along the length of the castings and connected by coaxial cables to a data logger, which recorded the temperature at a frequency of 5 Hz. Cooling rates and growth rates during solidification were determined using these temperature data. In addition, the Al-11 wt%Si-5 wt%Ni alloy was also poured in a well-insulated SiC crucible, air-cooled, with a view to permitting the transformation temperatures to be assessed by the resulting experimental cooling curve. In Fig. 1, the dashed vertical line indicates the composition of the examined eutectic Al–Si–Ni alloy in the partial Al-11 wt%Si-Xwt%Ni phase diagram determined by the computational thermodynamics Thermo-Calc software using the TTAL5 database. The liquidus temperatures of the binary Al-11 wt%Si and Al-5 wt%Ni alloys, 588 °C and 646 °C, respectively, were also determined by the Thermo-Calc software, using the TTAL5 database. The macrostructure of each alloy was etched with aqua regia solution (3:1 mixture of HCl and HNO3 , respectively). The dendritic arm spacings were measured on transverse and longitudinal sections of samples extracted along the length of the DS castings. The triangle method [26] was used in order to perform measurements of the primary dendritic arm spacing and the intercept method [26] for secondary and tertiary dendritic arm spacings (about 50, 200 and 300 measurements for each selected position, respectively). Some samples of the examined alloys were selected and subjected to a deep etching technique with HCl for 2 min. Chemical analyses along the length of DS Al-11 wt%Si-5 wt%Ni alloy casting were performed in an X-Ray Fluorescence (XRF) spectrometer (Rigaku-RIX3100) with a view to checking the distribution of solute elements. The phases forming the alloy microstructure were determined
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Fig. 2. Experimental cooling curves along the length of: (a) Al-11 wt%Si, (b) Al-5 wt%Ni and (c) Al-11 wt%Si-5 wt%Ni alloys castings.
by X-Ray Diffraction (XRD), using a PAN analytical X’Pert PRO Materials Research Diffractometer XL, Cu-K𝛼 radiation with a wavelength of 0.15406 nm. The microstructure was also investigated by scanning electron microscopy with field emission gun (SEM-FEG) FEI Inspect F50 and by energy dispersive X-ray spectroscopy (EDS). Vickers microhardness tests were performed on the cross sections of the samples using a test load of 1 kgf (9.8 N) and a dwell time of 10 s. The adopted hardness value of a representative position was the average of at least 30 measurements on each sample.
Table 2 Different solidification experiments (#) with the Al-11 wt%Si-5 wt%Ni eutectic alloy: position of macrostructural transition and related solidification thermal parameters. Experiment
#1
#2
#3
Transition [mm] Ṫ [°C/s] vE [mm/s] G= Ṫ/vE [°C/mm]
55–65 1.5–1.3 0.35–0.31 4.3–4.2
63–78 1.6–1.3 0.46–0.42 3.5–3.1
85–100 1.3–1.1 0.22–0.19 5.9–5.8
3. Results and discussion Fig. 2 shows the experimental cooling curves recorded at different positions along the length of the DS Al-11 wt%Si, Al-5 wt%Ni and Al11 wt%Si-5 wt%Ni alloys castings, which were used to determine the solidification thermal parameters: cooling rate (Ṫ) and growth rate (v). The directional solidification setup provides a wide range of Ṫ and v values in a single experiment. For the binary alloys (Al-11 wt%Si and Al-5 wt%Ni) the thermal parameters were evaluated with respect to the liquidus temperature (TL ), while for Al-11 wt%Si-5 wt%Ni the eutectic temperature (TE ) was used. In this regard, the subscripts □L and □E correspond to the hypoeutectic and eutectic alloys, respectively. As the bottom of the mold was made of steel, the chemical affinity between Ni and Fe seems to have improved the alloy/mold wettability of the Al5 wt%Ni alloy casting as compared to the two other alloy castings, thus resulting in higher Ṫ and v profiles, as shown in Fig. 3. On the other hand, Si seems to be detrimental to the wettability, because the eutectic ternary alloy reached intermediate values of Ṫ and v (Fig. 3). Columnar grains prevailed along the length of the DS binary alloys castings, without any evidence of CET (Figs. 4 and 5). According to previous studies with Al–Ni [16] and Al–Si [17] alloys, the critical solidification cooling rate at which CET occurs for these alloys is about 0.16–0.17 °C/s, values not achieved along the length of the DS castings examined in the present work. The as-solidified microstructures of both alloys were composed by an 𝛼-Al matrix of dendritic morphology with the eutectic mixture located in the interdendritic region, with primary Si crystals also appearing in the microstructure of the Al-11 wt%Si alloy. Fig. 6a–d shows the eutectic microstructures of the Al-11 wt%Si alloy (𝛼Al matrix and plate-like Si) and the Al-5 wt%Ni alloy (𝛼-Al matrix and Al3 Ni fibers) obtained through conventional metallography, as well as
the microstructures that resulted from deep etching used to reveal the three-dimensional morphology of the phases. Initially characterized by columnar grains, the macrostructure of the ternary DS alloy casting evolved to a transition between 55 and 65 mm with respect to the cooled base of the casting (Fig. 7). Visually, this transition could not be associated with a sharp structural change from columnar to equiaxed grains, but rather some columnar grains have lost their original direction of growth and new grains nucleated ahead the columnar front. Two additional experiments with the same alloy were performed to prove the macrostructural transition observed in the initial experiment. Table 2 shows the regions in the castings of each experiment in which this transition occurred and the corresponding solidification thermal parameters. The cooling rate seems to be the most realistic thermal parameter determining the macrostructural transition since it encompasses the narrower range of values (1.3–1.6 °C/s) when Ṫ, vE and G are compared, that is, a critical cooling rate that is more than 7 times higher than the CET critical value determined in the literature for Al–Ni [16] and Al–Si [17] alloys. The solute concentration along the length of the Al-11 wt%Si-5 wt%Ni alloy casting, determined by X-Ray Fluorescence (XRF) (Fig. 8a), did not indicate the occurrence of macrosegregation, with Si and Ni concentrations quite close to the nominal composition of the alloy along the length of the whole casting. Since the Al-11 wt%Si-5 wt%Ni alloy has a eutectic composition [27–29], primary Si crystals and 𝛼-Al dendritic branches were formed surrounded by the eutectic phase (𝛼-Al+Si+Al3 Ni) in the columnar region, as can be seen in Fig. 7. In Fig. 8b, X-Ray Diffraction (XRD) patterns of samples extracted from different positions along the length
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Fig. 3. Experimental profiles of (a) cooling rate and (b) growth rate against the position from the cooled surface of the DS castings for the Al-11 wt%Si, Al-5 wt%Ni and Al-11 wt%Si-5 wt%Ni alloys.
of the casting confirmed the aforementioned phases and Scanning Electron Microscopy / Energy Dispersive X-Ray Spectroscopy (SEM/EDS) analyses (Fig. 9) indicated absence of supersaturated solid solution in the 𝛼-Al phase and the atomic compositions were close to the stoichiometric formulas of each phase. The Al3 Ni intermetallics has a plate-like morphology instead of fibers due to the faceted growth of Si that inhibits
the formation of regular eutectic structures [29]. Fig. 6e–f shows the eutectic phase of Al-11 wt%Si-5 wt%Ni (matrix of 𝛼-Al with plate-like Si and Al3 Ni). From the position 15 mm (Ṫ = 5.8 °C/s and vE = 0.76 mm/s), that is, with the decrease in the solidification kinetics, a certain fraction of the Al3 Ni phase starts to lose its plate-like morphology, evolving
Fig. 4. Macrostructure of the DS Al-11 wt%Si alloy casting and longitudinal microstructures along the length of the casting.
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Fig. 5. Macrostructure of the DS Al-5 wt%Ni alloy casting and longitudinal microstructures along the length of the casting. Fig. 6. Eutectic of a) Al-11 wt%Si alloy with primary crystal of Si and b) deep etched sample showing irregular plate-like Si; c) Al-5 wt%Ni alloy and d) deep etched sample showing Al3 Ni fibers; and e) Al-11 wt%Si-5 wt%Ni alloy: dark phase is Si and gray phase is Al3 Ni and f) deep etched sample showing irregular plate-like Si and Al3 Ni.
Fig. 7. Macrostructure of the DS Al-11 wt%Si-5 wt%Ni alloy casting and longitudinal microstructures along the length of the casting.
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Fig. 8. (a) Solute distribution along the length of the DS Al-11 wt%Si-5 wt%Ni alloy casting and (b) typical XRD patterns.
Fig. 9. SEM-EDS results for samples from positions: (a) 5 mm and (b) 50 mm along the length of the Al-11 wt%Si-5 wt%Ni alloy casting.
to a fishbone morphology, as shown in Fig. 10. In the experiment #2 with the Al-11 wt%Si-5 wt%Ni alloy, such morphological transition occurred for similar solidification thermal parameters: Ṫ = 5.7 °C/s and vE = 0.81 mm/s. This kind of morphology has been reported to occur also in alloys of the Al–Mg–Mn–Fe, Al–Si–Cu–Ni–Mg–Co and Al–Si–Sc systems [30–32], being generally associated with high cooling rates (watercooled molds), which favor the formation of metastable phases [30,33]. Since the aforementioned Al–Mg–Mn–Fe system does not have Si in its composition and even so the formation of Al3 Ni with fishbone morphology occurs, it seems that both phases (Si and Al3 Ni) grow independently, and that the prevailing factor for the formation of the fishbone morphology is the cooling rate. The fishbone Al3 Ni phase occupies a wider area of the microstructure as compared to that occupied by the plate-like
one, and this seems to be responsible for the increase in the intensity of peaks in the XRD patterns related to the Al3 Ni phase that is formed at low cooling rates (Fig. 8b). Fig. 11a shows the cooling curve of the eutectic Al-11 wt%Si5 wt%Ni alloy obtained under slow cooling conditions (air-cooled: Ṫ∼0.2 °C/s), showing a eutectic temperature of 568 °C, which is in good agreement with results from the literature [27–29]. The corresponding microstructure, shown in Fig. 11b, is formed by a mixture of eutectic, whose area fraction is predominant, and primary phases. In a Bridgman steady-state solidification study [27], only the eutectic mixture (𝛼-Al+Si+Al3 Ni) was reported to form the microstructure of eutectic Al–Si–Ni alloy samples that were grown both at constant G (5.82 K/mm) and different v (4.60 to 35.14 μm/s), and constant v
R. Kakitani, C.B. Cruz and T.S. Lima et al.
Fig. 10. Morphological transition of Al3 Ni particles (gray phase) from plate-like to fishbone, with the decrease in the solidification cooling rate along the length of the Al-11 wt%Si-5 wt%Ni alloy casting.
(11.63 μm/s) and different G (2.11–5.82 K/mm). With G = 5.82 K/mm and increasing v (v> 173.61 μm/s), 𝛼-Al dendritic branches were also reported to occur, and once Ṫ=G∗ v, for Ṫ < 1.0 °C/s the microstructure was exclusively eutectic, and higher cooling rates induced the formation of dendrites. Therefore, similarly to the irregular Al–Si eutectic alloy, the Al–Si–Ni eutectic alloy, at the Al-rich side, has an asymmetric coupled zone in which a fully eutectic structure is obtained only for low Ṫ or vE , as indicated in Fig. 12. A deviation from the coupled zone results in an 𝛼-Al+E or E + X, where E is the eutectic and X a phase or an undefined structure. In the present work, a dendritic arrangement is justified by the coupled zone, however, there is no certainty as to the determinant factor that led to the formation of primary crystals of Si and fishbone Al3 Ni, that is, whether Ṫ or vE or even both. What is certain is that these phases arose due to the non-equilibrium conditions typical of transient solidification. The macrostructure transition of the eutectic Al–Si–Ni alloy occurred close to 1.0 °C/s, which is also the critical value associated with the microstructural transition between fully eutectic and eutectic/𝛼-Al den-
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drites, and this is probably connected to the nucleation of eutectic grains ahead the solidification front, which blocked the columnar growth. Some literature studies focusing on the nucleation of irregular eutectic Si have reported the occurrence of eutectic grains in near-eutectic Al– Si alloys [34–36]. The main theory is that a heterogeneous nucleation occurs induced by impurities, such as Fe and P that form intermetallic particles, or by 𝛼-Al dendrites [36]. Moreover, it has been reported that in ternary Al–Si–(Cu,Mg) alloys the number of eutectic grains has a bias dependence on the added element: Mg up to 0.5 wt% and Cu additions increase the density of eutectic nuclei [34]. According to a previous study [34], different ternary alloying elements added to the Al–Si alloy system imply different diffusion and partition coefficients (k). A faster diffusion rate and smaller k value affects the solute gradient at the solid/liquid interface and increases segregation, respectively, thus favoring the rate of eutectic nucleation. Thus, since the Al-11 wt%Si-5 wt%Ni is an irregular eutectic alloy, at high solidification cooling rates primary 𝛼-Al dendrites grew preferentially in relation to the eutectic phase [37], and the decrease in Ṫ favors the precipitation of eutectic grains with the simultaneous formation of dendrites at the grain boundaries [38]. Moreover, the Ni content may aid the precipitation of eutectic grains, considering that the partition coefficient of Ni in Al is low (k = 0.007) [39]. Kaya and Aker [40] added 2 wt%Ni to the eutectic Al-12.6 wt% alloy and reported a reduction in the Si interflake spacing that seems to be induced by higher number of eutectic nuclei. It is suggested that the addition of Ni has contributed to increase the rate of eutectic nucleation. Alloys with a eutectic composition have a resulting microstructure essentially constituted by (i) a eutectic arrangement, or (ii) eutectic mixture plus the presence of the primary phase of the preponderant element [1]. The coupled zone theory [1] reports that one or other of these morphologies will depend on the product between the thermal gradient and the growth rate, i.e., on the cooling rate. Hence, the further away from the equilibrium conditions the solidification occurs, more susceptible to the formation of the second type of morphology the system will be. Fig. 13 shows the experimental dendrite arm spacings (𝜆1 –3 ) determined in the present study as a function of cooling rate and growth rate, with the error bars representing the standard deviation. Experimental power laws are proposed having exponents that are in agreement with those usually reported in the literature [41,42]. As expected, the higher Ṫ the smaller the dendrite arm spacing. Single experimental power laws can correlate the primary dendrite arm spacing (𝜆1 ) with either Ṫ or vL /vE for the Al-11 wt%Si-5 wt%Ni and Al-11 wt%Si alloys (Fig. 13a–b).
Fig. 11. (a) Cooling curve of an air-cooled Al-11 wt%Si-5 wt%Ni alloy (Ṫ ∼ 0.2 °C/s) and (b) corresponding microstructure.
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Materialia 7 (2019) 100358
Fig. 12. Asymmetric coupled zone of the Al-11 wt%Si-5 wt%Ni eutectic as a function of Ṫ or vE , where E is the eutectic and X a phase or an undefined structure.
Different power laws are proposed to represent the Al-5 wt%Ni alloy, which has a more refined range of 𝜆1 values. The exponents −0.55 and −1.1 seem to be appropriate to relate 𝜆1 with the cooling rate and the growth rate, respectively. For the secondary dendrite arm spacing (𝜆2 ), single experimental laws (Fig. 13c–d) represent the evolution of 𝜆2 for the Al-11 wt%Si-5 wt%Ni and Al-5 wt%Ni alloys with either Ṫ or vL /vE , suggesting that Ni has a refining role on 𝜆2 , as compared to the results of the Al-11 wt%Si alloy. For the Al-11 wt%Si alloy, the same exponents derived for the 𝜆1 power laws, can also be applied to the experimental power laws relating the evolution of the tertiary dendrite arm spacing (𝜆3 ) with the solidification thermal parameters, as shown in Fig. 13e–f. When the −0.55 and −1.1 exponents were applied to correlate 𝜆3 with Ṫ and v, respectively, for the Al-11 wt%Si-5 wt%Ni alloy, the fit did not include most of the experimental data, suggesting that different exponents should be used. Theoretical models for the prediction of dendritic growth during solidification of multicomponent alloys as a function of solidification growth rate and cooling rate, cannot be found in the literature because of the difficulty of the task. On the other hand, different exponents for experimental growth laws in ternary alloy systems, such as Al–Si–Mg [43], Al–Cu–Ni [44] and Al–Si–Cu [45], have been reported when compared to the usual ones experimentally determined for binary alloys. Based on previous studies on Sn–Bi–(Cu,Ag) and Al-11 wt%Si-(3,4.5)wt%Cu ternary alloys [46,47], experimental power laws using −0.25 and −0.50 exponents were shown to be more appropriate to relate 𝜆3 to Ṫ and vE , respectively, as can be seen in Fig. 13e–f. No evidence of tertiary dendritic branches was found along the microstructure of the Al-5 wt%Ni alloy. These distinct behaviors on the evolution of dendrite arm spacings with Ṫ and vL/E , observed in Fig. 13, could be related to the differences in the amount of Si (11 wt%) and Ni (5 wt%) in the Al–Si–Ni alloy composition and to their densities with respect to that of liquid Al, which are: Al, Si and Ni: 2385 kg/m3 , 2340 kg/m3 and 8900 kg/m3 , respectively [48]. At the beginning of solidification of the eutectic Al–Si–Ni alloy, the rejection of Si to the molten liquid is more remarkable than that of Ni due to the higher amount of Si in the alloy composition. It is worth noting that the alloy solidified vertically upwards, i.e. the growth direction is opposite to that of the gravity vector. As Si has a density that is even lower than that of liquid Al, the Si rejected during solidification has a trend to float and concentrate along the solidification front provoking instabilities at this interface that lead to the formation of primary dendrite stalks, as shown schematically at the top row of Fig. 14. The enrichment of Si in the liquid around these dendritic stalks, could lead to local con-
centrations that could surpass the composition of the Al–Si eutectic, thus creating conditions for precipitation of primary crystals of Si. As Ni has a higher density than liquid Al, when rejected at the solidification interface it tends to move downward in the liquid channels located between the primary dendritic arms, thus provoking gradual lateral instabilities leading to the formation of secondary dendritic branches (central row of Fig. 14). In this sense, the growth of primary branches seems to be mainly controlled by the rejection of Si, whereas the secondary branches by the rejection of Ni, explaining the similar growth laws in Fig. 13: 𝜆1 (Al-11 wt%Si-5 wt%Ni and Al-11 wt%Si -Fig. 13a and b) and 𝜆2 (Al11 wt%Si-5 wt%Ni and Al-5 wt%Ni -Fig. 13c and d). The development of tertiary dendritic branches would be caused by the simultaneous action of atoms of Si and Ni (bottom row of Fig. 14). At regions in the Al-11 wt%Si-5 wt%Ni alloy casting associated with higher solidification kinetics, the rejected Si atoms could also initiate local instabilities leading to the formation of tertiary branches, and the local tertiary spacings would be like those of the Al-11 wt%Si alloy. With the decrease in the solidification kinetics, the growth of tertiary branches would be modified by the presence of Ni atoms, and the growth law relating 𝜆3 to Ṫ and vE for the Al-11 wt%Si-5 wt%Ni alloy casting, would progressively depart from the trend observed for the 𝜆3 results of the Al-11 wt%Si alloy casting (Fig. 13e–f). Fig. 15 shows the results of Vickers microhardness measured along the length of the DS castings against a representative microstructural spacing, 𝜆2 , i.e. the common dendrite spacing of higher order measured in all examined alloys (tertiary spacings have not been detected in the Al-5 wt%Ni alloy), which is reported to be responsible for a closer interaction among dendritic matrix, eutectic mixture and second phases [49]. The isolated reinforcing role of plate-like Si particles and Al3 Ni fibers for the binary Al-11 wt%Si and Al-5 wt%Ni alloys, respectively, seems to have similar performance in the blockage of dislocations all along the length of the DS castings, without any the influence of the scale of the dendritic Al-rich matrix, resulting in an average hardness of 48 HV. However, for the Al-11 wt%Si-5 wt%Ni alloy casting, the simultaneous interaction of both reinforcing phases located in the interdendritic spacings and the scale of the secondary dendritic arms is conducive to a remarkable increase in hardness with the decrease in 𝜆2 . A Hall–Petch-type correlation is proposed relating HV to 𝜆2 −1/2 . Along a comparative common range of 𝜆2 values of the ternary and binary alloys, the hardness of the ternary alloy increases from 46% to 77% with the decrease in 𝜆2 (increase in 𝜆2 −1/2 ) as compared to the mean hardness of the binary alloys (HV = 48).
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Fig. 14. Schematic representation of possible stages of formation of dendritic branches during the transient solidification of an Al-11 wt%Si-5 wt%Ni alloy, where S is the solid, L is the liquid and SF is the solidification front.
Fig. 13. Evolution of: (a, b) primary, (c, d) secondary and (e, f) tertiary dendritic spacings as a function of the growth rate (vE / L ) and of the cooling rate (Ṫ) for the Al-11 wt%Si-5 wt%Ni, Al-11 wt%Si and Al-5 wt%Ni alloys castings. R2 is the coefficient of determination.
Fig. 15. Vickers microhardness as a function of 𝜆2 −1/2 for the Al-11 wt%Si5 wt%Ni, Al-11 wt%Si and Al-5 wt%Ni alloys castings. R2 is the coefficient of determination.
R. Kakitani, C.B. Cruz and T.S. Lima et al.
4. Conclusions The following conclusions can be drawn from the present experimental study: • A columnar to equiaxed transition (CET) was shown to occur for the DS Al-11 wt%Si-5 wt%Ni alloy casting under solidification cooling rates in the range 1.3–1.6 °C/s. These critical cooling rate values are more than 7 times higher than the CET critical values determined in the literature for Al–Ni and Al–Si alloys. • The microstructure of the DS Al-11 wt%Si-5 wt%Ni eutectic alloy casting was characterized by primary Si crystals; 𝛼-Al dendritic matrix and 𝛼-Al+Si+Al3 Ni eutectic due to non-equilibrium transient solidification conditions. The Al3 Ni intermetallics was shown to have a plate-like morphology instead of fibers, however, for cooling rates Ṫ <5.7 °C/s and growth rates vE < 0.81 mm/s, it was shown to lose its plate-like morphology, evolving to a fishbone morphology. • Experimental growth laws relating the primary (𝜆1 ), secondary (𝜆2 ) and tertiary (𝜆3 ) dendritic spacings of the ternary Al-11 wt%Si5 wt%Ni and the corresponding binary Al-11 wt%Si and Al-5 wt%Ni alloys, to both the solidification cooling rate (Ṫ) and the growth rate (vE/L ) were shown to be given by: 𝜆1 = 112 vE/ L − 1.1 ; 𝜆1 = 405 Ṫ −0.55 for the Al-11 wt%Si-5 wt%Ni and Al-11 wt%Si alloys 𝜆1 = 54 vL −1.1 ; 𝜆1 = 149 Ṫ −0.55 for the Al-5 wt%Ni alloy 𝜆2 = 5.2 vE/ L − 2/3 ; 𝜆2 = 10.9 Ṫ −1/3 for the Al-11 wt%Si-5 wt%Ni and Al-5 wt%Ni alloys 𝜆2 = 9.1 vL −2/3 ; 𝜆2 = 19.2 Ṫ −1/3 for the Al-11 wt%Si alloy 𝜆3 = 5.5 vE −0.50 ; 𝜆3 = 9.9 Ṫ −0.25 for the Al-11 wt%Si-5 wt%Ni alloy 𝜆3 = 6.3 vL −1.1 ; 𝜆3 = 21.8 Ṫ −0.55 for the Al-11 wt%Si alloy where 𝜆1,2,3 [μm]; vE/L [mm/s] and Ṫ [ ̊C/s] • The Vickers microhardness (HV) along the length of both Al11 wt%Si and Al-5 wt%Ni alloys castings was shown not to be affected by 𝜆2 , with a mean HV = 48. In contrast, a Hall–Petch-type correlation has been proposed relating HV to 𝜆2 for the Al-11 wt%Si5 wt%Ni alloy casting: HV = 36+ 103 𝜆2 − 1/2 , representing a 46–77% increase with the decrease in 𝜆2 as compared to the hardness of the binary alloys. Acknowledgments The authors are grateful to FAPESP (São Paulo Research Foundation - grant 2017/15158-0), CNPq -National Council for Scientific and Technological Development and CAPES - Coordenação de Aperfeiçoamento de Pessoal de Nível Superior - Brasil (Finance Code 001) for their financial support and Brazilian Nanotechnology National Laboratory LNNano for the use of its facilities. Declaration of interest None. Supplementary material Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.mtla.2019.100358. References [1] W. Kurz, D.J. Fisher, Dendrite growth in eutectic alloys: the coupled zone, Int. Mater. Rev. 24 (1979) 177–204. https://doi.org/10.1179/imtr.1979.24.1.177. [2] L. Rátkai, A. Szállás, T. Pusztai, T. Mohri, L. Gránásy, Ternary eutectic dendrites: pattern formation and scaling properties, J. Chem. Phys. 142 (2015) 154501. https:// doi.org/10.1063/1.4917201.
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Full Length Article
Transient directional solidification of a eutectic Al–Si–Ni alloy: Macrostructure, microstructure, dendritic growth and hardness Rafael Kakitani a, Clarissa B. Cruz a, Thiago S. Lima a, Crystopher Brito b, Amauri Garcia a, Noé Cheung a,∗ a b
Department of Manufacturing and Materials Engineering, University of Campinas - UNICAMP, 13083-860 Campinas, SP, Brazil Campus of São João da Boa Vista, São Paulo State University - UNESP, 13876-750 São João da Boa Vista, SP, Brazil
a r t i c l e Keywords: Al–Si–Ni eutectic alloy Solidification Macrostructure Dendritic growth Hardness
i n f o
a b s t r a c t The Al-11 wt%Si-5 wt%Ni eutectic alloy was directionally solidified (DS) under transient heat flow conditions at cooling rates in a range of about 1–25 °C/s along the length of the casting. For comparison purposes binary Al-11 wt%Si and Al-5 wt%Ni alloys were also solidified under similar experimental conditions. Initially characterized by columnar grains, the macrostructure of the ternary DS alloy casting evolved during solidification to a columnar to equiaxed transition (CET) at critical cooling rates of about 1.3–1.6 °C/s, i.e. more than 7 times higher than the CET critical value determined in the literature for Al–Ni and Al–Si alloys. The microstructure of the ternary alloy casting is shown to be constituted by primary Si crystals and 𝛼-Al dendritic branches surrounded by the eutectic phase (𝛼-Al+Si+Al3 Ni), with the Al3 Ni phase having initially a plate-like morphology, which with the decrease in the solidification kinetics, evolved to a fishbone morphology. The pertinent scaling laws representing the dendritic growth of the eutectic ternary alloy, properly compared to those of the examined Al–Si and Al–Ni binary alloys are outlined. The scale of the dendritic 𝛼-Al phase is shown to affect the Vickers hardness (HV) along the length of the ternary eutectic alloy casting, and a Hall–Petch type equation is proposed relating HV to the secondary dendritic arm spacing.
1. Introduction Alloys of eutectic composition emerge as candidates to applications in phase change energy storage, electronic and electrical automatic control, welding etc. Over the last 70 years, this class of materials attracted the attention of several researchers looking for how to explain physicochemical features of the formation and growth of eutectic microstructures. The classical works carried out their approaches over binary systems [1], while the most current works have considered multicomponent systems and have tried to explain how the solute redistribution, associated with thermosolutal effects ahead of the solidification interface, govern the growth of the final microstructure [2,3]. During transient solidification conditions, the solute rejection at the solid/liquid interface is not as fast as the displacement of the solidification front, promoting instabilities that could lead to morphological transitions, precipitation of metastable phases and formation of supersaturated solutions. Some cases in which such transition was associated with the solidification cooling rate (Ṫ) have been reported in the literature. Brito et al. [4] reported a reverse cell to dendrite transition at cooling rates above 2.0 °C/s, that is, unexpected high cooling rate cells
∗
characterized the morphology of the 𝛼-Al matrix of an Al-3 wt%Mg1 wt%Si alloy directionally solidified under transient conditions, with the dendritic pattern being stable only for solidification cooling rates lower than 0.8 °C/s. In an Al–Cu–Mg multicomponent alloy with Fe and Ni additions, the increase in cooling rate suppressed the formation of flower-like and radiated morphologies of Al9 FeNi intermetallic particles, changing them to a morphology of small blocks [5]. Supersaturated solid solutions of Mn and Cr in Al of binary Al-based alloys can reach about 3 times the maximum equilibrium solid solubility when solidified at cooling rates of about 600 °C/s and for Zr, this value is higher than 4 times [6]. In the case of Al–Fe alloys, metastable Al6 Fe particles with fiber/needle-like morphologies were realized at Ṫ > 1 °C/s [7,8] in contrast to the phase diagram, which indicates the formation of the stable Al3 Fe intermetallics. As stated by the coupled zone theory [1], the final morphology of eutectic or near-eutectic ternary alloys is controlled by the growth velocity of each phase that form the eutectic mixture, i.e., depending on the solidification growth rate, the microstructure can be fully eutectic, single-phase dendritic with eutectic or two-phase dendritic with a eutectic component [9]. However, according to the Mullins–Sekerka
Corresponding author. E-mail address: [email protected] (N. Cheung).
https://doi.org/10.1016/j.mtla.2019.100358 Received 1 March 2019; Accepted 19 May 2019 Available online 22 May 2019 2589-1529/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
R. Kakitani, C.B. Cruz and T.S. Lima et al.
Materialia 7 (2019) 100358
Table 1 Chemical composition of elements used to prepare the alloys (wt%). Element
Al
Si
Ni
Cu
Zn
Sn
Fe
Pb
C
S
Co
Al Si Ni
Bal. 0.084 –
0.055 Bal. –
0.006 0.016 Bal.
0.01 – –
0.05 – –
0.005 – –
0.073 0.21 0.004
0.006 – 0.002
– – 0.004
– – 0.002
– – 0.017
interface stability criterion [10], the microstructural arrangement of the eutectic mixture can be modified, induced by non-equilibrium conditions of solidification. The Al-rich eutectic of the Al–Ni phase diagram has a symmetric coupled zone, so a fully eutectic morphology is expected to occur, which was confirmed by unidirectional steady-state solidification experiments [11,12], where aligned Al3 Ni fibers in an 𝛼-Al phase were shown to constitute the eutectic mixture. The unsteady-state solidification regime is responsible for the change from a simple and regular eutectic, to a bimodal structure, composed by eutectic colonies with 𝛼-Al + Al3 Ni in the inner part, and a coarse lamellar morphology at the boundaries [13]. In these two cases, despite the control of the microstructure fineness being exerted by the solidification cooling rate, the Mullins–Sekerka stability criterion was the main factor determining the final arrangement, since for a same Ṫ different microstructures resulted for steady and unsteady conditions of solidification. With respect to the morphology of the resulting macrostructure, some studies in the literature report that a columnar to equiaxed transition (CET) can occur according to critical values either of growth rate (v) or temperature gradient (G) in the liquid ahead the solidification front [14,15]. However, several works on aluminum binary alloys (Al– Cu, Al–Si, Al–Ni and Al–Sn) solidified under unsteady-state solidification conditions, state that this transition is imposed by the simultaneous action of G and v, that is for a critical value of the solidification cooling rate (Ṫ =Gv) for each alloy system, within a range comprised between 0.16 and 0.29 °C/s [16–18]. Al–Si–Ni alloys are one of the preferred materials for pistons, being recognized by their good castability, high mechanical strength, good wear resistance and low thermal expansion [19–21], and in addition, when close to the eutectic composition they are prone to form metallic glasses [22–24]. It is the intention of this work to develop a comparative investigation on the evolution of macrostructure, microstructure and dendritic growth as a function of a wide range of solidification thermal parameters (growth rate and cooling rate) during transient directional solidification of a ternary eutectic Al–Ni–Si alloy (Al-11 wt%Si-5 wt%Ni) and of the corresponding Al-5 wt%Ni and Al-11 wt%Si binary alloys. Experimental growth laws are envisaged relating the scale of the dendritic arrangements to v and Ṫ, as well as to the hardness along the length of the directionally solidified (DS) castings with a view to deriving useful information for the processing of these alloys through solidification. 2. Materials and methods The unsteady-state solidification of Al-11 wt%Si, Al-5 wt%Ni and Al11 wt%Si-5 wt%Ni alloys were carried out in a directional solidification apparatus, where radial electrical wiring allows to control the temperature in the melt and heat is extracted only through a water-cooled bottom made of AISI 1020 steel. A stainless steel mold was used, which has an internal diameter of 55 mm, a height of 110 mm, and a wall thickness of 5 mm. Once the melt temperature reached 10% above liquidus (for binary alloys) or eutectic temperatures (for the ternary alloy), the electric heaters were disconnected, and the water-flow system began to cool down permitting the onset of solidification. More details about the solidification setup have been described in a previous work [25]. Commercial pure elements, having the chemical compositions shown in Table 1, were used to prepare these alloys. The temperature during solidification was acquired by fine K-type thermocouples (1.6 mm
Fig. 1. Al-Xwt%Si-5 wt%Ni partial phase diagram calculated by the ThermoCalc software (TTAL5 database).
of external diameter) placed along the length of the castings and connected by coaxial cables to a data logger, which recorded the temperature at a frequency of 5 Hz. Cooling rates and growth rates during solidification were determined using these temperature data. In addition, the Al-11 wt%Si-5 wt%Ni alloy was also poured in a well-insulated SiC crucible, air-cooled, with a view to permitting the transformation temperatures to be assessed by the resulting experimental cooling curve. In Fig. 1, the dashed vertical line indicates the composition of the examined eutectic Al–Si–Ni alloy in the partial Al-11 wt%Si-Xwt%Ni phase diagram determined by the computational thermodynamics Thermo-Calc software using the TTAL5 database. The liquidus temperatures of the binary Al-11 wt%Si and Al-5 wt%Ni alloys, 588 °C and 646 °C, respectively, were also determined by the Thermo-Calc software, using the TTAL5 database. The macrostructure of each alloy was etched with aqua regia solution (3:1 mixture of HCl and HNO3 , respectively). The dendritic arm spacings were measured on transverse and longitudinal sections of samples extracted along the length of the DS castings. The triangle method [26] was used in order to perform measurements of the primary dendritic arm spacing and the intercept method [26] for secondary and tertiary dendritic arm spacings (about 50, 200 and 300 measurements for each selected position, respectively). Some samples of the examined alloys were selected and subjected to a deep etching technique with HCl for 2 min. Chemical analyses along the length of DS Al-11 wt%Si-5 wt%Ni alloy casting were performed in an X-Ray Fluorescence (XRF) spectrometer (Rigaku-RIX3100) with a view to checking the distribution of solute elements. The phases forming the alloy microstructure were determined
R. Kakitani, C.B. Cruz and T.S. Lima et al.
Materialia 7 (2019) 100358
Fig. 2. Experimental cooling curves along the length of: (a) Al-11 wt%Si, (b) Al-5 wt%Ni and (c) Al-11 wt%Si-5 wt%Ni alloys castings.
by X-Ray Diffraction (XRD), using a PAN analytical X’Pert PRO Materials Research Diffractometer XL, Cu-K𝛼 radiation with a wavelength of 0.15406 nm. The microstructure was also investigated by scanning electron microscopy with field emission gun (SEM-FEG) FEI Inspect F50 and by energy dispersive X-ray spectroscopy (EDS). Vickers microhardness tests were performed on the cross sections of the samples using a test load of 1 kgf (9.8 N) and a dwell time of 10 s. The adopted hardness value of a representative position was the average of at least 30 measurements on each sample.
Table 2 Different solidification experiments (#) with the Al-11 wt%Si-5 wt%Ni eutectic alloy: position of macrostructural transition and related solidification thermal parameters. Experiment
#1
#2
#3
Transition [mm] Ṫ [°C/s] vE [mm/s] G= Ṫ/vE [°C/mm]
55–65 1.5–1.3 0.35–0.31 4.3–4.2
63–78 1.6–1.3 0.46–0.42 3.5–3.1
85–100 1.3–1.1 0.22–0.19 5.9–5.8
3. Results and discussion Fig. 2 shows the experimental cooling curves recorded at different positions along the length of the DS Al-11 wt%Si, Al-5 wt%Ni and Al11 wt%Si-5 wt%Ni alloys castings, which were used to determine the solidification thermal parameters: cooling rate (Ṫ) and growth rate (v). The directional solidification setup provides a wide range of Ṫ and v values in a single experiment. For the binary alloys (Al-11 wt%Si and Al-5 wt%Ni) the thermal parameters were evaluated with respect to the liquidus temperature (TL ), while for Al-11 wt%Si-5 wt%Ni the eutectic temperature (TE ) was used. In this regard, the subscripts □L and □E correspond to the hypoeutectic and eutectic alloys, respectively. As the bottom of the mold was made of steel, the chemical affinity between Ni and Fe seems to have improved the alloy/mold wettability of the Al5 wt%Ni alloy casting as compared to the two other alloy castings, thus resulting in higher Ṫ and v profiles, as shown in Fig. 3. On the other hand, Si seems to be detrimental to the wettability, because the eutectic ternary alloy reached intermediate values of Ṫ and v (Fig. 3). Columnar grains prevailed along the length of the DS binary alloys castings, without any evidence of CET (Figs. 4 and 5). According to previous studies with Al–Ni [16] and Al–Si [17] alloys, the critical solidification cooling rate at which CET occurs for these alloys is about 0.16–0.17 °C/s, values not achieved along the length of the DS castings examined in the present work. The as-solidified microstructures of both alloys were composed by an 𝛼-Al matrix of dendritic morphology with the eutectic mixture located in the interdendritic region, with primary Si crystals also appearing in the microstructure of the Al-11 wt%Si alloy. Fig. 6a–d shows the eutectic microstructures of the Al-11 wt%Si alloy (𝛼Al matrix and plate-like Si) and the Al-5 wt%Ni alloy (𝛼-Al matrix and Al3 Ni fibers) obtained through conventional metallography, as well as
the microstructures that resulted from deep etching used to reveal the three-dimensional morphology of the phases. Initially characterized by columnar grains, the macrostructure of the ternary DS alloy casting evolved to a transition between 55 and 65 mm with respect to the cooled base of the casting (Fig. 7). Visually, this transition could not be associated with a sharp structural change from columnar to equiaxed grains, but rather some columnar grains have lost their original direction of growth and new grains nucleated ahead the columnar front. Two additional experiments with the same alloy were performed to prove the macrostructural transition observed in the initial experiment. Table 2 shows the regions in the castings of each experiment in which this transition occurred and the corresponding solidification thermal parameters. The cooling rate seems to be the most realistic thermal parameter determining the macrostructural transition since it encompasses the narrower range of values (1.3–1.6 °C/s) when Ṫ, vE and G are compared, that is, a critical cooling rate that is more than 7 times higher than the CET critical value determined in the literature for Al–Ni [16] and Al–Si [17] alloys. The solute concentration along the length of the Al-11 wt%Si-5 wt%Ni alloy casting, determined by X-Ray Fluorescence (XRF) (Fig. 8a), did not indicate the occurrence of macrosegregation, with Si and Ni concentrations quite close to the nominal composition of the alloy along the length of the whole casting. Since the Al-11 wt%Si-5 wt%Ni alloy has a eutectic composition [27–29], primary Si crystals and 𝛼-Al dendritic branches were formed surrounded by the eutectic phase (𝛼-Al+Si+Al3 Ni) in the columnar region, as can be seen in Fig. 7. In Fig. 8b, X-Ray Diffraction (XRD) patterns of samples extracted from different positions along the length
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Fig. 3. Experimental profiles of (a) cooling rate and (b) growth rate against the position from the cooled surface of the DS castings for the Al-11 wt%Si, Al-5 wt%Ni and Al-11 wt%Si-5 wt%Ni alloys.
of the casting confirmed the aforementioned phases and Scanning Electron Microscopy / Energy Dispersive X-Ray Spectroscopy (SEM/EDS) analyses (Fig. 9) indicated absence of supersaturated solid solution in the 𝛼-Al phase and the atomic compositions were close to the stoichiometric formulas of each phase. The Al3 Ni intermetallics has a plate-like morphology instead of fibers due to the faceted growth of Si that inhibits
the formation of regular eutectic structures [29]. Fig. 6e–f shows the eutectic phase of Al-11 wt%Si-5 wt%Ni (matrix of 𝛼-Al with plate-like Si and Al3 Ni). From the position 15 mm (Ṫ = 5.8 °C/s and vE = 0.76 mm/s), that is, with the decrease in the solidification kinetics, a certain fraction of the Al3 Ni phase starts to lose its plate-like morphology, evolving
Fig. 4. Macrostructure of the DS Al-11 wt%Si alloy casting and longitudinal microstructures along the length of the casting.
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Fig. 5. Macrostructure of the DS Al-5 wt%Ni alloy casting and longitudinal microstructures along the length of the casting. Fig. 6. Eutectic of a) Al-11 wt%Si alloy with primary crystal of Si and b) deep etched sample showing irregular plate-like Si; c) Al-5 wt%Ni alloy and d) deep etched sample showing Al3 Ni fibers; and e) Al-11 wt%Si-5 wt%Ni alloy: dark phase is Si and gray phase is Al3 Ni and f) deep etched sample showing irregular plate-like Si and Al3 Ni.
Fig. 7. Macrostructure of the DS Al-11 wt%Si-5 wt%Ni alloy casting and longitudinal microstructures along the length of the casting.
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Fig. 8. (a) Solute distribution along the length of the DS Al-11 wt%Si-5 wt%Ni alloy casting and (b) typical XRD patterns.
Fig. 9. SEM-EDS results for samples from positions: (a) 5 mm and (b) 50 mm along the length of the Al-11 wt%Si-5 wt%Ni alloy casting.
to a fishbone morphology, as shown in Fig. 10. In the experiment #2 with the Al-11 wt%Si-5 wt%Ni alloy, such morphological transition occurred for similar solidification thermal parameters: Ṫ = 5.7 °C/s and vE = 0.81 mm/s. This kind of morphology has been reported to occur also in alloys of the Al–Mg–Mn–Fe, Al–Si–Cu–Ni–Mg–Co and Al–Si–Sc systems [30–32], being generally associated with high cooling rates (watercooled molds), which favor the formation of metastable phases [30,33]. Since the aforementioned Al–Mg–Mn–Fe system does not have Si in its composition and even so the formation of Al3 Ni with fishbone morphology occurs, it seems that both phases (Si and Al3 Ni) grow independently, and that the prevailing factor for the formation of the fishbone morphology is the cooling rate. The fishbone Al3 Ni phase occupies a wider area of the microstructure as compared to that occupied by the plate-like
one, and this seems to be responsible for the increase in the intensity of peaks in the XRD patterns related to the Al3 Ni phase that is formed at low cooling rates (Fig. 8b). Fig. 11a shows the cooling curve of the eutectic Al-11 wt%Si5 wt%Ni alloy obtained under slow cooling conditions (air-cooled: Ṫ∼0.2 °C/s), showing a eutectic temperature of 568 °C, which is in good agreement with results from the literature [27–29]. The corresponding microstructure, shown in Fig. 11b, is formed by a mixture of eutectic, whose area fraction is predominant, and primary phases. In a Bridgman steady-state solidification study [27], only the eutectic mixture (𝛼-Al+Si+Al3 Ni) was reported to form the microstructure of eutectic Al–Si–Ni alloy samples that were grown both at constant G (5.82 K/mm) and different v (4.60 to 35.14 μm/s), and constant v
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Fig. 10. Morphological transition of Al3 Ni particles (gray phase) from plate-like to fishbone, with the decrease in the solidification cooling rate along the length of the Al-11 wt%Si-5 wt%Ni alloy casting.
(11.63 μm/s) and different G (2.11–5.82 K/mm). With G = 5.82 K/mm and increasing v (v> 173.61 μm/s), 𝛼-Al dendritic branches were also reported to occur, and once Ṫ=G∗ v, for Ṫ < 1.0 °C/s the microstructure was exclusively eutectic, and higher cooling rates induced the formation of dendrites. Therefore, similarly to the irregular Al–Si eutectic alloy, the Al–Si–Ni eutectic alloy, at the Al-rich side, has an asymmetric coupled zone in which a fully eutectic structure is obtained only for low Ṫ or vE , as indicated in Fig. 12. A deviation from the coupled zone results in an 𝛼-Al+E or E + X, where E is the eutectic and X a phase or an undefined structure. In the present work, a dendritic arrangement is justified by the coupled zone, however, there is no certainty as to the determinant factor that led to the formation of primary crystals of Si and fishbone Al3 Ni, that is, whether Ṫ or vE or even both. What is certain is that these phases arose due to the non-equilibrium conditions typical of transient solidification. The macrostructure transition of the eutectic Al–Si–Ni alloy occurred close to 1.0 °C/s, which is also the critical value associated with the microstructural transition between fully eutectic and eutectic/𝛼-Al den-
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drites, and this is probably connected to the nucleation of eutectic grains ahead the solidification front, which blocked the columnar growth. Some literature studies focusing on the nucleation of irregular eutectic Si have reported the occurrence of eutectic grains in near-eutectic Al– Si alloys [34–36]. The main theory is that a heterogeneous nucleation occurs induced by impurities, such as Fe and P that form intermetallic particles, or by 𝛼-Al dendrites [36]. Moreover, it has been reported that in ternary Al–Si–(Cu,Mg) alloys the number of eutectic grains has a bias dependence on the added element: Mg up to 0.5 wt% and Cu additions increase the density of eutectic nuclei [34]. According to a previous study [34], different ternary alloying elements added to the Al–Si alloy system imply different diffusion and partition coefficients (k). A faster diffusion rate and smaller k value affects the solute gradient at the solid/liquid interface and increases segregation, respectively, thus favoring the rate of eutectic nucleation. Thus, since the Al-11 wt%Si-5 wt%Ni is an irregular eutectic alloy, at high solidification cooling rates primary 𝛼-Al dendrites grew preferentially in relation to the eutectic phase [37], and the decrease in Ṫ favors the precipitation of eutectic grains with the simultaneous formation of dendrites at the grain boundaries [38]. Moreover, the Ni content may aid the precipitation of eutectic grains, considering that the partition coefficient of Ni in Al is low (k = 0.007) [39]. Kaya and Aker [40] added 2 wt%Ni to the eutectic Al-12.6 wt% alloy and reported a reduction in the Si interflake spacing that seems to be induced by higher number of eutectic nuclei. It is suggested that the addition of Ni has contributed to increase the rate of eutectic nucleation. Alloys with a eutectic composition have a resulting microstructure essentially constituted by (i) a eutectic arrangement, or (ii) eutectic mixture plus the presence of the primary phase of the preponderant element [1]. The coupled zone theory [1] reports that one or other of these morphologies will depend on the product between the thermal gradient and the growth rate, i.e., on the cooling rate. Hence, the further away from the equilibrium conditions the solidification occurs, more susceptible to the formation of the second type of morphology the system will be. Fig. 13 shows the experimental dendrite arm spacings (𝜆1 –3 ) determined in the present study as a function of cooling rate and growth rate, with the error bars representing the standard deviation. Experimental power laws are proposed having exponents that are in agreement with those usually reported in the literature [41,42]. As expected, the higher Ṫ the smaller the dendrite arm spacing. Single experimental power laws can correlate the primary dendrite arm spacing (𝜆1 ) with either Ṫ or vL /vE for the Al-11 wt%Si-5 wt%Ni and Al-11 wt%Si alloys (Fig. 13a–b).
Fig. 11. (a) Cooling curve of an air-cooled Al-11 wt%Si-5 wt%Ni alloy (Ṫ ∼ 0.2 °C/s) and (b) corresponding microstructure.
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Fig. 12. Asymmetric coupled zone of the Al-11 wt%Si-5 wt%Ni eutectic as a function of Ṫ or vE , where E is the eutectic and X a phase or an undefined structure.
Different power laws are proposed to represent the Al-5 wt%Ni alloy, which has a more refined range of 𝜆1 values. The exponents −0.55 and −1.1 seem to be appropriate to relate 𝜆1 with the cooling rate and the growth rate, respectively. For the secondary dendrite arm spacing (𝜆2 ), single experimental laws (Fig. 13c–d) represent the evolution of 𝜆2 for the Al-11 wt%Si-5 wt%Ni and Al-5 wt%Ni alloys with either Ṫ or vL /vE , suggesting that Ni has a refining role on 𝜆2 , as compared to the results of the Al-11 wt%Si alloy. For the Al-11 wt%Si alloy, the same exponents derived for the 𝜆1 power laws, can also be applied to the experimental power laws relating the evolution of the tertiary dendrite arm spacing (𝜆3 ) with the solidification thermal parameters, as shown in Fig. 13e–f. When the −0.55 and −1.1 exponents were applied to correlate 𝜆3 with Ṫ and v, respectively, for the Al-11 wt%Si-5 wt%Ni alloy, the fit did not include most of the experimental data, suggesting that different exponents should be used. Theoretical models for the prediction of dendritic growth during solidification of multicomponent alloys as a function of solidification growth rate and cooling rate, cannot be found in the literature because of the difficulty of the task. On the other hand, different exponents for experimental growth laws in ternary alloy systems, such as Al–Si–Mg [43], Al–Cu–Ni [44] and Al–Si–Cu [45], have been reported when compared to the usual ones experimentally determined for binary alloys. Based on previous studies on Sn–Bi–(Cu,Ag) and Al-11 wt%Si-(3,4.5)wt%Cu ternary alloys [46,47], experimental power laws using −0.25 and −0.50 exponents were shown to be more appropriate to relate 𝜆3 to Ṫ and vE , respectively, as can be seen in Fig. 13e–f. No evidence of tertiary dendritic branches was found along the microstructure of the Al-5 wt%Ni alloy. These distinct behaviors on the evolution of dendrite arm spacings with Ṫ and vL/E , observed in Fig. 13, could be related to the differences in the amount of Si (11 wt%) and Ni (5 wt%) in the Al–Si–Ni alloy composition and to their densities with respect to that of liquid Al, which are: Al, Si and Ni: 2385 kg/m3 , 2340 kg/m3 and 8900 kg/m3 , respectively [48]. At the beginning of solidification of the eutectic Al–Si–Ni alloy, the rejection of Si to the molten liquid is more remarkable than that of Ni due to the higher amount of Si in the alloy composition. It is worth noting that the alloy solidified vertically upwards, i.e. the growth direction is opposite to that of the gravity vector. As Si has a density that is even lower than that of liquid Al, the Si rejected during solidification has a trend to float and concentrate along the solidification front provoking instabilities at this interface that lead to the formation of primary dendrite stalks, as shown schematically at the top row of Fig. 14. The enrichment of Si in the liquid around these dendritic stalks, could lead to local con-
centrations that could surpass the composition of the Al–Si eutectic, thus creating conditions for precipitation of primary crystals of Si. As Ni has a higher density than liquid Al, when rejected at the solidification interface it tends to move downward in the liquid channels located between the primary dendritic arms, thus provoking gradual lateral instabilities leading to the formation of secondary dendritic branches (central row of Fig. 14). In this sense, the growth of primary branches seems to be mainly controlled by the rejection of Si, whereas the secondary branches by the rejection of Ni, explaining the similar growth laws in Fig. 13: 𝜆1 (Al-11 wt%Si-5 wt%Ni and Al-11 wt%Si -Fig. 13a and b) and 𝜆2 (Al11 wt%Si-5 wt%Ni and Al-5 wt%Ni -Fig. 13c and d). The development of tertiary dendritic branches would be caused by the simultaneous action of atoms of Si and Ni (bottom row of Fig. 14). At regions in the Al-11 wt%Si-5 wt%Ni alloy casting associated with higher solidification kinetics, the rejected Si atoms could also initiate local instabilities leading to the formation of tertiary branches, and the local tertiary spacings would be like those of the Al-11 wt%Si alloy. With the decrease in the solidification kinetics, the growth of tertiary branches would be modified by the presence of Ni atoms, and the growth law relating 𝜆3 to Ṫ and vE for the Al-11 wt%Si-5 wt%Ni alloy casting, would progressively depart from the trend observed for the 𝜆3 results of the Al-11 wt%Si alloy casting (Fig. 13e–f). Fig. 15 shows the results of Vickers microhardness measured along the length of the DS castings against a representative microstructural spacing, 𝜆2 , i.e. the common dendrite spacing of higher order measured in all examined alloys (tertiary spacings have not been detected in the Al-5 wt%Ni alloy), which is reported to be responsible for a closer interaction among dendritic matrix, eutectic mixture and second phases [49]. The isolated reinforcing role of plate-like Si particles and Al3 Ni fibers for the binary Al-11 wt%Si and Al-5 wt%Ni alloys, respectively, seems to have similar performance in the blockage of dislocations all along the length of the DS castings, without any the influence of the scale of the dendritic Al-rich matrix, resulting in an average hardness of 48 HV. However, for the Al-11 wt%Si-5 wt%Ni alloy casting, the simultaneous interaction of both reinforcing phases located in the interdendritic spacings and the scale of the secondary dendritic arms is conducive to a remarkable increase in hardness with the decrease in 𝜆2 . A Hall–Petch-type correlation is proposed relating HV to 𝜆2 −1/2 . Along a comparative common range of 𝜆2 values of the ternary and binary alloys, the hardness of the ternary alloy increases from 46% to 77% with the decrease in 𝜆2 (increase in 𝜆2 −1/2 ) as compared to the mean hardness of the binary alloys (HV = 48).
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Fig. 14. Schematic representation of possible stages of formation of dendritic branches during the transient solidification of an Al-11 wt%Si-5 wt%Ni alloy, where S is the solid, L is the liquid and SF is the solidification front.
Fig. 13. Evolution of: (a, b) primary, (c, d) secondary and (e, f) tertiary dendritic spacings as a function of the growth rate (vE / L ) and of the cooling rate (Ṫ) for the Al-11 wt%Si-5 wt%Ni, Al-11 wt%Si and Al-5 wt%Ni alloys castings. R2 is the coefficient of determination.
Fig. 15. Vickers microhardness as a function of 𝜆2 −1/2 for the Al-11 wt%Si5 wt%Ni, Al-11 wt%Si and Al-5 wt%Ni alloys castings. R2 is the coefficient of determination.
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4. Conclusions The following conclusions can be drawn from the present experimental study: • A columnar to equiaxed transition (CET) was shown to occur for the DS Al-11 wt%Si-5 wt%Ni alloy casting under solidification cooling rates in the range 1.3–1.6 °C/s. These critical cooling rate values are more than 7 times higher than the CET critical values determined in the literature for Al–Ni and Al–Si alloys. • The microstructure of the DS Al-11 wt%Si-5 wt%Ni eutectic alloy casting was characterized by primary Si crystals; 𝛼-Al dendritic matrix and 𝛼-Al+Si+Al3 Ni eutectic due to non-equilibrium transient solidification conditions. The Al3 Ni intermetallics was shown to have a plate-like morphology instead of fibers, however, for cooling rates Ṫ <5.7 °C/s and growth rates vE < 0.81 mm/s, it was shown to lose its plate-like morphology, evolving to a fishbone morphology. • Experimental growth laws relating the primary (𝜆1 ), secondary (𝜆2 ) and tertiary (𝜆3 ) dendritic spacings of the ternary Al-11 wt%Si5 wt%Ni and the corresponding binary Al-11 wt%Si and Al-5 wt%Ni alloys, to both the solidification cooling rate (Ṫ) and the growth rate (vE/L ) were shown to be given by: 𝜆1 = 112 vE/ L − 1.1 ; 𝜆1 = 405 Ṫ −0.55 for the Al-11 wt%Si-5 wt%Ni and Al-11 wt%Si alloys 𝜆1 = 54 vL −1.1 ; 𝜆1 = 149 Ṫ −0.55 for the Al-5 wt%Ni alloy 𝜆2 = 5.2 vE/ L − 2/3 ; 𝜆2 = 10.9 Ṫ −1/3 for the Al-11 wt%Si-5 wt%Ni and Al-5 wt%Ni alloys 𝜆2 = 9.1 vL −2/3 ; 𝜆2 = 19.2 Ṫ −1/3 for the Al-11 wt%Si alloy 𝜆3 = 5.5 vE −0.50 ; 𝜆3 = 9.9 Ṫ −0.25 for the Al-11 wt%Si-5 wt%Ni alloy 𝜆3 = 6.3 vL −1.1 ; 𝜆3 = 21.8 Ṫ −0.55 for the Al-11 wt%Si alloy where 𝜆1,2,3 [μm]; vE/L [mm/s] and Ṫ [ ̊C/s] • The Vickers microhardness (HV) along the length of both Al11 wt%Si and Al-5 wt%Ni alloys castings was shown not to be affected by 𝜆2 , with a mean HV = 48. In contrast, a Hall–Petch-type correlation has been proposed relating HV to 𝜆2 for the Al-11 wt%Si5 wt%Ni alloy casting: HV = 36+ 103 𝜆2 − 1/2 , representing a 46–77% increase with the decrease in 𝜆2 as compared to the hardness of the binary alloys. Acknowledgments The authors are grateful to FAPESP (São Paulo Research Foundation - grant 2017/15158-0), CNPq -National Council for Scientific and Technological Development and CAPES - Coordenação de Aperfeiçoamento de Pessoal de Nível Superior - Brasil (Finance Code 001) for their financial support and Brazilian Nanotechnology National Laboratory LNNano for the use of its facilities. Declaration of interest None. Supplementary material Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.mtla.2019.100358. References [1] W. Kurz, D.J. Fisher, Dendrite growth in eutectic alloys: the coupled zone, Int. Mater. Rev. 24 (1979) 177–204. https://doi.org/10.1179/imtr.1979.24.1.177. [2] L. Rátkai, A. Szállás, T. Pusztai, T. Mohri, L. Gránásy, Ternary eutectic dendrites: pattern formation and scaling properties, J. Chem. Phys. 142 (2015) 154501. https:// doi.org/10.1063/1.4917201.
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