1 Introduction

High-alloyed Al–Zn–Mg–Cu alloys are extensively used in aerospace and military industries because of their high specific strength, stress corrosion cracking resistance and fracture toughness [1, 2]. However, coarse intermetallic particles form in Al–Zn–Mg–Cu alloys with higher alloying element contents [3, 4]. The residual bulky intermetallic compounds (> 1 μm) typically first result in brittle fracture during deformation, which may be the source of crack initiation for Al–Zn–Mg–Cu alloys. Thus, these coarse residual phases can damage the toughness, fatigue properties [5] and corrosion properties of Al–Zn–Mg–Cu alloys [6]. In general, these large soluble non-equilibrium solidification intermetallic phases can be effectively eliminated through their re-dissolution into the matrix during homogenization [7]. However, the high-temperature stable Al2CuMg phase hardly dissolves into the matrix during homogenization [8]. A series of efforts have been made to remove this Al2CuMg phase, Deng et al. [9] developed new multi-step ultrahigh temperature homogenization technology to eliminate the Al2CuMg residual phase. Liu et al. [10] reported that the Al2CuMg phase was not observed after homogenization and was instead replaced by many residual AlZnMgCu phases, in an Al–Zn–Mg–Cu alloy with special high Zn content. Nevertheless, the over-burnt phenomenon may easily occur in actual industrial applications of Al–Zn–Mg–Cu alloys because of the ultrahigh temperature homogenization, and the presence of many residual AlZnMgCu phases in Al–Zn–Mg–Cu alloys with higher Zn content after homogenization may still worsen the performance of these alloys.

Zhang et al. [11] added Ce to an Al–18 wt% Si alloy and observed that the morphology of the primary Si and eutectic Si could be effectively modified from fibrous to lamellar shape. There have also been many reports on the effect of Ce on the morphology of intermetallic compounds of 7XXX series aluminum alloys, e.g., a refined dendritic structure with the precipitate morphology changing from spherical to needle shape in 7055Al alloy [12] and a refined eutectic microstructure in Al–6.7Zn–2.6Mg–2.6Cu (wt%) alloy [13]. The refinement or change of the morphology of the eutectic phase in as-cast alloys can affect the diffusion of segregation elements to the matrix during homogenization [10]. Therefore, the segregation elements, particularly Cu with its lower diffusion coefficient [8], in an appropriate eutectic phase morphology might also be favorably dissolved into the matrix during homogenization annealing. Meanwhile, the reason for adding 0.12 wt% Ce is that the addition of Ce to Al–Zn–Mg–Cu alloy increases the formation and stability of the AlCuCe enrichment phase [14, 15], which would effectively prevent the formation of the high-temperature stable Cu-containing Al2CuMg phases.

In binary Al–Sc hypereutectic alloys with Sc contents above 0.55 wt%, the first solidification phase to form during solidification is Al3Sc, which precipitates as fine particles in the liquid. The L12 crystal structure of these particles and their low lattice parameter mismatch with Al (1.6%) possibly make them effective nucleating agents [16]. However, the use of Sc is not cost-effective, which restricts its wide application. Similar to Sc, the rare-earth element Ce exhibits a liquid → α(Al)-Al11Ce3 eutectic reaction during solidification [17]. The Ce addition has been reported to affect the grain size of Al–Cu–Mg–Mn–Ag alloys [18] and refine the as-cast grains in Al–6.7Zn–2.6Mg–2.6Cu (wt%) alloy [13]. Nevertheless, it is difficult to find primary particles (Al11Ce3) that can act as heterogeneous nucleants at the center of Al grains for alloys with a rare-earth element addition except for Sc. Thus, it is generally accepted that grain refinement in other rare-earth containing alloys is not the result of the nucleation initiation of the first Al3Re phase because their lattice parameter mismatch is very large compared with Al [19].

Herein, the rare-earth element Ce was introduced to study its effect on the solidification behavior and elimination of Al2CuMg phase during homogenization in a high-alloyed Al–Zn–Mg–Cu alloy. The morphological and compositional changes of the eutectic phase in Ce-free and Ce-containing Al–Zn–Mg–Cu alloys after casting and subsequent annealing were investigated. Obvious grain refinement and dissolution of the Al2CuMg phase are observed during homogenization in the Al–Zn–Mg–Cu alloy with 0.12 wt% Ce addition. In addition, the main mechanisms responsible for the remarkable grain refinement during solidification and thorough elimination of the high-temperature residual Al2CuMg phase during homogenization in the Al–Zn–Mg–Cu alloy were proposed. This work provides new insight to aid in improving the performance of high-alloyed Al–Zn–Mg–Cu alloys through Ce addition.

2 Experimental

The experimental alloys were melted by induction heating and then cast into Cu molds to produce billets with 190 mm in length, 130 mm in width and 24 mm in thickness. After the alloy (American series 7136 alloy) or the addition of 0.12 wt% Ce being remelted directly, the measured composition is shown in Table 1, by means of inductively coupled plasma-atomic emission spectrometry (ICP-AES) method on BAIRD PS-6. The billets were homogenized at 440 °C for 12 h and 470 °C for 32 h, followed by air cooling to room temperature. The homogenization treatments were performed in a salt bath furnace with an accuracy of ± 2 K. Differential scanning calorimeter (DSC, METTLER TOLEDO DSC1/700) was performed at a heating rate of 10 K·min−1 to determine the melting points of the intermetallic phases. Phase identification was performed using X-ray diffractometer (XRD, Rigaku D/max-2400) with Cu Kα1 radiation. A combination of optical microscopy (OM, ZEISS Image M2m), scanning electron microscopy (SEM, Nova-Nano SEM450) and electron probe microanalysis (EPMA, JEOL JXA-8230, 20 kV) was used to characterize the as-cast and homogenized microstructures. SEM, in the backscattered electron (BSE) imaging mode, was performed using a scanning electron microscopy attached to EPMA instrument, operating at 20 kV. A quantitative X-ray wavelength-dispersive spectroscopy (WDS) system attached to EPMA instrument was used to analyze the chemistry of the intermetallic phases present in the as-cast and homogenized microstructures. Scanning transmission electron microscopy (STEM) analysis and elemental mapping using energy-dispersive X-ray spectroscopy (EDS) measurements were performed on a JEOL-2100F system.

Table 1 Chemical compositions of investigated alloys (wt%)
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3 Results

3.1 Microstructure of as-cast alloys

Figure 1 presents OM images of as-cast Ce-free and Ce-containing Al–Zn–Mg–Cu alloys. The addition of 0.12 wt% Ce clearly results in efficient grain refinement. The average grain size of the as-cast alloys is refined from ~ 300 μm in Ce-free alloy to 80 μm in Ce-containing alloy. Figure 1b also reveals that as refinement occurs, there is a complete lack of a dendritic substructure within the majority of the fine equiaxed grains which appear to have grown with a spherical front without exhibiting any internal phase or solute partitioning condition. However, a number of primary intermetallic particles are distributed within the fine equiaxed grains, as indicated by arrow in Fig. 1b.

Fig. 1
figure 1

OM images of as-cast alloys: a Al–Zn–Mg–Cu and b Al–Zn–Mg–Cu–0.12Ce

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Figure 2 presents SEM images of as-cast Ce-free and Ce-containing Al–Zn–Mg–Cu alloys. Eutectic solidified structures of Al–Zn–Mg–Cu and Al–Zn–Mg–Cu–0.12Ce alloys can be clearly observed in the enlarged images, respectively. A large number of finer lamellar eutectic network structures form during solidification in the Ce-modified Al–Zn–Mg–Cu alloy. Because of the substantial differences in brightness and morphology of the intermetallic microstructures in the two alloys, the eutectic phases can be classified into four types, marked as A, B, C and D. The corresponding quantitative WDS analysis results are presented in Table 2. The analysis reveals that the gray continuous phase (Area A in Fig. 2a) is AlZnMgCu phase and the gray lamellar eutectic structure phase (Areas B and C in Fig. 2) is a mixture of α(Al) and AlZnMgCu phases. Only Al, Cu and Ce are detected by WDS in the white phase (Area D in Fig. 2b), suggesting that it might be Ce segregated within Al2Cu phase [20].

Fig. 2
figure 2

BSE images and corresponding magnified images of rectangular boxes of as-cast alloys: a Al–Zn–Mg–Cu and b Al–Zn–Mg–Cu–0.12Ce

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Table 2 Chemical composition of intermetallic phases in Fig. 2 (wt%)
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As observed in Figs. 3 and 4, the main elements Zn, Mg and Cu are largely enriched in the dendritic boundary, and the concentration of the elements gradually decreases upon moving from the dendritic boundary to the interior. Cu and Ce tend to be concentrated together with an extreme deficiency of Mg in Al2(Cu, Ce) phase, as observed in Fig. 4. Another type of only Ce-enriched phase with a size of 18 μm may be the primary Al11Ce3 particles that normally form in hypereutectic Al–Ce alloy (encircle in Fig. 4a). Al11Ce3 particles have a typical square shape and a size of ~ 20 μm, similar to the characteristics of primary Al3Sc intermetallic particles [16].

Fig. 3
figure 3

SEM images of microstructure and distribution of main elements in as-cast Al–Zn–Mg–Cu alloy: a BSE image; elemental mappings of b Al, c Zn, d Mg and e Cu (numbers in top right corner being relative concentration of elements)

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Fig. 4
figure 4

SEM images of microstructure and distribution of main elements in as-cast Al–Zn–Mg–Cu–0.12Ce alloy: a BSE image; elemental mappings of b Al, c Zn, d Mg, e Cu and f Ce

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3.2 Microstructure of homogenized alloys

The microstructures of Ce-free and Ce-containing Al–Zn–Mg–Cu alloys after homogenization annealing are shown in Fig. 5. The dissolution of solidification phase occurs within the dendrite cell and along the dendrite boundary after homogenization at 440 °C for 12 h and 470 °C for 32 h. Meanwhile, a new gray intermetallic phase forms, as observed in Fig. 5a. The chemical compositions of the gray phase marked in Fig. 5a are 60.94 at% Al, 19.27 at% Cu and 19.79 at% Mg, as measured by quantitative WDS analysis. The new gray intermetallic phases are thought to be the undissolved Al2CuMg phase because of the Cu to Mg ratio close to 1:1 in this phase. After Ce addition, Al2CuMg phase completely dissolves into the matrix, and only a few microscale AlCuCe intermetallic particles and Al11Ce3 intermetallic particles (square in shape and 16–23 μm in size) are detected, as shown in Fig. 5b. Inside Al matrix, nanoscaled AlCuCe enrichment phases (indicated by arrow in Fig. 5c) in the homogenized Al–Zn–Mg–Cu–0.12Ce alloy are observed. The size of nanoscaled AlCuCe enrichment phases is 200–500 nm, and the particles are marked in Fig. 5c. The elemental distributions of Al, Cu, and Ce in these particles are intensively segregated together, as demonstrated in the mapping results in Fig. 5d. The large number of micro/nanoscale AlCuCe particles distributed within Al matrix in the Ce-modified Al–Zn–Mg–Cu alloy leads to considerable Cu atom trapping and hinders further formation of Al2CuMg phase during homogenization.

Fig. 5
figure 5

BSE images of homogenized alloys: a Al–Zn–Mg–Cu and b Al–Zn–Mg–Cu–0.12Ce; c STEM image of AlCuCe particles; d elemental mapping of Al, Cu and Ce in particle in square box in c

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Figure 6a, b presents DSC plots of Al–Zn–Mg–Cu and Al–Zn–Mg–Cu–0.12Ce alloys in the as-cast and homogenized states. The first endothermic peak (Peak A) at 478.4 °C is observed in both alloys. After homogenization, Peak A is nearly eliminated and another endothermic peak appears at 486.9 °C (Peak B) in both alloys. Peak B is attributed to the dissolution of Al2CuMg phase [9]. However, the intensity of Peak B in Fig. 6b is much less than that in Fig. 6a, indicating that only a small amount of undissolved Al2CuMg phases are formed in Al–Zn–Mg–Cu alloy with the addition of Ce. An additional endothermic peak appears at 578.9 °C (Peak C) in the as-cast Ce-containing alloy, as observed in Fig. 6b. After homogenization, Peak C shifts to even higher temperature (586.9 °C). Based on previous microstructural observations (Figs. 2b, 5b) and results in Ref. [21], the endothermic Peak C is attributed mainly to the melting of Ce-enriched phases, i.e., Al2(Cu, Ce) phases in the as-cast state and AlCuCe phase in homogenized state. The Ce-containing AlCuCe phase clearly exhibits an increased resistance to melting compared with Al2Cu phase, as demonstrated by DSC melting peak of AlCuCe phase appearing above 578 °C, whereas that of Al2Cu phase generally appears at 532 °C [5]. These AlCuCe phases with high-temperature stability are hardly re-dissolved during homogenization; thus, AlCuCe phase can trap many Cu atoms and effectively inhibits further formation of Al2CuMg phase during homogenization.

Fig. 6
figure 6

DSC curves of a Al–Zn–Mg–Cu and b Al–Zn–Mg–Cu–0.12Ce alloys

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XRD patterns of Ce-free and Ce-containing alloys in as-cast and homogenized states are presented in Fig. 7. XRD peaks that originate from MgZn2 (Mg(Zn, Al, Cu)2) phase could be distinguished in both alloys in the as-cast state. After homogenization, Mg(Zn, Al, Cu)2 phase is nearly eliminated, and XRD peaks originating from Al2CuMg phase are distinguished only in Ce-free alloy. Therefore, the elimination of Al2CuMg phase during homogenization in Ce-containing alloy is also supported by XRD results. However, the peak of Ce-containing Al2(Cu, Ce) phase is also observed in as-cast Al–Zn–Mg–Cu–0.12Ce alloy. The AlCuCe phase is not detected after homogenization because of its nanoscale size, which is the experimental limit for X-ray diffractometer used [20]. Based on above analysis, it should be noted that these nanoscaled AlCuCe dispersoids with high-temperature stability might effectively prevent recrystallization during subsequent thermomechanical processing by pinning grain and subgrain boundaries.

Fig. 7
figure 7

XRD patterns of Al–Zn–Mg–Cu and Al–Zn–Mg–Cu–0.12Ce alloys

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4 Discussion

4.1 Effect of Ce addition on grain refinement of as-cast Al–Zn–Mg–Cu alloy

The refinement of as-cast grain size of the high-alloyed Al–Zn–Mg–Cu alloy is far greater than that achieved using conventional grain refiners [22] and other rare-earth elements, e.g., Er [19] and Yb [23] in Al alloys. There is grain refinement with the addition of 0.12 wt% Ce accompanied by a change in the growth morphology. In detail, the large unrefined grains with dendritic growth transform into the fine spherical grains with divorced eutectic structure, which occurs on the grain boundaries in the refined casting (Fig. 1). Similar refinement phenomena have only been observed in Al alloys with Sc addition [16]. For Sc additions less than the eutectic composition of 0.55 wt% Sc, an unrefined large grain size is observed, whereas for Sc contents above 0.55 wt%, a dramatic reduction in grain size is observed. Analogously, Al–Ce binary phase diagram shows that a eutectic reaction occurs at 650 °C with a composition of 0.05 wt% Ce, resulting in the formation of Al11Ce3 intermetallic phase [24]. For the experimental alloy in this study, the addition of 0.12 wt% Ce (which is greater than the eutectic composition of 0.05 wt% Ce) results in a remarkable refinement of the grain size of aluminum casting (from ~ 300 to 80 μm), due to the formation of primary Al11Ce3 intermetallic phase during solidification (Figs. 1b, 2b, 4b). Interesting behavior is also observed when the substructures of the refined cast grains were examined in detail. In the Ce-free alloy, a typical dendritic substructure is observed (Fig. 1a). The first phase to solidify is clearly α(Al), as expected from phase diagram; this phase nucleates predominantly on the mold wall and grows inwards to form large columnar grains. The liquid ahead of the solidification front eventually becomes sufficiently supercooled to allow a region of equiaxed grains to form. In fact, the alloys in this study contain a large amount of alloying elements, resulting in the solute on the front of the solid/liquid interface being enriched during solidification. Thus, the actual temperature in front of the interface is lower than the solidification temperature determined by the solute distribution, which causes a large undercooling during solidification. Therefore, a large dendritic network is observed in the crystal cell in the Ce-free Al–Zn–Mg–Cu alloy. Conversely, in the Ce-containing alloy, as refinement occurs, there is a complete lack of a dendritic substructure within a few of the refine equiaxed grains, which appear to have grown with a spherical front exhibiting no solute partitioning except for a large number of primary internal intermetallic particles distributed within the fine equiaxed grains (Fig. 1b). This exceptional grain refinement could be attributed to the formation of intermetallic particles of Al11Ce3 phase in the melt, which may have acted as nucleation sites for the remaining α(Al) to grow outwards from the primary Al11Ce3 particle on which α(Al) is nucleated. This behavior is typical of that observed for a “divorced eutectic” [16] and is consistent with the difficulty of coupled growth, when the content of one phase is very small because of the dilute nature of the eutectic composition and when the growth kinetics of the Al11Ce3 phase is slow relative to that of α(Al). Nevertheless, the formation of a divorced eutectic would also require very efficient nucleation of α(Al) on the primary Al11Ce3 particle, as on the primary Al3Sc. It is apparent that the matching of the crystal structure and lattice parameter of Al11Ce3 phase to those of α(Al) is the key influencing factor. The nucleation of α(Al) on one of the crystal faces of the primary phase Al11Ce3 in the melt must be very efficient upon solidification. Fortunately, the Al11Ce3 phase has an orthorhombic crystal structure with lattice parameters of a = 0.4389 nm, b = 1.3025 nm and c = 1.0092 nm, and α(Al) has a fcc crystal structure with a lattice parameter of a = 0.40496 nm [25]. Thus, nucleation of α(Al) on one of the crystal faces of the primary phase Al11Ce3 would be very efficient.

4.2 Effect of Ce addition on intermetallic phase morphology during solidification

Mg(Zn, Al, Cu)2 solidification phase is formed in Al–Zn–Mg–Cu alloy when Zn/Mg mass ratio is higher than 2.2 [10]. The Zn/Mg mass ratio is ~ 4.34 or 4.37 in the alloys in the current work; thus, the main phases of the as-cast materials are α(Al) and Mg(Zn, Al, Cu)2 (Fig. 2b). It is well known that Al and Cu will substitute for Zn at the lattice position of Zn in the MgZn2 phase with the formation of Mg(Zn, Al, Cu)2; thus, the fraction of Mg(Zn, Al, Cu)2 is determined by Mg content [26]. Cu and Ce tend to be concentrated together at location with extremely Mg deficiencies, as observed for the Ce-containing alloy (Fig. 4). Therefore, the fraction of Mg(Zn, Al, Cu)2 in the current alloys is not affected by Ce addition. The effect of Mg/Cu mass ratio and Zn content on Mg(Zn, Al, Cu)2 phase after casting are summarized in Table 3 [8, 27,28,29,30,31,32]. The Mg(Zn, Al, Cu)2 phase morphology is significantly affected by Mg/Cu mass ratio. The eutectic transition from residual liquid to lamellar eutectic structured Mg(Zn, Al, Cu)2 phases occurs if Cu content of the residual liquid is sufficiently lower (Mg/Cu mass ratio > 1.1). After the precipitation of crystals of the aluminum solid solution (α(Al)) in the Al–Zn–Mg–Cu alloy begins, the solutes Zn, Mg and Cu are simultaneously diffused to the liquid at the front of solid/liquid interface. Cu atoms are more prone to moving into the residual liquid, whereas Zn and Mg are normally in the form of solid solution (Figs. 3, 4). Therefore, the concentration of the solute Cu in the residual liquid is the key factor determining the final morphology of the Mg(Zn, Al, Cu)2 phase. By combination of the solidification phase observation (Fig. 2b) and DSC analysis of Al–Zn–Mg–Cu–0.12Ce alloy (Fig. 6b), it is apparent that a Ce-containing Al2(Cu, Ce) phase forms easily during solidification. Thus, similar to that of Sc-containing Al–Zn–Mg–Cu alloy, crystallization of Ce-containing Al–Zn–Mg–Cu alloy begins with the precipitation of crystals in the aluminum solid solution followed by the formation of the binary “divorced” eutectic α(Al) + Al11Ce3. After most of the crystallizations of the “divorced” eutectic, crystallization of the binary eutectic α(Al) + Al2(Cu, Ce) first occurs followed by solidification of the alloys, resulting in the formation of α(Al) and Mg(Zn, Al, Cu)2 with lamellar eutectic structures. This result is attributed to the large number of Cu atoms trapped in the Al2(Cu, Ce) phase; therefore, Cu concentration in the final residual liquid is relatively lower, which is favorable for the formation of lamellar eutectic structures during solidification (Fig. 2b). Similar to the finding for a previous reported Er-containing Al–Zn–Mg–Cu alloy (Mg/Cu mass ratio of 0.89) [31], refined lamellar eutectic structures could still be observed in Al–Zn–Mg–Cu alloy with Ce addition, even with Mg/Cu mass ratio ≤ 0.80.

Table 3 Effect of Cu/Mg ratio on Mg(Zn, Al, Cu)2 phase morphology after casting
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4.3 Effect of Ce addition on elimination of undissolved Al2CuMg phase during homogenization

According to the microstructural analysis (Fig. 5) and corresponding compositional analysis of the residual phases after homogenization, it is concluded that the primary solidification phase Mg(Zn, Al, Cu)2 present in the as-cast microstructure disappears. Meanwhile, a new intermetallic phase Al2CuMg forms in the Ce-free alloy. However, the residual Al2CuMg phase is not detected in the Ce-containing alloy. These phenomena can be explained by the diffusion behavior of Cu which has a lower diffusion coefficient than Zn and Mg [8]. According to above analysis, the higher Cu content in the eutectics in Al–Zn–Mg–Cu alloy retards the diffusion process during homogenization heat treatment. Moreover, the higher supersaturation extent of Cu in the Ce-free alloy will provide more Cu solutes to form A12CuMg phase and lead to more stable A12CuMg particles. Hence, the Ce-free alloy with higher Cu content in the supersaturation solid solution contains more A12CuMg particles during homogenization. However, the stable AlCuCe phase in the Ce-containing alloy can trap many Cu atoms and effectively inhibits further formation of the Al2CuMg phase during homogenization.

In addition, comparison of the morphologies of Mg(Zn, Al, Cu)2 in the two Al–Zn–Mg–Cu alloys reveals that the transition from Mg(Zn, Cu, Al)2 to Al2CuMg phases may not occur if Cu is more easily dissolved in Al matrix. Therefore, the significant effect of Mg(Zn, Al, Cu)2 phase morphology can possibly be explained by the diffusion behavior of Cu. The effect of the dendritic thickness (m) on the homogenization time (τs) can be expressed as follows [10]:

$$\tau_{\text{s}} = am^{b}$$
(1)

where a and b are constants specified by the alloys used. A slight increase in m will lead to a significant increase in the time required for complete homogenization. Although m for the two alloys is almost the same, the dendritic thickness inside the refined network of the eutectic Mg(Zn, Cu, Al)2 of Al–Zn–Mg–Cu–0.12Ce alloy (m′) only ranges from 0.2 to 0.4 μm, as observed in Fig. 2d. The phase elimination can be greatly accelerated by decreasing the dendritic thickness. Therefore, Mg(Zn, Al, Cu)2 phase with finer dendritic network structure could be dissolved more rapidly during homogenization. Moreover, Al2CuMg phase in the Al–Zn–Mg–Cu alloy could be sufficiently eliminated with the addition of Ce.

5 Conclusion

The solidification behavior and intermetallic phase evolution during homogenization annealing of an Al–Zn–Mg–Cu alloy with 0.12 wt% Ce addition were examined. The addition of 0.12 wt% Ce results in clear grain refinement (300–80 μm), which is attributed to efficient α(Al) nucleation on one of the crystal faces of the Al11Ce3 primary phase. A large number of finer lamellar eutectic network structures form during solidification in the Ce-modified Al–Zn–Mg–Cu alloy. This result is attributed to the large number of Cu atoms trapped in Al2(Cu, Ce) phase; therefore, Cu concentration in the final residual liquid is relatively lower. In addition, after homogenization, sufficient elimination of Al2CuMg phase is achieved with 0.12 wt% Ce addition. This result is due to the large number of Cu atoms trapped in the stable AlCuCe phase and the formation of finer lamellar eutectic structures, which favor dissolution during homogenization.