# Tensile properties and fracture toughness of chemically-architECTured high-entropy alloys

Yazid Madi, A-F. Gourgues-Lorenzon, B. Villeroy, M. Laurent-Brocq

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Yazid Madi, A-F. Gourgues-Lorenzon, B. Villeroy, M. Laurent-Brocq. Tensile properties and fracture toughness of chemically-architECTured high-entropy alloys. Materials Science and Engineering: A, 2025, 944, pp.148922. 10.1016/j.msea.2025.148922. hal-05222743

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# Tensile properties and fracture toughness of chemically-architECTured high-entropy alloys

Y. Madi<sup>a, ID</sup>, A-F. Gourgues-Lorenzon<sup>a, ID</sup>, B. Villeroy<sup>b</sup>, M. Laurent-Brocq<sup>b,*</sup>

## ARTICLE INFO

**Keywords:**

Chemically-architECTured alloys

High-entropy alloys

Spark plasma sintering

Fracture toughness

Tensile properties

Microstructural characterization

## ABSTRACT

Chemically-architECTured materials combine a Ni matrix with a high-entropy alloy (HEA) phase to introduce a controlled chemical gradient. These materials are currently processed through spark plasma sintering (SPS) and are only available in limited quantities. In this work, their tensile properties and fracture toughness were determined for the first time using miniaturized specimens. A tensile yield strength of 476 MPa, a fracture elongation of 23%, and a fracture toughness of 41 MPa·√m were measured for a chemically-architECTured alloy exhibiting a steep chemical gradient. For a smoother gradient, the yield stress decreased to 338 MPa while the tensile elongation and fracture toughness increased to 44% and 125 MPa·√m, respectively. Chemical architECTuration resulted in strengthening compared to a reference SPSed HEA, but at the expense of ductility and fracture resistance. Impurity particles located at prior HEA particle boundaries played a critical role in the ductile damage and fracture mechanisms. Earlier void nucleation and accelerated damage development were observed in the material with the steeper chemical gradient, while the softer Ni-rich matrix appeared as a preferred path for fatigue crack propagation. Based on these results, guidelines for further microstructural optimization of these alloys were proposed to improve their mechanical properties.

## 1. Introduction

Strengthening mechanisms are key in metallurgy to improve the mechanical resistance of structural alloys. The respective contributions of solid-solution, dislocations, grain boundaries and precipitates have long been known and are still active research areas [1]. Nanocrystalline metals [2] and high-entropy alloys [3,4] have revived the potential of grain boundary and solid-solution strengthening. However, enhancing strength alone is not sufficient: the trade-off with ductility must also be preserved [5], for instance, by promoting deformation mechanisms such as twinning-induced plasticity and transformation-induced plasticity [6].

Recently, mesoscale composition fluctuations have emerged as another promising strengthening mechanism to overcome the strength-ductility trade-off [7]. This concept draws inspiration from the early stages of spinodal decomposition, where chemical gradients between phases enhance the yield strength [8], likely through a mechanism involving coherency strain [9]. To prevent the brittleness often associated with spinodal microstructures [10], the concept is also inspired by architECTured materials, where an intermediate mesoscale is introduced between the microstructure and macrostructure [11]. ArchitECTured materials mostly present a heterogeneous and yet controlled grain

size distribution, like bimodal grains or harmonic structures [12]. Combining these two concepts, a new multi-scale microstructure was designed and named “chemical architECTuration” (Fig. 1). In this concept, two phases, hereafter referred to as A and B, are chosen to have the same crystal structure but different chemical compositions; they should belong to the same solid-solution domain in an equilibrium multicomponent phase diagram (Fig. 1d). At the micron scale (Fig. 1a), phases A and B are separated by a region with a continuous chemical gradient (Fig. 1c), referred to as the “interphase”. At the mesoscale (Fig. 1b), the phases are separated by a three-dimensional network of composition fluctuations. Alternatively, chemically-architECTured alloys can be considered as composite materials, yet with a major difference: the transition between phases A and B is smooth and continuous. It should be emphasized that chemically-architECTured alloys are metastable by definition. These chemical fluctuations are intended to introduce a new strengthening mechanism at the scale of the prior powder particles, which is coarser than the grain size. At the microscopic and atomic scales, grain boundaries and solid solution also contribute to conventional strengthening mechanisms.

It was chosen to apply this concept to the face centered cubic single-phase CoCrFeMnNi high-entropy alloy because the Co-Cr-Fe-Mn-Ni

<footnote><sup>a</sup> Mines Paris, PSL University, Centre des Matériaux, CNRS UMR 7633, BP 87, 91003 Evry Cedex, France</footnote>

<footnote><sup>b</sup> Université Paris-Est (UPE), ICMPE (UMR 7182), CNRS, UPEC, 2-8 rue Henri Dunant, 94320 Thiais, France</footnote>

<footnote>* Corresponding author.</footnote>

<footnote>E-mail address: mathilde.laurent-brocq@cnrs.fr (M. Laurent-Brocq).</footnote>
<figure><img src="image_3.png" /><figcaption>**Fig. 1.** Scheme of the concept of chemical architecture. Spatial distribution of the two phases A and B along with the interphase at (a) the micron scale and (b) the mesoscale, where a dashed white square indicates the area represented in (a). (c) Concentration profile across the interphase. (d) Simplified equilibrium phase diagram (isopleth section close to the A-rich region). Crosses labeled 1 to 4 are used to connect the spatial distribution to the concentration profile and the phase diagram (cross No. 1: B phase; cross No. 4: A phase).</figcaption></figure>

system exhibits thermodynamic properties which are favorable to build the aimed chemical gradients [13]. From a mechanical point of view, CoCrFeMnNi alloys exhibit a high ductility [14] and an excellent toughness [15] but a limited yield strength [16,17]. Thus they would benefit from an additional strengthening mechanism. By applying spark plasma sintering (SPS) to a (Ni)-(CoCrFeMnNi) powder mixture, chemically-architected high-entropy alloys, with a controlled interphase, were processed [7,18]. Using compression tests [18], the strengthening effect of the interphase was measured locally and the role of chemical gradient was highlighted thanks to finite element simulations [19]. However, the tensile and fracture properties of such alloys have never been reported in literature. The current study aimed to expand this understanding by evaluating tensile properties, ductility, and fracture resistance in chemically-architected HEAs produced via SPS. Three materials were studied: two chemically-architected HEAs with a thin and a thick interphase respectively, and thus with a steep and smooth chemical gradient respectively. A single-phase HEA was also processed in similar conditions for reference.

Miniaturized specimens offer a practical solution to characterize materials available only in limited quantities or when standard-sized specimens cannot be extracted from structural components [20]. Miniaturized flat and smooth tensile specimens are widely used to investigate anisotropy, size effects, and in-service degradation [21–24]. However, the development of miniaturized specimens for fracture toughness assessment is less advanced. In fracture mechanics, $J_{IC}$ serves as a specific measure of fracture toughness under plane-strain conditions, quantifying the resistance to crack propagation of a given material, under small-scale yielding conditions. Miniaturized toughness specimens, particularly Compact Tension (CT) types, have been widely used in the nuclear industry to study the ductile-to-brittle transition in irradiated

steels. With careful consideration of scale effects, these mini-specimens are effective for assessing fracture properties [25–27].

For materials produced by SPS and available in limited quantities, most studies in the literature focused on compression tests [18], flat tensile specimens [28], or indentation tests to estimate their toughness [28,29]. To the authors’ knowledge, fracture toughness evaluation using genuine fracture mechanics specimens has not been explored yet for these materials. In the present work, both smooth tensile and CT-type miniaturized specimens were used to investigate relationships between the microstructure, the tensile behavior and the fracture properties of chemically-architected alloys, which are critical for their efficient design.

## 2. Materials and experimental methods

### 2.1. Materials

Two commercial master powders were used in this study: (i) a high purity Ni powder (metal content > 99.996 wt.%) and (ii) a gas-atomized equimolar CoCrFeMnNi HEA powder. A low carbon contamination, of 480 and 50 ± 10 ppm respectively for the Ni and HEA powders, was measured by combustion infrared absorption. The oxygen content measured by reducing melting infrared absorption was 1600 and 1400 ± 100 ppm respectively for the Ni and HEA powders, respectively. Some surface contamination was observed on HEA powder particles (see Fig. 17 in the Appendix). Thus they were treated in hydrochloric acid prior to sintering. The HEA and Ni powders were sieved between 20 and 36 µm and below 36 µm respectively, then mixed in a 1:1 mass ratio during 30 min in a Turbula® mixer.

A spark plasma sintering apparatus (model Fuji 515-S), of the “Plateforme Ile-de-France de Frittage” was used. The SPS parameter
<figure><img src="image_4.png" /><figcaption>Fig. 2. Geometry of tensile and fracture mechanics specimens. (a) Schematic view of miniature tensile specimens. (b) Picture of a mDCT specimen cut from a 15 mm diameter pellet of the (HEA+Ni)-S850 material. (c) Schematic drawing of the mDCT specimen with side-grooving.</figcaption></figure>

optimization had already been performed [7,18]. Similar sample dimensions were used in this work: a thickness between 2.5 and 3 mm, and a diameter of either 10 or 15 mm. To vary the steepness of the chemical gradient, chemically-architECTured alloys were sintered at a maximum temperature of either 850 °C or 600 °C. They are named accordingly hereafter: (HEA+Ni)-S850 (or “smooth-gradient alloy”) and (HEA+Ni)-S600 (or “steep-gradient alloy”) respectively. In both cases, the temperature was increased from 100 °C min⁻¹ to the maximum temperature; then, a uniaxial pressure was applied to the pellet and both the pressure and the maximum temperature were maintained during 15 min under vacuum. To obtain a full densification, the maximum applied pressure was set to 100 MPa with a graphite mold for (HEA+Ni)-S850, and to 400 MPa with a WC/Co mold for (HEA+Ni)-S600. The same parameters were applied to the HEA powder alone but in that case, only the first conditions allowed to obtain a fully dense material (HEA-S850 hereafter). An oxygen content of 520 ± 40 ppm was measured in (HEA+Ni)-S850 after sintering. This significant reduction in oxygen content relative to the master powders indicates an effective deoxidation powder treatment, potentially reinforced by SPS’s reducing effect.

## 2.2. Mechanical behavior

In this study, two types of miniaturized tensile specimens were employed (Fig. 2a): smooth ones (mST) and notched ones (mNT). The mST specimens had a 1.2 mm gauge diameter ($Φ_0$) and a 2 mm gauge length ($L_0$), while the mNT specimens featured a minimum diameter of 1.4 mm and a notch radius ($R_0$) of 0.56 mm. At the notch root, the initial cross-sectional area ($S_0$) was 1.54 mm², encompassing over 1500 HEA powder particles. These miniaturized specimens were carefully extracted from sintered pellets via electrical discharge machining (EDM) and meticulously machined to a surface roughness of 0.35 µm (Ra, arithmetic mean roughness), to minimize residual surface stresses.

Mini Disk-shaped Compact Tension (mDCT) specimens, with dimensions of 7.4 mm in width (W) and 3 mm in thickness (B), were also prepared (Fig. 2b–c). They were machined from SPS pellets and polished to a surface roughness of Ra = 1 µm to allow for fatigue precrack propagation tracking, following ASTM E1820 guidelines [30]. Due to their reduced size, the precracking procedure was carefully controlled to avoid influencing the subsequent fracture results. Initially, a maximum stress intensity factor ($K_{max}$) of 14 MPa√m was applied, with a load ratio (R) of 0.1 and a frequency of 35 Hz. The load was then adjusted to reduce $K_{max}$ to 11 MPa√m by the end of precracking. During each interruption, the specimen was flipped to minimize misalignment effects along the load line, enhancing crack propagation symmetry. After precracking, side grooves were added by EDM with a groove tip root radius of 50 µm, creating a net thickness ($B_n$) of 2.4 mm to promote a straight crack front, essential for reliable fracture toughness testing.

All tests were conducted in an air environment using a dynamic MTS servo-hydraulic machine equipped with an optical extensometer employing the Edge Tracing (ET) technique, as described in [20, 31]. The setup included a high-precision displacement sensor and a 5 kN load cell, allowing real-time communication between the Keyence XG-X2800 image controller and the MTS machine controller. This communication enabled adaptive displacement rate adjustments based on optical measurements of radial contraction in tensile specimens. Both the elongation and the radial contraction were measured using an optical extensometry method as detailed in [20]. To characterize the tensile behavior beyond the onset of necking, the radial contraction strain $ΔΦ/Φ_0$ — with $Φ_0$ denoting the minimum diameter within the gauge length — was preferred over the conventional axial elongation. This choice allows for a better emphasis on the localized deformation behavior of mST specimens and enables comparison with the response of mNT specimens. Furthermore, the consistent trend observed in [20], in the engineering stress $σ$ versus radial contraction strain $ΔΦ/Φ_0$ across specimens of different geometries suggests that the local plastic behavior remains largely unaffected by specimen size or machining. The ET technique also monitored key fracture toughness testing parameters, such as the Load Line Opening Displacement (LLOD) and the Opening Angle ($α$). Further details on the methodology can be found in [20].

Tensile tests were conducted at a nominal strain rate ($ε̇$) of 2.5 × 10⁻⁴ s⁻¹. Although the tensile machine was displacement-controlled, a control loop was implemented to continuously adjust the crosshead speed based on radial contraction measurements. This adjustment ensured that the radial strain rate ($ΔΦ/Φ_0$) remained consistent with the target strain rate in the region of interest of the specimen, according to the following relationship: $ΔΦ/Φ_0 × 2L_0 = ε̇ × L_0$, where $L_0$ corresponds to the gauge length for smooth specimens (mST), or to twice the notch root radius (2$R_0$) for notched specimens (mNT). One or two tests per material and specimen geometry (depending on material availability) were carried out until specimen fracture.

For mDCT specimens, the crack growth was monitored using the compliance method. A displacement rate of 50 µm min⁻¹ was applied to produce a 20 µm load-line displacement, corresponding to an initial $K̇$ rate of 0.53 MPa√m s⁻¹. After holding the displacement for 30 s to stabilize relaxation effects, the specimen was unloaded to 65% of the peak load and reloaded to the same force. This allowed crack length estimation based on the compliance equations defined in ASTM E1820 [30]. The cycle was repeated until a crack extension of 0.5–1.2 mm was reached, depending on the sample, after which the load was reduced to zero. Two valid toughness tests were performed per material. One specimen per condition was subsequently fractured by fatigue to reveal the initial ($a_0$) and final ($a_f$) crack lengths. Optical and compliance-based measurements agreed within 1%, confirming the reliability of the method. Each $J-Δa$ curve was analyzed according to the ASTM E1820 procedure. A power law of the form $J = A · (Δa)^n$ was fitted over the full crack extension range. The blunting line $J =$
<figure><img src="image_5.png" /><figcaption>Fig. 3. Light optical micrographs of the three materials. Nomarski contrast. (a) HEA-S850; (b) (HEA+Ni)-S850 (yellow regions: Ni-rich matrix; gray regions: HEA; white regions: chemically graded regions); (c) (HEA+Ni)-S600 material (light gray regions: HEA; yellow regions: Ni-rich matrix). Some HEA, Ni matrix and interphase regions are indicated with arrows and some of them are delineated in (b) and (c). Small black dots are impurity particles. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)</figcaption></figure>

$2\sigma_y \cdot (\Delta a - \Delta a_i)$ was defined with offsets $\Delta a_i = 0.20$ mm and 1.00 mm, the latter being used to determine the ductile tearing modulus. The flow strength $\sigma_y$ was taken as the average of the 0.2% proof stress ($\sigma_{yS0.2}$) and the ultimate tensile strength (UTS). The intersection of the power law with the blunting line at $\Delta a = 0.2$ mm gave $J_Q$, and the intersection at $\Delta a = 1$ mm gave $J_1$. To evaluate the validity of the test, especially fulfillment of small scale yielding conditions, the parameter $J_{max} = \sigma_y \cdot \min(W - a_0, B)/10$ was calculated, as prescribed by ASTM E1820. When the experimental $J_Q$ value remains below this limit, it can be considered a valid estimate of the fracture toughness, denoted as $J_{IC}$. The ductile tearing modulus, representing the slope of the $J-\Delta a$ curve, was computed between $J_Q$ and $J_1$. A larger value indicates a higher resistance to crack growth and a higher crack propagation stability.

## 2.3. Microstructures and physical fracture mechanisms

As in [18,32], the microstructures of the chemically-architected materials were characterized by Scanning Electron Microscopy (SEM), coupled with Energy Dispersive Spectroscopy (EDS) to determine the fractions of Ni, HEA, and interphase regions. Seven maps selected on at least two pellets per material, were recorded. Each map covered an area of $144 \times 108 \ \mu\text{m}^2$ with a pixel size of 0.56 µm. Each pixel corresponded to a high-resolution EDS spectrum, from which an accurate composition could be calculated. A pixel was assigned to the HEA phase or to the Ni phase if its Ni content was below 28 at.% or above 90 at.%, respectively. Table 1 reports the average and standard deviation values over the seven maps, considered as representative of the whole sample, as already established [19]. Phase maps were also generated, and the width of the interphase was measured at ten locations on each map. Here again, the reported value and uncertainty correspond to the average and standard deviation of all measurements. SEM coupled with Electron Backscatter Diffraction (EBSD) was also performed for the three materials. The step size was 150 nm for (HEA+Ni)-S600 and 500 nm for the other samples. A minimum of 10 pixels was required to define a grain. Grain boundaries were defined using a crystallographic disorientation of at least 10° while $\Sigma3$ twin boundaries were excluded. The grain size was determined by calculating the average diameter over at least 500 grains per sample, with the reported uncertainties corresponding to the 95% confidence interval. This uncertainty may have been smaller than the width of the grain size distribution, as it will be discussed in Section 3. The hardness was measured with a micro-indenter under a 1000 gf load. At least 10 measurements were performed on each sample. At least 4 pellets for each chemically-architected alloy were tested. The hardness of the Ni powder sintered alone with a maximum temperature of 850 °C had already been measured as 88 (±1) HV [18].

Typical views of the microstructures after conventional polishing and a colloidal silica finish are shown in Fig. 3.

After the tests, all fracture surfaces, including the fatigue precrack of mDCT specimens, were observed in a Zeiss Sigma 300 field emission gun SEM, using two secondary electron detectors, namely, the

conventional Everhardt-Thornley device and an in-lens detector. The working distance was adapted to the required spatial resolution, from 10.5 mm down to 4 mm. Local EDS chemical analysis was conducted on selected regions using a SDD Bruker XFlash 5030 device, both in point and map analysis configurations. The high voltage was set to 5 kV for imaging, and to 20 kV for chemical analysis. The counting time for EDS measurements was about 30 s for each point analysis and the aperture was adjusted to get a 30%-deadtime.

The relationships between the microstructure, the crack path, and the damage development mechanisms were investigated on longitudinal cross-sections of mST, mNT, and unbroken mDCT specimens, after cold mounting in resin, grinding down to mid-thickness, conventional polishing using diamond pastes, and a manual colloidal silica finish. Both bright field and Nomarski contrast were used with a Leica CTR 6000 optical microscope, in order to determine the crack path with respect to the different microstructural regions, namely, the HEA powder particles, the Ni-rich matrix, and the chemically-graded region that are already illustrated in Fig. 1b. Special attention was paid to the mDCT specimens that were not fully broken after the tests, to examine the extent and nature of damage development around the crack tip after the fracture toughness tests.

# 3. Results

## 3.1. Microstructural characterization

The main microstructural characteristics are reported in Table 1 and additional data illustrating the chemical gradient are given in Appendix.

The microstructures are first described based on Fig. 3b–c. The round particles, with a diameter of approximately 20 µm, corresponded to the HEA phase. They were surrounded by a percolating matrix, appearing in yellow, which was the Ni phase. Transition areas (in white in Fig. 3b and in gray in Fig. 3c) with chemical gradients corresponded to the interphase, which is the main novelty of chemically-architected alloys. Its chemically continuous nature, area fraction and spatial distribution are further illustrated in Fig. 18 of Appendix.

As expected, the two chemically-architected alloys exhibited different interphase thicknesses (i.e., a more or less steep chemical gradient) and volume fractions. (HEA+Ni)-S850 featured an interphase with an average thickness of 10 µm and a volume fraction of 0.57, whereas (HEA+Ni)-S600 had a thinner interphase (2.6 µm) and a lower volume fraction of 0.33 (Table 1 and Fig. 4). It resulted in a steeper chemical gradient. The grain size was also different in both samples: 7.0 and 1.6 µm in average respectively in (HEA+Ni)-S850 and (HEA+Ni)-S600. Qualitatively, the grain distributions were different (Fig. 4-a3-b3). The (HEA+Ni)-S850 exhibited equiaxed recrystallized grains, with a homogeneous size distribution. This was very similar to the grain distribution of the HEA-S850 reference sample, whose grain size was slightly smaller (5 µm, Fig. 19 in the Appendix). Moreover,
<table><caption>**Table 1**

Processing and microstructural description of the sintered samples. Values in parentheses are the uncertainty on the last digit.</caption><thead><tr><th rowspan="2">Sample name</th><th rowspan="2">Sintering temperature (°C)</th><th colspan="3">Phase fraction</th><th rowspan="2">Interphase thickness (μm)</th><th rowspan="2">Grain size (μm)</th><th rowspan="2">Hardness (HV)</th></tr><tr><th>HEA</th><th>Ni</th><th>Interphase</th></tr></thead><tbody><tr><td>HEA-S850</td><td>850</td><td>1</td><td>0</td><td>0</td><td>-</td><td>5.1 (2)</td><td>183 (1)</td></tr><tr><td>(HEA+Ni)-S850</td><td>850</td><td>0.31 (5)</td><td>0.12 (6)</td><td>0.57 (3)</td><td>10 (1)</td><td>7.0 (3)</td><td>172 (5)</td></tr><tr><td>(HEA+Ni)-S600</td><td>600</td><td>0.38 (6)</td><td>0.29 (5)</td><td>0.33 (1)</td><td>2.6 (4)</td><td>1.6 (1)</td><td>196 (5)</td></tr></tbody></table>

<figure><img src="image_6.png" /><figcaption>**Fig. 4.** SEM-EDS and EBSD microstructural architecture of the chemically-architected HEA. (a1, b1, a2, b2) Phase maps, based on EDS. (a3, b3) EBSD maps showing the crystallographic orientation of the sample vertical direction, coded according to the standard stereographic triangle. a2 and a3 as well as b2 and b3 correspond to the same area and are a slightly higher magnification compared to a1 and b1. The scale bar is 20 µm for all pictures.</figcaption></figure>
<figure><img src="image_7.png" /><figcaption>**Fig. 5.** SEM images at low and medium magnifications of (a) the HEA powder after sieving and surface acid treatment and the sintered (b) HEA-S850 and (c) (HEA+Ni)-S850. In (a), the bottom image corresponds to a cross section of a powder particle.</figcaption></figure>

the microtexture of this material was completely uncorrelated from the local chemistry. On the contrary, in (HEA+Ni)-S600, the HEA phase contained some larger grains, and some very small grains were sometimes found in the interphase (Fig. 4-b3).

The contamination in the powder and subsequently in the sintered samples was also characterized (Fig. 5). Numerous submicronic precipitates were detected within the HEA powder particles (Fig. 5a). Similar precipitates were found after sintering, spread in the whole sample for HEA-S850 (Fig. 5b) and only in the dark gray areas (i.e. the HEA phase) for (HEA+Ni)-S850 (Fig. 5c). As it is emphasized in the magnified image of (HEA+Ni)-S850, some precipitates were more densely located at the prior HEA particle boundaries (PPBs), but this was not systematic. According to EDS point measurements, the precipitates were enriched in Cr, Mn, and O in the three materials.

## 3.2. Tensile behavior

In the following, the engineering stress is defined as $\sigma = F/S_0$, where F is the applied load and $S_0$ is the initial cross-sectional area. As previously mentioned, this work focused on analyzing $\sigma$ vs. $\Delta\Phi/\Phi_0$ curves as shown in Fig. 6. The mechanical properties of the three materials summarized in Table 2 were derived from conventional engineering stress-strain curves.

All three materials exhibited a macroscopically ductile behavior, even in the presence of a severe notch. From optical measurements, the necking coefficient at fracture, $Z_{opt}$, was calculated as follows: $Z_{opt} = 1 - (1 - \Delta\Phi/\Phi_0)^2$. This coefficient is compared in Table 3 with $Z_{frac}$ for mST specimens, determined using the same equation but using SEM images of the fracture surfaces. These two coefficients were clearly comparable due to isotropic post-necking behavior leading to round-shaped fracture surfaces. In the following, only $Z_{opt}$ was considered due to its higher reliability: it offered calibration control, avoided distortion effects in low-magnification SEM imaging, and was unaffected by cup-and-cone fracture morphology for detecting the minimum diameter.

Regarding mST specimens, the HEA-S850 material exhibited the best mechanical properties, including ultimate tensile strength (UTS), uniform elongation (UE), and elongation at rupture (RE), as shown in Fig. 6a and Table 2. It demonstrated substantial strain hardening, with an UTS/$\sigma_{y,S0.2}$ ratio exceeding 1.6, a uniform elongation close to 33%, and high fracture ductility ($Z_{opt} = 58\%$). The (HEA+Ni)-S850 material (Fig. 6b) also showed excellent tensile properties, though

<caption>**Table 2**</caption>

Mechanical properties of the three materials, including the 0.2% and 0.5% proof stresses ($\sigma_{y,S0.2}$ and $\sigma_{y,S0.5}$), uniform elongation (UE), and ultimate tensile strength (UTS), all determined using mST specimens. The maximum load normalized by the initial cross-sectional area (F/S$_0$), measured from mNT specimens, is also reported.

<table><thead><tr><th>Material</th><th>σ<sub>y,S0.2</sub><br>(MPa)</th><th>σ<sub>y,S0.5</sub><br>(MPa)</th><th>UE<br>(%)</th><th>UTS<br>(MPa)</th><th>F/S<sub>0</sub> max (mNT)<br>(MPa)</th></tr></thead><tbody><tr><td>HEA-S850</td><td>422</td><td>437</td><td>33.2</td><td>736</td><td>909</td></tr><tr><td>(HEA+Ni)-S850</td><td>420</td><td>453</td><td>32.5</td><td>698</td><td>846</td></tr><tr><td>(HEA+Ni)-S600</td><td>338</td><td>404</td><td>29.7</td><td>685</td><td>887</td></tr><tr><td></td><td>476</td><td>516</td><td>17.3</td><td>653</td><td></td></tr></tbody></table>

slightly lower than HEA-S850, with a ductility $Z_{opt}$ of 46%. Conversely, the (HEA+Ni)-S600 material (Fig. 6c) had the lowest mechanical performance, despite its highest yield strength, with $Z_{opt} = 23\%$ only.

Regarding mNT specimens, the same ranking as for the mST specimens was observed (Fig. 6 and Tables 2 and 3). Tracking the minimum cross-sectional area (as illustrated in Fig. 6c) enabled a consistent comparison between different specimen types (Figs. 6 and 7). The increase in stress triaxiality (mNT geometry) led to increased maximum load and the ductility levels decreased by about one-third. The ranking among the three materials was more clearly highlighted than using the mST specimens.

## 3.3. Crack propagation behavior

Load vs. LLOD curves are reported in Fig. 8. Consistency was observed across tests for each material, although the initial crack length $a_0$ influenced both the initial elastic slope and the maximum load. The final opening angle is specified for each test and is shown in the last image acquired, after final unloading. This angle correlated with the final LLOD reached at the end of each test: a higher LLOD corresponded to a larger opening angle for a given specimen. Higher values indicate that more plastic deformation was required (thus consuming more energy and requiring a higher load) to propagate the crack over a given distance, suggesting that the specimen exhibited greater fracture toughness.

The J-Δa curves, representing the crack propagation resistance, are plotted for each test in Fig. 8d. All tests were valid, with the maximum value $J_{max}$, obtained for HEA-S850, reaching 194 kJ/m$^2$. Despite
<figure><img src="image_8.png" /><figcaption>Fig. 6. Tensile curves of (a) the HEA-S850, (b) the (HEA+Ni)-S850, and (c) the (HEA+Ni)-S600 materials; mST labels refer to uniaxial tension, mNT labels refer to notched tensile specimens. Inset in (c): example of the post-necking shape of a mST specimen with the detection of the minimum diameter $\Phi_{min}$ using the ET technique.</figcaption></figure>

<figure><img src="image_9.png" /><figcaption>Fig. 7. Reduction of area at fracture ($Z_{opt}$) obtained by optical measurements for the three materials and the two kinds of tensile specimens.</figcaption></figure>

some scatter for a given material, a clear ranking was observed: HEA-S850 showed the highest toughness, followed by (HEA+Ni)-S850, while (HEA+Ni)-S600 exhibited a lower toughness. These curves allowed to extract the full set of toughness properties, which are summarized in Table 3 for each material as an average over the two tested specimens.

The values of $J_{IC}$ revealed that the toughness of the HEA-S850 was more than twice that of (HEA+Ni)-S850, resulting in a $K_{JIC}$ value approximately 50% higher for the HEA-S850 material: 190 MPa$\sqrt{m}$ on average for HEA-S850, compared to 125 MPa$\sqrt{m}$ for (HEA+Ni)-S850. The (HEA+Ni)-S600 material, with $K_{JIC} = 41$ MPa$\sqrt{m}$, exhibited a toughness three times lower than that of (HEA+Ni)-S850. Notably, the material with the highest toughness also demonstrated the greatest crack propagation resistance, indicated by its tearing modulus: 159 N/mm$^2$ on average for HEA-S850, compared to 60 N/mm$^2$ for (HEA+Ni)-S850 and to 22 N/mm$^2$ for (HEA+Ni)-S600. This trend aligns with the observed opening angles for the three materials, also shown in Fig. 8a–c, with the largest angle for HEA-S850 (12.2° in test mDCT-1) compared to 4.5° in (HEA+Ni)-S600 (test mDCT-1) for a similar crack extension close to 1 mm.

### 3.4. Tensile fracture mechanisms

The fracture surface appearance of the tensile specimens varied with the material, but for a given material, it was the same for the mST, mNT, and mDCT geometries, i.e., whatever the stress triaxiality.

HEA-S850 specimens (Fig. 9) exhibited ductile fracture surfaces covered with dimples. Some larger dimples, a few micrometers in size (Fig. 9b), nucleated from non-metallic inclusions. Finer, submicrometric dimples (Fig. 9c) resulted from interfacial decohesion between the matrix and particles approximately 100–300 nm in size. This indicated a relatively high inclusion density, typically $1–2 \ \mu m^{-2}$, consistent with the high density of point-like features observed in optical micrographs (Fig. 3a).

Typical fracture surfaces of the (HEA+Ni)-S850 specimens (Fig. 10) showed two distinct regions: spherical “ball-like” features, and highly stretched matrix regions. The balls were covered with dimples (Fig. 10b), similar in size and morphology to those in HEA-S850 (compare Fig. 10c with Fig. 9c). The stretched regions fractured via internal necking and also occasionally displayed larger dimples (Fig. 10d). EDS measurements (Fig. 10e) revealed that the balls had a composition close to the HEA phase, while the matrix was richer in Ni. Elemental maps (Fig. 10e) confirmed these trends, despite shadowing artefacts. The size and shape of the balls were comparable to those of the master HEA powder particles.

In the (HEA+Ni)-S600 material (Fig. 11), the balls were even more prominent and in some cases nearly detached from the surrounding matrix (Fig. 11a). Their surfaces also exhibited dimples (Fig. 11b,c), which were shallower and finer than those in the other two materials. Inside these dimples, particles of 50–200 nm in size were observed, either isolated or clustered.

To further explore the damage mechanisms, longitudinal cross-sections of fractured specimens were examined via optical microscopy. At the macroscopic scale (insets of Fig. 12), the highest reduction in area at fracture was observed for HEA-S850 (Fig. 12a), while (HEA+Ni)-S850 displayed diffuse necking (Fig. 12b) and (HEA+Ni)-S600 exhibited limited necking (Fig. 12c), in agreement with the values reported in Table 3.

Damage initiation mechanisms were material-dependent but consistent across specimen types. In HEA-S850, very few microcracks were observed (Fig. 12a). In contrast, in both (HEA+Ni)-S850 and (HEA+Ni)-S600 materials, microcracks were frequently found at the HEA PPBs (Fig. 12b,c). EDS line profiles (Fig. 13) confirmed that cracks propagated within the interphase regions. In agreement with prior diffusion studies [18], Mn diffused more extensively than other elements, whose gradients only extended over 2.6–10 $\mu m$. The microcracks nucleated in the region with the steepest chemical gradient at HEA PPBs in alignment with the distribution of impurities shown in Fig. 3. Similar damage features were observed at crack tips in unbroken mDCT specimens (Fig. 14).
<figure><img src="image_10.png" /><figcaption>**Fig. 8.** Load vs. load line opening displacement (LLOD) of all mDCT specimens: (a) HEA-S850, (b) (HEA+Ni)-S850, (c) (HEA+Ni)-S600 materials. (d) *J* vs. *Δa* curves of all materials, *J*<sub>IC</sub> being the plane-strain fracture toughness.</figcaption></figure>

<caption>**Table 3**</caption>

Ductility and toughness of the three materials, including elongation at fracture (RE), reduction of area using (SEM) images (*Z*<sub>frac</sub>) for mST, reduction of area using optical measurements (*Z*<sub>opt</sub>) for mST and mNT, toughness *J*<sub>IC</sub>, critical stress intensity factor *K*<sub>IC</sub>, ductile tearing modulus *dJ/da*, and the coefficients *A* and *n* of the power law *J* = *A* · *Δa*<sup>*n*</sup>.

<table><thead><tr><th>Material</th><th>RE (%)</th><th>Z<sub>frac</sub> (mST) (%)</th><th>Z<sub>opt</sub> (mST) (%)</th><th>Z<sub>opt</sub> (mNT) (%)</th><th>J<sub>IC</sub> (kJ/m<sup>2</sup>)</th><th>K<sub>IC</sub> (MPa<sup>1/2</sup>/m)</th><th>dJ/da (N/mm<sup>2</sup>)</th><th>A</th><th>n</th></tr></thead><tbody><tr><td>HEA-S850</td><td>58.1</td><td>62.8</td><td>61.0</td><td>39</td><td>164</td><td>190</td><td>159</td><td>277</td><td>0.49</td></tr><tr><td>(HEA+Ni)-S850</td><td>53.3</td><td>59.9</td><td>55.2</td><td>29</td><td>71</td><td>125</td><td>60</td><td>117</td><td>0.38</td></tr><tr><td>(HEA+Ni)-S600</td><td>44.3</td><td>48.0</td><td>45.8</td><td>17</td><td>23</td><td>41</td><td>22</td><td>41</td><td>0.37</td></tr></tbody></table>

<figure><img src="image_11.png" /><figcaption>**Fig. 9.** Dimpled ductile fracture surface of a mST specimen of the HEA-S850 material. (a) General view; (b) enlarged view showing the two population of dimples; (c) detail of submicrometric particles inside the finer dimples.</figcaption></figure>

### 3.5. *Fatigue fracture mechanisms*

The fracture surfaces of broken mDCT specimens (Fig. 15a) revealed four successive regions: a smooth fatigue precrack zone, a rougher ductile tearing zone, a second fatigue propagation region, and a zone of final ductile fracture. Fatigue precracking was performed under decreasing stress intensity factors, from 14 to 11 MPa√m. The average crack growth rate during this stage was determined using the number of applied fatigue cycles and the observed extent of fatigue crack propagation. It ranged from 0.02 to 0.2 µm/cycle, which is consistent

with the striation spacing observed in the HEA-S850 material (Fig. 15b). Namely, in this material, the fatigue fracture surface appeared relatively flat with well-defined striations, typically spaced by about 0.2 µm (see inset).

In contrast, the fatigue fracture surface of the (HEA+Ni)-S850 material was rougher (Fig. 15c). It contained both regions with a striation pattern similar to that observed in HEA-S850, and regions composed of smooth facets a few micrometers in size (Fig. 15d). These features were consistent with the average grain size of this material (7.0 µm, Table 1).
<figure><img src="image_12.png" /><figcaption>**Fig. 10.** Fracture surfaces of the (HEA+Ni)-S850 material under tensile loading. (a–d) mNT specimen: (a) general view, (b) enlarged view of the region framed in (a); (c) detail of the fine-dimpled region at the surface of a ball-shaped feature; (d) detail of a large dimple in the stretched region. (e) Chemical analysis of the fracture surface of a mST specimen: general view, average chemical composition in regions A and B, and elemental maps of region C.</figcaption></figure>

<figure><img src="image_13.png" /><figcaption>**Fig. 11.** Dimpled ductile fracture surface of a mNT specimen of the (HEA+Ni)-S600 material. (a) General view; (b) enlarged view showing two ball-shaped features; (c) detail of the surface of a ball-shaped feature showing very fine, shallow dimples nucleated from tiny particles.</figcaption></figure>

<figure><img src="image_14.png" /><figcaption>**Fig. 12.** Longitudinal cross-section views of fractured tensile specimens. Optical microscopy, bright field contrast; HEA in light gray, Ni-matrix in yellow. (a) HEA-S850, (b) (HEA+Ni)-S600, (c) (HEA+Ni)-S600 materials. The arrows indicate microcracks at prior boundaries of HEA particles. The global views of fractured specimen halves are added as insets. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)</figcaption></figure>
<figure><img src="image_15.png" /><figcaption>Fig. 13. SEM-EDS chemical analysis profile across a damage microcrack in a mST specimen of the (HEA+Ni)-S850 material. (a) Secondary electron view and location of measurement points (small circles) along the profile (thick arrow). (b) Corresponding measured chemical compositions.</figcaption></figure>

<figure><img src="image_16.png" /><figcaption>Fig. 14. Longitudinal cross-section views of the crack tip of unbroken mDCT specimens. Optical microscopy, bright field (a–b) and Nomarski contrast (c). (a) HEA-S850, (b) (HEA+Ni)-S850, (c) (HEA+Ni)-S600 materials. Damage development around the HEA powder particles is indicated with arrows.</figcaption></figure>

To further investigate the nature of these facets, matching fracture surfaces, i.e., corresponding areas on both halves of the broken specimens were observed (Fig. 15e–f). Sets of parallel, thin lines were present on the facet surfaces, possibly corresponding to dislocation slip traces. Moreover, the orientations of these sets did not mirror one another for matching facets, suggesting that they belonged to different grains and that the facets originated from intergranular fracture. Some interfacial decohesion between the matrix and small particles was also observed (Fig. 15e–f), though it did not appear to drive the intergranular failure phenomenon.

A longitudinal cross-section from an unbroken mDCT specimen of the (HEA+Ni)-S850 material (Fig. 16a) confirmed that intergranular fracture occurred within the Ni-rich matrix. Hence, fatigue cracks in this material most probably propagated transgranularly across the HEA-rich regions and intergranularly along the grain boundaries of the Ni-rich matrix.

In the (HEA+Ni)-S600 material, the fatigue fracture surface exhibited similar ball-shaped features as those seen in the fracture surfaces of mST and mNT specimens (Fig. 15g). A close-up view (Fig. 15h) revealed numerous small holes but no clear dimples, suggesting that these voids may correspond to the same population of tiny particles previously observed within dimples after tensile fracture (Fig. 11c). The crack path (Fig. 16b) confirmed that fatigue cracks bypassed the HEA-rich regions and propagated around the HEA particles instead of across them.

As in the (HEA+Ni)-S850 material, matching surfaces (not shown here) revealed that the fatigue crack also propagated along the grain boundaries of the Ni-rich matrix. This was also suggested by the presence of small intergranular facets, Fig. 15(i–j), in agreement with the finer grain size of this material (1.6 µm in Table 1). Therefore, in the (HEA+Ni)-S600 material, fatigue cracks propagated along HEA PPBs and along grain boundaries within the Ni-rich matrix.

# 4. Discussion

## 4.1. *Comparison between the results from the different mechanical tests*

Due to the limited availability of the investigated exploratory materials, only one test per material was conducted, except for fracture toughness and for uniaxial tension of the HEA-S850 material, for which two tests were performed. The measurements for the HEA-S850 were compared to literature data, although for such newly developed chemically-architected alloys, no comparison with literature data was possible. Otto et al. [17] and Gludovatz et al. [15], reported about a CoCrFeMnNi HEA processed by melting, casting, and subsequent thermo-mechanical treatments. The grain sizes were 4.4 µm and 6 µm respectively, i.e., very similar to those found in the present study. The values of 0.2% proof stress, ultimate tensile stress, and fracture elongation were 346–400 MPa, 650–750 MPa, and 0.50–0.57, respectively, for [17] – for [15], compared to 420–422 MPa, 698–736 MPa, and 0.32–0.33 for HEA-S850, respectively (Table 2). Yield and ultimate strengths were in good agreement, while the lower elongation of HEA-S850 may result from oxide contamination, which is more easily avoided in the melting and casting processes used in [15,17]. A stress intensity factor $K_{IC}$ of 217 MPa$\sqrt{m}$ was measured in [17], slightly higher than that of HEA-S850 (190 MPa$\sqrt{m}$), again likely due to powder contamination. Overall, properties measured on minaturized HEA-S850 specimens well aligned with literature data from conventional HEA specimens.

Despite the limited number of repetitions, consistent trends could be identified. Yet, significant scatter was observed between the two replicates, both in elastic–plastic flow behavior (Fig. 6a, even before the onset of necking) and in crack propagation resistance (Fig. 8). This variability could partly be due to differences in SPS processing from
<figure><img src="image_17.png" /><figcaption>**Fig. 15.** Fatigue fracture surfaces of the mDCT specimens. (a) Typical macrograph, (HEA+Ni)-S600 material. (b) HEA-S850 material, fatigue striation, some of them indicated with white arrows. (c–f) (HEA+Ni)-S850 material: (c) general view, (d) a mixture of transgranular regions with fatigue striations and groups of facets, (e–f) matching surfaces showing that matching regions (similar symbols) did not exhibit mirror slip line patterns, suggesting intergranular fracture; white arrows indicate particle-matrix interfacial decohesion. (g–j) (HEA+Ni)-S600 material: (g) general view; (h) non-ductile fracture of HEA regions; (i) mixture of non-ductile fracture and facets; (j) details of facets and secondary cracking.</figcaption></figure>
<figure><img src="image_18.png" /><figcaption>Fig. 16. Fatigue crack path in (a) (HEA+Ni)-S850 and (b) (HEA+Ni)-S600 specimens. Some crack branches are indicated with arrows.</figcaption></figure>

one pellet to another. Here, the contribution of the multiscale and heterogeneous microstructure of the chemically-architECTured (HEA+Ni) materials is further analyzed.

The fine grain size ensured that a large number of grains were sampled in each tensile specimen. With a 20 µm average size for the HEA powder particles and a much finer grain size (<10 µm), the gauge section of a mST specimen included over one million grains and more than 4000 former HEA powder particles. This ensured that the elastic-plastic behavior was not affected by sampling effects. In fact, the two tensile curves of HEA-S850 overlapped up to 2% plastic strain (Fig. 6a). Moreover, the yield strength values obtained from both mST and mNT specimens followed the same ranking for small strain levels: (HEA+Ni)-S850 was softest, and (HEA+Ni)-S600 was strongest.

Experimental scatter was more pronounced in the fracture toughness tests (Fig. 8, Table 3). The amount of material sampled by the process zone in front of the crack in mDCT specimens increases with toughness. The crack front length was approximately 2.4 mm (distance between side grooves), i.e., 2400 times the maximum grain size (10 µm) and 1200 times the HEA powder particle size. Even considering a representative feature size of 100 µm (i.e., an aggregate of former HEA particles in the Ni-rich matrix), the crack front intersected at least 20 such zones.

Perpendicular to the crack front, the process zone size, $r_p$ can be estimated as $J_{IC}/\sigma_y$, where $\sigma_y$ is the average between the 0.2% proof stress and UTS. The results (Table 4) show that $r_p$ was of the same order of magnitude as key microstructural features such as aggregates of HEA particles. This made fracture toughness sensitive to sampling effects. As a result, the scatter in fracture toughness values, which was minimal for the HEA-S850 material, was more pronounced when $r_p$ became comparable to or smaller than these structural heterogeneities, as it was the case in (HEA+Ni)-S850 and even more so in (HEA+Ni)-S600.

Despite this variability, a consistent ranking of the fracture resistance of the three materials was established using miniature specimens (Table 3). Notched tensile mNT specimens already effectively distinguished the fracture behaviors, while relying on a simpler experimental procedure than the precracked mDCT fracture toughness experiments. They thus appeared to be an efficient tool for further material development. Although additional testing would be required for quantitative comparisons, the current results already demonstrated robust differences in behavior. Moreover, the observed damage and fracture mechanisms — along with their correlation to impurity distribution and interphase thickness — provided valuable insights for future microstructural optimization.

## 4.2. Effect of the chemical gradient and of the microstructure on tensile properties

Chemical architecture was designed to increase the resistance of the materials to plastic deformation. The effective extra-hardening induced by the interphase was already established by local nanoindentation measurements and full-field numerical simulations taking into account chemical gradients and strain gradient plasticity [14]. Despite grain refinement and solute diffusion during sintering [13], the Ni-rich matrix remained softer than the HEA phase, with a nanohardness lower by at least 30 pct in the (HEA+Ni)-S850 material [7]. Compression tests, which were interrupted for an engineering strain close to 0.2, measured the macroscopic resistance to plastic deformation and allowed a first ranking. The (HEA+Ni)-S600 alloy was identified as the most promising one, although a reduced work-hardening ability was noticed. The present tensile tests confirmed this ranking regarding their yield strength: the (HEA+Ni)-S600 material exhibited the highest value (based on $\sigma_{y,S0.2}$ for mST specimens), while the (HEA+Ni)-S850 material showed the lowest one. However, the fracture elongation and tensile strength, measured here for the first time, led to a very different ranking: both chemically-architECTured alloys were less ductile than the reference (HEA)-S850 alloy. Due to this strong decrease in ductility, even the ultimate tensile stress did not obey the same ranking as the compressive strength.

The present work demonstrated that hardness and compression tests alone were not able to give a complete picture of the resistance to plastic deformation of architECTured materials, especially those with a coarse-scale phase distribution. In the present tensile tests, the (HEA+Ni)-S600 material exhibited the highest resistance to ductile fracture (based on mST elongation and mNT reduction of area), while the (HEA+Ni)-S850 material showed the lowest one. This ranking was consistent with previous observations regarding work hardening [18]. Namely, the lower work hardening ability of the (HEA+Ni)-S600 material accelerated the onset of strain localization and fracture. This may be partly attributed to its higher fraction of softer Ni-rich matrix (28% vs. 18% in the (HEA+Ni)-S850 alloy), combined with a finer grain size both contributing to its earlier necking and reduced strain hardening capacity.

Although all three materials displayed high tensile ductility, both the reduction in area at fracture and the fracture toughness were significantly affected by the presence of impurity particles at prior HEA powder particle boundaries. These inclusions played a dominant role in the fracture mechanisms of the chemically-architECTured alloys, enabling early void nucleation and ductile cracking around ball-shaped features that correspond to former HEA powder particles. Eliminating
<caption>**Table 4**</caption>

Fracture-related parameters measured for the three materials, including average toughness $J_{IC}$, yield strength $\sigma_y$, and characteristic plastic zone size $r_p$. Ratios comparing $r_p$ to powder particle size (20 µm) and typical size of HEA particle aggregates (100 µm) are also provided.

<table><thead><tr><th>Material</th><th>HEA-S850</th><th>(HEA+Ni)-S850</th><th>(HEA+Ni)-S600</th></tr></thead><tbody><tr><td>Average value of J<sub>IC</sub> (kJ/m<sup>2</sup>)</td><td>164</td><td>71</td><td>23</td></tr><tr><td>σ<sub>y</sub> (MPa)</td><td>569</td><td>511</td><td>564</td></tr><tr><td>r<sub>p</sub> (µm)</td><td>288</td><td>139</td><td>41</td></tr><tr><td>Ratio r<sub>p</sub>/powder particle size (20 µm)</td><td>14.4</td><td>6.9</td><td>2.0</td></tr><tr><td>Ratio r<sub>p</sub>/HEA aggregate size (100 µm)</td><td>2.9</td><td>1.4</td><td>0.4</td></tr></tbody></table>

<figure><img src="image_19.png" /><figcaption>**Fig. 17.** SEM characterization of the raw HEA powder particles (i.e.: before acid treatment). (b) Surface of a particle.</figcaption></figure>

<figure><img src="image_20.png" /><figcaption>**Fig. 18.** SEM-EDS characterization of the sintered chemically-architected alloy with a large interphase (HEA+Ni)-S850. (a) EDS mappings of Ni concentration. (b) Concentration profile across the interphase indicated by an arrow on (a). The black dashed lines indicate the concentration and spatial limits of the interphase. (c) Phase mappings reconstructed from EDS mappings, after [32]. The scale is the same on (a) and (c).</figcaption></figure>

<figure><img src="image_21.png" /><figcaption>**Fig. 19.** EBSD orientation mappings of two sintered samples (a) a single-phase HEA for reference, (b) a chemically-architected HEA with a large interphase. The stereographic triangle is given, coding the sample vertical direction. The scale is the same on (a) and (b).</figcaption></figure>
these impurity particles — responsible for early damage initiation and for the reduced fracture resistance observed in the (HEA+Ni)-S850 and (HEA+Ni)-S600 materials — represents a critical step toward improving their ductility and toughness.

The difference in fracture behavior between the two (HEA+Ni) alloys can be rationalized by the steepness of their chemical gradients across the interphase around HEA-rich regions. In the (HEA+Ni)-S600 material, the gradient was approximately four times as steep as in (HEA+Ni)-S850, extending over 2.6 µm versus 10 µm (for all solute atoms except Mn) according to Table 1. A steeper chemical gradient leads to sharper mechanical mismatches near the interphase, increasing local stress concentrations and promoting earlier void nucleation. This interpretation is supported by lower $J_{0.2}$ values (Table 3), and by the observation of shallower dimples in the fracture surface of the (HEA+Ni)-S600 material, indicating faster void coalescence. The lower tearing modulus $dJ/da$ further confirms this trend.

Altogether, the steeper chemical gradient in the (HEA+Ni)-S600 material, combined with a high density of impurity particles, resulted in more localized damage and earlier fracture initiation, significantly reducing its resistance to ductile crack growth compared to the other two materials.

### 4.3. Effect of the chemical gradient and of the microstructure on the fatigue crack propagation behavior

Although fatigue crack growth rates were not quantitatively measured during the precracking procedures nor during post-tearing, the fractographic analysis revealed meaningful trends. In particular, the steepness of the chemical gradient appeared to play a major role in the physical fatigue crack propagation mechanisms.

In the HEA-S850 material, where no chemical gradient was present, the fatigue crack propagated in a transgranular way with well-defined striations. Impurity particles had little to no visible influence on crack propagation, in contrast to their impact on monotonic tensile fracture. Introducing a chemical gradient in the (HEA+Ni)-S850 material did not affect fracture within the HEA-rich regions, but led to intergranular fracture within the Ni-rich matrix, suggesting a lower grain boundary cohesion. This might be due to a less efficient sintering of the Ni powder. Indeed, previous studies have shown that achieving full densification using SPS requires more severe conditions for pure Ni powder than for HEA powder [7]. The chemical gradient might also have an effect, both close to the interphase and by introducing Mn in the bulk Ni-rich phase. This effect became more pronounced in the (HEA+Ni)-S600 material, where the steeper gradient redirected the crack path along the 3D impurity particle network, bypassing the HEA phase domains. These observations indicate that a steeper chemical gradient increased the sensitivity of fatigue crack propagation to both microstructural heterogeneities and impurity particles.

The underlying mechanism of intergranular fracture in the Ni-rich matrix remains unclear. While the observed facets showed evidence of crystal plasticity (slip bands), no secondary brittle phases were detected. A similar fracture mode in pure Ni has only been reported at room temperature in the presence of hydrogen gas [33]. Manganese segregation at grain boundaries may also play a role. The solubility limit of Mn in Ni is around 12 at.% at room temperature [34], so that at equilibrium, manganese atoms are expected to remain in solid solution. Nevertheless, grain boundaries are preferred diffusion paths at the considered sintering temperatures, so that Mn enrichment at grain boundaries of the Ni-matrix cannot be excluded. On the other hand, local Mn enrichment could disturb the magnetic order of Ni [35], potentially reducing boundary cohesion. This hypothesis is supported by prior work showing enhanced intergranular cracking in Mn-enriched austenitic steels [36].

Further investigation is needed to determine whether similar segregation or magnetic destabilization mechanisms are active in Mn-enriched Ni-rich matrices. If confirmed, tailoring the chemistry of the Ni-rich phase at grain boundaries could provide a way to improve its resistance to fatigue crack propagation.

## 5. Conclusions

For the first time, a consistent set of mechanical properties, under various loading conditions, was obtained for two exploratory chemically-architected high-entropy alloys (with either a steep or a smooth chemical gradient) and a reference single-phase high-entropy alloy (HEA), only available in limited quantities. Miniaturized specimen testing allowed determination of the tensile strength, ductility and fracture toughness. The key findings are summarized below:

* Chemically architecturing a HEA using Ni and powder metallurgy enhanced its resistance to plastic deformation at low strains (<10%). The steep-gradient alloy exhibited the highest yield strength (476 MPa), the reference HEA an intermediate value (420 MPa), and the smooth-gradient alloy the lowest one (338 MPa), consistently with previous compression results [18].

* Chemical architecturation reduced the maximum uniform elongation, affecting the tensile strength. The reference HEA exhibited the best performance (tensile strength >700 MPa, uniform elongation >32%), followed by the smooth-gradient alloy (685 MPa, 29.7%), then by the steep-gradient alloy (653 MPa, 17%).

* The fracture toughness decreased with chemical architecture, with values of 190, 125, and as low as $41 \text{ MPa}\sqrt{m}$ for the reference, smooth-gradient, and steep-gradient alloys, respectively.

* The interaction between chemical gradients and impurity particles at prior HEA particle boundaries triggered the early development of ductile damage despite the high tensile ductility of the Ni-rich matrix.

* Prior HEA particle boundaries acted as preferred fatigue crack paths in the steep-gradient alloy, whereas fatigue cracks propagated across HEA particles in the smooth-gradient alloy. In both architectured alloys, the Mn-enriched Ni-rich matrix was sensitive to intergranular fatigue crack propagation.

Up to now, chemical architecturation demonstrated an unprecedented strengthening effect through the addition of a softer phase. It enabled satisfactory ductility and damage resistance; however, the strength-ductility trade-off has not yet been overcome. Based on the present results, several opportunities for microstructural improvement can be identified, namely, reducing contamination in the parent HEA powder, further optimizing the steepness of the chemical gradient, and improving the strain hardening ability of the Ni-rich matrix and its resistance to intergranular fatigue crack propagation.

### CRediT authorship contribution statement

**Y. Madi:** Writing – review & editing, Writing – original draft, Validation, Supervision, Methodology, Investigation, Conceptualization.
**A-F. Gourgues-Lorenzon:** Writing – review & editing, Writing – original draft, Validation, Supervision, Methodology, Investigation, Conceptualization.
**B. Villeroy:** Writing – review & editing, Methodology, Investigation.
**M. Laurent-Brocq:** Writing – review & editing, Writing – original draft, Validation, Supervision, Methodology, Investigation, 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 paper.
## Acknowledgments

This work was supported by two French government grants managed by the ANR: the CAMEL JCJC project, focused on Chemically-Architected Metallic Alloys, and the Chair program MESSIAH (ANR-20-CHIN-0003), focused on Mini-tests for In-Service Monitoring of Structures with Application to Hydrogen Transport. We would also like to thank Léa Denax for her assistance with material processing, as well as Siebe Bongaerts, Conrad Quazza, and Anastasiia Suvorova for their assistance with cross-section specimen preparation.

## Appendix. Complementary microstructural data

See Figs. 17–19.

## Data availability

Data will be made available on request.

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