# Investigation of the high thermal ageing resistance of the 2219 aluminium alloy

Thomas Perrin, Arthur Després, Pierre Heugue, Alexis Deschamps, Frédéric De Geuser

► To cite this version:

Thomas Perrin, Arthur Després, Pierre Heugue, Alexis Deschamps, Frédéric De Geuser. Investigation of the high thermal ageing resistance of the 2219 aluminium alloy. Materials & Design, 2025, 256, pp.114373. 10.1016/j.matdes.2025.114373. hal-05380114

HAL Id: hal-05380114

https://hal.science/hal-05380114v1

Submitted on 24 Nov 2025

HAL is a multi-disciplinary open access
archive for the deposit and dissemination of sci-
entific research documents, whether they are pub-
lished or not. The documents may come from
teaching and research institutions in France or
abroad, or from public or private research centers.

L'archive ouverte pluridisciplinaire HAL, est
destinée au dépôt et à la diffusion de documents
scientifiques de niveau recherche, publiés ou non,
émanant des établissements d'enseignement et de
recherche français ou étrangers, des laboratoires
publics ou privés.

<figure><img src="image_1.png" /></figure>

Distributed under a Creative Commons Attribution 4.0 International License
<figure><img src="image_2.png" /></figure>

# Investigation of the high thermal ageing resistance of the 2219 aluminium alloy

Thomas Perrin<sup>a,* ID</sup>, Arthur Després<sup>a</sup>, Pierre Heugue<sup>b ID</sup>, Alexis Deschamps<sup>a ID</sup>, Frédéric De Geuser<sup>a ID</sup>

## ARTICLE INFO

Keywords:

Aluminium alloys

High-throughput

Precipitation

Atom probe tomography

Interfacial segregation

## ABSTRACT

The 2219 aluminium alloy is widely used in high-temperature applications, particularly in aerospace. Thermal ageing during duty can lead to the coarsening of θ' hardening precipitates, which are responsible for the mechanical properties of the alloy. In this study, temperature gradient heat treatments were employed to evaluate its resistance to thermal ageing over a continuous temperature range of 165 °C to 245 °C for up to 10,000 h. Despite its composition being primarily that of an Al-Cu binary alloy, hardness mappings revealed a limited decrease in mechanical properties. Microstructural characterisation using high-throughput scattering techniques (SAXS & WAXS) showed that this stability is due to a remarkable resistance of the θ' precipitates to coarsening and limited θ' to θ transformation, as validated by a precipitate hardening model. This stability suggests that minor alloying elements play a crucial role in thermal resistance. Atom probe tomography (APT) analyses showed that this stability may arise from mechanisms similar to those observed in recently developed ACMZ alloys, as Mn segregations were observed at the interfaces of the θ' precipitates. However, unlike ACMZ alloys, at 300 °C, the stabilising mechanism is insufficient to preserve mechanical properties, suggesting a strong sensitivity of these mechanisms to compositional variations.

## 1. Introduction

The 2219 alloy is a widely used aluminium alloy known for its excellent weldability, corrosion resistance, and strong resistance to thermal ageing. These properties make it suitable for high-performance applications, such as the pressurised hulls of the Columbus and Tranquility modules on the International Space Station. As a result, extensive studies have been conducted on this alloy, covering aspects such as welding properties [1], the effects of quenching rates [2], work hardening behaviour [3], fracture characteristics [4] and optimisation of thermo-mechanical heat treatments [5,6]. However, very few studies have explored the evolution of mechanical properties during prolonged over-ageing for this alloy, and in these cases the evolution was only studied for one given temperature [7]. Understanding over-ageing behaviour is crucial to predict the lifespan for applications where the material must endure thermal ageing during service.

The 2219 alloy is primarily an aluminium-copper binary alloy with small additions of other elements, leading to the precipitation of phases

from the θ family. The precipitation sequence from the supersaturated solid solution (SSS) is given by:

$$SSS \rightarrow GP \ \text{zones} \rightarrow \theta^{\ddagger} \rightarrow \theta \rightarrow \theta$$

with GP zones standing for Guinier-Preston zones. In the peak-aged condition, only θ' and θ precipitates are present [8]. The θ' precipitates form as platelets in the (100)$_{Al}$ planes, resulting in three variants. These precipitates exhibit a body-centred tetragonal crystal structure with parameters a$_{θ'}$ = 4.04 Å and c$_{θ'}$ = 5.80 Å, and an Al$_{2}$Cu composition [9]. The orientation relationship between the θ' precipitates and the aluminium matrix is (001)$_{θ'}$//((001)$_{Al}$ and [100]$_{θ'}$// [100]$_{Al}$ [9,10]). The broad interface is coherent with the aluminium matrix, while the rim of the platelet is semi-coherent [11]. This contributes to the formation of precipitates with high aspect ratios, typically ranging from 10 to 100, enhancing their strengthening effect [8]. The crystallographic structure of the broad interface is locally different from the bulk structure, with a GP-like interface and copper segregations observed [10]. The θ precipitates are stable and incoherent precipitates.

<footnote><sup>a</sup> Univ. Grenoble Alpes, CNRS, Grenoble INP, SIMaP, F-38000 Grenoble, France</footnote>

<footnote><sup>b</sup> DT/MPE, Safran Transmission Systems, 18 Boulevard Louis Seguin, Colombes 92707, France</footnote>

<footnote>* Corresponding author.</footnote>

<footnote>E-mail address: thomas.perrin@grenoble-inp.pro (T. Perrin).</footnote>
Due to their typically large size and low aspect ratio, they have close to no hardening effect. They exhibit a globular morphology with a tetragonal structure and an Al₂Cu composition [12]. Delaying the θ' to θ transition is therefore essential to maintain the best precipitation hardening effect possible.

The addition of several slow-diffusing elements in low amounts has been shown to enhance the stability and coarsening resistance of the θ' precipitates, which delays the θ' to θ transition. Rosalie and Bourgeois [13] demonstrated that adding Ag to the Al-Cu system promotes the formation of Ag segregations on the coherent interfaces of the θ' precipitates, which inhibits their lateral growth. Similarly, Poplawsky et al. studied an Al-Cu-Mn-Zr (ACMZ) alloy using atom probe tomography (APT) and observed that the addition of Mn in the Al-Cu alloy results in Mn segregations initially forming at the semi-coherent interface, and later at the coherent interface of the θ' precipitates during high-temperature heat treatments at 300 °C. Alloys containing Mn exhibited greater resistance to thermal ageing at high temperatures (up to 300 °C), suggesting that these segregations limit the coarsening of the precipitates [14]. Further additions of slower diffusing elements, such as Ti or Zr, improved this resistance even more, attributed to the formation of Al₃(Zrₓ, Ti₁-x) L₁₂ phase at the interface between θ' precipitates and matrix, which reduce the interfacial energy and impede coarsening. However, adding only Zr and Ti proved ineffective in enhancing ageing resistance. This can be explained by the slow diffusion of these elements, preventing them from segregating at the θ' interface before the phase transition to the θ phase occurs. In the presence of Mn, the Mn segregations sufficiently slow down coarsening, allowing subsequent Zr and Ti segregations to further inhibit the process. As a consequence, the thickness of the θ' platelets within the ACMZ alloy was shown to remain below 10 nm even after 5,000 h at 300 °C [15]. As the 2219 alloy contains amount of Mn, Zr and Ti close the composition of the ACMZ alloy studied by Poplawsky et al., these phenomena could potentially take place during ageing.

The study of over-ageing in metallic alloys typically involves creating ageing matrices with various ageing durations at one or a few specific temperatures [16–20]. Although these studies provide valuable insights into the trends of over-ageing, the data obtained are limited, and direct conclusions about the evolution of mechanical properties under complex ageing conditions, such as those encountered in service, should be drawn with caution. Ageing heat treatments in a temperature gradient have been shown to be an effective approach for studying the over-ageing of aluminium alloys [21], as they produce graded samples that can be characterised using high-throughput techniques, giving access to quasi-continuous data as a function of ageing temperature. Diffraction techniques such as small-angle and wide-angle X-ray scattering (SAXS and WAXS) are effective for characterising the evolution of precipitation along graded samples. SAXS provides information on the size and volume fraction of nano-precipitates when appropriate models are applied, while WAXS offers insights similar to powder X-ray diffraction, with information about the nature of the phases in the sample.

In this study, we investigate the over-ageing of the 2219 alloy using such a high-throughput approach that involves ageing samples in a temperature gradient ranging from 165 °C to 245 °C for up to 10,000 h. The samples are characterised by hardness maps and simultaneous SAXS/WAXS acquisitions to investigate the correlations between microstructural and mechanical evolutions. Time-temperature equiva- lences were implemented to extrapolate the evolution of the mechanical properties for a given temperature. Given the observed high resistance to thermal ageing, additional analyses using APT were conducted to examine potential segregations at the interfaces of the θ' precipitates.

## 2. Material and methods

The studied alloy was produced as thick plates by Kaiser Aluminium with the composition of the alloy measured by inductively coupled

plasma atomic emission spectroscopy (ICP) given in Table 1. It was received in a T851 initial state, corresponding to a solution heat treatment at 535 °C followed by water quenching, 0.5–3 % cold stretching and artificial ageing 16 h to 20 h at 175 °C. All samples were extracted from the same depth within a single thick plate (1300 × 420 × 200 mm³) to limit possible composition fluctuations between samples.

Concerning the preliminary single-temperature ageing treatments (see section 3.1), TC25 cylindrical tensile test specimens were machined from each sample. They were tested following the NF EN 2002-001 standard. The strain rate was set at 8.3 × 10⁻⁵ s⁻¹ up to 1 % strain followed by a strain rate of 8.3 × 10⁻⁴ s⁻¹. Six samples were tested for each condition.

Vickers hardness measurements were conducted with a 500 g load. Thirty indents were measured for each condition from the single temperature heat treatments. For the graded samples, hardness profiles were conducted with a 5 mm step which correspond to a 0.5 °C to 1 °C increment depending on the position in the sample. Five indents were measured per position along the gradient to obtain averaged values.

The gradient temperature heat treatments were performed using a Carbolite-Gero TG3 12/60/600 tubular gradient furnace and a cooler, as described in [21]. The samples had dimensions of 325 × 7 × 3.5 mm³. For each heat treatment duration, a different sample was used. To achieve this, a sample support with 15 sample locations was placed inside the gradient furnace and simultaneously connected to a cooler to increase the temperature gradient. Instrumented samples (with thermocouples) were used to ensure that the temperature gradient was similar in the different sample locations. Multiple heat treatments durations can thus be performed simultaneously, with sacrificial samples taking place of aged samples to keep the thermal load constant. This configuration offers a stable temperature gradient from 165 °C to 245 °C with a maximum deviation of ± 2 °C at each position, even for heat treatments lasting up to 10,000 h.

Simultaneous SAXS/WAXS acquisitions were carried out at the D2AM-BM02 beamline of the European Synchrotron Radiation Facility (ESRF) in Grenoble. Measurements were taken along the graded samples with a step size ranging from 1 mm to 1.5 mm, depending on the sample. To enhance data reliability, especially considering the large grain size of the alloy and the anisotropic nature of the precipitates, sample rotations were implemented during acquisitions, as detailed in [22]. SAXS fits were confirmed to be reliable using this technique for the 2219 alloy under study, as a good correlation was obtained between SAXS and TEM/APT measurements. A semi-continuous approach was employed, with a 3°/s rotation and 1-second acquisition, resulting in each data point covering a 3° range. Azimuthal integration was performed using the PyFAI package from [23]. For each position in the graded sample, the acquisitions were averaged along the -45°/+45° range, considering appropriate normalisation for each 3° step. The resulting I(q) plots were then averaged across ten positions. A beam size of 100 × 100 µm² and an 18 keV energy were used. The SAXS detector (IMXPAD D5) was positioned 3.5 m from the sample, while the WAXS detector (IMXPAD WOS) was located 15 cm away.

A JEOL 2100F with a field gun operated at 200 kV was used for TEM observations in STEM mode. The samples were prepared by electro-polishing using a solution of 33 % nitric acid in methanol at -25 °C.

APT acquisitions were performed on a LEAP6000 XR from CAMECA in laser mode, with a specimen temperature of 60 K. The laser energy was set at 150 pJ and the pulse frequency at 125 kHz. The IVAS software was used for the reconstruction and analysis of the tip specimen. These specimens were prepared using a two-step electropolishing with a solution of 25 % perchloric acid in acetic acid and then a solution of 2 %

<table><caption>**Table 1**

ICP measurement of the composition of the alloy.</caption><thead><tr><th>wt%</th><th>Al</th><th>Cu</th><th>Mn</th><th>Zr</th><th>Fe</th><th>V</th><th>Ti</th><th>Si</th><th>Zn</th></tr></thead><tbody><tr><td></td><td>B.</td><td>6.4</td><td>0.30</td><td>0.11</td><td>0.08</td><td>0.06</td><td>0.04</td><td>0.04</td><td>0.03</td></tr></tbody></table>
perchloric in 2-butoxyethanol at room temperature.

### 3. Evolution of mechanical properties during ageing

Since the high-throughput methodology employed in this study relies on graded samples, it is not feasible to machine tensile specimens, limiting the measurements to hardness profiles along the samples. Consequently, classical single-temperature heat treatments were initially used to establish correlations between hardness and tensile properties and to investigate the evolution of tensile behaviour during ageing.

#### 3.1. Preliminary single-temperature heat treatments

Three temperatures (150 °C, 175 °C, and 200 °C) were investigated, with ageing durations ranging from 1,000 to 15,000 h. The evolution of a few tensile curves is shown in Fig. 1 (a). Similar trends were observed for ultimate tensile strength in true stress (UTS), yield strength ($\sigma_y$), and hardness (HV), as illustrated in Fig. 1 (b). For each temperature, a rapid evolution is observed during the first 1,000 h of ageing, while further ageing impacts to a smaller extent the mechanical properties. Strong linear correlations were found between these properties, as demonstrated in Fig. 1 (c). Therefore, equivalent tensile properties can be computed based on hardness for the graded samples using these correlations.

#### 3.2. Temperature gradient heat treatments

A temperature gradient ranging from 165 °C to 245 °C was used to characterise the evolution of the mechanical properties during ageing. Hardness profiles were measured to assess this evolution over six different durations, ranging from 120 h to 10,000 h (Fig. 2), with one

<figure><img src="image_3.png" /><figcaption>Fig. 2. Hardness profiles and equivalent yield strength ($\sigma_y$) for the samples aged in the temperature gradient compared to the initial T851 state.</figcaption></figure>

<figure><img src="image_4.png" /><figcaption>Fig. 1. Results of the single-temperature heat treatments with (a) examples of true stress – true strain curves for samples aged at 200 °C, (b) evolution of the ultimate tensile strength (UTS), yield strength ($\sigma_y$) and hardness (HV) during ageing, and (c) linear correlations between these parameters.</figcaption></figure>
sample aged per duration. A rapid evolution of mechanical properties is observed within the first 120 h of ageing, followed by a stabilisation. For samples aged for 1,000 h or longer, the resulting hardness shows minimal dependency on both the ageing duration and temperature within the investigated range. A plateau is therefore observed and resulting hardness of 100 HV is obtained for the most extreme ageing condition (corresponding to a 200 MPa yield strength).

A time-temperature equivalence approach was applied to plot the experimental data as a function of an equivalent ageing time at a reference temperature. This equivalence assumes that the ageing process is a phenomenon governed by a single thermally activated mechanism. The equivalent ageing time ($t_{eq}$) for a chosen ageing temperature ($T_{ref}$) is given by relation (1):

$$t_{eq} = \exp\left(\frac{E_a}{R}\left(\frac{1}{T_{ref}} - \frac{1}{T}\right)\right) \quad (1)$$

where t and T represent the actual ageing duration and temperature, R is the gas constant and $E_a$ is the activation energy of the thermally activated mechanism.

The activation energy can be fitted to create a single master curve from the various experimental data. A good correlation is achieved when using the activation energy of copper diffusion, which is 133 kJ/mol according to ThermoCalc. Doing so, Fig. 3 is obtained where hardness results from the different graded samples are plotted as a function of the equivalent ageing time at 200 °C, calculated using relation (1) for each point. A plateau is reached after 10,000 h at 200 °C. An additional sample was aged for 1,000 h at 300 °C in a conventional furnace, and its results are also presented in Fig. 3. The deviation between this result and the plateau observed for the graded samples indicates that the time-temperature equivalence does not hold for ageing temperatures beyond the studied range. This suggests that another mechanism, such as phase transition from θ' to θ, or a significant change in the equilibrium volume fraction of θ' precipitates, becomes dominant at higher temperatures, invalidating the primary assumption of the equivalence. Equilibrium predictions using ThermoCalc software indeed indicate a significant decrease in the equilibrium volume fraction of the θ' precipitates in the present alloy when comparing values at 240 °C and 300 °C, dropping from 6.8 % to 5.9 %. However, within the 165 °C to 245 °C temperature range, ThermoCalc predicts only minor variations in equilibrium fraction, from 7.0 % to 6.8 %.

## 4. Evolution of precipitates during overageing

To evaluate the evolution of the hardening precipitation in regards with the stabilisation of the mechanical properties, simultaneous SAXS and WAXS were used to perform high-throughput characterisation. Additional TEM observations were conducted across various conditions to further analyse the evolution of the morphology of the nano-

precipitates.

Exploiting SAXS/WAXS acquisitions for anisotropic precipitates such as θ' platelets presents challenges due to the anisotropic nature of the signal, especially in materials with large grain sizes such as our alloy. To overcome this, sample rotations were implemented as discussed in the methods section, as they were shown to improve the reliability of the results for these precipitates.

To identify the precipitates within the alloy and qualitatively monitor the evolution of phases, WAXS was employed. For ease of comparison with traditional X-ray diffraction, the scattered intensity is plotted as a function of an equivalent double diffraction angle for a copper anode (2θCu). Peaks corresponding to the θ' and θ precipitates were identified (Fig. 4 (a)), showing a clear evolution between the initial T851 state and the condition aged for 1,000 h at 200 °C. The strongest peaks of these precipitates—specifically, the (110) peak for both phases—were fitted to estimate changes in peak area and full width at half maximum (FWHM). The (110) peak of the θ phase was fitted with a Gaussian function, while the (110) peak of the θ' phase was fitted with a Lorentzian function, as this provided a better fit. The peak area is known to be proportional to the volume fraction of the phase, while the FWHM is inversely proportional to the mean size of the phase. The evolution of the area and FWHM of the fitted peaks were measured and are plotted for each phase in Fig. 4 (b), (c), (d) and (e).

A significant increase in the peak area of the (110) peak for the θ phase is observed after 1,000 h, indicating an increase in the volume fraction of this phase (Fig. 4 (b)). No substantial changes are noted between 1,000 and 5,000 h of ageing, and varying the ageing temperature within the considered range does not have a clear impact. The FWHM of the θ phase does not exhibit notable changes during ageing, regardless of temperature, suggesting that either the average precipitate size remains constant or the upper limit of the WAXS setup for size measurement has been reached (Fig. 4 (c)).

For the θ' phase, a similar trend is observed in terms of peak area (Fig. 4 (d)), with a rapid increase during the first 1,000 h of ageing followed by stabilisation. However, the FWHM of the θ' peak shows a distinct evolution (Fig. 4 (e)). A significant reduction in the FWHM is observed during the first 1,000 h (as also visible in Fig. 4 (a)), with temperature having an impact on this characteristic as higher temperatures result in an even lower FWHM. For temperatures below 200 °C, extending the ageing duration further decreases the FWHM, while for higher temperatures, the effect of ageing duration is minimal, reaching a plateau. This suggests that, during the first 1,000 h of ageing, the θ' precipitates exhibit both an increase in volume fraction and a significant growth in size. For longer ageing times, the volume fraction remains relatively constant, regardless of the ageing temperature, while higher temperatures result in larger precipitates for ageing temperatures below 200 °C. For ageing at temperatures above 200 °C, the size of the precipitates reaches a plateau after 1,000 h and does not change with

<figure><img src="image_5.png" /><figcaption>Fig. 3. Time temperature equivalence of the evolution of hardness during ageing at 200 °C using an activation energy of 133 kJ/mol.</figcaption></figure>
<figure><img src="image_6.png" /><figcaption>Fig. 4. (a) Examples of the WAXS I(2θ<sub>Cu</sub>) plots obtained for the T851 initial state and the condition aged 1,000 h at 200 °C with the strongest peaks of the θ and θ' phases identified. (b) and (c) represent the evolution of the peak area and FWHM for the θ peak while (d) and (e) represent the evolution of the peak area and FWHM for the θ' peak. The standard deviation of the results from 10 different positions from the sample in the T851 initial state is indicated by a grey box.</figcaption></figure>

<figure><img src="image_7.png" /><figcaption>Fig. 5. SAXS characterisation of the evolution of the microstructure of the alloy during ageing with (a) examples of I(q) plots for three ageing conditions at 200 °C compared to the T851 initial state, and fit results for (b) the mean thickness and (c) volume fraction of the θ' platelets. The standard deviation of the SAXS fit for 10 different positions from the T851 sample are represented by a grey box in (b) and (c).</figcaption></figure>
further ageing. Increasing the temperature up to 245 °C does not affect the size of the precipitates beyond this plateau. The other peaks identified on the WAXS plots are related to non-hardening precipitates and do not evolve during ageing, regardless of ageing temperature and duration.

More quantitative results regarding the evolution of the θ' precipitates can be obtained using SAXS and fitting the data with an appropriate model considering the platelet morphology of the precipitates.

Examples of the SAXS I(q) obtained for different conditions are given in Fig. 5 (a). The qualitative evolution of these SAXS I(q) plots aligns well with WAXS observations, as a significant evolution is visible between the T851 initial state and the sample aged for 1,000 h at 200 °C and minimal differences are observed among the samples aged for 1,000, 2,000, and 5,000 h at 200 °C (Fig. 5 (a)).

All the I(q) plots were fitted with one family of platelets precipitates using the model from Hammouda [24]. Within the investigated q-range, only the average small dimension (thickness) of the precipitates can be

accurately fitted, as the signal related to the larger dimension (length) is not within the observed q-range.

Assuming the composition of the θ' precipitates is Al₂Cu and remains constant, the volume fraction of these precipitates can be determined for each acquisition. Fig. 5 (b) and (c) show the evolution of the fit results (average thickness and volume fraction of θ' platelets) for the T851 initial state and the different samples.

A rapid increase in the mean thickness of the precipitates is observed during the first 1,000 h of ageing (Fig. 5 (b)), with an average rise of more than 50 % (from 48 Å to about 80 Å, depending on the ageing temperature). A slight increase in thickness is noted when increasing the ageing temperatures (from 75 Å to 95 Å for the sample aged 1,000 h for instance). Between 1,000 and 5,000 h of ageing, the thickness shows minimal change, with negligible evolution for ageing temperatures above 200 °C. Regarding the volume fraction (Fig. 5 (c)), a slight increase is observed during the first 1,000 h, but no significant differences are measurable among the various ageing durations. Additionally, a minor decrease in volume fraction is noted when increasing the ageing

<figure><img src="image_8.png" /><figcaption>**Fig. 6.** Example of STEM dark field observations obtained for (a) the T851 initial state and the samples aged (b) 2,000 h, and (c) 10,000 h at 200 °C. Evolution of (d) the mean thickness, (e) mean diameter and (f) mean aspect ratio of the θ' platelets measured by TEM, with comparison to SAXS results for the evolution of the thickness. The error bars correspond to the standard deviation.</figcaption></figure>
temperatures, which may be attributed to the temperature-dependent equilibrium volume fraction, as well as some solute transfer to the θ phase, leading to some dissolution of the θ' precipitates at elevated temperatures. These results are in good agreement with equilibrium predictions from ThermoCalc using TCAL9 database.

These quantitative results align well with the qualitative WAXS results. A strong correlation is observed when comparing the fitted thickness from SAXS with the inverse of the FWHM measured in WAXS, as shown in the plot in Appendix A. After significant evolution during the first 1,000 h of ageing, the θ' precipitates exhibit a strong resistance to coarsening as their size does not significantly evolve for further ageing. The WAXS data further validates this resistance and also shows that the θ' to θ phase transition does not occur significantly after 1,000 h of ageing. However, the slow evolution of the volume fraction of the θ' precipitates is not clearly identified in WAXS when looking at the evolution of the area of the (110)θ' peak. This discrepancy can be explained by the relatively lower precision of the WAXS results compared to the SAXS results when looking at this parameter. Specifically, the standard deviation of the peak area measured at 10 different positions within the same T851 sample represents 16 % of the average value in the WAXS results, whereas the standard deviation for the fitted volume fraction obtained from SAXS at the same positions is only 6 % of the average.

TEM observations were performed on the initial T851 state and on samples aged for 1,000, 2,000, 5,000, and 10,000 h at 200 °C, coming from the preliminary single-temperature heat treatments. Examples of these observations are given in Fig. 6. Although these observations were not conducted on the same samples as the SAXS acquisitions, they provide measurements of both the thickness and length of the platelets (Fig. 6). More than 100 precipitates were measured for each condition and a good agreement was obtained between the thickness measurements obtained by TEM observations and by SAXS acquisitions. TEM observations also offer valuable insights into the evolution of the aspect ratio during ageing. For the T851 initial state, an aspect ratio of 31 ± 7 was measured while a mean aspect ratio of 20 ± 7 was obtained during ageing at 200 °C, with no significant evolution as a function of time. θ precipitates were also observed but no measurements were performed for these stable coarse precipitates as they do not contribute to precipitation hardening.

## 5. Modelling of precipitation hardening

To estimate the effect of the microstructural evolution on the mechanical properties, a model of precipitation strengthening was implemented. Considering the results from WAXS and TEM observations, only the evolution of the θ' precipitates was assumed to have a significant effect on the mechanical properties during ageing, and other contributions to the yield strength such as solid solution strengthening, grain boundary strengthening, work hardening and friction stress were considered as constant. During ageing at the investigated temperatures, the grain size and thus the Hall-Petch contribution does not evolve and the work-hardening contribution also stays constant. Solid solution strengthening may however evolve. APT measurements (that will be detailed in Section 6) were performed on a sample from the initial T851 state and a sample aged 15,000 h at 200 °C. The composition of the solid solutions was then computed for these two states using the DIAM method [25]. The solid solution composition displays minimal variations between these conditions as the Cu content evolve from 0.26 at% to 0.22 at% while Mn and Zr contents stay constant at 0.11 at% and 0.01 at% respectively. Consequently, only the evolution of precipitation hardening will be considered, while all other strengthening contributions are assumed to remain constant.

Therefore, we can model the yield strength as:

$$ \sigma_y = \sigma_0 + \sigma_{ppt} = \sigma_0 + M\tau_{ppt} \quad (2) $$

with $\sigma_{ppt}$ the contribution of precipitation hardening to the yield strength, $\sigma_0$ the contribution of all the other mechanisms, M the Taylor

coefficient of the alloy and $\tau_{ppt}$ the contribution of the hardening precipitates to the critical resolved shear stress.

Considering the work hardening plots, shown in Fig. 7, the high initial work-hardening for the different conditions, much larger than μ/20, suggests that precipitates are non-shearable. This aligns well with findings from Sharma et al. [3] that demonstrated that the transition from shearable to non-shearable precipitates is visible on the work-hardening curves before reaching the T8 condition. Therefore, a micro-mechanical model for by-passing was applied for the θ' hardening precipitates, considering their plate-like morphology.

Zhu and Starke [26] developed such a law in a computer experimental study. The contribution of non-shearable plate-like precipitates in {100} planes was estimated to be:

$$ \tau_{ppt} = 0.13 \frac{Gb}{\sqrt{D_p t_p}} \left( \sqrt{f_v} + 0.75 \sqrt{\frac{D_p}{t_p} f_v + 0.14 \frac{D_p f_v^3}{t_p}} \right) \ln \left( \frac{0.079 D_p}{r_0} \right) \quad (3) $$

with G the Shear modulus, b the Burgers vector, $r_0$ the inner cut-off radius for the calculation of the dislocation line tension, $D_p$ and $t_p$ the mean diameter and thickness of the precipitates and $f_v$ their volume fraction.

This model was therefore implemented considering the SAXS results giving $t_p$ and $f_v$, and assuming that the aspect ratio of the platelets ($D_p/t_p$) remains constant during ageing and equal to 20 (as measured by TEM). The estimated values of the different parameters are given in Table 2.

An excellent correlation between the equivalent yield strength (measured by hardness profiles) and the yield strength model using SAXS data is observed for $\sigma_0 = 80$ MPa as illustrated in Fig. 8.

The correlation between the evolution of the modelled and experimental yield strength suggests that the mechanical resistance to ageing is directly related to the high thermal stability of the θ' precipitates in the temperature range considered for ageing durations exceeding 1,000 h. For the T851 initial state, the average aspect ratio was measured to be 30 and this value was used to predict the yield strength accordingly. The predicted yield strength based on the SAXS results for the T851 initial is 290 MPa, which is in the good range considering the experimentally measured yield strength of 305 MPa for this condition.

## 6. Investigation of the origin of thermal stability

To investigate the thermal stability of the metastable θ' precipitates, additional characterisation was carried out.

<figure><img src="image_9.png" /><figcaption>Fig. 7. Work-hardening plots for the T851 initial state and samples aged 1,000 h, 5,000 h and 15,000 h at 200 °C (these correspond to the same samples as in Fig. 1 (a)).</figcaption></figure>
<table><caption>**Table 2**
Values of the different parameters for the yield strength model.</caption><thead><tr><th>Symbol</th><th>M</th><th>B</th><th>G</th><th>σ<sub>0</sub></th><th>D<sub>p</sub>/t<sub>p</sub></th><th>σ<sub>0</sub></th></tr></thead><tbody><tr><td>Value</td><td>2.8</td><td>2.86 Å</td><td>28 GPa</td><td>5.72 Å</td><td>20</td><td>80 MPa</td></tr></tbody></table>

## 6.1. Segregations at the interface of the θ' precipitates

APT acquisitions were performed on a sample in the T851 initial state and on a sample aged for 15,000 h at 200 °C, in order to examine potential segregations at the precipitate interfaces. Investigating segregations at the interfaces of platelet precipitates using APT can be challenging. Araullo-Peters et al. [27] demonstrated that, depending on the probing direction of the atom probe, segregations of Mg and Ag at the coherent interface of T₁-Al₂CuLi platelets may or may not be detected. When the surface investigated for segregations is perfectly perpendicular to the probing direction, segregations are more easily detected, while moving away from this ideal orientation reduces precision in the depth coordinate, making the identification of segregations more challenging [28].

Although several θ' precipitates were detected within the APT tips, due to the precipitate size and the limited volume of a typical APT tip, no θ' platelet was fully captured within a single tip (Fig. 9 (a) and (b)). Nonetheless, concentration profiles across the coherent (C) and semi-coherent (SC) interfaces of these precipitates were computed. The results are presented in Fig. 9 (c) and (d). In the T851 initial state, a low level of Mn segregation (0.5–1 at%) is visible at the C interface of the θ' precipitates, while stronger segregation of up to 3 at% is identified at the SC interface, although the uncertainty is higher for the SC interface due to the smaller analysed surface area. Low Zr segregation (0.2 at%) can be observed at the C interface, while no discernible Zr segregation is detected at the SC interface. For the C interface, a significant effect of the platelet orientation relative to the probing direction is evident, as higher Mn and Zr segregations are observed at the top interface compared to the bottom interface. This variation can be attributed to the platelet orientation and the pronounced difference in the evaporation field between the Al-matrix and the θ'-Al₂Cu precipitate. This implies that the spatial resolution on the bottom interface is worse than for the top interface, and increases the uncertainty of the maximum Mn concentration measured. No other segregation (e.g. in Si) was detected, and no significant level of Mn and Zr was found in the precipitate core.

For the sample aged 15,000 h at 200 °C, the orientation of the C interface is more favourable for better accuracy. Mn segregation is observed of the same order than in the T851 state (~0.5 at%), and the precipitates are enriched in Mn (0.2 at%) compared to the aluminium matrix (approximately 0.1 at%). Additionally, the SC interface shows higher Mn segregation (approximately 5 at%) compared to the initial state (3 at%). Zr segregation is observed at the C interface in both

conditions, albeit at lower levels than Mn (approximately 0.2 at%), while no Zr segregation is discernible at the SC interface. No conclusion can be drawn for Zr segregation at this interface due to the limited statistics. The profile on the C interface of the aged specimen also indicates significant Zr enrichment in the precipitates compared to the matrix (~0.15 at%), which was also observed by Jiang et al. [29]. Segregation of V, Ti or Si were not observed at the interfaces of the θ' precipitates, and these elements were not plotted to avoid overcrowding the figure. The higher Mn segregations on the SC interface than the C interface align well with findings from Poplawsky et al. [30].

## 6.2. Anisothermal heat treatments

Anisothermal heat treatments were conducted to assess whether a stabilising mechanism is developed during the first hours of ageing, that could improve ageing resistance for further ageing. As Poplawsky et al. demonstrated that Mn and Zr segregations at the θ' interfaces resulted in slow evolution at 300 °C, a first sample aged 15,000 h at 200 °C was further aged at 300 °C for 1,000 h. This sample exhibited similar hardness as the sample aged 1,000 h at 300 °C from the T851 initial state, with a strong loss of hardness showing that no specific stabilisation mechanism for such high temperature was developed during ageing at 200 °C.

Additional anisothermal heat treatments were implemented. Samples aged 1,000 h at 150 °C, 175 °C and 200 °C, 2,000 h at 150 °C and 175 °C and 5,000 h at 150 °C were subsequently put in a furnace at 200 °C for 1,000 h. The resulting hardness values were measured, and the equivalent ageing time at 200 °C for each condition was calculated using equation (1), by adding the equivalent ageing time of each step (in this case, adding 1,000 h to the initial equivalent ageing time for each condition). These results are plotted in Fig. 3. The data from the anisothermal ageing treatments show good alignment with the results from the graded samples, suggesting that calculating an equivalent ageing time for conditions subjected to complex ageing treatments with multiple steps can accurately predict the resulting mechanical properties, provided the ageing temperatures remain within the studied range.

# 7. Discussion

In the Al-4Cu alloy, a rapid evolution of the θ' precipitates is typically observed during over-ageing [31], which is consistent with phase field simulations [32]. After about 200 h at 200 °C, a phase transition from θ' to θ is expected [32], leading to a continuous decrease in the volume fraction of θ' precipitates with further ageing. Chen et al. [33] also reported a significant increase in the length of θ' precipitates during ageing at 225 °C in an Al-4.8Cu alloy, with an average length reaching 580 nm after just 72 h. Similarly, Merle and Fouquet [34] studied the evolution of the thickness of θ' platelets in three Al-Cu binary alloys

<figure><img src="image_10.png" /><figcaption>**Fig. 8.** Comparison between the predicted yield strength using SAXS data and the experimental equivalent yield strength measured by hardness for the different ageing conditions.</figcaption></figure>
<figure><img src="image_11.png" /><figcaption>**Fig. 9.** APT acquisitions for (a) the T851 initial state and (b) a specimen aged 10,000 h at 200 °C with aluminium atoms in blue, 20 at%Cu isosurfaces in orange to identify θ' platelets and concentration profiles along the coherent (C) and semi-coherent (SC) interface for (c) the T851 initial state and (d) the specimen aged 15,000 h at 200 °C.</figcaption></figure>

(Fig. 10). After ageing at 225 °C, the three alloys exhibited thickness from 15 nm to 20 nm, with the thickness increasing as the Cu content increased. However, in the 2219 alloy, the θ' precipitate volume fraction remains stable after 1,000 h of ageing at temperatures up to 245 °C, with the average precipitate length not exceeding 250 nm after 10,000 h at 200 °C and thickness not exceeding 10 nm even after 5,000 h at 245 °C, indicating the presence of stabilising mechanisms (Fig. 10).

Shower et al. [35] explored potential stabilising mechanisms involving interfacial segregations. Four phenomena were proposed to explain the enhanced stability of the θ' precipitates in presence of interfacial segregations: (1) the formation of a new interface, which could reduce the interfacial energy, (2) interactions between solute elements and θ'/Al coherency strains, (3) solute atoms lowering the mobility of Cu atoms, and (4) a solute drag effect. According to Shower et al. [35], given the potential stabilising effect of each mechanism and based on the morphology of Mn segregations typically observed in the 2219 alloy and other ACMZ alloys, the solute drag and interfacial energy reduction mechanisms are likely responsible for the stabilisation provided by Mn segregations. In contrast, Zr segregations appear to stabilise

the precipitates by reducing lattice strain.

Various studies on ACMZ alloys have explored this high thermal stability [14,15,29,30,32,35–38]. Although the composition of the studied 2219 alloy is very similar to that of the alloy investigated by Poplawsky et al. [30], the 2219 alloy exhibits lower stability at high temperature. While the ACMZ alloy of Poplawsky et al. demonstrated minimal evolution after thermal ageing at 300 °C, the present alloy shows a significant decline in mechanical properties after 1,000 h of ageing at 300 °C. Specifically, the ACMZ alloy conserved a hardness greater than 100 HV after 2,000 h at 300 °C, while the 2219 alloy exhibited a hardness of 87 HV after 1,000 h at 300 °C. SEM observations of the aged sample (available in Appendix B) revealed a high density of the equilibrium θ phase, indicating that the θ' precipitates coarsened sufficiently to undergo a θ' to θ phase transition, which explains the substantial loss in mechanical properties. Therefore, it can be assumed that even minor differences in composition can significantly impact the stabilising mechanisms of the θ' precipitates. Shyam et al. [14] demonstrated that the Si content could have a considerable negative effect on the stabilising mechanisms, as Si compete with Mn for similar
<figure><img src="image_12.png" /><figcaption>**Fig. 10.** Comparison between the evolution of the thickness of θ' platelets in the 2219 alloy measured by SAXS and the θ' thickness observed by Merle and Fouquet [34] in Al-Cu binary alloys after 1,000 h at 225 °C.</figcaption></figure>

positions at the Al/θ' interfaces. This was also observed by Biswas et al. [39]. As Si diffuses faster than Mn, the Mn segregations can be reduced by increased Si content. The Si content in the present alloy (0.04 wt%) is higher than the ACMZ alloy from Poplawsky et al. [30] (0.01 wt%), however APT observations did not reveal any level of segregation of Si at the precipitate interfaces. The 2219 alloy also contains a small amount of vanadium, primarily intended to form dispersoids that help control the grain structure [40]. This vanadium content may also play a role in the stability of the θ' precipitates, although we detected no specific segregation of this element either. Another difference to note is the variation in the initial heat treatment: for the ACMZ alloy, the heat treatment was carried out at 240 °C for 4.5 h, which may have promoted a better development of the stabilising mechanism compared to the 2219 alloy, where the initial heat treatment was performed at 175 °C. The diffusivities of Mn and Zr are more sensitive to temperature variations than that of Cu, as their activation energies (212 kJ/mol for Mn and 242 kJ/mol for Zr) are significantly higher than that of Cu (133 kJ/mol). This could result in the formation of better stabilising segregations during the initial, higher temperature heat treatment, with minimal coarsening of the precipitates.

Concerning the results of the anisothermal heat treatments, the validity of time-temperature equivalences for these samples further supports the selected activation energy for the ageing process. This suggests that the stabilising mechanism is likely governed by a thermally activated process, with an activation energy comparable to that of copper diffusion. While the activation energies for Cu and Mn diffusion differ (133 kJ/mol versus 212 kJ/mol), this may initially seem to contradict the proposed stabilising mechanism of Mn segregations. This fact may either mean that Mn transport during coarsening remains local (within a given precipitate) whereas Cu moves at longer distance, or that the Mn stabilising effect plays mostly an energetic role and not a kinetic role.

Coming back to the experimental facts, the important features about segregation evolution are:

i. significant increase of Mn segregation at the SC interface.

ii. no significant change of Mn segregation at the C interface.

iii. significant enrichment of the precipitation core both in Mn and Zr.

Similar evolutions during ageing were reported by Jiang et al. [29] regarding the transformation of the SC interface, and by Gwalani et al. [38] concerning the enrichment of the precipitate core. These facts, together with the rapid decrease of strength in the initial times of ageing followed by stabilisation, strongly suggest that the segregation at the C

interface initially present in the T851 state is not sufficient to prevent rapid coarsening, and that the stabilisation at longer ageing times is related to the segregation at the SC interface and to the solute enrichment within the precipitates.

Finally, the observed time-temperature equivalence provides valuable insight for real-world applications, as the evolution of mechanical properties under complex thermal loadings can be estimated by calculating the equivalent ageing time at 200 °C (by integrating the different steps) and using Fig. 3 to predict the resulting yield strength. This equivalence should theoretically remain valid for heat treatments at temperatures below 165 °C, as no additional phenomena are expected in this range and the solubility of the alloying elements is unlikely to be affected. However, for ageing temperatures above 245 °C, the time-temperature equivalence should be applied with caution, as our results indicate different mechanisms for higher temperatures.

## 8. Conclusions

A high-throughput methodology was employed to characterise the resistance to over-ageing of the 2219 alloy across a continuous range of temperatures from 165 °C to 245 °C for durations up to 10,000 h. Initially, rapid changes in mechanical properties and precipitation hardening were observed within the first 1,000 h. However, a high thermal stability was noted thereafter, with a yield strength of 200 MPa remaining for the most extreme heat treatments. This stability is attributed to the resistance to coarsening of the θ' precipitates, as SAXS and WAXS analyses indicated minimal growth of the θ' precipitates after 1,000 h and no significant phase transformation from θ' to θ, for prolonged ageing beyond 1,000 h.

Additional ageing treatments indicated that this stabilisation does not persist beyond the studied temperature range, as samples aged at 300 °C exhibited significantly lower mechanical properties.

To investigate this high stability, APT acquisitions were conducted and revealed segregations of Mn at the interface of the θ' precipitates for both the T851 initial state and overaged state, as well as solute enrichment (Mn and Zr) during overageing. This segregation and enrichment can slow down coarsening and be responsible for the thermal stability. The initial very low segregation at the precipitate edge could potentially explain the fact that a rapid evolution is seen for the first 1,000 h of ageing before stabilisation.

Time-temperature equivalences were implemented to extrapolate the evolution of mechanical properties during ageing at a given temperature. These equivalences provided satisfactory results for an activation energy close to that of copper diffusion in the alloy, though higher activation energies also yielded acceptable outcomes. Anisothermal heat treatments demonstrated that integrating the thermal history of a sample as an equivalent ageing time at a specific temperature produced reliable results. Therefore, the plausible stabilisation mechanism should also be governed by a thermally activated phenomenon with an activation energy close to that of copper diffusion.

### CRediT authorship contribution statement

**Thomas Perrin:** Writing – review & editing, Writing – original draft, Visualization, Methodology, Investigation, Conceptualization. **Arthur Després:** Writing – review & editing, Visualization, Supervision, Investigation, Conceptualization. **Pierre Heugue:** Writing – review & editing, Supervision, Resources, Project administration. **Alexis Deschamps:** Writing – review & editing, Visualization, Supervision, Investigation, Conceptualization. **Frédéric De Geuser:** Writing – review & editing, Supervision, Software, Project administration, Investigation.

### Funding

Project carried out with the support of the DGAC, the ‘France Relance’ national stimulus plan and the ‘Nextgeneration EU’ European
support program. This work also received support from the DIAMS project of the PEPR DIADEM (ANR-22-PEXD-0005).

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.

Acknowledgement

Authors are thankful to the technical and scientific team of the D2AM-BM02 beamline at the ESRF synchrotron for their assistance for the SAXS/WAXS experiments.

<figure><img src="image_13.png" /><figcaption>Appendix 2. SEM observations of the sample aged 1,000 h at 300 °C</figcaption></figure>

<figure><img src="image_14.png" /></figure>

Data availability

Data will be made available on request.

References

[1] Y.C. Chen, J.C. Feng, H.J. Liu, Precipitate evolution in friction stir welding of 2219-T6 aluminum alloys, Mater. Charact. 60 (2009) 476–481, https://doi.org/10.1016/j.matchar.2008.12.002.

[2] E.M. Elgallad, Z. Zhang, X.-G. Chen, Effect of quenching rate on precipitation kinetics in AA2219 DC cast alloy, Phys. B Condens. Matter 514 (2017) 70–77, https://doi.org/10.1016/j.physb.2017.03.039.

[3] V.M.J. Sharma, K. Sree Kumar, B. Nageswara Rao, S.D. Pathak, Studies on the work-hardening behavior of AA2219 under different aging treatments, Metall. Mat. Trans. A 40 (2009) 3186–3195, https://doi.org/10.1007/s11661-009-0062-4.

[4] V.M.J. Sharma, K.S. Kumar, B.N. Rao, S.D. Pathak, Effect of microstructure and strength on the fracture behavior of AA2219 alloy, Mater. Sci. Eng. A 502 (2009) 45–53, https://doi.org/10.1016/j.msea.2008.11.024.

[5] B. Xie, L. Huang, Z. Wang, X. Li, J. Li, Microstructural evolution and mechanical properties of 2219 aluminum alloy from different aging treatments to subsequent

electromagnetic forming, Mater. Charact. 181 (2021) 111470, https://doi.org/10.1016/j.matchar.2021.111470.

[6] T.R. Prabhu, Effects of ageing time on the mechanical and conductivity properties for various round bar diameters of AA 2219 Al alloy, Eng. Sci. Technol., An Int. J. 20 (2017) 133–142, https://doi.org/10.1016/j.jestch.2016.06.003.

[7] N. Bello, C. Larignon, J. Douin, Long-term thermal ageing of the 2219-T851 and the 2050-T84 Al-Cu alloys, Mater. Today Commun. 29 (2021) 102834, https://doi.org/10.1016/j.mtcomm.2021.102834.

[8] A.M. Hassan, O.M. Bataineh, K.M. Abed, The effect of time and temperature on the precipitation behavior and hardness of Al-4wt%Cu alloy using design of experiments, J. Mater. Process. Technol. 204 (2008) 343–349, https://doi.org/10.1016/j.jmatprotec.2007.11.047.

[9] J.M. Silcock, T.J. Heal, H.K. Hardy, Structural ageing characteristics of binary aluminium-copper alloys, J. Inst. Metals (1954) https://www.osti.gov/biblio/4400209.

[10] L. Bourgeois, C. Dwyer, M. Weyland, J.-F. Nie, B.C. Muddle, Structure and energetics of the coherent interface between the θ' precipitate phase and aluminium in Al-Cu, Acta Mater. 59 (2011) 7043–7050, https://doi.org/10.1016/j.actamat.2011.07.059.

[11] W.M. Stobbs, G.R. Purdy, The elastic accommodation of semicoherent θ' in Al-4wt. %Cu alloy, Acta Metall. 26 (1978) 1069–1081, https://doi.org/10.1016/0001-6160(78)90135-9.
[12] C. Wolverton, V. Ozolins, Entropically favored ordering: the metallurgy of Al 2 Cu revisited, Phys. Rev. Lett. 86 (2001) 5518-5521, https://doi.org/10.1103/PhysRevLett.86.5518.

[13] J.M. Rosalie, L. Bourgeois, Silver segregation to θ' (Al2Cu)-Al interfaces in Al-Cu-Ag alloys, Acta Mater. 60 (2012) 6033-6041, https://doi.org/10.1016/j.actamat.2012.07.039.

[14] A. Shyam, S. Roy, D. Shin, J.D. Poplawsky, L.F. Allard, Y. Yamamoto, J.R. Morris, B. Mazumder, J.C. Idrobo, A. Rodriguez, T.R. Watkins, J.A. Haynes, Elevated temperature microstructural stability in cast AlCuMnZr alloys through solute segregation, Mater. Sci. Eng. A 765 (2019) 138279, https://doi.org/10.1016/j.msea.2019.138279.

[15] S. Bahl, L. Xiong, L.F. Allard, R.A. Michi, J.D. Poplawsky, A.C. Chuang, D. Singh, T. R. Watkins, D. Shin, J.A. Haynes, A. Shyam, Aging behavior and strengthening mechanisms of coarsening resistant metastable θ' precipitates in an Al-Cu alloy, Mater. Des. 198 (2021) 109378, https://doi.org/10.1016/j.matdes.2020.109378.

[16] A. Assadiki, V.A. Esin, R. Martinez, W.J. Poole, G. Cailletaud, Modelling precipitation hardening in an A356+0.5 wt%Cu cast aluminum alloy, Mater. Sci. Eng.: A 819 (2021) 141450, https://doi.org/10.1016/j.msea.2021.141450.

[17] C. Rockenhäuser, C. Rowolt, B. Milkereit, R. Darvishi Kamachali, O. Kessler, B. Skrotzki, On the long-term aging of S-phase in aluminum alloy 2618A, J. Mater. Sci. 56 (2021) 8704-8716, https://doi.org/10.1007/s10853-020-05740-x.

[18] Q. Li, J. Qin, J. Daiping, Y. Danqing, W. Bin, Precipitate coarsening and mechanical properties in 6082 aluminium alloy during long-term thermal exposure, J. Alloy. Compd. 909 (2022) 164819, https://doi.org/10.1016/j.jallcom.2022.164819.

[19] Y. Ge, Z. Que, K. Lindgren, N. Hytönen, M. Thuvander, Effect of thermal aging on microstructure and carbides of SA508/Alloy 52 dissimilar metal weld, Mater. Charact. 200 (2023) 112880, https://doi.org/10.1016/j.matchar.2023.112880.

[20] A. Deschamps, M. Garcia, J. Chevy, B. Davo, F. De Geuser, Influence of Mg and Li content on the microstructure evolution of Al Cu Li alloys during long-term ageing, Acta Mater. 122 (2017) 32-46, https://doi.org/10.1016/j.actamat.2016.09.036.

[21] T. Perrin, A. Després, P. Heugue, A. Deschamps, F. De Geuser, High-throughput characterisation of the long-term ageing of an A357+1wt%Cu cast aluminium alloy using temperature gradient, Materialia 39 (2025) 102378, https://doi.org/10.1016/j.mtla.2025.102378.

[22] T. Perrin, G.A. Chahine, S. Arnaud, A. Després, P. Heugue, A. Deschamps, F. De Geuser, Improving the reliability of small- and wide-angle X-ray scattering measurements of anisotropic precipitates in metallic alloys using sample rotation, J. Appl. Cryst. 57 (2024), https://doi.org/10.1107/S1600576724009294.

[23] G. Ashiotis, A. Deschildre, Z. Nawaz, J.P. Wright, D. Karkoulis, F.E. Picca, J. Kieffer, The fast azimuthal integration Python library: pyFAI, J. Appl. Cryst. 48 (2015) 510-519, https://doi.org/10.1107/S1600576715004306.

[24] B. Hammouda, A new Guinier-Porod model, J. Appl. Cryst. 43 (2010) 716-719, https://doi.org/10.1107/S0021889810015773.

[25] F.D. De Geuser, W. Lefebvre, Determination of matrix composition based on solute-solute nearest-neighbor distances in atom probe tomography, Microscopy Res & Technique 74 (2011) 257-263, https://doi.org/10.1002/jemt.20899.

[26] A.W. Zhu, E.A. Starke, Strengthening effect of unshearable particles of finite size: a computer experimental study, Acta Mater. 47 (1999) 3263-3269, https://doi.org/10.1016/S1359-6454(99)00179-2.

[27] V. Araullo-Peters, B. Gault, F.D. Geuser, A. Deschamps, J.M. Cairney, Microstructural evolution during ageing of Al-Cu-Li-x alloys, Acta Mater. 66 (2014) 199-208, https://doi.org/10.1016/j.actamat.2013.12.001.

[28] D. Blavette, P. Duval, L. Letellier, M. Guttmann, Atomic-scale APFIM and TEM investigation of grain boundary microchemistry in Astroloy nickel base superalloys, Acta Mater. 44 (1996) 4995-5005, https://doi.org/10.1016/S1359-6454(96)00087-0.

[29] L. Jiang, B. Rouxel, T. Langan, T. Dorin, Coupled segregation mechanisms of Sc, Zr and Mn at θ' interfaces enhances the strength and thermal stability of Al-Cu alloys, Acta Mater. 206 (2021) 116634, https://doi.org/10.1016/j.actamat.2021.116634.

[30] J.D. Poplawsky, B.K. Milligan, L.F. Allard, D. Shin, P. Shower, M.F. Chisholm, A. Shyam, The synergistic role of Mn and Zr/Ti in producing θ'/L12 co-precipitates in Al-Cu alloys, Acta Mater. 194 (2020) 577-586, https://doi.org/10.1016/j.actamat.2020.05.043.

[31] J.D. Boyd, R.B. Nicholson, The coarsening behaviour of θ' and θ' precipitates in two Al-Cu alloys, Acta Metall. 19 (1971) 1379-1391, https://doi.org/10.1016/0001-6160(71)90076-9.

[32] P. Shower, J.R. Morris, D. Shin, B. Radhakrishnan, L.F. Allard, A. Shyam, Temperature-dependent stability of θ'-Al2Cu precipitates investigated with phase field simulations and experiments, Materialia 5 (2019) 100185, https://doi.org/10.1016/j.mtla.2018.100185.

[33] Z. Chen, Y. Zhao, Z. Zhang, Theoretical and experimental study of precipitation and coarsening kinetics of θ' phase in Al-Cu alloy, Vacuum 189 (2021) 110263, https://doi.org/10.1016/j.vacuum.2021.110263.

[34] P. Merle, F. Fouquet, Coarsening of θ' plates in al-cu alloys—I. experimental determination of mechanisms, Acta Metall. 29 (1981) 1919-1927, https://doi.org/10.1016/0001-6160(81)90029-8.

[35] P. Shower, J. Morris, D. Shin, B. Radhakrishnan, J. Poplawsky, A. Shyam, Mechanisms for stabilizing θ'(Al2Cu) precipitates at elevated temperatures investigated with phase field modeling, Materialia 6 (2019) 100335, https://doi.org/10.1016/j.mtla.2019.100335.

[36] Y.H. Gao, P.F. Guan, R. Su, H.W. Chen, C. Yang, C. He, L.F. Cao, H. Song, J. Y. Zhang, X.F. Zhang, G. Liu, J.F. Nie, J. Sun, E. Ma, Segregation-sandwiched stable interface suffocates nanoprecipitate coarsening to elevate creep resistance, Mater. Res. Lett. 8 (2020) 446-453, https://doi.org/10.1080/21663831.2020.1799447.

[37] Y.H. Gao, C. Yang, J.Y. Zhang, L.F. Cao, G. Liu, J. Sun, E. Ma, Stabilizing nanoprecipitates in Al-Cu alloys for creep resistance at 300°C, Mater. Res. Lett. 7 (2019) 18-25, https://doi.org/10.1080/21663831.2018.1546773.

[38] B. Gwalani, J. Liu, S. Lambeets, M. Olszta, J. Poplawsky, A. Shyam, A. Devaraj, Rapid assessment of interfacial stabilization mechanisms of metastable precipitates to accelerate high-temperature Al-alloy development, Mater. Res. Lett. 10 (2022) 771-779, https://doi.org/10.1080/21663831.2022.2102947.

[39] A. Biswas, D.J. Siegel, C. Wolverton, D.N. Seidman, Precipitates in Al-Cu alloys revisited: atom-probe tomographic experiments and first-principles calculations of compositional evolution and interfacial segregation, Acta Mater. 59 (2011) 6187-6204, https://doi.org/10.1016/j.actamat.2011.06.036.

[40] Y. Meng, J. Cui, Z. Zhao, Y. Zuo, Effect of vanadium on the microstructures and mechanical properties of an Al-Mg-Si-Cu-Cr-Ti alloy of 6XXX series, J. Alloy. Compd. 573 (2013) 102-111, https://doi.org/10.1016/j.jallcom.2013.03.239.