# Effect of a Zr source addition on the microstructure of Al-SiC composites elaborated by the Laser Powder Bed Fusion (L-PBF) process

Marie-Reine Manlay, Camille Flament, Stéphane Gossé, Mathieu Soulier, Jean-Paul Garandet

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Marie-Reine Manlay, Camille Flament, Stéphane Gossé, Mathieu Soulier, Jean-Paul Garandet. Effect of a Zr source addition on the microstructure of Al-SiC composites elaborated by the Laser Powder Bed Fusion (L-PBF) process. Materials Characterization, 2024, 218, pp.114472. 10.1016/j.matchar.2024.114472. hal-04753052

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# Effect of a Zr source addition on the microstructure of Al-SiC composites elaborated by the Laser Powder Bed Fusion (L-PBF) process

Marie-Reine Manlay<sup>a</sup>, Camille Flament<sup>a</sup>, Stéphane Gossé<sup>b</sup>, Mathieu Soulier<sup>a</sup>, Jean-Paul Garandet<sup>a,*</sup>

## ARTICLE INFO

### Keywords:

SiC decomposition

Nucleation

Microstructure

L-PBF

Transmission electron microscopy

Calphad

## ABSTRACT

The present work focuses on the effect of the addition of a zirconium source on the microstructure of Al-SiC composites produced by Laser Powder Bed Fusion (L-PBF). More specifically, the aim is to address the issue of the SiC decomposition into the water-soluble aluminum carbide Al₄C₃ in Al-SiC composites produced by L-PBF, with the objective of limiting its formation by adding another element to the system. To this end, AlSi₇Mg₀.₆-SiC-ZrO₂ composite powders are successfully prepared and printed in a standard L-PBF equipment. The resulting parts are then thoroughly characterized, in order to understand the physico-chemical mechanisms involved during the L-PBF process. The results show a decrease in the Al₄C₃ amount by ZrC formation. Another important result is that bulk composites exhibit a fully equiaxed microstructure attributed to the τ₁ (Al,Si)₃Zr ternary phase, with all the characteristics of a good nucleating agent for aluminum phase. To support these microstructure experimental results, a first version of a quaternary Al-Zr-Si-C thermodynamic database was developed using the Calphad method. These calculations enable to establish a solidification path providing information on the phases that may form after heat treatment of L-PBF materials.

## 1. Introduction

Laser Powder Bed Fusion (L-PBF) is one of the most mature methods in the rapidly developing field of metal additive manufacturing. It consists in the repetition of the selective laser fusion and solidification of metal powder beds according to a pre-defined scanning path, until the desired part is grown. This technique addresses the growing demand for high-density metal parts with complex geometries, fine and controlled microstructures. A major scientific and technical challenge, particularly in the transportation industry, is to reduce the weight of structures while maintaining the functional properties of parts, such as mechanical properties. In this respect, aluminum alloys are known to offer an excellent compromise between lightness and mechanical strength. The L-PBF process fits into this context of lightening structures by further topological optimization of parts. Moreover, many aluminum alloys, especially those with high silicon content, are well adapted for the implementation of this technique, since they are not prone to hot cracking issues.

According to the target application, it may be necessary to mechanically strengthen aluminum alloys by introducing reinforcements to form metal matrix composites. Aluminum matrix composites reinforced with silicon carbide particles are widely used in industrial applications, particularly in the automotive, aerospace and defense sectors [1]. They are attractive because of their low density, high mechanical strength and Young modulus, and corrosion resistance [2]. However, during the conventional production of these composites by stir-casting or infiltration, it is known that a chemical reaction can occur above 660 °C between the liquid aluminum and the SiC reinforcements, leading to the formation of an aluminum carbide, Al₄C₃, according to the following reaction [3].

$$4 \text{Al} + 3 \text{SiC} \rightarrow \text{Al}_4\text{C}_3 + 3 \text{Si} \quad (1)$$

This carbide alters the mechanical properties and parts lifetime, due to its brittleness and tendency to hydrolyze [4]. In casting processes, a strict control of the temperature cycle and of the system chemistry allow to limit SiC decomposition. Unfortunately, the much higher

<footnote><sup>a</sup> Univ. Grenoble Alpes, CEA, LITEN, DTNM, F-38000 Grenoble, France</footnote>

<footnote><sup>b</sup> Université Paris-Saclay, CEA, Service de Recherche en Corrosion et Comportement des Matériaux, 91191 Gif-sur-Yvette, France</footnote>

<footnote>* Corresponding author.</footnote>

<footnote>E-mail addresses: camille.flament@cea.fr (C. Flament), stephane.gosse@cea.fr (S. Gossé), mathieu.soulier@cea.fr (M. Soulier), jean-paul.garandet@cea.fr (J.-P. Garandet).</footnote>
temperatures associated with L-PBF prevent these strategies from being effective. On the other hand, the kinetics of this process being much faster than those of casting or infiltration, our previous work investigated the issue of SiC decomposition in an AlSi<sub>7</sub>Mg<sub>0.6</sub>-SiC composite produced by L-PBF, with the aim of estimating whether this decomposition could be limited by the kinetics of the process [5]. Through a coupled experimental and numerical analysis, this previous study showed that submicron SiC particles could not survive the L-PBF process, resulting in the formation of Al<sub>4</sub>C<sub>3</sub>.

In this context, the objective of the present work is to develop an alternative approach to counteract the negative effects of SiC decomposition in Al-SiC composites produced by L-PBF. This consists in adding an element to an AlSi<sub>7</sub>Mg<sub>0.6</sub>-SiC composite powder to form another carbide than Al<sub>4</sub>C<sub>3</sub>. Based on a binary carbide stability diagram [6], our choice was zirconium because ZrC has a more negative Gibbs free energy than Al<sub>4</sub>C<sub>3</sub>. Although such stability diagrams of binary carbides do not take into account the ternary interactions of these carbides, they allow to identify the most energetically stable carbides, and can be used for purposes of first selection based on thermodynamic criteria. At this point, it should be stated that the Al-Zr-C ternary system studied by Schuster and Nowotny [7] indicates the absence of equilibrium between liquid aluminum and ZrC at 700 °C and 1000 °C. However, composition conditions can be locally satisfied to allow the formation of ZrC.

Indeed, Long et al. [6] showed that the mechanical properties reduction of a laser weld joint of two Al-SiC composite parts containing 10 vol% SiC, due to the expected degradation of SiC into Al<sub>4</sub>C<sub>3</sub>, could be alleviated by the addition of a 50 µm-thick zirconium foil between the two composite plates. With such an approach, the amount of porosity in the joint was reduced, and its mechanical strength increased. This effect is associated with the absence of Al<sub>4</sub>C<sub>3</sub> precipitates and the formation of ZrC precipitates of about 5–10 µm in size, allowed by the local composition conditions around the zirconium foil, according to the following reaction:

$$ \mathrm{SiC} + \mathrm{Zr} \rightarrow \mathrm{ZrC} + \mathrm{Si} \quad (2) $$

Therefore, the results of Long et al. [6] represented an interesting proof of concept of the possibility to limit the amount of Al<sub>4</sub>C<sub>3</sub> by the in situ formation of ZrC during the process. Our objective in the present work is to adapt this strategy of Zr addition to the elaboration of Al-SiC based composites by L-PBF. In order to satisfy the objective of promoting the in situ formation of ZrC during the process in order to avoid the formation of Al<sub>4</sub>C<sub>3</sub>, it is necessary to introduce into the system a molar fraction of zirconium at least equal to the molar fraction of carbon from silicon carbide. Moreover, it could be profitable to introduce a molar excess of Zr, which could promote the formation of Al<sub>3</sub>Zr, known to be an excellent nucleating agent in aluminum alloys [8–11].

Regarding the paper organization, the materials and methods used to perform this research work are detailed in Section 2. The results obtained are then presented and analyzed in Section 3. Finally, Section 4 is dedicated to a discussion of the findings, providing clues regarding the physico-chemistry of a new composite elaborated by L-PBF.

## 2. Materials and methods

### 2.1. Raw materials

The host material used in this study is gas atomized spherical AlSi<sub>7</sub>Mg<sub>0.6</sub> powder acquired from TLS Technik GmbH & Co. The particles have a range size of 15–53 µm and a median diameter of 29 µm.

<table><caption>**Table 1**
Chemical composition of the AlSi<sub>7</sub>Mg<sub>0.6</sub> alloy powder (wt%).</caption><thead><tr><th>Composition</th><th>Al</th><th>Si</th><th>Mg</th><th>Fe</th><th>Cu</th><th>Zn</th><th>Ti</th><th>Mn</th></tr></thead><tbody><tr><td>Specification</td><td>Bal.</td><td>6.50–7.50</td><td>0.45–0.70</td><td>≤0.19</td><td>≤0.05</td><td>≤0.07</td><td>0.08–0.25</td><td>≤0.10</td></tr><tr><td>Powder</td><td>Bal.</td><td>7.23</td><td>0.62</td><td>0.18</td><td>0.005</td><td>&lt;0.001</td><td>0.18</td><td>0.003</td></tr></tbody></table>

Their chemical composition is reported in Table 1. The powder density, as measured by Helium pycnometry, is 2.66 g·cm⁻³.

As mentioned in the introduction, the aim of the present work is to add a zirconium source to an AlSi<sub>7</sub>Mg<sub>0.6</sub>-SiC composite powder. To this end, SiC nanoparticles with a mean size of 35 nm are used. These nanoparticles, provided by Nanomakers, are produced by laser pyrolysis.

Several zirconium sources can be selected to meet the desired objectives. In line with a previous study based on Al 6061 powders [8], a nanosized ZrO<sub>2</sub> powder with a characteristic particle size of 80 nm supplied by US Research Nanomaterials, was selected for its large availability and easy handling in order to produce an AlSi<sub>7</sub>Mg<sub>0.6</sub>-SiC-ZrO<sub>2</sub> ternary composite powder. Since the objective in the present work is to clarify complex physico-chemistry issues rather than to produce application-oriented materials, we decided to set the total content of reinforcements at 5 wt% in order to form a significant density of precipitates, making it easier the analysis of X-ray diffractograms. Two mixtures with different ratios R = [Zr]/[C] were prepared, one with R = 1 and the other with an excess of Zr, with R = 3. Their composition is presented in Table 2. These composite powders will be respectively named R1 and R3 afterwards.

### 2.2. Nanocomposite powders elaboration

The aim of the mixing step is to perform dry coating of nano-sized reinforcements on matrix particles. Dry coating is based on an electrostatic assembly of nanoparticles at the surface of micrometric powders [14]. The potential of this approach to develop nanocomposite powders for the L-PBF process has already been demonstrated in a previous work [15]. Composite powders mixtures with 5 wt% reinforcements are produced using a 5 L capacity Cyclomix high-shear mixer, with a rotation speed of 15 m·s⁻¹ and a mixing duration of 30 min.

Composite powders properties are investigated through several characterization techniques. Powders granulometry is estimated using static image analysis (Malvern Morphologi G3), allowing to give information on the electrostatic mixture stability. The analysis is performed on 150,000 particles at a x20 magnification with a detection limit for particles smaller than 1 µm. The quality of electrostatic assembly of powders is also evaluated on secondary electron images taken from Scanning Electron Microscopy (SEM) using a Field Emission Gun (FEG) SEM Zeiss Merlin. For the determination of mass balances, Instrumental Gas Analysis (IGA) and Inductively Coupled Plasma (ICP) analysis, performed by Elektrowerk Weisweiler GmbH, were also used. IGA measurements for carbon are performed on 0.10–0.40 g samples, while ICP analysis are performed on 1 g samples.

### 2.3. Laser Powder Bed Fusion (L-PBF) process

Bulk composites are produced by Laser Powder Bed Fusion with a SLM Solutions machine (SLM 125 HL), equipped with a 400 W laser and

<table><caption>**Table 2**
Composition of composite powder mixtures.</caption><thead><tr><th>R = [Zr]/[C]</th><th>R = 1</th><th>R = 3</th></tr></thead><tbody><tr><td>Composite name</td><td>R1</td><td>R3</td></tr><tr><td>SiC content</td><td>1.23 wt%</td><td>0.49 wt%</td></tr><tr><td>ZrO<sub>2</sub> content</td><td>3.77 wt%</td><td>4.51 wt%</td></tr></tbody></table>
a spot diameter of 70 µm, using the same procedure and the same pro-
cess parameters than our previous work [5].

Based on these parameters, samples are identified according to their
Energy per Built Volume (EBV) [16]. The density of bulk composites is
then measured using the standard Archimedes' method in ethanol. In
order to study the physico-chemistry of the composites produced, their
microstructure is thoroughly characterized using a variety of methods.
The different phases constituted the composite are identified by means
of X-Ray Diffraction (XRD) (Bruker D8 Advance) with a variation of 2θ
angle from 28° to 42°. Such a relatively restricted range was selected for
time saving purposes. Nevertheless, we checked using larger acquisition
intervals that it was adequate to identify the relevant phases within our
samples. At an intermediate spatial scale, grains morphology and
orientation are characterized tanks to SEM and Electron Backscatter
Diffraction (EBSD). Energy Dispersive X-rays spectroscopy (EDX) ana-
lyses are performed to characterize precipitation. For SEM analyses, the
electron acceleration voltage is set at 5 kV and the current at 1 nA for
images and EDX analyses. EBSD maps are acquired by tilting the sample
by 70° and applying a voltage of 20 kV and a current of 7 nA, with a pixel
size of 100 nm. For a more detailed characterization, Transmission
Electron Microscopy (TEM) analyses are performed using a FEG Tecnai
Osiris from Thermofisher operating at 200 kV. The images are recorded
by conventional TEM, or by S-TEM with High Angle Annular Dark Field
(HAADF) and Bright Field (BF) detectors. EDX data are acquired in
STEM mode with the super-X detector equipped with four Silicon Drift
Detectors (SSD). Data are treated with Esprit 1.9 from Bruker Phase and
orientation maps are performed by recording the diffraction pattern at
each point with a camera thanks to ACOM-TEM. These patterns are
automatically indexing using ASTAR software from NanoMEGAS. For
details on this technique, the interested reader is referred to references
[17, 18]. Acquisitions are conducted with a precessed beam with an
angle of 1.2°. Beam size is of about 1 nm and 5 nm pixel. Samples for
TEM analyses are prepared using a cross-beam Focused Ion Beam 550
from Zeiss (FIB). The characterized TEM lamellae are lifted out ac-
cording to the perpendicular plane to building direction.

2.4. Thermodynamic assessment of the quaternary Al-Zr-Si-C system

In order to relate the experimental results obtained with thermody-
namics, a first version of a thermodynamic database of the quaternary
Al-Zr-Si-C system was developed using the Calphad method and
Thermo-Calc package (2022b version). Although other chemical ele-
ments are present in the system, we selected to consider only the in-
teractions between major elements Al, Si, Zr and C. The Calphad method
allows to calculate phase diagrams and thermodynamic properties from
the description of the Gibbs energy functions of all phases (solid, liquid
and gas) in a given chemical system, and to make predictive calculations
to investigate as yet unstudied temperature or composition domains.

The quaternary description is based six binary and four ternary sub-
systems, i.e. the Al—C, Al—Si, Al—Zr, C—Si, C—Zr and Si—Zr binaries
and the Al-C-Si, Al-C-Zr, Al-Si-Zr and C-Si-Zr ternaries. Among these
systems, all were already thermodynamically assessed but Al-C-Zr is
poorly described experimentally. The description of the ternary Al-C-Si
and sub-systems comes from Gröbner et al. [19]. This assessment con-
siders two ternary intermetallics: Al₄C₄Si and Al₈C₇Si; both these phases
are stoichiometric. The Al—Zr, Si—Zr and Al-Si-Zr thermodynamic de-
scriptions come from Hirano et al. [12]. This study considers three
ternary intermetallics noted τ₁, τ₂ & τ₃. They correspond to both line
compounds (Al,Si)₃Zr and (Al,Si)Zr and to the stoichiometric phase
Al₄Si₅Zr₃. The C—Zr thermodynamic model is taken from Fernandez
Guillermet [20]. No particular excess parameters were considered in the
ternary C-Si-Zr.

The Al-C-Zr system is poorly described experimentally. Because no
Calphad model exists for this ternary, a specific database was developed
from limited information from the literature. The ternary isothermal
sections were established from the diagrams by Schuster and Nowotny

[7] who identified three intermetallics: Zr<sub>5</sub>Al<sub>3</sub>C, ZrAlC<sub>2-x</sub>, and Zr<sub>2</sub>Al<sub>3</sub>C<sub>5-x</sub>. Due to the lack of data, these two last phases were considered stoichiometric in the Calphad model. The formation enthalpy of these MAX phases were taken from DFT data of Poulou et al. [21]. No supplementary Al-C-Zr ternary phases were considered.

As the thermodynamic properties and stability domains of these Al-
Zr-C ternary phases are not well known, this ternary system needs to be
further investigated and considered with caution in view of the limited
data available. This is also the case for the Al-Zr-Si ternary system,
calculated from the Hirano et al. database [12], in which the description
of ternary intermetallics appears to be only based on experimental re-
sults at 700 °C. On the other hand, the thermodynamic modeling of the
other two ternary systems Al-Si-C [19] and Si-Zr-C [22] are considered
very satisfying when compared to the experimental data.

It is important to keep in mind that the AlSi<sub>7</sub>Mg<sub>0.6</sub> matrix used in this work has a high silicon content, which could lead to the formation of the τ₁ ternary line compound with (Al,Si)<sub>3</sub>Zr composition, rather than Al<sub>3</sub>Zr [12,13]. In the quaternary system Al-Si-Zr-C, anticipating on the focus that will be on the τ₁ and ZrC phases, as well as their interaction with liquid aluminum, the ternary diagrams of interest are Al-Si-Zr and Al-Zr-C. Fig. 1 shows the isothermal cross-sections of Al-Zr-C and Al-Si-Zr ternary diagrams calculated at 1000 °C. The quaternary system is then extrapolated from the thermodynamic description of the four ternary subsystems.

Considering the isothermal section of the Al-Zr-C ternary diagram, it
is interesting to note that among the ternary phases, two compounds,
Zr₂Al₃C₅ and ZrAlC₂, are in equilibrium with the liquid phase,
depending on its composition. The carbide ZrC₁₋ₓ is also in equilibrium
with these ternary phases, depending on its stoichiometry, and zirco-
nium aluminides, including Al₃Zr.

Considering the Al-Zr-Si isothermal cross-section, the presence of τ₁ and τ₂ ternary phases can be noted at 1000 °C, respectively defined as Zr (Al₁ₓSiₓ)₃ with 0.1 ≤ x ≤ 0.4 and Zr(Al₇Si₁₁ₓ) with 0.2 ≤ x ≤ 0.25 [13]. At this temperature, the main phases in equilibrium with the liquid phase, according to the liquid composition, are Al₃Zr, τ₁ and a Si₂(Al,Zr) solid solution. This τ₁ phase is notably in equilibrium with the Al₃Zr phase, as well as τ₂, depending on the composition of τ₁.

2.5. *Heat treatments*

Thermal analyses are also performed to complement the calculations
from the Calphad database, and to make the link with the results ob-
tained after L-PBF. Differential thermal analysis (DTA) tests are used to
estimate the formation temperatures of the various compounds. These
measurements are carried out using a SETARAM TAG24 DTA, with a
temperature ramp from 10 °C·min⁻¹ from the ambient up to 1000 °C and
an argon flow rate of 30 L·h⁻¹.

Moreover, heat treatment tests are carried out on composite powders
in a Nabertherm tube furnace under argon flow. The temperature rise
and cooling ramp is 250 °C·h⁻¹, and the sample is held at working
temperature, i.e. 650 °C and 800 °C, for one hour.

3. Results

3.1. *Composite powders characterization*

Before focusing on the physico-chemistry of ternary AlSi<sub>7</sub>Mg<sub>0.6</sub>-SiC-
ZrO<sub>2</sub> composites produced by L-PBF, it is necessary to characterize the
electrostatic assembly of the SiC and ZrO<sub>2</sub> nanoparticles which ensures
the stability and the homogeneity of the powder mixture.

The methodology used to characterize composite powders will not be
detailed here. Further information are given in our previous work [5].
The aim is to ensure that the mixture is of sufficient quality for subse-
quent physico-chemical studies on composite parts. SEM images provide
a first qualitative approach to qualify the electrostatic assembly. On the
SEM images of R1 and R3 powders shown in the Supplementary data 1,
<figure><img src="image_3.png" /><figcaption>Fig. 1. (a) Al-Zr-C and (b) Al-Si-Zr isothermal cross-sections at 1000 °C.</figcaption></figure>

the nanoparticles appear to completely cover the surface of the aluminum particles, as evidenced by their fluffy surface appearance. No segregation of free nanoparticles is observed in the powder samples analyzed by SEM. To complete these results, granulometry results show a monomodal size distribution of the composite powders, as that of the matrix powder, indicating a complete electrostatic assembly (Supplementary data 2). Indeed, if this were not the case, the size distribution of the powder would be modified by the presence in the mixture of SiC and ZrO₂ ungrafted particles, self-agglomerated because of their high specific area, detectable by a second peak.

Table 3 displays IGA and ICP measurements on C and Zr elements in the composite powders, compared with their theoretical quantity introduced into the mixture.

The values presented for the R1 composite powder correspond to an average value obtained from three samples taken from a 250 mL powder volume, while a single sample was analyzed for the R3 composite. From the three measurements carried out on the R1 composite powder, standard deviations in absolute values of respectively of 0.01 wt% and 0.08 wt% were obtained for the elements C and Zr. As such a dispersion is of the order of the error linked to the measurement techniques themselves, which are 0.01 pds% for carbon and 0.06 pds% for zirconium, it is not possible to highlight any differences that might be due to the heterogeneous distribution of SiC and ZrO₂ powders. In conclusion, the composite powders produced appear suitable for investigating the fate of these particles after L-PBF.

## 3.2. Microstructure characterization

Density measurements using the Archimedes' method (Supplementary data 3) allow to select the sample printed at a EBV of 32.6 J·mm⁻³ among the densest for the subsequent microstructural study, measured at 2.53 g·cm⁻³. Since no relevant reference density values can be obtained in such complex multiphase systems, the results are presented in absolute, i.e. g·cm⁻³. The porosity rate is then estimated by optical

<table><caption>**Table 3**
Analysis of carbon and zirconium content in the composite powders compared to their expected content.</caption><thead><tr><th rowspan="2">Composite<br>Element</th><th colspan="2">R1</th><th colspan="2">R3</th></tr><tr><th>C</th><th>Zr</th><th>C</th><th>Zr</th></tr></thead><tbody><tr><td>Powder (wt%)</td><td>0.35</td><td>2.30</td><td>0.12</td><td>3.14</td></tr><tr><td>Theory (wt%)</td><td>0.37</td><td>2.79</td><td>0.15</td><td>3.38</td></tr></tbody></table>

microscopy on polished sections of samples R1 and R3, by image binarization on an analysis area of 13.5 mm². This analysis indicates a porosity rate of around 6 % in both composites. Numerous porosities can be observed, both irregular and spherical in shape. However, although the composites would require process optimization to reduce the porosity, the microstructure seems to be similar between the porosities, allowing to carry out the study of the physicochemistry of composites on these areas.

According to EBSD analysis of Al-SiC bulk composites from our previous study, the microstructure of Al-SiC composites elaborated by L-PBF mainly consists of columnar grains textured (100) in the building direction (Supplementary data 4). On the contrary, R1 and R3 composites with the addition of ZrO₂ display a completely equiaxed microstructure, as shown in Fig. 2. This figure also shows the presence of numerous intragranular and intergranular Zr-rich precipitates. Compared with Al-SiC composites and their columnar grains, the microstructure is highly refined, with a linear grain size of around 1 µm. The formation of this totally equiaxed microstructure may implies the presence of a nucleating agent for the aluminum phase. The goal now is to identify the precipitation in the bulk samples, and to identify the germinating agent involved.

Fig. 3 displays the X-ray diffractograms obtained on the bulk composites R1 and R3 obtained at a EBV of 32.6 J·mm⁻³.

The diffractograms of composites R1 and R3 both exhibit peaks characteristic of the diffraction of the ZrC phase. This result reflects the successful in situ synthesis of ZrC during the L-PBF thermal cycle undergone by the material. The objective of trapping the released carbon with zirconium has thus been at least partially achieved. However, the absence of peaks characteristic of Al₄C₃ diffraction does not certify the absence of Al₄C₃ in the material, as it can be present in insufficient amounts to be detected. Nevertheless, the detection of the ZrC phase necessarily implies a reduction in Al₄C₃ formation. Both composites also show peaks associated with the diffraction of the ternary Al-Zr-Si τ₁ phase.

The τ₁ phase appears to have been formed preferentially to the Al₃Zr phase, as this phase widely documented in the literature is undetected and not observed by electron microscopy. This appears to be due to the high silicon content present in the AlSi₇Mg₀.₆ matrix, since similar tests carried out with pure Al or Al 6061 matrix show Al₃Zr precipitation [23]. This result was also highlighted by Gao et al. [24], although in another manufacturing process. Furthermore, the fact that the τ₁ phase is identified in the R1 composite, which contains as much zirconium as carbon, means that there are less zirconium atoms available to form ZrC,
<figure><img src="image_4.png" /><figcaption>Fig. 2. (a) SEM-BSE image of R1 microstructure, (b) SEM-BSE image of R3 microstructure, (c) EDX analysis for Zr in R3 composite.</figcaption></figure>

<figure><img src="image_5.png" /><figcaption>Fig. 3. Composites diffractograms at an EBV of 32.6 J·mm⁻³.</figcaption></figure>

so the free carbon has probably recombined with aluminum to form Al₄C₃. The TEM analysis shown in Fig. 4 reveals ZrC precipitates by STEM-EDX and electron diffraction, with interreticular distance measurements coinciding with the (2,0,-2), (0,2,-2) et (-2,2,0) planes of the face-centered cubic structure (space group Fm-3 m) of ZrC, with lattice parameter a = 4.68 Å [25]. These precipitates form a necklace of intergranular agglomerated spherical-like nanoprecipitates. The agglomeration of these precipitates supports the hypothesis that only local compositional conditions can enable the formation of the ZrC phase during thermal cycling.

These analyses also reveal residual Al₄C₃ precipitates, with the detection of a carbon-rich precipitate with no Zr or Si signature, as well as both intergranular and intragranular τ₁ phase precipitates, with a morphology mostly in the form of platelets.

The aim now is to identify which of these precipitates is responsible for aluminum germination. As the ZrC phase precipitates as agglomerates, with a heterogeneous and mostly intergranular repartition throughout the material, these precipitates are therefore unlikely to be responsible for germination. On the other hand, as TEM analysis has shown the presence of τ₁ intragranular precipitates, further analysis is carried out to investigate whether the precipitate presents all the criteria for being a good nucleating agent [26,27], starting with a potential orientation relationship between them and the matrix.

This investigation of an orientation relationship is based on a

<figure><img src="image_6.png" /><figcaption>Fig. 4. (a) STEM-HAADF image of precipitates in R1 composite, (b)-(e) STEM-EDX analysis of the same area, (f) Diffraction pattern of a ZrC precipitate oriented along [111] zone axis in the same zone.</figcaption></figure>

comparison of the electron diffraction pattern of the precipitate with that of the matrix surrounding it, using ASTAR [17,18]. Automatic indexing of the diffraction patterns revealed an orientation relationship between the matrix and an intragranular τ₁ precipitate in R3 composite. Fig. 5 provides the EDX analysis of this precipitate, displaying the Si and Zr chemical signatures. The diffraction pattern indexation of the precipitate enables the identification of the (0,-2,0), (-3,-1,2) et (-3,1,2) planes of the quadratic structure of τ₁ phase (space group I4/mmm), whose orientation is close to the [203] zone axis. On the other hand, the face-centered cubic structure of aluminum is oriented close to the [103] zone axis.
<figure><img src="image_7.png" /><figcaption>Fig. 5. STEM-EDX analysis of a τ₁ precipitate in R3 composite, with diffraction pattern of the matrix surrounding the precipitate and the τ₁ precipitate.</figcaption></figure>

According to the zone axes identified, the orientations of each of these planes families are plotted in pole figures in Fig. 6 in order to highlight the presence of parallel planes between both structures. The parallel planes between the matrix and the precipitate, highlighted in the pole figures and already suspected from the diffraction images, allow to establish the following orientation relationship: $[(103)_{Al}]/[(203)_{τ1}, (0,1,0)_{Al}/((0,1,0)_{τ1}, (-3,-1,1)_{Al}/((-3,-1,2)_{τ1}, (-3,1,1)_{Al}/((-3,1,2)_{τ1}))]$.

Although encouraging, this observation is as such insufficient to guarantee that the τ₁ phase is a good nucleating agent for the aluminum phase according this orientation relationship. Indeed, the literature indicates several crystallographic criteria that have to be respected [26,27]. Firstly, the parallel planes between both structures must have a small interplanar misfit. Moreover, both structures must have a common dense atomic direction contained within the parallel planes, with a low interatomic distance misfit along this direction. The expression of these misfits is somewhat arbitrarily defined by Wang et al. [26], who propose a limit value of 10 %. To quantify matters, the Vesta software [28] is used to study this orientation relationship from a crystallographic point of view.

The first line in Fig. 7 represents the structures in their respective experimental orientations, each containing the parallel planes previously identified. As the structures are oriented along zone axes, the planes are parallel to the reader's view. The second line of Fig. 7 shows measurements in the (0,1,0) planes of the matrix and precipitate, while the last line concerns the (-3,-1,1) plane of the matrix and the (-3,-1,2) plane of the precipitate. These last two lines show the structures in an orientation that allows the relevant plane to face the reader. On each of these planes, the densest atomic direction is identified, and annotated on the figure. To facilitate the understanding of the figure, the atoms contained in the highlighted planes are circled in dotted lines. Table 4 then presents the values of interplanar and interatomic distance misfits, respectively called *f*<sub>d</sub> and *f*<sub>r</sub>. It reveals that all misfits are below 10 %, and therefore that the orientation relationship detected by TEM allows the τ₁ phase to be a good nucleating agent for the aluminum phase according to the orientation relationship written above.

It is important to mention that this orientation relationship has only been identified for a single precipitate. Nevertheless, it is reasonable to assume that intragranular τ₁-phase precipitates are germinating. On the one hand, electron microscopy analyses have sometimes shown the presence of several Zr-rich precipitates within the same grain, which necessarily implies that all but one of them are not responsible for germination. Furthermore, these analyses showed the presence of Zr-rich precipitates within most grains. Even though such precipitates do not show an orientation relationship with the matrix, they could nevertheless contribute to the reduction of the nucleation energy barrier as foreign particles usually do. In this respect, it should be noted that no other phases that could be candidates for germinating agents were identified. In the same line of thought, the recently proposed Icosahedral Short Range Order (ISRO) nucleation model can be safely ruled out, since TEM analyses did not show the expected high twin densities associated with such a mechanism [29,30].

<figure><img src="image_8.png" /><figcaption>Fig. 6. Pole figures of planes families {1,0,0} for both matrix and precipitate (left), and of planes families {3,1,1} for matrix and {3,1,2} for the precipitate (right).</figcaption></figure>
<figure><img src="image_9.png" /><figcaption>Fig. 7. Crystallographic study of the orientation relationship between matrix and precipitate.</figcaption></figure>

<table><caption>**Table 4**
Interplanar and interatomic distance misfits.</caption><thead><tr><th>Parallel planes</th><th>(0,1,0)<sub>Al</sub>/(0,1,0)<sub>τ1</sub></th><th>(-3,-1,1)<sub>Al</sub>/(-3,-1,2)<sub>τ1</sub></th></tr></thead><tbody><tr><td>f<sub>d</sub></td><td>3.41 %</td><td>2.26 %</td></tr><tr><td>f<sub>r</sub></td><td>3.67 %</td><td>3.67 %</td></tr></tbody></table>

### 3.3. Thermodynamics in the Al-Zr-Si-C system

#### 3.3.1. Solidification path

The Al-Zr-Si-C database can be used to follow the formation path of phases during system cooling and solidification. Here, the initial composition is set to match the composition of the R3 composite powder in terms of mass fraction of Si, Zr and C, supplemented by aluminum, as shown in the Table 5.

By calculating the solidification path, the fraction of each phase that will form as the system cools can be predicted. The calculations can be carried out under global (macroscopic) or local (solid-liquid interface) equilibrium conditions. In the latter case, the Scheil-Gulliver model that assumes a homogeneous liquid, no diffusion in the solid and the release of solutes ahead of the growth front, is used. The result presented here relies on the Scheil-Gulliver model, expected to be more representative of rapid solidification conditions, such as those encountered in L-PBF.

The Scheil-Gulliver solidification path is shown in Fig. 8. The first crystalline phase formed during cooling, albeit only to a very small solid fraction, is Zr₂Al₃C₅. Next, the τ₁ phase is solidified from 1100 °C to circa 800 °C, before the appearance of the τ₃ phase, defined as Zr₃Al₄Si₅

<table><caption>**Table 5**
Elemental composition (wt%) considered for solidification path calculation.</caption><thead><tr><th>Element</th><th>Al</th><th>Si</th><th>Zr</th><th>C</th></tr></thead><tbody><tr><td>Mass fraction</td><td>Bal.</td><td>7 10<sup>-2</sup></td><td>3.34 10<sup>-2</sup></td><td>1.47 10<sup>-2</sup></td></tr></tbody></table>

<figure><img src="image_10.png" /><figcaption>Fig. 8. Solidification path in Scheil-Gulliver condition for R3 composition.</figcaption></figure>

[13]. The calculation then indicates solidification of the Al₄C₃ and Si₂Zr phases from 660 °C, followed by solidification of the aluminum-rich phase.

Regarding the comparison with the experimental results obtained, this solidification path predicts phases that are not observed in practice (Zr₂Al₃C₅, τ₃), and also shows that some experimentally formed phases like ZrC are not expected from thermodynamic arguments in the current state of the database. There is therefore no perfect correspondence between experimentally obtained phases and thermodynamics, although
some important similarities are observable, notably through the formation of the τ₁ and Al₄C₃ phases.

Before further analyzing these thermodynamic calculation results in relation to the experimental data previously obtained, heat treatment tests with lower heat transfer kinetics compared to L-PBF process are carried out on the R3 composite powder to support the solidification path results.

### 3.3.2. Heat treatments tests with lower kinetics

Firstly, differential thermal analysis (DTA) tests were carried out on the R3 composite powder, as well as on the reference AlSi₇Mg₀.₆ powder, in order to estimate the formation temperatures of the various phases. The results are qualitatively similar, except for the presence of an exothermic peak observed around 742 °C on the R3 composite powder (see Supplementary data 5).

Then, interrupted thermal treatments on R3 composite at 650 °C and 800 °C allow an XRD identification of the phase formed in this temperature range (Fig. 9).

The τ₁ phase is detected at 800 °C, but not at 650 °C, indicating that the exothermic peak observed at 742 °C is characteristic of the formation of this τ₁ phase, in agreement with the solidification path in Scheil-Gulliver conditions that predicts its precipitation around 800 °C (Fig. 8). Moreover, the ZrSi₂ compound, with an orthorhombic crystallographic structure, is detected after both heat treatments, as expected at the end of the solidification path presented above. During heat treatment at 800 °C, the ZrSi₂ phase is observed with lower peak intensities than at 650 °C, implying a smaller amount of this compound. The carbides SiC, Al₄C₃ or ZrC are not detected, but it is not however possible to conclude whether this is due to their total absence or their presence in too small quantities. No conclusions can therefore be drawn as to the fate of carbon in these experiments.

## 4. Discussion

Before turning to a discussion of the relation between experimental observations with the predictions of the thermodynamic model, a number of questions needs to be addressed. In chronological process terms, the first regards the becoming of the ZrO₂ particles and the associated issue of the availability of metallic Zr to form alternative phases needs to be addressed. Indeed, according to the Ellingham diagram, Zr should be a better reducer than Al, meaning that ZrO₂ could be expected to be stable within an aluminum based melt. However,

preliminary thermodynamic calculations within the (Al-Zr-O) system show that liquid Al is in equilibrium with both Al₂O₃ and ZrO₂ at high temperatures, so that Zr dissolution within the melt is likely. Another option to account for the release of Zr is to invoke the presence of Mg, a very potent reducer, within the alloy composition. As a matter of fact, Mg based oxides had been observed in a previous work [8] on LPBF manufacturing of a related Al 6061 alloy.

Finally a last mechanism would be to consider that at least some of the ZrO₂ particles could melt. The question of the maximum temperature reached during the process is again a tricky one, but values close to the ZrO₂ melting point (circa 2715 °C) are not unreasonable. Furthermore, since the presence of Al in the vicinity of ZrO₂ leads to a decrease of the liquidus temperature, a contribution of the melting mechanism to the solubilisation of Zr appears quite likely. In any case, from a practical standpoint, it can be stated that we are fairly confident that a sufficient amount of Zr will be released within the melt. Turning again to our previous work on LPBF manufacturing of a related Al 6061 alloy [8], it was observed that even at a much lower ZrO₂ content within the powder, a large number of Al₃Zr precipitates could be formed in the fabricated parts. Moreover, the observed increase of the intensity of the Al₂O₃ peak on the XRD diffractograms is compatible with the trapping by Al of oxygen released from ZrO₂.

Another point that needs to be stated is that, as mentioned before and discussed in detail in our previous publication, a full transformation of SiC into Al₄C₃ was observed for the reference AlSi₇Mg₀.₆-SiC composite material without ZrO₂ addition. Regarding a potential modification of the laser absorption due to the addition of ZrO₂ particles, we did not find evidence of an increase with respect to the reference AlSi₇Mg₀.₆-SiC powders, that would e.g. result in an increase of keyhole porosities at a given energy density. Therefore, it can be safely stated that the reduction of the amount of Al₄C₃ phase after adding ZrO₂ is related to the fact that at least some of the released carbon has been trapped within Zr precipitates.

Regarding the question of whether the amount of intragranular τ₁ phase is sufficient to achieve the purpose of grain refinement, characterizations performed on bulk materials did not allow to accurately quantify the density of intra-granular precipitates, a major issue in this respect being that the position of the grain boundaries was often not very clear on the 3D FIB-SEM data. Nevertheless, our in depth observations of a large number of images allow to ensure that τ₁ precipitates are visible in most grains, with in some instances more than one precipitate within a given grain. Besides, turning once more to our previous work on LPBF manufacturing of a related Al 6061 alloy [8], it was observed that even at a much lower ZrO₂ content within the powder, the number of Al₃Zr precipitates was largely sufficient to promote equiaxed grain growth. Therefore our opinion is that the amount of intragranular τ₁ phase is indeed sufficient for grain refinement purposes.

We now turn to a discussion of the relation between the experimental observations with the predictions of the thermodynamic model, with the objective to better understand the physico-chemical mechanisms involved. Table 6 lists the results obtained in terms of phases in the three situations presented above: L-PBF, heat treatment at 800 °C and thermodynamic calculations.

Firstly, in situ ZrC synthesis is only observed after L-PBF. In this respect, it is important to emphasize that the solidification path calculation is based on the overall composition of the system, and thus may not be able to account for local variations in composition. In the system implemented here, involving a micrometric aluminum powder coated with SiC and ZrO₂ nanoparticles in a process with fast kinetics, this is a limitation to keep in mind. As the hypothesis of ZrC formation is based on local compositional conditions, it was expected that its formation would not be predicted by thermodynamic calculations. Furthermore, the results of experiments aimed at achieving thermodynamic equilibrium (heat treatments) do not show the formation of ZrC.

It is therefore likely that the formation of ZrC in L-PBF was promoted by local compositional conditions on the one hand, and by the kinetic

<figure><img src="image_11.png" /><figcaption>Fig. 9. Heat treated R3 composite powder diffractograms.</figcaption></figure>
<table><caption>**Table 6**

Comparative table of phases obtained with the three methods used.</caption><thead><tr><th>Method</th><th>L-PBF</th><th>Heat treatment (800 °C)</th><th>Thermodynamic calculations</th></tr></thead><tbody><tr><td>ZrC</td><td>✔</td><td>✗</td><td>✗</td></tr><tr><td>τ<sub>1</sub></td><td>✔</td><td>✔</td><td>✔</td></tr><tr><td>Al<sub>4</sub>C<sub>3</sub></td><td>✔</td><td>✗</td><td>✔</td></tr><tr><td>ZrSi<sub>2</sub></td><td>✗</td><td>✔</td><td>✔</td></tr><tr><td>Zr<sub>2</sub>Al<sub>3</sub>C<sub>5</sub></td><td>✗</td><td>✗</td><td>✔</td></tr><tr><td>τ<sub>3</sub></td><td>✗</td><td>✗</td><td>✔</td></tr></tbody></table>

aspect of the process on the other, which did not allow enough time for the system to reach its thermodynamic equilibrium. A detailed investigation of the convecto-diffusive mass transport mechanisms within the melt pool is far outside the scope of the present work. Nevertheless, it can be thought these mechanisms are not fast enough to instantaneously homogenize the melt, and that the intimate contact between the ZrO2 and SiC additives on top of the AlSi7Mg host particles favors the establishment of process conditions where the formation of ZrC is thermodynamically favored. The 1000 °C isothermal section of the Al-Zr-C phase diagram of Fig. 1(a) indeed shows that for the both R1 and R3 experimental compositions, ZrC is expected to be in equilibrium with Al—Zr intermetallics up to 75 mol% Al.

Again, neither the temperature nor the local compositions within the process can be accurately specified, but it can be safely thought that over a large temperature range the local Al content remains below 75 mol%, thus justifying the formation of ZrC. For industrial applications, it should be noted that, although the formation of ZrC necessarily implies a reduction in the Al<sub>4</sub>C<sub>3</sub> content in the parts, its heterogeneous distribution in the material, in the form of agglomerates, will not have the expected impact in terms of mechanical reinforcement.

Then, in line with the predictions of solidification path calculations, the τ₁ phase is observed experimentally after L-PBF, but also after the one-hour heat treatment at 800 °C. Thanks to its role as nucleant, a fine equiaxed microstructure is obtained, instead of a coarser columnar microstructure. Such an equiaxed microstructure could have a beneficial impact on the mechanical properties of composites, promoting isotropic behavior, and with higher performances associated to the numerous grain boundaries acting as obstacles to dislocation movement.

Regarding the ZrSi₂ phase, its formation is observed following heat treatment, confirming the results of thermodynamic calculations in the quaternary database. On the other hand, the presence of this ZrSi₂ phase was not detected in the samples after L-PBF, which can probably be attributed to the fast kinetics of the process, which did not give this phase time to nucleate, or at least to grow to a measurable amount.

Table 6 then shows the presence of Al<sub>4</sub>C<sub>3</sub> after L-PBF, a phase also expected from the solidification path. Local compositional conditions that led to the formation of ZrC, necessarily decreased the available amount of Al<sub>4</sub>C<sub>3</sub>, but aluminum carbide formation remains expected for the overall system composition. However, as previously indicated, no conclusion can be drawn on the formation of carbide phases during heat treatment due to their absence from XRD diffractograms, which requires characterization at finer but more local scale, complicating the understanding of the composites physicochemistry.

Questions remain about the phases calculated but not obtained or experimentally detected, neither by L-PBF and heat treatments, namely τ₃ and Zr₂Al₃C₅. A first assumption is again that they may be present in quantities too small to be detected. Their formation may also be restricted by diffusion kinetics. However, the fact they could not be

observed after heat treatments could also point to deficiencies in the database. Indeed, as previously mentioned, the Al-Zr-Si and Al-Zr-C ternary systems suffers from a lack of reliable phase diagram and thermodynamic data.

Finally, the L-PBF results, in line with the solidification path calculation, and by comparison with the heat treatment results, indicate the solidification of the τ₁ phase as well as ZrC and Al₄C₃. On the other hand, the fast kinetics of the L-PBF process limit the formation of the ZrSi₂ phase since this phase is observed when the kinetics are slower. Note that the low-temperature equilibrium system does not include the τ₁ phase, so this metastable phase is expected to disappear during thermal annealing in favor of a stable phase at lower temperatures. To check this point, a R3 bulk sample was heat treated for two hours at 500 °C, and the XRD analysis indeed showed the disappearance of the τ₁ phase and the formation of the ZrSi₂ phase.

Based on these results, it would be interesting to propose a reaction scheme for the L-PBF processing of an Al-SiC-ZrO₂ composite powder. However, the complexity of the system and the different elements involved in precipitate formation, particularly Zr, make it difficult to take any position. Nevertheless, one could imagine this phases formation reaction scheme in two different situations, depending on the local composition: one with ZrC, the other with Al₄C₃, with τ₁ in both cases. Various reaction mechanisms could be imagined to obtain this final system.

## 5. Conclusion

This work takes place in the context of the L-PBF elaboration of Al-SiC composites, and more specifically of the thermodynamic instability of SiC in liquid aluminum, leading to the formation of Al₄C₃, a water-soluble carbide that reduces parts lifetime. Our objective was to check the effect of the addition of a source of zirconium to the system to promote the formation of ZrC and thus limit the formation of Al₄C₃. To do so, AlSi₇Mg₀.₆-SiC-ZrO₂ composite powders were elaborated by dry coating and printed in a standard L-PBF equipment. The parts obtained were then thoroughly characterized. In parallel, a first version of a thermodynamic database for the Al-Zr-Si-C system has been set up to help understand the thermochemistry of the composites produced. The main conclusions of this work are summarized as follows:

1. L-PBF processing of AlSi₇Mg₀.₆-SiC-ZrO₂ powder reduces the amount of Al₄C₃ through the formation of ZrC. However, ZrC formation is only possible under local compositional conditions, leading to the formation of ZrC agglomerates, which are a priori not advantageous for the mechanical properties of the parts. Furthermore, Al₄C₃ precipitates can still be observed, even if in smaller quantities.

2. Bulk composites exhibit a fully equiaxed microstructure with a linear grain size of around 1 µm, explained by the precipitation of a nucleating Al-Zr-Si ternary phase named τ₁. Indeed, it displays all characteristics of a good nucleating agent according to the following orientation relationship with the matrix: $[103]_{Al} // [203]_{τ₁}$, $(0,1,0)_{Al} // ((0,1,0)_{τ₁}, (-3,-1,1)_{Al} // (-3,-1,2)_{τ₁})$.

3. The L-PBF composites microstructures show the precipitation of ZrC, Al₄C₃ and τ₁ phases. Coupled with heat treatment tests with slower kinetics, thermodynamic calculations allow to determine the phases that could be obtained at equilibrium, ZrSi₂ and Al₄C₃, after heat treatment of L-PBF materials. These results revealed that the formation of the expected ZrSi₂ phase was limited by the L-PBF process kinetics, and that the metastable τ₁ phase disappears in favor of ZrSi₂ during heat treatment of a bulk sample from L-PBF.

Regarding perspectives, it would be interesting to study the effect these L-PBF microstructures may have on the mechanical properties of these composites, and to compare them with those of an Al-SiC composite produced by casting and LPBF. These properties could also be compared with those of an AlSi₇Mg₀.₆-ZrO₂ composite by L-PBF, which
would probably allow to obtain an equiaxed microstructure similar to that of the composites developed in this work, without ZrC agglomerates and residual Al₄C₃ precipitates.

Nevertheless, such an objective would require an optimization of the density of the bulk materials obtained after L-PBF. Among the options to be pursued is the drying of the powders before printing, since moisture is known to lead to the formation of occluded gases porosity [31]. In addition, the reinforcement ratio could be optimized, since it has been observed that it is more difficult to obtain dense materials with a high reinforcement ratio. A lower reinforcement ratio is thus expected to lead to density improvements. A more extensive parametric study could also be carried out.

Beyond the scope of this study that focused on the addition of a Zr source, it would also be very interesting to study the effect that the addition of a Ti source could have, since TiC also has a more negative Gibbs energy than Al₄C₃ [6].

## Funding

Financial support from the cross-cutting CEA program "Matériaux et Procédés" is gratefully acknowledged. This work, carried out on the Platform for Nanocharacterisation (PFNC), was supported by the "Recherche Technologique de Base" and "France 2030 - ANR-22-PEEL-0014" programs of the French National Research Agency (ANR).

## CRediT authorship contribution statement

**Marie-Reine Manlay:** Writing – original draft, Software, Methodology, Investigation, Data curation, Conceptualization.
**Camille Flament:** Writing – review & editing, Supervision, Methodology, Investigation.
**Stéphane Gossé:** Writing – original draft, Validation, Software, Methodology.
**Mathieu Soulier:** Supervision, Resources, Project administration, Investigation, Funding acquisition, Conceptualization.
**Jean-Paul Garandet:** Writing – original draft, Supervision, Methodology, Funding acquisition, Conceptualization.

## 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 article.

## Data availability

Data will be made available on request.

## Acknowledgements

The authors would like to thank Laurent Guillem-Guerrero for the 3D printing experiments, Nathalie Ladrat for the FIB lamella preparation and Laurent Chaffron for ATD tests. The authors also acknowledge many fruitful discussions with SIMAP, Univ. Grenoble Alpes, researchers.

## Appendix A. Supplementary data

Supplementary data to this article can be found online at https://doi.org/10.1016/j.matchar.2024.114472.

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