Development of High-Strength Al–Cu–Mg Alloy by Combined Application of High-Pressure Torsion and Aging Treatment

Abstract

Solution treated Al–Cu–Mg alloys with three different compositions were subjected to high-pressure torsion (HPT) for 1 to 50 turns, and then aged at 423 K. By conducting HPT process, the hardness of the three alloys significantly increased after 50 turns, and the average grain sizes were refined to 130–140 nm. After aging treatment, the hardness was further increased to 271 HV, 288 HV and 273 HV in the peak aging condition for the Al–4Cu–1.5Mg, Al–4Cu–3Mg and Al–5Cu–3Mg (in mass%) alloys, respectively. The contribution of several strengthening mechanisms was quantitatively evaluated in terms of grain boundary hardening, dislocation hardening, solid solution hardening and cluster/precipitation hardening. It is shown from the quantitative evaluation that simultaneous strengthening due to grain refinement and nanoscale precipitates is successfully achieved by combined application of HPT process and subsequent aging treatment.

Fig. 7 Hardness variation of the HPT-processed samples of (a) 2A, (b) 2B and (c) 2C alloys during aging at 423 K.

1. Introduction

Precipitation hardening is a useful mean for strengthening aluminum alloys, particularly Al–Cu–Mg (2xxx series), Al–Mg–Si (6xxx series) and Al–Zn–Mg–Cu (7xxx series) alloys,^1–3) in addition to the strengthening by grain refinement. In recent years, it is confirmed that significant grain refinement of metallic materials can be achieved by severe plastic deformation (SPD) such as HPT (high-pressure torsion), HPS (high-pressure sliding), ARB (accumulative roll bonding) and ECAP (equal-channel angular pressing).⁴⁾ Horita et al. reported that grain sizes of high-purity aluminum (99.99%) are reduced to ∼1 µm through ECAP, ARB, HPS and HPT processes.^5–8) Among them, HPT process was quite effective for grain refinement compared to other SPD processes,^9,10) because much greater strain can be imposed depending on the number of turns of HPT process.⁵⁾

Although SPD processes have a high potential for significant increase in strength due to the refined grain sizes, coarse precipitates are preferentially precipitated on grain boundaries during subsequent aging treatment, and thus the contribution of precipitation hardening is likely reduced in ultrafine-grained alloys.¹¹⁾ Hirosawa et al. proposed three strategies to achieve simultaneous strengthening due to grain refinement and nanoscale precipitates; i.e. lowering aging temperatures, utilization of microalloying elements and taking advantage of spinodal decomposition, and such effectiveness was demonstrated for Al–Mg–Si, Al–Ag, Al–Li–Cu–Mg and Al–Zn–Mg–Cu alloy systems.^12,13) Recently, Masuda et al. achieved ultra-high strengthening to a level as high as 1 GPa for the commercial A2024 alloy by combined application of HPT process and aging treatment. Their results showed that the segregation of Cu and Mg atoms onto grain boundaries plays an important role in the enhancement in strength of the ultrafine-grained and peak aged alloy.¹⁴⁾

In this study, Cu and Mg contents were increased from those of the commercial A2024 alloy, and HPT process and subsequent aging treatment were applied to develop new high-strength Al–Cu–Mg alloys. Then, the strengthening mechanisms were discussed by quantifying the contribution of grain boundary hardening, dislocation hardening, solid solution hardening and cluster/precipitation hardening, respectively.

2. Experimental Procedures

The chemical compositions of Al–Cu–Mg alloys used in this study are listed in Table 1. The 2B and 2C alloys were designed to the compositions of Al–4Cu–3Mg and Al–5Cu–3Mg (mass%) by increasing Cu and Mg contents of the 2A alloy with the same composition as the commercial A2024 Al–4Cu–1.5Mg alloy. 1 mm-thick sheet of each alloy was provided from research subcommittee of high-strength aluminum alloys in the Japan Institute of Light Metals.

Table 1 Chemical compositions of the investigated Al–Cu–Mg alloys (mass%).

Disk samples with 10 mm diameter were cut from the sheets and subjected to solution treatment at 778 K for 1.8 ks and then quenched into water (ST samples). HPT process was conducted for the ST samples at room temperature for four different turns of N = 1, 5, 10 and 50 under an applied pressure of 6 GPa with a rotation speed of 1 rpm. After the HPT process, the thickness of the HPT-processed samples was reduced to 0.8–0.85 mm, and thus equivalent strain ε_eq introduced by HPT process was estimated through   

\begin{equation} \varepsilon_{\textit{eq}} = 2\pi rN/\sqrt{3}t. \end{equation} (1)

Here, r is the distance from the disk center and t is the thickness of the HPT-processed samples.^5,14,15) From this formula, ε_eq in the edge regions of the HPT-processed samples are calculated to be 8, 44, 88, and 440 after N = 1, 5, 10, and 50 turns, respectively. Furthermore, some HPT-processed samples were then aged in oil baths at 373 and 423 K.

After ground by abrasive papers and polished to a mirror-like surface, the Vickers microhardness of the ST and HPT-processed samples was measured during aging by applying 500 gf for duration of 15 s using Matsuzawa MMT-Xi tester. The hardness measurements were made along 8 radial directions with an interval of 0.5 mm as shown in Fig. 1.

Fig. 1

Schematic illustration of disk sample for HPT process. The locations for hardness measurement and TEM observation are also indicated.

X-ray diffraction (XRD) analyses were conducted for the ST, HPT-processed (N = 50) and HPT(N = 50)-Peak aged samples using CuKα radiation at an accelerating voltage of 40 kV and a current of 45 mA with a scanning speed of 0.5°/min and a scanning step of 0.01°. The corresponding dislocation densities were also evaluated from their peak broadening using the Williamson-Hall method.^16,17) Four different full widths at half maximum obtained from fundamental peaks of (111)_α, (200)_α, (220)_α and (311)_α planes were used for the estimation after calibrating the instrument.¹⁸⁾

For transmission electron microscopy (TEM) observation, disks with 3 mm diameter were punched out from the ST-Peak aged, HPT-processed (N = 50) and HPT(N = 50)-Peak aged samples (Fig. 1), ground to a thickness of 0.15 mm, and subjected to a twin-jet electro-chemical polishing to obtain electron-transparent field of view with a solution of 20 HNO₃-80 CH₃OH (in vol%) at a temperature of 253 K under 9–12 V. The bright-field images, dark-field images and selected area electron diffraction (SAED) patterns were recorded using a JEM-2100F transmission electron microscope at an accelerating voltage of 200 kV. Average grain sizes of the HPT-processed (N = 50) and HPT(N = 50)-Peak aged samples were determined from dark-field images by measuring at least 100 different grains surrounded by high-angle grain boundaries (HAGBs).

3. Experimental Results

3.1 Hardness and microstructures after solution and aging treatments

Figure 2(a) shows the isothermal phase diagram of the Al–Cu–Mg alloy system at 778 K. It is expected from Fig. 2(a) that the 2A alloy almost consists of α-Al phase after solution treatment at 778 K, while the 2B and 2C alloys contain small amount of S (Al₂CuMg) phase in the α-Al matrix. The corresponding XRD profiles of the ST samples well supported such a prediction as illustrated in Fig. 2(b).

Fig. 2

(a) Isothermal phase diagram of the Al–Cu–Mg alloy system at 778 K. (b) XRD profiles of the ST samples.

Figure 3 plots the variation of Vickers microhardness during aging of the ST samples at 423 K. The three Al–Cu–Mg alloys exhibit significant age hardening, and the hardness reached 131 HV (2A), 137 HV (2B) and 138 HV (2C) in the peak aging condition from the ST levels of 79 HV (2A), 88 HV (2B) and 93 HV (2C). Thus, it was found that the increase in Cu and Mg contents from the conventional A2024 alloy leads to further strengthening not only just after quenching but also after peak aging.

Fig. 3

Hardness variations of the ST samples during aging at 423 K.

Figure 4 shows TEM bright-field images and the corresponding SAED patterns taken from a beam direction of ⟨001⟩_α for the ST-Peak aged samples. The lath-like precipitates are present in the α-Al matrix and identified to be S′ phase from $0 1 \bar{3}$, 013, $0 2 \bar{1}$ and 021 diffraction spots in the SAED patterns.^19–21) It is well known that major precipitates observable in the Al–Cu–Mg alloy system depend on the Cu/Mg ratios; i.e. S′ phase is preferentially formed when the Cu/Mg ratio is in the range of 1.5 to 4.²⁰⁾ The Cu/Mg rations of the three alloys used in this study are 2.64 (2A), 1.5 (2B) and 1.89 (2C), and thus the formation of S′ phase appears to be reasonable.

Fig. 4

TEM bright-field images (left) and selected area electron diffraction patterns (right) for the ST-Peak aged samples of (a) 2A, (b) 2B and (c) 2C alloys at 423 K.

3.2 Hardness and microstructures after HPT process

Table 2 shows Vickers microhardness variation of the three alloys with increasing the number of turns of HPT process. The hardness at a position of 3.5 mm from the disk center significantly increased even by 1 turn, and 248 HV (2A), 269 HV (2B) and 249 HV (2C) were eventually obtained after 50 turns, respectively. Such increased hardness of the HPT-processed samples is almost twice the hardness of the ST-Peak aged samples (i.e. 131 HV (2A), 137 HV (2B) and 138 HV (2C)), and the effect of HPT process was the largest in the 2B alloy among the three alloys.

Table 2 Vickers microhardness variation of the HPT-processed samples with increasing the number of turns of HPT process N.

Figure 5 shows TEM bright-field images, dark-field images and the corresponding SAED patterns of the HPT-processed samples after 50 turns. From the SAED patterns with a series of diffraction rings, fine-grained structures with HAGBs were found to well develop after HPT process with average grain sizes of 130 nm (2A), 140 nm (2B) and 140 nm (2C). According to the results of XRD analyses in Fig. 6(a), the intensities of S phase of the 2B and 2C alloys were drastically reduced from the ST levels (Fig. 2(b)), suggesting that S phase is forcibly dissolved into the α-Al matrix by HPT process as previously suggested by Feng et al.²²⁾ Furthermore, the corresponding dislocation densities of the HPT-processed (N = 50) samples were estimated to be 6.6 × 10¹⁵ m⁻² (2A), 6.6 × 10¹⁵ m⁻² (2B) and 3.8 × 10¹⁵ m⁻² (2C), respectively.

Fig. 5

TEM bright-field images (left), dark-field images (right) and selected area electron diffraction patterns (inset) for the HPT-processed (N = 50) samples of (a) 2A, (b) 2B and (c) 2C alloys.

Fig. 6

XRD profiles of (a) HPT-processed (N = 50) and (b) HPT(N = 50)-Peak aged samples.

3.3 Hardness and microstructures after HPT process and aging treatment

Figure 7 plots the variation of Vickers microhardness during aging at 423 K for the HPT-processed (N = 50) samples. It is clearly shown that the noticeable increase in hardness is achieved by combined application of HPT process and aging treatment, regardless of the chemical compositions of the three alloys. For example, the highest hardness of 271 HV (2A), 288 HV (2B) and 273 HV (2C) was obtained for the HPT(N = 50)-Peak aged samples, although the increment in hardness during aging is only 23 HV (2A), 19 HV (2B) and 24 HV (2C), respectively. Remember that the increment in hardness during aging of the ST-Peak aged samples are all about 50 HV, as illustrated in Fig. 3.

Fig. 7

Hardness variation of the HPT-processed samples of (a) 2A, (b) 2B and (c) 2C alloys during aging at 423 K.

Figure 8 shows TEM bright-field images, dark-field images and the corresponding SAED patterns for the HPT(N = 50)-Peak aged samples. The average grain sizes were measured as 160 nm (2A), 170 nm (2B) and 180 nm (2C) confirming that ultrafine grains are well maintained even after peak aging at 423 K. As for the formation of precipitates during aging, on the other hand, similar microstructures as the HPT-processed (N = 50) samples (Fig. 5) disabled us from identifying the strengthening phases, but volume fractions of S phase were found to be increased after aging from the results of XRD analyses for the HPT(N = 50)-Peak aged samples (Fig. 6(b)). Furthermore, the corresponding dislocation densities were decreased to 8.7 × 10¹⁴ m⁻² (2A), 1.7 × 10¹⁵ m⁻² (2B) and 1.4 × 10¹⁵ m⁻² (2C), from those of the HPT-processed (N = 50) samples; 6.6 × 10¹⁵ m⁻² (2A), 6.6 × 10¹⁵ m⁻² (2B) and 3.8 × 10¹⁵ m⁻² (2C), respectively.

Fig. 8

TEM bright-field images (left), dark-field images (right) and selected area electron diffraction patterns (inset) for the HPT(N = 50)-Peak aged samples of (a) 2A, (b) 2B and (c) 2C alloys at 423 K.

4. Discussion

4.1 Strengthening mechanism

In this study, the hardness of the three Al–Cu–Mg alloys successfully reached 271 HV (2A), 288 HV (2B) and 273 HV (2C) by combined application of HPT process for 50 turns and subsequent aging treatment at 423 K. Thus, the elucidation of the strengthening mechanisms becomes essential for developing new high-strength Al–Cu–Mg alloys. The contribution of several strengthening mechanisms was evaluated by quantifying grain boundary hardening (Δσ_gb), dislocation hardening (Δσ_dis), solid solution hardening (Δσ_ss), cluster hardening (Δσ_clust) and precipitation hardening (Δσ_pre), and expressed by the increase in Vickers microhardness ΔHV through ΔHV = 0.3Δσ.^23,24)

4.1.1 Grain boundary hardening

Higher density of grain boundaries introduced by HPT process can inhibit dislocation motion by preventing its propagation to neighboring grains through the Hall-Petch relation:^25,26)   

\begin{equation} \varDelta \sigma_{\textit{gb}} = \sigma - \sigma_{0} = \frac{k_{y}}{\sqrt{d}} \end{equation} (2)

where k_y is a constant depending on the alloy system. In this study, 0.08 MPa m^1/2 for an Al–4%Cu alloy was taken as k_y,²⁷⁾ and measured average grain sizes of 130 nm (2A), 140 nm (2B) and 140 nm (2C) or 160 nm (2A), 170 nm (2B) and 180 nm (2C) were used for d in the HPT-processed (N = 50) or HPT(N = 50)-Peak aged samples. Thus, ΔHV_gb was evaluated as 65 HV (2A), 63 HV (2B) and 63 HV (2C) for the HPT-processed (N = 50) samples, or 59 HV (2A), 57 HV (2B) and 55 HV (2C) for the HPT(N = 50)-Peak aged samples, respectively (see ΔHV_gb in Table 3).

Table 3 Estimated contribution of each strengthening mechanism for the HPT-processed (N = 50) and HPT(N = 50)-Peak aged samples of 2A, 2B and 2C alloys. The measured hardness is also listed for comparison.

4.1.2 Dislocation hardening

The contribution of dislocation hardening was evaluated by the Bailey-Hirsch equation;²⁸⁾   

\begin{equation} \varDelta \sigma_{\textit{dis}} = \alpha MGb\sqrt{\rho}. \end{equation} (3)

Here, α is a constant of 0.2 for aluminum,^29,30) M is Taylor factor of 3.06 for fcc alloys,^23,24) G and b are shear modulus (25.5 GPa) and Burgers vector (0.286 nm) of aluminum, respectively. The experimentally measured ρ of 6.6 × 10¹⁵ m⁻² (2A), 6.6 × 10¹⁵ m⁻² (2B) and 3.8 × 10¹⁵ m⁻² (2C) or 8.7 × 10¹⁴ m⁻² (2A), 1.7 × 10¹⁵ m⁻² (2B) and 1.4 × 10¹⁵ m⁻² (2C) were used, and thus ΔHV_dis was quantified as 107 HV (2A), 107 HV (2B) and 87 HV (2C) or 39 HV (2A), 53 HV (2B) and 50 HV (2C) for the HPT-processed (N = 50) or HPT(N = 50)-Peak aged samples, respectively (see ΔHV_dis in Table 3).

4.1.3 Solid solution hardening

Solid solution hardening occurs in the investigated Al–Cu–Mg alloys because Cu and Mg atoms are partially present in solid solution. In this study, the contributions of solid solution hardening by trace elements of Si, Mn, Fe, Cr, Zn, and Ti were ignored, and thus the two main elements of Cu and Mg were considered for the calculation. To evaluate solid solution hardening Δσ_ss, we used the following equations.^31–33)   

\begin{equation} \varDelta \sigma^{i}{}_{\textit{ss}} = \frac{0.9MGb}{L_{\textit{ss}}^{i}}\cos \left(\frac{\varphi_{\textit{ss}}^{i}}{2}\right)^{\frac{3}{2}}\left(1 - \cfrac{\cos \biggl(\cfrac{\varphi_{\textit{ss}}^{i}}{2}\biggr)^{5}}{6}\right) \end{equation} (4)

  

\begin{equation} L^{i}{}_{\textit{ss}} = \frac{3^{1/4}}{2\sqrt{C_{i}}}b \end{equation} (5)

Here, $L_{ss}^{i}$ and C_i are the equivalent square spacing and the concentration of solute i atom, and $\varphi_{ss}^{i}$ is dislocation breaking angle for shearing solute i atom (178° for i = Cu and 177° for i = Mg^32,33)). Then we applied the Pythagorean addition law to unify the contribution of solid solution hardening by Cu and Mg;^31–34)   

\begin{equation} \varDelta \sigma_{\textit{ss}} = \sqrt{\varDelta\sigma^{\textit{Cu}}{}_{\textit{ss}}{}^{2} + \varDelta \sigma^{\textit{Mg}}{}_{\textit{ss}}{}^{2}}. \end{equation} (6)

From the results of TEM observation and XRD analyses for the HPT-processed (N = 50) samples, only a small amount of S phase was found to exist in the 2B and 2C alloys. Thus, it was reasonably assumed that after HPT process, all the Cu and Mg atoms in the three alloys were dissolved into the α-Al matrix with C_Cu = 0.0167 and C_Mg = 0.0165 (2A), C_Cu = 0.0190 and C_Mg = 0.0331 (2B) or C_Cu = 0.0230; C_Mg = 0.0318 (2C). Thus, ΔHV_ss for the HPT-processed (N = 50) samples was estimated as 20 HV (2A), 27 HV (2B) and 27 HV (2C), respectively. As for ΔHV_ss of the HPT(N = 50)-Peak aged samples, on the other hand, smaller values of 15 HV (2A), 22 HV (2B) and 22 HV (2C) were approximated because S phase was newly precipitated during aging treatment, as suggested in Fig. 6 (see ΔHV_ss in Table 3).

4.1.4 Cluster hardening

Recent studies have shown that solute atom clusters formed even just after quenching play an important role in the strengthening of aluminum alloys.^14,31,32) In this study, the contribution of cluster hardening Δσ_clust was evaluated from strengthening mechanisms of the ST samples;^31,32)   

\begin{equation} \textit{HV}_{\textit{ST}} = \textit{HV}_{0} + \varDelta \textit{HV}_{\textit{ss}} + \varDelta \textit{HV}_{\textit{clust}}, \end{equation} (7)

where HV_ST is the hardness of the three alloys in the as-quenched (A.Q.) condition (i.e. 79 HV (2A), 88 HV (2B) and 93 HV (2C) in Fig. 3), and ΔHV_ss was calculated from predicted concentrations of dissolved Cu and Mg in Fig. 2(a). Remember that ΔHV_gb and ΔHV_dis were ignored in the ST samples because of much smaller dislocation densities and much larger grain sizes than those of the HPT-processed samples. Thus, ΔHV_clust of the HPT-processed (N = 50) samples was evaluated as 39 HV (2A), 41 HV (2B) and 46 HV (2C), respectively (see ΔHV_clust in Table 3). As for the contribution of cluster hardening in the HPT(N = 50)-Peak aged samples, on the other hand, ΔHV_clust was evaluated together with the contribution of precipitation hardening ΔHV_pre because of the indistinctive progress from clusters to strengthening phases during aging.

4.1.5 Precipitation hardening

In this study, the distribution of strengthening phases was not well detected in the HPT(N = 50)-Peak aged samples (Fig. 8), although the noticeable increase in hardness of 23 HV (2A), 19 HV (2B) and 24 HV (2C) was observed during aging (Fig. 7). For an ultrafine-grained alloy, however, the presence of only a few precipitate particles within a grain is reported to exert a significant strengthening effect^15,35) in accordance with the Orowan’s theory;³⁶⁾   

\begin{equation} \varDelta \tau_{\textit{pre}} = \frac{\textit{Gb}}{\lambda'}. \end{equation} (8)

Here, λ′ is the inter-particle spacing defined as λ′ = L − d_P, L is the distance between particles and d_P is the particle size. Thus, when two precipitate particles with d_P = 10 nm are supposed to form within each grain of the HPT(N = 50)-Peak aged samples, λ′ = 50 nm (2A), 52 nm (2B) and 55 nm (2C) are calculated using the measured average grain sizes of 160 nm (2A), 170 nm (2B) or 180 nm (2C), resulting in ΔHV_pre + ΔHV_clust = 146 HV (2A), 140 HV (2B) and 133 HV (2C), respectively (see ΔHV_pre in Table 3).

4.1.6 Summation of contribution of each strengthening mechanism

Table 3 summarizes the evaluated contribution of each strengthening mechanism for the HPT-processed (N = 50) and HPT(N = 50)-Peak aged samples, and the strengthening mechanisms of the two samples are interpreted as follows. (i) The contribution of grain boundary hardening ΔHV_gb is little changed after aging because ultrafine grains are well maintained even in the HPT(N = 50)-Peak aged samples. (ii) The contribution of dislocation hardening ΔHV_dis is significantly decreased after aging, but the 2B and 2C alloys with increased Cu and Mg contents exhibit smaller decreases in ΔHV_dis due to the prevented annihilation of dislocations. (iii) Precipitation hardening gives a dominant contribution to the peak hardness of the HPT(N = 50)-Peak aged samples, while the contribution of cluster hardening is quantified about 40 HV in the HPT-processed (N = 50) samples.

The summation of all the contribution of strengthening mechanisms;³⁷⁾   

\begin{align} \textit{HV}_{\textit{total}} &= \textit{HV}_{0} + \varDelta \textit{HV}_{\textit{gb}} + \varDelta \textit{HV}_{\textit{dis}} + \varDelta \textit{HV}_{\textit{ss}} \\ &\quad + \varDelta \textit{HV}_{\textit{clust}} + \varDelta \textit{HV}_{\textit{pre}}, \end{align} (9)

is compared with the measured hardness HV_measured (Here, 20 HV was adopted as HV₀³⁸⁾ based on the lattice friction stress of aluminum) in Table 3. A slight difference between HV_Total and HV_Measured may arise from the contribution of precipitation hardening ΔHV_pre, where two precipitate particles are assumed to form within each grain although there are possibilities that the number of particles may be more or less in the actual microstructures. In addition, solute segregation on grain boundaries may be also responsible for the slight difference because of its strong contribution to the strengthening of ultrafine-grained and aged alloys.^14,39) However, it is well demonstrated to a great extent that the investigated Al–Cu–Mg alloys achieve simultaneous strengthening due to grain refinement and nanoscale precipitates by combined application of HPT process and subsequent aging treatment.

4.2 Effect of lowered aging temperature

As mentioned above, the strength of HPT-processed samples of the three Al–Cu–Mg alloys was increased after aging at 423 K (Fig. 7), but the hardness increment was lower than that of the ST samples (Fig. 3). In this study, to establish more suitable heat treatment process, aging temperature was lowered to 373 K as previously proposed by the authors,¹²⁾ and then a peak hardness of 293 HV was obtained together with a hardness increment of 31 HV for the 2B alloy. These values of hardness are higher than those of the HPT(N = 50)-Peak aged sample (i.e. 288 HV and 19 HV after aging at 423 K), and the corresponding average grain size and dislocation density (i.e. d = 170 nm and ρ = 1.7 × 10¹⁵ m⁻²) were found to be almost the same as those after aging at 423 K. Thus, the contribution of precipitation hardening could be increased by lowering aging temperatures, while the contributions of grain boundary hardening and dislocation hardening are remained.

5. Summary and Conclusions

1. (1)    The average grain sizes of the investigated Al–4Cu–1.5Mg, Al–4Cu–3Mg and Al–5Cu–3Mg (in mass%) alloys were well reduced to submicron levels of 130 nm, 140 nm and 140 nm after 50 turns of HPT process. These refined gains together with the increased dislocation densities attribute to the significantly increased hardness of 248 HV, 269 HV and 249 HV in the HPT-processed samples, respectively.
2. (2)    The hardness was further increased to 271 HV, 288 HV and 273 HV when aging treatment at 423 K was applied to the HPT-processed samples. Such strengthening is considered due to the formation of nanoscale precipitates of S phase within ultrafine grains.
3. (3)    The contributions of several strengthening mechanisms were quantitatively evaluated. After aging of the HPT-processed samples, grain boundary hardening was little changed, while dislocation hardening was significantly decreased, but precipitation hardening gave a dominant contribution to the increased hardness of the HPT-Peak aged samples.
4. (4)    It is demonstrated that lowering of aging temperatures is useful to further strengthen the Al–Cu–Mg alloys due to more simultaneous strengthening of grain refinement and nanoscale precipitates.
5. (5)    The effect of the increased Mg content from the A2024 Al–4Cu–1.5Mg is more obvious than that of the increased Cu content on the enhancement in strength. Thus, the Al–4Cu–3Mg alloy is recommended as a new high-strength Al–Cu–Mg alloy if subjected to HPT process and aging treatment.

Acknowledgements

The authors would like to acknowledge the Japan Institute of Light Metals, the Light Metal Educational Foundation, Inc. and the Japan Aluminium Association for their generous support to this study. One of the authors (TM) also acknowledges a Grant-in-Aid for Young Scientists from MEXT, Japan (JP21K14436).

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