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Hardening behavior of Al-0.25Sc and Al-0.25Sc-0.12Zr alloys during isothermal annealing
Journal Pre-proof Hardening behavior of Al-0.25Sc and Al-0.25Sc-0.12Zr alloys during isothermal annealing Yuhong Luo, Qinglin Pan, Yuqiao Sun, Shuhui Liu, Yuanwei Sun, Liang Long, Xinyu Li, Xiaoping Wang, Mengjia Li PII:
S0925-8388(19)34168-4
DOI:
https://doi.org/10.1016/j.jallcom.2019.152922
Reference:
JALCOM 152922
To appear in:
Journal of Alloys and Compounds
Received Date: 23 July 2019 Revised Date:
14 October 2019
Accepted Date: 4 November 2019
Please cite this article as: Y. Luo, Q. Pan, Y. Sun, S. Liu, Y. Sun, L. Long, X. Li, X. Wang, M. Li, Hardening behavior of Al-0.25Sc and Al-0.25Sc-0.12Zr alloys during isothermal annealing, Journal of Alloys and Compounds (2019), doi: https://doi.org/10.1016/j.jallcom.2019.152922. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.
Hardening behavior of Al-0.25Sc and Al-0.25Sc-0.12Zr Alloys during isothermal annealing Yuhong Luoa, Qinglin Pana,b, Yuqiao Suna, Shuhui Liub, Yuanwei Suna, Liang Longb, Xinyu Lib, Xiaoping Wangc, Mengjia Lia* a
School of Materials Science and Engineering, Central South University, Changsha 410083, china
b
Light Alloy Research Institute, Central South University, Changsha 410083, china c
Hunan Dongfang scandium industry Co., Ltd. Changsha 410083, china
Abstract The hardening behavior of Al-0.25Sc and Al-0.25-0.12(wt.%) alloys during isothermal annealing were studied by transmission electron microscopy (TEM), scanning electron microscopy (SEM), optical microscopy (OM), and hardness measurements. The results show that Sccontaining aluminum alloy contained fine dispersion of spherical secondary Al3Sc particles, which developed during the annealing process. Al3(Sc, Zr) precipitates in the Al-0.25Sc-0.12Zr alloys are core-shell structure. Different isothermal annealing temperatures can lead to different strengthening effects on the alloys. Al-0.25Sc and Al-0.25Sc0.12Zr alloys both exhibit obvious annealing hardening behavior, and the ternary Al-0.25Sc-0.12Zr alloy shows better thermal stability than Al0.25Sc alloy. Key words: hardening behavior, precipitation strengthening, thermal stability, secondary Al3Sc particles, Sc-containing aluminum alloy 1. Introduction The addition of Sc can greatly improve the mechanical properties, tensile strength, plasticity, weld ability, corrosion resistance of aluminum alloy [1-3]. It has been found that Sc-containing industrial aluminum alloys can show annealing hardening behavior. The addition of Sc and Zr can greatly improves the thermal stability of the alloys during annealing. Supersaturated solid solution containing Al and Sc in the alloy during *
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solidification. The crystal structure of primary Al3Sc phase particles is same to the α-Al matrix, the lattice mismatch between the primary phase particles and the aluminum matrix is 0.27%, and these precipitates act as nucleation sites for Al matrix, which can lead to the grain refining effect in Al alloys [4]. In the subsequent annealing process, secondary Al3Sc particles (L12 type) are precipitated, which can produce precipitation strengthening effect and pin the grain boundaries [5,6]. Al3(Sc, Zr) particles with L12 structure can precipitate in the Al alloys after adding trace amounts of Sc and Zr [7-9]. The lattice parameters difference between Al3(Sc, Zr) and Al3Sc particles are negligible. The precipitation phase of Al3(Sc, Zr) can maintain the coherent relationship with the matrix of α-Al during annealing, which can strongly hinder the movements of dislocation and grain boundaries, leading to higher strength and recrystallization temperature of the alloy. Lohar et.al [4] reported that the combined addition of Sc and Zr can maintain the stability of the particles at a high temperature (reducing the coarsening rate of the particles) compared to adding Sc alone, which can improve the grain refinement effect and thermal stability of the Al alloys. The microstructure evolution of the Sc-containing aluminum alloy during annealing can lead to the variation mechanical properties in the alloy [1013]. The annealing hardening behavior has been widely reported in Sccontaining aluminum alloys [14-19]. However, there is no systematic discussion on the mechanism of microscopic organization on the precipitation strengthening. In order to reduce the influence of other factors on the annealing hardening behavior of Sc-containing aluminum alloys, Al-Sc binary alloy and Al-Sc-Zr ternary alloy are used for detailed experiment. This paper aims to research the effect of strengthening mechanism of precipitation phase and coarsening of Al3Sc, Al3(Sc, Zr) precipitated particles on the annealing hardening behavior of Al-Sc and Al-Sc-Zr alloys. 2. Experimental 2.1 Material preparation Aluminum alloys (Al-0.25Sc, Al-0.25Sc-0.12 wt. %Zr) were prepared by melting Al with commercial purity and Al-2.2Sc and Al4.0Zr master alloy in a muffle furnace with graphite crucible. Firstly, the ingots were heated to 780°C, then Al-4.0Zr master alloy and Al-2.2Sc master alloy were added. After the molten alloys were thorough stirring, the molten alloys were allowed to stand for 5 min. Then the melt is
poured into the water-cooling copper mold. The samples are annealed at 260°C, 290°C, 340°C, and 390°C for 0.1 h, 0.5 h, 1 h, 3 h, 5 h, 10 h, 16 h, and 24 h, respectively. 2.2 Hardness measurements The hardness test used a 401 MVD TM digital micro Vickers hardness tester with a load of 0.3 KN and time of indentation of 15 s. The hardness values reported are taken from the average of the five measurements values. Hardness measurements were prepared by cutting sample pieces of 12 mm*12 mm*4 mm (the test surface size was 12mm*12mm) from the annealing treated materials and by mechanically grinding them. 2.4 Microstructure characterization The microstructure of the as-cast Al-0.25Sc and Al-0.25Sc-0.12Zr alloys were observed at 340°C for 3 hours and 24 hours, respectively, under a metallographic microscope with the instrument model Leica DFC295. The samples were cut by electrical discharge machining, smoothed by water sanding paper and metallographic sandpaper, and mechanically polished. Then the polished samples were treated in a coating liquid (30ml of fluoroboric acid + 200 ml of distilled water), the voltage is 15~20V, the current is 0.1~0.5mA, and the operation time is 1~2 min at room temperature. The grain shape and size of alloy samples were observed by metallographic microscope. A Tecnai G2 F20 transmission electron microscope (TEM) with an accelerating voltage of 200 KV was used for investigating the general microstructure and the distribution of Al3Sc or Al3(Sc, Zr) precipitates. Thin foils for TEM were prepared by cutting sample pieces from the annealing treated materials and by mechanically grinding them down to ~100um. Discs with 3mm in diameter were punched and dimpled down to 60um, followed by twin-jet electro polishing at 100mA. A solution of 1/3(volume fraction) nitric acid and 2/3 methanol at -30°C was applied during the electro polishing [20,21]. The average radius size and volume fraction of the precipitated phases in the alloys were statistically analyzed by Image-J software [22]. In a large number of TEM dark field phase diagrams, three TEM dark phase images which contain over 100 precipitates were chosen for obtain the precipitate average radius. The distribution range of the precipitates can be obtained. The detailed method to calculate the size of precipitates is shown as below. Import the image into the software first, set the ruler, and then adjust the threshold
color of the image then use the analysis precipitate function in the software to analyze and finally get the volume fraction. 3. Results 3.1 Microhardness of Al-0.25Sc and Al-0.25Sc-0.12Zr alloys during annealing The hardness change curves of as-cast Al-0.25Sc alloys during annealing at different temperatures is shown in Fig.1. It can be found that the alloy’s hardness is significantly increased after annealing. Under the annealing condition of 260°C, the hardness value of the alloy increases with the annealing time extended from 0 to 24 hours. The hardness value of the alloy reach 68.1 HV after annealing for 24 hours. When the annealing temperature is 290°C, the hardness of the alloy reaches to the highest value after 0.5 hour, and the peak hardness value is 68.3 HV. The hardness of the alloy does not change much when further increase the annealing time. During annealing at 340°C, the hardness of the alloy reaches a peak value of 63 HV at 0.5 hour. When increase the annealing time to 10 hours, the hardness of the alloy nearly does not change, while the hardness of the alloy begins to decrease after annealing for 16 hours. When annealing at 390°C, the alloy peak at 0.5 hours and the hardness begins to decrease after that. It can be found that the hardening effect of the alloy is not so obviously after annealing at higher temperatures. The peak hardness of the as-cast alloy after 390°C treatment is much lower than other samples. The hardness curves of the as-cast Al-0.25Sc-0.12Zr alloys during annealing at different temperatures is shown in Fig.2. The hardness of alloys is significantly improved after annealing. When the annealing is 260°C, the variation of hardness of the alloy is similar to that of the Al0.25Sc alloy, the hardness increases with the annealing time. Other curves exhibit roughly the same trend of hardness change. Initially, the hardness of the alloys increases with increasing annealing time sharply from the starting point until reaching the highest value. With the prolongation of annealing time, the hardness values of the three alloys do not decline obviously. When the alloy is annealed at 290°C, the hardness of the alloy reaches its peak at 3 h, and the peak value is 70.7 HV. When the annealing treatment temperature is 340°C and 390°C, the alloys reach a peak at 0.5 h, the peak values of the alloys are 64.9 HV and 65.4 HV, respectively.
3.2 Microstructure of Al-0.25Sc and Al-0.25Sc-0.12Zr alloys Image J software is used to measure the average size of crystal grains by the line cutting method. By cutting the grain with a straight line of a certain length, it can be obtained that the average length of the straight line can intercept one grain, that is, the average grain size. If the length of the straight line is L and the total number of cut grains is N, the average grain size can be expressed as (1): Average Grain Size =
N
(1)
The optical microscopy investigations from the anodized specimen of the Al-0.25Sc alloys is shown in Fig. 3. It reveals that the average grain size is about 88 µm, the shape of the equiaxed grains is not uniform and the grain size is inconsistent in the structure after annealing at 340°C for 3h (Fig.3(a)). When alloy is annealed for 24 hours, the shape and size of the equiaxed crystals are slightly more uniform in the structure, and the grain size is 92 µm (Fig.3 (b)). The grain size of the alloy annealing at 340°C for 3 hours decreases to 58 µm after the addition of Zr, and the dendrites disappears, completely replaced by equiaxed grains (Fig.3(c)). The average grain size does not change much with annealing time increasing, the average grain size is about 55 µm (Fig.3 (d)). It can be found that the microstructure of the alloys is significantly refine by the addition of Zr. It is indicating that the combined addition of Sc and Zr can lead to better refinement for microstructure of the aluminum alloy. Dark-field TEM micrographs of precipitates in Al-0.25Sc alloys annealing at 290~390°C are shown in Fig. 4. It can be seen from the Fig.4 that many white approximately spherical precipitates are dispersed in the α-Al matrix. The selected area diffraction (SAD) pattern is inserted in Fig. 4(b). As the annealing temperature increases, the radius of Al3Sc particles increases with the annealing time, and its number density of precipitates decreases significantly. After annealing at 290°C for 3 h (Fig. 4 (a)), Al3Sc particles are uniformly dispersed in a large amount, and the average radius of the precipitates is about 1.5 nm. After annealing at 340°C for 3 h (Fig. 4 (c)), the number density of Al3Sc precipitates decrease, the average radius of the precipitates increase to about 2.5 nm, when the annealing time is extended to 24 h (Fig. 4 (d)) the Al3Sc precipitates are obviously coarsening, the average radius of the precipitates is increase to about 4 nm and the number density of Al3Sc precipitates is reduces. After annealing at 390°C for 3 h (Fig. 4 (e)), the
number of precipitated phases of Al3Sc become very small and the average radius is increase. The dark-field TEM micrographs of Al3(Sc, Zr) precipitates in Al0.25Sc-0.12Zr alloys are shown in Fig.5. After annealing at 340°C for 24 h (Fig. 5 (a)), the average radius of the precipitated phase is about 1.5 nm. After annealing at 390°C for 3 h (Fig. 5 (c)), the number density of precipitates is higher, and the average radius of the precipitated phase is about 2 nm. The particle radius does not increase much when the annealing temperature increases. The particles of Al3(Sc, Zr) shows very excellent resistance towards coarsening. Comparing the TEM dark field phase of the two alloys, it can be seen that the Al3(Sc, Zr) particles in the Al-0.25Sc-0.12Zr alloy have higher number density and smaller size than Al3Sc in the Al-0.25Sc alloy. The TEM bright field images of Al-0.25Sc alloy and Al-0.25Sc0.12Zr alloy after annealing at 290~390°C are shown in Fig.6 and Fig.7, respectively. In general, the coherency between Al3Sc precipitate phase and the matrix is observed by HREM analysis, but this method is too inefficient in analyzing many particle coherences. Therefore, Iwanura and Miura et al. [23] proposed the use of Ashby-Brown contrast to determine the coherent relationship of the precipitated phase to the matrix. It is observed by TEM that fine nanoscale particles with bean-shaped disperse in annealing samples. The direction of the no-contrast line of the beanlike center is perpendicular to the operation vector g. This bean-like contrast is generally observed in alloys with Ashby-Brown characteristics, indicating that the precipitated phase is coherent with the matrix. The HRTEM micrograph of Al3Sc precipitates are shown in Fig.8. In the HRTEM micrograph, no dislocations can be observed at the interface, and it can be clearly seen that the particle maintains a coherent relationship with the matrix. The radius of Al3Sc precipitates after annealing at 290°C for 3 hours are about 2.5 nm (Fig. 8(a)), and the radius increase to about 5 nm after the alloy annealing at 390°C for 3 hours (Fig. 8 (b)). Figure 9 shows the HRTEM micrograph of Al3(Sc, Zr) precipitates in the Al-0.25Sc-0.12Zr alloy. The radius of Al3(Sc, Zr) precipitates are about 1.5 nm of the samples annealing at 340°C for 24 hours (Fig. 9 (a)), and the radius of the Al3(Sc, Zr)precipitates increase to about 2 nm after the alloy is annealed for 3 hours at 390°C (Fig. 9 (b)).
4. Discussions 4.1 Microhardness evolution It can be seen from Fig. 1 and Fig. 2 that both alloys have annealing hardening behavior during the annealing process, which due to the addition of Sc. When annealing at low temperatures, the hardness of the alloy increased slower than at higher annealing temperature, which is due to that the precipitation of Al3Sc and Al3(Sc, Zr) precipitates is controlled by the diffusion of Sc and Zr atoms [24,25]. Studies on the annealing of aluminum alloys containing Sc at a content of 0.2~0.5% at temperatures below 340°C indicate that the increase in hardness of the alloy is due to the formation of nano sized Al3Sc particles during annealing [7]. After reaching the hardness peak during annealing, the hardness begins to decrease with time, as the Al3Sc particles begin to coarsen (~40 nm) [26]. According to the related report [7], it can be known that the heterogeneous nucleation of Al3Sc particles is promoted under high temperature annealing, thus resulting in a lower volume fraction and a larger radius size of the particles, so that the hardness value of the alloy is lowered after annealing at high temperature. In the solidification process of Sc-containing aluminum alloy, primary Al3Sc particles are precipitated. Since the lattice constant of the precipitated Al3Sc particles does not differ much from the lattice constant of the Al matrix, the phase of the primary Al3Sc can provide more nucleation point during solidification of the alloy, which can significantly refine the grains of the alloy. In the subsequent annealing process, the alloy will precipitate a large amount of dispersed, fine secondary Al3Sc particles, which can lead to precipitation strengthening. However, as the annealing temperature increases, the secondary Al3Sc particles begin to grow and coarsen, so the hardness value of the alloys begins to decrease. Because the Al3Sc particles have a very high coherency mismatch, it is clear that lattice strain can prevent the movement of dislocation and prevent grain growth. However, under higher temperature annealing conditions, the coarsening of Al3Sc particles is unavoidable, and the weakening of Al-0.25Sc alloy becomes more and more obvious. When Sc and Zr elements are added into the aluminum alloy, the Al3Sc particles will precipitate first. As the annealing time increase, Zr will be enriched in the outer layer of Al3Sc to form the core-shell structure of Al3(Sc, Zr) [27-30]. Tolley et al. [4] used the EDS in transmission electron microscopy to analyze the precipitation phase of the Al-0.61Sc-0.40Zr(wt.%) alloy during annealing. The results show that the precipitated phase is Al3(Sc, Zr) phase, the structure of the precipitated phase particles is core rich in Sc, and the shell is rich in Zr. Forbord et al. [24] used a three-dimensional atom probe to study the
formation process of Al3(Sc, Zr) precipitated phase in Al-Sc-Zr alloy. The results show that the Al3(Sc, Zr) precipitated phase mainly includes Al and Sc atoms in the early nucleation, while the Zr atom is concentrated on the precipitated phase at a later stage. Since the diffusion rate of Zr in Al is lower than Sc and mainly distributed in the outer layer of Al3(Sc, Zr), the growth rate of the precipitated phase is slowed down, and the thermal stability of the Al alloy is effectively improved [13]. Therefore, the hardness of the Zr-containing ternary alloy decrease slowly with time at higher annealing temperature. It can be seen from the dark field phase diagram of the alloys (Fig. 4) that the density and radius of the Al3Sc particles are different. The differences are caused by different annealing conditions. Strengthening effect is influenced by the radius and density of precipitated phase. As a result, the hardness of the alloys changed. Sc can be melted into Al to form a supersaturated solid solution. In the subsequent thermal processing and heat treatment, the supersaturated solid solution will overcome the energy barrier and dissolvent decomposition forming an Al3Sc precipitation phase. The Al3Sc precipitation phase is closely related to the heat treatment temperature. The higher the heat treatment temperature, the larger the driving force of the second phase precipitation, and the faster the precipitation rate of the Al3Sc precipitation phase, and the number density will decrease. When the holding time is extended or the annealing temperature is increased, the precipitation phase of Al3Sc will become larger. The nucleation and coarsening of the precipitated phase are greatly depended on the temperature. This process can be explained by the phase change driving force (2): Fαβ = −
1 Cβ − Cα C ∙ ∙ RT ln Vβ 1 − Cα Cα
(2)
Where Vβ is the volume of the β phase （ m3/g ） , C0 is the concentration of the B component in the supersaturated solid solution, Cα is the concentration of the B component in the equilibrium α phase, and Cβ is the equilibrium of the B component in the β phase. R is the ideal gas constant and T is the absolute temperature at which the phase change occurs. It can be seen that the atomic diffusion rate increases with temperature, the phase change driving force also increases, so the precipitated phase is more likely to precipitate out from the matrix, as a result the higher the annealing temperature of the alloy, the shorter the peak time. However, the number density of precipitation reduced when increase the annealing temperature, which leads to the decrease of the peak value of the alloys when increase the annealing temperature. The
higher the temperature, the faster the precipitation phase grows, moreover, the strengthening of the alloy also decay faster. Combined with the particle size distribution diagram in Fig. 10, it can be seen that the size of the Al3Sc particles is larger than Al3(Sc, Zr) particles in general. Therefore, the precipitation strengthening effect of Al3(Sc, Zr) particles is larger than Al3Sc. The lattice structure of Al3(Sc,Zr) belongs to L12 structure, which lattice parameters is slightly different from Al3Sc particles, resulting in slower particle coarsening rate. Forbord et al. [24] and Tolley et al. [31] suggested that the presence of Zr weakens the phenomenon of particle coarsening, and the size of the precipitates are smaller, with larger quantity. Forbord et al. [24] used atom probes to study the formation of Al3(Sc, Zr) precipitates in Al-Sc-Zr alloys. It shows that Al3(Sc, Zr) precipitates mainly include Al and Sc in the early nucleation. The Zr atom is concentrated on the precipitation phase after that. The addition of Zr also increases the supersaturating of the aluminum-based solid solution, which in turn increases the dynamic and quantitative density of the precipitated phase of Al3(Sc, Zr). Since the resistance towards coarsening performance of the Al3(Sc, Zr) particles is excellent, the precipitates after annealing remain fine and dispersed, so the hardness of alloys can maintain the peak value when further increase annealing time. It can be concluded that the addition of Zr element to the alloy can lead to better strengthening properties than the alloy with addition of Sc alone. 4.2 Strengthening mechanism of precipitation phase The evolution of microhardness is related to the growth of precipitated phase and strengthening mechanism. There are two main types of strengthening mechanisms for the precipitation phase: 1) shearing mechanism; 2) the strengthening of the Orowan bypass mechanism. When the shearing mechanism dominates, the yield strength increment is proportional to the square of the precipitation phase radius. The larger the precipitation phase size is, the more obvious the strengthening effect will be. When the Orowan bypass mechanism dominates, the increase in yield strength is inversely proportional to the radius of the precipitated phase. The domination of mechanism depends on the radius of the precipitation phase. When the radius of the precipitation phase is small, the shearing mechanism is the main strengthening mechanism. The yield strength increment can be estimated based on order strengthening ∆σord, modulus mismatch strengthening ∆σmod, and coherency strengthening ∆σcoh.
The order strengthening t ∆σord can be calculated according to the equation (3) [29,33]:
γ$%& 3π∅ +⁄, (3) ∆σ = 0.81M 2b 8 Where M=3.06 is the Taylor mean matrix orientation factor for Al, γ$%& =0.5 J/m2 is an average value of the Al3Sc antiphase boundary (APB) energy for the (111) plane, b=0.286 nm is the magnitude of the Al Burgers vector, and ∅ is the volume fraction of the precipitated phase. The modulus mismatch strengthening ∆σmod can be calculated according to the equation (4) [32,33]:
2∅ + , r 5.⁄,6+ (4) ∆σ. = 0.0055M0∆G2 b3 4 Gb , b Where G=25.4 GPa is the shear modulus of Al matrix, ∆G is the difference between the shear modulus of the precipitated phase and that of Al matrix, and the shear modulus of the precipitation phase of Al3Sc and Al3(Sc,Zr) is 68 GPa, r is the precipitation phase average radius, m = 0.85 is a constant. +⁄,
⁄
The coherency strengthening ∆σcoh can be calculated according to the equation (5) [32,33]:
r∅ + , (5) ∆σ7 8 = Mαε 0.5Gb Where αε = 2.6 is a constant and ε=δ 2⁄3 is the mismatch parameter, δ=1.3% is the lattice mismatch for the Al3(Sc1-xZrx) precipitated, which is estimated from the composition-dependent lattice parameters. Coherency strengthening occurs simultaneously with modulus mismatch strengthening. 0Gε25⁄,
⁄
The yield strength increment caused by the shearing mechanism increases with the average radius r until the strengthening mechanism is switched to the Orowan bypass mechanism. The Orowan bypass mechanism strengthening ∆σor can be calculated according to the equation below (6) [32,33]: =
r<2⁄3 ln ; = b 0.4MGb
(6)
λ π√1 − ν Where ν = 0.34 is the Poisson's ratio of Al and λ is the spacing of the edge-to-edge interprecipitate spacing (7) [32,33]: ∆σ
3π λ = >? − 1.64A r 4ϕ
(7)
The above model of yield strength increment is used to calculate the theoretical yield strength increment and compare with the experimental values. The experimental yield strength increment is estimated by using ∆HV/3 [6,32,34], where ∆HV is the microhardness increment of the ascast alloy after annealing. Which formula is valid for aluminum alloys, except pure Al [6]. Generally, the unit of microhardness increment is HV, and the unit of yield strength is MPa. For the sake of comparison, the unit of microhardness is converted by empirical formula 1 HV=10 MPa [35], and the calculation results are shown in Table 2. From the data in Table 2, it can be seen that the increase in yield strength (∆HV/3) of the Al-0.25Sc alloy after annealing at 340°C for 24 hours and 390°C for 3 hours is close to the value of the yield strength predicted by the Orowan bypass mechanism (∆σor), indicating that the mechanism of alloy precipitation phase strengthening at this time is mainly dominated by the Orowan bypass mechanism. For the Al-0.25Sc alloy, after annealing at 290°C for 3 hours and 340°C for 3 hours, the yield strength (∆HV/3) is close to the value of ∆σord and ∆σor. Because the critical radius of the Al3Sc precipitation phase strengthening mechanism transition in the literature [6] is 2.1 nm, the average radius of the precipitated particles of the Al-0.25Sc alloy after annealing at 290°C for 3 hours is about 1.5 nm, and the radius is less than the critical radius according to Table 1. It shows that the strengthening mechanism of the Al3Sc precipitation phase in the Al-0.25Sc alloy at 290°C for 3 hours is mainly dominated by the shearing mechanism. The average radius of Al3Sc particles after annealing at 340°C for 3 hours in Al-0.25Sc alloy is about 2.5 nm, which is close to the critical radius of 2.1 nm of the Al3Sc precipitation phase strengthening mechanism transition as reported by Seidman et al. [6]. Therefore, it can be inferred that the critical radius of the strengthening mechanism of the Al3Sc precipitation phase in the Al0.25Sc alloy is about 2.1 nm. The precipitation strengthening is dominated by the shearing mechanism and the Orowan bypass mechanism. The yield strength increment (∆HV/3) of Al-0.25Sc-0.12Zr alloy after annealing at 340°C for 24 hours is close to the yield strength value (∆σord) calculated by the shearing mechanism, indicating the strengthening mechanism of this samples is mainly dominated by the shearing mechanism at this time. For the Al-0.25Sc-0.12Zr alloy, the yield increase is close to that of the shearing mechanism and the Orowan bypass mechanism after annealing at 390°C for 3 hours. According to Table 1, the average radius of the precipitated phase of Al-0.25Sc-0.12Zr
alloy after annealing at 390°C for 3 hours is about 2 nm, which is close to the critical radius of 2.12 nm of the Al3(Sc, Zr) precipitation mechanism transition as reported by Seidman et al. [6]. Therefore, it can be inferred that the critical radius of the Al3(Sc, Zr) precipitation phase strengthening mechanism in the Al-0.25Sc-0.12Zr alloy is about 2 nm, and the precipitation strengthening is dominated by the Orowan mechanism and the shearing mechanism simultaneously. 4.3 Coarsening of precipitated phase Since the precipitation phase plays an important role in determining the final properties of the alloy, many researches focused on creating a precipitation phase evolution model for predicting the coarsening behavior of the precipitated phase. The theoretical models currently used to predict the coarsening kinetics of Al3Sc precipitates are mainly Lifshitz-Slyozov-Wagner theory (LSW) [7,36] and Kuehmann and Koorhees theory (KV) [37], which are applicable to binary and ternary alloys, respectively. As the Al3Sc precipitation phase grows, the strain energy caused by the coherent interface increases. When the precipitation phase grows to a certain size, dislocations are introduced at the interface between the Al3Sc precipitation phase and the α-Al matrix to reduce the mismatch. The occurrence of interface dislocation indicates the coherence of the Al3Sc precipitation phase and the destruction of the α-Al matrix. At this time, the radius of the precipitated phase is the critical radius r of the coherent/semi-coherent transition, which can be roughly expressed as (8) [38]: r=b/2δ (8) Where b is the Bernian vector of the aluminum matrix, and δ is the lattice misalignment of Al3Sc precipitation phase and α-Al matrix. In the report of the Al-Sc binary alloy, the critical radius of the precipitated phase and the matrix remains about 20 nm [8,32]. In general, the coherence of the precipitation phase of Al3Sc particles with the matrix is observed by high-resolution transmission electron microscopy. It can be seen from Fig. 8(a) that after annealing at 340°C for 3 hours, the radius of the precipitated phase particles are 2.5 nm, the transition at the interface is smooth and no misfit dislocations are observed, so it is apparently coherent with the aluminum matrix. After annealing at 390°C for 3 hours (Fig. 8(b)), the Al3Sc particles size is significantly larger and the radius of the Al3Sc particles are about 5 nm. It can be observed that Al3Sc particles still maintain a good coherent relationship with the α-Al matrix. As the HRTEM shows, the particle size increases with annealing temperature, although the particles are
significantly coarsened during annealing, the coherent relationship with the matrix remains for a long time. Fig. 9 shows that the radius of Al3(Sc, Zr) particle is small spherical and coherent with the α-Al matrix. As the annealing temperature increases and the annealing time prolongs, the particle size does not change significantly. Therefore, the resistance of Al3(Sc, Zr) particles towards coarsening is more excellent, which indicates that the thermal stability of Al3(Sc, Zr) particles are better than Al3Sc. 5. Conclusions (1) The annealing hardening phenomenon in Al-Sc and Al-Sc-Zr alloy is caused by Al3Sc and Al3(Sc, Zr) precipitates. The Al-0.25Sc alloy is markedly strengthened when annealed below 340°C, the hardness of the alloy decreases continuously with the increase of annealing time over 340°C. While in the Al-0.25Sc-0.12Zr, the decrease of the hardness value is not so obvious. The hardness decrease in Al-Sc alloy is mainly caused by the coarsening of Al3Sc particle. (2) Since the diffusion coefficient of Zr is smaller than that of Sc, the Al3(Sc, Zr) phase has better thermal stability than the Al3Sc phase. The radius of the Al3(Sc,Zr) particles can maintain its value when annealed at higher temperature for longer time. (3) The critical radius of the Al3Sc precipitation phase strengthening mechanism transformation in Al-0.25Sc alloy is about 2.1 nm. The critical radius of the Al3(Sc, Zr) precipitation phase strengthening mechanism transformation in Al-0.25Sc-0.12Zr alloy is about 2 nm. (4) Under various annealing conditions, both the Al3Sc precipitated phase and the Al3(Sc, Zr) precipitated phase are coherent with the matrix. Acknowledgements This work was supported by the Science and technology Major Project of Hunan Province [grant number 2017GK4002].
References [1] K.V. Yang, Y.J. Shi, F. Palm, X.H. Wu, P. Rometsch, Columnar to equiaxed transition in Al-Mg(-Sc)-Zr alloys produced by selective laser melting, Scripta Mater. 145 (2018) 113-117.
[2] M.J. Li, Y.J. Shi, Q.L. Pan, Y. Zhang, G.X. Lua, S.K. Guan, N. Birbilisc, Low anisotropy of fatigue crack growth in Al-5.8Mg-0.25Sc, Int. J. Fatigue.125(2019)170-178. [3] Y.J. Shi, K. Yang, S.K. Kairy, F. Palm, X.H. Wu, P.A. Rometsch, Effect of platform temperature on the porosity, microstructure and mechanical properties of an Al-Mg-Sc-Zr alloy fabricated by selective laser melting, Mater. Sci. Eng. A 732 (2018) 41-52. [4] A.K. Lohar, B. Mondal, D. Rafaja, V. Klemm, S.C. Panigrahi, Microstructural investigations on as-cast and annealed Al–Sc and Al–Sc– Zr alloys, Mater. Charact. 60 (2009) 1387-1394. [5] V.G. Davydov, T.D. Rostova, V.V. Zakharov, Y.A. Filatov, V.I. Yelagin, Scientific principles of making an alloying addition of scandium to aluminium alloys, Mater. Sci. Eng. A 280 (2000) 30-36. [6] Y. Altinsel, Y. Topkaya, Ş. Kaya, B. Şentürk, Extraction of Scandium from Lateritic Nickel-Cobalt Ore Leach Solution by Ion Exchange: A Special Study and Literature Review on Previous Works, TMS. 2018 1545-1554. [7] Weitao Zhao, Desheng Yang, Lijian Rong, TEM observation of annealing microstructure of deformed Al-Mg-Sc-Zr alloy, Acta Metall. Sin. (China). 41 (2005) 1150-1154. [8] L.M. Dougherty, I.M. Robertson, J.S. Vetrano, Direct observation of the behavior of grain boundaries during continuous dynamic recrystallization in an Al–4Mg–0.3Sc alloy, Acta Mater. 51 (2003) 43674378. [9] D.N. Seidman, E.A. Marquis, D.C. Dunand, Precipitation strengthening at ambient and elevated temperatures of heat-treatable Al (Sc) alloys, Acta Mater. 50 (2002) 4021-4035. [10] E.A. Marquis, D.N. Seidman, Nanoscale structural evolution of Al3Sc precipitates in Al (Sc) alloys, Acta Mate. 49 (2001) 1909-1919. [11] G.B. Teng, C.Y. Liu, Z.Y. Ma, W.B. Zhou, L.L. Wei, Y. Chen, J. Li, Y.F. Mo, Effects of minor Sc addition on the microstructure and mechanical properties of 7055 Al alloy during aging, Mater. Sci. Eng. A 713 (2018) 61-66. [12] M. Vlach, J. Čížek, B. Smola, O. Melikhova, M. Vlček, V. Kodetová, H. Kudrnová, P. Hruška, Heat treatment and age hardening of Al–Si– Mg–Mn commercial alloy with addition of Sc and Zr, Mater. Charact. 129 (2017) 1-8.
[13] Z.H. Jia, J. Røyset, J.K. Solberg, Q. Liu, Formation of precipitates and recrystallization resistance in Al–Sc–Zr alloys, Trans. Nonferrous Metals Soc. China 22 (2012) 1866-1871. [14] S. Costa, H. Puga, J. Barbosa, A.M.P. Pinto, The effect of Sc additions on the microstructure and age hardening behaviour of as cast Al–Sc alloys, Mater. Des. 42 (2012) 347-352. [15] Y. Miyake, Y. Sato, R. Teranishi, K. Kaneko, Effect of heat treatments on the microstructure and formability of Al–Mg–Mn–Sc–Zr alloy, Micron. 101 (2017) 151-155 [16] C. Suwanpreecha, J. P. Toinin, R.A. Michi, P. Pandee, D.C. Dunand, C. Limmaneevichittr, Strengthening mechanisms in Al-Ni-Sc alloys containing Al3Ni microfibers and Al3Sc nanoprecipitates, Acta Mater. 164 (2019) 334-346. [17] J. Taendl, A. Orthacker, H. Amenitsch, G. Kothleitner, C. Poletti, Influence of the degree of scandium supersaturation on the precipitation kinetics of rapidly solidified Al-Mg-Sc-Zr alloys, Acta Mater. 117 (2016) 43-50. [18] P. Xu, F. Jiang, Z. Tang, N. Yan, J. Jiang, X. Xu, Y. Peng, Coarsening of Al3Sc precipitates in Al-Mg-Sc alloys, J Alloy Compd. 781 (2019) 209-215. [19] P. Xu, F. Jiang, M. Tong, Z. Tang, J. Jiang, N. Yan, Y. Peng, Precipitation characteristics and morphological transitions of Al3Sc precipitates, J Alloy Compd. 790(2019) 509-516. [20] W.S. Lee, T.H. Chen, C.F. Lin, M.S. Chen, Impact deformation behavior and dislocation substructure of Al–Sc alloy, J Alloy Compd. 493(2010) 580-589. [21] J. Liu, P. Yao, N. Q. Zhao, C. S. Shi, H. J. Li, X. Li, D. S. Xi, S. Yang, Effect of minor Sc and Zr on recrystallization behavior and mechanical properties of novel Al-Zn-Mg-Cu alloys, J Alloy Compd. 657(2016) 717-725. [22] L. Fu, Y. Li, F.Q. Jiang, J.W. Huang, G.F. Xu, Z.M. Yin, On the role of Sc or Er micro-alloying in the microstructure evolution of Al-Mg alloy sheets during annealing, Mater. Charact. 157 (2019) 109918. [23] S. Iwamura, Y. Miura, Loss in coherency and coarsening behavior of Al3Sc precipitates, Acta Mater. 52 (2004) 591-600.
[24] B. Forbord, W. Lefebvre, F. Danoix, H. Hallem, K. Marthinsen, Three dimensional atom probe investigation on the formation of Al3(Sc,Zr)-dispersoids in aluminium alloys, Scripta Mater. 51 (2004) 333-337. [25] G.B. Teng, C.Y. Liu, Z.Y. Ma, W.B. Zhou, L.L. Wei, Y. Chen, J. Li, Y.E. Mo, Effects of minor Sc addition on the microstructure and mechanical properties of 7055 Al alloy during aging, Mater. Sci. Eng. A 713(2008) 61-66. [26] C. Watanabe, T. Kondo, R. Monzen, Coarsening of Al3Sc precipitates in an Al-0.28 wt pct Sc alloy, Metall. Mater. Trans. A 35 (2004) 3003-3008. [27] C.B. Fuller, J.L. Murray, D.N. Seidman, Temporal evolution of the nanostructure of Al (Sc,Zr) alloys: Part I-Chemical compositions of Al3(Sc1－x Zrx) precipitates, Acta Mater. 53(2005) 5401-5413. [28] C.B. Fuller, D.N. Seidman, Temporal evolution of the nanostructure of Al (Sc,Zr) alloys: Part II-coarsening of Al3(Sc1 － x Zrx) precipitates, Acta Mater. 53(2005) 5415-5428. [29] W. Lefebvrea, F. Danoixa, H. Hallem, Precipitation kinetic of Al3(Sc,Zr) dispersoids in aluminium, J Alloy Compd. 470(2009) 107-110. [30] E. Clouet, L. Laé, Ludovic, T. Épicier, W. Lefebvre, M. Nastar, A. Deschamps, Complex precipitation pathways in multicomponent alloys, Nature Mater. 5(2006) 482-488. [31] A. Tolley, V. Radmilovic, U. Dahmen, Segregation in Al3(Sc,Zr) precipitates in Al–Sc–Zr alloys, Scripta Mater. 52 (2005) 621-625. [32] K.E. Knipling, R.A. Karnesky, C.P. Lee, D.C. Dunand, D.N. Seidman, Precipitation evolution in Al–0.1Sc, Al–0.1Zr and Al–0.1Sc– 0.1Zr (at. %) alloys during isochronal aging, Acta Mater. 58 (2010) 51845195. [33] Booth-Morrison Christopher, D.C. Dunand, D.N. Seidman, Coarsening resistance at 400°C of precipitation-strengthened Al–Zr–Sc– Er alloys, Acta Mater. 59 (2011) 7029-7042. [34] D. Tabor, Br J Appl Phys 7(1956) 159-166. [35] C.L. TANG, D.J. ZHOU, Precipitation hardening behavior of dilute binary Al-Yb alloy, Transactions of Nonferrous Metals Society of China, 24 (2014) 2326-2330.
[36] G.M. Novotny, A.J. Ardell, Precipitation of Al3Sc in binary Al–Sc alloys, Mater. Sci. Eng. A, 318 (2001) 144-154. [37] R.A. Karnesky, D.C. Dunand, D.N. Seidman, Evolution of nanoscale precipitates in Al microalloyed with Sc and Er, Acta Mater. 57 (2009) 4022-4031. [38] Y. Zhang, K. Gao, S. Wen, H. Huang, Z. Nie, D. Zhou, The study on the coarsening process and precipitation strengthening of Al3Er precipitate in Al–Er binary alloy, J Alloy Compd. 610 (2014)7-34.
Table 1. Measured average radii and volume fraction of precipitates Sc-containing aluminum alloy(wt.%) Al-0.25Sc
Al-0.25Sc-0.12Zr
Φx10-3
Annealing condition
r (nm)
290°C /3h
1.5±0.28 0.40±0.01
340°C /3h
2.5±0.50 0.35±0.03
340°C /24h
4±0.67
0.20±0.04
390°C /3h
5±0.92
0.17±0.07
340°C /24h
1.5±0.28 0.42±0.05
390°C /3h
2±0.53
0.47±0.03
Table 2. Experimental and calculated strength increments Sc-containing aluminum alloy(wt.%) Al-0.25Sc
Annealing condition
∆HV/3 (MPa)
290°C /3h
149.9±2.8 148.7 126.4±5.1 139.1
340°C /3h
340°C /24h 73.9±3.4 390°C /3h 71.5±6.3 Al-0.25Sc0.12Zr
∆σord( MPa)
∆σcoh+ ∆σor( ∆σmod MPa) (MPa) 297.6 149.3 324.7
112.7
105.2
283.2
64.9
97.0
279.6
52
304.9
153.2
351.7
146.2
340°C /24h 139.8±2.4 143.0 390°C /3h 157.3±2.4 161.2
Fig.1. Effects of annealing temperature and time on the microhardness of the Al-0.25Sc alloy.
Fig.2. Effects of annealing temperature and time on the microhardness of the Al-0.25Sc-0.12Zr alloy.
Fig.3. Low magnification optical micrographs:(a) Al-0.25Sc annealed at 340°C for 3h; (b) Al-0.25Sc annealed at 340°C for 24h; (c) Al-0.25Sc0.12Zr annealed at 340°C for 3h; (d) Al-0.25Sc-0.12Zr annealed at 340°C for 24h.
Fig.4. Dark field TEM micrographs of Al3Sc precipitates (utilizing a 110 superlattice reflection): (a) 290°C for 3h; (b) SADP of (a); (c) 340°C for 3h; (d) 340°C for 24h; (e) 390°C for 3h.
Fig.5. Dark field TEM micrographs of Al3(Sc,Zr) precipitates (utilizing a 001 superlattice reflection): (a) 340°C for 3h; (b) SADP of (a); (c) 390°C for 3h.
Fig.6. Bright field TEM micrographs of Al3Sc precipitates in Al-0.25Sc alloy annealed at : (a) 290°C for 3h; (b) 340°C for 3h; (c) 340°C for 24h; (d) 390°C for 3h.
Fig.7. Bright field TEM micrographs of Al3(Sc,Zr) precipitates in Al0.25Sc-0.12Zr alloy annealed at : (a) 290°C for 3h; (b) 340°C for 3h; (c) 340°C for 24h; (d) 390°C for 3h.
Fig.8. HRTEM micrographs of Al3Sc precipitates in Al-0.25Sc alloy annealed at (a)340°C for 3h; (b) 390°C for 3h.
Fig.9. HRTEM micrographs of Al3(Sc, Zr) precipitates in Al-0.25Sc0.12Zr alloy annealed at (a) 340°C for 24h and (b) 390°C for 3h.
Fig.10. the distribution range of the precipitates radius: Al3Sc precipitates in Al-0.25Sc alloy annealed at (a)290°C for 3h; (b) 340°C for 3h; (c) 340°C for 24h; (d) 390°C for 3h; Al3(Sc, Zr) precipitates in Al-0.25Sc0.12Zr alloy annealed at (e) 340°C for 24h and (f) 390°C for 3h.
Highlights: 1 The diffusion rate of Sc atoms is inhibited by Zr atoms, which can prevent the coarsening of Al3(Sc, Zr). 2 The hardness decrease in Al-Sc alloy is mainly caused by the coarsening of Al3Sc particle. 3 The radius of the Al3(Sc,Zr) particles can maintain thier value when annealed at higher temperature(>340 °C) for longer time(~24 h).
Declaration of interests The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
S0925-8388(19)34168-4
DOI:
https://doi.org/10.1016/j.jallcom.2019.152922
Reference:
JALCOM 152922
To appear in:
Journal of Alloys and Compounds
Received Date: 23 July 2019 Revised Date:
14 October 2019
Accepted Date: 4 November 2019
Please cite this article as: Y. Luo, Q. Pan, Y. Sun, S. Liu, Y. Sun, L. Long, X. Li, X. Wang, M. Li, Hardening behavior of Al-0.25Sc and Al-0.25Sc-0.12Zr alloys during isothermal annealing, Journal of Alloys and Compounds (2019), doi: https://doi.org/10.1016/j.jallcom.2019.152922. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.
Hardening behavior of Al-0.25Sc and Al-0.25Sc-0.12Zr Alloys during isothermal annealing Yuhong Luoa, Qinglin Pana,b, Yuqiao Suna, Shuhui Liub, Yuanwei Suna, Liang Longb, Xinyu Lib, Xiaoping Wangc, Mengjia Lia* a
School of Materials Science and Engineering, Central South University, Changsha 410083, china
b
Light Alloy Research Institute, Central South University, Changsha 410083, china c
Hunan Dongfang scandium industry Co., Ltd. Changsha 410083, china
Abstract The hardening behavior of Al-0.25Sc and Al-0.25-0.12(wt.%) alloys during isothermal annealing were studied by transmission electron microscopy (TEM), scanning electron microscopy (SEM), optical microscopy (OM), and hardness measurements. The results show that Sccontaining aluminum alloy contained fine dispersion of spherical secondary Al3Sc particles, which developed during the annealing process. Al3(Sc, Zr) precipitates in the Al-0.25Sc-0.12Zr alloys are core-shell structure. Different isothermal annealing temperatures can lead to different strengthening effects on the alloys. Al-0.25Sc and Al-0.25Sc0.12Zr alloys both exhibit obvious annealing hardening behavior, and the ternary Al-0.25Sc-0.12Zr alloy shows better thermal stability than Al0.25Sc alloy. Key words: hardening behavior, precipitation strengthening, thermal stability, secondary Al3Sc particles, Sc-containing aluminum alloy 1. Introduction The addition of Sc can greatly improve the mechanical properties, tensile strength, plasticity, weld ability, corrosion resistance of aluminum alloy [1-3]. It has been found that Sc-containing industrial aluminum alloys can show annealing hardening behavior. The addition of Sc and Zr can greatly improves the thermal stability of the alloys during annealing. Supersaturated solid solution containing Al and Sc in the alloy during *
Corresponding author
E-mail address: [email protected]
solidification. The crystal structure of primary Al3Sc phase particles is same to the α-Al matrix, the lattice mismatch between the primary phase particles and the aluminum matrix is 0.27%, and these precipitates act as nucleation sites for Al matrix, which can lead to the grain refining effect in Al alloys [4]. In the subsequent annealing process, secondary Al3Sc particles (L12 type) are precipitated, which can produce precipitation strengthening effect and pin the grain boundaries [5,6]. Al3(Sc, Zr) particles with L12 structure can precipitate in the Al alloys after adding trace amounts of Sc and Zr [7-9]. The lattice parameters difference between Al3(Sc, Zr) and Al3Sc particles are negligible. The precipitation phase of Al3(Sc, Zr) can maintain the coherent relationship with the matrix of α-Al during annealing, which can strongly hinder the movements of dislocation and grain boundaries, leading to higher strength and recrystallization temperature of the alloy. Lohar et.al [4] reported that the combined addition of Sc and Zr can maintain the stability of the particles at a high temperature (reducing the coarsening rate of the particles) compared to adding Sc alone, which can improve the grain refinement effect and thermal stability of the Al alloys. The microstructure evolution of the Sc-containing aluminum alloy during annealing can lead to the variation mechanical properties in the alloy [1013]. The annealing hardening behavior has been widely reported in Sccontaining aluminum alloys [14-19]. However, there is no systematic discussion on the mechanism of microscopic organization on the precipitation strengthening. In order to reduce the influence of other factors on the annealing hardening behavior of Sc-containing aluminum alloys, Al-Sc binary alloy and Al-Sc-Zr ternary alloy are used for detailed experiment. This paper aims to research the effect of strengthening mechanism of precipitation phase and coarsening of Al3Sc, Al3(Sc, Zr) precipitated particles on the annealing hardening behavior of Al-Sc and Al-Sc-Zr alloys. 2. Experimental 2.1 Material preparation Aluminum alloys (Al-0.25Sc, Al-0.25Sc-0.12 wt. %Zr) were prepared by melting Al with commercial purity and Al-2.2Sc and Al4.0Zr master alloy in a muffle furnace with graphite crucible. Firstly, the ingots were heated to 780°C, then Al-4.0Zr master alloy and Al-2.2Sc master alloy were added. After the molten alloys were thorough stirring, the molten alloys were allowed to stand for 5 min. Then the melt is
poured into the water-cooling copper mold. The samples are annealed at 260°C, 290°C, 340°C, and 390°C for 0.1 h, 0.5 h, 1 h, 3 h, 5 h, 10 h, 16 h, and 24 h, respectively. 2.2 Hardness measurements The hardness test used a 401 MVD TM digital micro Vickers hardness tester with a load of 0.3 KN and time of indentation of 15 s. The hardness values reported are taken from the average of the five measurements values. Hardness measurements were prepared by cutting sample pieces of 12 mm*12 mm*4 mm (the test surface size was 12mm*12mm) from the annealing treated materials and by mechanically grinding them. 2.4 Microstructure characterization The microstructure of the as-cast Al-0.25Sc and Al-0.25Sc-0.12Zr alloys were observed at 340°C for 3 hours and 24 hours, respectively, under a metallographic microscope with the instrument model Leica DFC295. The samples were cut by electrical discharge machining, smoothed by water sanding paper and metallographic sandpaper, and mechanically polished. Then the polished samples were treated in a coating liquid (30ml of fluoroboric acid + 200 ml of distilled water), the voltage is 15~20V, the current is 0.1~0.5mA, and the operation time is 1~2 min at room temperature. The grain shape and size of alloy samples were observed by metallographic microscope. A Tecnai G2 F20 transmission electron microscope (TEM) with an accelerating voltage of 200 KV was used for investigating the general microstructure and the distribution of Al3Sc or Al3(Sc, Zr) precipitates. Thin foils for TEM were prepared by cutting sample pieces from the annealing treated materials and by mechanically grinding them down to ~100um. Discs with 3mm in diameter were punched and dimpled down to 60um, followed by twin-jet electro polishing at 100mA. A solution of 1/3(volume fraction) nitric acid and 2/3 methanol at -30°C was applied during the electro polishing [20,21]. The average radius size and volume fraction of the precipitated phases in the alloys were statistically analyzed by Image-J software [22]. In a large number of TEM dark field phase diagrams, three TEM dark phase images which contain over 100 precipitates were chosen for obtain the precipitate average radius. The distribution range of the precipitates can be obtained. The detailed method to calculate the size of precipitates is shown as below. Import the image into the software first, set the ruler, and then adjust the threshold
color of the image then use the analysis precipitate function in the software to analyze and finally get the volume fraction. 3. Results 3.1 Microhardness of Al-0.25Sc and Al-0.25Sc-0.12Zr alloys during annealing The hardness change curves of as-cast Al-0.25Sc alloys during annealing at different temperatures is shown in Fig.1. It can be found that the alloy’s hardness is significantly increased after annealing. Under the annealing condition of 260°C, the hardness value of the alloy increases with the annealing time extended from 0 to 24 hours. The hardness value of the alloy reach 68.1 HV after annealing for 24 hours. When the annealing temperature is 290°C, the hardness of the alloy reaches to the highest value after 0.5 hour, and the peak hardness value is 68.3 HV. The hardness of the alloy does not change much when further increase the annealing time. During annealing at 340°C, the hardness of the alloy reaches a peak value of 63 HV at 0.5 hour. When increase the annealing time to 10 hours, the hardness of the alloy nearly does not change, while the hardness of the alloy begins to decrease after annealing for 16 hours. When annealing at 390°C, the alloy peak at 0.5 hours and the hardness begins to decrease after that. It can be found that the hardening effect of the alloy is not so obviously after annealing at higher temperatures. The peak hardness of the as-cast alloy after 390°C treatment is much lower than other samples. The hardness curves of the as-cast Al-0.25Sc-0.12Zr alloys during annealing at different temperatures is shown in Fig.2. The hardness of alloys is significantly improved after annealing. When the annealing is 260°C, the variation of hardness of the alloy is similar to that of the Al0.25Sc alloy, the hardness increases with the annealing time. Other curves exhibit roughly the same trend of hardness change. Initially, the hardness of the alloys increases with increasing annealing time sharply from the starting point until reaching the highest value. With the prolongation of annealing time, the hardness values of the three alloys do not decline obviously. When the alloy is annealed at 290°C, the hardness of the alloy reaches its peak at 3 h, and the peak value is 70.7 HV. When the annealing treatment temperature is 340°C and 390°C, the alloys reach a peak at 0.5 h, the peak values of the alloys are 64.9 HV and 65.4 HV, respectively.
3.2 Microstructure of Al-0.25Sc and Al-0.25Sc-0.12Zr alloys Image J software is used to measure the average size of crystal grains by the line cutting method. By cutting the grain with a straight line of a certain length, it can be obtained that the average length of the straight line can intercept one grain, that is, the average grain size. If the length of the straight line is L and the total number of cut grains is N, the average grain size can be expressed as (1): Average Grain Size =
N
(1)
The optical microscopy investigations from the anodized specimen of the Al-0.25Sc alloys is shown in Fig. 3. It reveals that the average grain size is about 88 µm, the shape of the equiaxed grains is not uniform and the grain size is inconsistent in the structure after annealing at 340°C for 3h (Fig.3(a)). When alloy is annealed for 24 hours, the shape and size of the equiaxed crystals are slightly more uniform in the structure, and the grain size is 92 µm (Fig.3 (b)). The grain size of the alloy annealing at 340°C for 3 hours decreases to 58 µm after the addition of Zr, and the dendrites disappears, completely replaced by equiaxed grains (Fig.3(c)). The average grain size does not change much with annealing time increasing, the average grain size is about 55 µm (Fig.3 (d)). It can be found that the microstructure of the alloys is significantly refine by the addition of Zr. It is indicating that the combined addition of Sc and Zr can lead to better refinement for microstructure of the aluminum alloy. Dark-field TEM micrographs of precipitates in Al-0.25Sc alloys annealing at 290~390°C are shown in Fig. 4. It can be seen from the Fig.4 that many white approximately spherical precipitates are dispersed in the α-Al matrix. The selected area diffraction (SAD) pattern is inserted in Fig. 4(b). As the annealing temperature increases, the radius of Al3Sc particles increases with the annealing time, and its number density of precipitates decreases significantly. After annealing at 290°C for 3 h (Fig. 4 (a)), Al3Sc particles are uniformly dispersed in a large amount, and the average radius of the precipitates is about 1.5 nm. After annealing at 340°C for 3 h (Fig. 4 (c)), the number density of Al3Sc precipitates decrease, the average radius of the precipitates increase to about 2.5 nm, when the annealing time is extended to 24 h (Fig. 4 (d)) the Al3Sc precipitates are obviously coarsening, the average radius of the precipitates is increase to about 4 nm and the number density of Al3Sc precipitates is reduces. After annealing at 390°C for 3 h (Fig. 4 (e)), the
number of precipitated phases of Al3Sc become very small and the average radius is increase. The dark-field TEM micrographs of Al3(Sc, Zr) precipitates in Al0.25Sc-0.12Zr alloys are shown in Fig.5. After annealing at 340°C for 24 h (Fig. 5 (a)), the average radius of the precipitated phase is about 1.5 nm. After annealing at 390°C for 3 h (Fig. 5 (c)), the number density of precipitates is higher, and the average radius of the precipitated phase is about 2 nm. The particle radius does not increase much when the annealing temperature increases. The particles of Al3(Sc, Zr) shows very excellent resistance towards coarsening. Comparing the TEM dark field phase of the two alloys, it can be seen that the Al3(Sc, Zr) particles in the Al-0.25Sc-0.12Zr alloy have higher number density and smaller size than Al3Sc in the Al-0.25Sc alloy. The TEM bright field images of Al-0.25Sc alloy and Al-0.25Sc0.12Zr alloy after annealing at 290~390°C are shown in Fig.6 and Fig.7, respectively. In general, the coherency between Al3Sc precipitate phase and the matrix is observed by HREM analysis, but this method is too inefficient in analyzing many particle coherences. Therefore, Iwanura and Miura et al. [23] proposed the use of Ashby-Brown contrast to determine the coherent relationship of the precipitated phase to the matrix. It is observed by TEM that fine nanoscale particles with bean-shaped disperse in annealing samples. The direction of the no-contrast line of the beanlike center is perpendicular to the operation vector g. This bean-like contrast is generally observed in alloys with Ashby-Brown characteristics, indicating that the precipitated phase is coherent with the matrix. The HRTEM micrograph of Al3Sc precipitates are shown in Fig.8. In the HRTEM micrograph, no dislocations can be observed at the interface, and it can be clearly seen that the particle maintains a coherent relationship with the matrix. The radius of Al3Sc precipitates after annealing at 290°C for 3 hours are about 2.5 nm (Fig. 8(a)), and the radius increase to about 5 nm after the alloy annealing at 390°C for 3 hours (Fig. 8 (b)). Figure 9 shows the HRTEM micrograph of Al3(Sc, Zr) precipitates in the Al-0.25Sc-0.12Zr alloy. The radius of Al3(Sc, Zr) precipitates are about 1.5 nm of the samples annealing at 340°C for 24 hours (Fig. 9 (a)), and the radius of the Al3(Sc, Zr)precipitates increase to about 2 nm after the alloy is annealed for 3 hours at 390°C (Fig. 9 (b)).
4. Discussions 4.1 Microhardness evolution It can be seen from Fig. 1 and Fig. 2 that both alloys have annealing hardening behavior during the annealing process, which due to the addition of Sc. When annealing at low temperatures, the hardness of the alloy increased slower than at higher annealing temperature, which is due to that the precipitation of Al3Sc and Al3(Sc, Zr) precipitates is controlled by the diffusion of Sc and Zr atoms [24,25]. Studies on the annealing of aluminum alloys containing Sc at a content of 0.2~0.5% at temperatures below 340°C indicate that the increase in hardness of the alloy is due to the formation of nano sized Al3Sc particles during annealing [7]. After reaching the hardness peak during annealing, the hardness begins to decrease with time, as the Al3Sc particles begin to coarsen (~40 nm) [26]. According to the related report [7], it can be known that the heterogeneous nucleation of Al3Sc particles is promoted under high temperature annealing, thus resulting in a lower volume fraction and a larger radius size of the particles, so that the hardness value of the alloy is lowered after annealing at high temperature. In the solidification process of Sc-containing aluminum alloy, primary Al3Sc particles are precipitated. Since the lattice constant of the precipitated Al3Sc particles does not differ much from the lattice constant of the Al matrix, the phase of the primary Al3Sc can provide more nucleation point during solidification of the alloy, which can significantly refine the grains of the alloy. In the subsequent annealing process, the alloy will precipitate a large amount of dispersed, fine secondary Al3Sc particles, which can lead to precipitation strengthening. However, as the annealing temperature increases, the secondary Al3Sc particles begin to grow and coarsen, so the hardness value of the alloys begins to decrease. Because the Al3Sc particles have a very high coherency mismatch, it is clear that lattice strain can prevent the movement of dislocation and prevent grain growth. However, under higher temperature annealing conditions, the coarsening of Al3Sc particles is unavoidable, and the weakening of Al-0.25Sc alloy becomes more and more obvious. When Sc and Zr elements are added into the aluminum alloy, the Al3Sc particles will precipitate first. As the annealing time increase, Zr will be enriched in the outer layer of Al3Sc to form the core-shell structure of Al3(Sc, Zr) [27-30]. Tolley et al. [4] used the EDS in transmission electron microscopy to analyze the precipitation phase of the Al-0.61Sc-0.40Zr(wt.%) alloy during annealing. The results show that the precipitated phase is Al3(Sc, Zr) phase, the structure of the precipitated phase particles is core rich in Sc, and the shell is rich in Zr. Forbord et al. [24] used a three-dimensional atom probe to study the
formation process of Al3(Sc, Zr) precipitated phase in Al-Sc-Zr alloy. The results show that the Al3(Sc, Zr) precipitated phase mainly includes Al and Sc atoms in the early nucleation, while the Zr atom is concentrated on the precipitated phase at a later stage. Since the diffusion rate of Zr in Al is lower than Sc and mainly distributed in the outer layer of Al3(Sc, Zr), the growth rate of the precipitated phase is slowed down, and the thermal stability of the Al alloy is effectively improved [13]. Therefore, the hardness of the Zr-containing ternary alloy decrease slowly with time at higher annealing temperature. It can be seen from the dark field phase diagram of the alloys (Fig. 4) that the density and radius of the Al3Sc particles are different. The differences are caused by different annealing conditions. Strengthening effect is influenced by the radius and density of precipitated phase. As a result, the hardness of the alloys changed. Sc can be melted into Al to form a supersaturated solid solution. In the subsequent thermal processing and heat treatment, the supersaturated solid solution will overcome the energy barrier and dissolvent decomposition forming an Al3Sc precipitation phase. The Al3Sc precipitation phase is closely related to the heat treatment temperature. The higher the heat treatment temperature, the larger the driving force of the second phase precipitation, and the faster the precipitation rate of the Al3Sc precipitation phase, and the number density will decrease. When the holding time is extended or the annealing temperature is increased, the precipitation phase of Al3Sc will become larger. The nucleation and coarsening of the precipitated phase are greatly depended on the temperature. This process can be explained by the phase change driving force (2): Fαβ = −
1 Cβ − Cα C ∙ ∙ RT ln Vβ 1 − Cα Cα
(2)
Where Vβ is the volume of the β phase （ m3/g ） , C0 is the concentration of the B component in the supersaturated solid solution, Cα is the concentration of the B component in the equilibrium α phase, and Cβ is the equilibrium of the B component in the β phase. R is the ideal gas constant and T is the absolute temperature at which the phase change occurs. It can be seen that the atomic diffusion rate increases with temperature, the phase change driving force also increases, so the precipitated phase is more likely to precipitate out from the matrix, as a result the higher the annealing temperature of the alloy, the shorter the peak time. However, the number density of precipitation reduced when increase the annealing temperature, which leads to the decrease of the peak value of the alloys when increase the annealing temperature. The
higher the temperature, the faster the precipitation phase grows, moreover, the strengthening of the alloy also decay faster. Combined with the particle size distribution diagram in Fig. 10, it can be seen that the size of the Al3Sc particles is larger than Al3(Sc, Zr) particles in general. Therefore, the precipitation strengthening effect of Al3(Sc, Zr) particles is larger than Al3Sc. The lattice structure of Al3(Sc,Zr) belongs to L12 structure, which lattice parameters is slightly different from Al3Sc particles, resulting in slower particle coarsening rate. Forbord et al. [24] and Tolley et al. [31] suggested that the presence of Zr weakens the phenomenon of particle coarsening, and the size of the precipitates are smaller, with larger quantity. Forbord et al. [24] used atom probes to study the formation of Al3(Sc, Zr) precipitates in Al-Sc-Zr alloys. It shows that Al3(Sc, Zr) precipitates mainly include Al and Sc in the early nucleation. The Zr atom is concentrated on the precipitation phase after that. The addition of Zr also increases the supersaturating of the aluminum-based solid solution, which in turn increases the dynamic and quantitative density of the precipitated phase of Al3(Sc, Zr). Since the resistance towards coarsening performance of the Al3(Sc, Zr) particles is excellent, the precipitates after annealing remain fine and dispersed, so the hardness of alloys can maintain the peak value when further increase annealing time. It can be concluded that the addition of Zr element to the alloy can lead to better strengthening properties than the alloy with addition of Sc alone. 4.2 Strengthening mechanism of precipitation phase The evolution of microhardness is related to the growth of precipitated phase and strengthening mechanism. There are two main types of strengthening mechanisms for the precipitation phase: 1) shearing mechanism; 2) the strengthening of the Orowan bypass mechanism. When the shearing mechanism dominates, the yield strength increment is proportional to the square of the precipitation phase radius. The larger the precipitation phase size is, the more obvious the strengthening effect will be. When the Orowan bypass mechanism dominates, the increase in yield strength is inversely proportional to the radius of the precipitated phase. The domination of mechanism depends on the radius of the precipitation phase. When the radius of the precipitation phase is small, the shearing mechanism is the main strengthening mechanism. The yield strength increment can be estimated based on order strengthening ∆σord, modulus mismatch strengthening ∆σmod, and coherency strengthening ∆σcoh.
The order strengthening t ∆σord can be calculated according to the equation (3) [29,33]:
γ$%& 3π∅ +⁄, (3) ∆σ = 0.81M 2b 8 Where M=3.06 is the Taylor mean matrix orientation factor for Al, γ$%& =0.5 J/m2 is an average value of the Al3Sc antiphase boundary (APB) energy for the (111) plane, b=0.286 nm is the magnitude of the Al Burgers vector, and ∅ is the volume fraction of the precipitated phase. The modulus mismatch strengthening ∆σmod can be calculated according to the equation (4) [32,33]:
2∅ + , r 5.⁄,6+ (4) ∆σ. = 0.0055M0∆G2 b3 4 Gb , b Where G=25.4 GPa is the shear modulus of Al matrix, ∆G is the difference between the shear modulus of the precipitated phase and that of Al matrix, and the shear modulus of the precipitation phase of Al3Sc and Al3(Sc,Zr) is 68 GPa, r is the precipitation phase average radius, m = 0.85 is a constant. +⁄,
⁄
The coherency strengthening ∆σcoh can be calculated according to the equation (5) [32,33]:
r∅ + , (5) ∆σ7 8 = Mαε 0.5Gb Where αε = 2.6 is a constant and ε=δ 2⁄3 is the mismatch parameter, δ=1.3% is the lattice mismatch for the Al3(Sc1-xZrx) precipitated, which is estimated from the composition-dependent lattice parameters. Coherency strengthening occurs simultaneously with modulus mismatch strengthening. 0Gε25⁄,
⁄
The yield strength increment caused by the shearing mechanism increases with the average radius r until the strengthening mechanism is switched to the Orowan bypass mechanism. The Orowan bypass mechanism strengthening ∆σor can be calculated according to the equation below (6) [32,33]: =
r<2⁄3 ln ; = b 0.4MGb
(6)
λ π√1 − ν Where ν = 0.34 is the Poisson's ratio of Al and λ is the spacing of the edge-to-edge interprecipitate spacing (7) [32,33]: ∆σ
3π λ = >? − 1.64A r 4ϕ
(7)
The above model of yield strength increment is used to calculate the theoretical yield strength increment and compare with the experimental values. The experimental yield strength increment is estimated by using ∆HV/3 [6,32,34], where ∆HV is the microhardness increment of the ascast alloy after annealing. Which formula is valid for aluminum alloys, except pure Al [6]. Generally, the unit of microhardness increment is HV, and the unit of yield strength is MPa. For the sake of comparison, the unit of microhardness is converted by empirical formula 1 HV=10 MPa [35], and the calculation results are shown in Table 2. From the data in Table 2, it can be seen that the increase in yield strength (∆HV/3) of the Al-0.25Sc alloy after annealing at 340°C for 24 hours and 390°C for 3 hours is close to the value of the yield strength predicted by the Orowan bypass mechanism (∆σor), indicating that the mechanism of alloy precipitation phase strengthening at this time is mainly dominated by the Orowan bypass mechanism. For the Al-0.25Sc alloy, after annealing at 290°C for 3 hours and 340°C for 3 hours, the yield strength (∆HV/3) is close to the value of ∆σord and ∆σor. Because the critical radius of the Al3Sc precipitation phase strengthening mechanism transition in the literature [6] is 2.1 nm, the average radius of the precipitated particles of the Al-0.25Sc alloy after annealing at 290°C for 3 hours is about 1.5 nm, and the radius is less than the critical radius according to Table 1. It shows that the strengthening mechanism of the Al3Sc precipitation phase in the Al-0.25Sc alloy at 290°C for 3 hours is mainly dominated by the shearing mechanism. The average radius of Al3Sc particles after annealing at 340°C for 3 hours in Al-0.25Sc alloy is about 2.5 nm, which is close to the critical radius of 2.1 nm of the Al3Sc precipitation phase strengthening mechanism transition as reported by Seidman et al. [6]. Therefore, it can be inferred that the critical radius of the strengthening mechanism of the Al3Sc precipitation phase in the Al0.25Sc alloy is about 2.1 nm. The precipitation strengthening is dominated by the shearing mechanism and the Orowan bypass mechanism. The yield strength increment (∆HV/3) of Al-0.25Sc-0.12Zr alloy after annealing at 340°C for 24 hours is close to the yield strength value (∆σord) calculated by the shearing mechanism, indicating the strengthening mechanism of this samples is mainly dominated by the shearing mechanism at this time. For the Al-0.25Sc-0.12Zr alloy, the yield increase is close to that of the shearing mechanism and the Orowan bypass mechanism after annealing at 390°C for 3 hours. According to Table 1, the average radius of the precipitated phase of Al-0.25Sc-0.12Zr
alloy after annealing at 390°C for 3 hours is about 2 nm, which is close to the critical radius of 2.12 nm of the Al3(Sc, Zr) precipitation mechanism transition as reported by Seidman et al. [6]. Therefore, it can be inferred that the critical radius of the Al3(Sc, Zr) precipitation phase strengthening mechanism in the Al-0.25Sc-0.12Zr alloy is about 2 nm, and the precipitation strengthening is dominated by the Orowan mechanism and the shearing mechanism simultaneously. 4.3 Coarsening of precipitated phase Since the precipitation phase plays an important role in determining the final properties of the alloy, many researches focused on creating a precipitation phase evolution model for predicting the coarsening behavior of the precipitated phase. The theoretical models currently used to predict the coarsening kinetics of Al3Sc precipitates are mainly Lifshitz-Slyozov-Wagner theory (LSW) [7,36] and Kuehmann and Koorhees theory (KV) [37], which are applicable to binary and ternary alloys, respectively. As the Al3Sc precipitation phase grows, the strain energy caused by the coherent interface increases. When the precipitation phase grows to a certain size, dislocations are introduced at the interface between the Al3Sc precipitation phase and the α-Al matrix to reduce the mismatch. The occurrence of interface dislocation indicates the coherence of the Al3Sc precipitation phase and the destruction of the α-Al matrix. At this time, the radius of the precipitated phase is the critical radius r of the coherent/semi-coherent transition, which can be roughly expressed as (8) [38]: r=b/2δ (8) Where b is the Bernian vector of the aluminum matrix, and δ is the lattice misalignment of Al3Sc precipitation phase and α-Al matrix. In the report of the Al-Sc binary alloy, the critical radius of the precipitated phase and the matrix remains about 20 nm [8,32]. In general, the coherence of the precipitation phase of Al3Sc particles with the matrix is observed by high-resolution transmission electron microscopy. It can be seen from Fig. 8(a) that after annealing at 340°C for 3 hours, the radius of the precipitated phase particles are 2.5 nm, the transition at the interface is smooth and no misfit dislocations are observed, so it is apparently coherent with the aluminum matrix. After annealing at 390°C for 3 hours (Fig. 8(b)), the Al3Sc particles size is significantly larger and the radius of the Al3Sc particles are about 5 nm. It can be observed that Al3Sc particles still maintain a good coherent relationship with the α-Al matrix. As the HRTEM shows, the particle size increases with annealing temperature, although the particles are
significantly coarsened during annealing, the coherent relationship with the matrix remains for a long time. Fig. 9 shows that the radius of Al3(Sc, Zr) particle is small spherical and coherent with the α-Al matrix. As the annealing temperature increases and the annealing time prolongs, the particle size does not change significantly. Therefore, the resistance of Al3(Sc, Zr) particles towards coarsening is more excellent, which indicates that the thermal stability of Al3(Sc, Zr) particles are better than Al3Sc. 5. Conclusions (1) The annealing hardening phenomenon in Al-Sc and Al-Sc-Zr alloy is caused by Al3Sc and Al3(Sc, Zr) precipitates. The Al-0.25Sc alloy is markedly strengthened when annealed below 340°C, the hardness of the alloy decreases continuously with the increase of annealing time over 340°C. While in the Al-0.25Sc-0.12Zr, the decrease of the hardness value is not so obvious. The hardness decrease in Al-Sc alloy is mainly caused by the coarsening of Al3Sc particle. (2) Since the diffusion coefficient of Zr is smaller than that of Sc, the Al3(Sc, Zr) phase has better thermal stability than the Al3Sc phase. The radius of the Al3(Sc,Zr) particles can maintain its value when annealed at higher temperature for longer time. (3) The critical radius of the Al3Sc precipitation phase strengthening mechanism transformation in Al-0.25Sc alloy is about 2.1 nm. The critical radius of the Al3(Sc, Zr) precipitation phase strengthening mechanism transformation in Al-0.25Sc-0.12Zr alloy is about 2 nm. (4) Under various annealing conditions, both the Al3Sc precipitated phase and the Al3(Sc, Zr) precipitated phase are coherent with the matrix. Acknowledgements This work was supported by the Science and technology Major Project of Hunan Province [grant number 2017GK4002].
References [1] K.V. Yang, Y.J. Shi, F. Palm, X.H. Wu, P. Rometsch, Columnar to equiaxed transition in Al-Mg(-Sc)-Zr alloys produced by selective laser melting, Scripta Mater. 145 (2018) 113-117.
[2] M.J. Li, Y.J. Shi, Q.L. Pan, Y. Zhang, G.X. Lua, S.K. Guan, N. Birbilisc, Low anisotropy of fatigue crack growth in Al-5.8Mg-0.25Sc, Int. J. Fatigue.125(2019)170-178. [3] Y.J. Shi, K. Yang, S.K. Kairy, F. Palm, X.H. Wu, P.A. Rometsch, Effect of platform temperature on the porosity, microstructure and mechanical properties of an Al-Mg-Sc-Zr alloy fabricated by selective laser melting, Mater. Sci. Eng. A 732 (2018) 41-52. [4] A.K. Lohar, B. Mondal, D. Rafaja, V. Klemm, S.C. Panigrahi, Microstructural investigations on as-cast and annealed Al–Sc and Al–Sc– Zr alloys, Mater. Charact. 60 (2009) 1387-1394. [5] V.G. Davydov, T.D. Rostova, V.V. Zakharov, Y.A. Filatov, V.I. Yelagin, Scientific principles of making an alloying addition of scandium to aluminium alloys, Mater. Sci. Eng. A 280 (2000) 30-36. [6] Y. Altinsel, Y. Topkaya, Ş. Kaya, B. Şentürk, Extraction of Scandium from Lateritic Nickel-Cobalt Ore Leach Solution by Ion Exchange: A Special Study and Literature Review on Previous Works, TMS. 2018 1545-1554. [7] Weitao Zhao, Desheng Yang, Lijian Rong, TEM observation of annealing microstructure of deformed Al-Mg-Sc-Zr alloy, Acta Metall. Sin. (China). 41 (2005) 1150-1154. [8] L.M. Dougherty, I.M. Robertson, J.S. Vetrano, Direct observation of the behavior of grain boundaries during continuous dynamic recrystallization in an Al–4Mg–0.3Sc alloy, Acta Mater. 51 (2003) 43674378. [9] D.N. Seidman, E.A. Marquis, D.C. Dunand, Precipitation strengthening at ambient and elevated temperatures of heat-treatable Al (Sc) alloys, Acta Mater. 50 (2002) 4021-4035. [10] E.A. Marquis, D.N. Seidman, Nanoscale structural evolution of Al3Sc precipitates in Al (Sc) alloys, Acta Mate. 49 (2001) 1909-1919. [11] G.B. Teng, C.Y. Liu, Z.Y. Ma, W.B. Zhou, L.L. Wei, Y. Chen, J. Li, Y.F. Mo, Effects of minor Sc addition on the microstructure and mechanical properties of 7055 Al alloy during aging, Mater. Sci. Eng. A 713 (2018) 61-66. [12] M. Vlach, J. Čížek, B. Smola, O. Melikhova, M. Vlček, V. Kodetová, H. Kudrnová, P. Hruška, Heat treatment and age hardening of Al–Si– Mg–Mn commercial alloy with addition of Sc and Zr, Mater. Charact. 129 (2017) 1-8.
[13] Z.H. Jia, J. Røyset, J.K. Solberg, Q. Liu, Formation of precipitates and recrystallization resistance in Al–Sc–Zr alloys, Trans. Nonferrous Metals Soc. China 22 (2012) 1866-1871. [14] S. Costa, H. Puga, J. Barbosa, A.M.P. Pinto, The effect of Sc additions on the microstructure and age hardening behaviour of as cast Al–Sc alloys, Mater. Des. 42 (2012) 347-352. [15] Y. Miyake, Y. Sato, R. Teranishi, K. Kaneko, Effect of heat treatments on the microstructure and formability of Al–Mg–Mn–Sc–Zr alloy, Micron. 101 (2017) 151-155 [16] C. Suwanpreecha, J. P. Toinin, R.A. Michi, P. Pandee, D.C. Dunand, C. Limmaneevichittr, Strengthening mechanisms in Al-Ni-Sc alloys containing Al3Ni microfibers and Al3Sc nanoprecipitates, Acta Mater. 164 (2019) 334-346. [17] J. Taendl, A. Orthacker, H. Amenitsch, G. Kothleitner, C. Poletti, Influence of the degree of scandium supersaturation on the precipitation kinetics of rapidly solidified Al-Mg-Sc-Zr alloys, Acta Mater. 117 (2016) 43-50. [18] P. Xu, F. Jiang, Z. Tang, N. Yan, J. Jiang, X. Xu, Y. Peng, Coarsening of Al3Sc precipitates in Al-Mg-Sc alloys, J Alloy Compd. 781 (2019) 209-215. [19] P. Xu, F. Jiang, M. Tong, Z. Tang, J. Jiang, N. Yan, Y. Peng, Precipitation characteristics and morphological transitions of Al3Sc precipitates, J Alloy Compd. 790(2019) 509-516. [20] W.S. Lee, T.H. Chen, C.F. Lin, M.S. Chen, Impact deformation behavior and dislocation substructure of Al–Sc alloy, J Alloy Compd. 493(2010) 580-589. [21] J. Liu, P. Yao, N. Q. Zhao, C. S. Shi, H. J. Li, X. Li, D. S. Xi, S. Yang, Effect of minor Sc and Zr on recrystallization behavior and mechanical properties of novel Al-Zn-Mg-Cu alloys, J Alloy Compd. 657(2016) 717-725. [22] L. Fu, Y. Li, F.Q. Jiang, J.W. Huang, G.F. Xu, Z.M. Yin, On the role of Sc or Er micro-alloying in the microstructure evolution of Al-Mg alloy sheets during annealing, Mater. Charact. 157 (2019) 109918. [23] S. Iwamura, Y. Miura, Loss in coherency and coarsening behavior of Al3Sc precipitates, Acta Mater. 52 (2004) 591-600.
[24] B. Forbord, W. Lefebvre, F. Danoix, H. Hallem, K. Marthinsen, Three dimensional atom probe investigation on the formation of Al3(Sc,Zr)-dispersoids in aluminium alloys, Scripta Mater. 51 (2004) 333-337. [25] G.B. Teng, C.Y. Liu, Z.Y. Ma, W.B. Zhou, L.L. Wei, Y. Chen, J. Li, Y.E. Mo, Effects of minor Sc addition on the microstructure and mechanical properties of 7055 Al alloy during aging, Mater. Sci. Eng. A 713(2008) 61-66. [26] C. Watanabe, T. Kondo, R. Monzen, Coarsening of Al3Sc precipitates in an Al-0.28 wt pct Sc alloy, Metall. Mater. Trans. A 35 (2004) 3003-3008. [27] C.B. Fuller, J.L. Murray, D.N. Seidman, Temporal evolution of the nanostructure of Al (Sc,Zr) alloys: Part I-Chemical compositions of Al3(Sc1－x Zrx) precipitates, Acta Mater. 53(2005) 5401-5413. [28] C.B. Fuller, D.N. Seidman, Temporal evolution of the nanostructure of Al (Sc,Zr) alloys: Part II-coarsening of Al3(Sc1 － x Zrx) precipitates, Acta Mater. 53(2005) 5415-5428. [29] W. Lefebvrea, F. Danoixa, H. Hallem, Precipitation kinetic of Al3(Sc,Zr) dispersoids in aluminium, J Alloy Compd. 470(2009) 107-110. [30] E. Clouet, L. Laé, Ludovic, T. Épicier, W. Lefebvre, M. Nastar, A. Deschamps, Complex precipitation pathways in multicomponent alloys, Nature Mater. 5(2006) 482-488. [31] A. Tolley, V. Radmilovic, U. Dahmen, Segregation in Al3(Sc,Zr) precipitates in Al–Sc–Zr alloys, Scripta Mater. 52 (2005) 621-625. [32] K.E. Knipling, R.A. Karnesky, C.P. Lee, D.C. Dunand, D.N. Seidman, Precipitation evolution in Al–0.1Sc, Al–0.1Zr and Al–0.1Sc– 0.1Zr (at. %) alloys during isochronal aging, Acta Mater. 58 (2010) 51845195. [33] Booth-Morrison Christopher, D.C. Dunand, D.N. Seidman, Coarsening resistance at 400°C of precipitation-strengthened Al–Zr–Sc– Er alloys, Acta Mater. 59 (2011) 7029-7042. [34] D. Tabor, Br J Appl Phys 7(1956) 159-166. [35] C.L. TANG, D.J. ZHOU, Precipitation hardening behavior of dilute binary Al-Yb alloy, Transactions of Nonferrous Metals Society of China, 24 (2014) 2326-2330.
[36] G.M. Novotny, A.J. Ardell, Precipitation of Al3Sc in binary Al–Sc alloys, Mater. Sci. Eng. A, 318 (2001) 144-154. [37] R.A. Karnesky, D.C. Dunand, D.N. Seidman, Evolution of nanoscale precipitates in Al microalloyed with Sc and Er, Acta Mater. 57 (2009) 4022-4031. [38] Y. Zhang, K. Gao, S. Wen, H. Huang, Z. Nie, D. Zhou, The study on the coarsening process and precipitation strengthening of Al3Er precipitate in Al–Er binary alloy, J Alloy Compd. 610 (2014)7-34.
Table 1. Measured average radii and volume fraction of precipitates Sc-containing aluminum alloy(wt.%) Al-0.25Sc
Al-0.25Sc-0.12Zr
Φx10-3
Annealing condition
r (nm)
290°C /3h
1.5±0.28 0.40±0.01
340°C /3h
2.5±0.50 0.35±0.03
340°C /24h
4±0.67
0.20±0.04
390°C /3h
5±0.92
0.17±0.07
340°C /24h
1.5±0.28 0.42±0.05
390°C /3h
2±0.53
0.47±0.03
Table 2. Experimental and calculated strength increments Sc-containing aluminum alloy(wt.%) Al-0.25Sc
Annealing condition
∆HV/3 (MPa)
290°C /3h
149.9±2.8 148.7 126.4±5.1 139.1
340°C /3h
340°C /24h 73.9±3.4 390°C /3h 71.5±6.3 Al-0.25Sc0.12Zr
∆σord( MPa)
∆σcoh+ ∆σor( ∆σmod MPa) (MPa) 297.6 149.3 324.7
112.7
105.2
283.2
64.9
97.0
279.6
52
304.9
153.2
351.7
146.2
340°C /24h 139.8±2.4 143.0 390°C /3h 157.3±2.4 161.2
Fig.1. Effects of annealing temperature and time on the microhardness of the Al-0.25Sc alloy.
Fig.2. Effects of annealing temperature and time on the microhardness of the Al-0.25Sc-0.12Zr alloy.
Fig.3. Low magnification optical micrographs:(a) Al-0.25Sc annealed at 340°C for 3h; (b) Al-0.25Sc annealed at 340°C for 24h; (c) Al-0.25Sc0.12Zr annealed at 340°C for 3h; (d) Al-0.25Sc-0.12Zr annealed at 340°C for 24h.
Fig.4. Dark field TEM micrographs of Al3Sc precipitates (utilizing a 110 superlattice reflection): (a) 290°C for 3h; (b) SADP of (a); (c) 340°C for 3h; (d) 340°C for 24h; (e) 390°C for 3h.
Fig.5. Dark field TEM micrographs of Al3(Sc,Zr) precipitates (utilizing a 001 superlattice reflection): (a) 340°C for 3h; (b) SADP of (a); (c) 390°C for 3h.
Fig.6. Bright field TEM micrographs of Al3Sc precipitates in Al-0.25Sc alloy annealed at : (a) 290°C for 3h; (b) 340°C for 3h; (c) 340°C for 24h; (d) 390°C for 3h.
Fig.7. Bright field TEM micrographs of Al3(Sc,Zr) precipitates in Al0.25Sc-0.12Zr alloy annealed at : (a) 290°C for 3h; (b) 340°C for 3h; (c) 340°C for 24h; (d) 390°C for 3h.
Fig.8. HRTEM micrographs of Al3Sc precipitates in Al-0.25Sc alloy annealed at (a)340°C for 3h; (b) 390°C for 3h.
Fig.9. HRTEM micrographs of Al3(Sc, Zr) precipitates in Al-0.25Sc0.12Zr alloy annealed at (a) 340°C for 24h and (b) 390°C for 3h.
Fig.10. the distribution range of the precipitates radius: Al3Sc precipitates in Al-0.25Sc alloy annealed at (a)290°C for 3h; (b) 340°C for 3h; (c) 340°C for 24h; (d) 390°C for 3h; Al3(Sc, Zr) precipitates in Al-0.25Sc0.12Zr alloy annealed at (e) 340°C for 24h and (f) 390°C for 3h.
Highlights: 1 The diffusion rate of Sc atoms is inhibited by Zr atoms, which can prevent the coarsening of Al3(Sc, Zr). 2 The hardness decrease in Al-Sc alloy is mainly caused by the coarsening of Al3Sc particle. 3 The radius of the Al3(Sc,Zr) particles can maintain thier value when annealed at higher temperature(>340 °C) for longer time(~24 h).
Declaration of interests The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.