# Deconvoluting cracking mechanisms in fusion processing of steel-copper multi-materials via Operando X-ray characterisation

Andaç Özsoy, William A Hearn, Steve Gaudez, Rijuta Jeswani, Yunhui Chen, Alexander Rack, Zoltan Hegedüs, Nicola Casati, Roland E Logé, Steven van Petegem

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Andaç Özsoy, William A Hearn, Steve Gaudez, Rijuta Jeswani, Yunhui Chen, et al. Deconvoluting cracking mechanisms in fusion processing of steel-copper multi-materials via Operando X-ray characterisation. Virtual and Physical Prototyping, 2025, 20 (1), 10.1080/17452759.2025.2526798. [hal-05157585](https://hal.science/hal-05157585v1)

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# Deconvoluting cracking mechanisms in fusion processing of steel-copper multi-materials via Operando X-ray characterisation

Andaç Özsoy, William A. Hearn, Steve Gaudez, Rijuta Jeswani, Yunhui Chen, Alexander Rack, Zoltan Hegedüs, Nicola Casati, Roland E. Logé & Steven Van Petegem

**To cite this article:** Andaç Özsoy, William A. Hearn, Steve Gaudez, Rijuta Jeswani, Yunhui Chen, Alexander Rack, Zoltan Hegedüs, Nicola Casati, Roland E. Logé & Steven Van Petegem (2025) Deconvoluting cracking mechanisms in fusion processing of steel-copper multi-materials via *Operando* X-ray characterisation, Virtual and Physical Prototyping, 20:1, e2526798, DOI: 10.1080/17452759.2025.2526798

**To link to this article:** https://doi.org/10.1080/17452759.2025.2526798

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© 2025 The Author(s). Published by Informa UK Limited, trading as Taylor & Francis Group

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Published online: 10 Jul 2025.

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# Deconvoluting cracking mechanisms in fusion processing of steel-copper multi-materials via Operando X-ray characterisation

Andaç Özsoy ID a,b, William A. Hearn ID a, Steve Gaudez a, Rijuta Jeswani a, Yunhui Chen ID c, Alexander Rack d, Zoltan Hegedüs ID e, Nicola Casati ID a, Roland E. Logé b and Steven Van Petegem a

aLaboratory for Condensed Matter, PSI Center for Photon Science, Villigen, Switzerland; bLaboratory of Thermomechanical Metallurgy (LMTM), École Polytechnique Fédérale de Lausanne (EPFL), Neuchatel, Switzerland; cRMIT Centre for Additive Manufacturing, RMIT University, Melbourne, Australia; dESRF – The European Synchrotron, Grenoble, France; eDeutsches Elektronen-Synchrotron (DESY), Hamburg, Germany

## ABSTRACT

This study investigates various cracking mechanisms and their prevalence in fusion processing of steel-copper multi-materials using operando X-ray diffraction and imaging during laser powder-bed fusion (LPBF) of 316L-CuCrZr multi-material. During this investigation, three main types of cracking were identified: (i) solidification cracking, (ii) metal-induced embrittlement (MIE), and (iii) liquation cracking. All cracking types are closely related to phase formation during processing and stem from two underlying mechanisms. First, liquid-liquid phase separation (LLPS) and the monotectic reaction in the 316L-CuCrZr system cause two liquids with vastly different solidification ranges to form, leading to solidification cracking. Second, LLPS and the monotectic reaction uniformly distribute Cu-rich liquid between the Fe-rich dendrites, leading to MIE and/or liquation cracking. Conducted based on the insights gained from the operando characterisation, further experiments showed that cracking can be drastically reduced by avoiding phase separation. However, the complete elimination of cracking necessitates chemical alterations in the material feedstock, indicating that while process adjustments can mitigate cracking, they may fail to fully prevent it. These findings serve as a guideline for understanding the underlying causes of cracking in steel-copper multi-materials, how process optimisation can effectively mitigate cracking, and to what extent such adjustments in processing can achieve this outcome.

## ARTICLE HISTORY

Received 16 April 2025
Accepted 21 June 2025

## KEYWORDS

Multi-material; Cracking;
Additive manufacturing;
Operando X-ray; Functionally
graded materials (FGM)

## 1. Introduction

Advancements in engineering have driven the adoption of multi-materials, commonly referred to as functionally graded materials (FGM), which integrate the properties of multiple materials into a single component. This allows for the creation of site-specific properties such as hardness, strength, thermal and electrical conductivity, and corrosion resistance. The applications of such multi-materials span across diverse fields such as aerospace [1–3] and medical industries [4,5].

Among multi-materials, the combination of steel and copper alloys has gained noticeable interest due to the complimentary properties of both alloys. Copper alloys provide high thermal and electrical conductivity together with excellent malleability and weldability. On the other hand, steels provide a combination of hardness, toughness, strength, corrosion resistance, wear resistance, and biocompatibility along with processability. The contrasting properties

of the two alloys make them an attractive option for multi-material combinations, especially for applications that require strong thermal and electrical properties coupled with high strength and corrosion resistance. Such applications include plastic injection moulds [6] and nuclear reactor components [7,8].

Although steel-copper multi-materials are an attractive combination, the stark differences in the thermophysical properties of the two materials reduce their processability. The most common conventional attempts to join the two alloys include fusion-based techniques such as welding and brazing. Here, there have been issues reported when copper alloys and steels are co-processed, with the most prominent issue being cracking at the joining interface of the two materials. This cracking is attributed to various mechanisms in the relevant literature including the mismatch in thermophysical properties such as coefficient of thermal

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expansion (CTE) [9,10], solidification cracking or hot tearing [10–14], copper contamination cracking (CCC) or liquid metal embrittlement (LME) [15–20], and liquation cracking [14,21]. While some of these studies point at the specifics of crack formation, some make rather general descriptions of the mechanism of cracking. For instance, solidification cracking refers to cracking at the terminal solidification stage, especially during welding and directional solidification of alloys with a wide freezing range [22]. Similarly, LME is described as the loss of strength when a dissimilar liquid metal is in contact with a parent metal. This is often observed due to the infiltration of the parent metal grain boundaries by the dissimilar liquid [23]. Being closely related to LME, CCC is the intergranular hot cracking in steels subjected to copper contamination prior to welding [19]. Having a similar appearance to LME in terms of the end result, liquation cracking refers to cracking due to the melting of low-melting-point phases often observed at the grain boundaries [22]. Rather than being a direct cracking mechanism, the CTE mismatch between two dissimilar materials creates thermally induced stresses at the interface during heating or cooling because of the difference between their thermal expansions, which may lead to cracking. Overall, a variety of explanations have been provided in the literature for the same or similar observations of cracking in steel-copper multi-materials. Although vast majority of studies report some degree of cracking, there are few exceptions that reported sound joints using conventional methods as well [15,24].

Following the conventional methods, the development of additive manufacturing (AM) technologies has opened a new window for the adoption of multi-materials, among which laser powder-bed fusion (LPBF), laser directed energy deposition (L-DED) and wire-arc additive manufacturing (WAAM) being the flagship technologies. As these technologies allow for precise control over the material feedstock during manufacturing, processing multiple materials has become feasible. Moreover, another dimension has been unlocked for engineering designs as locally adjusted compositions became possible within a single part. As can be imagined, AM of steel and copper alloys has received great interest in the research community as well. However, since they mostly rely on fusion joining, similar issues to those mentioned for conventional fusion processing have been reported.

Table 1 lists the literature review on processing of copper alloys and steels using both conventional and AM techniques. Irrespective of the AM technology used, the same three main mechanisms are proposed for cracking: (i) solidification cracking [25–29], (ii) liquid metal embrittlement [30,31], and (iii) liquation cracking [26]. As in the

conventional case, there are other explanations that usually describe process-related phenomena rather than the exact crack formation mechanism [32]. As seen in Table 1, multiple explanations exist for cracking together with observations such as phase separation, grain boundary infiltration, cracking at the grain boundaries or melt pools, and the dependence of cracking on composition. Moreover, many studies report multiple mechanisms or use different mechanisms interchangeably despite each describing different phenomena. The common point between these studies in the literature is that they rely on post-mortem and *ex situ* observations. Therefore, depending on the characterisation techniques and interpretation of the researchers, different conclusions have been drawn from the same observations. This indicates that the proposed mechanisms are not mutually exclusive, and that cracking may be incorrectly attributed to one mechanism while multiple mechanisms are active. Therefore, developing a thorough understanding of why, when, and how certain cracking mechanisms dominate is necessary to provide further insights. In particular, the use of *in situ* and *operando* monitoring would provide important information that could better decouple these mechanisms.

*In situ* and *operando* synchrotron X-ray characterisation of AM has become increasingly important as it allows for direct observation of complex phenomena during AM [40]. The capabilities of such methods include, but are not limited to, investigating keyhole and pore formation [41–45], cracking [46–49], phase evolution [50,51], thermal history evolution [52–55], algorithm training [56,57] and model validation [58]. As it provides real-time data, such experiments give deeper insights into complex processes and allow researchers to validate or discard competing hypotheses.

In this study, *operando* X-ray diffraction (XRD) and imaging were employed to distinguish various cracking mechanisms during multi-material LPBF of 316L stainless steel and CuCrZr copper alloy. The experimental insights were complemented with thermodynamic calculations and a range of post-mortem and *ex situ* analyses, including scanning electron microscopy (SEM), energy-dispersive X-ray spectroscopy (EDS), electron back-scatter diffraction (EBSD), and X-ray computed tomography (CT). The results revealed three distinct types of cracking that occur at different stages of the melting and solidification sequence, primarily driven by two main mechanisms. The relative contribution of these mechanisms was assessed under various processing conditions. Based on these insights, a novel approach has been explored as proof-of-concept to mitigate cracking. The challenges associated with fusion-based processing of steel-copper multi-materials exhibit similar characteristics across different techniques, and LPBF is no exception.
<figure><img src="image_12.png" /><figcaption>**Table 1.** Literature review on co-processing of copper alloys and steels using conventional and AM methods.</figcaption></figure>

<table><thead><tr><th>Process</th><th>Material combination</th><th>Observations</th><th>Proposed mechanism</th><th>Ref.</th></tr></thead><tbody><tr><td>Arc Welding</td><td>S355 Steel/CuAl10, CuSi3Mn1</td><td>Cracking at the steel grain boundaries, no cracking at high Si contents</td><td>CCC/LME</td><td>[15]</td></tr><tr><td>Arc Welding</td><td>1103 Steel/Cu</td><td>Phase separation and cracking</td><td>Solidification Cracking</td><td>[11]</td></tr><tr><td>Arc Welding</td><td>Various Steels/Cu</td><td>Cracking at the steel grain boundaries.</td><td>CCC/LME</td><td>[16]</td></tr><tr><td>Brazing</td><td>Various Stainless Steels/Cu-based Brazing Alloys</td><td>Cracking at the grain boundaries. No cracking for ferritic stainless steel.</td><td>CCC/LME</td><td>[17]</td></tr><tr><td>Brazing-Welding</td><td>S06 SS/S201 filler wire/CuCr0.8</td><td>Prevention of cracking by formation of a ferrite band</td><td>-</td><td>[33]</td></tr><tr><td>Stress-rupture test</td><td>Various Steels/Cu</td><td>Cracking at the steel grain boundaries/No cracking for certain compositions</td><td>CCC/LME</td><td>[19]</td></tr><tr><td>Furnace Melting of Cu on Steel</td><td>Various Steels/Cu</td><td>Cracking at the steel grain boundaries. More prominent cracking close to melting point of Cu due to lower interface energy</td><td>CCC/LME</td><td>[20]</td></tr><tr><td>Electron Beam Welding</td><td>EUROFER97 Steel/CuCrZr</td><td>Phase separation without cracking</td><td>-</td><td>[24]</td></tr><tr><td>Electron Beam Welding</td><td>304, 304L, 316L SS/Cu</td><td>Phase separation and cracking at the grain boundaries</td><td>CCC/LME</td><td>[18]</td></tr><tr><td>Laser Welding</td><td>201 SS/Cu</td><td>Phase separation and cracking</td><td>CTE mismatch</td><td>[9]</td></tr><tr><td>Laser Welding</td><td>303 SS/Cu</td><td>Cracking at the melt pools</td><td>Solidification Cracking/CTE mismatch</td><td>[10]</td></tr><tr><td>Laser Welding</td><td>304 SS/Cu</td><td>Phase separation and Cu infiltration into the steel part</td><td>Hot Cracking</td><td>[12]</td></tr><tr><td>Laser Welding</td><td>304 SS/Cu</td><td>Cracking at the steel grain boundaries</td><td>Liquation Cracking</td><td>[21]</td></tr><tr><td>Laser Welding</td><td>304 SS/Cu</td><td>Cracking for compositions between 10 and 50 wt.% Cu</td><td>Solidification Cracking</td><td>[13]</td></tr><tr><td>Laser Welding</td><td>304 SS/Cu</td><td>Cracking at the steel grain boundaries</td><td>Grain Boundary Liquation/ Solidification Cracking</td><td>[14]</td></tr><tr><td>LPBF</td><td>316L SS/Cu</td><td>Copper-filled cracks</td><td>Solidification Cracking</td><td>[25]</td></tr><tr><td>LPBF</td><td>316L SS/CuCrZr</td><td>Phase separation and cracking at the interface along the build direction</td><td>-</td><td>[34]</td></tr><tr><td>LPBF</td><td>316L SS/CuSn10</td><td>Phase separation and cracking for compositions between 5-50 wt.% CuSn10-316L</td><td>Liquation and Solidification Cracking</td><td>[26]</td></tr><tr><td>LPBF</td><td>316L SS/CuSn10</td><td>Phase separation and cracking at the interface along the build direction</td><td>Excessive Heat Input</td><td>[32]</td></tr><tr><td>LPBF</td><td>316L SS/CuSn10</td><td>Phase separation and cracking at the interface along the build direction</td><td>LME</td><td>[30]</td></tr><tr><td>LPBF</td><td>316L SS/CuCrZr</td><td>Cracking at the interface. Mitigation of cracks when using a ring-mode laser beam shape due to shallower melt pools</td><td>CCC/LME</td><td>[31]</td></tr><tr><td>LPBF</td><td>904L SS/CuSn10</td><td>Cracking at the Fe-rich side of the interface with phase separation</td><td>Solidification Cracking/CTE mismatch</td><td>[28]</td></tr><tr><td>L-DED</td><td>316L, H13/CuBe</td><td>Copper network and cracking at the grain boundaries</td><td>Solidification Cracking</td><td>[29]</td></tr><tr><td>L-DED</td><td>316L/Monel 400</td><td>Copper segregation and cracking at the grain boundaries</td><td>Solidification Cracking</td><td>[27]</td></tr><tr><td>L-DED</td><td>15-5 PH SS/CuAl10Ni5Fe5</td><td>No cracking, formation of Al-Cu intermetallics at the interface</td><td>-</td><td>[35]</td></tr><tr><td>WAAM</td><td>HSLA (ER-120S-G)/CuAl10</td><td>No cracking, formation of Al-Cu intermetallics at the interface</td><td>-</td><td>[36]</td></tr><tr><td>WAAM</td><td>C25 Steel/CuAl8</td><td>No Cracking, limited mixing at the interface</td><td>-</td><td>[37]</td></tr><tr><td>WAAM</td><td>HSLA (ER-120S-G)/ CuAl10Ni5Fe5</td><td>No Cracking, precipitation and grain refinement at the interface</td><td>-</td><td>[38]</td></tr><tr><td>WAAM</td><td>C25 Steel/CuAl8</td><td>Prevention of cracking using bypass current</td><td>-</td><td>[39]</td></tr></tbody></table>

Therefore, the findings of this study are expected to provide valuable insights and potential new strategies for improving the fusion-based processing of steel and copper alloys across various processing routes.

## 2. Experimental procedure

### 2.1. Raw materials

Gas-atomised CuCrZr powder with $D_{10}=38$ µm, $D_{50}=49$ µm and $D_{90}=61$ µm, 316L stainless steel powder with $D_{10}=20$ µm, $D_{50}=32$ µm and $D_{90}=51$ µm particle size distributions were used in the study together with baseplates made of 316L. CuCrZr powder was supplied by m4p Material Solutions GmbH (Magdeburg, Germany) and 316L powder was supplied by Carpenter Additive (Philadelphia, USA). Table 2 shows the chemical compositions of the powders and the baseplate.

### 2.2. LPBF equipment

Operando X-ray diffraction and imaging experiments, as well as *ex situ* prints have been realised using the miniSLM, the miniature LPBF equipment developed at Paul Scherrer Institute (PSI), Switzerland. The equipment has been used previously for *in situ* and *operando* X-ray diffraction and imaging experiments [47,52]. The technical details of the miniSLM can be found elsewhere [59].

<table><caption>**Table 2.** Elemental composition of the material feedstock.</caption><thead><tr><th>Material</th><th colspan="6">Elemental composition (wt.%)</th></tr></thead><tbody><tr><td>316L SS (Baseplate)</td><td>Cr</td><td>Ni</td><td>Mo</td><td>Mn</td><td>Si</td><td>C</td><td>Fe</td></tr><tr><td></td><td>17.7</td><td>9.8</td><td>2.4</td><td>1.8</td><td>0.7</td><td>&lt;0.03</td><td>Bal.</td></tr><tr><td>316L SS (Powder)</td><td>Cr</td><td>Ni</td><td>Mo</td><td>Mn</td><td>Si</td><td>C</td><td>Fe</td></tr><tr><td></td><td>18.9</td><td>12.8</td><td>2.7</td><td>1.2</td><td>1.0</td><td>&lt;0.03</td><td>Bal.</td></tr><tr><td>CuCrZr</td><td>Cr</td><td>Zr</td><td>O</td><td>Cu</td><td></td><td></td><td></td></tr><tr><td></td><td>0.9</td><td>0.07</td><td>0.04</td><td>Bal.</td><td></td><td></td><td></td></tr></tbody></table>
<figure><img src="image_13.png" /></figure>

## 2.3. Synchrotron X-ray characterisation

### 2.3.1. Operando X-ray diffraction

Two different operando XRD experiments were conducted, as schematically shown in Figure 1(a,b). The first diffraction experiments shown in Figure 1(a) were done at Materials Science (MS) beamline at Swiss Light Source (SLS) in Villigen, Switzerland. The X-ray energy and nominal beam size were set to 16.2 keV and 150 × 80 µm² (H × V), respectively. A thin layer of CuCrZr powder was recoated on the 316L baseplate for successive layers, and the X-ray beam was positioned at the top surface of the powder. During printing, diffraction patterns were collected in reflection mode using an Eiger1M detector (PSI, Switzerland) at a frequency of 20 kHz (exposure time = 0.05 ms), where the baseplate was set at an angle of 18° with respect to the X-ray beam (around X-axis). A rectangular area of 4 × 4 mm² (X × Y) was scanned in each layer using the parameters given in Table 3, using 67° rotation between layers.

The second set of diffraction experiments shown in Figure 1(b) were done at P21.2 beamline at PETRA III, Deutsches Elektronen-Synchrotron (DESY) in Hamburg, Germany. The X-ray energy and nominal beam size were set to 52 keV and 150 × 30 µm² (H × V),

respectively. Before the experiment, a rectangular 316L part of 3 × 0.5 × 3 mm³ (X × Y × Z) was built on the 316L baseplate. Then, for a total of 3 layers, CuCrZr powder was recoated on top of the part using the recoater of the miniSLM, the baseplate was raised and cleaned from the surrounding powder, and the X-ray beam was positioned right below the powder layer so that the diffraction rings of both materials were visible. Diffracton patterns were collected in transmission mode at 4 kHz (exposure time = 0.25 ms) using two Eiger1M (DECTRIS, Switzerland) detectors that were positioned side-by-side, where the miniSLM was tilted 5° with respect to the X-ray beam (around X-axis). The laser was positively defocused to obtain a large and divergent beam to minimise penetration into the 316L part. Then, gradual melting of CuCrZr was done by scanning the same area over 100 repetitions without depositing powder. The area was scanned bidirectionally along the long edge of the rectangular part (X-axis). It was aimed to melt CuCrZr while avoiding melting 316L part below, see Table 3 for details. The process parameters provided have been determined over several trials during the beamtime, where it was ensured that FCC-Cu peaks partially disappeared while FCC-Fe peaks remained fully visible.

<figure><img src="image_14.png" /><figcaption>**Figure 1.** Schematic configurations of (a) operando X-ray diffraction experiments in reflection (MS, SLS) and (b) transmission modes (P21.2, PETRA III). (c) Schematic configuration of the operando X-ray imaging experiments (ID19, ESRF).</figcaption></figure>

<table><caption>**Table 3.** LPBF process parameters used in the experiments.</caption><thead><tr><th>Experiment</th><th>Material</th><th>Power (W)</th><th>Scan Speed (mm/s)</th><th>Hatch Distance (µm)</th><th>Layer Thickness (µm)</th><th>Spot Size (µm)</th><th>Reps.</th></tr></thead><tbody><tr><td>Diffraction - MS</td><td>CuCrZr</td><td>460</td><td>600</td><td>70</td><td>30</td><td>45</td><td>-</td></tr><tr><td rowspan="2">Diffraction - P21.2</td><td>316L</td><td>150</td><td>250</td><td>50</td><td>30</td><td>45</td><td>-</td></tr><tr><td>CuCrZr</td><td>75</td><td>250</td><td>50</td><td>30</td><td>350</td><td>100</td></tr><tr><td rowspan="2">Imaging - ID19</td><td>316L</td><td>140</td><td>250</td><td>50</td><td>30</td><td>45</td><td>-</td></tr><tr><td>CuCrZr</td><td>350 (L)/460 (H)</td><td>600</td><td>70</td><td>30</td><td>45</td><td>-</td></tr><tr><td rowspan="2">X-ray CT - ID19</td><td>316L</td><td>200</td><td>800</td><td>50</td><td>30</td><td>45</td><td>-</td></tr><tr><td>CuCrZr</td><td>460</td><td>600</td><td>70</td><td>30</td><td>45</td><td>-</td></tr><tr><td rowspan="3">Controlled Melting</td><td>316L</td><td>200</td><td>800</td><td>50</td><td>30</td><td>45</td><td>-</td></tr><tr><td>CuCrZr (Control)</td><td>460</td><td>600</td><td>70</td><td>30</td><td>45</td><td>-</td></tr><tr><td>CuCrZr (Method Validation)</td><td>200</td><td>800</td><td>50</td><td>30</td><td>350</td><td>36</td></tr></tbody></table>

Spot size is defined based on the 1/e² intensity level. Repetitions are on a per-layer basis without powder deposition.
<figure><img src="image_15.png" /></figure>

The Python library pyFAI [60] was used to azimuthally integrate the collected 2D diffraction patterns to obtain 1D XRD spectra. NaCl and LaB₆ powders were used as the references to calibrate detector position and tilts for MS and P21.2 beamtimes, respectively. Background subtraction was performed by fitting a sixth-degree polynomial to the smallest of every Δ2θ = 0.095° interval.

The elastic lattice strain ε during printing was calculated considering the relative change in peak positions as:

$$ \epsilon = -\cot\left(\frac{2\theta}{2}\right) \frac{2\theta - 2\theta_0}{2} \quad (1) $$

where 2θ is the centre of mass of the peak and 2θ₀ is the reference position of the peak at room temperature. Under the assumption that the lattice strain is mainly influenced by thermal effects at high temperature [61], the average temperature T in the probed volume can be estimated by:

$$ T = \alpha_L(\epsilon) \quad (2) $$

where α_L(ε) is the strain-dependent thermal expansion function to convert strain into temperature. α_L(ε) was determined by fitting a third-degree polynomial (R² = 1) to the data presented in [62] and [63] for 316L and CuCrZr (assumed pure Cu). For the phases formed during solidification with intermediate compositions, ThermoCalc was used to generate CTEs. (Cu, Fe, Cr and Ni were considered for simplicity). Fitted functions are available in the Appendix together with the corresponding temperature vs. strain graph, see Figure A1.

### 2.3.2. Operando X-ray imaging and computed tomography

Operando X-ray imaging (radiography) experiments shown in Figure 1(c) were done at ID19 beamline at the European Synchrotron Radiation Facility (ESRF) in Grenoble, France. A white X-ray beam with a peak energy of 45 keV was used. A CMOS-based fast camera (Photron, type: SAZ, Japan) lens-coupled (10× magnification, 0.28 numerical aperture, Mitutoyo, Japan) to a 250 µm thick LuAG:Ce crystal scintillator was used for image acquisition. Acquisition rate was set to 40,000 images/s and the images were collected with a resolution of 2.2 µm/pixel during printing.

Before the experiment, a rectangular 316L part of 10 × 0.5 × 3 mm³ (X × Y × Z) was built on the 316L base-plate. This part was then polished using hand files and sandpaper (#1200) to remove powder stuck on the side walls to obtain clear images. Then, the powder was changed to CuCrZr, and recoating was done using the recoater of miniSLM to control the layer height. Powder around the parts was removed and the build plate was raised for imaging. Two different parameters

with low (L) and high (H) power as given in Table 3 were respectively used to melt CuCrZr powder on separate regions on the 316L part. Laser scanning was done bidirectionally along the short edge (Y) of the parts to observe crack formation. FIJI (ImageJ) [64] was used for flat field correction, sharpening and brightness/contrast adjustments on the collected images.

X-ray computed tomography (CT) was also performed on a 316L-CuCrZr part of dimensions 3.5 × 0.75 × 2 mm³ (X × Y × Z), ex situ printed using the parameters in Table 3, where only seven layers of CuCrZr were printed in total. A peak X-ray energy of 65 keV and an sCMOS camera (pco.edge, type: 5.5, PCO AG, Germany) lens-coupled (10× magnification, 0.28 numerical aperture, Mitutoyo, Japan) to a 25-µm-thick LuAG:Ce scintillator was used for acquisition. These yielded a voxel size of 0.6³ µm³. The tomograms were reconstructed using the Tomwer software (v1.3.13). Flat field correction and normalisation were used for reconstruction. DragonFly (v2024.1) was used to remove ring artefacts and to segment cracks and pores. Cracks were separated from pores using a sphericity upper limit of 0.5. Crack length was calculated from the segmented data using the longest length function in DragonFly.

### 2.4. Ex situ prints

To validate the hypotheses built on the results obtained from the synchrotron experiments, two ex situ parts were printed to observe the difference between their cracking behaviour. For both, two 316L parts of 4 × 4 × 4 mm³ were initially built using the set of parameters given in Table 3. The control sample was built using the optimal parameter set developed for printing CuCrZr with the highest possible density. The method validation sample was printed by positively defocusing the laser and gradually melting CuCrZr over many repetitions within a single layer, similarly to the P21.2 sample. Ten layers of CuCrZr were built for both samples using the parameters given in Table 3. The parameters for method validation build have been determined by several trials to ensure that sufficient bonding was achieved at the interface while minimising Cu-infiltration and cracking.

### 2.5. Computational thermodynamics

To understand the expected phase evolution during printing, ThermoCalc® [65] with TCFE12 database was utilised to calculate the 316L-CuCrZr material-material isopleth and perform Scheil-Gulliver solidification calculations using the classical model. These calculations were done with 10 wt.% intervals considering the
<figure><img src="image_16.png" /></figure>

mixtures of 316L and CuCrZr to predict the changes in the solidification path with respect to the composition under fast cooling. FCC, BCC and liquid phases were taken into consideration based on the experimentally observed phases from the *operando* XRD measurements. Fe, Cu, Cr, Ni, Mo, Mn, and Si were considered in the calculations as the main alloying elements. TCFE12 database includes the complete set of thermodynamic data for the interactions between Fe, Cr, Ni, Mo, Mn, and Si within the composition range used in the study. Although the recommended maximum composition of Cu is given as 5 wt.% for this database, all binary combinations of the selected elements and ternary combinations Cr-Cu-Fe, Cr-Cu-Mo, Cr-Cu-Ni, Cu-Fe-Mn, Cu-Fe-Mo, Cu-Fe-Ni, Cu-Fe-Si, Cu-Mn-Ni have been assessed in the database. Therefore, the calculations presented in this study are expected to be accurate with these specific elements and the selected compositions.

## 2.6. Post-mortem characterisation

Samples were cut from the baseplates using a wire saw and hot mounted in resin to be observed from their side (X-Z plane). Grinding using sandpapers up to P4000 and polishing using diamond suspensions up to 1 µm were applied. Further polishing was applied using 0.025 µm oxide polishing suspension followed by Ar-ion polishing (Hitachi, IM4000, Japan) for 5 min at 6 kV and a tilt angle of 8°.

Optical microscopy was done using VHX-5000 (Keyence, Japan) and DRMX (Leica, Germany) optical microscopes. SEM characterisation was conducted using an ULTRA 55 SEM (Zeiss, Germany) utilising both in-lens and external (Everhart-Thornley, ETD) detectors. Moreover, an EBSD detector (EDAX, Hikari, Japan) coupled with an EDS detector (EDAX, Hikari, Japan) was used to simultaneously obtain EBSD and EDS maps. Another EDS detector (Oxford, EXTREME, United Kingdom) was utilised to obtain EDS analyses with better precision. An accelerating voltage of 20 kV and step sizes between 0.1 and 0.2 µm were used for EBSD-EDS coupled mapping, whereas 10 kV accelerating voltage was used for SEM-EDS characterisation using the second (Oxford) EDS detector. EBSD data were cleaned by grain dilation and neighbour phase correlation using TSL OIM Analysis 8.1.0, where less than 10% of the points were altered per ASTM E2627 [66].

# 3. Results and discussion

## 3.1. Calculated phase formation

In order to understand different types of cracking in the 316L-CuCrZr system, a foundation is required regarding

the expected phase evolution during fusion-processing, on which the main hypotheses of the study will be built in the following sections.

To present an overall view of the expected phase formation under equilibrium, Figure 2 shows the calculated 316L-CuCrZr metastable equilibrium isopleth, which will be referred to as 'phase diagram' henceforth. According to the phase diagram, below 1650°C and for a large composition range of 20-90 wt.% CuCrZr, the liquid phase separates into Cu-rich and Fe-rich liquids (i.e. L(Cu) and L(Fe)). In the Fe-Cu binary phase diagram [67], LLPS is not observed, but fast-cooled Fe-Cu binary alloys exhibit LLPS through spinodal decomposition due to the metastable miscibility gap [68]. However, the addition of elements such as Cr and Mo, which have high enthalpy of mixing with Cu in liquid state [69], drives the system to exhibit LLPS under equilibrium in addition to fast cooling, together with a monotectic reaction at lower temperatures (~1360°C) as seen in Figure 2. Therefore, for a large portion of the phase diagram, solidification starts with two distinct liquid phases of different compositions, which are subsequently retained as separate phases at lower temperature.

To better represent the behaviour of the 316L-CuCrZr system under fast cooling conditions of LPBF, Scheil-Gulliver calculations were performed, which have been successfully utilised to predict solidification paths in AM of multi-materials [70,71]. Based on these, Figure 3(a) depicts the solidification behaviour expected for CuCrZr-316L mixtures with 10 wt.% increments in CuCrZr composition. Notably, all compositions shown in Figure 3(a) exhibit a similar solidification sequence, where a stepwise solidification behaviour is observed. In such a sequence, the initial solid (FCC or BCC) forms at high temperature (~1380°C) and solidification progresses very little until reaching a much lower temperature (~1085°C). Figure 3(b) shows the compositional details of such a phase formation sequence for 50 wt.% CuCrZr-316L mixture as a reference, whereas the elemental composition of each phase during solidification of the same mixture is depicted in Figure 3(c). As seen in Figure 3(b), solidification starts with the formation of FCC(Fe) in presence of L(Fe) and L(Cu) at 1375°C. FCC(Fe) formation nearly completes at 1350°C where the total solid fraction reaches 0.5. Afterwards, there is a small amount of additional FCC(Fe) forming due to the rejected elements from L(Cu), Figure 3(b), such as Fe, Cr and Mo, which have lower solubility in L(Cu) at lower temperatures, see Figure 3(c). Therefore, throughout a large temperature range (1380-1085°C), FCC(Fe) and L(Cu) co-exist without any significant changes.

As soon as the liquidus temperature of L(Cu) is reached (~1080°C), the FCC(Cu) phase starts to form.
<figure><img src="image_17.png" /></figure>

<figure><img src="image_18.png" /><figcaption>**Figure 2.** Metastable equilibrium isopleth of 316L-CuCrZr system.</figcaption></figure>

However, FCC(Cu) forms as almost pure Cu, with the other elements still remaining within the L(Cu). Therefore, solidification completes when L(Cu) transforms into FCC(Cu) and an Fe-Cr-rich BCC phase that accommodates the remaining elements. For all compositions presented in Figure 3(a), a similar solidification sequence is predicted by Scheil-Gulliver calculations, except that the first-forming solid can be BCC(Cr-Fe) instead of FCC(Fe) as the percentage of CuCrZr increases.

Considering the expected behaviour presented in the phase diagram and Scheil-Gulliver calculations, the most important insights regarding phase formation and evolution can be summarised as follows:

* Two liquids with respectively Fe-rich and Cu-rich compositions and significantly different solidification ranges are expected to form under fast cooling for all calculated compositions.

* A stepwise solidification sequence is expected for the vast majority of compositions, where L(Fe) and L(Cu) with distinct compositions and solidification temperature ranges solidify into their respective solids, i.e. FCC(Fe) and FCC(Cu). Here, FCC(Fe) forms first and exists for a large temperature range with L(Cu) without a significant change.

## 3.2. Melt pool centreline cracks

The first type of cracking observed in the experiments is distinguished by the regularly spaced lines at the 316L

side of the interface as shown in Figure 4(a), where the interface of the 316L-CuCrZr sample from the MS beam-time experiments is shown. The observed cracks are at the centre of the melt pools and separated by approximately one hatch distance (~70 µm) as visible in Figure 4(b). Upon closer examination, several features can be observed. First, grey droplets are observed within the Cu-colored regions, whereas in the grey regions, Cu-colored droplets are visible. These indicate the existence of liquid-liquid phase separation (LLPS), where the separated droplets assume a spherical shape to minimise surface energy as observed in other examples exhibiting LLPS [68,72,73]. This result corroborates the calculated phase behaviour in the phase diagram, see Figure 2. Second, the cracks at the centre of the melt pools contain Cu along the melt pool lines. Although some lines visible in Figure 4(b) have not yet separated into cracks, examples of branching cracks out of such lines can be readily observed in Figure 4(a).

Figure 5 shows the combined SEM-EBSD-EDS analyses to relate the elemental distribution and grain formation to the observed centreline cracks in the MS beam-time sample. As seen in Figure 5(a), the cracks follow the Cu-rich regions. The EBSD image in Figure 5(b) shows that the cracks took place at the high-angle grain boundaries (HAGBs), especially in regions where the grains oriented towards the centre of the melt pools meet, which are visible in the EDS maps in Figure 5(c,d). In addition, some micro-cracks branch out of these channels, following along the HAGBs. Moreover, it is observed in Figure 5(c,d) that there are Cu-rich
<figure><img src="image_19.png" /></figure>

<figure><img src="image_20.png" /><figcaption>**Figure 3.** (a) Scheil-Gulliver solidification calculations for CuCrZr-316L mixtures of different mixture percentages, and calculations showing (b) the phase formation sequence and (c) composition of the phases for the 50 wt.% CuCrZr-316L mixture. For each graph in (c), the minor elements (< 10 wt.%) are shown on the right axis with dashed lines.</figcaption></figure>

droplets within an Fe-rich matrix, confirming the presence of LLPS. EDS maps from Supplementary Figure S1 and corresponding EDS point analyses in Supplementary Table S1 further support this claim, where it is shown that the alloying elements of 316L are concentrated in the Fe-rich regions and Cu-rich regions consist of more than 90 wt.% Cu. Among the minor alloying elements from the two alloys, Cr and Mo retain in the Fe-rich regions due to their low solubility in Cu. While Ni exhibits a similar behaviour and mostly retains in the Fe-rich

regions, some Ni remains within the Cu-rich regions after LLPS as also predicted by the Scheil-Gulliver calculations in Figure 3(c).

Operando X-ray diffraction results given in Figure 6 reveal the solidification path of the alloy mixture at the interface during printing. Figure 6(a) shows the evolution of XRD spectra with time at the approximate probing position shown in Figure 4(a) while that layer was being printed. Upon melting and mixing, the diffraction lines disappear as the mixture becomes liquid. They
<figure><img src="image_21.png" /></figure>

<figure><img src="image_22.png" /><figcaption>**Figure 4.** (a) Low- and (b) high-magnification side-view optical microscope images of the sample from the MS diffraction experiments showing the whole interface with the approximate position of the X-ray, and microstructural features at the interface, respectively. Dashed line indicates a melt pool and blue arrows indicate the build direction.</figcaption></figure>

form two distinct FCC phases upon cooling as two side-by-side lines, see Figure 6(a). This is also indicative of prior LLPS in the liquid. Upon closer examination of the melting timeframes, FCC(Fe) starts forming at $t_{Fe} = 325$ ms whereas FCC(Cu) starts forming at $t_{Cu} = 352$ ms, totalling to a difference of 27 ms between the initial

formation of the two. Figure 6(c) depicts the temperature estimations obtained from FCC(Fe) and FCC(Cu) peaks respectively. It can be observed that FCC(Fe) starts to form at approximately 1380°C, while FCC(Cu) both disappears (melts) and re-solidifies at temperatures close to 1100°C.

<figure><img src="image_23.png" /><figcaption>**Figure 5.** (a) SEM and (b) EBSD inverse pole figure (IPF) + image quality (IQ) map with corresponding (c) Cu and (d) Fe EDS maps obtained from the sample from MS diffraction experiments. Blue arrow indicates the build direction. High-angle grain boundaries (misorientation angle > 15°) are highlighted by black outlines in b.</figcaption></figure>
<figure><img src="image_24.png" /></figure>

<figure><img src="image_25.png" /><figcaption>**Figure 6.** (a) Time-resolved XRD plot obtained at the X-ray position shown in Figure 4(a), indicating the initial heating up, melting and cooling stages together with (b) the close-up view showing the appearances of FCC(Fe) and FCC(Cu) peaks together with their formation sequence marked on the time axis. (c) Shows the estimated temperature evolution for the two phases. The colour bar is common for (a) and (b).</figcaption></figure>
<figure><img src="image_26.png" /></figure>

The temperature estimations may deviate from reality due to alterations in chemistry, as well as phase transformation and differential thermal contraction strains between the two phases. While the compositional changes are important at all temperatures, all other sources of strain become important at relatively lower temperatures, and thermal strains dominate at high temperature [61]. As can be seen in Figure 3(c), the compositions of FCC(Fe) and FCC(Cu) do not tend to change significantly after their formation. Therefore, since the thermal expansion functions have already been determined based on the expected compositions of each phase, the temperature estimations can be considered accurate at temperatures where solidification starts. Consistently, the cooling curves calculated independently for FCC(Fe) and FCC(Cu) match closely. Finally, the observations from Figure 6(c) agree with the expected phase formation sequence and phase formation temperatures predicted by Scheil-Gulliver calculations given in Figure 3.

The last evidence showing when and where cracking takes place is presented in Figure 7, which shows snapshots from Supplementary Video 1. First, the laser creates a deep keyhole distinguished by white contrast in the 316L part due to the laser parameters with relatively high energy density. The deep keyhole mixes the two materials via laser recoil pressure and Marangoni convection. After the laser passes, the melt pool starts

shrinking from the edges towards the centre as solidification begins. Right before complete solidification, a crack with white contrast appears at the centreline of the melt pool and remains visible.

The key findings regarding melt pool centreline cracking can be summarised to understand the mechanism of cracking. First, Figure 7 directly shows that cracking can occur at the end of solidification at the melt pool centrelines. Second, Figure 5 shows that the cracks are related to the Cu-filled channels at the centrelines, where directionally grown dendrites meet. Finally, in line with the Scheil-Gulliver predictions, Figure 6 shows that FCC(Cu) solidifies significantly later than FCC(Fe). These in situ and ex situ observations indicate that the observed centreline cracks indicate solidification cracking.

To better represent this mechanism of cracking, a schematic is presented in Figure 8. Upon laser melting, the two materials mix in the melt pool under strong forced convection to form a single liquid phase. However, the mixture may not be entirely homogeneous. Upon cooling, LLPS takes place, which creates distinct zones and droplets of L(Cu) and L(Fe) dictated by the miscibility gap. It must be noted that depending on the processing conditions (undercooling), LLPS may take place through spinodal decomposition or nucleation-growth type of formation [74,75]. In the former case, an interconnected network of the two liquids can be observed initially, whereas the latter

<figure><img src="image_27.png" /><figcaption>**Figure 7.** X-ray radiographs from the ID19 sample (H) showing the formation of a crack at the centreline of the melt pool upon melting and solidification. Dashed yellow arrow shows the keyhole induced by the laser, dashed white lines schematically show the melt pool boundary, and white arrows show the crack. Scan direction is in and out of the page/screen. The scale bar is common to all images. For a visual representation of the dynamics, refer to Supplementary Video 1.</figcaption></figure>
<figure><img src="image_28.png" /></figure>

<figure><img src="image_29.png" /><figcaption>**Figure 8.** Schematic representation of solidification cracking in steel-copper multi-materials. (1) Liquid mixing in the melt pool, (2) liquid-liquid phase separation, (3) competitive growth of Fe-rich dendrites, (4) Cu-rich liquid segregation between dendrite boundaries, (5) crack initiation at the dendrite boundaries, (6) crack growth following the Cu-infiltrated centreline.</figcaption></figure>

directly leads to the formation of droplets. As the separated liquids evolve, they tend to ripen and assume a spherical shape to minimise surface energy. Therefore, as for the case of nucleation-growth, LLPS occurred through spinodal decomposition is expected to result in droplet morphology especially under strong convection [68], as is the case for LPBF. Following LLPS, the formation of FCC(Fe) begins. The FCC(Fe) dendrites then start growing towards the centre of the melt pool due to competitive growth. Here, the grains of the previous layer at the edge of the melt pool potentially act as a substrate for the freshly forming grains since solidification initiates from the edge of the melt pool in Figure 7 and progresses towards the centre, which is consistent with the appearance of the directional grain growth towards the melt pool centres observed in Figure 5(b). Then, as predicted by Scheil-Gulliver calculations for the whole composition range, see Figure 3(a), there is limited solid formation for a large temperature range (~300°C) where FCC(Fe) and L(Cu) co-exist. As the

growth of FCC(Fe) dendrites continues towards the centreline of the melt pool, the remaining liquid goes to the solidification front and between the dendrites as usually observed in directional growth [22]. As cooling proceeds, tensile thermal stresses keep accumulating on FCC(Fe) dendrites due to non-homogeneous thermal gradients [22,76,77]. When the centreline L(Cu) starts solidifying into FCC(Cu), the phase transformation shrinkage creates additional tensile stress [22,76,77]. Therefore, at the last stage of solidification, these stresses acting on FCC(Fe) dendrites separate the mushy zone composed of L(Cu) and FCC(Cu) at the centreline into cracks. At this stage, if the Cu-rich droplets can supply further L(Cu) into these cracks, they can become back-filled as also known in welding [22], leading to the appearance of Cu-infiltrated lines observed in OM images in Figure 4(b) and the EDS map in Figure 5(c).

The solidification cracking mechanism described above takes place only under certain conditions. The first prerequisite is that the first forming solid,
<figure><img src="image_30.png" /></figure>

FCC(Fe), must form dendrites growing towards the centre of the melt pool. For instance, compositions approximately above 50 wt.% CuCrZr lead to a high fraction of L(Cu), which becomes the main matrix in which FCC(Fe) forms as separated droplets [26], an example of which is the grey droplets shown in Figure 4(b). Therefore, solidification cracking does not take place for these compositions as FCC(Fe) cannot form a connected body with the base part that can crack due to the restraint. The second prerequisite for solidification cracking is the existence of L(Cu) between the FCC(Fe) dendrites. This can only happen for compositions sufficiently rich in CuCrZr, because for very low concentrations, all Cu can be dissolved in FCC(Fe) without forming any L(Cu) as can be seen in the phase diagram in Figure 2. Hence, solidification cracking is avoided as solidification takes place within a relatively narrow freezing range comparable to 316L itself, which is not known for solidification cracking. Following this understanding, solidification cracking should always be observed for compositions between 5 and 50 wt.% CuCrZr. This matches the observations in Figure 4(a), Figure 5 and Supplementary Figure S1, where cracking was observed at the 316L side of the interface instead of the CuCrZr side.

Another important consideration is the nature of soli- dification cracking for hypomonotectic compositions (< 20 wt.% CuCrZr) in Figure 2, where no LLPS is expected, and the process instead starts with a single liquid of Fe- rich composition. One may expect no solidification cracking in this case, as there is no initial L(Cu) that can infiltrate the centrelines. However, a small amount L(Cu) is expected to form as the result of the monotectic reaction at the dendrite boundaries where L(Fe) and FCC(Fe) meet, as has been observed for directionally soli- dified monotectic alloys [78]. This introduces L(Cu) right at the solidification front and between the dendrites, making these compositions particularly susceptible to cracking due to very little amount of disconnected L(Cu) regions that can cause cracks [78].

The observations regarding the compositional dependence of solidification cracking are consistent with the literature where solidification cracking was pro- posed as the main mechanism of cracking for various steel-copper alloy combinations, presented in Table 1. Notably, Noecker et al. [11] have shown that the crack susceptible compositions lie between 5.4 and 43.3 wt.% Cu-steel during welding by putting emphasis on the freezing range and the amount of terminal Cu liquid. Rinne et al. reported hot cracking for compo- sitions above 10 wt.% Cu [13]. Zhao et al. showed that compositions above 9 wt.% Cu become susceptible to solidification cracking while cracking stops at 50 wt.%

Cu when co-processed with both H13 and 316L steel grades [29]. Similarly, Wen et al. [26] used Kou's crack susceptibility model [77] for the full range of compo- sitions for 316L-Cu10Sn mixtures. Their calculations showed that the highest crack susceptibility was between 5 and 50 wt.% Cu10Sn with the peak being at 5 wt.% Cu10Sn, which was the only range where their experiments resulted in cracking.

When steel and copper are extensively melted and mixed, a gradient of composition is established within the melt pool. Specifically for laser-based techniques, a high heat input is often required to melt Cu alloys due to their low laser absorptivity in the infrared range (the predominant wavelength range used in the literature) and high thermal conductivity. Hence, horizontal inter- faces where Cu alloys are placed on top of steels are expected to have deep melt pools and pronounced mixing. In such conditions, the gradual transition of composition within the melt pool likely favours both phase separation and solidification cracking. Therefore, solidification cracking together with phase separation is often reported in cases where extensive mixing in the melt pool is observed [11,14,28]; the same applies when a pre-mixed material feedstock with such a com- position is used [27]. As solidification cracking is favoured at Fe-rich compositions (5–50 wt.% Cu), it can also be observed when steel is fused on top of copper, even if liquid mixing is limited [10]. In addition, some research specifically reports solidification cracking due to the Cu network between the Fe-rich grains without any observation of phase separation, hinting a relation to the solidification cracking due to the formation of monotectic L(Cu) at the dendrite boundaries as explained previously [13].

In summary, the centreline cracking occurs due to the formation of two liquids of vastly different composition, and thus, freezing range. These two liquids can form either because of LLPS or the monotectic reaction. Cracking takes place at the Fe-rich dendritic grain boundaries, hence, requires the overall composition to fall within the 5–50 wt.% CuCrZr range, which is appli- cable for other copper-steel alloy combinations.

### 3.3. Non-centreline cracks

Apart from the larger, Cu-infiltrated centreline cracks, other micro-cracks and Cu-infiltrated channels were observed within the multi-material samples, see Figure 5 and Supplementary Figure S1. Moreover, it can be seen by comparing Figure 5(b–d) that these channels follow HAGBs, but they do not necessarily lie along the build direction within the centreline of the melt pools and can instead traverse melt pool boundaries. While
<figure><img src="image_31.png" /></figure>

the Cu-infiltrated centreline channels and cracks are between 100 and 200 µm in length, these are much shorter, between 10 and 50 µm. These indicate that another cracking mechanism is at play for these cracks.

Figure 9(a,b) show snapshots from two cracking events in the ID19 sample (L) as observed in Supplementary Video 2 and 3, respectively. First, in Figure 9(a) and Supplementary Video 2, the laser creates a deep keyhole as laser scans over the sample, which then forms the melt pool. While the melt pool shrinks towards its centre, a crack appears in a location approximately 500 µm away from the melt pool before the solidification is complete. This crack then grows slightly upon further cooling and subsequently remains within the material. This observation indicates cracking due to the stresses induced at the crack site during cooling. Secondly, as seen in Figure 2(b) and Supplementary Video 3, the laser passing over the track creates a keyhole. However, differently from the previous scenario, a crack initiates right next to the melt pool upon melting.

To provide further insights, Figure 10(a) shows an SEM image from the surface of the ID19 sample (L)

exhibiting longitudinal cracks at the melt pool edges and transverse cracks crossing the melt pools. Upon closer examination, see Figure 10(b), the directionally grown dendrites are visible on the surface, where cracks follow the dendritic grain boundaries. While it is not possible to distinguish the secondary dendrite arms in the bulk as they coalesce, see Figure 5(b), they are visible on the surface on the sample in Figure 10(b). The primary and secondary dendrite arm spacings of the dendrites are respectively 0.85 ± 0.15 µm and 0.70 ± 0.10 µm. These correspond to a cooling rate of approximately 2×10⁵°C/s following Belsvik et al. [79], which is consistent with the previous reports on LPBF [80]. EDS maps presented in Figure 10(c) clearly indicate that the dendrites are richer in alloying elements of 316L whereas Cu is heavily segregated towards the inter-dendritic regions. Segregation of Cu leads to a greater cracking tendency as molecular dynamics (MD) simulations have shown that the presence of Cu alters the Fe-Fe bonds to Cu-Fe and Cu-Cu bonds at the Fe grain boundaries, leading to a lower grain boundary separation (fracture) energy [81]. Moreover, the strength of these bonds

<figure><img src="image_32.png" /><figcaption>**Figure 9.** X-ray radiographs from the ID19 sample (L) showing the formation of cracks (a) far away from the melt pool and (b) directly next to the melt pool boundary. Yellow arrows show the keyhole induced by the laser, dashed white lines schematically show the melt pool boundary, and white arrows show the cracks. The scale bar is common to all images. Scan direction is in and out of the page/screen. Please refer to Supplementary Video 2 and 3 for clear understanding.</figcaption></figure>
<figure><img src="image_33.png" /></figure>

rapidly deteriorates at high temperatures, especially Cu-
Cu bonds at temperatures near the Cu melting point.
Hence, tensile thermal stresses induced on the grain
boundaries can easily lead to cracking once Cu is distrib-
uted between the dendrites. Therefore, it is crucial to
understand how Cu decorates the grain boundaries of
Fe to prevent cracking.

As is the case for many embrittling metals, Cu tends
to wet Fe grain boundaries. As reported by Savage
et al. [16], the wetting tendency increases with tempera-
ture, reaching a peak close to the melting point of Cu,
and then reduces again at even higher temperatures.
Thus, Cu close to its melting point readily infiltrates
steel grain boundaries. This is observed in Figure 5,
where branches of Cu filled channels stem from Cu-
rich droplets and channels. However, it is unlikely that
this homogeneous distribution of Cu is attained
between FCC(Fe) dendrites only by wetting from Cu-
rich regions. Considering that FCC(Cu) starts forming at

a temperature approximately 300°C below the formation
temperature of FCC(Fe), L(Cu) can easily infiltrate the
inter-dendritic regions of FCC(Fe) dendrites during
their growth. Furthermore, further L(Cu) can form at
the dendritic grain boundaries because of the monotec-
tic reaction L(Fe) → FCC(Fe) + L(Cu) where L(Fe) and
FCC(Fe) dendrites meet, as was observed for other
monotectic alloys with the second liquid wetting the
forming solid [78]. As the solubility of Cu in both L(Fe)
and FCC(Fe) does not change significantly with tempera-
ture, see Figure 3(c), L(Cu) between the dendrites
remains as-is until solidifying into FCC(Cu) later on,
making these the probable causes of the homogeneous
distribution of L(Cu) observed in Figure 10(c).

The cracking due to the strength loss at the grain
boundaries of steel in presence of Cu is commonly
referred to as copper contamination cracking (CCC) or
liquid metal embrittlement (LME). In conventional pro-
cessing, Cu penetrates the grain boundaries of steel

<figure><img src="image_34.png" /><figcaption>**Figure 10.** (a) Low- and (b) high-magnification SEM images of the surface of ID19 sample (L) showing the longitudinal and transverse cracks at the surface together with (c) the corresponding EDS maps showing the elemental distribution at the dendrites and dendrite boundaries. Blue arrow indicates the build direction.</figcaption></figure>
<figure><img src="image_35.png" /></figure>

from a Cu source such as a coating or brazing filler that melts during processing. However, the formation cracks is not dictated by the physical state of Cu. Instead, the physical state determines the mobility of Cu atoms and how easily they migrate to steel grain boundaries. This ultimately dictates how fast cracking tendency rises. Therefore, cracking can be observed earlier when the embrittling element is a liquid compared to solid. This is because mass transport takes place via bulk liquid flow (fast) for the liquid, whereas the mass transport of the solid depends on diffusion (slow) [82]. However, solid metal embrittlement (SME) can still be observed for many elements that are known liquid embritlers in steel such as Zn, Pb, Cd, Sn and In [83] as well as for Cu itself [84,85].

To summarise, the main evidence regarding the observed micro-cracking are as follows:

* Cracking takes place at the dendritic grain boundaries where extensive Cu segregation is observed, see Figure 10.

* Unlike solidification cracks that take place at the centres, these cracks take place either far away from the melt pools during printing, or right next to the melt pools in the HAZ, see Figure 9, Supplementary Video 2 and 3.

Based on this information and considering the weakening effect of Cu on Fe grain boundaries [81], the mechanism behind the cracks observed far away from the melt pool can be referred to as metal-induced embrittlement (MIE), see Supplementary Video 2 and Figure 9(a). This naming is used as the umbrella term for both SME and LME for steel-copper multi-materials,

although mass transport of Cu mostly takes place in liquid state in this specific case. Considering the cracking event in Figure 9(a), the fact that cracking takes place in solid state and away from the melt pool points at the loss of GB strength due to MIE, and subsequent cracking takes place when thermal stresses are induced on the GBs during printing in addition to ever-present residual stresses.

On the other hand, the cracking observed in Supplementary Video 3 and Figure 9(b) resembles a typical liquation crack, where the low-melting-point species at the grain boundaries, Cu in this case, melt at the HAZ of the adjacent melt pool and separate into cracks due to the thermal stresses induced during printing. Different than MIE, liquation cracking takes place when Cu at the GBs becomes liquid with virtually zero strength due to reheating. Therefore, although both MIE and liquation cracking stem from the Cu segregation between the GBs, liquation cracking necessitates local temperature to be above the melting point of Cu, unlike MIE.

The above-mentioned mechanism is schematically shown in Figure 11 with both types of cracking. In both cases, micro-segregation of Cu at the dendritic grain boundaries constitutes the root cause of weakening, and whether MIE or liquation causes cracking is essentially stochastic.

To assess the prevalence of distinct types of cracking, X-ray CT measurements have been conducted. The 3D view of the X-ray CT sample shown in Figure 12(a) demonstrates the complex crack network within the sample. As visible in the side cross-section in Figure 12(b), large cracks lie in Z-direction and are positioned just below the Cu-rich melt pools, which can be

<figure><img src="image_36.png" /><figcaption>**Figure 11.** Schematic representation of liquation and metal-induced embrittlement cracking in steel-copper multi-materials produced via fusion-based methods. (1) dendritic solidification and Cu infiltration and segregation into grain boundaries, (2) cracking at the grain boundaries due to MIE and liquation at the HAZ.</figcaption></figure>
<figure><img src="image_37.png" /></figure>

<figure><img src="image_38.png" /><figcaption>**Figure 12.** (a) X-ray CT images of a 316L-CuCrZr sample showing the segmented crack network (in red) with corresponding (b) side and (c) top cross-section views. (d) Shows the crack length distribution where the blue dashed line shows the median value.</figcaption></figure>

distinguished by their white contrast, exhibiting a similar appearance to Figure 5(a). On the other hand, the top cross-section view presented in Figure 12(c) shows that most cracks are associated to the part surface and penetrate the thickness (Y-direction) while showing smaller side branches along the length (X-direction). For better understanding, the frequency of crack lengths is shown in Figure 12(d). As seen in this distribution, majority (~75%) of cracks are shorter than 75 µm, the median value being 51 µm. These smaller cracks are different both in their length and location compared to the long cracks (100–200 µm) observed in Figure 12(b) as well as the centreline cracks shown in Figure 4. Instead, most cracks in Figure 12 appear like the types observed in Figure 10 by their length, location, and orientation. Thus, these should stem from MIE or liquation cracking mechanisms. This also explains why these cracks are more frequent. While solidification

cracking only takes place in the terminal stage of solidification at the melt pool centrelines, MIE/liquation cracking at the grain boundaries can occur at various locations for a wide combination of temperatures and stress states once Cu is distributed between the GBs. Since there are more dendritic grain boundaries than there are melt pools, the probability of observing micro-cracks formed due to MIE or liquation is naturally higher than observing large cracks formed due to solidification cracking. Therefore, it is possible to distinguish solidification and MIE/liquation cracks based on their length, location, and alignment with respect to the melt pools. While *operando* X-ray radiography clearly shows their difference, no *ex situ* observations could be made that can distinguish MIE and liquation cracks. Here, it must also be noted that although distinguishing the type of cracking is possible in controlled experiments that offer valuable insights, it becomes rather
<figure><img src="image_39.png" /></figure>

difficult to understand the type of cracking in the case of printing complex multi-material components. In such cases where all three cracking mechanisms may be active throughout the build, complex and interconnected crack networks with various lengths, locations and alignments can be expected. Therefore, the descriptions based on the presented criteria are to be taken as guidelines rather than absolute rules that can be applied to all possible cases.

Consistent with the presented findings, studies utilising similar alloy combinations reported both solidification and MIE/liquation cracking under regular LPBF conditions as the composition gradient at the interface can favour both mechanisms [25, 26, 30, 32, 34]. Depending on the nature of the fusion-process, phase separation can be observed with Cu-infiltrated cracks (MIE/liquation) stemming from Cu-rich regions, rather than solidification cracks if the composition in the melt pool does not favour the latter [9, 12, 18, 21], resembling an appearance depicted in Figure 11. These examples all exhibit Cu segregation at Fe-rich grain boundaries, which likely stems from LLPS or the monotectic reaction upon extensive melting, liquid mixing and solidification.

On the other hand, some studies report cracking exclusively related to Cu infiltration between the grain boundaries in absence of phase separation, where the nature of the process does not necessitate liquid mixing. For instance, cracking is observed at the HAZ during GTAW of Cu-plated steels, where cracks form outside of the fusion zone where the base metal does not melt [16]. Furnace melting of Cu and Cu-based brazing alloys on various steel types leads to cracking at the grain boundaries due to Cu infiltration [17, 20]. Similarly, cracking in presence of liquid Cu is observed during high temperature stress-rupture tests of Cu-plated steels, especially under tensile stress [19]. Moreover, micro-crack formation at the steel grain boundaries was reported for diffusion bonding of copper and steel, even though the system was completely in solid state [84, 85]. What should be noted in these examples is that the timescales for the observation of cracking are much longer ($10^0-10^3$ s instead of $10^{-2}-10^0$ s) compared to the case of liquid mixing and Cu redistribution into the grain boundaries, as Cu migration from a local source takes significantly longer. This difference can be taken advantage of for mitigating cracks and will be discussed in the following section.

### 3.4. Crack mitigation by controlled melting

As presented in previous sections, both solidification and MIE/liquation cracking are closely related to the phase formation during melting and solidification.

Specifically, LLPS and the monotectic reaction directly cause solidification cracking at the melt pool centrelines while amplifying MIE/liquation cracking at the grain boundaries. Therefore, avoiding the formation of the two liquids constitutes the first step towards mitigating cracks in copper-steel multi-materials. The prerequisite for not having separate liquids of L(Cu) and L(Fe) is not to melt CuCrZr and 316L above the monotectic reaction isotherm (~1360°C). Therefore, it was aimed in P21.2 diffraction experiments to melt only CuCrZr powder on top of the 316L part without melting the latter by staying in the temperature range 1100-1360°C. This was ensured by probing the interface and monitoring the diffraction peaks of FCC(Fe) and FCC(Cu), from which the temperature profiles were estimated.

Figure 13(a) shows the heating and cooling stages of the sample with the side-by-side peaks of FCC(Fe) and FCC(Cu). The heating stage takes approximately 2.5 s and is composed of short, repeated scans as can be observed from the serrations in the diffraction lines in Figure 13(b). The peak positions of both phases shift to smaller scattering vectors (Q) in the repeated scanning stages during heating, where FCC(Cu) peaks occasionally disappear due to melting while the peak related to FCC(Fe) remains visible. The simultaneous existence of the diffraction lines from both FCC(Fe) and FCC(Cu) confirms that the X-ray beam was correctly positioned at the multi-material interface. As seen in the estimated temperature profile of FCC(Fe) shown in Figure 13(c), the steel part remains below 900°C. On the other hand, Figure 13(d) shows that the peak temperature of FCC(Cu) reaches up to ~1050°C, which is close to the melting point of CuCrZr (1065°C). As CuCrZr was in powder form which moves around during printing, the diffraction lines of FCC(Cu) involve contributions from different powder particles as well as previously printed CuCrZr regions in different locations and melting-cooling stages. Therefore, there are occasional abnormal spikes in temperature. Still, it can be reasonably claimed that CuCrZr powders have at least partially melted during the experiment, whereas the 316L part remained below its melting point (~1400°C).

It is important to mention the source of errors in the temperature estimations to support the hypothesis above. The first assumption was that the observed peak shifts only stem from thermal strains, whereas elastic strains were also present. This leads to an error since elastic strains appear as temperature changes. However, this error is negligible at high temperatures, and only becomes important at moderate to low temperatures [61]. Moreover, the 316L part initially contains some elastic strain. Therefore, assuming the initial FCC(Fe) peaks as the strain-free reference is another
<figure><img src="image_40.png" /></figure>

<figure><img src="image_41.png" /><figcaption>**Figure 13.** (a) Time-resolved XRD plot obtained at the 316L-CuCrZr interface from the P21.2 experiment indicating the heating and cooling stages together with (b) the close-up view showing FCC Fe and FCC Cu peaks. (c) and (d) show the estimated temperature evolution for FCC Fe and FCC Cu, respectively. The colour bar is common for (a) and (b).</figcaption></figure>

source of error. Still, considering that the initial CuCrZr powder can be assumed strain-free, the difference between the estimated temperatures of FCC(Cu) and FCC(Fe) during cooling does not exceed 50°C, see Figure 13(c,d). Based on these assessments, potential errors in the temperature estimations do not exclude the hypothesis that the 316L part never melted.

The optical micrograph of the P21.2 sample given in Supplementary Figure S2a shows a dense part except the cracks visible at the top region together with some Cu-rich regions distinguished by the colour contrast. Upon closer examination the cracks are closely related to the Cu-rich regions, see Supplementary Figure S2b. There are no Cu- or Fe-rich droplets indicating phase separation. This observation supports the findings presented in Figure 13 that the monotectic temperature was not exceeded, which would have induced phase separation. Figure 14(a) shows an SEM image of one of

the crack networks at the top layer. As seen in the EBSD IPF map in Figure 14(b), the crack shown in Figure 14(a) exclusively follows HAGBs. As shown in Figure 14(c,d), there is a significant Cu segregation at the grain boundaries where cracking occurred, whereas the grains themselves are Fe-rich as expected. Since the 316L part was never melted together with CuCrZr, there is no phase separation or extensive diffusion observed between the two regions.

Comparison between the OM images of the control and method validation samples from the *ex situ* controlled melting experiment given in Figure 15 corroborate the findings presented above. As seen in the OM image of the control specimen given in Figure 15(a), there are many large and small cracks visible. Following the observations presented previously, it can be seen that the larger cracks at the centre of the melt pools given in Figure 15(b) point at solidification cracking
<figure><img src="image_42.png" /></figure>

<figure><img src="image_43.png" /><figcaption>**Figure 14.** (a) SEM image and (b) EBSD inverse pole figure (IPF) + ETD detector signal map with corresponding (c) Cu and (d) Fe EDS maps obtained from the sample from P21.2 diffraction experiment. Blue arrow indicates the build direction.</figcaption></figure>

mechanism, whereas smaller cracks right below the interface indicate MIE/liquation cracking, see Figure 15(c). Moreover, the droplets of Fe-rich regions are visible in Figure 15(c), confirming that phase separation

took place. On the other hand, a much cleaner interface is observed for the method validation sample, see Figure 15(d), without any evidence of significant phase separation. Although the interface appears crack-free, some

<figure><img src="image_44.png" /><figcaption>**Figure 15.** (a) Side-view optical microscopy image of the reference sample of the controlled melting build with close-up views showing (b) solidification and (c) metal-induced embrittlement/liquation cracks, respectively. (d) Side-view optical microscopy image of the method validation sample with (e) a high-magnification image showing Cu infiltration lines. Black arrows indicate the cracks, and white arrows indicate the build direction.</figcaption></figure>
<figure><img src="image_45.png" /></figure>

lines of Cu infiltration are still visible at higher magnification, as shown in Figure 15(e). As the method validation sample was not built under temperature monitoring, it is not possible to ascertain that 316L part was not melted during laser scanning. Still, this appears as a valid assumption as neither melt pools nor clear phase separation are visible in Figure 15(d). It must be noted that several parameters have been experimented with for the method validation build by changing the laser power, layer thickness and repetition count. While insufficient bonding was observed due to insufficient melting of CuCrZr for most cases, excessive cracking was observed for parameters with higher repetition counts as they allowed for grain boundary penetration. Only a small set of experiments yielded continuous interfaces with minor cracking. As a complete optimisation requires a broader research effort, the full set of experiments is not provided in this study. However, it was concluded that optimal processing can be achieved by letting CuCrZr melt and solidify as fast as possible while ensuring steel side of the interface remains solid.

Based on these, it can be claimed that avoiding extensive mixing and subsequent phase separation greatly reduces cracking. In addition, it can be observed by comparing Figure 6(c) and Figure 13(d) that the average cooling rate was reduced when gradual melting was applied instead of regular printing. This might have reduced cracking due to a lower residual stress build up, as smaller thermal gradients are generated under slower cooling [86,87]. Residual stresses are particularly important when MIE is considered. Solidification and liquation cracking primarily occur due to the thermal stresses induced because of rapid heating and cooling. In addition, the former requires complete melting of the multi-material system whereas the latter requires reheating up to the melting point of Cu. On the other hand, as long as Cu infiltration into the grain boundaries takes place, MIE can occur at any region within the sample due to the buildup of residual stresses, in addition to thermal stresses. This distinction makes MIE relatively frequent. Accordingly, although Supplementary Video 2 shows only a single instance of MIE, several instances of it were observed during ID19 experiments.

The presented method in the controlled melting build comprises the first step towards obtaining sound interfaces in steel-copper multi-material parts. As cracking only takes place in vicinity of the steel part, one only needs to ensure a sufficiently thick Cu layer by slow melting. Afterwards, the rest of the Cu-based portion can be printed using regular printing parameters. Therefore, the presented method can be utilised to obtain a

sound multi-material interface without a substantial compromise in build rate. One drawback of this approach is that it relies on Cu alloy being on top in a horizontal interface. Moreover, although a continuous interface was established and cracking was significantly mitigated via this approach, the actual bonding strength is still to be verified. Therefore, the applicability of gradual melting of the Cu alloy at various vertical and horizontal interfaces requires further investigation.

### 3.5. Dominant factors affecting cracking and mitigation strategies

As has been previously detailed, cracking in steel-copper multi-materials can be categorised by two main mechanisms: (i) solidification cracking due to LLPS and the monotectic reaction leading to a wide solidification range; (ii) MIE/liquation cracking due to Cu infiltration into the steel grain boundaries which may or may not be reinforced by the presence of liquid mixing or phase separation. Overall, some degree of cracking appears unavoidable in most cases. Still, there have been reports that presented successful results, which will be discussed in detail below.

The first key consideration for achieving sound joints between steel and copper is to minimise extensive mixing and subsequent phase separation, which is the dominant cause of solidification cracking, and amplifies MIE/liquation cracking. As a rule of thumb, the processing parameters must be adjusted such that a continuous interface with good metallurgical bonding is established, while the materials melt and solidify as fast as possible to avoid extensive liquid mixing and grain boundary penetration. Accordingly, it can be expected that higher heat inputs lead to more cracking because: (i) extensive mixing in a large melt pool can easily establish a local composition prone to solidification cracking, (ii) the liquid remains alive for a longer period that allows for more grain boundary infiltration, (iii) a larger HAZ forms that can lead to more liquation cracking. Literature reports exploring the link between cracking and processing parameters across different fusion-processes corroborate these. For instance, severe cracking was observed in LPBF of 316L/CuSn10 multi-materials when lower scan speeds and higher laser powers were used [88]. Similar observations were reported in LPBF of copper and maraging steel, where a decrease in scan speed below the optimal value for sufficient interface bonding led to more cracking at the interface [89]. Severe cracking was observed in steel-copper dissimilar welds when laser power is increased and welding speed is decreased, which led to extensive mixing between the two materials [13].
<figure><img src="image_46.png" /></figure>

Moreover, the same study reported that the degree of mixing at the interface can be reduced by using an over-focused laser beam [13], similar to the approach presented in the current study. While avoiding LLPS and the monotectic reaction certainly eliminates solidification cracking, it can be avoided even in presence of phase separation by ensuring a sufficiently rich Cu-rich composition (>50 wt.%) in the melt pool in a vertical interface welding configuration, [24]. In this case, whether MIE/liquation cracking will occur depends on how long liquid Cu stays in contact with the steel part. Therefore, this approach is exclusive to welding as the local composition in the melt pool is set during a single melting event. Moreover, the reproducibility of this route depends heavily on the interface configuration, as well as the strict control of the processing conditions. In addition to the common process parameter adjustments, cracking can be mitigated, for instance, by controlling the laser beam shape, although MIE/liquation cracking remains difficult to eliminate completely [31]. The degree of cracking via MIE/liquation mainly depends on how fast Cu atoms migrate to the grain boundaries. The faster Cu melts and solidifies; the fewer cracks can be expected. This is evident when comparing Supplementary Figure S2 and Figure 15(d), where melting CuCrZr over more repetitions increased the time where Cu-rich liquid is in contact with the steel part, and thus, the degree of cracking.

Beyond adjustments of the process parameters, other routes to obtain sound copper-steel interfaces run through the adjustment of the phase constitution or chemical composition. For instance, early reports have shown that steels with a ferritic microstructure at the melting point of Cu have higher resistance against copper infiltration compared to the ones with an austenitic microstructure [20]. This mainly stems from the difference between the grain boundary energies of austenite (γ) and ferrite (α), as well as their respective interface energies with liquid Cu. Complete wetting is expected at the grain boundaries when the solid-liquid interface energy is smaller than half of the GB energy ($2\sigma_{SL} < \sigma_{GB}$), which is favoured at higher temperatures [90]. At 1105°C, y/γ GB energy and y/liquid Cu interface energy are respectively given as 0.85 and 0.43 J/m² [91], which shows a high wetting tendency. On the other hand, α/α GB energy at 1150°C is approximately 10–15% lower compared to y/γ GB [91]. Assuming similar values for the interface energies of y/liquid Cu and α/liquid Cu, it can be expected that the tendency of liquid Cu to penetrate α/α GBs is lower compared to the y/γ GBs, since $\sigma_{\alpha} < \sigma_{\gamma}$. Thus, cracking tendency is expected to be lower for α-Fe compared to γ-Fe. Several studies have leveraged this behaviour in arc welding by creating a ferrite band

at the interface between liquid Cu and the steel piece, which was achieved by the addition of Si as the ferrite stabiliser [15,33]. Another important element common to the research that report successful joints is Al. Early studies have shown that the addition of ~12 wt.% Al to Cu can eliminate LME in steels [19]. Conformingly, many recent research report successful welding of Cu-Al bronzes with various steel types [35–39]. A recent study utilising MD simulations has shown that the beneficial effect of Al stems from its high segregation potential into the steel grain boundaries, which retards the infiltration of Cu by altering the electronic structure of the grain boundaries [81]. As shown in the study, Al slows down the penetration of Cu rather than preventing it completely. Moreover, grain boundary fracture energy is relatively higher when Al is co-doped with Cu into the grain boundaries, compared to the Cu-only case. However, although to a lower extent compared to Cu, Al still lowers the grain boundary cohesion [81]. Accordingly, it can be observed that if time spent in liquid state at the interface is not strictly controlled, cracking can still occur even if a Cu-Al bronze is used with steel. [15,39].

In summary, prevention of cracking in copper-steel joints presents a multi-faceted problem. This issue must be solved using a combined approach of adjusting the processing conditions as well as the chemical composition of the materials. As a first course of action, solidification cracking can be avoided and MIE/liquation cracking can be minimised by ensuring reduced mixing at the interface via process control. However, MIE/liquation cracking may persist in many cases due to grain boundary wetting. In this case, extending the chemical composition space beyond the common remedies such as adding Al or Si is required. For this aim, the literature on similar embrittling elements can be considered. For instance, as shown in a study on Zn-induced LME in steels, embrittling elements such as Cu and Zn have similar segregation potentials and weakening effect on α-Fe grain boundaries [92]. Moreover, as in the case of Cu, Al and Si may have a slight positive effect against embrittlement when co-doped with Zn into steel grain boundaries [92,93]. B also appears to have a promising effect due to its repulsive interaction with Zn [92], which helps boron steels to reduce their susceptibility to LME [94,95]. On these grounds, borrowing remedies against other embrittling elements appears to be a viable route for future research and ultimate solution to cracking in copper-steel combinations.

## 4. Conclusions

Steel and copper alloys constitute an attractive set of multi-materials for applications that require a
<figure><img src="image_47.png" /></figure>

combination of strength, corrosion resistance and thermal or electrical conductivity. However, this combination can suffer from cracking when processed via fusion-based techniques. To understand the underlying causes of cracking in steel-copper multi-materials and assess their relative importance, *operando* X-ray diffraction and imaging experiments have been conducted during LPBF of 316L and CuCrZr. These experiments were then combined with auxiliary characterisation methods and thermodynamic calculations. The results provided real-time insights into the cracking mechanisms. Based on these, the main findings of this work are as follows:

* Upon melting CuCrZr and 316L in the melt pool, liquid-liquid phase separation (LLPS) takes place due to the miscibility gap at high temperatures. This stems from the large positive enthalpy of mixing between Cu and Fe, in addition to Cr and Mo.

* Cu- and Fe-rich liquids form upon LLPS over a wide range of compositions. These two liquids solidify into their respective solids of Cu- and Fe-rich compositions with vastly different solidification ranges.

* The extension of the solidification range due to having two liquids of different composition leads to solidification cracking. Solidification cracking occurs at the Fe-rich side of the interface in the terminal stage of solidification, where cracking initiates from the mushy zone within the Cu-rich regions at the centrelines of melt pools.

* Severe micro-cracking was observed at the Fe-rich dendritic grain boundaries, together with extensive Cu segregation. These micro-cracks occur via two routes: (i) metal-induced embrittlement (MIE) in the solid state far away from the fusion zone, (ii) liquation cracking in the heat-affected zone around the melt pool.

* The main reason for Cu segregation was identified as the redistribution of Cu-rich liquid between the Fe-rich grain boundaries during solidification, which occurs due to LLPS as well as the monotectic reaction. Cu at the GBs deteriorates GB strength and acts as a low-melting-point phase, respectively leading to MIE and liquation cracking.

* Micro-cracks were shorter compared to solidification cracks, but more frequent. While solidification cracks at the melt pool centrelines lied along the build direction, micro-cracks exhibited a random orientation, primarily following high-angle grain boundaries.

* Overall, two main mechanisms of cracking have been identified, that cause three different types of cracking: (i) solidification cracking due to LLPS and the monotectic reaction leading to an extended freezing range, (ii) redistribution of Cu between the Fe-rich dendrites during solidification subsequently leading to either MIE or liquation cracking during processing.

* An experiment was devised to avoid cracking where CuCrZr powder was gradually melted on top of a 316L part by ensuring that the latter did not melt. It was shown that by avoiding LLPS and the monotectic reaction, cracking can be greatly mitigated.

* It was demonstrated that most cracking events reported in the literature can be explained within the scope of the two main mechanisms proposed in this study. It was also shown that the crack-free examples either corroborate the main findings of this study or involve significant alterations to the composition of the alloys which reduce the cracking tendency.

In conclusion, two main mechanisms have been presented for cracking in steel-copper multi-materials, which mainly stem from the metallurgical behaviour of the two alloy types during fusion-based processing. As the issues related to such combinations are quite similar irrespective of the fusion process, the findings presented in this study are expected to lay a foundation for avoiding cracking in fusion-processing of steel and copper alloys in the future.

## Acknowledgments

We acknowledge the European Synchrotron Radiation Facility (ESRF) for provision of synchrotron radiation facilities under proposal number HC-5703 and we would like to thank Jean Paul Valade and Diego Hopson Safatli for assistance and support in using beamline ID19. We acknowledge DESY (Hamburg, Germany), a member of the Helmholtz Association HGF, for the provision of experimental facilities. Parts of this research were carried out at PETRA III and we would like to thank Sven Gutschmidt for assistance in using P21.2. Beamtime was allocated for proposal I-20240114 EC.

## Author contributions

CRediT: **Andaç Özsoy**: Conceptualization, Data curation, Formal analysis, Investigation, Methodology, Validation, Visualization, Writing – original draft; **William A. Hearn**: Conceptualization, Investigation, Methodology, Writing – review & editing; **Steve Gaudez**: Investigation, Software, Writing – review & editing; **Rijuta Jeswani**: Data curation, Investigation; **Yunhui Chen**: Data curation, Investigation, Resources; **Alexander Rack**: Data curation, Investigation, Resources; **Zoltan Hegedüs**: Data curation, Formal analysis, Resources, Software; **Nicola Casati**: Data curation, Investigation, Resources; **Roland E. Logé**: Supervision, Writing – review & editing; **Steven Van Petegem**: Data curation, Funding acquisition, Investigation, Project administration, Resources, Supervision, Writing – review & editing.
<figure><img src="image_48.png" /></figure>

## Disclosure statement

No potential conflict of interest was reported by the author(s).

## Funding

A.Ö. and S.V.P. acknowledge financial support by Board of the Swiss Federal Institutes of Technology (ETH Domain) within the SFA-AM MULTIMAT project. W.A.H. and S.V.P. thank Swiss National Science Foundation (SNSF) for financial support under grant agreement No 193799. This project has received funding from the European Union's Horizon Europe research and innovation programme under the Marie Skłodowska-Curie Actions grant agreement No 884104.

## Data availability statement

The data that support the findings of this study are openly available in https://zenodo.org/records/15222606. Derived data supporting the findings of this study are available from the corresponding author S.V.P on request.

## ORCID

Andaç Özsoy ID http://orcid.org/0000-0002-7745-2211

William A. Hearn ID http://orcid.org/0000-0001-9106-7945

Yunhui Chen ID http://orcid.org/0000-0002-1684-2507

Zoltan Hegedüs ID http://orcid.org/0000-0003-2691-8111

Nicola Casati ID http://orcid.org/0000-0002-4206-9239

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Appendix

<figure><img src="image_52.png" /><figcaption>Figure A1. Temperature-strain function of 316L, Cu and FCC(Fe).</figcaption></figure>

Polynomial functions fitted to Figure A1 are given below, where T is temperature and ϵ is strain.

$$ 
\begin{aligned}
T^{\text{FCC(Fe)}}(\epsilon) &= 4.26 \times 10^5 \epsilon^3 - 1.21 \times 10^5 \epsilon^2 + 5.36 \times 10^4 \epsilon + 2.00 \times 10^1 \\
T^{\text{316L}}(\epsilon) &= 6.42 \times 10^5 \epsilon^3 - 2.01 \times 10^5 \epsilon^2 + 5.37 \times 10^4 \epsilon + 2.52 \times 10^1 \\
T^{\text{Cu}}(\epsilon) &= 4.25 \times 10^5 \epsilon^3 - 6.32 \times 10^5 \epsilon^2 + 5.96 \times 10^4 \epsilon + 1.98 \times 10^1
\end{aligned}
$$