Microstructures and Tensile Fracture Behavior of 2219 Wrought Al–Cu Alloys with Different Impurity of Fe

: The Fe-rich intermetallic phases have a broadly detrimental effect on the mechanical properties of Al–Cu alloy. In this paper, the continuous evolution of Fe-rich intermetallics and their effects on mechanical properties, especially the tensile fracture behavior of 2219 wrought Al–Cu alloys as a function of Fe content against different processing approaches (i.e., as-cast, homogenization, multidirectional forging, and solution-peak aging treatment) were investigated using optical microscopy, scanning electron microscopy, and tensile tests. The results indicated that needle-like Al 7 Cu 2 Fe or Al 7 Cu 2 (Fe, Mn) intermetallics mainly presented in the ﬁnal microstructures of all alloys with various Fe contents. The size and number of Al 7 Cu 2 Fe/Al 7 Cu 2 (Fe, Mn) intermetallics increased with the increase of Fe content. The increase of Fe content had little inﬂuence on the ultimate tensile strength and yield strength, while obvious deterioration in the elongation, because fracture initiators mainly occurred at the Al 7 Cu 2 Fe/Al 7 Cu 2 (Fe, Mn) particles or particles–matrix interface. Therefore, the 2219 Al–Cu alloy with 0.2 wt.% Fe content presented relatively low tensile ductility. The tensile fracture mechanism has been discussed in detail.


Introduction
The heat-treatable type 2xxx Al-Cu alloys have been widely used in transportation and aerospace industry owing to its high specific strength and good weldability. Such as the employment of 2219 Al-Cu alloy in fabricating propellant tank of large launch vehicle [1,2] and the utilization of 2519 Al-Cu alloy as structure component of armored vehicle [3]. It is well known that the solubility of impurity Fe in Al-Cu alloys is very poor. As a result, Fe-rich intermetallics, including needle-like morphology Al 7 Cu 2 Fe/Al 7 Cu 2 (Fe, Mn) (named as β-Fe) [4] and Al 3 (FeMn) [5] or Chinese script-shaped Al 15 (FeMn) 3 (CuSi) 2 (named as α-Fe) [6], Al 6 (FeMn) [7], and Al m Fe [8] phases, can be formed readily during the solidification process. Generally, these Fe-rich intermetallics are hard to dissolve into the α-Al matrix during solution treatment and can act as crack initiation sites due to its incompatible plastic deformation with the softer α-Al matrix during deformation, thereby deteriorating the mechanical properties and in particular, the fracture properties of the alloy [2,8]. Therefore, a low upper limit of Fe content can be set in wrought aluminum alloy. For instance, Fe content is usually limited to a maximum amount of 0.3 wt.% in 2219 wrought Al-Cu alloy.
The needle-like β-Fe particles have been thought of as the most detrimental Fe-rich intermetallics. Thus, it is necessary to modify, refine, and even inhibit the formation of the needle-like Fe-rich intermetallics. Several effective methods have been employed to Metals 2021, 11, 174 3 of 16 observed that the coarse constituents were modified and its area fraction could be decreased from the center (69.07%) to the edge (22.10%) in the ultrasonicated ingot. In the work of He et al. [1,23], the number of Al 2 Cu coarse particles was clearly decreased, and the characteristic of Al 2 Cu particles exhibited a more spheroidized shape with increasing the temperature of multi-directional forging. They also observed that increasing the cold predeformation before solution treatment can promote the dissolution of coarse Al 2 Cu particles during solution treatment and a similar phenomenon was reported in the literature [24]. Dong et al. [3] employed cryogenic deformation to 2219 Al alloy forgings prior to solution treatment and found that the coarse Al 2 Cu particles were apt to dissolution because the cryogenic deformation caused more dislocation near the coarse particles. As a large number of dislocation formed around the particles during deformation, these particles could be in a higher energy state and were activated readily. In addition, dislocations also can act as fast channels for atomic diffusion. Thus, more coarse particles could dissolve into the matrix. All of these aforementioned conclusions were developed based on the refinement of the coarse Al 2 Cu particles. However, the evolution of Fe-rich intermetallic particles are rarely reported.
In the present work, we investigated the continuous evolution of Fe-rich intermetallic particles in 2219 wrought Al-Cu alloys as a function of Fe content across casting, homogenization, multidirectional forging (MDF), and solution-peak aging treatment. In addition, a comparative study of corresponding mechanical properties between casting, homogenization, MDF, and solution-peak aging treatment was carried out. Finally, the formation of Fe-rich intermetallic particles and correlation between these particles and tensile fracture behavior were discussed in detail.

Materials and Methods
The 2219 Al-Cu alloys with different Fe contents were melted in an electrical resistance furnace and the final chemical compositions were listed in Table 1. The temperature of the melt was kept at 750-760 • C, followed by gentle stirring, degassing by C 2 Cl 6 , and filtering. The temperature of the melt dropped and held at 720-730 • C for 30 min, and then it was poured into a cylindrical steel mold with a diameter of 100 mm. The as-cast alloys were firstly homogenized at 525 ± 2 • C for 24 h, and then cubic samples measuring 90 mm × 90 mm × 150 mm were cut off from the homogenized ingots. Subsequently, the samples were subjected to complex thermo-mechanical treatment (TMT) process, including 2 cycles of MDF by changing the axis through 90 • (X, Y, Z), adding solution treatment at 545 ± 2 • C for 4 h in the middle of the MDF, and 1 pass of shape-forging with a height reduction of 50%. Before the MDF process, the cubic samples were kept at 450 • C for 1.5 h. The temperature of MDF was around 450 ± 20 • C and the compression speed of MDF was 6 mm/s by using a numerically controlled hydraulic press (YH27-500T; Hefei Forging Machine Co. Ltd., Hefei, China). The schematic diagram of the MDF process was shown in Figure 1. In each upset forging, the samples were compressed along different direction (A, B, C in order) at a forging ratio of 50%. Then samples were flipped around the Y axis and forced to stretch along the un-deformed direction. After MDF, the samples were solution-treated at 537 ± 2 • C for 4 h and immediately quenched in water at room temperature, followed by cold pre-deformation by 3% and artificial aging at 165 • C for 24 h.  The microstructures and fracture morphologies were examined by using a Leica DM4000M optical microscopy (OM; Leica Microsystems, Wizz, German) and a Nova Nano SEM230 scanning electron microscopy (SEM; FEI Co., Hillsboro, OR, USA) equipped with an energy dispersive X-ray spectrometer (EDS). After that, the size of the second-phase particles was the arithmetic mean of at least ten measured values using Image J analyzer software at the basis of ten SEM images from different positions for each sample. The tensile specimens were machined from different processing approaches according to the standard of GB/T228-2002. The gauge dimensions of tensile samples were 70 mm in length, 10 mm in width, and 2 mm in thickness. Following mechanical grinding, the tensile tests were performed in an Instron 3369 electronic universal testing machine (Instron Co., Canton, MA, USA) at room temperature. The extension was 2 mm/min. The data reported below were an average value of at least three independent tensile samples. Figure 2 shows the microstructures of as-cast alloys with Fe content varying from 0.03 to 0.20 wt.%. A typical dendrite features can be observed in all as-cast alloys, while a new phase obviously appeared as needle-like in 0.10 wt.% Fe alloy. The black needle-like intermetallics were observed to increase significantly in both size and amount with the increase of Fe content from 0.1 to 0.2 wt.% (as indicated by the arrows). To confirm these constituents (as marked by the crosses in Figure 2c), EDS was employed (as seen in Figure 3). The results indicated that the Cu-rich intermetallics (Point A) contained 71.74 at.% Al and 28.26 at.% Cu, which agreed with the Al2Cu phase, while the composition of the needle-like phase (Point B) was close to Al7Cu2(Fe, Mn). It is worth noticing that the needle-like Al7Cu2(Fe, Mn) phase mainly distributed across the dendrite, due to its formation through a peritectic reaction. The microstructures and fracture morphologies were examined by using a Leica DM4000M optical microscopy (OM; Leica Microsystems, Wizz, German) and a Nova Nano SEM230 scanning electron microscopy (SEM; FEI Co., Hillsboro, OR, USA) equipped with an energy dispersive X-ray spectrometer (EDS). After that, the size of the second-phase particles was the arithmetic mean of at least ten measured values using Image J analyzer software at the basis of ten SEM images from different positions for each sample. The tensile specimens were machined from different processing approaches according to the standard of GB/T228-2002. The gauge dimensions of tensile samples were 70 mm in length, 10 mm in width, and 2 mm in thickness. Following mechanical grinding, the tensile tests were performed in an Instron 3369 electronic universal testing machine (Instron Co., Canton, MA, USA) at room temperature. The extension was 2 mm/min. The data reported below were an average value of at least three independent tensile samples. Figure 2 shows the microstructures of as-cast alloys with Fe content varying from 0.03 to 0.20 wt.%. A typical dendrite features can be observed in all as-cast alloys, while a new phase obviously appeared as needle-like in 0.10 wt.% Fe alloy. The black needle-like intermetallics were observed to increase significantly in both size and amount with the increase of Fe content from 0.1 to 0.2 wt.% (as indicated by the arrows). To confirm these constituents (as marked by the crosses in Figure 2c), EDS was employed (as seen in Figure 3). The results indicated that the Cu-rich intermetallics (Point A) contained 71.74 at.% Al and 28.26 at.% Cu, which agreed with the Al 2 Cu phase, while the composition of the needlelike phase (Point B) was close to Al 7 Cu 2 (Fe, Mn). It is worth noticing that the needle-like Al 7 Cu 2 (Fe, Mn) phase mainly distributed across the dendrite, due to its formation through a peritectic reaction.  Figure 4 shows the microstructures of homogenized alloys with Fe content varying from 0.10 to 0.20 wt.%. It can be seen that the dendrite features still existed, whereas they became thinner and clearer. However, the overall morphology of the needle-like Al7Cu2(Fe, Mn) intermetallics does not change (as seen in Figure 4d).  Figure 4 shows the microstructures of homogenized alloys with Fe content varying from 0.10 to 0.20 wt.%. It can be seen that the dendrite features still existed, whereas they became thinner and clearer. However, the overall morphology of the needle-like Al7Cu2(Fe, Mn) intermetallics does not change (as seen in Figure 4d).  Figure 4 shows the microstructures of homogenized alloys with Fe content varying from 0.10 to 0.20 wt.%. It can be seen that the dendrite features still existed, whereas they became thinner and clearer. However, the overall morphology of the needle-like Al 7 Cu 2 (Fe, Mn) intermetallics does not change (as seen in Figure 4d).  Figure 5 shows the microstructures of aged alloys with Fe content varying from 0.03 to 0.20 wt.%. It can be seen that the average grain sizes were large and the un-dissolved coarse particles mainly exhibited inside the grain. Meanwhile, the amount and size of the coarse particles decreased obviously (as indicated by the arrows). During the multidirectional forging process, the coarse particles were subjected to stress concentration then can act as crack initiators because of the different elastic modulus of the coarse particle and its matrix counterpart. Moreover, the high density of dislocations can accumulate at their interfaces. These particles could be in a higher energy state and were activated readily, inducing the cracking of these coarse particles. Thus, the dynamic fragmentation of these coarse particles was expected to take place during forging. During the subsequent solution treatment, most of the fragmented Al2Cu particles were dissolved in the alloy matrix, while the Al7Cu2(Fe, Mn) particles were un-dissolved in the alloy matrix due to the different solubility of Cu and Fe in the Al matrix.  Figure 5 shows the microstructures of aged alloys with Fe content varying from 0.03 to 0.20 wt.%. It can be seen that the average grain sizes were large and the un-dissolved coarse particles mainly exhibited inside the grain. Meanwhile, the amount and size of the coarse particles decreased obviously (as indicated by the arrows). During the multidirectional forging process, the coarse particles were subjected to stress concentration then can act as crack initiators because of the different elastic modulus of the coarse particle and its matrix counterpart. Moreover, the high density of dislocations can accumulate at their interfaces. These particles could be in a higher energy state and were activated readily, inducing the cracking of these coarse particles. Thus, the dynamic fragmentation of these coarse particles was expected to take place during forging. During the subsequent solution treatment, most of the fragmented Al 2 Cu particles were dissolved in the alloy matrix, while the Al 7 Cu 2 (Fe, Mn) particles were un-dissolved in the alloy matrix due to the different solubility of Cu and Fe in the Al matrix. In order to evaluate the un-dissolved particles, the microstructures of aged alloys with different Fe contents were further observed by SEM, as shown in Figure 6. The corresponding EDS analysis as shown in Table 2 reveals that the more granular particles were Al2Cu phase (as marked by point A and D) and the rod-like particles with sharp edges were Al7Cu2(Fe, Mn) phase (as marked by point B and C in Figure 6). The reason was that compressive stress can be exerted on Al2Cu and Al7Cu2(Fe, Mn) particles during multidirectional forging process, then the particle can be fragmented when the applied stress was greater than its strength limit. In the subsequent solution treatment, partial Al2Cu particles would be dissolved into the Al matrix, however, the fragmented Al7Cu2(Fe, Mn) particles could not be dissolved into the Al matrix due to the insolubility of Fe in Al-Cu alloy. Namely, compared with the refinement mechanism of Al2Cu particles, only fragmentation took place in Al7Cu2(Fe, Mn) particles. In order to evaluate the un-dissolved particles, the microstructures of aged alloys with different Fe contents were further observed by SEM, as shown in Figure 6. The corresponding EDS analysis as shown in Table 2 reveals that the more granular particles were Al 2 Cu phase (as marked by point A and D) and the rod-like particles with sharp edges were Al 7 Cu 2 (Fe, Mn) phase (as marked by point B and C in Figure 6). The reason was that compressive stress can be exerted on Al 2 Cu and Al 7 Cu 2 (Fe, Mn) particles during multidirectional forging process, then the particle can be fragmented when the applied stress was greater than its strength limit. In the subsequent solution treatment, partial Al 2 Cu particles would be dissolved into the Al matrix, however, the fragmented Al 7 Cu 2 (Fe, Mn) particles could not be dissolved into the Al matrix due to the insolubility of Fe in Al-Cu alloy. Namely, compared with the refinement mechanism of Al 2 Cu particles, only fragmentation took place in Al 7 Cu 2 (Fe, Mn) particles. Table 2. SEM-EDS analysis results of the un-dissolved particles shown in Figure 6 (at. %).    Figure 7 shows the Transmission electron microscopy (TEM) images and corresponding diffraction spots of aged alloys with Fe contents. As seen from the diffraction spots, there were two kinds of perpendicular precipitates, named as θ' phase (coarse precipitates with a mean length of~99 nm) and θ" phase (fine precipitates with a mean length of 26 nm), respectively. It was found that the number and size of these two precipitates did not change obviously with the increase of impurity Fe content. Calculated carefully by the particle-diameter analysis software, the area fraction of the precipitates reduced from 6.5% to 6.1% with increasing the Fe content from 0.03 to 0.20 wt.%. Compared with the tensile tested results with different Fe contents, the change of Fe content had little influence on strength, as shown in Figure 8.    Figure 8 illustrates the mechanical properties against different processing approaches, i.e., as-cast, homogenization, MDF, and solution-peak aging treatment. All of the samples show the same trend of ultimate tensile strength (UTS), yield strength (YS), and elongation (EL) variation from the processes of as-cast to peak aging stage. It also can be seen that under the same processing condition, similar tendency of UTS, YS, and EL variation with increasing the Fe content. The as-cast samples presented relatively low values of UTS/YS/EL, i.e., 165. 35 Figure 9 shows the tensile fracture morphology of aged alloys with different Fe contents. It can be seen that the fracture of tensile specimens with 0.03 wt.% Fe was predominantly dominated by inter-granular fracture (as seen in Figure 9a). Higher magnification observations of a rough surface (as marked by the ellipse in Figure 9a), trans-granular frac-ture caused by the particle-matrix interface decohesion also existed, as shown in Figure 9b. The dimples around the fine spherical Al 2 Cu particles were considerably shallow in the size less than 6 um. As the Fe content increased to 0.10 wt.%, a typical bimodal dimple size distribution was observed on the rough surface and the main fracture mode was diverted to trans-granular fracture, as shown in Figure 9c. The first population of dimples, named as larger primary dimples, were formed by the fracture of Fe-rich constituent particles, whereas the other population of dimples with finer sizes occupied the ligaments between the primary dimples were the result of particle-matrix interface de-cohesion. The similar fracture feature was found in the tensile specimen with 0.20 wt.% Fe (as seen in Figure 9d). However, the population of primary dimples caused by coarse constituent particles was observed to be more and the size of partial dimples increased to 20 µm. In brief, with an increase in Fe content, the un-dissolved Fe-rich impurity particles increased as a result, and tended to crack formation because a higher stress concentration existed at the larger particles during deformation, which was in accord with the tensile tested results.

Fe-Rich Intermetallics Analysis
Referring to Al-Cu-Fe-Mn quaternary phase diagram [25] (as indicated in Figure 10), it can be seen that the Fe-rich phases including Al 3 Fe, Al 6 (FeMn) and Al 7 Cu 2 Fe phases may appear in Al-Cu cast alloys. For the 2219 Al-Cu alloy, the solubility of Mn in Al-Cu alloy can reach 2%, form a supersaturate solid solution as a result, which leads to the formation of Al 20 Mn 3 Cu 2 dispersoids during homogenization. With the increasing of Fe content to a certain extent, the solubility of Mn in Al-Cu alloy significantly reduces and Al 6 (FeMn) phase can form readily, which is in line with the results calculated by software JMatPro 7, as shown in Figure 11. This phenomenon was related to the strong segregation tendency of Fe. Since the solubility of Fe in Al-Cu alloy is quite low, almost all of the Fe segregates to the interface front during solidification and combines with a small amount of Mn to form Al 6 (FeMn) by eutectic reaction (L→α-Al+Al 6 (FeMn)). With a further increase in the content of Fe, the amounts of Al 6 (FeMn) increase, meanwhile, the temperature of eutectic reaction also rises (from about 577 to 608 • C), as shown in Figure 11b-d. But no Al 6 (FeMn) phase existed in 2219 Al-Cu alloy (as shown in Figure 3), which would be explained by the solid-state transformation from Al 6 (FeMn) phase into Al 7 Cu 2 Fe or Al 7 Cu 2 (Fe, Mn) phase during solidification, as shown in Table 3. As a result, the Al 7 Cu 2 Fe or Al 7 Cu 2 (Fe, Mn) phases increased with the increase of the Fe content. In addition, the content of impurity Fe in 2219 Al-Cu alloy is generally controlled below 0.30 wt.%. So, only needle-like Al 7 Cu 2 Fe or Al 7 Cu 2 (Fe, Mn) phase in 2219 Al-Cu cast alloy, which is agreement with the reference of Tseng et al. [7,21].

Reactions T ( • C)
L→α-Al 651-648 L→α-Al+Al 6 (FeMn) 608-577 L+ Al 6 (FeMn)→α-Al+Al 20 Mn 3 Cu 2 +Al 7 Cu 2 Fe 597-576 L→α-Al +Al 20   Based on the Al-Cu-Fe-Mn quaternary phase diagram (as seen in Figure 10), the calculation results by the software JMatPro 7 (seen in Figure 11), the metallographic observations with SEM, and those reported by Liu et al. [13,26,27], the possible solidification reactions are listed in Table 3. During the solidification, the liquid alloys are nucleated at about 648-651 • C and the α-Al dendrite network forms, then the eutectic and peritectic reactions shall take place. If there is only a small amount of Fe existing in Al-Cu-Mn alloy, the solidification ends with the formation of Al 7 Cu 2 Fe, Al 2 Cu, and Al 20 Mn 3 Cu 2 though the ternary eutectic reactions L→α-Al + Al 2 Cu + Al 20 Mn 3 Cu 2 (e 1 -E line) and L→α-Al + Al 2 Cu + Al 7 Cu 2 Fe (e 2 -E line). However, the amount of Al 7 Cu 2 Fe phase was too small to be found in 0.03 wt.% Fe cast alloy (as shown in Figure 2a). With the increasing of Fe content to 0.10 wt.%, Al 6 (FeMn) intermetallic phase shall precipitate firstly at about 600 • C, and then the peritectic transformation L+Al 6 (FeMn)→α-Al+Al 20 Mn 3 Cu 2 +Al 7 Cu 2 Fe (p 1 -P line or p 2 -P line) leads to form the Al 7 Cu 2 Fe intermetallics, which always distribute across the dendrite network (as shown in Figure 2b). The solidification ends with the formation of the eutectic reactions L→α-Al+Al 2 Cu+Al 20 Mn 3 Cu 2 +Al 7 Cu 2 Fe. For a further increase in Fe content, the solidification sequence of 0.15 wt.% Fe alloy and 0.20 wt.% Fe alloy remains unchanged. However, the precipitation temperature for Al 6 (FeMn) intermetallic phase rises, and its volume fraction also goes up, which leads to the increase of Al 7 Cu 2 Fe intermetallic phase (as shown in Figure 2c,d). Figure 8 indicated that both MDF and heat treatment could improve the tensile fracture behavior. For a metallographic point of view, the fracture mechanisms of 2219 wrought Al-Cu alloys are simultaneously dominated by the multi-scale second-phase particles, and grains and grain boundaries [28][29][30].

Effect of the Fe-Rich Intermetallic Particles on Tensile Fracture Behavior of 2219 Al-Cu Alloys
Firstly, as mentioned in Figures 2, 4 and 5, coarse grains with an average size of about 380 ± 40 um (average of at least 100 grains) were observed in 2219 Al-Cu-xFe alloys under different processing condition. The large initial grain size was controlled by the solidification condition. However, during MDF at 450 • C, dynamic recovery readily occurred and therefore the deformation energy was timely relieved for grain growth during subsequent solution treatment. Thus, coarse grains were obtained after solution-peak aging treatment, which was agreement with the reference of Dong et al. [1][2][3]19,20,31]. Therefore, the effect of grains and grain boundaries on tensile fracture behavior of 2219 Al-Cu alloys as a function of Fe content can be considered to be identical.
Secondly, as mentioned in Figure 7, large amounts of precipitates (θ' and θ") were observed in 2219 Al-Cu-xFe (x = 0.03, 0.20 wt.%) alloys. In general, volume fraction, size and characteristic of precipitates were of significant roles in influencing the tensile fracture properties. The UTS and YS values of solution-peak aging stage 2219 Al-Cu-xFe alloys were higher than those of as cast, homogenized and MDF processed 2219 Al-Cu-xFe alloys, while their EL values in solution-peak aging stage were relatively lower than MDF processed alloys due to strengthening precipitates could reduce the ductility of alloys. Compared Figure 7a with Figure 7b, the number, size and area fraction of precipitates (θ' and θ") did not change obviously with the increase of impurity Fe content. Consequently, the change of Fe content had little influence on values of strength in peak aging heat treatment.
Finally, as mentioned in Figures 2, 4 and 5, the number, size, area fraction, and characteristic of coarse intermetallic particles decreased obviously against different processing approaches (i.e., as-cast, homogenization, MDF, and solution-peak aging treatment). The coarse intermetallics are always considered as hard-brittle phases and therefore those particles have no deformation abilities. Under external service loading, they were easy to dehisce or separate from the matrix and so acted as crack initiators. Therefore, the decrease of primary coarse intermetallics could also improve the tensile fracture properties of alloys.
As mentioned in Figures 2 and 3, the main constituents in as-cast 2219 Al-Cu-xFe alloys were Al 2 Cu and Al 7 Cu 2 (Fe, Mn) phases. Through different processing approaches, Al 2 Cu particles exhibited an ellipse or spherical shape due to the interactions with fragmentation, dissolution, and diffusion, whereas, Al 7 Cu 2 (Fe, Mn) particles were just fragmented into rod-like with sharp edges due to the insolubility of Fe in Al-Cu alloy, as shown in Figures 4-6. To further understand the mechanical characteristics of coarse second-phase particles in the alloys, Pugh [32,33] proposed a method to predict the ductility of the second particles based on their B/G and υ. Here, B and G is bulk modulus and shear modulus, respectively, and the values of B and G were calculated using the Voigt-Reuss-Hill approximation from the elastic constant of the second-phase particles. Poisson's ratio (υ) can be derived from B and G using formula υ = (3B − 2G)/2(3B + G). In the study of Tian et al. [34][35][36], the B/G ratio and υ values of Al 2 Cu phase are 2.65 and 0.332, respectively. However, the B/G ratio and υ values of Al 7 Cu 2 (Fe, Mn) phase are 1.31 and 0.154, respectively. This means that Al 7 Cu 2 (Fe, Mn) particles are easier to act as crack initiators than Al 2 Cu particles. Therefore, the greater the number of Al 7 Cu 2 (Fe, Mn) particles was, the lower the plastic nature (as shown in Figure 8). In addition, rod-like Al 7 Cu 2 (Fe, Mn) particles with sharp edges would fracture readily because they were subjected to higher stress concentration, and then cracks propagated along themselves (as shown in Figure 9c,d).

Conclusions
In this paper, the microstructures evolution of Fe-rich intermetallics, mechanical properties of 2219 Al-Cu alloys with different Fe content against different processing approaches (i.e., as-cast, homogenization, MDF, and solution-peak aging treatment) were studied. The main conclusions are as follows.
(1) When the Fe content was less than 0.03 wt.%, the main constituents were Al 2 Cu intermetallics. As the Fe content increased to 0.10 wt.%, a new needle-like Al 7 Cu 2 Fe or Al 7 Cu 2 (Fe, Mn) phase presented. Further increase in the Fe content, the characteristic of the needle-like Al 7 Cu 2 Fe or Al 7 Cu 2 (Fe, Mn) intermetallics did not change, whereas their sizes became longer and wider. (2) The fragmented Al 7 Cu 2 Fe or Al 7 Cu 2 (Fe, Mn) intermetallics were obtained during multidirectional forging process. However, they were un-dissolved in the α-Al matrix in subsequent solution treatment due to the low tolerance of Fe in Al-Cu alloys. The sharp edges of the fragmented Al 7 Cu 2 Fe or Al 7 Cu 2 (FeMn) particles can act as crack initiators and then as crack propagation paths because they were subjected to higher stress concentrations during deformation.