Effect of Deformation on Mechanical Properties and Microstructure of Al-1%Cu-0.96%Mg-0.36%Si (mass%) Alloy

Abstract

In this study, we investigated the effect of a combination of deformation and aging after solution heat treatment of an Al-1%Cu-0.96%Mg-0.36%Si (mass%) alloy on the precipitation during subsequent artificial aging. The combination of microstructural studies and hardness tests allowed for a detailed characterization of the precipitation in the preformed material. The main phase in the deformed condition is the S′ phase. Aging time-hardness curves, DSC analysis, microstructural examination, and heat treatment optimization were performed to understand the precipitation process in more detail. The results showed that high hardness can be achieved in a short period of time. The aging process is also analyzed to compare the peak aging period and the hardness level of the material. We then compared the thermal aging process with three different cold-rolling conditions and concluded which condition has the most suitable properties for commercial purposes.

Fig. 7 Bright-field TEM image of three cold-rolling conditions which are pre-aging at 100°C, then cold rolling, artificial aging at 160°C peak aging (a), (b) 0%CR, (c), (d) 30%CR, (e), (f) 60%CR, (g), (h) 80%CR and SAED pattern corresponds to each condition in a horizontal row.

1. Introduction

Al-Cu-Mg-Si alloys are age-hardenable alloys, and depending on their composition, follow the precipitation order of Al-Cu-Mg (2XXX series) and Al-Mg-Si (6XXX) alloys. Thermomechanical processing is commonly used to produce high-strength products from these alloys. After solution treatment, quenching, pre-aging, cold working, and artificial aging are performed in sequence. For most wrought aluminum alloys, cold working by rolling or drawing will accelerate the precipitation process and strengthen [1]. Al-Cu-Mg(-Si) alloy belongs to the class of heat-treatable aluminum alloys developed for structural applications and used in the aerospace and other industries [2]. It is known that the aging temperature and alloy composition change the order of the precipitated phases, which affects the competitive relationship between the phases. When the aging temperature is below 180°C, the S′ phase is formed before the GPB zone occurs, and when the aging temperature is above 180°C, both the S′ phase and the GPB zone occur simultaneously [3]. It has the advantage of being able to achieve high strength. Two-stage aging is often carried out to enhance the precipitation hardening effect. The two-stage aging process uses the nucleation and formation of precipitates through low-temperature pre-aging. Generally, in these alloys, the precipitation process includes the following reactions:

S.S.S.S (Supersaturate solid solution) → Cluster/GPB zones → S′ phase → equilibrium S phase or S.S.S.S → Cluster/GP zones → θ′ phase → equilibrium θ phase.

Adding a small amount of Si improves the stability of the GPB zones and prevents the precipitation of the S′ phase. This also improves the heat resistance and the rapid aging response of the alloy [4, 5].

Most studies on early-stage precipitation of Al-Cu-Mg alloys to date have focused on alloys in the T4 (solution-treated and naturally aged) or T6 (solution-treated and artificially aged) states [6]. However, deformation after quenching, which is common in commercial applications, has a significant impact on the properties and precipitation of the alloy. Understanding the effects of deformation on the microstructure and mechanical properties of these alloys is important for optimizing their performance in different manufacturing processes. The initial microstructure of the Al-Cu-Mg-Si series consists of a solid solution matrix with dispersed reinforcing precipitates. Deformation changes the morphology, size, and distribution of these precipitates, which affect the strength, hardness, and other mechanical properties such as ductility and fatigue resistance. Several detailed studies on the precipitation behavior of these deformed and aged materials have been reported [6, 7, 15, 16]. In this paper, we combine hardness measurement (HV), differential scanning calorimetry (DSC), and transmission electron microscopy (TEM) techniques to analyze quenched, pre-aging (PA), cold-rolled (CR), and then artificially aged (AA) to provide comprehensive insight into the precipitation behavior of Al-Cu-Mg-Si alloy. A previous study [24] demonstrated that natural aging (NA) and 80% cold-rolled change the types of precipitates formed during artificial aging. Specifically, a single aging process produces two types of precipitates: the L phase and the GPB zones. In contrast, the aging treatment after deformation produces four types of precipitates: S′, C, E, and L. This study is a continuation of the previous study and has two main objectives. First and foremost, we investigated the effects of varying the levels of deformation from 0%, 30%, 60% to 80%, and investigated the types of precipitates present in both undeformed and deformed samples. Next, in this study, we observed the evolution of the S′ phase and the GPB zones during deformation. This may provide insight into the mechanism that generates the precipitation and dislocation behavior in the Al-Cu-Mg-Si alloy subjected to this deformation process.

2. Experimental Procedure

The chemical composition of the alloy used was Al-1.0%Cu-0.96%Mg-0.36%Si (mass%) supplied by YKK corporation, and the wire had a length of approximately ∅10 mm × 1 m. The heat treatment diagram is shown in Fig. 1. The wire was cut into 10 mm rod, solution treated at 505°C for 3 h, and then quenched in ice-water (0°C). It was then pre-aged for 1 week at 100°C. After this heat treatment, the rod was cold rolled at room temperature with 30, 60, and 80% in reduction to form plate shape with a thickness of 0.2 mm. Finally, the cold-worked plates were aged at 160°C for different times in an oil furnace. Hardness was measured using a Mitutoyo HM-101 by applying a load of 0.98 N and holding for 15 s. The hardness test typically uses an average of 10 data from 12 points. The samples for TEM observation were thinned by mechanical polishing to a thickness of 0.2 mm to 0.08 mm and then electropolished using a twin-jet method. The temperature of the electrolyte containing a mixture of 1/3 nitric acid (HNO3) and 2/3 methanol (CH3OH) was maintained between −20°C and −30°C during polishing. TEM observations of the samples were performed using a transmission electron microscope (TOPCON EM-002B). All observations were made in the ⟨110⟩ and ⟨100⟩ matrix directions. DSC runs were performed using a Perkin-Elmer apparatus at a constant heating rate of 10°C/min and a temperature range of 25°C to 500°C.

Fig. 1

Thermomechanical treatment procedures diagram used for Al-1.0%Cu-0.96%Mg-0.36%Si (mass%).

3. Results and Discussions

3.1 Mechanical properties

Figure 2 shows the hardness measurements of the cold-rolled alloy during aging. With 0%, 30%, 60% and 80% cold rolling, the hardness increased from 85 HV (after deformation) to 129, 135 and 139 HV, respectively. All cold-rolled (CR) samples had significantly higher hardness than the 0%CR sample. The hardness of the cold-rolled samples increased gradually to the peak aging and then decreased. In addition, the highest aging was achieved most quickly in the 80%CR sample. It can be concluded that cold rolling caused deformation before aging, which affected the hardening properties of the alloy during subsequent artificial aging, resulting in a significant decrease in hardness when subjected to over-aging. However, the hardness after deformation does not change much at 30% CR, 60% CR and 80% CR, so it is called asR. This is because the hardness of this alloy is affected by the deformation mechanism that occurs during cold rolling. During cold rolling, the alloy is plastically deformed, causing dislocation movement and rearrangement. At low deformation levels, such as 30% CR, the mobility of the dislocations increases, and the hardness increases significantly. On the other hand, when the deformation increases to 60% CR or 80% CR, the dislocations become more entangled and form a complex network [39]. As a result, the mobility is limited, and the hardness is saturated at this state. From Fig. 2, it can be seen that 30%CR reaches its maximum hardness after about 1000 min, and both 60%CR and 80%CR reach their maximum hardness after about 64 min. The undeformed 0%CR sample does not reach its maximum hardness after 10k min, which is the maximum time studied. This suggests that the presence of dislocations accelerates the precipitation rate and reduces the artificial aging time required to reach the peak aging condition. The decrease in the aging response of the undeformed samples is attributed to the lack of vacancies at the dislocation sites [8, 9]. On the other hand, the hardness curves show that the age-hardness of the cold-rolled samples when rolled ≥30% occurs in a different way, with a slight hardening followed by softening after the peak aging. Increasing deformation levels then reduces the time to reach the maximum hardness and the higher deformation level after rolling tends to increase the hardness, followed by the aging treatment [6]. After deformation, in the subsequent artificial aging treatment, the formation of GPB zones is inhibited and under the peak aging condition, almost all the precipitation is attributed to the S′ phase and the density of heterogeneous nucleation sites caused by dislocations is greatly reduced, so that they are uniformly distributed in the matrix [7].

Fig. 2

Age hardening curves demonstration evolution of Al-1.0%Cu-0.96%Mg-0.36%Si (mass%), alloy Vickers microhardness of pre-aging at 100°C, cold-rolled, and then artificial aging at 160°C.

3.2 DSC measurements

Figure 3 shows DSC curves for this alloy under four conditions. All heat treatment conditions for DSC are listed in Table 1. Analysis of the traces obtained at the lowest quenching temperature of 505°C can be used as a starting point for the discussion of the results. The thermal effects marked A, B and C can be easily attributed to cluster formation, GPB zones formation and formation of the main precipitated S′ phase [27, 28, 38]. Peak C shows a DSC trace with very typical characteristics of the large peak effect shown in Fig. 2 of the paper by S.P. Ringer et al. [10]. They also assert that their results “indicate the formation of GPB regions at all temperatures below 200°C and even with short aging times”. Therefore, the additional D, E, and F peaks that appeared with increasing solution temperature may be due to the precipitation of Q′, θ′, and β′ phases, respectively [25, 26, 38]. Based on the general DSC/TEM analysis, the exothermic effects observed in this temperature range may be related to the S′ and θ′ phases, although it is not possible to precisely assign each peak to a specific phase [11]. The S′ phase was the main peak in the cold-rolled samples under all conditions. This peak temperature decreased with increasing deformation levels. This was attributed to the increase in the number of heterogeneous S′ phase nucleation sites (mainly dislocations) [12]. The peak temperature shifted to lower values with increasing deformation levels, indicating that dislocations effectively promoted precipitation. The activation energy associated with the peak was calculated using the Kissinger equation [13]:

  

\begin{equation} \ln \frac{\alpha}{T_{p}^{2}} = - \frac{\text{Q}}{\text{R}T_{p}} + C \end{equation} (1)

where α is the heating rate, T_p is the peak-to-peak temperature, Q is the activation energy, and R is the universal gas constant [13]. The activation energy can be obtained from the slope of the fitting curve by using eq. (1) shown in Table 3 and Fig. 4.

Fig. 3

DSC thermogram at 10°C/min of Al-1.0%Cu-0.96%Mg-0.36%Si (mass%) alloy in S.T and S.T + Cold rolling conditions.

Table 1 Heat treatment conditions for DSC sample.

Fig. 4

Kissinger plots between 1/(RT_p) and $\ln (\alpha /T_{p}^{2})$ of each peak from DSC curve under four conditions. (online color)

Table 2 shows the activation energies of the S′ and θ′ phases calculated by DSC using the same method as in the literature [29–36]. The activation energies reported in the studies cited in Table 2 are comparable to those of the S′ and θ′ phases in this study. The activation energies are very similar. According to the report [21], Cu can enter the cluster and occupy part of the Mg atoms, making the cluster composition very close to that of the subsequent precipitate. The occupation of Cu in the precipitate due to the difference in cluster size can also reduce the activation energy and facilitate the growth of the precipitate. In addition, Cu atoms also serve as nucleation sites for the formation of other reinforcing phases. This promotes the development of a fine and uniform microstructure, while further reducing the activation energy required for dislocation migration and deformation. In general, the Cu effect in Al-Cu-Mg-Si alloys is attributed to the presence of Cu element, which promotes diffusion, nucleation, and microstructural refinement, reducing the activation energy [40].

Table 2 The activation energy for S′, θ′ phase formation in different alloys by using DSC measurement.

Table 3 shows the activation energies of all peaks of this alloy. The activation energies of these peaks for the cold rolled samples were lower than those for the non-cold rolled samples, indicating that as the deformation levels increases, the activation energy can be reduced by the presence of a large number of dislocations. Therefore, dislocations, which are crystal defects characterized by high activation energy, facilitate the existence of solute atoms and effectively reduce the overall activation energy [13]. In general, dislocation caused by material deformation can significantly affect various properties, including the diffusion activation energy [22]. For Al-Cu-Mg-Si alloys, dislocations can affect the activation energy associated with the diffusion of different types of atoms in the material. By acting as preferential migration routes, dislocations facilitate the diffusion of atoms [23]. If an atom is located near a dislocation line, the energy barrier to move along the dislocation is lower than that of moving in the crystal lattice. This can reduce the activation energy for atomic diffusion along the dislocation line. In Al-Cu-Mg-Si alloys commonly used in structural applications, the presence of dislocations can reduce the activation energy for the diffusion of alloying elements such as Cu, Mg. This reduction in activation energy can increase the diffusivity of these elements in the material, resulting in changes in microstructure and mechanical properties [40]. It is important to note that the exact details of the effect of dislocations on the activation energy of Al-Cu-Mg-Si alloys can vary depending on the specific composition and material processing.

Table 3 The activation energy of all peaks (kJ/mol) for this alloy under different conditions in this research.

3.3 TEM observation

3.3.1 Microstructure of single aging

Figure 5 shows TEM images of single aging condition at 35°C, 100°C, and 160°C. Clusters/GPB zones are seen after single aging at 35°C and 100°C, while the precipitates are more densely at 100°C. As a result, the single aged at 100°C precipitates are referred to as small rod-shaped GPB zones along the ⟨100⟩Al. According to E. Thronsen et al. [24, 37], the image (Fig. 5(c)) appears to show that the L phase is formed at 160°C, since the structural components of the L phase and the GPB zones are mixed in the undeformed sample. Due to their extreme fineness, these fine precipitates lack a clearly defined shape, which suggests that they are completely connected with the matrix and lack any distinguishable structure. Clusters/GPB zones dissolve and act as the core of θ′ and S′ phases as the artificial aging. In Fig. 5(d), diffraction spots can be seen in the corresponding SAED pattern, which is attributed to the monolayer or few layers of clusters reported by L. Kovarik et al. [14]. Based on the above results, it can be concluded that the type of precipitation of single aging depends on both the aging temperature and time. If the temperature is lower and aged for a longer time, it mainly consists of the GPB zones, while if the temperature is higher and aged for a shorter time, it mainly consists of the S′ phase [3].

Fig. 5

Bright-field TEM images of alloys with various single aging temperature: (a) 35°C for 1 week, (b) 100°C for 1 week, (c) 160°C for 1 week, and (d) the corresponding SAED pattern recorded along ⟨001⟩Al of (a) and illustration of a SAED pattern. (online color)

3.3.2 Microstructure of deformed sample

Figures 6(a), (b), (c) shows the bright-field TEM images of three conditions, immediately after cold rolling with increasing cold rolling reductions of 30%CR, 60%CR, and 80%CR, respectively. In each condition, the presence of dislocations becomes increasingly pronounced as increasing the deformation levels. The dislocation is visible but relatively dispersed within Al matrix, reflecting the moderate level of deformation introduced. In Fig. 6(d), a TEM image under the 80%CR condition provides a high magnification on the microstructure, showing no detectable presence of the S′ phase. As dislocations dominate the microstructure, this observation aligns with this study suggesting that applying deformation may promote the formation nucleation sites of S′ phase after artificial aging at 160°C [24]. To further examine the dislocation distribution, Fig. 6(e) shows an inverse FFT image from yellow square in Fig. 6(d) which the higher contrast regions correspond to the (110)Al while lower contrast regions reveal the locations of dislocation within Al matrix.

Fig. 6

Bright-field TEM images of three conditions, immediately after cold rolling: (a) 30%CR, (b) 60%CR, (c) 80%CR, (d) TEM image from 80%CR with high magnification and (e) Inverse FFT image from yellow square in (d). (online color)

Figure 7 shows TEM images of pre-aged samples for 1 week at 100°C, then cold rolled at four conditions: 0%CR, 30%CR, 60%CR, 80%CR and artificial aging at 160°C for the peak aging. The images show typical precipitation for four conditions and SAED pattern corresponds to each condition in a horizontal row. In contrast, from Figs. 7(c), (e), (g), a large number of dislocations were generated by CR, showing typical bright field TEM micrographs of the rolled sample, while Fig. 7(a) no dislocations is visible. The presence of dislocations after deformation depends on the initial hardness on the hardness curve in Fig. 2. The increase in hardness is a result of work hardening [15].

Fig. 7

Bright-field TEM image of three cold-rolling conditions which are pre-aging at 100°C, then cold rolling, artificial aging at 160°C peak aging (a), (b) 0%CR, (c), (d) 30%CR, (e), (f) 60%CR, (g), (h) 80%CR and SAED pattern corresponds to each condition in a horizontal row.

As shown in these figures, the main strengthening phase of the alloy used in this study is S′ (Al2CuMg). The precipitation order of this alloy is SSSS → clusters/GPB zone → S′ phase → equilibrium S phase. The S′ phase is said to precipitate preferentially when dislocations occur [24]. In Figs. 7(b), (d), (f), (h), when the sample is tilted near the Al ⟨110⟩ zone axis, nanometer-sized precipitates can be clearly observed. However, although the S′ phase precipitation formation occurs heterogeneously at the dislocation sites, with more aging time, it may tend towards a uniform precipitate distribution in the matrix. This is due to the after initial nucleation at dislocations, the growth of precipitation becomes diffusion-controlled. The atoms required for S′ phase formation diffuse from matrix toward precipitates, allowing them to grow uniformly across around dislocations [43]. During aging, the alloy may undergo recovery process, where dislocations rearrange or annihilate, reducing dislocation density. This lowers the encouraging a more even distribution of precipitation [44]. The S′ phase is an important precipitate when determining the hardness of Al-Cu-Mg-Si alloys. This occurs due to aging and is a strengthening precipitate. The increase in hardness occurs because the S′ phase interacts with the dislocations and promotes solid solution strengthening. By using appropriate heat treatment conditions, the formation and distribution of the S′ phase can be controlled. Therefore, the S′ phase plays a major role in the hardness of this alloy and is therefore an important focus in alloy design and processing. Also, the formation of GPB zones phase occurs and Cu-Mg clusters act as precursors for this alloy.

Furthermore, the S′ phase is refined and reduced in size as shown in Fig. 8. Both of these are necessary for effective strengthening. However, there is a limit to how much dislocations and precipitation density contribute to hardness once the matrix reaches a certain saturation point. With increasing deformation, dislocations and precipitate density reach a level where further increases have a limited impact on hardness due to strain saturation, leading to similar hardness levels [45]. Also, cold rolling introduces a significant dislocation density that itself contributes to initial hardness. This dislocation network already provides a substantial baseline hardening effect as explained in “Mechanical properties” section. The relationship between hardness increments, time required to reach peak aging with average length of S′ phase and its number of densities as shown in Table 4. When increasing deformation levels from 0%CR, 30%CR, 60%CR, and 80%CR, the average length of the S′ phase decreases from 9.05 nm to 7.23 nm, 5.81 nm, and 5.04 nm, respectively and number densities of precipitation is increased, suggesting that the dislocations promote the formation of nucleation sites for the S′ phase. In addition, a significant part of the overall strengthening is accomplished by the increase in precipitate size associated with the development of well-defined deformation regions. It is important to note that strength anisotropy appears when dislocation is present, and deformation is applied in the preferred direction [16]. The deformation can promote the nucleation of the S′ phase by creating additional sites for precipitation. Dislocations and other defects generated during deformation can serve as nucleation sites for S′ phase particles and lead to an increase in precipitate density, as shown in Fig. 8. Deformation can also accelerate the diffusion of alloying elements such as Cu, Mg, and Si, which are required for the formation of the S′ phase [40]. The increased diffusion rate due to deformation increases the supply of solute atoms to potential nucleation sites and promotes precipitation. On the other hand, deformation can fragment the existing S′ phase. This fragmentation increases the number of nucleation sites and leads to the formation of finer precipitates. Deformation can cause lattice distortion and introduce lattice defects, which can lead to loss of coherency between the S′ phase and the Al matrix [24]. Loss of coherency can affect the growth and stability of precipitation behavior. Deformation can also alter the precipitation rate by changing the supersaturation levels of the alloying elements. Applying deformation increases the effective driving force for precipitation, increasing the precipitation rate and accelerating the formation of the S′ phase.

Fig. 8

Length distribution of S′ phase in the Figs. 7(b), (d), (f), (h), and its number of densities.

Table 4 Summary of hardness increments, time required to reach peak aging, average length of S′ phase and its number of densities.

Figure 9 shows the TEM image and corresponding FFT pattern of the rod-shaped precipitate. Figure 9(b) shows the presence of S′ phase in (210)Al // (100)S′. This type of precipitate belongs to the orthorhombic system with lattice parameters a = 4.04 nm, b = 9.23 nm, and c = 7.14 nm and has the orientation relationship of (001)S′ // (021)Al, [100]S′ // [100]Al. Our results are in good agreement with the results of A.K. Gupta et al. [19]. The atomic arrangement is also shown in the S′ precipitate model in Fig. 9(d) reported by Z.R. Liu et al. [41], where Cu and Mg atoms are represented by red and yellow circles, respectively. The present study did not reveal any systematic differences between precipitates that clearly nucleated with such defects and those in which the nucleation sites were less obvious.

Fig. 9

The S′ phase in 80%CR sample after peak aging observed along ⟨100⟩Al axis: (a) TEM image; (b) the corresponding FFT; (c) the corresponding SAED pattern and (d) modeling of atomic from (b). (online color)

The TEM image and associated FFT pattern of the small rod-shaped GPB region are shown in Fig. 10 with the long axis parallel to the [100]Al viewing direction. As a result, the FFT image shows a cross-section of the GPB zones, and each common bright spot represents an atomic column projected parallel to the viewing direction in Fig. 10(b). The atomic arrangement is illustrated in the GPB zones precipitation model reported by C. Pan et al. [42] and shown in Fig. 10(d). The column period is assumed to be the period of aluminum and the Cu and Mg atoms are represented by red and yellow circles, respectively. The central column of the central GPB zones in Fig. 10(b) can be Cu-rich, suggesting that it may be a growing GPB zones, in which Cu atoms are replaced by Mg and Al during artificial aging [20].

Fig. 10

The GPB zones in 80%CR sample after peak aging observed along ⟨100⟩Al axis: (a) TEM image; (b) the corresponding FFT, (c) the corresponding SAED pattern and (d) modeling of atomic from yellow square in (b). (online color)

The hardness curve of the cold-rolled samples continues to increase as deformation levels increases, as shown in Fig. 2. The fine precipitate, large density and small precipitate size are responsible for the increase in hardness. The interaction between the particles of the second phase and the dislocation moment determine the strength of the material. The strength of the cold-rolled samples is enhanced by the strengthening mechanism in the fine precipitate phase [17]. In addition, this increase is due to the dense S′ precipitation generated during the aging process in Fig. 7. It is clear that the contribution of the S′ precipitation to the strength is due to the effect of the high S′ precipitation not only compensating for the loss of strength due to growth but also further increasing the hardness and strengthening [18].

4. Conclusion

In this study, we investigated how deformation affects the microstructure and mechanical properties of Al-Cu-Mg-Si alloys. The main findings can be summarized as follows.

1. (1)    In the single sample, no dislocations are observed, and the number density of precipitates is lowest. It was confirmed that dislocations and the number of high-density precipitates appeared after cold rolling and artificial aging.
2. (2)    When cold rolling from 0% to 80%, the length of the S′ phase decreased with increasing density. When cold rolling with 0%, 30%, 60% and 80% in reduction, the microhardness increased from 85 HV to 129, 135 and 139 HV.
3. (3)    When a large number of dislocations are present, the activation energy of each peak formed in the DSC decreases.
4. (4)    The peak aged alloys exhibit a continuous distribution of S′ precipitates and GPB zones. In contrast, rapid aging of cold-rolled samples results in a uniform distribution of nano-sized S′ precipitates and improved properties.

Acknowledgments

We would like to express our sincere gratitude to the MEXT (Ministry of Education, Culture, Sports, Science and Technology) scholarship for providing fincancial support during Mr. Vu Ngoc Hai’s studies at University of Toyama. We are also thankful to YKK Corporation for supplying the alloy used in this research.

REFERENCES

• 1)  I. Polmear, D. StJohn, J.F. Nie and M. Qian: Light Alloys, (Metallurgy of the Light Metals, Butterworth-Heinemann, 2017) pp. 19–27.
• 2)   J.C.  Williams and  E.A.  Starke: Progress in structural materials for aerospace systems, Acta Mater. 51 (2003) 5775–5799. doi:10.1016/j.actamat.2003.08.023
• 3)   M.J.  Yin,  J.H.  Chen,  S.B.  Wang,  Z.R.  Liu,  L.M.  Cha,  S.Y.  Duan and  C.L.  Wu: Anisotropic and temperature-dependent growth mechanism of S-phase precipitates in Al-Cu-Mg alloy in relation with GPB zones, Trans. Nonferrous Met. Soc. China 26 (2016) 1–11. doi:10.1016/S1003-6326(16)64082-7
• 4)   C.R.  Hutchinson and  S.P.  Ringer: Precipitation processes in Al-Cu-Mg alloys microalloyed with Si, Metall. Mater. Trans. A 31 (2000) 2721–2733. doi:10.1007/BF02830331
• 5)   C.  Li,  G.  Sha,  J.  Xia,  Y.  Liu and  S.P.  Ringer: Si-induced precipitation modification and related age-hardening response of an Al-4Mg-1Cu-0.5Si alloy, Mater. Chem. Phys. 193 (2017) 421–426. doi:10.1016/j.matchemphys.2017.01.041
• 6)   I.S.  Zuiko and  R.  Kaibyshev: Ageing response of cold-rolled Al-Cu-Mg alloy, Mater. Sci. Eng. A 781 (2020) 139148. doi:10.1016/j.msea.2020.139148
• 7)   F.J.  Niu,  J.H.  Chen,  S.Y.  Duan,  W.Q.  Ming,  J.B.  Lu,  C.L.  Wu and  Z.  Le: The effect of pre-deformation on the precipitation behavior of AlCuMg(Si) alloys with low Cu/Mg ratios, J. Alloy. Compd. 823 (2020) 153831. doi:10.1016/j.jallcom.2020.153831
• 8)   T.  Masuda,  Y.  Takaki,  T.  Sakurai and  S.  Hirosawa: Combined effect of pre-straining and pre-aging on bake-hardening behavior of an Al-0.6 mass% Mg-1.0 mass% Si alloy, Mater. Trans. 51 (2010) 325–332. doi:10.2320/matertrans.L-M2009831
• 9)   Y.  Birol: Pre-straining to improve the bake hardening response of a twin-roll cast Al-Mg-Si alloy, Scr. Mater. 52 (2005) 169–173. doi:10.1016/j.scriptamat.2004.10.001
• 10)   S.P.  Ringer,  S.K.  Caraher and  I.J.  Polmear: Response to comments on cluster hardening in an aged Al-Cu-Mg alloy, Scr. Mater. 39 (1998) 1559–1567. doi:10.1016/S1359-6462(98)00364-9
• 11)   G.  Riontino,  S.  Abis and  C.  Bottero: A DSC Study of Precipitation in Al-Cu-Mg Alloys, Aluminum Alloys 2 (1998) 903–908.
• 12)   N.  Sen and  D.R.F.  West: Some factors influencing S precipitation in Al-Cu-Mg and Al-Cu-Mg-Ag alloys, J. Inst. Met. 97 (1969) 87.
• 13)   M.  Cui,  Y.H.  Jo,  S.H.  Kayani,  H.W.  Kim and  J.H.  Lee: Effects of Cu additions on the precipitation activation energy and mechanical properties of prestrained Al-Mg-Si alloys, J. Mater. Res. Technol. 20 (2022) 2629–2637. doi:10.1016/j.jmrt.2022.08.037
• 14)   L.  Kovarik,  P.I.  Gouma,  C.  Kisielowski,  S.A.  Court and  M.J.  Mills: A HRTEM study of metastable phase formation in Al-Mg-Cu alloys during artificial aging, Acta Mater. 52 (2004) 2509–2520. doi:10.1016/j.actamat.2004.01.041
• 15)   H.Z.  Li,  R.M.  Liu,  X.P.  Liang,  D.E.N.G.  Min,  H.J.  Liao and  L.  Huang: Effect of pre-deformation on microstructures and mechanical properties of high purity Al-Cu-Mg alloy, Trans. Nonferrous Met. Soc. China 26 (2016) 1482–1490. doi:10.1016/S1003-6326(16)64253-X
• 16)   V.  Kulitskiy,  S.  Malopheyev and  R.  Kaibyshev: Effect of extensive isothermal rolling on microstructure and mechanical properties of an Al-Mg-Sc alloy, Mater. Sci. Forum 794–796 (2014) 1187–1192. doi:10.4028/www.scientific.net/MSF.794-796.1187
• 17)   M.R.  Roshan,  M.  Mirzaei and  S.J.  Jahromi: Microstructural characteristics and tensile properties of nano-composite Al 2014/4 wt.% Al2O3 produced from machining chips, J. Alloy. Compd. 569 (2013) 111–117. doi:10.1016/j.jallcom.2013.03.182
• 18)   S.  Cheng,  Y.H.  Zhao,  Y.T.  Zhu and  E.  Ma: Optimizing the strength and ductility of fine structured 2024 Al alloy by nano-precipitation, Acta Mater. 55 (2007) 5822–5832. doi:10.1016/j.actamat.2007.06.043
• 19)   A.K.  Gupta,  P.  Gaunt and  M.C.  Chaturvedi: The crystallography and morphology of the S′-phase precipitate in an Al(CuMg) alloy, Philos. Mag. A 55 (1987) 375–387. doi:10.1080/01418618708209875
• 20)   X.  Chen,  C.D.  Marioara,  S.J.  Andersen,  J.  Friis,  A.  Lervik,  R.  Holmestad and  E.  Kobayashi: Precipitation processes and structural evolutions of various GPB zones and two types of S phases in a cold-rolled Al-Mg-Cu alloy, Mater. Des. 199 (2021) 109425. doi:10.1016/j.matdes.2020.109425
• 21)   M.W.  Zandbergen,  A.  Cerezo and  G.D.W.  Smith: Study of precipitation in Al-Mg-Si Alloys by atom probe tomography and influence of Cu additions, Acta Mater. 101 (2015) 149–158. doi:10.1016/j.actamat.2015.08.018
• 22)   J.C.  Guía-Tello,  C.G.  Garay-Reyes,  H.M.  Medrano-Prieto,  M.A.  Ruiz-Esparza-Rodriguez,  M.C.  Maldonado-Orozco,  G.  Rodriguez-Cabriales and  R.  Martínez-Sánchez: Effect of Plastic Deformation on the Activation Energy for Precipitation of S′ Phase in Al₂₀₂₄ Alloy, Microsc. Microanal. 26 (2020) 2888–2889. doi:10.1017/S1431927620023119
• 23)  A.P. David, E.E. Kenneth and Y.S. Mohamed: Phase Transformations in Metals and Alloys, 4th ed., (Boca Raton, 2021) pp. 265–275.
• 24)   E.  Thronsen,  H.  Mørkeseth,  C.D.  Marioara,  K.  Minakuchi,  T.  Katsumi,  K.  Marthinsen and  R.  Holmestad: The effect of small additions of Fe and heavy deformation on the precipitation in an Al-1.1 Mg-0.5 Cu-0.3 Si At. Pct alloy, Metall. Mater. Trans. A 53 (2022) 3296–3310. doi:10.1007/s11661-022-06744-9
• 25)   L.  He,  H.  Zhang and  J.  Cui: Effects of pre-ageing treatment on subsequent artificial ageing characteristics of an Al-1.01 Mg-0.68 Si-1.78 Cu alloy, J. Mater. Sci. Technol. 26 (2010) 141–145. doi:10.1016/S1005-0302(10)60023-0
• 26)   V.G.  Arigela,  N.R.  Palukuri,  D.  Singh,  S.K.  Kolli,  R.  Jayaganthan,  P.  Chekhonin and  W.  Skrotzki: Evolution of microstructure and mechanical properties in 2014 and 6063 similar and dissimilar aluminium alloy laminates produced by accumulative roll bonding, J. Alloy. Compd. 790 (2019) 917–927. doi:10.1016/j.jallcom.2019.03.231
• 27)   S.P.  Ringer,  S.K.  Caraher and  I.J.  Polmear: Response to comments on cluster hardening in an aged Al-Cu-Mg alloy, Scr. Mater. 39 (1998) 1559–1567. doi:10.1016/S1359-6462(98)00364-9
• 28)   V.A.  Esin,  L.  Briez,  M.  Sennour,  A.  Köster,  E.  Gratiot and  J.  Crépin: Precipitation-hardness map for Al-Cu-Mg alloy (AA2024-T3), J. Alloy. Compd. 854 (2021) 157164. doi:10.1016/j.jallcom.2020.157164
• 29)   P.  Bassani,  E.  Gariboldi and  G.  Vimercati: Calorimetric analyses on aged Al-4.4 Cu-0.5 Mg-0.9 Si-0.8 Mn alloy (AA2014 grade), J. Therm. Anal. Calorim. 87 (2007) 247–253. doi:10.1007/s10973-006-7836-3
• 30)   A.  Hayoune,  S.G.  Shabestari and  D.  Hamana: Study on the structural evolution during non-isothermal aging of an Al-Cu-Mg-Si alloy by means of thermal analysis, Trans. Indian Inst. Met. 69 (2016) 1529–1535. doi:10.1007/s12666-015-0726-2
• 31)   V.S.  Aigbodion and  S.B.  Hassan: The study of the precipitation process in Al-Cu-Mg/bagasse ash particulate composites, J. Alloy. Compd. 501 (2010) 104–109. doi:10.1016/j.jallcom.2010.04.052
• 32)   A.K.  Jena,  A.K.  Gupta and  M.C.  Chaturvedi: A differential scanning calorimetric investigation of precipitation kinetics in the Al-1.53 wt% Cu-0.79 wt% Mg alloy, Acta Metall. 37 (1989) 885–895. doi:10.1016/0001-6160(89)90015-1
• 33)   M.  Fatmi,  A.  Ouali,  A.  Djemli,  T.  Chihi,  M.A.  Ghebouli,  F.  Sahnoune and  B.  Barka: Effect of Mg contents on the mechanical proprieties and precipitation kinetics in Al-3.3 wt.% Cu alloy, Metall. Mater. Eng. 23 (2017) 335–345. doi:10.30544/244335
• 34)   S.K.  Ghosh: Influence of cold deformation on the aging behaviour of Al-Cu-Si-Mg Alloy, J. Mater. Sci. Technol. 27 (2011) 193–198. doi:10.1016/S1005-0302(11)60048-0
• 35)   J.  Wang,  J.  Xie,  Z.  Mao,  T.  Liang,  A.  Wang,  W.  Wang and  S.  Hao: Microstructure evolution and mechanical properties of the Al-Cu-Mg-Ag alloy during non-isothermal aging, J. Alloy. Compd. 942 (2023) 169031. doi:10.1016/j.jallcom.2023.169031
• 36)   P.  Heugue,  D.  Larouche,  F.  Breton,  R.  Martinez and  X.G.  Chen: Evaluation of the Growth Kinetics of θ′ and θ-Al2Cu precipitates in a binary Al-3.5Wt Pct Cu alloy, Metall. Mater. Trans. A 50 (2019) 3048–3060. doi:10.1007/s11661-019-05227-8
• 37)   E.  Thronsen,  C.D.  Marioara,  J.K.  Sunde,  K.  Minakuchi,  T.  Katsumi,  I.  Erga and  R.  Holmestad: The effect of heavy deformation on the precipitation in an Al-1.3Cu-1.0Mg-0.4 Si wt.% alloy, Mater. Des. 186 (2020) 108203. doi:10.1016/j.matdes.2019.108203
• 38)  G. Riontion and S. Abis: Aluminum alloys, Vol. 2.
• 39)   S.  Tsuchida,  K.  Tokuno,  M.  Mitsuhara,  R.  Fukuoka,  R.  Tsuboi,  J.  Miyamoto and  M.  Hagino: Dislocation cell structures formed inside dislocation channels of rapid-cooled and tensile-deformed aluminum single crystals, Mater. Trans. 62 (2021) 1109–1117. doi:10.2320/matertrans.MT-M2021021
• 40)   N.  Afify,  A.  Gaber,  A.  Am and  A.D.  Ym: Precipitation kinetics in supersaturated Al-2.0 at% Cu-1.0 at% Mg alloy, Mater. Trans. 51 (2010) 317–320. doi:10.2320/matertrans.MC200926
• 41)   Z.R.  Liu,  J.H.  Chen,  S.B.  Wang,  D.W.  Yuan,  M.J.  Yin and  C.L.  Wu: The structure and the properties of S-phase in AlCuMg alloys, Acta Mater. 59 (2011) 7396–7405. doi:10.1016/j.actamat.2011.08.009
• 42)   C.  Pan,  Y.  Yang,  S.  Wang,  Y.  Liu,  S.  Hu,  Z.  Wang and  P.  Shen: Atomistic building blocks of one-dimensional Guinier-Preston-Bagaryatsky zones in Al-Cu-Mg alloys, Mater. Des. 187 (2020) 108393. doi:10.1016/j.matdes.2019.108393
• 43)   D.E.  Coates: Diffusion-controlled precipitation growth in ternary systems I, Metall. Trans. 3 (1972) 1203–1212. doi:10.1007/BF02642453
• 44)   A.A.  Kohnert and  L.  Capolungo: The kinetics of static recovery by dislocation climb, npj Comput. Mater. 8 (2022) 104. doi:10.1038/s41524-022-00790-y
• 45)   Y.  Zheng,  N.  Guo,  B.  Tang,  B.  Su and  Q.  Zhou: Uncovering Dislocation-and-Precipitation-Induced Viscoplastic Damage in Al-Zn-Mg alloy, Materials 16 (2023) 3769. doi:10.3390/ma16103769