# Directed energy deposition of aluminium bronze onto H13 steel - microstructure in multi-material additive manufacturing

Håkon Linga, Constantinos Hatzoglou, Yubin Zhang, Vegard Brøtan, Ida Westermann, Xiaobo Ren, Bjørn Holmedal

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Håkon Linga, Constantinos Hatzoglou, Yubin Zhang, Vegard Brøtan, Ida Westermann, et al. Directed energy deposition of aluminium bronze onto H13 steel - microstructure in multi-material additive manufacturing. 2025. hal-04924408

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# Directed energy deposition of aluminium bronze onto H13 steel - microstructure in multi-material additive manufacturing

Håkon Linga<fnref n="1" />, Constantinos Hatzoglou<fnref n="1" />, Yubin Zhang<fnref n="2" />, Vegard Brøtan<fnref n="3" />, Ida Westermann<fnref n="1" />, Xiaobo Ren<fnref n="4" />, and Bjørn Holmedal<fnref n="1" />

E-mail: hakon.linga@ntnu.no

**Abstract.** Multi-material additive manufacturing has over the years had increasing interest due to the possibility of combining attributes of different alloys, to create parts with tailored properties. This study investigates the combination of aluminium bronze and AISI H13 tool steel, aiming to utilize the thermal properties of bronze and the wear resistance of tool steel through laser-based directed energy deposition with powder feedstock. The characterization was done using a variety of techniques, including optical- and secondary electron microscopy, atom probe tomography and X-ray microtomography. The deposition of bronze onto steel resulted in an interface nearly free of imperfections between the steel and bronze, as well as a mixing zone where nearly spherical iron-rich particles were suspended in a bronze matrix. The miscibility gap in iron-copper liquid primarily caused the formation of these spherical iron-rich particles, and the separation of elements in the liquid phase is believed to prevent the formation of intermetallic particles. Additionally, the region around the mixing zone exhibited increased hardness, approximately 1 mm above and below the bronze-steel interface, compared to the bulk bronze and -steel.

## 1. Introduction

Significant effort has been devoted to combining different materials, such as steel and copper alloys, over the years, with techniques like laser welding [1] and electroplating [2]. The potential for multi-material joining has greatly expanded with the broader adoption of additive manufacturing. Among the various additive manufacturing methods, laser-based directed energy deposition (DED-LB) stands out as the most practical for multi-material applications, often utilizing powder as feedstock [3]. Although powder bed fusion (PBF) has also been employed for multi-material additive manufacturing [4], it presents various practical challenges compared DED-LB, in addition to size limitation and deposition rate usually related to PBF.

One advantage of multi-material additive manufacturing is the ability to create parts with tailored properties that would be challenging to achieve with a single material. For instance, in

<footnote n="1">¹ Department of Materials Science and Engineering, Norwegian University of Science and Technology, Sem Sælands vei 14, 7034, Trondheim, Norway.</footnote>

<footnote n="2">² Department of Civil and Mechanical Engineering, Technical University of Denmark, Produktionstorvet, 425, 2800 Kgs. Lyngby, Denmark.</footnote>

<footnote n="3">³ SINTEF Manufacturing AS, S.P. Andersens veg 3, 7031, Trondheim, Norway.</footnote>

<footnote n="4">⁴ SINTEF Industry, Forskningsveien 1, N 0370 Oslo, Norway</footnote>
injection moulds, tool steels are typically used for their excellent wear resistance and mechanical properties, but have issues with heat dissipation. Combining tool steel with a copper-based alloy can address this issue effectively, due to the high thermal conductivity of copper-based alloys. Specifically, this study will examine the deposition of aluminium bronze containing approximately 10 wt% aluminium onto AISI H13 tool steel.

The phase diagram of aluminium bronze is quite complex, especially around 10-20 wt% aluminium in copper, where many different phases are possible [5]. At 10 wt% aluminium the equilibrium structure at room temperature consists of two phases, the face centered α, and a phase usually referred to as γ₂ [6], which is Al₄Cu₉ structured [7]. However, at higher cooling rates, a martensitic phase emerges instead. This forms when the body centered β phase is quenched and transforms into a face centered cubic structure [8].

Although H13 can have many different microstructures [9], it becomes close to 100% martensitic with sufficient cooling rate, such as in directed energy deposition [10]. The martensite is a large contributing factor to its superior mechanical properties, although the formation of various carbides, such as M₂C₆, M₇C₃, M₂C and MC contribute greatly to abrasive resistance and hardness [11]. Due to this, heat treatment is usually needed to reach peak hardness for conventionally manufactured H13 tool steels [12]. This is, however, not the case when it is produced with DED-LB, where the 'as-built' state is often optimal [13].

Even though the mentioned combination of material properties are promising, one should be vary when combining iron with copper. Copper is usually not used as an alloying element in steels, since the solid solution solubility of copper in iron matrix is very low. This leads to copper being pushed to grain boundaries, and the solidification shrinkage of the copper causing hot cracking [14]. Copper and iron also forms a miscibility gap [15, 16, 17], which can lead to similar issues [18]. Additionally, due to the sheer amount of elements present in the two alloys, there is the possibility of intermetallic particles. These are often unwanted since they are often brittle and lead to weakening of the material interface [19].

Even with these limitations, H13 steel and various copper alloys have previously been successfully combined, both as simple deposition onto a substrate, and as functionally graded material [20, 21]. The same applies to aluminium bronze to various steel alloys [22, 23]. Additionally, a preliminary study with aluminium bronze deposited onto H13 tool steel showed that there is good cohesion between the two materials, even under significant load [24]. The current research goes more in depth in the microstructural features present in and around the H13 steel to bronze interface, and discusses both the formation of the features and the effect they have on the mechanical properties. This is done with the bigger aim of acquiring a better understanding of the solidification process in multi-material additive manufacturing.

## 2. Materials and methods

The material examined in study is aluminium bronze deposited onto AISI H13 tool steel. All samples were produced using the same DED-LB cell, specifically a TRUMPF TruLaser Cell 3000, with powder as feedstock. Approximately 5 mm of H13 tool steel was first deposited onto a carbon steel substrate, before aluminium bronze was immediately deposited onto the H13 tool steel without any delay when switching powders. Another thin layer of steel was deposited onto the 5 mm thick layer of bronze, creating a sandwich type structure. The total dimensions of the sample, excluding the initial substrate, was therefore approximately 12 mm in height, with a base area of 50x50 mm. However, only the first interface between steel and bronze and the
surrounding area was focused on in this work, illustrated in Figure 1. The building direction is upwards in the figure, and the bi-directional alternating scan path is shown.

<figure><img src="image_2.png" /><figcaption>**Figure 1.** An overview of buildup of the sample regions and the deposition directions.</figcaption></figure>

The same building parameters were used throughout the steel and bronze, respectively. These are listed in Table 1, where the laser power is 500 W and the deposition speed is 1000 mm/min. The only difference in deposition parameters between the two materials was that the steel has a higher spot size and powder feed rate, with 1000 µm and 5 g/min, compared to the 800 µm and 3.8 g/min used for the bronze. Both materials were deposited at a hatch spacing of approximately 0.6 mm, programmed, and also experimentally measured in an optical microscope.

<table><caption>**Table 1.** Deposition parameters for H13 tool steel and aluminium bronze</caption><thead><tr><th>Material</th><th>Composition<br>[wt%]</th><th>Particle size<br>distribution<br>[µ]</th><th>Laser<br>power<br>[W]</th><th>Spot<br>size<br>[µ]</th><th>Deposition<br>speed<br>[mm/min]</th><th>Powder<br>feed rate<br>[g/min]</th></tr></thead><tbody><tr><td>H13<br>tool steel</td><td>Cr 5.40<br>C 0.38<br>Mo 1.38<br>Si 1.18<br>V 1.09<br>Mn 0.41<br>Fe balance</td><td>D10 53.8<br>D50 88.1<br>D90 144.0</td><td>500 W</td><td>1000</td><td>1000</td><td>5.0</td></tr><tr><td>Aluminium<br>bronze</td><td>Al 8.5 - 10.75<br>Fe 0.5 - 2.0<br>Cu balance</td><td>-125 +45</td><td>500 W</td><td>800</td><td>1000</td><td>3.8</td></tr></tbody></table>

In addition to the sample described above, another sample was made with the exact same building parameters, but where the H13 surface was ground down with an angle grinder before deposition of the bronze. This created a more reflective surface in the steel, altering the substrate for the first layer of bronze. The sample was used to investigate if the surface had an impact on the microstructure on the subsequent deposited bronze.
The material was initially cut to inspect cross-sections from various orientations. For microstructure characterization, each sample was embedded in epoxy and then ground with increasingly finer SiC paper. Following this, the samples were polished with diamond suspensions and finished with OP-U polishing. Samples intended for electron backscattered diffraction (EBSD) underwent additional vibration polishing to achieve a sufficient surface finish. To reveal the microstructure in light optical microscope, two different etchants were used. The first was an iron chloride based etchant made up of 10 g $FeCl_3$, 50 mL HCl and 200 mL $H_2O$, and the other was Beraha's tint etchant for tool steels, made up of 1 g potassium metabisulfite, 1 mL HCl and 100 mL $H_2O$ [25].

Optical microscopy was done with a Leica MEF4 M, whereas the electron microscopy was carried out with a ZEISS Ultra 55 LE FEGSEM, equipped with a Bruker Quantax energy-dispersive spectroscopy (EDS) detector and a NORDIF EBSD detector. Matlab was used for data processing of EBSD orientation data, by using the toolbox MTEX [26, 27]. The sample used for X-ray microtomography was fully made by manual polishing, into a rod shaped sample with approximately 0.5 mm x 0.5 mm square cross section, and the data was recorded with a ZEISS Xradia 520 Versa. Absorption contrast images were recorded, at 966 incident angles with 100 seconds of exposure time each. The distance between X-ray source and the sample and distance between sample and detector were approximately 9500 mm and 10000 mm, respectively. The optical magnification was 20X, the pixel size was 0.66 µm and the binning was set to 2.

Specimens for atom probe tomography (APT) analysis were made from aluminum bronze grains and iron particles by using an FEI Helios G4 UX focused ion beam (FIB). The APT specimens were prepared from site-specific lift-outs following procedures described in detail in [28]. APT experiments were carried out on a Local Electrode Atom Probe (LEAP) 5000XS from Cameca Instruments. It is a straight flight path instrument with a detection efficiency of 80 %. For all analysis, the base temperature was set to 50 K, and the laser energy was adjusted to obtain an equivalent pulse fraction of 20 % of the direct current voltage to avoid preferential evaporation [29] (i.e., around 100 pJ). A detection rate of 0.5 % and a pulse frequency of 200 kHz were used for all samples analyzed. Reconstructions were made dynamically [30] using the structural information according to [31] by using the Integrated Visualization and Analysis Software (IVAS) by Cameca Instruments Inc. The identification of precipitates was carried out using the iso-position method (IPM) [31, 32]. Their size (i.e., radius) was measured based on the radius of gyration in the evaporation direction [33], to limit biases caused by trajectory aberrations in the plane perpendicular to the evaporation direction [34, 35]. Their core chemical composition was measured using a radial distribution function (RDF) centered on the center of mass of each precipitate identified by IPM.

### 3. Results

When aluminium bronze is deposited onto H13 tool steel, some of the substrate has to be melted to attain proper cohesion. The melting steel mixes into the first deposited layer of aluminium bronze, creating the region labeled the primary mixing zone in Figure 1. This zone, depicted in the optical microscopy images in Figure 2, is located above the H13 substrate and below the secondary mixing zone, containing mostly aluminium bronze. Its thickness ranges from 0.3 to 0.4 mm, depending on the measurement location. From the optical microscopy it is clear that the primary mixing zone consists of iron rich particles in an aluminium bronze matrix. Furthermore, the scan lines are clearly visible from the density of particles, shown from a collection of particles along the previous melting pools in Figure 2.
<figure><img src="image_3.png" /><figcaption>**Figure 2.** Optical micrographs of primary mixing zone from three different perspectives. The dotted red line shows the approximate location of the plane perpendicular to the building direction in respect to the planes perpendicular to the scan- and transverse directions.</figcaption></figure>

A wave-like pattern emerges in the primary mixing area in Figure 2 due to collection of iron particles, and is also shown in three dimensions in Figure 3, as isosurface X-ray microtomography data. Figure 3 was created from computed microtomography absorption data, where atomic number contrast was measured and a threshold was set to distinguish the iron-based particles from the surrounding bronze. Although a general trend in the iron particle placement is revealed, the resolution is not high enough to distinguish all the individual particles. This is both due to small particle size and high quantity of particles, which cannot be resolved based on current microtomography resolution of 0.66 µm. Despite this, some large particles can be distinguished, especially above the primary mixing zone, where the density of particles is lower.

<figure><img src="image_4.png" /><figcaption>**Figure 3.** Isosurface plot where a threshold in X-ray microtomography absorption data was set to visually distinguish iron-rich and copper-rich areas within primary mixing zone.</figcaption></figure>

While it is difficult to determine the morphology of the iron particles in the different zones in
Figures 2 and 3, this is more clearly revealed in the mixing zone in Figure 4 with backscattered electron- and secondary electron images. Figure 4 a) shows that a majority of the particles are spherical of varying radius, whereas more irregularly shaped particles are found along previous melt pool boundaries. A closer look in Figure 4 b) shows that iron dendrites grow from the substrate below, in addition to forming between particles. Some particles are also fused together, and larger particles are found to have copper-rich dispersoids within them, as revealed by the secondary electron image in Figure 4 c).

<figure><img src="image_5.png" /><figcaption>**Figure 4.** a) Backscattered electron image of primary mixing zone. b) Backscattered electron image of interface between H13 tool steel and primary mixing area. c) Secondary electron image of a spherical steel particle in the primary mixing area. All three images are perpendicular to the scan direction in the primary mixing zone, with the building direction upwards.</figcaption></figure>

Elemental mapping of the major alloying elements– iron, chromium, copper and aluminium– for the particle in the primary mixing zone highlighted in Figure 4 c), is shown in Figure 5. This mapping indicates minimal mixing of elements between the two phases. Although the elemental mapping in Figure 5 shows a higher aluminium concentration within the iron particles, this is attributed to a lower signal-to-noise ratio. Point measurements within the particles result in approximately 10-12 wt% copper and 2-3 wt% aluminium, evenly distributed within the particles. The same phenomenon explains the seemingly larger detection of chromium in the bronze matrix. The presented EDS measurements could, however, not discern the individual particles that were shown in Figure 4 c), due to a too large interaction volume between the electron beam and the sample.
<figure><img src="image_6.png" /><figcaption>**Figure 5.** EDS elemental mapping of iron, chromium, copper and aluminium in the spherical particle shown in Figure 4 c).</figcaption></figure>

APT samples were fetched from the bronze grains between the iron-rich particles distributed in the primary mixing zone (Fig. 6 a). These contained a fine (Fig. 6 c), dense ((7.6 ± 0.3)·10²³ m⁻³), and homogeneous (Fig. 6 a, S1 to S3) distribution of spherical nano-precipitates (average radius of 1.7 nm) that are primarily enriched in Fe, but also in Cr and Al, compared to the matrix (Fig. 6 c and Table 2). The nominal composition is very close to the reference alloy, i.e., aluminium bronze (Table 1). However, trace amounts (i.e., < 0.1 at.%) of other chemical elements, particularly Cr, are also present and enrich the precipitates. A variation in the chemical composition of the precipitates with their size is also observed. For example, the Fe content ranges from 40 at.% for smaller precipitates (< 1nm) to nearly 80 at.% for larger ones (> 8nm, Fig. 6 a, S2). Since these are high-field precipitates (i.e., the evaporation field of the precipitates is higher than that of the surrounding matrix), there is no chemical composition bias in APT measurements due to local magnification effects [36, 37, 38]. This variation in precipitate composition with size is therefore not an APT artifact, but a genuine change.

The microstructure observed in the iron particles is much more complex (Fig. 6 b). Indeed, even during the preparation of APT samples via SEM, a backscattered electron image of such a needle (Fig. 6 b) shows different phase contrasts. By overlaying the APT analyses with the SEM images, the bright contrasts correspond to carbides (Fig. 6 b, S1), while the dark contrasts correspond to aluminum bronze phases (Fig. 6 b, S4). However, the main phase is a phase whose chemical composition (Table 2) closely resembles that of the reference H13 steel. This phase also contains a fine (Fig. 6 c) and dense ((5.3 ± 0.1)·10²³ m⁻³) distribution of spherical nano-precipitates (average radius of 1.7 nm), enriched in Cu and Al compared to the matrix (Table 2), and conversely depleted in Fe and Cr (Fig. 6 c). Numerous nanometric carbides are also present in this phase in all analyses (from S1 to S4, Fig. 6 b). Near these carbides, a precipitate-free zone (PFZ) of CuAl-enriched precipitates is observed in all cases. As seen previously, the composition of these precipitates varies with their size. For example, the Cu content ranges from 15 at.% for smaller precipitates (< 1nm) to nearly 75 at.% for larger ones (> 8nm). However, unlike the previous case, these CuAl-enriched precipitates are low-field precipitates (i.e., the evaporation field of the precipitates is lower than that of the surrounding matrix). This can potentially lead to biases in the composition of matrix atoms (mainly Fe, Table 2) due to ion trajectory overlaps caused by local magnification effects [39]. By applying the APT chemical composition correction model developed by Hatzoglou et al. [40], the composition of these precipitates is then consistent across sizes and very close to that of the aluminum bronze phase (i.e., 85 at.% Cu and 15 at.% Al). In this case, the variation in chemical composition with precipitate size was an APT artifact. As for the carbides, whether they are visible by SEM contrast or nanometric carbides in the APT volumes, they are all enriched in Mo, V, and Cr.
<figure><img src="image_7.png" /></figure>

<figure><img src="image_8.png" /></figure>

<figure><img src="image_9.png" /></figure>

<figure><img src="image_10.png" /></figure>

<figure><img src="image_11.png" /><figcaption>Distance from the precipitate core (nm)</figcaption></figure>

<caption>**Figure 6.** APT reconstruction from aluminum bronze grains and iron particles. (a) Atomic distribution in aluminum bronze grains, highlighting Fe-rich precipitates in samples S1 to S3. (b) SEM image of an APT sample showing phase contrast between multiple phases. Atomic distributions within iron particles primarily reveal Cu-rich precipitates and carbides across samples S1 to S4. In sample S4, the presence of an aluminum bronze phase containing Fe-rich precipitates is also observed. (c) Size distribution and composition profiles of Fe-rich precipitates in aluminum bronze grains and Cu-rich precipitates in iron particles. Composition profiles are obtained using radial distribution functions (RDF) centered on the mass centers of the precipitates.</caption>
For example, the carbide in Fig. 6 b S1 is composed of 40 at.% V, 30 at.% C, and 15 at.% Cr and Mo (likely of type $M_7C_3$ or $M_2C$). The aluminum bronze phase observed by APT within the iron particles (Fig. 6 b S4) has a chemical composition very close to the reference phase (Table 1), but it also contains trace amounts of all the chemical elements of H13 steel. Additionally, the same Fe-enriched precipitates (in size and chemical composition) as those observed in the aluminum bronze grains are also present. However, in this case, the distribution within this phase is not homogeneous, as they are only located near the interface with the H13-type phase, with a PFZ nearby.

<table><caption>**Table 2.** Elemental composition (at.%) in aluminum bronze grains and iron particles, comparing nominal (all collected atoms), matrix, and precipitate compositions. The statistical uncertainty is below 0.05 at.%. Elements marked as ‘n.s.’ (not significant) indicate negligible concentrations (i.e., < 0.1 at.%), while ‘-’ denotes elements not detected.</caption><thead><tr><th rowspan="2"></th><th colspan="3">Aluminium bronze grains</th><th colspan="3">Iron particles</th></tr><tr><th>Nominal</th><th>Matrix</th><th>Precipitate</th><th>Nominal</th><th>Matrix</th><th>Precipitate</th></tr></thead><tbody><tr><td>Al</td><td>14.73</td><td>14.75</td><td>15.54</td><td>7.04</td><td>6.58</td><td>12.42</td></tr><tr><td>Cu</td><td>81.68</td><td>83.25</td><td>31.69</td><td>6.76</td><td>2.09</td><td>20.50</td></tr><tr><td>Fe</td><td>3.36</td><td>1.85</td><td>51.22</td><td>77.33</td><td>81.59</td><td>60.60</td></tr><tr><td>Cr</td><td>0.10</td><td>0.04</td><td>1.29</td><td>5.10</td><td>5.39</td><td>3.32</td></tr><tr><td>Mn</td><td>n.s.</td><td>n.s.</td><td>n.s.</td><td>0.39</td><td>0.37</td><td>0.32</td></tr><tr><td>Si</td><td>n.s.</td><td>n.s.</td><td>n.s.</td><td>1.79</td><td>1.88</td><td>1.77</td></tr><tr><td>V</td><td>n.s.</td><td>n.s.</td><td>n.s.</td><td>0.67</td><td>0.72</td><td>0.29</td></tr><tr><td>Mo</td><td>-</td><td>-</td><td>-</td><td>0.83</td><td>0.88</td><td>0.58</td></tr><tr><td>C</td><td>-</td><td>-</td><td>-</td><td>0.41</td><td>0.44</td><td>0.09</td></tr></tbody></table>

Figure 7 a) shows the crystallographic orientation of both iron particles and aluminium bronze as measured by EBSD. In this case the solid color areas are iron particles, whereas striped areas are bronze grains. The iron particles are single grained, ignoring nanoparticles within the particles that are too small to be discerned with EBSD. While indexed to ferrite, the nano sized carbides inside make them resemble coalesced bainite [41]. The orientation of these particles seem to be random, shown from the pole figures in Figure 7 b), although particles close in proximity have a tendency to have similar orientations. The copper rich grains were also found to have approximately random orientation.

The size distribution of the particles is given in Figure 7 c), where the measured diameter follows approximately a log-normal distribution with a mode of 4 µm. This distribution was not found from the EBSD data in Figure 7 a), instead it was measured from an optical micrograph. It should be noted that while this gives an indication of the distribution of the size of the particles, the true 3D size of each particle is not found here. Since the measurements were made from a two dimensional cross section, the measured size is dependent on both the size of the particle and where on each particle the cross section cuts. Therefore, the true size is larger, but has likely less variance.

Although the orientations of the copper grains seems to be random, some grains are elongated diagonally between the building- and transverse directions. This is shown to occur in the center of previous melt pools, as in Figure 7 a), but is more commonly observed along the boundaries of the melt pools. This is shown in Figure 8, where only the orientations of copper grains are given. Many of these long diagonal grains seem to cross the boundaries between melt pools. This means that the existing grains from the previously deposited bronze continues to grow as
<figure><img src="image_12.png" /><figcaption>**Figure 7.** a) EBSD grain orientation map in the center of a primary mixing zone melt pool, with building direction as projection axis. Steel particles are shown as solid color, bronze grains are striped, and areas not indexed are white. b) Pole figures for the iron particles in the (100)- and (110) directions. c) Distribution of iron particle diameter in primary mixing zone. Measurements were done in a cross section perpendicular to the scan direction, and the distribution is made up of 543 particles.</figcaption></figure>

the consecutive line is deposited, in contrast to nucleating new grains which has to occur near the steel substrate.

The primary mixing zone can readily be distinguished in both optical- and electron microscope. However, the adjacent material, both below and above, are also influenced. From the overview of the sample seen in Figure 1, the secondary mixing zone is located above the primary mixing zone. In this area, some of the H13 steel that was included into the primary mixing zone is remelted, and solidifies along with the bronze. A small portion of this area can be observed in the top part of Figure 8, revealing a much finer grain structure. A closer look at the secondary mixing zone is given in Figure 9 a) and b). The EBSD map in Figure 9 b) shows small equiaxed grains with diameters of 1-2 µm, with small amounts of iron particles of similar size distributed throughout the bronze. Additionally, the X-ray microtomography data in Figure 3 shows some larger particles evenly distributed in the secondary mixing zone, with diameter greater than 10 µm. This contrasts the martensitic microstructure that is achieved in the bulk bronze, shown in Figure 9 c) and d).
<figure><img src="image_13.png" /><figcaption>**Figure 8.** EBSD grain orientation map of the bronze in the primary mixing zone, projected along the building direction. The dotted line shows the boundary between two melt pools.</figcaption></figure>

<figure><img src="image_14.png" /><figcaption>**Figure 9.** Optical micrographs and EBSD maps of aluminium bronze at different areas of the sample. a) and b) are in the secondary mixing area, c) and d) are in the bulk bronze. The bronze was etched with ferric chloride for the optical micrographs. Areas not indexed are shown as white in b) and d), and iron particles are outlined with thick lines.</figcaption></figure>
Measurements of microhardness in both the primary- and secondary mixing zones show that the mixing of iron into the bronze increases the mechanical strength. This is shown in Figure 10, where microhardness is plotted against the distance from the H13-aluminium bronze interface. Compared to the hardness measurements of around 170 HV3 in the bulk bronze, the hardness in both mixing zones exceeds 200 HV3, reaching around 220 HV3 in the primary mixing zone.

<figure><img src="image_15.png" /><figcaption>**Figure 10.** Microhardness around the H13 - aluminium bronze interface. Measurements are done with loads of 3 kg and 1 kg, marked as blue stars and red dots, respectively.</figcaption></figure>

In addition to an increased hardness in the primary- and secondary mixing zone above the steel - bronze interface, the steel below also has a significantly increased hardness, around 550 HV3 compared to 500 HV3 in the bulk steel. This area is around 1 mm thick and is denoted the affected zone in Figure 10. One clear effect the deposited bronze has on the substrate, is to introduce copper and aluminium into the steel. This is shown in Figure 11 b) and c) with EDS elemental mapping, along with backscattered electron image of the same area in Figure 11 a). In addition to copper rich particles suspended into the steel, the elemental mapping shows an even distribution of copper and aluminium within approximately 15 µm below the interface.

Point measurements for the marked areas in Figure 11 a) are given in Table 3, showing that the copper concentration is around 8-12 wt% and aluminium 2-3 wt% throughout the area, giving it approximately the same concentration of copper and aluminium as the spherical iron particles within the primary mixing zone. The only exception is point 7, where the interaction volume is likely large enough to include an adjacent copper particle. However, below this area virtually no copper nor aluminium is found.

However, the hardness increase seen in the affected zone in Figure 10 cannot be explained by
<figure><img src="image_16.png" /><figcaption>**Figure 11.** a) Backscattered electron image from interface between H13 steel and primary mixing area. b) EDS elemental mapping of copper. c) EDS elemental mapping of aluminium.</figcaption></figure>

<table><caption>**Table 3.** EDS elemental measurements for points given in Figure 11. All measurements given in weight percentage.</caption><thead><tr><th>Point</th><th>1</th><th>2</th><th>3</th><th>4</th><th>5</th><th>6</th><th>7</th><th>8</th><th>9</th></tr></thead><tbody><tr><td>Fe</td><td>77.71</td><td>78.09</td><td>77.65</td><td>78.17</td><td>81.38</td><td>81.93</td><td>65.47</td><td>77.96</td><td>79.75</td></tr><tr><td>Cr</td><td>4.74</td><td>4.88</td><td>4.19</td><td>4.98</td><td>4.92</td><td>4.88</td><td>5.33</td><td>4.69</td><td>4.76</td></tr><tr><td>Mo</td><td>1.64</td><td>1.92</td><td>1.24</td><td>1.71</td><td>1.56</td><td>1.47</td><td>2.19</td><td>1.61</td><td>1.68</td></tr><tr><td>V</td><td>1.10</td><td>1.19</td><td>0.86</td><td>1.17</td><td>1.04</td><td>1.06</td><td>1.73</td><td>1.09</td><td>1.08</td></tr><tr><td>Si</td><td>0.89</td><td>0.97</td><td>0.78</td><td>0.91</td><td>0.78</td><td>0.73</td><td>0.82</td><td>0.99</td><td>0.92</td></tr><tr><td>Mn</td><td>0.22</td><td>0.27</td><td>0.18</td><td>0.25</td><td>0.28</td><td>0.33</td><td>0.24</td><td>0.21</td><td>0.30</td></tr><tr><td>Cu</td><td>11.00</td><td>10.03</td><td>11.65</td><td>10.45</td><td>8.13</td><td>6.83</td><td>21.22</td><td>10.64</td><td>9.15</td></tr><tr><td>Al</td><td>2.70</td><td>2.65</td><td>3.43</td><td>2.36</td><td>1.90</td><td>1.71</td><td>3.02</td><td>2.80</td><td>2.37</td></tr></tbody></table>

the copper and aluminium concentration shown in Figure 11, since the zone width is two orders of magnitude smaller. The H13 tool steel below the primary mixing zone is shown in Figure 12, etched with Beraha's tint etchant. This type of etchant is often used to distinguish between different phases in steels by creating sulfate layers of different thicknesses [25, 42]. The gradient in color displayed therefore corresponds to a gradient in the thickness of this layer, progressing towards a brown to orange color that is present in the remaining bulk material. This in turn suggests a change in reaction kinetics of the etching due to changing microstructure. The thickness of the area is roughly equal to the area with elevated measured hardness, shown in Figure 10.

The H13 steel is martensitic, and contains some retained austenite. This can be seen in the bulk steel as colored grains in Figure 13 c). The retained austenite can be seen in the affected zone in Figure 13 b) as grains outlined by black lines. Although the number of retained austenite grains per area is similar in the two cases, it is clear that the austenite grains in the bulk steel are vastly larger. This makes the total area percentage of austenite much larger in the bulk steel, even with seemingly similar grain structure otherwise in the martensite.

By polishing the substrate, a rather flat interface such as the one shown in Figure 14 a) is achieved. With this treatment the first layer of aluminium bronze is deposited without signif-
<figure><img src="image_17.png" /><figcaption>**Figure 12.** Three optical micrographs stitched together of area below steel - bronze interface. Etched with Beraha's etchant.</figcaption></figure>

<figure><img src="image_18.png" /><figcaption>**Figure 13.** a) Image quality map from EBSD data around the H13 - aluminium bronze interface. b) EBSD map of marked area shown in a). c) Calculated grains from EBSD data in bulk steel. Colored grains show austenite grains, and blue lines show subgrain boundaries.</figcaption></figure>

icant steel remelting. The EBSD map here shows a grain structure close to the one attained in the secondary mixing shown in Figure 9 b). However, a drawback is that the interface is prone to common additive manufacturing imperfections, such as lack of fusion [43] and gas pore formation [44, 45], clearly shown in Figure 14 b). While the H13 steel substrate had increased reflectivity due to polishing, the edges of the sample were not polished. The effect of this can be seen in Figure 14 c), where an amount of remelted steel is clearly visualized in this region.

An increase in austenite percentage was found in the steel just below the interface, with area percentage measured to approximately 14%. This contrasts the decrease that was found below
<figure><img src="image_19.png" /><figcaption>**Figure 14.** a) EBSD map of H13 - aluminium bronze interface without a large mixing zone. b) Optical micrograph of H13 - aluminium bronze interface showing various imperfections. c) Optical micrograph of edge of sample where mixing between H13 and aluminium bronze is achieved.</figcaption></figure>

the primary mixing zone, shown in Figure 13. Hardness measurements were done around this interface as well, and while the aluminium bronze showed a slight increase in hardness in close proximity to the interface, a similar increase in hardness in the H13 steel was not observed. Instead it remained approximately the same as the bulk steel.

**4. Discussion**

One of the first things one notices when inspecting the primary mixing zone is the spherical iron particles. This is not unique to this system, instead it seems to be common when aluminium bronze is deposited onto steel [23]. While the dendrites between the particles form during the solidification of the primary mixing zone, it is harder to see how the spherical particles can solidify in such a short time. Instead, it is likely that they have formed due to a miscibility gap, and therefore separated already in the liquid. If the particles observed in this work is compared to other particles formed due to liquid phase separation, in the copper-iron system, there are clear similarities in morphology of the particles [16, 18].

The random texture found in both the spherical iron particles and bronze matrix may at first seem a bit strange. Often in additively manufactured samples, grains are oriented so that (100)-directions are oriented parallel to the temperature gradient [46]. The temperature gradient is dependent on the shape of the melt pool, but is on average somewhere between the scan direction and building direction. However, this only occurs if nucleation is not necessary, in other words when grains can continue growing from previously nucleated grains. Since the substrate is H13 steel, this is not the case. The only exception is the boundary to the previous melt pool in the primary mixing zone. Figure 8 shows this exact case, where new grains do not
nucleate, but rather grows from previous grains. Therefore there is likely a weak texture in the bronze matrix in the primary mixing zone.

Although the iron particles also were shown to have random orientation, there is in many cases similar orientation in adjacent particles, as shown in Figure 7 a). This could indicate that there is some kind of orientation relationship between the iron particles and copper grains. However, many of the iron particles are big enough to be adjacent to several copper grains, and vice versa. Therefore it is extremely hard to resolve any such relationship, and would additionally require three dimensional EBSD data. If there is an orientation relationship between the steel particles and bronze matrix, however, that would mean that the bronze matrix solidifies first. However, the melting point of aluminium bronze is significantly lower than that of H13 tool steel, this would require substantial undercooling of the steel melt if bronze were to solidify first. This has been shown to occur in nano-sized particles [47, 48], but since the ones in this work are in the size range of 5 to 10 µm, it seems unlikely.

Another explanation, and perhaps a more probable one, could be that some of the spherical particles may be fused together already in the melt, even when they appear to be separated in the EBSD scan. The reason for this is the same reason why the particle size distribution in Figure 7 c) does not show the full picture. The EBSD data are taken from a polished plane in the sample, so the particles are cut at random cross sections. At the thickest part they may be fused together, while they appear to not be so in the inspected cross section. This is shown particularly in Figure 3, where individual particles are difficult to distinguish. Although this is partially due to the small particle size and high density of particles, overlaps would make distinguishing them difficult.

One important aspect of the interface between the H13 steel and the aluminium bronze is the lack of intermetallic particles. With the exception of nanoparticles, such as the ones shown in Figure 6, no intermetallic particles could be detected. This is positive, as these are usually brittle and detrimental to the mechanical properties of the material. One reason for this is that the main components of the two alloys that are mixed, iron and copper, do not create intermetallic phases alone [49]. There are, however, many elements combinations that can create intermetallics, such as iron and aluminium [50], in addition to the countless combinations that are possible with the other elements present. Despite this, EDS elemental mapping shows that the elements belonging to the H13 tool steel is mainly contained within the areas rich in iron, and aluminium is mainly found in the aluminium bronze grains. A significant contributing factor to this is that the two materials are separated already in the liquid before solidifying, leading to less opportunity for formation of intermetallic phases. The miscibility gap between copper and iron is therefore in this case positive, which is usually not the case. The miscibility gap and the low solubility of copper in solid iron are the primary reasons why copper is usually not used as alloying element in steels, as it can lead to hot cracking [14, 51, 52].

Regarding the increase in hardness in the affected zone of the steel, there are two potential causes. The first is a chemical difference, where copper and aluminium diffuse into the steel, leading to hardness increase through solid solution hardening. However, this can largely be dismissed due to the absence of copper and aluminium 10 - 15 µm beneath the interface. Furthermore, the deposition happens quickly, so any large scale diffusion is unlikely to occur during the deposition. In fact, the 10 - 15 µm of enrichment of copper and aluminium is likely not due to diffusion, as there is no gradual decline in concentration. Instead, this part of the steel was likely melted along with the primary mixing zone, especially considering the composition is nearly identical to the spherical iron particles.
It should be noted that the investigation does not take into account the carbon content, as accurate measurements of carbon are not possible with EDS. Carbon is a small element, taking interstitial positions in steel and could therefore diffuse faster than copper and aluminium. One could imagine a scenario where the steel and bronze melted into a single liquid phase instead of separating into two. Since addition of copper in iron melt drastically reduces the solubility of carbon [15, 53], the carbon from the steel would evaporate or dissolve into the substrate. This could result in a higher concentration of carbon in the area below the primary mixing zone. However, it is unlikely that the liquid was ever a single phase, since the low solubility of carbon in an iron-copper melt is a reason for the existence of the miscibility gap between iron and copper in the first place [54]. APT measurements nominal concentration of carbon within the iron particles reinforces that carbon does not diffuse into the substrate. The measured carbon content within the particles is approximately equal to the carbon content in the feedstock powder.

Instead, it is more realistic to assume that the main contributing factor to the hardness increase is the heat input from the bronze deposition. Isolated, depositing bronze introduces heat into the H13 steel substrate. This would lead to annealing of the martensite and a reduction in hardness. However, it is important to note that each layer of the bulk steel has effectively been annealed by the following deposited layer. Therefore, one has to compare the heat input from depositing a layer of bronze to depositing a layer H13 steel to understand the hardness increase. While the laser power was equal while depositing bronze and steel, bronze is generally more reflective, leading to less laser absorption and thus less effective heat input [55]. This in turn means that the martensitic steel is less annealed than if another layer of steel was to be deposited. This is likely the interpretation of the tint etching in Figure 13, where the kinetics behind the reaction is dependent on the degree of annealing of the martensite.

An explanation for the varying austenite percentage follows the same logic. Deposition of additional layers could cause the preceding material to reach not only annealing temperatures, but could also reach the temperatures where austenite becomes thermodynamically stable. This could lead to growth of austenite grains. A smaller heat input when depositing bronze compared to steel would result in growth to a less degree. While it may seem strange that the austenite percent is higher in the sample without significant mixing, it is important to remember that the steel here has effectively been heated an additional time, since some of the deposited steel was removed by grinding. Additionally, depositing a layer of material where the heat input is too low to remelt the preceding H13 steel can counter-intuitively lead to a higher austenite percentage. Since none of the heat is used to melt the steel, leading to new rapid solidification, the heat instead contributes to the austenite transformation [56].

The reason why the area below the steel-bronze interface is denoted the affected zone in Figure 1 and 10, and not the heat affected zone, is that the term heat affected zone is a bit misleading in this case. While it per definition is a heat affected zone, all of the deposited steel is at some point affected by the heat of the subsequent layer. The affected zone is in fact affected by less heat than the remaining material.

The hardness increase in the secondary mixing is more straight forward to grasp, as this is due to the fine grain size that is achieved. This grain refinement is likely due to suppression of the bronze martensite, since iron is known to occupy nucleation sites for the high temperature BCC phase in the aluminium bronze system [8].
## 5. Conclusions

By depositing aluminium bronze onto H13 a layer of densely packed spherical steel particles and dendrites in aluminium bronze matrix is built. The spherical steel particles are single grained ferrite, with nano sized carbides and copper rich particles inside. The particles have random texture, but particles in close proximity tend to have similar orientation. No intermetallic particles are present in the mixing zone, which is attributed to the miscibility gap between iron and copper separating the two materials before solidification. The layer of H13 below the deposited aluminium bronze is harder than the bulk steel, likely due to less annealing than if steel was deposited onto it. Aluminium bronze can be deposited onto H13 tool steel without forming a layer of spherical steel particles and dendrites. This can be achieved using the same deposition parameters, but by simply altering the surface reflectivity through machining. However, maintaining a mixing zone between the materials is often preferred to avoid lack of fusion defects.

## Acknowledgments

This research is based on the research activites within the center for research-based innovation, SFI Manufacturing in Norway and is partially funded by the Research Council of Norway (NFR) under contract number 237900. A special thanks to Nordic Additive Manufacturing AS and the national research project IPN NextMoldMaker hosted by Skriverform AS (Project No.: 281967, NFR) that have contributed with test parts, case description and AM parameter development. The Research Council of Norway (RCN) is acknowledged for its support to the Norwegian Micro- and Nano-Fabrication Facility, NorFab (No. 295864) for the use of the FIB facilities, the Norwegian Laboratory for Mineral and Materials Characterization, MiMaC (No. 269842) for the use of the Atom Probe facility. The authors acknowledge the use of 3D Imaging Center (3DIM) at the Technical University of Denmark. Furthermore, Y Zhang acknowledges the support by a research grant (VIL54495) from the VILLUM FONDEN.

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