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

HAL
open science

# Influence of ageing on the high-temperature mechanical strength of the Al-4Mn-3Ni-2Cu-1Zr alloy produced by laser powder bed fusion

Tristan Lenoir, Béchir Chéhab, Isabelle Mouton, Florence Robaut, Frédéric de Geuser, Baptiste Gault, Arthur Després, Jean-Jacques Blandin, Guilhem Martin

<figure><img src="image_2.png" /></figure>

## ► To cite this version:

Tristan Lenoir, Béchir Chéhab, Isabelle Mouton, Florence Robaut, Frédéric de Geuser, et al.. Influence of ageing on the high-temperature mechanical strength of the Al-4Mn-3Ni-2Cu-1Zr alloy produced by laser powder bed fusion. Materials Science and Engineering: A, 2025, 948, pp.149372. 10.1016/j.msea.2025.149372. hal-05380100

HAL Id: hal-05380100

https://hal.science/hal-05380100v1

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_3.png" /></figure>

Distributed under a Creative Commons Attribution 4.0 International License
# Influence of ageing on the high-temperature mechanical strength of the Al-4Mn-3Ni-2Cu-1Zr alloy produced by laser powder bed fusion

Tristan Lenoir<sup>a,b, ID</sup>, Béchir Chéhab<sup>b</sup>, Isabelle Mouton<sup>a</sup>, Florence Robaut<sup>a</sup>, Frédéric De Geuser<sup>a, ID</sup>, Baptiste Gault<sup>c,d</sup>, Arthur Després<sup>a,**</sup>, Jean-Jacques Blandin<sup>a</sup>, Guilhem Martin<sup>a,e, * ID</sup>

## ARTICLE INFO

### Keywords:

Aluminium alloys

Laser powder bed fusion

High temperature

Mechanical strength

Nanoprecipitation

Solid solution

Intermetallics

## ABSTRACT

The mechanical properties of additively manufactured aluminium alloys at room temperature are often optimized through post-fabrication heat treatment (peak ageing). However, at high temperatures, peak-ageing does not necessarily result in the best mechanical properties. This raises the question of whether there is an optimal microstructure depending on the operating temperature. Herein, we examine the influence of ageing in an Al-4Mn-3Ni-2Cu-1Zr (wt.%) alloy processed by laser powder bed fusion on the tensile yield strength from room temperature to 300 °C at 10⁻² and 10⁻⁵ s⁻¹. In the stress-relieved conditions, the material shows a higher yield strength than in the peak-aged conditions at temperatures above 150 °C at 10⁻² s⁻¹ and as low as 100 °C at 10⁻⁵ s⁻¹. This evolution was not expected since nanoscale L1₂-Al₃Zr thermally stable precipitates form in the peak-aged conditions. The microstructure evolutions upon ageing are characterized across all scales and provide insights to rationalize the evolution of the alloy strength with temperature. The role of different microstructural features such as grain size, intermetallic morphology, nanoscale precipitates and solid solution composition is discussed. While the nanoscale Zr-rich precipitates contribute greatly to the room-temperature yield strength, their effect at elevated temperatures, typically >150 °C, is minor. On the contrary, the supersaturation in Mn in the as-built material that is preserved after stress-relief is found to have a large impact on the tensile yield strength at elevated temperature. The globularisation of the elongated intermetallics observed in the stress-relieved microstructure is also identified to lower the elevated-temperature yield strength.

## 1. Introduction

Currently, new opportunities are raised by leveraging the out-of-equilibrium solidification conditions of laser beam powder bed fusion (LB-PBF) to develop processable aluminium alloys specifically dedicated to high-temperature applications, see e.g. Refs. [1–10]. Various alloy design strategies have been employed in the literature to develop new compositions with good mechanical properties at elevated temperatures [11]. Michi et al. have defined three different alloy families depending on the strategy employed [11]. The first approach involves promoting the formation of a high-volume fraction of fine, thermally stable

intermetallics [1–4]. Developing near-eutectic compositions has been identified as a promising strategy for producing microstructures reinforced by a fine network of intermetallic particles. The second family uses ceramic particles for reinforcement instead of intermetallics [12–14]. While the idea is analogous, from an industrial point of view, using ceramics as reinforcing particles presents processing challenges such as the incorporation of the reinforcing particles into the powder feedstock. The third family aims to promote solute trapping of elements that can produce a high-volume fraction of strengthening precipitates during post-fabrication heat treatment. Zr and Sc have been the two most widely used solutes to develop alloys belonging to this family.

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

<footnote><sup>b</sup> Constellation Technology Center, Parc économique Centr'Alp, 725, rue Aristide-Berges, CS 10027, Voreppe cedex, 38341, France</footnote>

<footnote><sup>c</sup> Max-Planck Institute for Sustainable Materials, Max-Planck-Str. 1, Düsseldorf, 40237, Germany</footnote>

<footnote><sup>d</sup> Department of Materials, Imperial College London, South Kensington, SW7 2AZ, UK</footnote>

<footnote><sup>e</sup> Institut Universitaire de France, France</footnote>

<footnote>* Corresponding author. Univ. Grenoble Alpes, CNRS, Grenoble INP, SIMaP, F-38000 Grenoble, France.</footnote>

<footnote>** Corresponding author. Univ. Grenoble Alpes, CNRS, Grenoble INP, SIMaP, F-38000 Grenoble, France.</footnote>

<footnote>E-mail addresses: arthur.despres@grenoble-inp.fr (A. Després), guilhem.martin@grenoble-inp.fr (G. Martin).</footnote>
Nanoscale thermally stable L1₂-Al₃(Zr,Sc) precipitates form during the decomposition of the supersaturated solid solution upon ageing [2,5–7, 15]. Finally, a fourth family could be added, which is not explicitly defined in Ref. [11]. The objective here would be to retain a high concentration of slow-diffusing solutes in the matrix, well beyond the equilibrium solubility limit in the solid state, via solute trapping. These solutes have to interact strongly with dislocations to hinder their motion (glide and climb). This concept is similar to that used for developing solid-solution-strengthened superalloys. It is worth highlighting that these strategies are not mutually exclusive. Novel Al-alloy compositions for high-temperature applications can be developed by leveraging simultaneously several of these alloy design strategies to increase the elevated-temperature mechanical performances of these new alloys.

In Al alloys designed to achieve a Zr supersaturated solid solution in the as-printed condition, the mechanical properties are improved through the nanoscale precipitation of the L1₂-Al₃Zr phase during the post-fabrication ageing treatment. The ageing response is usually investigated through hardness measurements carried out at room temperature to determine the peak-aged condition [15–18]. However, heat treatment that leads to peak ageing at room temperature may also result in other microstructural changes, such as variations in grain size and intermetallic particles size and morphology, as well as changes in the composition of the solid solution. The relative contribution of the various strengthening mechanisms may change with temperature. As a result, the optimal microstructure for elevated-temperature applications may be different than the room-temperature peak-aged condition. For example, Griffiths et al. reported that the Al-2.9Mg-2.1Zr (wt.%) LB-PBF alloy exhibits a higher yield strength above 150 °C in the as-built (AB) condition than after 8h at 400 °C peak-aged condition (150 MPa at 200 °C in the AB condition vs. 90 MPa in the peak-aged condition, i.e. 8h/400 °C) [19]. This result raises the question of the effect of ageing on the mechanical properties as a function of the testing temperature.

In this work, we study the Al-4Mn-3Ni-2Cu-1Zr alloy (wt.%) designed for LB-PBF and recently introduced [10,15,20]. This composition was specifically developed to target high mechanical strength from room temperature up to 300 °C by taking advantage of a large volume fraction of intermetallics and the formation upon ageing of nanoscale thermally stable L1₂-Al₃Zr precipitates. The tensile mechanical properties of this alloy at room temperature and the improvements achieved through ageing at 400 °C have recently been reported [15]. The peak-aged condition was determined after 4h at 400 °C. The evolution of the room-temperature mechanical properties upon ageing was clarified based on the observed microstructural evolutions during ageing at 400 °C. The increase in strength after ageing at 400 °C results from the large strengthening contribution brought by the nanoscale precipitation of the L1₂-Al₃Zr phase. However, the elevated-temperature mechanical properties of this new alloy after ageing have not yet been reported, despite the fact that this material was intentionally designed to maintain a high yield strength up to 300 °C. Additionally, it is unclear whether the microstructure in the peak-aged state would result in the highest yield strength at elevated temperatures due to the temperature dependency of the various strengthening mechanisms.

It is worth mentioning that a few alloys showing similarities with the composition of the alloy studied here, namely a large Mn-content (~4–5 wt.%) combined with the presence of Zr and/or Sc, have been introduced in the additive manufacturing literature.

* Alloys based on the Al-Mn-Zr system and its variants with Cr and/or Mg additions [21–23];

* Alloys derived from the Al-Mn-Mg-Sc-Zr system [9,24–26].

The tensile properties of both Al-5Mn-0.8Cr-0.6Zr-0.17Fe-0.24Si and Al-4.8Mn-0.9Cr-1.1Zr-1.2Mg-0.19Fe-0.20Si (in wt.%) in their peak-aged conditions, respectively 14 and 8h at 375 °C, have been recently determined between room temperature and 300 °C at a strain rate of 10⁻⁴ s⁻¹ [23]. However, comparisons with their as-printed counterparts

were not established and the solid solution was not characterized quantitatively making difficult to fully establish microstructure-property relationships. In addition, possible heterogeneities of the solid solution composition induced by the variations of local solidification conditions at the melt pool scale were not considered. When it comes to the Al-4.46Mn-1.09Mg-0.80Sc-0.64Zr-0.08Fe (wt.%), its high temperature tensile properties and creep performances have recently been investigated in the peak-aged condition (5h at 300 °C) [9, 26], but again no comparison with the properties of the as-printed material was established. This composition has particularly retained our attention because it has a double interest for us. On the one hand, this Al-Mn-Mg-Sc-Zr alloy has a similar Mn-content as the one introduced in our Al-4Mn-3Ni-2Cu-1Zr alloy and contains Zr. However, on the other hand, it also shows significant differences. Firstly, our alloy does not contain Sc, whereas Sc has a diffusivity that is three orders of magnitude greater than Zr and nearly two orders of magnitude greater than that of Mn at 400 °C in Al [11]. This could significantly affect the ageing response, as L1₂-Al₃Zr precipitates require a higher temperature and/or longer ageing times than L1₂-Al₃Sc precipitates to be formed. Secondly, the large amounts of Ni (3 wt%) and Cu (2 wt%) introduced into the composition under investigation result in a higher volume fraction of intermetallic particles in the as-printed conditions, about 20 % [15] in volume fraction against 5–6 % in the Al-Mn-Mg-Sc-Zr alloy [24].

Herein, we compare the evolution of the yield strength with temperature of the Al-4Mn-3Ni-2Cu-1Zr (wt.%) alloy in the stress-relief conditions (4h at 300 °C) with the ones measured after different ageing conditions at different strain rates. Through microstructural characterizations conducted from the grain scale down to the atomic scale, we also characterize the microstructural evolutions occurring upon ageing at both 300 °C and 400 °C. While a few microstructure observations look similar to the ones presented in Ref. [15], particularly the EBSD maps and SEM images, we have collected new data specifically for this work. More importantly, the results have been supplemented and consolidated. The ageing response at 300 °C has not yet been reported, and the APT measurements have been significantly enriched both statistically with additional samples and numerically with the addition of three conditions (as-built, 400 °C for 4h, and 300 °C for 192h) and a systematic analysis conducted at two different locations within the melt pool to capture the effect of variations in local solidification conditions. In light of these microstructural evolutions, we aim to clarify the influence of ageing on the elevated-temperature strength by discussing the role of the grain structure, solid solution, nanoscale thermally stable precipitation, and intermetallic size and morphology.

## 2. Experimental procedures

### 2.1. Powder and LB-PBF processing conditions

The composition of the pre-alloyed powder, measured by inductively coupled plasma atomic emission spectrometry (ICP-AES) is given in Table 1 in both wt.% and at.%. This composition was developed by Constellium and is commercially available under the trademark Ahead®HT1. It was produced by gas atomization. The particle size distribution is between 20 µm and 60 µm. Wall samples (dimensions of 60 mm × 60 mm × 11 mm) were fabricated by melting the pre-alloyed powder in an EOS M290 LB-PBF machine operating under a high-purity argon atmosphere. Fabrication was carried out on a build plate heated at

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

Chemical composition of the pre-alloyed powder batch measured by ICP-AES (in wt.% and at.%).</caption><thead><tr><th></th><th>Al</th><th>Mn</th><th>Ni</th><th>Cu</th><th>Zr</th></tr></thead><tbody><tr><th>wt.%</th><td>Bal.</td><td>3.87</td><td>3.14</td><td>1.86</td><td>1.02</td></tr><tr><th>at.%</th><td>Bal.</td><td>2.01</td><td>1.53</td><td>0.84</td><td>0.32</td></tr></tbody></table>
a constant temperature of 100 °C. The power was kept constant at 370 W and the laser scan speed was set at 1400 mm/s with a laser spot size of 80 µm. The processing parameters were optimized for a layer thickness of 60 µm to produce samples free of lack-of-fusion defects. A bidirectional scanning strategy with a 110 µm hatching distance and a rotation of 67° of the scanning pattern between two consecutive layers was used.

One sample was removed from the build plate to study the as-built microstructure (AB) while other samples were subjected to a stress-relief (SR) heat treatment at 300 °C for 4h in air before being cut from the build plate. Following the SR heat treatment, some samples were aged at 300 °C or 400 °C for different a time ranging from 0.5 to 500h, to study the ageing response of this alloy and determine the peak-aged conditions.

## 2.2. Microstructural characterization

Samples for scanning electron microscopy (SEM) observations were prepared by standard polishing procedures: mechanical grinding with SiC abrasive papers up to grade P1200, followed by successive polishing with a 9 µm, 3 µm and 1 µm diamond suspensions. A final polish was performed with a 0.03 µm colloidal silica solution for 15 min. SEM observations were made in the plane containing the building direction using a Zeiss Gemini SEM 500 FEG equipped with an electron back scattered diffraction (EBSD) camera. The fraction and sizes of the intermetallic particles were estimated based on SEM back scattered electron (BSE) micrographs. The images were segmented using the ImageJ software. A threshold was then applied to measure the area fraction of intermetallics and their size was estimated using the Feret diameter. The variations measured on at least three different SEM-BSE micrographs taken in the SR condition were considered to estimate the error bars. Low magnification EBSD maps (440× 340 µm) were acquired using a 200 nm step size. The EBSD maps were cleaned using a neighbour orientation correlation followed by a grain dilation. These options are available in the EDAX OIM Analysis software. From the EBSD data, columnar grains were segmented by considering only the grains with an aspect ratio <0.7 and a size >2 µm. The average columnar grain size was then estimated using the average equivalent diameter (diameter of a sphere whose projected area is the same as that of the average columnar grain) and the surface fraction of columnar grains determined by the fraction of points in the partition. Error bars were estimated based on the variations measured on three different EBSD maps collected in the SR condition. High magnification EBSD maps (55× 42 µm) were acquired with a step size of 50 nm and then cleaned using the same procedure. Equiaxed grains were segmented by considering only the grains with an aspect ratio >0.9 and a size <1 µm. The average equiaxed grain size was then estimated using the average equivalent diameter.

Thin foils for transmission electron microscopy (TEM) observations in the undeformed samples were prepared by electropolishing using a Tenupol 5 operating at -25 °C under 15 V in a solution consisting of 34 % nitric acid and 66 % methanol. A JEOL 2100F TEM operating at 200 kV was employed for automated crystallographic orientation mapping (ACOM) using the hardware and software toolset ASTAR™ from Nanomegas [27], with an 8 nm step size.

Atom probe tomography (APT) analysis was used to investigate the evolution of the solid solution composition and the nanoscale precipitation upon ageing. APT specimens were extracted from the AB and SR (4h/300 °C) samples using an FEI Helios plasma focused ion beam (PFIB) and then analysed by a local electrode atom probe (LEAP) 5000 XR™ from Cameca Instrument Inc. The different steps of the APT sample preparation using FIB is summarized in Fig. S1 provided in Supplementary Materials. First, the region of interest (ROI) was identified and protected thanks to a Pt-deposition at the surface. Then the material underneath the Pt-protection was extracted using FIB milling and turned into a needle-shaped specimen by annular FIB milling. Additional APT specimens were also taken from the SR+1h/400 °C, SR+4h/400 °C and

SR+192h/300 °C samples using a Thermo Fisher Helios 5 CX PFIB and subsequently analysed using a LEAP 6000 XR™ from Cameca Instrument Inc.

The specimens were systematically field evaporated in voltage-pulsed mode, at a temperature of 40 K with a pulsing rate of 125 kHz, a pulse fraction of 20 %, and a detection rate of 1 ion detection per 1000 pulses on average. The data were analysed and reconstructed using the AP Suite 6.1 software and the composition of the solid solution was determined using the distribution of isolated atoms for the determination of the matrix composition (DIAM method) described in Ref. [28]. The DIAM method is particularly relevant to measure the solid solution composition in presence of nanoscale precipitates. It is based on the calculation of the relative amount of isolated solute atoms (i.e. not surrounded by similar atoms) as a function of a distance d in the APT dataset. First, the bulk composition of the specimen is extracted from the AP Suite 6.1 software. Then, the first nearest neighbour distribution is calculated and extracted for each solute. This distribution is then compared to a random distribution of solute atoms in order to determine the matrix composition. The interested reader is referred to Ref. [28] to get a more detailed description of the DIAM method. The main advantage of this method is that it does not depend on any user parameter which contrasts with other methods relying on cluster selection algorithms. Besides, it works equally well regardless of the size, morphology and distribution of the precipitates, since it only recognizes them as local solute-rich regions. The average solid solution composition was determined from the results of at least 3 specimens from the two different locations: fine-grained zones (FZ) and coarse-grained zone (CZ) within the melt pool for each condition. Note that the solid solution composition can be locally affected by the presence of grain boundaries or large intermetallic particles. In this case, a region of interest is defined in the APT sample away from such microstructural features. A schematic illustration of a typical region of interest within the fine-grained zone (FZ) and the coarse-grained zone (CZ) from which APT samples were extracted using FIB is presented in Fig. S2 provided in Supplementary Materials. Particular attention has been paid the extraction of APT samples from the centre of a FZ and from a columnar grain located in the CZ, respectively.

In the aged samples, the volume fraction of the nanoscale $L_{1_2-Al_3Zr}$ precipitates was estimated by taking the average concentration measured in the APT samples extracted from the FZ and CZ and using the lever rule as follows:

$$f_{V}^{Al_{3}Zr} = \frac{C_{V}^{Zr} - C_{SS}^{Zr}}{C_{Al_{3}Zr}^{Zr} - C_{SS}^{Zr}} \quad (1)$$

where $C_V^Zr$ is the average Zr concentration in the volume analysed, $C_{SS}^Zr$ the composition in Zr in the solid solution and measured using the DIAM method [28], and $C_{Al_3Zr}^Zr$ the average Zr concentration in the nanoscale $L_{1_2-Al_3Zr}$ precipitates estimated using proximity histograms. These proxigrams were averaged on the ten largest Zr-rich precipitates for each aged specimen (based on their volume).

## 2.3. Room- and elevated-temperature mechanical testing

Vickers hardness measurements were carried out on a Tester FALCON 455 using a load of 30 kg and a dwell time of 10 s. A minimum of 10 measurements were taken for each condition. Different samples were used to probe the hardness during isothermal ageing at 300 °C and 400 °C.

For tensile tests from room temperature to 350 °C, flat specimens with a dog-bone geometry (10 mm gauge length, 2 mm width and 1 mm thickness) were machined by electro-discharge machining from the LB-PBF wall samples. Tensile tests were carried out at two different strain-rates $\dot{e} = 10^{-2}s^{-1}$ and $\dot{e} = 10^{-5}s^{-1}$ on an ADAMEL DY35 tensile machine equipped with a 2 kN load cell and a resistive furnace. The temperature was regulated with a $\pm 2$ °C precision by 4 thermocouples located in the
furnace near the tensile specimen. The specimen was fixed in the grips, heated to the tensile test temperature and held at this temperature for 10 min before loading. Tensile tests were performed to failure with the loading direction aligned with the build direction. The applied load and crosshead displacement were then used to calculate the engineering stress and strain, which enabled the determination of the yield strength at 0.2 % of elongation (σ₀.₂ %). The reproducibility of the tensile response was verified by performing three tests under some specific conditions.

### 3. Experimental results

#### 3.1. Microstructure evolution during the stress-relief heat treatment (4h/300 °C)

Fig. 1a–b shows EBSD-IPF maps including several melt pools and collected respectively in the AB and SR samples. These maps exhibit the typical bimodal grain structure previously reported for this alloy [10,15,20]. Such a bimodal grain structure is typical of Al alloys processed by LB-PBF containing substantial additions of Zr or Sc [6,9,19]. The mechanism responsible for the fine-grained regions has been detailed previously [10,12,18,19,29] and has been mainly attributed to the formation of primary L1₂-Al₃Zr precipitates in the liquid, which act as an

<figure><img src="image_4.png" /><figcaption>Fig. 1. IPF-EBSD maps collected in (a) the As-Built (AB) sample, and (b) the Stress-Relief (SR) sample using a 200 nm step size. Black regions associated with FZ were segmented using the following criteria: a grain size >2 µm and an aspect ratio <0.7. (c) Crystallographic orientation map for the FCC-Aluminium collected in the FZ of the SR sample (8 nm step size) using ACOM-TEM and reconstructed using the ASTAR™ software. Intergranular intermetallics are displayed in black.</figcaption></figure>

<figure><img src="image_5.png" /><figcaption>Fig. 2. SEM-BSE micrographs with intermetallics in bright and the Al matrix in dark taken at the melt-pool scale, in the FZ, and in the CZ: (a) in the AB condition, and (b) after the SR heat treatment. The yellow dashed line represents the melt pool boundary in the AB sample. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)</figcaption></figure>
<figure><img src="image_6.png" /><figcaption>**Fig. 3.** Examples of 3D APT reconstructions revealing the homogeneity of the solute distribution at the atomic scale and the absence of nanoscale precipitation in the AB condition: (a) in the FZ, and (b) in the CZ.</figcaption></figure>

efficient nucleating agent for Al grains due to a low misfit parameter [10,12]. A CZ, representing about 66 and 62 ± 5 % of the grain structure respectively in the AB and SR conditions, with large columnar grains elongated along the thermal gradient direction (grain width of about 10 µm) is found in the melt-pool centre, while the region near the melt-pool boundaries consists of a FZ with submicron grains. These grains coloured in black in Fig. 1a–b were segmented using a grain size >2 µm and an aspect ratio <0.7.

Fig. 1c shows an orientation map collected in the FZ sample using ACOM in TEM with an 8 nm step size. This map reveals fine equiaxed grains with an average diameter of 500–600 nm and decorated with a high-volume fraction of intermetallics (displayed in black in Fig. 1c). These intermetallics are mostly distributed at the grain boundaries and their nature was already characterized in details in previous studies [10, 15]. The average grain size in the CZ, $d_{CZ}$ does not evolve during the SR heat treatment: $d_{CZ} = 10 \pm 4$ µm in the AB sample and $d_{CZ} = 8 \pm 4$ µm in the SR sample. This shows that the grain structure is stable after 4h at 300 °C.

Fig. 2a–b shows SEM-BSE micrographs taken in the AB and SR samples at different magnifications. The low-magnification images give an overview of the microstructure at the melt pool scale. Enlarged views of the microstructure in the FZ and CZ are also provided in Fig. 2a–b. In the AB sample, the submicron grains are decorated with a high-volume fraction of fine intermetallics in both FZ and CZ (average area fraction of about 20 ± 2 %). However, these intermetallics have very different morphologies depending on their location in the melt pool, as evidenced in Fig. 2a–b. They are mostly intergranular and globular in the FZ (200 ± 50 nm in diameter). Only a few particles are found within the equiaxed grains in the FZ. In the CZ, the intermetallics are distributed in the interdendritic regions and are elongated in an interconnected network (100 ± 50 nm in width). These observations agree well with previous reports that had characterised these intermetallics in details [10,15]. In summary, three families of intermetallic particles were found in the AB

and SR conditions: (Ni,Cu)-rich particles, Mn-rich particles, and Zr-rich particles. The (Ni,Cu)-rich are the ones present in the largest proportion in terms of volume fraction. The (Ni,Cu)-rich and Mn-rich particles were present in both CZ and FZ whereas the Zr-rich ones were only found in the FZ. Two different phases were identified among the (Ni,Cu)-rich intermetallics: the monoclinic $Al_9(Ni,Cu)_2$ phase and the hexagonal $Al_3(Ni,Cu)_2$ phase. The Mn-rich intermetallics are associated with the orthorhombic $Al_{60}Mn_{11}Ni_4$ phase, while the Zr-rich ones exclusively found in the core of the equiaxed grains in the FZ correspond to the primary $L1_2-Al_3Zr$ phase. The latter is responsible for the formation of the submicron equiaxed grains in the FZ [10,15]. TEM-ACOM and EDX analyses are provided in the Supplementary Materials (Fig. S3) and confirm our previous observations summarized above.

The SR heat treatment does not result in changes regarding the volume fraction (20 ± 2 % in both AB and SR conditions), size (200 ± 50 nm in diameter in FZ and 100 ± 50 nm in width in CZ) and morphology of the intermetallic population in both the FZ and the CZ, see the comparisons in Fig. 2a–b.

Fig. 3a–b shows 3D reconstructions of APT specimens extracted in the FZ and CZ from the AB sample. These reconstructions reveal a homogeneous distribution of the different solutes at the nanoscale. The 3D atomic distributions of Mn, Zr, Cu and Ni are homogeneous and reveal the absence of nanoscale precipitates within the limits of resolution of APT [30]. Similar observations have been made in the SR condition, see Fig. S4 in Supplementary Materials.

The composition of the solid solution and its evolution after the SR heat treatment is summarized in Fig. 4a–d. In the AB sample, the solid solution is supersaturated in Mn and Zr in the FZ and CZ as illustrated in Fig. 4a–b. Their concentrations exceed the maximum equilibrium solubility in solid state taken from the binary systems [31,32] and are superior to the nominal concentration in the case of Zr. This is possible because the microstructure consists of nearly 20 % of second-phase particles and most of those intermetallics are free of Zr [10]. Ni is
<figure><img src="image_7.png" /><figcaption>**Fig. 4.** Composition of the different solutes retained in the Al solid solution in the AB and SR conditions: (a) Manganese, (b) Zirconium, (c) Nickel and (d) Copper. Error bars for the FZ and CZ were calculated using the standard deviation of the bulk composition for all the tips extracted respectively in the FZ and in the CZ.</figcaption></figure>

nearly absent from the solid solution (<0.03 at.%) in agreement with its low solubility in FCC-Al while about 0.3 at.% of Cu is retained in the solid solution. These two solutes (Cu and Ni) are mostly found in the different intermetallic particles identified in this alloy [10]. The Zr concentration in the solid solution is very close in the CZ and FZ. This means that the formation of the primary $L_{1_2}-Al_3Zr$ precipitates in the FZ (not detected in APT) does not consume a significant amount of Zr relatively to its overall concentration. The degree of supersaturation in Mn differs significantly between the FZ and CZ. The Mn concentration in the Al matrix in the CZ (1.7 at.%) is nearly twice that measured in the FZ (0.9 at.%). This difference in Mn supersaturation between the FZ and CZ is due to the local solidification conditions at the melt-pool scale. Near the melt-pool-boundaries, i.e. in the FZ, the solidification velocity is slower than in the CZ [33,34]. Consequently, solute trapping is thought to be more efficient in the CZ than in the FZ, meaning that a higher concentration of Mn is retained in solid solution in the CZ.

The solid solution composition was nearly unaffected by the SR heat treatment as shown in Fig. 4a–d. The only difference with the AB sample is the slight depletion in Cu. This is consistent with the higher diffusion rate of Cu in Al compared to the slow diffusing species such as Mn and Zr.

## 3.2. Ageing response at 300 and 400 °C

Comparison of the hardness in the AB (140 HV30) and SR (130 HV30) samples shows that there is a slight decrease in hardness. Since only minor microstructural changes were observed from the mesoscale down to the atomic scale, this difference in hardness is attributed to the relaxation of residual stresses inherited from the fabrication or possibly

the slight recovery that may happen at 300 °C.

Fig. 5 shows the ageing response of the Al-4Mn-3Ni-2Cu-1Zr alloy at 300°C and 400°C, following an initial SR heat treatment at 300°C for 4h. Upon ageing at 300°C, the hardness slightly decreases between 1h and 10h, before increasing to reach the peak-age at nearly 200h. The slight decrease in hardness during the first 10h at 300°C can be attributed to static recovery and/or the possible decrease of residual stresses. At 400°C, the precipitation kinetics is faster, the hardness increases rapidly and reaches the peak-ageing condition after only 4h (nearly 200h are required to reach the peak-age at 300°C). Then, for longer ageing times, the hardness starts decreasing indicating the overageing regime. Note that the hardness measured after 192h/300°C and 1h/400°C are almost similar, around 150 HV30.

Now that the ageing response of the material was determined at both 300 °C and 400 °C using hardness measurements, we will study the evolution of the yield strength with temperature for selected conditions, see those pointed out by coloured arrows in Fig. 5, namely the SR, $SR+192h/300°C$, $SR+1h/400°C$, and $SR+4h/400°C$.

## 3.3. Impact of ageing on the high-temperature mechanical strength

The evolutions of the yield strength for temperatures ranging from room temperature up to 300 °C, at two different strain rates, $10^{-2} s^{-1}$ and $10^{-5} s^{-1}$, are given in Fig. 6a and b, respectively. The evolutions of the yield strength of the SR, $SR+1h/400 °C$ and $SR+4h/400 °C$ conditions are compared over a large range of temperatures. For the peak-aged at 300 °C, we only conducted tests at room temperature and at 300 °C. The data were extracted from the engineering stress-strain curves of tensile tests, which are shown in Fig. S5 in Supplementary
<figure><img src="image_8.png" /><figcaption>Fig. 5. Evolution of the Vickers hardness of the Al-4Mn-3Ni-2Cu-1Zr alloy upon ageing at respectively 300 °C and 400 °C. Note that the Stress-Relief heat treatment (4h at 300 °C) was systematically applied before ageing at 300 or 400 °C.</figcaption></figure>

Materials. Note that at a given temperature, the material exhibits a higher strength at $10^{-2}$ than at $10^{-5}$ s$^{-1}$ as classically observed when entering in the viscoplastic regime, where the flow stress becomes strain rate-sensitive. However, this alloy exhibits a relatively low strain rate sensitivity $m$ in the SR condition with ($m$ = 0.03 at 200 °C), whereas $m$ is doubled at 300 °C ($m$ = 0.06). After ageing at 400 °C, $m$ increases to 0.08 and remains relatively stable between 200 °C and 300 °C. Consequently, there is a larger difference in flow stress between the specimens deformed at $10^{-2}$ and $10^{-5}$ s$^{-1}$ after ageing. In addition, a ductility loss is observed when reducing the strain rate (elongation to failure lower at $10^{-5}$ s$^{-1}$ than at $10^{-2}$ s$^{-1}$). The origin of this ductility loss was recently investigated in Ref. [35] and was attributed to a change in deformation mechanism in the region consisting of submicron equiaxed grains that causes early intergranular damage in this specific region. In this work, we focus on the evolution of the yield strength as a function of the temperature under different ageing conditions.

As expected, the peak-aged at 400 °C (SR+4h/400 °C) shows the highest yield strength at both $10^{-2}$ (480 MPa) and $10^{-5}$ s$^{-1}$ (430 MPa) at room temperature (Fig. 5a–b). This represents a significant increase compared to the SR condition where the yield strength is about 380 MPa regardless of the strain rate. However, the most striking observation is that the yield strength of the samples aged at 400 °C decreases more

rapidly than the yield strength of the SR sample. This trend remains true at high ($10^{-2}$ s$^{-1}$) and low ($10^{-5}$ s$^{-1}$) strain rates. At a strain rate of $10^{-2}$ s$^{-1}$, the materials aged at 400 °C show a superior yield strength compared to the material in its SR condition below 150 °C. From 200 to 300 °C, the SR microstructure outperforms the aged materials. For example, at 300 °C, the yield strength of the SR sample is about 230 MPa while that of the samples aged at 400 °C is about 150–160 MPa (Fig. 6a). When decreasing the strain rate to $10^{-5}$ s$^{-1}$, this change occurs at lower temperature and the SR sample has a yield strength higher than the aged materials above 100 °C. For example, the yield strength of the SR sample reaches 200 MPa at 250 °C, but is about 120–130 MPa in the materials aged at 400 °C.

The evolutions of the yield strength as a function of temperature for the materials aged 1h and 4h at 400 °C are similar as illustrated by the overlap of the curves in Fig. 6a–b. The case of the sample aged at 300 °C for 192h shows a yield strength lower than its SR counterpart at 300 °C but higher than the ones of the materials aged at 400 °C for 1 or 4h, see the star symbols displayed in orange in Fig. 6a–b.

To clarify why the decrease in yield strength is larger in the samples aged at 300 and 400 °C compared to that of the SR sample, the following sections will focus on the microstructural changes occurring upon ageing.

<figure><img src="image_9.png" /><figcaption>Fig. 6. Evolution of yield strength as a function of temperature for SR, 1h/400 °C, 4h/400 °C and 192h/300 °C conditions and two different strain rates: (a) at $10^{-2}$ s$^{-1}$, and (b) at $10^{-5}$ s$^{-1}$. The yield strength at 0.2 % of deformation was extracted from the engineering stress-strain tensile curves provided in Fig. S5 in Supplementary Materials.</figcaption></figure>
<figure><img src="image_10.png" /><figcaption>**Fig. 7.** Low magnification IPF-EBSD maps acquired using a 200 nm step size in different conditions: (a) SR reproduced from Fig. 2 for clarity, (b) SR+192h/300 °C, (c) SR+1h/400 °C and (d) SR+4h/400 °C. Black regions associated with FZ were segmented using the following criteria: grain size >2 µm and aspect ratio <0.7.</figcaption></figure>

<figure><img src="image_11.png" /><figcaption>**Fig. 8.** High magnification IPF-EBSD maps acquired using a 50 nm step size to reveal the grain structure in the FZ in: (a) the SR condition, and (b) in the SR+4h/400 °C sample. Grain size within the FZ was estimated after segmentation using the following criteria: grain size <1 µm and aspect ratio >0.9.</figcaption></figure>

### 3.4. Microstructure evolution upon ageing

#### 3.4.1. Evolution of the grain structure

Large EBSD-IPF maps containing several melt pools acquired respectively in the SR, SR+192h/300 °C, SR+1h/400 °C and SR+4h/400 °C samples are shown in Fig. 7a-d. Once again, the areas displayed in black correspond to the fine-grained regions which were segmented using a grain size >2 µm and an aspect ratio <0.7. The average surface fraction of CZ obtained by segmentation remains stable for all conditions (about 62, 57, 59, and 63 % ± 5 % for respectively the SR, SR+192h/300 °C, SR+1h/400 °C and SR+4h/400 °C conditions). The average grain size in the CZ, estimated using the definition of the equivalent diameter based on the maps shown in Fig. 7a-d, is stable regardless of the condition examined (between 7 and 8 ± 4 µm), the changes in grain size measured between the SR and the SR+4h/400 °C samples are within the margin of error. The grain structure can thus be considered unchanged. This result is not necessarily surprising given the high fraction of intermetallic particles that help to prevent grain boundary migration upon ageing.

Fig. 8a-b shows high magnification EBSD-IPF maps revealing the

grain structure in the FZ in the SR and SR+4h/400 °C samples. The average grain size in the FZ, estimated using the equivalent diameter, remains stable upon ageing at 4h/400 °C ($d_{FZ}$ = 550 ± 200 nm in the SR sample vs. $d_{FZ}$ = 650 ± 200 nm in the SR+4h/400 °C sample). Regardless of the FZ or CZ, the grain structure can therefore be considered stable upon ageing since the variation in grain size are systematically found within the margin of error.

#### 3.4.2. Evolution of the intermetallic population

SEM-BSE micrographs taken in the FZ and CZ of the SR. SR+192h/300 °C, SR+1h/400 °C and SR+4h/400 °C samples are shown in Fig. 9a-d and Fig. 10a-d, respectively. Similar SEM-BSE micrographs were taken at the melt-pool scale for each condition and are summarized in Fig. S6 in Supplementary Materials.

In the FZ, the intermetallics remain globular regardless of the ageing treatment (Fig. 9a-d). While no evidence of coarsening is detected after ageing at 300 °C for 192h, a slight coarsening of the intergranular intermetallics is observed after ageing 1h at 400 °C, and it becomes more apparent after 4h at 400 °C (Fig. 9c-d). After ageing, the presence of a higher density of intragranular precipitates is evidenced, see examples
<figure><img src="image_12.png" /><figcaption>**Fig. 9.** SEM-BSE micrographs with intermetallics in bright and the Al matrix in dark taken inside the FZ in different conditions: (a) SR reproduced from Fig. 3 for clarity, (b) SR+192h/300 °C, (c) SR+1h/400 °C and (d) SR+4h/400 °C. Yellow arrows point towards intragranular precipitates forming upon ageing. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)</figcaption></figure>

<figure><img src="image_13.png" /><figcaption>**Fig. 10.** SEM-BSE micrographs with intermetallics in bright and the Al matrix in dark taken inside the CZ in different conditions: (a) SR reproduced from Fig. 3 for clarity, (b) SR+192h/300 °C, (c) SR+1h/400 °C and (d) SR+4h/400 °C. Orange arrows indicate the slight coarsening of the particles located at the grain boundaries. Yellow arrows point towards intragranular precipitates forming upon ageing. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)</figcaption></figure>

pointed out with yellow arrows. These intragranular precipitates have been found to be enriched in Mn contrary to the Zr-rich primary precipitates already present in the SR condition, see Ref. [15]. Overall, the evolution of the intermetallic particles in the FZ is rather limited in size (respectively 200 ± 50 nm and 250 ± 50 nm in diameter in the SR and 4h/400 °C conditions) and morphology even after 4h at 400 °C.

In the CZ, more significant morphological changes are observed, especially after ageing at 400 °C (Fig. 10a–d). After 192h/300 °C, the first signs of globularisation of the intermetallics with an elongated morphology are observed only locally (Fig. 10b). It can therefore be considered that the size and morphology of the intermetallics are stable

compared to the SR sample. However, after ageing at 400 °C, there is a drastic morphological change of the intermetallics in the CZ, see Fig. 10c and d. The intermetallic network consisting of elongated particles has been completely globularised with a slight coarsening of the particles located at the grain boundaries, indicated with orange arrows in Fig. 10c and d (respectively 100 ± 50 nm and 250 ± 50 nm in width in the SR and 4h/400 °C conditions). Upon ageing, intragranular precipitates (Mn-rich as shown in Ref. [15]), once again pointed with yellow arrows, have been observed and were absent in the SR conditions.

In a previous study [15], we showed that the nature of the intermetallic phases did not change after 4h at 400 °C. The (Ni,Cu)-rich
intermetallics were the less thermally stable and tend to coarsen significantly after 96h at 400 °C. Here, the few intermetallic particles subjected to globularisation in the CZ and slight coarsening are thus considered to be (Ni,Cu)-rich particles.

### 3.4.3. Decomposition of the solid solution and nanoscale precipitation

Fig. 11a–f shows examples of APT reconstructions of specimens extracted in the FZ and CZ from the aged samples (SR+1h/400 °C, SR+192h/300 °C, and SR+4h/400 °C). In these reconstructions, Mn atoms appear in pink and the presence of Zr-rich nanoscale precipitates is revealed using iso-concentration surfaces delineating regions containing over 2.5 at.% Zr. The typical composition of such Zr-rich precipitates is given by the composition profile calculated as a function of the distance to this set of iso-surfaces – referred to as proximity histograms or proxigrams [36] – shown in Fig. 12a–c. These proxigrams were taken across the ten largest Zr-rich nanoscale precipitates in the samples aged respectively at 192h/300 °C, 1h/400 °C and 4h/400 °C. Regardless of the ageing conditions, those precipitates have the same composition with a stoichiometry corresponding to the Al₃Zr phase, which was previously identified by TEM as L1₂-Al₃Zr [15]. In some specimens, iso-surfaces based on the Mn composition (5 at.% Mn) or Cu composition (5 at.% Cu) allow the observation of the coarser intermetallics described in the previous section. Composition profiles across these intermetallic particles are plotted in Fig. S7, providing more detailed information on their chemical composition.

Fig. 13a–d exhibit the composition in solutes of the solid solution and its evolution after ageing in the three following conditions: SR+192h/300 °C, SR+1h/400 °C and SR+4h/400 °C. The composition of the solid solution is determined in both FZ and CZ based on the 3D APT reconstructions using the DIAM method [28]. The precipitation of nanoscale Zr-rich phases revealed in Fig. 11a–f is associated with a depletion in Zr in the solid solution regardless of the region of interest (FZ or CZ), see Fig. 13a–d. This depletion in Zr in the solid solution is almost similar

in the three ageing conditions examined with only limited differences between the FZ and CZ (within the error bars). Similarly, the decomposition of the solid solution goes along with a strong depletion in Mn in both FZ and CZ upon ageing regardless of the temperature 300 or 400 °C. Indeed, the Mn concentration in solid solution is divided by a factor of 4 in the CZ and by a factor of 2 in the FZ after ageing at 400 °C for 1h compared to the SR condition. Interestingly, the difference in Mn supersaturation between the FZ and CZ in the SR condition has vanished after ageing regardless the temperature (300 and 400 °C) and time (1 or 4h at 400 °C). This depletion in Mn upon ageing is associated to a precipitation of Mn-rich secondary phases, which has already been reported in a previous study [15]. When it comes to Ni and Cu, only minor variations are observed. We recall here that Ni and Cu are mainly found in the composition of the intermetallic particles. Overall, the solid solution composition does not differ significantly when comparing the three ageing conditions.

Assuming that the Zr concentration in solid solution after ageing is, on average, about 0.18 at.% (Fig. 13b) and given the composition of the nanoscale Zr-rich precipitates measured using the proxigrams shown in Fig. 12a–c (around 23 at.% Zr on average in the precipitates), the volume fraction of precipitates was estimated using the lever rule to 0.8 %. This value is in relatively good agreement with the values measured by SAXS in our previous work (0.6 % after 1h at 400 °C and 0.7 % after 4h at 400 °C) [15]. The average radius of the Zr-rich nanoscale precipitates was found to vary between 1 and 1.5 nm regardless of the ageing conditions, see Fig. 11a–f. The size of the Zr-rich precipitates does not seem to vary significantly between the FZ and CZ regardless of the ageing conditions. Although this value is close to the effective spatial resolution of APT [30], it is in relatively good agreement with the values measured by SAXS in the same alloy in our previous work (average precipitate radius of 1.3 nm after 1h/400 °C and 1.5 nm after 4h/400 °C) [15].

<figure><img src="image_14.png" /><figcaption>Fig. 11. Examples of 3D APT reconstructions with Mn atoms displayed in pink and 2.5 at.% Zr, 5 at.% Mn and 5 at.% Cu iso-concentration surfaces drawn respectively in green, pink and orange: (a) FZ and (b) CZ after 192h/300 °C; (c) FZ and (d) CZ after 1h/400 °C; (e) FZ and (f) CZ after 4h/400 °C. All reconstructions reveal the nanoprecipitation of spherical Zr-rich phases after ageing. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)</figcaption></figure>
<figure><img src="image_15.png" /></figure>

<figure><img src="image_16.png" /></figure>

<figure><img src="image_17.png" /><figcaption>**Fig. 12.** Proximity histograms showing the elemental distribution across the five biggest Zr-rich nanoscale precipitates in: (a) CZ after 192h/300 °C, (b) CZ after 1h/400 °C, and (c) CZ after 4h/400 °C, confirming the L1₂-Al₃Zr nanoprecipitation upon ageing.</figcaption></figure>
<figure><img src="image_18.png" /><figcaption>Fig. 13. Composition of the different solutes retained in the Al solid solution after SR heat treatment, SR+192h/300 °C, SR+1h/400 °C and SR+4h/400 °C: (a) Manganese, (b) Zirconium, (c) Nickel and (d) Copper.</figcaption></figure>

# 4. Discussion

## 4.1. Thermal stability of the microstructure after 4h at 300 °C (SR)

The results presented in Section 3.1 show that no significant microstructural evolution occurs, even at the atomic scale. The only variation detected was a slight decrease in hardness that was attributed to the relaxation of residual stresses and static recovery. The thermal stability of the microstructure, exemplified by the solid solution composition that remains nearly unchanged after 4h at 300 °C (Fig. 4a–d) and the absence of nanoscale Mn or Zr-rich precipitates (Fig. 3a–b), is related to the nature of the supersaturated solutes. Mn and Zr are slow-diffusers while solutes having a higher diffusivity in Al such as Ni and Cu are mostly consumed by the intermetallic particles, see Fig. S3. This result is worth highlighting because a heat treatment at 300 °C for 4h in traditional Al alloys from the 2xxx, 6xxx, or 7xxx series would, on the contrary, severely alter the solid solution composition and precipitation state. The comparison with the Al-Mn-Mg-Sc-Zr alloy from the literature is interesting because such a heat treatment (4h at 300 °C) would nearly correspond to the peak-aged. Indeed, the peak-aged was reached after 5–6h at 300 °C [9,24–26]. This was correlated to the profuse precipitation of nanoscale L1₂-Al₃Sc with a limited Zr content entering in their composition, and the concentration in Mn in the solid solution that was reported to be equal to that measured in the as-built microstructure [25]. Such considerations show that microstructural evolutions observed upon ageing may thus be drastically different in the Al-Mn-Mg-Sc-Zr alloy compared to the Al-Mn-Ni-Cu-Zr alloy. For example, in the latter Sc-free alloy, one needs 192h at 300 °C to reach the hardness peak caused by the formation of nanoscale L1₂-Al₃Zr precipitates. This difference in ageing response between the Al-Mn-Mg-Sc-Zr and our Al-Mn-Ni-Cu-Zr alloy results from the higher diffusivity of Sc compared to Zr in FCC-Al. For example, the diffusion of

Sc in FCC-Al is four orders of magnitude greater than that of Zr at 300 °C, see Ref. [32]. This means that L1₂-Al₃Zr nanoscale precipitates require a higher temperature and/or longer ageing times than L1₂-Al₃Sc precipitates to be formed.

## 4.2. Origin of the decrease in yield strength after ageing at elevated temperatures

As expected for a precipitation-strengthened alloy, the yield strength trend at room temperature is such that the yield strength reaches its highest value in the peak-ageing condition. In the presence of nanoscale Zr-rich precipitates (ageing conditions: SR+1h/400 °C, SR+4h/400 °C, and SR+192h/300 °C), the yield strength is higher than in the SR condition where no Zr-rich nanoscale precipitate was observed. However, as the temperature increases above 150 °C, the behaviour is more surprising. The decrease in yield strength with temperature of the aged microstructures strengthened by thermally stable L1₂-Al₃Zr precipitates is larger compared to the same material without strengthening precipitates. These results raise the question of the origin of this large decrease in yield strength at elevated temperature after ageing. Therefore, we aim to discuss the different microstructural features that may be at the origin of this decrease in yield strength at elevated temperatures in the light of the microstructural evolutions observed after ageing at 300 and 400 °C.

### 4.2.1. Role of the grain size

The bimodal grain structure of the Al-4Mn-3Ni-2Cu-1Zr alloy was found to be thermally stable during ageing. No major changes in grain size and morphology were detected after 4h at 400 °C and 192h at 300 °C (Fig. 7a–d). This implies that the high-volume fraction of intergranular intermetallic particles is efficient in pining the grain boundaries. Therefore, the grain size cannot be held directly responsible for
the superior performance of the SR microstructure at elevated temper-
atures compared to the peak-aged samples.

The presence of approximately 35 % of submicron grains can also potentially indirectly affect the material's mechanical behaviour at elevated temperatures, due to a change in deformation mechanism as the temperature increases. In particular, the activation of grain boundary sliding (GBS) can be promoted in the presence of submicron grains. For example, the occurrence of GBS in ultra-fine grained aluminium alloys has been reported in several studies at temperatures as low as 200 °C and a strain rate of 10⁻⁵ s⁻¹ [37,38]. Here, the decoration of the grain boundaries with numerous particles is thought to inhibit GBS and is most likely one of the reasons accounting for the good mechanical strength of the Al-4Mn-3Ni-2Cu-1Zr alloy in the temperature range 150–300 °C. Interestingly, in a recent work, Griffiths et al. [19] reported a similar trend in an Al-2.90Mg-2.10Zr (wt.%) ternary alloy processed by LB-PBF. At temperatures above 100 °C, the yield strength of the peak-aged material was lower than its as-fabricated counterpart. The authors of this study attributed this particular behaviour to the coarsening of the intergranular particles during ageing at 400 °C. Their idea was that the coarsening of the intergranular intermetallics reduced their effectiveness in inhibiting GBS, explaining the drop in yield strength after ageing despite the presence of thermally stable Zr-rich nanoscale precipitates. However, the scenario proposed in Ref. [19] does not seem to apply here. First, SEM observations shown in Fig. 9a–d indicate that very limited coarsening of the intergranular intermetallics occurs in the FZ after ageing, even after 4h at 400 °C. Thus, the decoration of the grain boundaries with a high density of intermetallics after ageing is still believed to be efficient in preventing GBS. Second, the same trend is observed at 10⁻² s⁻¹ and 10⁻⁵ s⁻¹: larger reduction in yield strength at elevated temperatures for the aged materials compared to the SR condition. However, it is unlikely that the GBS activity is significant enough at 10⁻² s⁻¹, especially at 200 °C. This is supported by the study of Liu and Ma [37] who have summarized different studies reporting low-temperature superplasticity in Al alloys. While superplasticity can be achieved for submicron grain sizes at temperatures as low as 200 °C, it often requires a strain rate of the order of 10⁻⁴ s⁻¹, see e.g. Ref. [37]. It is therefore concluded that the better performance of the SR material at elevated temperatures compared to the aged samples cannot be related to a change in deformation mechanism in FZ as proposed in the case of the Al-2.90Mg-2.10Zr (wt.%) ternary alloy [19].

4.2.2. Role of the thermally stable L1₂-Al₃Zr nanoprecipitates

Zr is a key solute in Al alloys designed for AM. Firstly, Zr improves the processability of compositions that are susceptible to cracking by refining the grains via the precipitation of primary L1₂-Al₃Zr particles in the liquid. These primary precipitates are responsible for the formation of submicron equiaxed grains observed near the melt pool boundaries, see e.g. Fig. 1c [10,16]. Second, the slow diffusivity of Zr in liquid aluminium and in its FCC structure make it a good candidate for supersaturation in the solid solution via solute trapping, as illustrated in Fig. 4b. This enables the material to be strengthened during a post-fabrication heat treatment by the precipitation of nanoscale secondary L1₂-Al₃Zr precipitates, as shown in Fig. 11a–f. The improvement of the room-temperature yield strength due to the formation of nanoscale secondary L1₂-Al₃Zr precipitates has been widely reported in Zr-containing AM alloys, see e.g. Refs. [2,5–7,9,18,19,39]. Finally, the beneficial role of these nanoscale secondary L1₂-Al₃Zr precipitates in cast alloys on high-temperature strength and creep resistance has often been emphasized in the literature because of their high thermal stability compared to the traditional strengthening precipitates found in precipitation-hardened aluminium alloys [32,40]. For example, it was reported that the microhardness and creep resistance of cast Al-Sc-Zr alloys increases when increasing their volume fraction. However, the optimum precipitate size depends on the usage temperature [40]. Precipitates with a radius $r \le 2$ nm are preferable for room-temperature applications while larger precipitates (radii typically between 4 and 9

nm) are more desirable for creep resistance at 300 °C [40].

In a recent study, we investigated the mechanical properties of this Al-4Mn-3Ni-2Cu-1Zr alloy at room temperature after ageing [15]. We demonstrated that the Zr-rich precipitates significantly contribute to the yield strength at room temperature. However, their role at high temperatures has not yet been studied in this material. As emphasized earlier, these precipitates are known for their high thermal stability [32, 40]. Therefore, promoting their formation may be desirable where high-temperature strength is required. Here, we show that forming a volume fraction of approximately 0.8 % of Zr-rich spherical precipitates, with radii ranging from 1 to 1.5 nm, does not result in an improvement in yield strength at temperatures above 150 °C. Rather, we observe a decrease in yield strength after ageing, irrespective of the ageing conditions. Although there may be slight differences in volume fraction and size of the precipitates between the sample aged 1h at 400 °C and the one aged for 4h (not detected by APT but evidenced using SAXS, see Ref. [15]), the evolution of their yield strength with temperature is very similar (Fig. 6a–b). These findings lead us to question the beneficial role of these precipitates in impeding dislocation motion at elevated temperatures compared to their high efficiency at room temperature, especially when their volume fraction is of the order of 1 % and their average radius of 1–1.5 nm. We hypothesize that the threshold stress induced by such precipitates is small in relation to the stress required to cause the material to yield. This reduced threshold stress at elevated temperatures can be related to thermally activated deformation mechanisms such as cross-slip and dislocation climb. Given their size, these nanoscale Zr-rich precipitates are expected to be cut by dislocations; however, the increase in temperature reduces their ability to alter dislocation motion due to the aforementioned two thermally activated deformation mechanisms.

4.2.3. Role of Mn in solid solution

The measurements of the evolution of the matrix composition using APT upon ageing provide key insights to discuss the origin of the higher yield strength of the SR sample compared to that of the aged specimens at temperatures >150 °C. The concentration of Mn retained in solid solution decreased drastically after ageing, by a factor of 2 in the FZ, and by a factor of 4 in the CZ. This depletion in Mn in the matrix is accompanied by the formation of Mn-rich precipitates [15], which correspond to the intragranular precipitates observed in the SEM images of the aged samples (Fig. 9b–d and Fig. 10b–d). This suggests that Mn can be much more beneficial for the elevated-temperature strength when trapped in solid solution as in the SR sample, than when Mn is consumed by the formation of precipitates. This key result also suggests that Mn atoms in solid solution strongly affects dislocation motion at elevated temperature.

The recent work of Farkoosh et al. [41] dedicated to the creep properties of age-hardened cast Al-Mn-Zr alloys at 300 °C is particularly relevant to the discussion of the present results. In their study, the authors have compared three alloys to discriminate the effect of different microstructural features. The Al-0.5Mn-0.09Zr-0.05Er-0.05Si (at.%) alloy was compared with its Mn-free (Al-0.11Zr-0.005Er at.%) and Zr-free (Al-0.5Mn-0.02Si at.%) counterparts [41]. In the peak-aged condition, the Mn-free Al-0.11Zr-0.005Er (at.%) alloy exhibited a reduced creep resistance at 300 °C compared to the Al-0.5Mn-0.09Zr-0.05Er-0.05Si (at.%) alloy. Both alloys had L1₂ precipitates of the same size (average radius of 2.5 nm). This provided an initial insight into the key role of Mn in relation to creep properties. In addition, the Zr-free Al-0.5Mn-0.02Si (at.%) alloy in its as-cast state (single phase solid solution alloy) showed a better creep resistance than its Mn-free counterpart (Al-0.11Zr-0.005Er) in its peak-aged state with L1₂ precipitates whose radius was reported to be equal to 2.5 nm [41]. On the opposite, at room temperature, the Mn-free Al-0.09Zr-0.01Er-0.02Si (at.%) alloy was weaker than its as-cast Al-0.5Mn-0.02Si (at.%) alloy. This result further highlights that Mn, when in the solid solution, can have a stronger effect on the creep resistance than the L1₂
precipitates. This observation reported in Ref. [41] is in line with the
decrease in yield strength at elevated temperatures after ageing reported
in this work.

The remarkable effect of Mn in solid solution on the creep properties
was attributed to the Mn-induced retardation of climbing [41]. Two
conditions have to be fulfilled for a solute to hinder the climbing process:
(i) the presence of a solute atmosphere, this implies that a strong
interaction between Mn atoms and the dislocations exists; and (ii) a
slow-diffusivity relative to the Al atoms. These two conditions must be
met because a slow diffuser that would not interact strongly with the
dislocations may not significantly affect their mobility. Interestingly, Mn
satisfies these two conditions. First-principles studies reported that
indeed Mn has a strong solute-dislocation interaction energy (0.49 eV at
0K), more than three times higher than Zr (0.14 eV at 0K) [42]. Mn is
also a slow diffuser in Al ($D_{Al}^{Mn} = 6.24 \times 10^{-19} m^2s^{-1}$ vs. $D_{Al}^{Al} = 3.25 \times 10^{-15} m^2s^{-1}$ for the auto-diffusion of Al at 400 °C) [32].

Finally, it is worth highlighting that the non-equilibrium nature of
the solidification during LB-PBF leads to a highly supersaturated solid
solution, well above the equilibrium solubility, up to 1.6 at.% in Mn in
the CZ of the SR sample, see Fig. 4a. This makes the SR material much
stronger than peak-aged materials (4h/400 °C) at temperatures above
150 °C, since the Mn concentration in solid solution decreases to an
average of about 0.4 at.% after ageing. Mn is therefore a good candidate
to be introduced in Al alloys designed for high-temperature applications.
Here, the ability to retain a large amount Mn in the solid solution by
leveraging solute trapping in the LB-PBF process enables us to further
enhance the high-temperature strength of the Al-4Mn-3Ni-2Cu-1Zr

alloy. Although L1₂-precipitates are often considered beneficial for
high-temperature strength, in the present case with a volume fraction of
the order of 1 % and a radius about 1–1.5 nm, their effect seems minor
compared to the large effect of Mn in solid solution.

4.2.4. Role of the intermetallic network

Finally, the possible role of the intermetallic particles on the high-
temperature mechanical strength is discussed in the light of the behav-
iour of the sample aged at 300 °C for 192h. This sample is interesting
because it has an intermetallic network almost unchanged compared to
the SR condition with no coarsening of the intergranular intermetallics
in the FZ (Fig. 9b), and a limited globularisation (Fig. 10b) of the
elongated intermetallics in the CZ. However, at a finer scale, the sample
aged 192h at 300 °C is significantly different from the SR sample and
shows similar characteristics as the samples aged at 400 °C, i.e. a
reduced Mn concentration in the solid solution and the presence of
nanoscale Zr-rich precipitates. As a result, the decrease in yield strength
at 300 °C when comparing the SR sample with the one aged at 300 °C for
192h is thought to be caused by the reduction of Mn in solid solution that
is not counterbalanced by the nanoscale Zr-rich precipitates, as dis-
cussed previously.

However, the yield strength at 300 °C of the sample aged at 300 °C for 192h remains higher than that of the samples aged at 400 °C for 1 and 4h. As the main microstructural change between these samples is the globularization of the intermetallic network in the CZ, it has been concluded that preserving a continuous intermetallic network inside the CZ, representing about two thirds of the entire grain structure by volume

<figure><img src="image_19.png" /><figcaption>**Fig. 14.** STEM images in bright field taken in zone axis to reveal the dislocations in the gauge length of the deformed specimens at: (a,b) 200 °C and (c,d) 300 °C at 10⁻² s⁻¹ in the CZ of the SR condition. The elongated intermetallics in the CZ of the SR microstructure are pointed out with red arrows. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)</figcaption></figure>
fraction, is beneficial for elevated-temperature strength. This observation is consistent with other recent results in the literature indicating that elongated intermetallics exhibit superior mechanical properties at elevated temperatures compared to fragmented ones, see e.g. Refs. [2, 43]. Two mechanisms are often invoked in the literature to explain why elongated particles offer better resistance at high temperatures than globular particles. On the one hand, some authors establish analogies with metal matrix composites, proposing that load transfer from the matrix to the reinforcement increases in the presence of elongated particles compared to particulate reinforcements, see e.g. Ref. [43]. On the other hand, given the ultrafine size of such intermetallic particles, other studies rather suggest that their elongated morphology make dislocation motion more difficult [1,2], elongated particles being more difficult to be bypassed by dislocation climb compared to globular isolated particles.

In our opinion, even if load transfer from the matrix to the elongated particles does occur, this does not seem to be the most plausible explanation for the higher yield strength observed in the CZ in the presence of elongated intermetallic particles. The microstructure is more similar to a dispersion-strengthened material than to a metal matrix composite, in which the reinforcement consists of larger particles or fibres. Furthermore, elongated particles with a length no greater than the melt pool depth (approximately 200 µm) are only present in the CZ, accounting for around 60 % of the material. Here, we hypothesize that the beneficial role of elongated particles evidenced in this work (Fig. 6a–b) rather results from their ability to limit dislocation motion at high temperature. Indeed, such elongated particles would be more difficult to be bypassed by dislocation climb compared to globular isolated particles. Preliminary TEM observations conducted in the gauge length of tensile specimens deformed at 200 and 300 °C at a strain rate of 10⁻² s⁻¹ in the SR condition and shown in Fig. 14a–d suggest that dislocations are efficiently blocked by elongated intermetallic particles in the CZ. Further investigation would be required to consolidate this idea, but we believe it was worth reporting.

### 4.3. Recommendations for high strength at elevated temperatures in AM alloys

The present results show that the relative contribution of the different strengthening mechanisms varies with temperature. Consequently, the microstructure must be optimized depending on the operating temperature. This work yields two recommendations for achieving high strength at elevated temperatures: (i) preserving the elongated morphology of the intermetallic network and (ii) retaining a significant amount of slow-diffusing solutes that strongly interact with dislocations in the solid solution. Manganese (Mn) was found to be a good candidate because it is a slow diffuser in aluminium, has a high interaction energy with dislocations and can be retained in solid solution in significant amounts, well beyond its maximum equilibrium solubility, via solute trapping by leveraging the LB-PBF processing conditions.

Interestingly, the formation of thermally stable L1₂-Al₃Zr nanoscale precipitates has long been recognised as a way to increase the high-temperature resistance of castable Al alloys [32,40]. However, in the present study their contribution to elevated-temperature yield strength appears relatively limited compared to the other two mechanisms previously described, especially when their radius is < 2 nm because of thermally activated deformation mechanisms such as cross-slip and dislocation climb. Therefore, depending on the relative importance of intermetallics, solid solution and nanoscale precipitates, promoting the formation of nanoscale precipitates with a radius typically <2 nm via ageing heat treatment at 400 °C could be indirectly detrimental, as this process leads to the globularisation of elongated intermetallic particles and a decrease in Mn in the matrix. These two factors were identified as providing the greatest strengthening effect between 150 and 300 °C in the Al-4Mn-3Ni-2Cu-1Zr (wt.%) alloy.

The negative effect of the peak-ageing treatment on the high-

temperature strength evidenced in this work in the Al-4Mn-3Ni-2Cu-1Zr (wt.%) may not necessarily apply to all aluminium alloys designed for LB-PBF. It will depend on the nature of the different solutes, especially (i) their ability to be retained in solid solution in the aluminium matrix after a peak-ageing treatment, (ii) their ability to interact with dislocations and hinder dislocation climb, and (iii) their relative diffusivity relatively to that of Zr, which is consumed to form nanoscale precipitates upon ageing. However, the present results demonstrate that the heat treatment that leads to the best room-temperature mechanical properties is not necessarily the same treatment that leads to the best mechanical properties at elevated temperatures.

## 5. Conclusion

In this work, the influence of ageing on the high-temperature yield strength of the Al-4Mn-3Ni-2Cu-1Zr (wt.%) alloy was investigated and the links with the microstructural evolutions occurring upon ageing have been established. The main conclusions drawn from this work can be summarized as follows.

* The as-built microstructure remains stable after a stress-relief heat treatment at 300 °C for 4h. No major microstructural change was observed from the melt pool size down to the nanoscale (the solid solution composition is unchanged compared to the AB condition).

* At room temperature, the materials aged at 400 °C have a superior yield strength in comparison with their SR counterpart, as expected for an alloy strengthened by L1₂-Al₃Zr precipitation. However, at elevated temperatures, typically >150 °C, the SR microstructure outperforms the three studied ageing conditions. The reduction in yield strength with temperature is larger for the peak-aged microstructures compared to their stress-relieved counterpart (absence of nanoscale Zr-rich precipitates).

* Nanoscale L1₂-Al₃Zr precipitates (radii ranging between 1 and 1.5 nm) bring a substantial contribution to the room-temperature strength. However, their contribution to the high-temperature strengthening appears to be minor.

* The supersaturation of the solid solution in Mn and the elongated morphology of the network of intermetallics in the CZ in the SR condition are the two microstructural characteristics identified as being beneficial for maintaining a high yield strength at elevated temperatures in this LB-PBF-designed alloy.

The conclusions drawn from this study emphasize the importance of optimizing the heat treatment of Al alloys designed for AM depending on their operating temperature. It is also crucial to pay greater attention to the supersaturations resulting from out-of-equilibrium solidification conditions in LB-PBF, particularly with regard to solutes such as Mn that interact strongly with dislocations and can hinder dislocation climb.

### CRediT authorship contribution statement

**Tristan Lenoir:** Data curation, Formal analysis, Investigation, Methodology, Writing – original draft, Writing – review & editing.

**Béchir Chéhab:** Conceptualization, Funding acquisition, Project administration, Supervision, Validation, Writing – review & editing.

**Isabelle Mouton:** Data curation, Methodology, Visualization, Resources. **Florence Robaut:** Methodology, Resources. **Frédéric De Geuser:** Data curation, Investigation, Methodology, Validation, Writing – review & editing. **Baptiste Gault:** Data curation, Validation, Writing – review & editing. **Arthur Després:** Conceptualization, Investigation, Methodology, Supervision, Validation, Writing – review & editing.

**Jean-Jacques Blandin:** Conceptualization, Supervision, Validation, Writing – review & editing. **Guilhem Martin:** Conceptualization, Funding acquisition, Investigation, Project administration, Supervision, Validation, Writing – review & editing.
Declaration of competing interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

Acknowledgements

Tristan Lenoir acknowledges CONSTELLIUM for the ANRT n°2022/0760 funding in the framework of the CNRS joint laboratory between SIMaP and CONSTELLIUM C-TEC. The authors are also grateful to the Région Auvergne Rhône-Alpes research network Initiative3D. This work has benefited from the characterization equipment of the Grenoble INP - CMTC platform supported by the Centre of Excellence of Multifunctional Architectured Materials “CEMAM” noANR-10-LABX-44-01 funded by the Investments for the Future program. The authors are grateful for the kind support of Marti Lopez Freixes for the preparation and analyses of APT specimens at the then MPIE – now MPI SusMat. He and Baptiste Gault are grateful for Uwe Tezins & Andreas Sturm for their support of the FIB & APT facilities at the institute.

Appendix A. Supplementary data

Supplementary data to this article can be found online at https://doi.org/10.1016/j.msea.2025.149372.

Data availability

Data will be made available on request.

References

[1] R.A. Michi, K. Sisco, S. Bahl, Y. Yang, J.D. Poplawsky, L.F. Allard, R.R. Dehoff, A. Plotkowski, A. Shyam, A creep-resistant additively manufactured Al-Ce-Ni-Mn alloy, Acta Mater. 227 (2022) 117699, https://doi.org/10.1016/j.actamat.2022.117699.

[2] S. Bahl, T. Wu, R.A. Michi, K. An, D. Yu, L.F. Allard, J.U. Rakhmonov, J. D. Poplawsky, C.M. Fancher, D.C. Dunand, A. Plotkowski, A. Shyam, An additively manufactured near-eutectic Al-Ce-Ni-Mn-Zr alloy with high creep resistance, Acta Mater. 268 (2024) 119787, https://doi.org/10.1016/j.actamat.2024.119787.

[3] W. Wang, N. Takata, A. Suzuki, M. Kobashi, M. Kato, Design of Al-Fe-Mn alloy for both high-temperature strength and sufficient processability of laser powder bed fusion, Addit. Manuf. 68 (2023) 103524, https://doi.org/10.1016/j.addma.2023.103524.

[4] M.T. Pérez-Prado, A. Martin, D.F. Shi, S. Milenkovic, C.M. Cepeda-Jiménez, An Al-5Fe-6Cr alloy with outstanding high temperature mechanical behavior by laser powder bed fusion, Addit. Manuf. 55 (2022) 102828, https://doi.org/10.1016/j.addma.2022.102828.

[5] S. Bahl, K. Sisco, Y. Yang, F. Theska, S. Primig, L.F. Allard, R.A. Michi, C. Fancher, B. Stump, R. Dehoff, A. Shyam, A. Plotkowski, Al-Cu-Ce-Zr) alloys with an exceptional combination of additive processability and mechanical properties, Addit. Manuf. 48 (2021) 102404, https://doi.org/10.1016/j.addma.2021.102404.

[6] A. Shyam, A. Plotkowski, S. Bahl, K. Sisco, L.F. Allard, Y. Yang, J.A. Haynes, R. R. Dehoff, An additively manufactured AlCuMnZr alloy microstructure and tensile mechanical properties, Materialia 12 (2020) 100758, https://doi.org/10.1016/j.mtla.2020.100758.

[7] T. Kimura, T. Nakamoto, T. Ozaki, T. Miki, I. Murakami, Y. Hashizume, A. Tanaka, Microstructural development and aging behavior of Al-Cr-Zr heat-resistant alloy fabricated using laser powder bed fusion, J. Mater. Res. Technol. 15 (2021) 4193-4207, https://doi.org/10.1016/j.jmr2021.10.027.

[8] T. Wu, J.D. Poplawsky, L.F. Allard, A. Plotkowski, A. Shyam, D.C. Dunand, Microstructure and strengthening of Al-6Ce-3Ni-0.7Fe (wt%) alloy manufactured by laser powder-bed fusion, Addit. Manuf. 78 (2023) 103858, https://doi.org/10.1016/j.addma.2023.103858.

[9] Y. Yan, C. Lu, Z. Chen, Y. Zhuo, C. Wang, Q. Jia, Microstructures and high temperature tensile behaviours of laser powder bed fusion fabricated al-mn-sc alloy, Metals 13 (2023) 788, https://doi.org/10.3390/met13040788.

[10] M. Buttard, B. Chehab, R. Shahani, F. Robaut, G. Renou, C. Tassin, E. Rauch, P. Donnadieu, A. Deschamps, J.-J. Blandin, G. Martin, Multi-scale microstructural investigation of a new Al-Mn-Ni-Cu-Zr aluminium alloy processed by laser powder bed fusion, Materialia 18 (2021) 101160, https://doi.org/10.1016/j.mtla.2021.101160.

[11] R.A. Michi, A. Plotkowski, A. Shyam, R.R. Dehoff, S.S. Babu, Towards high-temperature applications of aluminium alloys enabled by additive manufacturing, Int. Mater. Rev. 67 (2022) 298-345, https://doi.org/10.1080/09506608.2021.1951580.

[12] J.H. Martin, B.D. Yahata, J.M. Hundley, J.A. Mayer, T.A. Schaedler, T.M. Pollock, 3D printing of high-strength aluminium alloys, Nature 549 (2017) 365-369, https://doi.org/10.1038/nature23894.

[13] G. Del Guercio, F. Bosio, C. Phutela, S. Robertson, N.T. Aboulkhair, Additive manufacturing of novel aluminium matrix composites with enhanced strength and processability via boron nitride functionalization, Addit. Manuf. Lett. 11 (2024) 100237, https://doi.org/10.1016/j.addlet.2024.100237.

[14] G. Li, B. Tunca, S. Senol, M. Casata, Y. Wu, Z. Chen, K. Vanmeensel, Revealing the precipitation behavior of crack-free TiB2/Al-Zn-Mg-Cu composites manufactured by laser Powder Bed fusion, Addit. Manuf. 66 (2023) 103460, https://doi.org/10.1016/j.addma.2023.103460.

[15] M. Buttard, M.L. Freixes, C. Josserond, P. Donnadieu, B. Chehab, J.-J. Blandin, B. Gault, F. De Geuser, G. Martin, Ageing response and strengthening mechanisms in a new Al-Mn-Ni-Cu-Zr alloy designed for laser powder bed fusion, Acta Mater. 259 (2023) 119271, https://doi.org/10.1016/j.actamat.2023.119271.

[16] J.R. Croteau, S. Griffiths, M.D. Rossell, C. Leinenbach, C. Kenel, V. Jansen, D. N. Seidman, D.C. Dunand, N.Q. Vo, Microstructure and mechanical properties of Al-Mg-Zr alloys processed by selective laser melting, Acta Mater. 153 (2018) 35-44, https://doi.org/10.1016/j.actamat.2018.04.053.

[17] R.A. Michi, K. Sisco, S. Bahl, L.F. Allard, K.B. Wagner, J.D. Poplawsky, D. N. Leonard, R.R. Dehoff, A. Plotkowski, A. Shyam, Microstructural evolution and strengthening mechanisms in a heat-treated additively manufactured Al-Cu-Mn-Zr alloy, Mater. Sci. Eng. A 840 (2022) 142928, https://doi.org/10.1016/j.msea.2022.142928.

[18] C. Pauzon, M. Buttard, A. Després, F. Charlot, M. Fivel, B. Chehab, J.-J. Blandin, G. Martin, Direct ageing of LPBF Al-1Fe-1Zr for high conductivity and mechanical performance, Acta Mater. 258 (2023) 119199, https://doi.org/10.1016/j.actamat.2023.119199.

[19] S. Griffiths, J.R. Croteau, M.D. Rossell, P. Erni, A. De Luca, N.Q. Vo, D.C. Dunand, C. Leinenbach, Coarsening- and creep resistance of precipitation-strengthened Al-Mg-Zr alloys processed by selective laser melting, Acta Mater. 188 (2020) 192-202, https://doi.org/10.1016/j.actamat.2020.02.008.

[20] M. Buttard, B. Chehab, C. Josserond, F. Charlot, P. Lhuissier, X. Bataillon, A. Deschamps, J. Villanova, M. Fivel, J.-J. Blandin, G. Martin, Influence of microstructure heterogeneity on the tensile response of an aluminium alloy designed for laser powder bed fusion, Acta Mater. 269 (2024) 119786, https://doi.org/10.1016/j.actamat.2024.119786.

[21] B. Mehta, K. Frisk, L. Nyborg, Advancing novel Al-Mn-Cr-Zr based family of alloys tailored for powder bed fusion-laser beam process, J. Alloys Compd. 967 (2023) 171685, https://doi.org/10.1016/j.jallcom.2023.171685.

[22] B. Mehta, K. Frisk, L. Nyborg, Effect of precipitation kinetics on microstructure and properties of novel Al-Mn-Cr-Zr based alloys developed for powder bed fusion - laser beam process, J. Alloys Compd. 920 (2022) 165870, https://doi.org/10.1016/j.jallcom.2022.165870.

[23] B. Mehta, S. Bengtsson, D. Riabov, E. Natesan, K. Frisk, J. Ahlström, L. Nyborg, Mechanical properties and microstructural characterisation including high-temperature performance of Al-Mn-Cr-Zr based alloys tailored for additive manufacturing, Mater. Des. 244 (2024) 113160, https://doi.org/10.1016/j.matdes.2024.113160.

[24] D. Bayoumy, K. Kwak, T. Boll, S. Dietrich, D. Schliephake, J. Huang, J. Yi, K. Takashima, X. Wu, Y. Zhu, A. Huang, Origin of non-uniform plasticity in a high-strength Al-Mn-Sc based alloy produced by laser powder bed fusion, J. Mater. Sci. Technol. 103 (2022) 121-133, https://doi.org/10.1016/j.jmst.2021.06.042.

[25] Q. Jia, P. Rometsch, P. Kürsteiner, Q. Chao, A. Huang, M. Weyland, L. Bourgeois, X. Wu, Selective laser melting of a high strength Al Mn Sc alloy: alloy design and strengthening mechanisms, Acta Mater. 171 (2019) 108-118, https://doi.org/10.1016/j.actamat.2019.04.014.

[26] Q. Jia, Y. Zhuo, Y. Yan, C. Lu, Z. Chen, Y. Cheng, C. Wang, Tensile creep mechanisms of Al-Mn-Sc alloy fabricated by additive manufacturing, Addit. Manuf. 79 (2024) 103910, https://doi.org/10.1016/j.addma.2023.103910.

[27] E.F. Rauch, M. Véron, Automated crystal orientation and phase mapping in TEM, Mater. Char. 98 (2014) 1-9, https://doi.org/10.1016/j.matchar.2014.08.010.

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

[29] M. Genc, P. Eloi, J.-J. Blandin, C. Pascal, P. Donnadieu, F. De Geuser, P. Lhuissier, C. Desrayaud, G. Martin, Optimization of the strength vs. conductivity trade-off in an aluminium alloy designed for laser powder bed fusion, Mater. Sci. Eng. A 858 (2022) 144139, https://doi.org/10.1016/j.msea.2022.144139.

[30] F. De Geuser, B. Gault, Metrology of small particles and solute clusters by atom probe tomography, Acta Mater. 188 (2020) 406-415, https://doi.org/10.1016/j.actamat.2020.02.023.

[31] M. Buttard, G. Martin, X. Bataillon, G. Renou, P. Lhuissier, J. Villanova, B. Chehab, P. Jarry, J.-J. Blandin, P. Donnadieu, Towards an alloy design strategy by tuning liquid local ordering: what solidification of an Al-alloy designed for laser powder bed fusion teaches Us, Addit. Manuf. 61 (2023) 103313, https://doi.org/10.1016/j.addma.2022.103313.

[32] K.E. Knipling, D.C. Dunand, D.N. Seidman, Criteria for developing castable, creep-resistant aluminium-based alloys - a review, Z Met (2006), https://doi.org/10.1515/jmr-2006-0042.

[33] A. Chouhan, L. Mådler, N. Ellendt, Modeling of rapid solidification in laser powder Bed fusion processes, Comput. Mater. Sci. 238 (2024) 112918, https://doi.org/10.1016/j.commatsci.2024.112918.

[34] X. Tang, H. Zhang, P. Xue, L. Wu, J. Zhang, L. Zhang, F. Liu, D. Ni, B. Xiao, Z. Ma, Ultrahigh strength heat-resistant Al-Fe-V-Si-Sc alloy fabricated by laser powder bed
fusion, J. Mater. Sci. Technol. 239 (2025) 299-306, https://doi.org/10.1016/j.jmst.2025.02.060.
[35] T. Lenoir, M. Buttard, B. Chehab, S. Coindeau, A. Després, J.-J. Blandin, G. Martin, Influence of the microstructure on the high temperature ductility of the Al-4Mn-3Ni-2Cu-1Zr aluminium alloy designed for laser powder bed fusion, Mater. Sci. Eng. A 940 (2025) 148559, https://doi.org/10.1016/j.msea.2025.148559.
[36] O.C. Hellman, J.A. Vandenbroucke, J. Rüsing, D. Isheim, D.N. Seidman, Analysis of three-dimensional atom-probe data by the proximity histogram, Microsc. Microanal. 6 (2000) 437-444, https://doi.org/10.1007/S100050010051.
[37] F.C. Liu, Z.Y. Ma, Contribution of grain boundary sliding in low-temperature superplasticity of ultrafine-grained aluminum alloys, Scr. Mater. 62 (2010) 125-128, https://doi.org/10.1016/j.scriptamat.2009.10.010.
[38] Z.Y. Ma, F.C. Liu, R.S. Mishra, Superplastic deformation mechanism of an ultrafine-grained aluminum alloy produced by friction stir processing, Acta Mater. 58 (2010) 4693-4704, https://doi.org/10.1016/j.actamat.2010.05.003.

[39] C.N. Ekaputra, J.U. Rakhmonov, E. Senvardarli, D. Weiss, J.-E. Mogonye, D. C. Dunand, Combining solution-, precipitation- and load-transfer strengthening in a cast Al-Ce-Mn-Sc-Zr alloy, Acta Mater. 266 (2024) 119683, https://doi.org/10.1016/j.actamat.2024.119683.
[40] C.B. Fuller, D.N. Seidman, D.C. Dunand, Mechanical properties of Al(Sc,Zr) alloys at ambient and elevated temperatures, Acta Mater. 51 (2003) 4803-4814, https://doi.org/10.1016/S1359-6454(03)00320-3.
[41] A.R. Farkoosh, Solute-induced strengthening during creep of an aged-hardened Al-Mn-Zr alloy, Acta Mater. (2021).
[42] S. Vannarat, M.H.F. Sluiter, Y. Kawazoe, First-principles study of solute-dislocation interaction in aluminum-rich alloys, Phys. Rev. B 64 (2001) 224203, https://doi.org/10.1103/PhysRevB.64.224203.
[43] A.F. Whitehouse, Creep of metal matrix composites, in: Compr. Compos. Mater., 2000, pp. 371-418, https://doi.org/10.1016/B0-08-042993-9/00014-0.