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Journal of Alloys and Compounds 823 (2020) 153831
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The effect of pre-deformation on the precipitation behavior of AlCuMg(Si) alloys with low Cu/Mg ratios F.J. Niu a, J.H. Chen a, S.Y. Duan a, W.Q. Ming a, J.B. Lu b, C.L. Wu a, *, Z. Le a a b
Center for High-Resolution Electron Microscopy, College of Materials Science and Engineering, Hunan University, Changsha, Hunan, 410082, China School of Physics and Information Technology, Shaanxi Normal University, Xi’an, 710119, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 23 September 2019 Received in revised form 17 December 2019 Accepted 12 January 2020 Available online 13 January 2020
The addition of a trace amount of Si significantly affects the precipitation sequence of the AleCueMg alloy with a low Cu/Mg ratio, such as forming Si-modified Guinier-Preston-Bagaryatsky (GPB) zones to suppress the precipitation of the S phase. The introduction of pre-deformation also affects the precipitation behavior during the subsequent aging process. In this paper, the atomic-resolution high-angleannular-dark-field (HAADF) imaging technique and first-principles calculation are used to study the effect of pre-deformation on the precipitation behavior of an Al-3.0Cu-1.8Mg-0.5Si alloy during aging at 180 C. In order to learn the effect of minor Si-addition, a Si-free alloy with similar chemical composition is also studied. It is found that for Si-free alloy, the 6% pre-deformation promotes the precipitation of the S phase dispersedly during the subsequent aging, leading to an increase of the peak hardness. For Sicontaining alloy, the 6% pre-deformation prior to aging suppresses the precipitation of Si-modified GPB zone acting as the original main strengthening precipitates and promotes the formation of successive composite precipitates including the S phase, various GPB zones, C phase and sub-units of C/Qʹ phase. The successive composite precipitates are easy to grow up and coarsen, which weakens the contribution of precipitation strengthening compared to the small Si-modified GPB zone. As such, the 6% pre-deformation prior to aging at 180 C provides a positive strengthening effect for Si-free alloy, but a negative strengthening effect for the Si-containing alloy. © 2020 Elsevier B.V. All rights reserved.
Keywords: Aluminum alloys Pre-deformation Precipitation Transmission electron microscopy First-principles calculation
1. Introduction AleCueMg alloys belong to the heat-treatable aluminum alloys developed for structural applications, which are widely used in the aerospace and other industries due to their high strength, good ductility and excellent creep strength at high temperature [1]. The strengthening precipitates of AleCueMg alloy with a low Cu/Mg ratio mainly include S phase (Al2CuMg) and GPB zone (GuinierPreston-Bagaryatsky) [2e4]. The S phase, which has a face-centered orthorhombic structure with the lattice parameters aS ¼ 0.400 nm, bS ¼ 0.923 nm and cS ¼ 0.714 nm [5,6], always forms on {012}Al habit planes and grows mainly along the <100>Al direction [3]. The GPB zone usually consists of agglomerated one-dimensional (1D) crystals that have a periodicity of 4.05 Å along the <100>Al direction [7,8]. The competitive relationship between the S phase and GPB zone in aluminum alloys with different compositions has been
* Corresponding author. E-mail address: [email protected] (C.L. Wu). https://doi.org/10.1016/j.jallcom.2020.153831 0925-8388/© 2020 Elsevier B.V. All rights reserved.
focused [9e12]. It is generally believed that the aging temperature [10] and composition of alloys [13] can adjust the precipitation sequence and affect the competitive relationship between the phases. The formation of the S phase precedes that of the GPB zone when the aging temperature is below 180 C, while the S phase and the GPB zone appear simultaneously when aging temperature is above 180 C [10]. The addition of trace amounts of Si can promote the uniform precipitation of Si-modified GPB zone to improve the stability of the GPB zone and inhibit the precipitation of the S phase, thereby improving the aging response speed and heatresistance of the alloy [14e17]. In addition, dislocations also have significant influence on the subsequent precipitation. It is believed that the formation of zigzag S precipitates is related to the quenching dislocations in some cases [18e20]. For 2024 alloy strengthened by the zigzag-arranged S phase and GPB zone, predeformation prior to aging promotes the fine rod-shaped S phase and more excellent U phase to precipitate in the subsequent aging process, and suppress the precipitation of the GPB zone [21e23]. Generally, the combination of deformation and aging can improve
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the mechanical strength of the aluminum alloys, where the dislocation strengthening, grain boundary strengthening and precipitation strengthening are applied to achieve a good combination of strength and plasticity through adjusting deformation and aging [24e26]. For example, dislocations produced by pre-deformation can promote the precipitation of a stripe wall-like phase with a width of 1e2 nm and a habit plane of (100)Al or (110)Al in an Al0.75Mg-0.75Si alloys, leading to higher strength and better elongation combined with work hardening [25]. But inappropriate thermo-mechanical treatment may not get the expected results, such as for 7050 alloys [27] and Al-5.74Cu-0.44Mg-0.88Ag alloys [28]. Pre-deformation followed by artificial aging results in the coarser precipitates and change in type and volume fraction of precipitates, which leads to a decrease in strength compared to the T6 aged alloys without pre-deformation. The addition of Si may cause some precipitation of Si-containing precipitates in the some alloys, such as bʺ, bʹ and Qʹ phase [13,29]. Moreover, the participation of Si atoms can improve the stability of the GPB zone by forming the so-called Si-modified GPB zone [14,15]. Therefore, the effect of pre-deformation prior to aging on the precipitation sequence of Si-containing AleCueMg alloys may be more complicated than that of AleCueMg alloys with similar composition [30,31]. For example, a variety of precipitates were observed along dislocation lines in an Al-4.9Cu-0.74Mg-0.51Si0.48Mn (wt. %) alloy with the small deformation prior to aging [32]. A new precipitate, E phase was discovered in the predeformed Al-1.3Cu-1.0Mg-0.4Si (wt. %) alloy during the aging process [33]. In this paper, high-resolution transmission electron microscopy (TEM), especially high-angle-annular-dark-field (HAADF) imaging in scanning transmission electron microscopy (STEM) was used to study the precipitation behaviors of Al-3.0Cu-1.8Mg-0.5Si (wt. %) alloy with or without pre-deformation in the process of aging at 180 C. For comparison, the effect of pre-deformation on the precipitation behaviors of Al-3.1Cu-1.9 Mg alloy without Si addition was also investigated. It is shown that the pre-deformation causes the Si-containing alloy to produce more complex precipitates in the process of aging at 180 C.
aging treatment at 180 C, which is named T8 heat treatment.
2.2. The HRTEM instrument The TEM characterization was carried out on FEI Tecnai G2 F20 electron microscope and Titan Cubed Themis G2 300 spherical aberration correction electron microscope, whose resolutions are 0.2 nm and 0.06 nm in HAADF-STEM mode, respectively. The F20 electron microscope was operated at 200 kV with a convergence semi-angle of 10 mrad and collection angle range of 36e200 mrad. The spherical aberration correction electron microscope was operated at 300 kV with a convergence semi-angle of 25.6 mrad and collection angle range of 67e200 mrad.
2.3. Image simulation analysis To verify the proposed atomic structure, image simulation was performed by multislice algorithm with Debye model implemented in an in-house software [34]. In image simulations, large enough supercells including the Al matrix and the precipitates were used. It should be noted that the comparison with experimental image is not quantitative since (1) the absolute value of STEM signal in this experiment is unknown and (2) STEM simulation costs too much time. The imaging parameters used in the simulation are: 300 kV, Cs ¼ 1.5 mm, df ¼ 0 nm, semi-convergent angle: 25.6 mrad, and the detector range: 67e200 mrad. The Debye Waller factors are all set 0.83338 which is the standard value of Al at 300 K [35], since all atoms are embedded in Al matrix. To account for residual probe aberrations and finite source size, a Gaussian function, with a full width at half maximum (FWHM) of 0.06 nm, was applied. The sample thickness is set about 4 nm according to the relative contrast between the simulated and experimental images. We also simulated the STEM image of sample with thickness of 10 nm and 16 nm, and found that the results are very similar.
2.4. First-principles energy calculations 2. Experiments and methods 2.1. Alloy samples and thermal aging The experimental two alloys used in this paper are self-made Al3.0Cu-1.8Mg-0.5Si (wt. %) alloy and Al-3.1Cu-1.9 Mg (wt. %) alloy, which can be called the Si-containing alloy and the Si-free alloy. In addition, the two alloys contain some trace elements, such as 0.15 wt % Zn, 0.1 wt % Ti and 0.1 wt % Fe. The alloys were melted in vacuum arc melting furnace and cast into some ingots. After cutting the both end of ingots, the ingots were measured by the nondestructive flaw detection to insure no obvious defects. Chemical analysis was applied to measure the contents of solute elements which was described above. The qualified ingots were homogenized at 510 C for 24 h in an air-cycling furnace, hot-rolled at 450 C from 20 mm to 5 mm, and then cold-rolled to sheets of 2 mm in thickness. Small pieces with the size of 10 mm 10 mm 2 mm were wire-cut from the coldrolled sheets. The solution temperature is 535 C for Si-containing alloy and 505 C for Si-free alloy, respectively, according to the differential scanning calorimetry of the incipient melting point. All samples were solution heat treated for 1 h and then quenched into the water at room temperature. Then some samples were immediately aged at 180 C in a constant temperature oil bath furnace, which is called T6 heat treatment. Some samples underwent 6% pre-deformation after solution treatment, followed by the same
In this paper, the first-principles calculation within density functional theory is used to investigate the formation enthalpy of the supercell composed of the precipitate and matrix. The lower the formation enthalpy, the higher the stability of the precipitate. The PAW (Projector Augmented Wave) potential with PBE (PerdewBurke-Ernzerhof) exchange correlation function, namely PAW-PBE, is chosen in calculation. The convergence test shows that the cutoff energy of 355 eV can obtain high-precision calculation results for the AlCuMgSi alloy system. For all structures, k-point sampling was performed using the Monkhorst-Pack method to achieve an accuracy of 0.1 kJ/mol [36]. It must be ensured that supercells in all calculations are large enough to include interfacial energy and strain energy. It is worth noting that there is an a/2 shift between the aluminum matrix at the two sides of the new phase along the [001] direction, so vacuum layers must be introduced at both sides of the supercell containing the new phase. Because only in this case the formation enthalpy is negative and the fully relaxed structure matches the experimental HAADF image perfectly. The thickness of vacuum layer is 13.0 Å to prevent the interaction between the slab and its periodic images. The surface energy resulting from the matrix and vacuum must be considered when the formation enthalpy is calculated, so the calculation formula of formation enthalpy for the supercell composed of AlaCubMgcSid and containing vacuum layers can be defined as follows [37,38]:
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DHðAla Cub Mgc Sid Þ ¼
EðAla Cub Mgc Sid Þ aEðAlÞ bEðCuÞ cEðMgÞ dEðSiÞ 2As bþcþd
Among them, E(AlaCubMgcSid) is the total energy of the supercell, E(Al), E(Cu), E(Mg), E(Si) are the energy of a single atom in the elemental state of four elements. Among them, “a”, “b”, “c” and “d” are the number of corresponding atoms belonging to the calculated supercell. The unit of formation enthalpy is kJ/mol solute atom1. s is on behalf of surface energy, referring to the surface energy of the (110)Al plane of the aluminum in this article. A is the area of the aluminum matrix in contact with the vacuum layers. The calculation method of surface energy is based on reference [38]. It is calculated that the Al(001) and Al(011) slabs with more than six layers can converge to 0.91 J/m2 and 0.96 J/m2, which are almost in agreement with the previous calculation results, 0.90 J/m2 and 0.97 J/m2, respectively[38,39]. 3. Results and discussion Fig. 1 shows the hardness curves for the Si-free alloy and the Sicontaining alloy with different heat treatment. The hardness of the Si-free alloy treated by T6 temper reached the peak at 12 h, and the peak hardness was about 134 HV. The T8-treated samples reached the peak at 9 h, and the peak hardness was about 146 HV, which means the introduction of pre-deformation improved the peakhardness value of the Si-free alloy. The hardness of the Si-containing alloy reached the peak at 48 h after T6 artificial aging, and the peak hardness was approximately 153 HV. While the T8-treated samples reached the peak after aging for 18 h, and the peak hardness was about 140 HV, indicating the introduction of pre-deformation lowered the hardness of the Sicontaining alloy processed by T6 heat treatment. The microstructure characterization of the peak-aged Si-free alloys are shown in Fig. 2. The peak-aged samples processed by T6 heat treatment is mainly strengthened by the zigzag-arranged S precipitates and the uniformly distributed GPB zones. After 6% predeformation, the nucleation of the GPB zone is suppressed during the subsequent aging process. So almost all of the precipitates in the peak-aged T8 samples without Si addition belong to the S phase, which is uniformly distributed in the matrix due to the high density of the heterogeneous nucleation sites induced by the dislocations. Figs. 3e6 display the microstructure of the Si-containing alloys processed by T6 and T8 tempers . Judging from HAADF-STEM images in Fig. 3aec, the main precipitates of the peak-aged Si-containing alloys treated by T6 process are a great deal of needle-like
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Si-modified GPB zone and a small amount of rod-like S phase. In the peak-aged samples processed by T8 heat treatment, the number density of the uniformly distributed precipitates including the Simodified GPB zone and rod-like S phase in the undeformed bulk areas decreases. But a large number of successive precipitates are formed, which can be referred to as the “zig-zag composite precipitates” in this paper. Some parts of the zig-zag composite precipitates belong to the S phase, and others are some unknown structures labeled by the frames in Fig. 3f. Furthermore, the rod-like S phase in Fig. 3c and the lath-like S phase in Fig. 3f possess the same crystal structure but different habit plane. The lath-like S phases have the classic orientation relationship ([100]S//[100]Al, (001)S//(021)Al), and the rod-like S phases have an orientation relationship rotated by 2e6 [40e42]. The habit planes of the S precipitates are parallel to (120)Al (or close to (120)Al), but the zig-zag composite precipitates, which are connected into curved walls, include some periodic structures lying on (110)Al and (100)Al planes. These unknown structures are so complex that it’s difficult to clearly identify them using the Tecnai F20 transmission electron microscope with a resolution of 0.2 nm. Therefore, the spherical aberration-corrected electron microscopy is applied to confirm the structure of the zig-zag composite precipitates. Fig. 4aeb further illustrate that the large parts of the zig-zag composite precipitates belong to the S phase, and other parts consist of a variety of precipitates. The C phase (AlCu0.7Mg4Si3.3) which belongs to the P21/m space group with lattice parameters a ¼ 1.032 nm, b ¼ 0.81 nm, c ¼ 0.405 nm, g ¼ 101 [43] can be frequently observed in the zig-zag composite precipitates. The phase relationship between C phase and the matrix is: [100]Al// [100]C, [010]Al//[150]C. Furthermore, it is found that the precipitates with a habit plane of (100)Al are all determined to be C phase in this Si-containing alloy. Besides, there are a lot of normal GPB zones or sub-units of GPB zone appearing at the interface between the precipitates and matrix. In addition, there are some darker regions inside the zig-zag composite precipitates and an unknown phase lying on the (110)Al plane of matrix, which are shown in Figs. 5 and 6, respectively. As shown in Fig. 5, the interior of some zig-zag composite precipitates appears to be long-range disordered and looks like atomic clusters without periodic structure. However, compared with the unit cells of the C phase and the Qʹ phase (AlxCu2Mg12-xSi7) [44], these clusters turn out to be the sub-units constituting the C phase or Qʹ phase, which are composed of Cu, Mg and Si atoms and named as “sub-unit of C/Qʹ phase” in this paper. Due to the instability of the sub-units, there are some GPB zones or S precipitates at the interface between the sub-units and matrix, which serve to reduce the interfacial energy.
Fig. 1. The age hardening curves of Si-free alloy (a) and Si-containing alloy (b) with different heat treatment. Samples selected for TEM investigations are marked with blue circles. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
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Fig. 2. TEM images of the peak-aged Si-free alloys processed by T6 heat treatment (a-b) and T8 heat treatment (c-d). All images are viewed along the [001]Al direction.
Fig. 3. HAADF-STEM images of the peak-aged Si-containing alloys processed by T6 heat treatment (a-c) and T8 heat treatment (d-f), and (f) is the magnified image of a black framed area of (e). The insets in (c) and (f) are the high-magnification images of the corresponding red framed areas. The S phase in (c) and (f) were identified to possedd the same crystal structure but different habit planes. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
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Fig. 4. HAADF-STEM images of the zig-zag composite precipitates in peak-aged samples of Si-containing alloys processed by T8 heat treatment, (c) and (d) are high-magnification images of the framed areas labeled in (a) and (b), respectively.
A rectangular phase with a coherent interface of the (110)Al plane also appears frequently in some zig-zag composite precipitates as shown in Fig. 6. Comparing atomic brightness in the HAADF-STEM images, we establish all possible supercell models of the rectangular phase, one of which is shown in Fig. 6g. The atomic types at eight positions (symmetric positions are considered to be occupied by the same atomic type) labeled as Positions 1e8 in Fig. 6g can’t be determined immediately. The Al atomic columns in an individual HAADF-STEM image have similar brightness, so the atomic type belonging to the precipitate can be inferred by comparing the brightness of each atomic column with that of Al atomic columns through Gatan Digitalmicrogragh. When the brightness of an atomic site is close to or slightly darker than that of Al atom, the atomic site may be occupied by Al or Mg atom. When the one’s brightness is close to or slightly brighter than that of Al, the atomic site may be occupied by Al or Si atom. When the one’s brightness is much brighter than that of Al, the atomic site should be occupied by Cu atom. Here, it can be speculated that the possible atomic types at Position 1, 3, 4, 5 may be Al or Mg; the possible atomic type at Position 2 is Al or Si; the possible atomic type at Position 6 is Al or Mg; the possible atomic types at Position 7 and 8 are possible to be Al or Si. So there are 128 combinations and the lowest formation enthalpy of possible supercell models is determined using first-principles calculation. Next, the first-principles calculation is performed for the formation enthalpies of possible supercell models, and the calculation results are listed in Table 1. It has been accepted that the negative
formation enthalpy of supercells means the decrease in the energy of system through the formation of precipitate and represents the rationality of structure model. It is worth noting that the formation enthalpy is much lower when Position 6, 7 and 8 are occupied by Mg, Al and Al, respectively (the detailed calculation results of which are not listed due to the limited article length). So in the Table 1, we do not care the situation that Position 6, 7 and 8 are occupied by Al, Si and Si, respectively. Judging from Table 1, the formation enthalpy of supercell models is the lowest when the Positions 1e8 are occupied by Al, Si, Mg, Mg, Al, Mg, Al and Al respectively (No. 11), and the corresponding structure composed of Al33Cu4Mg8Si is the most stable as well. Image simulation was performed to verify whether this model fit the experimental results. Fig. 7 demonstrates the filtered atomicresolution HAADF-STEM images of the rectangular phase and simulated image based on the most stable structure as shown in Fig. 6g. By comparing the atomic contrast of the simulated image and the experimental image, the correctness of the structure can be proved by a perfect matching. The structure of rectangular phase is similar but different with E phase which has been discovered in AleCueMgeSi alloys before [33]. In addition, the formation enthalpies are also very low when the Position 2 or 3 are occupied by Al atoms in the similar conditions of No. 11 in Table 1, which is also possible to be observed in experimental results. Comparing this rectangular phase with the conventional GPB zone appearing in Fig. 6, we can find that the rectangular phase is similar with conventional GPB zone and can be
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Fig. 5. Sub-units constituting of C phase or Qʹ phase inside the zig-zag composite precipitates in peak-aged samples of Si-containing alloys processed by T8 heat treatment.
considered as a variant of the GPB zone, referred to as GPB zonevariant. It is worth noting that the formation enthalpy is the lowest when Si atom occupies the Position 2. The addition of Si plays an important role in the precipitation of the rectangular phase, considering that the rectangular GPB zone-variant cannot be found in the Si-free alloy. Generally, the nucleation and growth of the precipitate belong to solid phase transformation. The nucleation methods include uniform nucleation and heterogeneous nucleation. The nonequilibrium defects such as grain boundary, inclusion interface, sub-grain boundary and dislocation increase the total free energy of system, leading to the preference of the heterogeneous nucleation. Dislocations can also always serve as fast diffusion paths for solute atoms [45,46] and trap solute atoms to promote the precipitation process of AleCueMg alloys during aging. The combination of deformation and aging can significantly change the precipitation sequence and macroscopic properties of the alloy [47e49]. For the Si-free alloys, the peak-aged sample processed by T6 heat treatment is strengthened by uniformly distributed GPB zones and zigzag-shaped S precipitates and the latter plays a major role in strengthening. After solution treatment followed by the 6% predeformation, Cu and Mg atoms accumulate near dislocations to promote the precipitation of the S phase and inhibit the nucleation of GPB zone in the subsequent aging process. In this case, the higher density of S precipitates with more uniform distribution in association with work hardening improves the performance of Si-free alloys.
For the Si-containing alloy, the peak-aged sample treated by T6 aging process is strengthened mainly by a great deal of Si-modified GPB zone and a small amount of the S phase. For the deformed Sicontaining alloy, the dislocation also promotes the precipitation of the S phase and inhibits the precipitation of Si-modified GPB zone during the subsequent aging process. The Si-modified GPB zone as the main strengthening phase is suppressed, which must be detrimental to the mechanical properties of the Si-containing alloy. Furthermore, Si atoms also move to the dislocations like Cu and Mg atoms to form the C phase, sub-unit of C/Qʹ phase and GPB zone-variant in the subsequent aging process. The variant of GPB zone lying on the habit plane of (110)Al also contains Si atoms. These Si-containing precipitates are combined with S precipitates and form the zig-zag composite precipitates. In order to reduce the interfacial energy between zig-zag composite precipitates and matrix, GPB zones or sub-units of GPB zone form at the interface between these precipitates and the matrix. Therefore, a large number of successive composite precipitates, including S phase, C phase, conventional GPB zone, GPB zone-variant and sub-unit of C/ Qʹ phase appear in the Si-containing alloy after pre-deformation and aging. All precipitates induced by dislocations are arranged continuously in the form of zig-zag morphology and become much coarser with the extension of aging time, which is also detrimental to the performance. Therefore, the performance of the Sicontaining alloy processed by T8 treatment is degraded due to the coarse zig-zag composite precipitates and a serious reduction in the density of the Si-modified GPB zones despite the effect of work hardening.
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Fig. 6. The atomic-resolution HAADF-STEM images of the rectangular phase lying on the (110)Al plane formed in the peak-aged Si-containing alloys processed by T8 heat treatment (a-f) and the schematic diagram of a supercell model composed of the rectangular phase and matrix (g). Among them, (a) (b) are the same image and (c) (d) are the same image. The conventional GPB zone is also observed in (b), whose single-cell model is shown in (g). All images are viewed along the [001]Al direction.
Table 1 Formation enthalpies (DH) of different supercell models composed of the rectangular phase and matrix. No.
P1
P2
P3
P4
P5
P6
P7
P8
DH (kJ/mol solute atom1)
1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32
Mg Mg Mg Mg Al Al Al Al Al Al Al Al Mg Mg Mg Mg Mg Mg Mg Mg Al Al Al Al Al Al Al Al Mg Mg Mg Mg
Si Si Si Si Si Si Si Si Si Si Si Si Si Si Si Si Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al
Mg Mg Mg Mg Al Al Al Al Mg Mg Mg Mg Al Al Al Al Mg Mg Mg Mg Al Al Al Al Mg Mg Mg Mg Al Al Al Al
Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al
Mg Al Al Mg Mg Al Al Mg Mg Al Al Mg Mg Al Al Mg Mg Al Al Mg Mg Al Al Mg Mg Al Al Mg Mg Al Al Mg
Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg
Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al
Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al
15.74 13.86 19.87 10.38 17.65 17.73 22.56 11.41 19.02 18.82 23.85 13.68 14.88 12.82 18.58 8.35 11.28 12.52 17.31 7.91 14.56 17.39 22.23 10.29 15.25 18.64 22.49 11.93 10.84 11.35 17.11 6.46
Fig. 7. The filtered atomic-resolution HAADF-STEM images of the rectangular phase and the corresponding simulated image.
4. Conclusions Using TEM and HAADF-STEM in association with first-principles energy calculation, the precipitation behaviors of Al-3.0Cu-1.8Mg0.5Si (wt. %) and Al-3.1Cu-1.9 Mg (wt. %) alloy aged at 180 C with or without 6% pre-deformation were investigated. It can be concluded as follows: (1) For the Si-free alloy, the GPB zones and the zig-zag S phases are the main strengthening precipitates in the T6 peak-aged samples. After 6% pre-deformation followed by aging, a high density of the S precipitates with a uniform distribution become the main precipitates in the T8 peak-aged sample, which improve the hardness of the alloy. So the hardness of the T8-treated samples is higher than that of the T6-treated samples. (2) For the Si-containing alloy, the T6 peak-aged sample is mainly strengthened by Si-modified GPB zone. After 6% predeformation followed by aging, the homogenous
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precipitation of Si-modified GPB zones is suppressed, and the successive zig-zag composite precipitates become the main strengthening precipitates, leading to the degradation of performance eventually. So the hardness of the T8-treated samples is lower than that of T6-treated samples. (3) In T8-treated Si-containing alloys, the zig-zag composite precipitates include S phase, conventional GPB zone, GPB zone-variant, C phase, and the sub-units of the C/Qʹ phase. It is worth noting that a two-dimensional rectangular GPB zone-variant with a coherent habit-plane interface parallel to (110)Al was frequently observed, which reduced the energy of system and improved the stability of the zig-zag composite precipitates.
Author contributions section F.J. Niu: Conceptualization, Methodology, Validation, WritingOriginal Draft. C.L. Wu: Funding acquisition, Methodology, Writing-review and editing. J.H. Chen: Investigation, Project administration, Supervision. S.Y. Duan: Resources, Data curation. W.Q. Ming: Software, Visualization. J.B. Lu: Resources. Z. Le: Resources. Declaration of competing interest 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. Acknowledgements This work is supported by the National Natural Science Foundation of China (Nos. 11427806, 51831004, 51671082, 51471067) and the National Key Research and Development Program of China (No. 2016YFB0300801). References [1] J.C. Williams, E.A. Starke, Progress in structural materials for aerospace systems11The golden jubilee issuedselected topics in materials science and engineering: past, present and future, in: S. Suresh (Ed.), Acta Mater, 2003, pp. 5775e5799, 51(19). [2] S.B. Wang, J.H. Chen, M.J. Yin, Z.R. Liu, D.W. Yuan, J.Z. Liu, C.H. Liu, C.L. Wu, Double-atomic-wall-based dynamic precipitates of the early-stage S-phase in AlCuMg alloys, Acta Mater. 60 (19) (2012) 6573e6580. [3] S.C. Wang, M.J. Starink, Precipitates and intermetallic phases in precipitation hardening AleCueMge(Li) based alloys, Int. Mater. Rev. 50 (4) (2005) 193e215. [4] S.C. Wang, M.J. Starink, N. Gao, Precipitation hardening in AleCueMg alloys revisited, Scr. Mater. 54 (2) (2006) 287e291. [5] W.A. Perlitz H, Ark Kemi Mineral och Geol B 16 (1943) 13. [6] Z.R. Liu, J.H. Chen, S.B. Wang, D.W. Yuan, M.J. Yin, C.L. Wu, The structure and the properties of S-phase in AlCuMg alloys, Acta Mater. 59 (19) (2011) 7396e7405. [7] L. Kovarik, M.J. Mills, Structural relationship between one-dimensional crystals of GuinierePrestoneBagaryatsky zones in AleCueMg alloys, Scr. Mater. 64 (11) (2011) 999e1002. [8] L. Kovarik, M.J. Mills, Ab initio analysis of GuinierePrestoneBagaryatsky zone nucleation in AleCueMg alloys, Acta Mater. 60 (9) (2012) 3861e3872. [9] A. Charai, T. Walther, C. Alfonso, A.M. Zahra, C.Y. Zahra, Coexistence of clusters, GPB zones, S0-, S’- and s-phases in an Al±0.9% Cu±1.4% Mg alloy, Acta Mater. 48 (10) (2000) 2751e2764. [10] M.J. Yin, J.H. Chen, S.B. Wang, Z.R. Liu, L.M. Cha, S.Y. Duan, C.L. Wu, Anisotropic and temperature-dependent growth mechanism of S-phase precipitates in AleCueMg alloy in relation with GPB zones, Trans. Nonferrous Metals Soc. China 26 (1) (2016) 1e11. [11] S. Cheng, Y.H. Zhao, Y.T. Zhu, E. Ma, Optimizing the strength and ductility of fine structured 2024 Al alloy by nano-precipitation, Acta Mater. 55 (17) (2007)
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
The effect of pre-deformation on the precipitation behavior of AlCuMg(Si) alloys with low Cu/Mg ratios F.J. Niu a, J.H. Chen a, S.Y. Duan a, W.Q. Ming a, J.B. Lu b, C.L. Wu a, *, Z. Le a a b
Center for High-Resolution Electron Microscopy, College of Materials Science and Engineering, Hunan University, Changsha, Hunan, 410082, China School of Physics and Information Technology, Shaanxi Normal University, Xi’an, 710119, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 23 September 2019 Received in revised form 17 December 2019 Accepted 12 January 2020 Available online 13 January 2020
The addition of a trace amount of Si significantly affects the precipitation sequence of the AleCueMg alloy with a low Cu/Mg ratio, such as forming Si-modified Guinier-Preston-Bagaryatsky (GPB) zones to suppress the precipitation of the S phase. The introduction of pre-deformation also affects the precipitation behavior during the subsequent aging process. In this paper, the atomic-resolution high-angleannular-dark-field (HAADF) imaging technique and first-principles calculation are used to study the effect of pre-deformation on the precipitation behavior of an Al-3.0Cu-1.8Mg-0.5Si alloy during aging at 180 C. In order to learn the effect of minor Si-addition, a Si-free alloy with similar chemical composition is also studied. It is found that for Si-free alloy, the 6% pre-deformation promotes the precipitation of the S phase dispersedly during the subsequent aging, leading to an increase of the peak hardness. For Sicontaining alloy, the 6% pre-deformation prior to aging suppresses the precipitation of Si-modified GPB zone acting as the original main strengthening precipitates and promotes the formation of successive composite precipitates including the S phase, various GPB zones, C phase and sub-units of C/Qʹ phase. The successive composite precipitates are easy to grow up and coarsen, which weakens the contribution of precipitation strengthening compared to the small Si-modified GPB zone. As such, the 6% pre-deformation prior to aging at 180 C provides a positive strengthening effect for Si-free alloy, but a negative strengthening effect for the Si-containing alloy. © 2020 Elsevier B.V. All rights reserved.
Keywords: Aluminum alloys Pre-deformation Precipitation Transmission electron microscopy First-principles calculation
1. Introduction AleCueMg alloys belong to the heat-treatable aluminum alloys developed for structural applications, which are widely used in the aerospace and other industries due to their high strength, good ductility and excellent creep strength at high temperature [1]. The strengthening precipitates of AleCueMg alloy with a low Cu/Mg ratio mainly include S phase (Al2CuMg) and GPB zone (GuinierPreston-Bagaryatsky) [2e4]. The S phase, which has a face-centered orthorhombic structure with the lattice parameters aS ¼ 0.400 nm, bS ¼ 0.923 nm and cS ¼ 0.714 nm [5,6], always forms on {012}Al habit planes and grows mainly along the <100>Al direction [3]. The GPB zone usually consists of agglomerated one-dimensional (1D) crystals that have a periodicity of 4.05 Å along the <100>Al direction [7,8]. The competitive relationship between the S phase and GPB zone in aluminum alloys with different compositions has been
* Corresponding author. E-mail address: [email protected] (C.L. Wu). https://doi.org/10.1016/j.jallcom.2020.153831 0925-8388/© 2020 Elsevier B.V. All rights reserved.
focused [9e12]. It is generally believed that the aging temperature [10] and composition of alloys [13] can adjust the precipitation sequence and affect the competitive relationship between the phases. The formation of the S phase precedes that of the GPB zone when the aging temperature is below 180 C, while the S phase and the GPB zone appear simultaneously when aging temperature is above 180 C [10]. The addition of trace amounts of Si can promote the uniform precipitation of Si-modified GPB zone to improve the stability of the GPB zone and inhibit the precipitation of the S phase, thereby improving the aging response speed and heatresistance of the alloy [14e17]. In addition, dislocations also have significant influence on the subsequent precipitation. It is believed that the formation of zigzag S precipitates is related to the quenching dislocations in some cases [18e20]. For 2024 alloy strengthened by the zigzag-arranged S phase and GPB zone, predeformation prior to aging promotes the fine rod-shaped S phase and more excellent U phase to precipitate in the subsequent aging process, and suppress the precipitation of the GPB zone [21e23]. Generally, the combination of deformation and aging can improve
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the mechanical strength of the aluminum alloys, where the dislocation strengthening, grain boundary strengthening and precipitation strengthening are applied to achieve a good combination of strength and plasticity through adjusting deformation and aging [24e26]. For example, dislocations produced by pre-deformation can promote the precipitation of a stripe wall-like phase with a width of 1e2 nm and a habit plane of (100)Al or (110)Al in an Al0.75Mg-0.75Si alloys, leading to higher strength and better elongation combined with work hardening [25]. But inappropriate thermo-mechanical treatment may not get the expected results, such as for 7050 alloys [27] and Al-5.74Cu-0.44Mg-0.88Ag alloys [28]. Pre-deformation followed by artificial aging results in the coarser precipitates and change in type and volume fraction of precipitates, which leads to a decrease in strength compared to the T6 aged alloys without pre-deformation. The addition of Si may cause some precipitation of Si-containing precipitates in the some alloys, such as bʺ, bʹ and Qʹ phase [13,29]. Moreover, the participation of Si atoms can improve the stability of the GPB zone by forming the so-called Si-modified GPB zone [14,15]. Therefore, the effect of pre-deformation prior to aging on the precipitation sequence of Si-containing AleCueMg alloys may be more complicated than that of AleCueMg alloys with similar composition [30,31]. For example, a variety of precipitates were observed along dislocation lines in an Al-4.9Cu-0.74Mg-0.51Si0.48Mn (wt. %) alloy with the small deformation prior to aging [32]. A new precipitate, E phase was discovered in the predeformed Al-1.3Cu-1.0Mg-0.4Si (wt. %) alloy during the aging process [33]. In this paper, high-resolution transmission electron microscopy (TEM), especially high-angle-annular-dark-field (HAADF) imaging in scanning transmission electron microscopy (STEM) was used to study the precipitation behaviors of Al-3.0Cu-1.8Mg-0.5Si (wt. %) alloy with or without pre-deformation in the process of aging at 180 C. For comparison, the effect of pre-deformation on the precipitation behaviors of Al-3.1Cu-1.9 Mg alloy without Si addition was also investigated. It is shown that the pre-deformation causes the Si-containing alloy to produce more complex precipitates in the process of aging at 180 C.
aging treatment at 180 C, which is named T8 heat treatment.
2.2. The HRTEM instrument The TEM characterization was carried out on FEI Tecnai G2 F20 electron microscope and Titan Cubed Themis G2 300 spherical aberration correction electron microscope, whose resolutions are 0.2 nm and 0.06 nm in HAADF-STEM mode, respectively. The F20 electron microscope was operated at 200 kV with a convergence semi-angle of 10 mrad and collection angle range of 36e200 mrad. The spherical aberration correction electron microscope was operated at 300 kV with a convergence semi-angle of 25.6 mrad and collection angle range of 67e200 mrad.
2.3. Image simulation analysis To verify the proposed atomic structure, image simulation was performed by multislice algorithm with Debye model implemented in an in-house software [34]. In image simulations, large enough supercells including the Al matrix and the precipitates were used. It should be noted that the comparison with experimental image is not quantitative since (1) the absolute value of STEM signal in this experiment is unknown and (2) STEM simulation costs too much time. The imaging parameters used in the simulation are: 300 kV, Cs ¼ 1.5 mm, df ¼ 0 nm, semi-convergent angle: 25.6 mrad, and the detector range: 67e200 mrad. The Debye Waller factors are all set 0.83338 which is the standard value of Al at 300 K [35], since all atoms are embedded in Al matrix. To account for residual probe aberrations and finite source size, a Gaussian function, with a full width at half maximum (FWHM) of 0.06 nm, was applied. The sample thickness is set about 4 nm according to the relative contrast between the simulated and experimental images. We also simulated the STEM image of sample with thickness of 10 nm and 16 nm, and found that the results are very similar.
2.4. First-principles energy calculations 2. Experiments and methods 2.1. Alloy samples and thermal aging The experimental two alloys used in this paper are self-made Al3.0Cu-1.8Mg-0.5Si (wt. %) alloy and Al-3.1Cu-1.9 Mg (wt. %) alloy, which can be called the Si-containing alloy and the Si-free alloy. In addition, the two alloys contain some trace elements, such as 0.15 wt % Zn, 0.1 wt % Ti and 0.1 wt % Fe. The alloys were melted in vacuum arc melting furnace and cast into some ingots. After cutting the both end of ingots, the ingots were measured by the nondestructive flaw detection to insure no obvious defects. Chemical analysis was applied to measure the contents of solute elements which was described above. The qualified ingots were homogenized at 510 C for 24 h in an air-cycling furnace, hot-rolled at 450 C from 20 mm to 5 mm, and then cold-rolled to sheets of 2 mm in thickness. Small pieces with the size of 10 mm 10 mm 2 mm were wire-cut from the coldrolled sheets. The solution temperature is 535 C for Si-containing alloy and 505 C for Si-free alloy, respectively, according to the differential scanning calorimetry of the incipient melting point. All samples were solution heat treated for 1 h and then quenched into the water at room temperature. Then some samples were immediately aged at 180 C in a constant temperature oil bath furnace, which is called T6 heat treatment. Some samples underwent 6% pre-deformation after solution treatment, followed by the same
In this paper, the first-principles calculation within density functional theory is used to investigate the formation enthalpy of the supercell composed of the precipitate and matrix. The lower the formation enthalpy, the higher the stability of the precipitate. The PAW (Projector Augmented Wave) potential with PBE (PerdewBurke-Ernzerhof) exchange correlation function, namely PAW-PBE, is chosen in calculation. The convergence test shows that the cutoff energy of 355 eV can obtain high-precision calculation results for the AlCuMgSi alloy system. For all structures, k-point sampling was performed using the Monkhorst-Pack method to achieve an accuracy of 0.1 kJ/mol [36]. It must be ensured that supercells in all calculations are large enough to include interfacial energy and strain energy. It is worth noting that there is an a/2 shift between the aluminum matrix at the two sides of the new phase along the [001] direction, so vacuum layers must be introduced at both sides of the supercell containing the new phase. Because only in this case the formation enthalpy is negative and the fully relaxed structure matches the experimental HAADF image perfectly. The thickness of vacuum layer is 13.0 Å to prevent the interaction between the slab and its periodic images. The surface energy resulting from the matrix and vacuum must be considered when the formation enthalpy is calculated, so the calculation formula of formation enthalpy for the supercell composed of AlaCubMgcSid and containing vacuum layers can be defined as follows [37,38]:
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DHðAla Cub Mgc Sid Þ ¼
EðAla Cub Mgc Sid Þ aEðAlÞ bEðCuÞ cEðMgÞ dEðSiÞ 2As bþcþd
Among them, E(AlaCubMgcSid) is the total energy of the supercell, E(Al), E(Cu), E(Mg), E(Si) are the energy of a single atom in the elemental state of four elements. Among them, “a”, “b”, “c” and “d” are the number of corresponding atoms belonging to the calculated supercell. The unit of formation enthalpy is kJ/mol solute atom1. s is on behalf of surface energy, referring to the surface energy of the (110)Al plane of the aluminum in this article. A is the area of the aluminum matrix in contact with the vacuum layers. The calculation method of surface energy is based on reference [38]. It is calculated that the Al(001) and Al(011) slabs with more than six layers can converge to 0.91 J/m2 and 0.96 J/m2, which are almost in agreement with the previous calculation results, 0.90 J/m2 and 0.97 J/m2, respectively[38,39]. 3. Results and discussion Fig. 1 shows the hardness curves for the Si-free alloy and the Sicontaining alloy with different heat treatment. The hardness of the Si-free alloy treated by T6 temper reached the peak at 12 h, and the peak hardness was about 134 HV. The T8-treated samples reached the peak at 9 h, and the peak hardness was about 146 HV, which means the introduction of pre-deformation improved the peakhardness value of the Si-free alloy. The hardness of the Si-containing alloy reached the peak at 48 h after T6 artificial aging, and the peak hardness was approximately 153 HV. While the T8-treated samples reached the peak after aging for 18 h, and the peak hardness was about 140 HV, indicating the introduction of pre-deformation lowered the hardness of the Sicontaining alloy processed by T6 heat treatment. The microstructure characterization of the peak-aged Si-free alloys are shown in Fig. 2. The peak-aged samples processed by T6 heat treatment is mainly strengthened by the zigzag-arranged S precipitates and the uniformly distributed GPB zones. After 6% predeformation, the nucleation of the GPB zone is suppressed during the subsequent aging process. So almost all of the precipitates in the peak-aged T8 samples without Si addition belong to the S phase, which is uniformly distributed in the matrix due to the high density of the heterogeneous nucleation sites induced by the dislocations. Figs. 3e6 display the microstructure of the Si-containing alloys processed by T6 and T8 tempers . Judging from HAADF-STEM images in Fig. 3aec, the main precipitates of the peak-aged Si-containing alloys treated by T6 process are a great deal of needle-like
3
Si-modified GPB zone and a small amount of rod-like S phase. In the peak-aged samples processed by T8 heat treatment, the number density of the uniformly distributed precipitates including the Simodified GPB zone and rod-like S phase in the undeformed bulk areas decreases. But a large number of successive precipitates are formed, which can be referred to as the “zig-zag composite precipitates” in this paper. Some parts of the zig-zag composite precipitates belong to the S phase, and others are some unknown structures labeled by the frames in Fig. 3f. Furthermore, the rod-like S phase in Fig. 3c and the lath-like S phase in Fig. 3f possess the same crystal structure but different habit plane. The lath-like S phases have the classic orientation relationship ([100]S//[100]Al, (001)S//(021)Al), and the rod-like S phases have an orientation relationship rotated by 2e6 [40e42]. The habit planes of the S precipitates are parallel to (120)Al (or close to (120)Al), but the zig-zag composite precipitates, which are connected into curved walls, include some periodic structures lying on (110)Al and (100)Al planes. These unknown structures are so complex that it’s difficult to clearly identify them using the Tecnai F20 transmission electron microscope with a resolution of 0.2 nm. Therefore, the spherical aberration-corrected electron microscopy is applied to confirm the structure of the zig-zag composite precipitates. Fig. 4aeb further illustrate that the large parts of the zig-zag composite precipitates belong to the S phase, and other parts consist of a variety of precipitates. The C phase (AlCu0.7Mg4Si3.3) which belongs to the P21/m space group with lattice parameters a ¼ 1.032 nm, b ¼ 0.81 nm, c ¼ 0.405 nm, g ¼ 101 [43] can be frequently observed in the zig-zag composite precipitates. The phase relationship between C phase and the matrix is: [100]Al// [100]C, [010]Al//[150]C. Furthermore, it is found that the precipitates with a habit plane of (100)Al are all determined to be C phase in this Si-containing alloy. Besides, there are a lot of normal GPB zones or sub-units of GPB zone appearing at the interface between the precipitates and matrix. In addition, there are some darker regions inside the zig-zag composite precipitates and an unknown phase lying on the (110)Al plane of matrix, which are shown in Figs. 5 and 6, respectively. As shown in Fig. 5, the interior of some zig-zag composite precipitates appears to be long-range disordered and looks like atomic clusters without periodic structure. However, compared with the unit cells of the C phase and the Qʹ phase (AlxCu2Mg12-xSi7) [44], these clusters turn out to be the sub-units constituting the C phase or Qʹ phase, which are composed of Cu, Mg and Si atoms and named as “sub-unit of C/Qʹ phase” in this paper. Due to the instability of the sub-units, there are some GPB zones or S precipitates at the interface between the sub-units and matrix, which serve to reduce the interfacial energy.
Fig. 1. The age hardening curves of Si-free alloy (a) and Si-containing alloy (b) with different heat treatment. Samples selected for TEM investigations are marked with blue circles. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
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Fig. 2. TEM images of the peak-aged Si-free alloys processed by T6 heat treatment (a-b) and T8 heat treatment (c-d). All images are viewed along the [001]Al direction.
Fig. 3. HAADF-STEM images of the peak-aged Si-containing alloys processed by T6 heat treatment (a-c) and T8 heat treatment (d-f), and (f) is the magnified image of a black framed area of (e). The insets in (c) and (f) are the high-magnification images of the corresponding red framed areas. The S phase in (c) and (f) were identified to possedd the same crystal structure but different habit planes. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
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Fig. 4. HAADF-STEM images of the zig-zag composite precipitates in peak-aged samples of Si-containing alloys processed by T8 heat treatment, (c) and (d) are high-magnification images of the framed areas labeled in (a) and (b), respectively.
A rectangular phase with a coherent interface of the (110)Al plane also appears frequently in some zig-zag composite precipitates as shown in Fig. 6. Comparing atomic brightness in the HAADF-STEM images, we establish all possible supercell models of the rectangular phase, one of which is shown in Fig. 6g. The atomic types at eight positions (symmetric positions are considered to be occupied by the same atomic type) labeled as Positions 1e8 in Fig. 6g can’t be determined immediately. The Al atomic columns in an individual HAADF-STEM image have similar brightness, so the atomic type belonging to the precipitate can be inferred by comparing the brightness of each atomic column with that of Al atomic columns through Gatan Digitalmicrogragh. When the brightness of an atomic site is close to or slightly darker than that of Al atom, the atomic site may be occupied by Al or Mg atom. When the one’s brightness is close to or slightly brighter than that of Al, the atomic site may be occupied by Al or Si atom. When the one’s brightness is much brighter than that of Al, the atomic site should be occupied by Cu atom. Here, it can be speculated that the possible atomic types at Position 1, 3, 4, 5 may be Al or Mg; the possible atomic type at Position 2 is Al or Si; the possible atomic type at Position 6 is Al or Mg; the possible atomic types at Position 7 and 8 are possible to be Al or Si. So there are 128 combinations and the lowest formation enthalpy of possible supercell models is determined using first-principles calculation. Next, the first-principles calculation is performed for the formation enthalpies of possible supercell models, and the calculation results are listed in Table 1. It has been accepted that the negative
formation enthalpy of supercells means the decrease in the energy of system through the formation of precipitate and represents the rationality of structure model. It is worth noting that the formation enthalpy is much lower when Position 6, 7 and 8 are occupied by Mg, Al and Al, respectively (the detailed calculation results of which are not listed due to the limited article length). So in the Table 1, we do not care the situation that Position 6, 7 and 8 are occupied by Al, Si and Si, respectively. Judging from Table 1, the formation enthalpy of supercell models is the lowest when the Positions 1e8 are occupied by Al, Si, Mg, Mg, Al, Mg, Al and Al respectively (No. 11), and the corresponding structure composed of Al33Cu4Mg8Si is the most stable as well. Image simulation was performed to verify whether this model fit the experimental results. Fig. 7 demonstrates the filtered atomicresolution HAADF-STEM images of the rectangular phase and simulated image based on the most stable structure as shown in Fig. 6g. By comparing the atomic contrast of the simulated image and the experimental image, the correctness of the structure can be proved by a perfect matching. The structure of rectangular phase is similar but different with E phase which has been discovered in AleCueMgeSi alloys before [33]. In addition, the formation enthalpies are also very low when the Position 2 or 3 are occupied by Al atoms in the similar conditions of No. 11 in Table 1, which is also possible to be observed in experimental results. Comparing this rectangular phase with the conventional GPB zone appearing in Fig. 6, we can find that the rectangular phase is similar with conventional GPB zone and can be
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Fig. 5. Sub-units constituting of C phase or Qʹ phase inside the zig-zag composite precipitates in peak-aged samples of Si-containing alloys processed by T8 heat treatment.
considered as a variant of the GPB zone, referred to as GPB zonevariant. It is worth noting that the formation enthalpy is the lowest when Si atom occupies the Position 2. The addition of Si plays an important role in the precipitation of the rectangular phase, considering that the rectangular GPB zone-variant cannot be found in the Si-free alloy. Generally, the nucleation and growth of the precipitate belong to solid phase transformation. The nucleation methods include uniform nucleation and heterogeneous nucleation. The nonequilibrium defects such as grain boundary, inclusion interface, sub-grain boundary and dislocation increase the total free energy of system, leading to the preference of the heterogeneous nucleation. Dislocations can also always serve as fast diffusion paths for solute atoms [45,46] and trap solute atoms to promote the precipitation process of AleCueMg alloys during aging. The combination of deformation and aging can significantly change the precipitation sequence and macroscopic properties of the alloy [47e49]. For the Si-free alloys, the peak-aged sample processed by T6 heat treatment is strengthened by uniformly distributed GPB zones and zigzag-shaped S precipitates and the latter plays a major role in strengthening. After solution treatment followed by the 6% predeformation, Cu and Mg atoms accumulate near dislocations to promote the precipitation of the S phase and inhibit the nucleation of GPB zone in the subsequent aging process. In this case, the higher density of S precipitates with more uniform distribution in association with work hardening improves the performance of Si-free alloys.
For the Si-containing alloy, the peak-aged sample treated by T6 aging process is strengthened mainly by a great deal of Si-modified GPB zone and a small amount of the S phase. For the deformed Sicontaining alloy, the dislocation also promotes the precipitation of the S phase and inhibits the precipitation of Si-modified GPB zone during the subsequent aging process. The Si-modified GPB zone as the main strengthening phase is suppressed, which must be detrimental to the mechanical properties of the Si-containing alloy. Furthermore, Si atoms also move to the dislocations like Cu and Mg atoms to form the C phase, sub-unit of C/Qʹ phase and GPB zone-variant in the subsequent aging process. The variant of GPB zone lying on the habit plane of (110)Al also contains Si atoms. These Si-containing precipitates are combined with S precipitates and form the zig-zag composite precipitates. In order to reduce the interfacial energy between zig-zag composite precipitates and matrix, GPB zones or sub-units of GPB zone form at the interface between these precipitates and the matrix. Therefore, a large number of successive composite precipitates, including S phase, C phase, conventional GPB zone, GPB zone-variant and sub-unit of C/ Qʹ phase appear in the Si-containing alloy after pre-deformation and aging. All precipitates induced by dislocations are arranged continuously in the form of zig-zag morphology and become much coarser with the extension of aging time, which is also detrimental to the performance. Therefore, the performance of the Sicontaining alloy processed by T8 treatment is degraded due to the coarse zig-zag composite precipitates and a serious reduction in the density of the Si-modified GPB zones despite the effect of work hardening.
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Fig. 6. The atomic-resolution HAADF-STEM images of the rectangular phase lying on the (110)Al plane formed in the peak-aged Si-containing alloys processed by T8 heat treatment (a-f) and the schematic diagram of a supercell model composed of the rectangular phase and matrix (g). Among them, (a) (b) are the same image and (c) (d) are the same image. The conventional GPB zone is also observed in (b), whose single-cell model is shown in (g). All images are viewed along the [001]Al direction.
Table 1 Formation enthalpies (DH) of different supercell models composed of the rectangular phase and matrix. No.
P1
P2
P3
P4
P5
P6
P7
P8
DH (kJ/mol solute atom1)
1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32
Mg Mg Mg Mg Al Al Al Al Al Al Al Al Mg Mg Mg Mg Mg Mg Mg Mg Al Al Al Al Al Al Al Al Mg Mg Mg Mg
Si Si Si Si Si Si Si Si Si Si Si Si Si Si Si Si Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al
Mg Mg Mg Mg Al Al Al Al Mg Mg Mg Mg Al Al Al Al Mg Mg Mg Mg Al Al Al Al Mg Mg Mg Mg Al Al Al Al
Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al Mg Al
Mg Al Al Mg Mg Al Al Mg Mg Al Al Mg Mg Al Al Mg Mg Al Al Mg Mg Al Al Mg Mg Al Al Mg Mg Al Al Mg
Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg Mg
Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al
Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al Al
15.74 13.86 19.87 10.38 17.65 17.73 22.56 11.41 19.02 18.82 23.85 13.68 14.88 12.82 18.58 8.35 11.28 12.52 17.31 7.91 14.56 17.39 22.23 10.29 15.25 18.64 22.49 11.93 10.84 11.35 17.11 6.46
Fig. 7. The filtered atomic-resolution HAADF-STEM images of the rectangular phase and the corresponding simulated image.
4. Conclusions Using TEM and HAADF-STEM in association with first-principles energy calculation, the precipitation behaviors of Al-3.0Cu-1.8Mg0.5Si (wt. %) and Al-3.1Cu-1.9 Mg (wt. %) alloy aged at 180 C with or without 6% pre-deformation were investigated. It can be concluded as follows: (1) For the Si-free alloy, the GPB zones and the zig-zag S phases are the main strengthening precipitates in the T6 peak-aged samples. After 6% pre-deformation followed by aging, a high density of the S precipitates with a uniform distribution become the main precipitates in the T8 peak-aged sample, which improve the hardness of the alloy. So the hardness of the T8-treated samples is higher than that of the T6-treated samples. (2) For the Si-containing alloy, the T6 peak-aged sample is mainly strengthened by Si-modified GPB zone. After 6% predeformation followed by aging, the homogenous
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precipitation of Si-modified GPB zones is suppressed, and the successive zig-zag composite precipitates become the main strengthening precipitates, leading to the degradation of performance eventually. So the hardness of the T8-treated samples is lower than that of T6-treated samples. (3) In T8-treated Si-containing alloys, the zig-zag composite precipitates include S phase, conventional GPB zone, GPB zone-variant, C phase, and the sub-units of the C/Qʹ phase. It is worth noting that a two-dimensional rectangular GPB zone-variant with a coherent habit-plane interface parallel to (110)Al was frequently observed, which reduced the energy of system and improved the stability of the zig-zag composite precipitates.
Author contributions section F.J. Niu: Conceptualization, Methodology, Validation, WritingOriginal Draft. C.L. Wu: Funding acquisition, Methodology, Writing-review and editing. J.H. Chen: Investigation, Project administration, Supervision. S.Y. Duan: Resources, Data curation. W.Q. Ming: Software, Visualization. J.B. Lu: Resources. Z. Le: Resources. Declaration of competing interest 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. Acknowledgements This work is supported by the National Natural Science Foundation of China (Nos. 11427806, 51831004, 51671082, 51471067) and the National Key Research and Development Program of China (No. 2016YFB0300801). References [1] J.C. Williams, E.A. Starke, Progress in structural materials for aerospace systems11The golden jubilee issuedselected topics in materials science and engineering: past, present and future, in: S. Suresh (Ed.), Acta Mater, 2003, pp. 5775e5799, 51(19). [2] S.B. Wang, J.H. Chen, M.J. Yin, Z.R. Liu, D.W. Yuan, J.Z. Liu, C.H. Liu, C.L. Wu, Double-atomic-wall-based dynamic precipitates of the early-stage S-phase in AlCuMg alloys, Acta Mater. 60 (19) (2012) 6573e6580. [3] S.C. Wang, M.J. Starink, Precipitates and intermetallic phases in precipitation hardening AleCueMge(Li) based alloys, Int. Mater. Rev. 50 (4) (2005) 193e215. [4] S.C. Wang, M.J. Starink, N. Gao, Precipitation hardening in AleCueMg alloys revisited, Scr. Mater. 54 (2) (2006) 287e291. [5] W.A. Perlitz H, Ark Kemi Mineral och Geol B 16 (1943) 13. [6] Z.R. Liu, J.H. Chen, S.B. Wang, D.W. Yuan, M.J. Yin, C.L. Wu, The structure and the properties of S-phase in AlCuMg alloys, Acta Mater. 59 (19) (2011) 7396e7405. [7] L. Kovarik, M.J. Mills, Structural relationship between one-dimensional crystals of GuinierePrestoneBagaryatsky zones in AleCueMg alloys, Scr. Mater. 64 (11) (2011) 999e1002. [8] L. Kovarik, M.J. Mills, Ab initio analysis of GuinierePrestoneBagaryatsky zone nucleation in AleCueMg alloys, Acta Mater. 60 (9) (2012) 3861e3872. [9] A. Charai, T. Walther, C. Alfonso, A.M. Zahra, C.Y. Zahra, Coexistence of clusters, GPB zones, S0-, S’- and s-phases in an Al±0.9% Cu±1.4% Mg alloy, Acta Mater. 48 (10) (2000) 2751e2764. [10] M.J. Yin, J.H. Chen, S.B. Wang, Z.R. Liu, L.M. Cha, S.Y. Duan, C.L. Wu, Anisotropic and temperature-dependent growth mechanism of S-phase precipitates in AleCueMg alloy in relation with GPB zones, Trans. Nonferrous Metals Soc. China 26 (1) (2016) 1e11. [11] S. Cheng, Y.H. Zhao, Y.T. Zhu, E. Ma, Optimizing the strength and ductility of fine structured 2024 Al alloy by nano-precipitation, Acta Mater. 55 (17) (2007)
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