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Toward the Equiaxed Grain Microstructure in
CrMnFeCoNi High-Entropy Alloy Fabricated by
Directed-Energy Deposition

Minsoo Jin, Yunhui Chen, Bogdan Dovgyy, Peter Lee, Minh-Son Pham

► To cite this version:

Minsoo Jin, Yunhui Chen, Bogdan Dovgyy, Peter Lee, Minh-Son Pham. Toward the Equiaxed Grain Microstructure in CrMnFeCoNi High-Entropy Alloy Fabricated by Directed-Energy Deposition. Advanced Engineering Materials, 2024, 26 (15), pp.2301969. 10.1002/adem.202301969. hal-05098744

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https://hal.science/hal-05098744v1

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# Toward the Equiaxed Grain Microstructure in CrMnFeCoNi High-Entropy Alloy Fabricated by Directed-Energy Deposition

Minsoo Jin,* Yunhui Chen, Bogdan Dovgyy, Peter Lee, and Minh-Son Pham*

The columnar grains in additively manufactured alloys increase the tendency to form solidification cracks and cause anisotropy. Studying the effect of process parameters on microstructure development helps to guide the manufacturing of the equiaxed grain microstructure. First, the effect of process conditions on the melt pool dimensions using in situ synchrotron X-ray imaging and thermal profile and solidification condition using finite element simulation and calculation of thermodynamics phase diagrams of CrMnFeCoNi high-entropy alloy fabricated by directed energy deposition is studied. Increasing the laser power reduces the thermal gradient to solidification rate ratio, pushing the solidification closer to the columnar-equiaxed transition. Nevertheless, the simulations still indicate the columnar microstructure for all scan conditions in contrast to the experimental observation that shows single-wall samples built at 200 W consisted of dominantly equiaxed grains, whereas columnar grains are dominant in samples built at 100 W. It is believed that in addition to the effect of thermal gradient and solidification rate, the chemical segregation (Mn and Ni) during solidification may promote dendrite detachment, hence assisting the transition to equiaxed grains. The multitrack deposition results in more solid beneath a new melt pool, increasing the thermal gradient that promotes more columnar grains in comparison to single tracks.

## 1. Introduction

Additive manufacturing (AM) is a layer-by-layer fabrication process that has received significant attention as it enables alloy components to be fabricated with complex shapes. Directed energy

deposition (DED) is one of the mainly used AM methods that received significant attention thanks to their high efficiency in large-scale manufacturing and capability for repairing.[1-4] AM not only offers new and effective ways to manufacture components with great flexibility and freedom, but it also results in interesting microstructure, hence mechanical properties. In particular, the constitutional supercooling with high solidification rates and extremely high thermal gradients during the repeated deposition cycles of AM process form microstructures such as cells (and/or dendrites) whose boundaries contain dense dislocation networks.[5,6] Such dense dislocation networks are believed to enhance the mechanical strength of alloys fabricated by DED.[7-10] In addition, DED is not constrained by the size of the build chamber, enabling the fabrication of large metallic components with good consolidation. The application of DED process also enables the production of structural metallic components with enhanced strength without compromising ductility. However, the

manufacturing process often results in columnar grains with strong textures due to epitaxial growth antiparallel to the highest thermal gradient/build direction.[6,9,11] The presence of columnar grains often increases the tendency to form solidification cracks that detrimentally affect the material performance and

M. Jin, B. Dovgyy, M.-S. Pham
Engineering Alloys
Department of Materials
Imperial College London
Exhibition Road, South Kensington, London SW7 2AZ, UK
E-mail: jmins94@kaist.ac.kr; son.pham@imperial.ac.uk

M. Jin
Nanostructured Materials & Advanced Characterisation Laboratory
Department of Materials Science and Engineering
Korea Advanced Institute of Science and Technology
Daejeon 34141, Republic of Korea

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The ORCID identification number(s) for the author(s) of this article can be found under https://doi.org/10.1002/adem.202301969.

© 2024 The Author(s). Advanced Engineering Materials published by Wiley-VCH GmbH. This is an open access article under the terms of the Creative Commons Attribution License, which permits use, distribution and reproduction in any medium, provided the original work is properly cited.

DOI: 10.1002/adem.202301969

Y. Chen, P. Lee
Mechanical Engineering
University College London
Torrington Place, London WC1E 7JE, UK

Y. Chen, P. Lee
ESRF - The European Synchrotron
Grenoble CS 40220, France

Y. Chen
Department of Materials
University of Manchester
Manchester M13 9PL, UK

Y. Chen
RMIT Centre for Additive Manufacturing
RMIT University
Melbourne 3000, Australia

B. Dovgyy
International Additive Manufacturing Group
Domaniewska 3, 05-800 Pruszków (Polska), Poland
result in anisotropic and undesirable mechanical behavior.[12,13]
Therefore, the current challenge of AM techniques, including DED process, is achieving the desired grain structure—equiaxed grains. Solidification mode is governed by thermal gradient (G) and growth rate (R), and achieving the columnar-to-equiaxed transition (CET) of grains during solidification requires controlling these parameters.[14] Consequently, numerous studies have been undertaken to manipulate process parameters such as scan speed and laser power to achieve CET in various metals produced via DED. Notable examples demonstrating the achievement of partial CET include the studies for AlSi10Mg,[15] multiple titanium alloys (Ti6Al4V and Ti-3Al-8V-6Cr-4Mo-4Zr),[16] a CrCoNi medium entropy alloy,[17] and a CrMnFeCoNi high-entropy alloy (HEA).[18,19] However, many of these studies have been confined to single-track wall configurations, typically comprising multiple layers (about 3 to 5). Multitrack deposition increases the solid beneath a new melt pool, significantly changing the thermal profile of a new melt pool, and altering the CET in multitrack builds. Thus, the optimized conditions determined from single-track specimens should be applied to fabricate multitrack builds, followed by subsequent microstructure investigation and mechanical testing to study the effect of multitrack (and multilayer) deposition on the microstructure.

Many types of alloys such as stainless steels, nickel-based superalloys, titanium alloys, and aluminum alloys have been additively manufactured thanks to the advantages of AM processes.[3,6,20–24] Unfortunately, not all alloys were found suitable for AM processes due to the formation of defects such as porosity (e.g., lack of fusion pores, keyhole pores, and entrapped gas pores) and cracking (solidification cracking and liquation cracking). The presence of such defects deteriorated mechanical performances, in particular fatigue and creep resistances, making the AM fabricated alloys less favorable in load-bearing applications compared to other manufacturing processes. For example, the cracking in AM nickel-based superalloys (e.g., CSMX-4, Inconel 625, and Inconel 718) and AM aluminum alloys has been the main problem.[25–28] In contrast, CrMnFeCoNi HEA is the most investigated HEA fabricated by AM thanks to its excellent printability.[8,9,18,29–32] A narrow freezing range (170 K) and a low solidification gradient of this HEA also make it less prone to solidification and liquation cracking and ensure the fabricated components with a good consolidation.[9,29] Moreover, the CrMnFeCoNi exhibits promising mechanical properties and corrosion resistance over a wide range of temperatures including cryogenic conditions.[33–48] The fatigue life at room and elevated temperatures (e.g., at 550 °C) is comparable to austenitic face-centred cubic (FCC) steels.[49,50] The resistance to fatigue crack propagation is also comparable to austenitic steels and twinning-induced plasticity steels at room temperature (20 °C) and –75 °C.[51,52] These excellent mechanical properties suggest that this HEA holds a great potential as a structural alloy that could be applied in many applications. There has been an increasing interest in fabricating the alloy with laser AM.[8,9,18,29–31]

However, it is worth to note that previous studies on the CrMnFeCoNi HEA consistently reported a strong texture and columnar grain microstructure. Despite the investigations conducted by Xiang et al.[18] and Li et al.[19] focusing on the CET behavior of the CrMnFeCoNi HEA, the number of studies remains very limited. Hence, there is a strong need to study

the effects of AM process parameters on solidification conditions and subsequent microstructure development to provide a basis for process tuning to achieve equiaxed grain microstructure. Consequently, this current study first examines the melt pool dimensions in response to varying the process parameters using in situ synchrotron X-ray imaging with a high resolution to capture the interaction between laser beam and melt pool during deposition.[53,54] The third-generation synchrotron radiation source with high flux and high energy[55] enables to fast (millisecond to microsecond) X-ray imaging and diffraction of the laser-matter interactions, allowing to quantify the impact of laser power and scan speed on melt pool dynamics and dimensions. Subsequently, the examination of the melt pool dimensions provides information to validate finite element (FE) simulations of the melt pool using Rosenthal’s analytical model[56] to the thermal profile in the melt pool. The chemical segregation, cellular dendrite, and grain structure were investigated using a scanning electron microscope (SEM) equipped with electron backscattered diffraction (EBSD) and energy-dispersive X-ray spectroscopy (EDX). The cooling rate predicted by FE simulations is validated using a measurement of primary dendrite spacing and of chemical segregation that was examined both experimentally and using the calculation of thermodynamics phase diagrams (CALPHAD). The validated FE simulations are subsequently used to calculate the thermal gradient and solidification rates which were not provided by the CALPHAD and dendrite spacing measurement. Furthermore, this study also studies the effect of multitrack deposition on the microstructure and subsequently mechanical behavior. All of these investigations using in situ synchrotron X-ray imaging, FEA, CALPHAD, and ex situ microstructure characterization carried out in this study provide detailed information of thermal gradient, solidification rate, and cooling rates in melt pool that are crucial for obtaining insights into solidification microstructure, explaining the transition of microstructure from columnar to equiaxed in DED builds in response to varied laser powers and scan speeds. Such insights will shed light on paths toward achieving equiaxed microstructure in DED of the high entropy alloy.

## 2. Experimental Section

### 2.1. Materials and DED-AM: Blown Powder AM Process Replicator

Prealloyed CrMnFeCoNi powder was provided by H.C. Starck Surface Technology & Ceramic Powders GmbH. The nominal chemical composition of the powders is shown in Table 1. The powder particle size distribution was measured using SEM imaging, showing a variation from 3 to 45 µm (Figure 1), and with an average size of about 7.5 µm. Samples were built using a blown powder AM process replicator (BAMPR) which replicates a commercial laser DED-AM system and allows the use of synchrotron

<table><caption>Table 1. Nominal chemical composition of the HEA powder.</caption><thead><tr><th>Element</th><th>Cr</th><th>Mn</th><th>Fe</th><th>Co</th><th>Ni</th></tr></thead><tbody><tr><td>wt%</td><td>18.2</td><td>17.9</td><td>19.5</td><td>Balance</td><td>21.2</td></tr></tbody></table>
<figure><img src="image_5.png" /><figcaption>Figure 1. The CrMnFeCoNi powder particle size distribution and cumulative distribution curve (gray line).</figcaption></figure>

beamlines simultaneously (e.g., the12: the Joint Engineering, Environmental, and Processing (JEEP) at Diamond Light Source). [53,54,57] The system operates with a stationary laser beam and powder feeder in an argon atmosphere, and a 316 L stainless steel substrate with ≈1 mm width moves at a preset scan speed. The powder feed rate was 1 to 3 g min⁻¹. The laser power (W) is comparatively low (100–200 W) to generate a melt pool that allows sufficient transmission of the X-ray beam for imaging. [53,54,57] However, a small beam diameter (200 µm) ensures that the BAMPR uses a laser power density (power per unit area, W/cm²) that is comparable to industrial practice. Details of the apparatus, including the powder delivery system, can be found in refs. [53,54,57]. To study the change in microstructure with varying process parameters, single-track walls (4 layers) were fabricated with two laser powers (100 and 200 W) and five scan speeds (1.0–2.5 mm s⁻¹ with an interval of 0.5 mm s⁻¹).

## 2.2. Synchrotron Radiography, Image Processing, and Melt Pool Volume Quantification

The imaging process was conducted using a monochromatic X-ray beam generated by synchrotron radiation with an approximate average energy of 53 keV. The scintillator converts the X-ray beam to visible light, then is captured by a high-speed camera with an internal dark-field correction (model name: PCO.edge). A framerate of 200 frames per second and a pixel size of 3.24 µm were selected to achieve an optimized signal-to-noise ratio (S/N) image to capture melt pool features. The acquired radiographs were processed using Image] [58] and MATLAB. A flat-field correction was applied with a normalization of 100 flat images [59] to remove artifacts, noise, and noise variation between each pixel during the image acquisition process.

## 2.3. Microstructure Characterization

Samples were ground using 1200, 2400, and 4000 grit SiC papers followed by polishing in an OP-S solution with 0.4 µm colloidal silica particles for EBSD. EBSD analyses were carried out using a

ZeissSigma 300 SEM equipped with a field emission gun electron source at an accelerating voltage of 20 kV. EBSD analyses were performed using a step size of 3 × 3 µm² per pixel and grain boundaries were defined when a misorientation larger than 5° between neighboring pixels was detected. An Oxford Instruments EDX system with INCA software was used for chemical analysis. Polished samples were further electrochemically etched in 10% oxalic acid solution for 30 s at 10 V to observe the solidification microstructure. The dendritic microstructure of samples was analyzed using a ZeissAuriga SEM.

## 2.4. Cooling Rate Calculation Based on CALPHAD Models and Dendrite Spacing Measurement

Solidification behavior and chemical segregation under Scheil-Guliver approximation and a diffusion-enabled cooling at various cooling rates were evaluated via the CALPHAD using ThermoCalc (TC) 2020a. Scheil-Guliver solidification was carried out with a step of 1 K and the TCHEA2 thermodynamics database. Diffusion-enabled module solidification was carried out in Dictra module in TC (using TCNI8 and MOBIN15 databases) for three different cooling rates (10³, 10⁴, and 10⁵ K s⁻¹) to evaluate the extent of chemical segregation. Simulations were performed from the liquid phase at 2500 K down to 500 K in 2D using a circular domain with a diameter of 9 µm (analogous to the measured dendrite spacing). The simulation results were compared to the experimental SEM EDX measurements. To compensate for the edge effects, the obtained SEM EDX line-scan data were normalized by dividing the number of counts of a particular element by the total number of counts from the selected measurement point. The best agreement of the chemical segregation between the CALPHAD and experimental observation was used to estimate the cooling rate. Additionally, the cooling rate was quantified via the measurement of the primary dendritic arm spacing. The spacing of dendrites $λ_c$ was measured from the etched samples and was related to thermal gradient G and solidification rate R as follows: [11,60,61]

$$ \lambda_c = aG^{-m}R^{-n} \qquad (1) $$

where m and n are material constants, and a is an alloy-dependent factor ($a = 35^{[11]}$). The material constants m and n were often reported to be the same and 0.3 for the FCC phase. [62,63] The CrMnFeCoNi HEA also consists of a single FCC phase, and thus, its material constants m and n were assumed to be 0.3 for this study. Therefore, the dendrite arm spacing is proportional to the cooling rate, which is the product of the thermal gradient (G) and solidification rate (R).

## 2.5. Melt Pool Thermal Profile Analysis Using the Rosenthal Solutions

Rosenthal 3D solution [56] has been used widely to estimate temperature and heat flow in a moving melt pool. Figure 2a is a schematic diagram showing a melt pool moving during the deposition process with the parameters used for the 3D Rosenthal solution labeled. The melt pool dimensions were estimated using the 3D Rosenthal solution according to
<figure><img src="image_6.png" /><figcaption>**Figure 2.** a) A schematic diagram showing the heat flow of a moving melt pool during deposition. The calculated temperature of the melt pool and heat affected zone were calculated using the 3D Rosenthal solutions displayed in b) a view from the top and c) a view from the side along the scan direction. d) Schematic diagram of two adjacent melt pools (solid lines: solidification front, dashed lines: melting front). $v_{scan}$ is the scan speed and R is the solidification rate.</figcaption></figure>

$$T = T_0 + \frac{\eta P}{2\pi kr} \exp\left(-\frac{v_{scan}[r + (x - vt)]}{2a}\right) \quad (2)$$

where $T_0$ is the base plate temperature (K), P is the laser power (W), $v_{scan}$ is the laser scan speed (m s⁻¹), η is the laser absorptivity, k is thermal conductivity (W m⁻¹ K⁻¹), x is the spatial coordinate along the scanning direction (m), t is the time (s), r is the distance to the heat source ($r = \sqrt{(x - vt)^2 + y^2 + z^2}$), and α is the thermal diffusivity (m² s⁻¹). Calculations were performed in MATLAB on a grid of 1 × 2 × 1 mm³ with a step between evaluation points of 10⁻³ mm. The dimensions of the region with $T \ge T_m$ (liquidus temperature) were taken as the melt pool size. Calculations of the thermal gradient and cooling rate were performed at the melt pool boundary along which $T = T_m$, by differentiating Equation (2) with respect to r and t, respectively.

Derivation of the differentiated equations can be found in the ref. [64]. The melt pool dimensions examined using in situ X-ray imaging were used to validate the Rosenthal simulations. In addition, the cooling rate predicted by this analysis is also validated by values of the cooling rate calculated using the primary dendrite spacing and CALPHAD. Furthermore, the solidification rate R was calculated along the liquidus line and within a steady-state melt pool simulated by FEA, where $R = v_{scan}cosθ$, as shown in Figure 2a.

## 2.6. Mechanical Testing

DED builds with bidirectional multitracks were printed at the laser powers of 150 and 200 W and two scan speeds 1.0 and 1.5 mm s⁻¹ for tensile tests. These scan conditions were chosen
thanks to the promising equiaxed grain microstructure observed
in single-track walls at the studied parameters. Furthermore, the
scan speeds of 1.0 and 1.5 mm s⁻¹ provided laser energy densi-
ties which were sufficient to produce the builds large enough for
tensile test specimens. Tensile tests were conducted with the
multitrack samples instead of single wall tracks as the small
dimensions of single-track walls were limiting preparation and
setup for the mechanical testing. Hence, dog-bone specimens
with a gauge length of 19 mm and a thickness of 1.5 mm were
cut from the DED builds using an electrical discharge machine.
Monotonic tensile tests were performed using an Instron 300DX
at 22 °C with a strain rate of 10⁻³ s⁻¹. Three tensile tests were
performed for each print condition. The deformation behavior
during the tensile tests was recorded in 1 s intervals using a
Nikon DSLR camera with a resolution of 30 MP, translating to
0.2 mm per pixel of the image taken. The deformation field in
the gauge was then analyzed by digital image correlation using
the GOM Correlate software.

3. Results

3.1. Effect of Print Parameters on Melt Pools and Solidification Microstructure

Synchrotron X-ray images and graphs (Figure 3) show the
increase in melt pool size and volumes with a reduction in
the scan speed and an increase in laser power. It is because a
higher laser power increases the laser energy density, creating
a larger volume of melt pool. In particular, the melt pool volume
increased from 8 to 11 times at the same scan speed as the
laser power increased from 100 to 200 W (namely, 11 times at
1.0 mm s⁻¹, 11 times at 1.5 mm s⁻¹, 8.5 times at 2.0 mm s⁻¹,
and 8 times at 2.5 mm s⁻¹). EBSD maps of single-track walls
(Figure 4) show the grain structures at 100 W were strongly
columnar, whereas the grain morphology in the samples built
at 200 W was more equiaxed. Laser energy density is inversely

<figure><img src="image_7.png" /><figcaption>Figure 3. Synchrotron X-ray images showing melt pool geometries and volumes at different scan speeds (1.0, 1.5, 2.0, and 2.5 mm s⁻¹) and laser powers (100 and 200 W).</figcaption></figure>

proportional to the speed of a laser beam. A slower scan speed provides a longer laser beam exposure which enables to generate a larger melt pool volume. At 200 W, the melt pool volume at the scan speed of 1.0 mm s⁻¹ was more than twice larger than that at 2.0 mm s⁻¹. At the laser power of 100 W the melt pool volume increased from 0.01 to 0.04 cm³ with decreasing scan speed from 2.5 to 1.0 mm s⁻¹. Despite the change in the melt pool volume with varying the scan speed, the grain microstructure at the laser power of 100 W was predominantly columnar without a distinctive difference, Figure 4. In contrast, the melt pool volume at 200 W increased from 0.07 to 0.44 cm³ as the scan speed reduced from 2.5 to 1.0 mm s⁻¹. The grain microstructure was much more noticeably equiaxed at 200 W for all the scan speeds. Increasing the scan speed appears to result in more equiaxed grains. Nevertheless, it is very interesting to note that more equiaxed grains are seen for a condition of 1.5 mm s⁻¹ and 200 W than that of 2.0 and 2.5 mm s⁻¹.

**3.2. Chemical Segregation and Cooling Rate Analysis**

The constitutional supercooling and high ratio of the thermal
gradient to the solidification rate in AM (including DED) result
in the growth of columnar dendrites during solidification. The
columnar dendritic microstructure was observed in the DED
HEA samples (Figure 5). SEM EDX analysis revealed that Ni
and Mn were segregated to dendritic boundaries, while Cr, Co,
and Fe were segregated at the core of dendrites, similar to the
segregation observed for this alloy fabricated by laser powder
bed fusion (PBF).[10] The normalized atomic percentage bars next
to the element maps show stronger partitioning of Mn to the
interdendritic boundaries compared to Ni. Cr and Fe also exhib-
ited stronger partitioning in dendrites than Co. Moreover, the
presence of oxides in the DED HEA was observed. Both chro-
mium and manganese oxides in the AM HEA were reported,
but the chemical map suggests that the chromium oxides were
more dominant. **Figure 6a–e** shows CALPHAD simulations
of chemical segregation in solidification at the cooling rates of
10³, 10⁶, and 10⁵ K s⁻¹ under diffusion calculated by the
ThermoCalc/Dictra module and the Scheil–Gulliver calculation.
For comparison, the result of a line-scan EDX from the core of
the dendrite toward the edge (Figure 6f) is shown in Figure 6g.
The normalized atomic concentration along the line scan
from the dendrite core to the interdendritic region is in good
agreement with EDX mapping (Figure 5). Furthermore, the
experimentally observed chemical segregation (Figure 6g) showed
the best match with the simulation conducted at the cooling rate of
10³ K s⁻¹, suggesting the cooling rate was most likely in the range
of 10³ K s⁻¹.

Figure 7a shows the spacing of primary dendritic arms in
single-track walls built at different scan speeds and laser powers.
The dendritic spacings were measured at two locations in depo-
sition tracks: one was near the free surface, and another was in
the center of the track. Dendrites located at the center had
smaller spacings compared to those located near the free surface.
It has been well accepted in solidification that the spacing of pri-
mary dendrites λ<sub>c</sub> is related to the thermal gradient G and the
growth speed R as described by Equation (1).[11,60,61] The den-
drite spacing has been used to calculate the cooling rate in
<figure><img src="image_8.png" /><figcaption>**Figure 4.** EBSD maps (Inverse Pole Figure -scan direction) of the DED HEA single-track walls printed at different scan speeds 1.0, 1.5, 2.0, and 2.5 mm s⁻¹ and laser power of 100 W (red circle) and 200 W (black square).</figcaption></figure>

fusion-based manufacturing processes, including AM.[11,65,66] Therefore, the measurement of dendrite spacing was used to calculate the cooling rate in the builds. Figure 7b displays the cooling rates of different scan conditions estimated using the measured dendritic spacings and Equation (2). Heat transfer was faster at the center of a track via conduction through the solids (e.g., previously printed layer or substrate), resulting in a higher cooling rate at the center that leads to finer dendrites. In contrast, heat transfer was slower at the free surface due to the presence of gaps between loose powder particles that limited conduction. Thereby, a relatively lower cooling rate of dendrites resulted in coarser dendrites near the free surface. Nevertheless, the cooling rates calculated via the dendrite spacing in both regions regardless of scan conditions were in the range of 10³ K s⁻¹ which was in good agreement with the estimate of cooling rate using the CALPHAD simulation. Interestingly, the differences in dendritic spacing at the scan speed of 2.0 and 2.5 mm s⁻¹ (for the two laser powers) were insignificant. Therefore, the cooling rates calculated based on the dendritic spacing of these process parameters were also similar (Figure 7a). In contrast, at the lower scan speed, the cooling rate decreased significantly as laser power increased from 100 to 200 W.

### 3.3. Solidification Map

**Figure 8** shows a solidification map of the CrMnFeCoNi HEA estimated by Li et al.[19] with solidification conditions (thermal gradient and solidification rate). Li et al.[19] calculated the thermal gradient G and solidification rate R using a FE model. The figure also includes the solidification and thermal conditions predicted using the Rosenthal model for moving melt pools corresponding to the processing conditions considered in this study: 100 and 200 W and 1.0 and 2.0 mm s⁻¹. Each data point in the curves

shown in Figure 8 represents a solidification condition at a location along the longitudinal axis from head to tail of a melt pool. For all scan conditions, the solidification rate increased from the heat source toward the tail indicating solidification front (dendrites) advanced fast from the tail. The predicted thermal gradients at the lower laser power were above ≈10⁷ K m⁻¹, higher than those at 200 W (≈10⁶ K m⁻¹). Consequently, the cooling rates (G × R) along the melt pool at 100 W (10²–10⁴ K s⁻¹, from head to tail) were also higher than those at 200 W (10¹–10³ K s⁻¹) (Figure 8). Considering the fast growth of dendrites at the tail, the predicted cooling rates at the tail of melt pools at 200 W are in agreement with those calculated using the dendritic spacing (≈10³ K s⁻¹), whereas the ones at 100 W were slightly overestimated to be ≈10⁴ K s⁻¹ (Table 2). Figure 8 suggests that the combination of G (smaller than 4 × 10⁶ K m⁻¹) and R (larger than 2 × 10⁻⁴ m s⁻¹) is required for the transition from columnar to equiaxed microstructure of the CrMnFeCoNi.[19] The predicted values of R near the tail for all scan conditions (≈10⁻³ m s⁻¹) were close to the range for the equiaxed condition. However, the predicted values G for all scan conditions (>10⁶ K m⁻¹) were much higher than the range for the equiaxed condition. Microstructure for all scan conditions was expected fully columnar based on the prediction using the Rosenthal solution. The processing condition that is closest to the CET was 200 W and 1.0 mm s⁻¹ (Figure 7).

### 3.4. Grain Microstructure and Mechanical Properties

Figure 7 suggests that the processing condition that is close to the CET would be 200 W with the scan speed from 1.0 to 2.0 mm s⁻¹. In addition, Figure 3 suggests that the scan speed of 1.5 mm s⁻¹ should give more equiaxed grains than that of 2.0 mm s⁻¹. Therefore, two conditions of 200 W and the laser
<figure><img src="image_9.png" /><figcaption>**Figure 5.** SEM EDX mapping analysis showing chemical segregation behavior in the DED CrMnFeCoNi HEA sample fabricated at 200 W and 1.0 mm s⁻¹ with normalized atomic percentage of elements and indication of oxides (indicated by red circles).</figcaption></figure>

speed of 1.5 and 1 mm s⁻¹ were employed to fabricate bidirectional multitrack (and multilayer) builds for mechanical testing. The grain microstructure was more columnar in the multitrack builds in single-track ones (Figure 9 vs Figure 3). Equiaxed grains in the multitrack builds were mainly located at the regions between two consecutive layers (white dashed ellipses in Figure 9a). The microstructure in the multitrack build of 200 W and 1.0 mm s⁻¹ consists of two sets of columnar: one is along the build direction, and another is along the scan direction. The second set of columnar grains was coarser than the first one (Figure 9b). Multitrack (and multilayer) builds corresponding to the print condition resulting in columnar grains in single tracks (Figure 3) were also fabricated. Figure 3 and 7 suggest that a laser power of less than 200 W with a scan speed from 1.0 to 2.0 mm s⁻¹ should result in columnar grains. Therefore, a condition of 1.5 mm s⁻¹ and 150 W (instead of 100 W because the melt pool at 100 W was insufficient to produce a good consolidation) was used to generate tensile samples. This process condition resulted in columnar grains whose major axis appeared to be limited by the layer thickness (Figure 9c). Among all the three multitrack conditions, the build printed at 200 W and 1.5 mm s⁻¹

exhibited the best mechanical properties (Figure 9d). The yield stress, fracture strength, and elongation of this scan condition were, respectively, larger by ≈50%, ≈50%, and ≈400% compared to the other samples (Table 3). Moreover, the elastic modulus of the samples built at 200 W and 1.5 mm s⁻¹ was higher by ≈30% than those fabricated at the other print conditions, suggesting that this print condition resulted in the best consolidation in comparison to the other two. Compared to the other studies that investigated the same alloy fabricated by DED, the samples built at 200 W and 1.5 mm s⁻¹ exhibited a similar yield stress.[18,67–70] Nevertheless, the fracture strength was lower than those reported.

## 4. Discussion

### 4.1. Effect of Print Parameters on Solidification Microstructure and Mechanical Properties

The effects of laser power on the melt pool dimensions and solidification microstructure are shown in Figure 4 and 7a. The laser energy density at 200 W was double that at 100 W. While the
<figure><img src="image_10.png" /><figcaption>**Figure 6.** a–d) Chemical segregation predicted at cooling rates of 10³, 10⁴, 10⁵ K, and Scheil condition, respectively. e) Solidification behavior at the four different conditions. f) SEM image of dendritic structure where line-scan was performed with g) the experimental chemical segregation result showing the normalized atomic concentration along the scanned line (red arrow in f).</figcaption></figure>

energy density only increased by a factor of two, the melt pool volume significantly increased by a factor of about 8 to 11 depending on scan speed (Figure 3), suggesting the melt pool dynamics is influential in the fusion and penetration of laser beam. The melt pool volume increased much more significantly at 200 W as laser energy density exceeded 100 J mm⁻², in particular when scan speed was reduced to below 1.5 mm s⁻¹ (corresponding to an energy density of about 130 J mm⁻²). It will be discussed later that the melt pool dimensions are important in the solidification mode thanks to their effect on the thermal gradient.

The cooling rate predicted by the Rosenthal model (Figure 8) is in agreement with CALPHAD and experimental measurements of dendrite spacing. The Rosenthal simulation shows that increasing the laser power reduced the thermal gradient from ≈10⁷ to 10⁶ K m⁻¹. Thereby, the thermal gradient in the liquid G and subsequent G/R ratio decreases with an increase in the laser power for a given scan speed. It is worth noting that G/R governs the solidification mode (thereby the grain microstructure). The decrease in G/R promotes the nucleation of crystals, favoring the equiaxed microstructure. Increasing the laser energy decreases G, hence G/R for the same laser speed, explaining why the grain structure of single-track walls built at 200 W was more equiaxed than at 100 W (Figure 3). In addition, the increase in the scan speed increases the solidification rate, R, which also promotes more equiaxed and finer grains as it lowers the G/R. Therefore, the grain microstructure becomes finer and more equiaxed grains at 200 W, especially when comparing the microstructure for the scan speed of 2.5 and 2.5 mm s⁻¹ versus 1.0 mm s⁻¹ (Figure 4).

Interestingly, the dendritic spacings at the scan speeds of 2.0 and 2.5 mm s⁻¹ were almost the same for both 100 and 200 W

(Figure 7a), suggesting that the cooling rates (i.e., G × R) were similar at those conditions. Nevertheless, the thermal gradient (and G/R) ratios at 200 W were much lower toward the equiaxed mode (Figure 8). In fact, Figure 4 shows the increase of the power from 100 to 200 W resulted in more equiaxed grains at 2.0 and 2.5 mm s⁻¹. Therefore, the solidification rate at 200 W should be higher than that at 100 W to get almost the same cooling rates at both laser powers at 2.0 and 2.5 mm s⁻¹, while achieving the grain growth mode closer to the CET zone. In general, at the same laser power, a faster scan speed reduces the laser energy density and generates a smaller melt pool, leading to a higher thermal gradient as compared to a larger melt pool. Although the solidification rate also increases with the laser scan speed, the increase is not sufficient to favor the equiaxed microstructure, in particular at the lower power due to the extremely high thermal gradient (Figure 8). At 100 W, the melt pool sizes under various scan speeds were relatively small and the melt pool volume only increased from 0.01 to 0.04 cm³ when scan speed was reduced from 2.5 to 1.0 mm s⁻¹ (Figure 3). Because the melt pool size was small while the surroundings were cool, the thermal gradients at the 100 W conditions were noticeably high (Figure 8). Hence, the grain microstructure remained columnar with no change in the solidification mode at 100 W. In contrast, significant increase in the laser energy density by reducing the scan speed at 200 W (Figure 7a). Thereby, the rise in melt pool volumes at 200 W is much more substantial (0.07 to 0.44 cm³) as the scan speed is reduced from 2.5 to 1.0 mm s⁻¹. The increase in melt pool volume became most significant as scan speed was reduced to below 1.5 mm s⁻¹ and laser energy density surpassed 100 J mm⁻². The volume of the melt pool at 1.0 mm s⁻¹ was the highest, sufficiently heating the surrounding, and hence lowering the thermal gradient, promoting more equiaxed at 200 W.
<figure><img src="image_11.png" /></figure>

<figure><img src="image_12.png" /><figcaption>**Figure 7.** a) Measured dendritic spacing versus linear laser energy density (laser power divided by scan speed) at different processing conditions. b) Cooling rates in regions at the center and near the free surface of deposition track calculated based on the measured dendritic spacing.</figcaption></figure>

From the microstructure examination (Figure 4) and solidification mode analysis, two conditions of 200 W and the laser speed of 1.5 and 1 mm s⁻¹ give more equiaxed grain microstructure while the 150 W and 1.5 mm s⁻¹ conditions gave more columnar grains. Therefore, these three conditions were used to fabricate tensile samples (using multitrack deposition) to study the effect of different microstructure on the mechanical behavior. Compared to the single-track builds, the grain microstructure in the multitrack builds was more columnar than the single tracks (Figure 9 versus Figure 4). This is believed due to an increase in the area of solid beneath a new melt pool thanks to the multitrack deposition. Such an increase allowed more effective heat conduction, increasing the thermal gradient in the melt pool, hence limiting the CET in the multitrack builds. Therefore, the equiaxed grains were less dominant in comparison to the single-track builds.

Figure 9 shows that the 200 W and 1 mm s⁻¹ exhibited much different grain microstructure in comparison to the other two. As

<figure><img src="image_13.png" /><figcaption>**Figure 8.** The solidification map of the CrMnFeCoNi HEA calculated using a FE model[19] with cooling rates labeled. The thermal gradients and solidification rate in moving melt pools predicted based on the Rosenthal solution analysis are plotted to understand the solidification mode.</figcaption></figure>

<table><caption>**Table 2.** The predicted and measured cooling rates near the tail of melt pools for all scan conditions.</caption><thead><tr><td>Scan conditions (laser power and scan speed)</td><td>Cooling rate at melt pool tail [K s<sup>-1</sup>]</td></tr></thead><tbody><tr><td>100 W 1.0 mm s<sup>-1</sup></td><td>≈10<sup>4</sup></td></tr><tr><td>100 W 2.0 mm s<sup>-1</sup></td><td>≈4 × 10<sup>4</sup></td></tr><tr><td>200 W 1.0 mm s<sup>-1</sup></td><td>≈10<sup>3</sup></td></tr><tr><td>200 W 2.0 mm s<sup>-1</sup></td><td>≈3 × 10<sup>3</sup></td></tr></tbody></table>

shown in Figure 3 and 7, this condition exhibited the largest and deepest melt pools because it has the highest laser energy density. A deeper melt pool results in a larger area and more vertical for the rear side of melt pool. Because the epitaxial growth is anti-parallel to the heat flux, which is perpendicular to the melt pool boundary, a deep melt pool promotes more epitaxial growth of crystals from the rear side of melt pool and the growth direction of such epitaxial growth is aligned with the scan direction. Amongst all three multitrack build conditions, the sample built at 200 W and 1.5 mm s⁻¹ contained more equiaxed grains—this might contribute to better tensile responses (Figure 9d and Table 3). In addition, the superior tensile properties observed in specimens fabricated at 200 W and 1.5 mm s⁻¹ (particularly when compared to those built at 150 W and 1.5 mm s⁻¹) could also be attributed to the better consolidation at this condition owing to the higher laser energy density. However, the mechanical properties of 200 W and 1.5 mm s⁻¹ still outperformed those produced at 200 W and 1.5 mm s⁻¹ (which results in higher energy density than the former one by 50%). It is noteworthy that a higher energy density might cause excessive heat input, typically leading to the formation of keyhole porosity. Lower apparent elastic modulus measured experimentally (Figure 9d) suggested that the consolidation of 200 W and 1.5 mm s⁻¹ was lower than
<figure><img src="image_14.png" /><figcaption>**Figure 9.** a) IPF EBSD maps (with respect to the scan direction) showing grain structures in as-printed tensile specimens (multitracks) built at a) 200 W and 1.5 mm s⁻¹ where white dashed ellipses indicate the presence of fine equiaxed grains, b) 200 W and 1.0 mm s⁻¹, and c) 150 W and 1.5 mm s⁻¹. d) Engineering stress–strain curves of samples printed at different conditions.</figcaption></figure>

that of 200 W and 1.5 mm s⁻¹. Therefore, the energy density at 200 W + 1.5 mm s⁻¹ was likely more optimized among the considered conditions both in terms of microstructure and consolidation.

**4.2. Chemical Segregation and CET**

The CET in DED process is, in general, identified by the combination of the thermal gradient *G* and solidification rate *R*.
<table><caption>**Table 3.** The tensile properties of multitrack DED builds at different conditions.</caption><thead><tr><td>Scan condition</td><td>Yield stress [MPa]</td><td>Fracture strength [MPa]</td><td>Elongation [%]</td></tr></thead><tbody><tr><td>150 W 1.5 mm s<sup>-1</sup></td><td>127 ± 16</td><td>161 ± 24</td><td>5.1 ± 1</td></tr><tr><td>200 W 1.5 mm s<sup>-1</sup></td><td>190 ± 17</td><td>246 ± 31</td><td>19 ± 5</td></tr><tr><td>200 W 1.0 mm s<sup>-1</sup></td><td>125 ± 7</td><td>147 ± 6</td><td>4.5 ± 1</td></tr></tbody></table>

Section 4.1 also showed a change in both variables by controlling print parameters and melt pool volumes led to more equiaxed microstructure. However, the Rosenthal solutions analysis predicted the G of all considered print conditions was still high for the CET, and columnar grains were predicted regardless of the scan conditions (Figure 8). This result opposes the change in grain structure experimentally observed in the single wall builds (Figure 3), suggesting that either the estimate of the solidification mode on the basis of G and R shown in Figure 8 needs some improvements or other factors are not accounted for in the determination of the solidification mode. One of the other factors that are not included in the identification of solidification mode and affect the microstructure development during solidification is the dendrite fragmentation and subsequently detachment, promoting the nucleation of crystals, thereby leading to more equiaxed grains.[71,72] Helmer et al.[73] and Karimi et al.[74] also showed the formation of equiaxed grains in Inconel 718 fabricated by electron beam-PBF through the dendrite fragmentation mechanism. However, the convection flow in liquids and associated stress might be insufficient to mechanically break metallic dendrites.[72] The solute enrichment[72,75–78] or recalescence[79] also contributes to the dendrite fragmentation. In general, solute atoms rejected from dendrites into liquid flow downwards to the roots of side-branching dendrites.[72] The solutes accumulate forming a constitutional supercooled region which lowers the melting point of solid-liquid interface and remelts dendrite roots, contributing to detachment of dendrites.[19,76] Such detached dendrites flow into liquid ahead of the solidification front and tend to rise up due to their low density than surrounding liquids.[75] Floating dendrites promote the nucleation of new grains.[75] In addition, the chemical segregation lowers melting point (i.e., constitutional supercooling). The sudden crystallization of undercooled liquid causes a local increase in entropy and subsequent heat release which can further promote the detachment of dendrites.[79] Figure 5 and 6 show Mn and Ni partition to the liquid in the CrMnFeCoNi HEA samples upon solidification. The segregation of these elements could contribute to the dendrite detachment. Furthermore, the relatively slow scan speed in the larger melt pools at 200 W and 1.0–1.5 mm s⁻¹ provided sufficient time for dendrites to become coarser and cause more rejection of Ni and Mn into the liquid. Such more segregation likely promoted the fragmentation and detachment of dendrites. In addition, the stronger convection and fluid flow at 200 W than at 100 W could transport dendrite fragments within the melt pools to promote equiaxed grains. Increasing the scan speed should promote more detachment. However, it also increases the G, negating the CET effect resulting from the dendrite detachment, explaining why the equiaxed grain was only seen at 200 W and 1.5 mm s⁻¹, but not at higher scan speed

(Figure 4). Therefore, it should be a window of process parameters that lower G while promoting the dendrite detachment, i.e., the CET is a combination of G/R and the dendrite detachment.

## 5. Conclusion

This study examines the effect of varying the process conditions in particular laser power and laser speed on melt pool dimensions, microstructure, and mechanical properties in the CrMnFeCoNi high-entropy alloy fabricated by DED. While the laser energy intensity increases only by a factor of 2, the melt pool volume significantly increases by a factor of 8 to 11 depending on the same scan speed. Single wall samples built at 200 W had a grain microstructure consisting of dominantly equiaxed grains whereas columnar grains were dominant in the samples built at 100 W. Increasing the laser power generated larger and deeper melt pools and lowered thermal gradient G. The effect of reducing scan speed on the melt pool volume was more apparent at the higher laser power (200 W). The melt pool volume increased from 0.07 to 0.44 cm³ as the scan speed decreased from 2.5 to 1.0 mm s⁻¹. The ratio of the thermal gradient to the solidification rate decreased when the laser power increased from 100 W to 200 W, favoring the transition from columnar to equiaxed in particular at lower scan speeds.

The prediction of solidification mode based on the thermal gradient and solidification rate indicated a columnar dendritic mode for all the considered process conditions. However, experimental observation of single-track builds showed strong evidence of equiaxed grains for the 200 W, suggesting other important factors needed to be considered in the microstructure development in solidification. One of these factors is the chemical segregation that causes the constitutional supercooling and the dendrite detachment by remelting dendrite roots. For the high entropy alloy, Mn and Ni are partitioned to the liquid. The coarse dendrites observed at 200 W showed that these elements segregated more at a lower scan speed. Chemical segregation is likely contributing to the dendrite detachment and equiaxed grains seen in the 200 W conditions. Multitrack deposition resulted in less equiaxed grain microstructure in comparison to single-track builds due to increased solids beneath a melt pool that increases the thermal gradient.

## Acknowledgements

M.J., B.D., and Y.C. contributed equally to this work. The authors acknowledge the financial support provided by the Imperial-Tsinghua Research and Innovation Seed Fund at Imperial College London. The authors acknowledge the Research Complex at Harwell for the use of facilities and thank the Diamond Light Source (DLS) for providing the beamtime at beamline 112: JEEP (MT20096). This research was supported under MAPP: EPSRC Future Manufacturing Hub Manufacture using Advanced Powder Processes (EP/P006566/1), a Royal Academy of Engineering Chair in Emerging Technology (CiET1819/10), and Rolls-Royce plc. via the Horizon 2020 Clean Sky 2 WP5.8.1 programs and through the support of L.S.’s studentship.

## Conflict of Interest

The authors declare no conflict of interest.
**Data Availability Statement**

The data that support the findings of this study are available from the
corresponding author upon reasonable request.

**Keywords**

additive manufacturing, columnar-to-equiaxed transition, directed-energy
deposition, high-entropy alloy, synchrotron X-ray imaging

Received: November 23, 2023

Revised: March 11, 2024

Published online: July 2, 2024

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