Investigation of Scaling and Materials’ Performance of EHLA-Fabricated Cladding in Simulated Geothermal Brine

Abstract

This study investigates the corrosion and scaling behaviour of Extreme High-speed Laser Application (EHLA)-fabricated corrosion-resistant alloy (CRA) claddings under simulated geothermal brine conditions. EHLA 316L stainless steel and alloy 625 coatings were produced and tested in simulated brine (chloride–carbonate–silica geothermal brine) at 70 °C for 720 h to evaluate the influence of additive manufacturing (AM) microstructures on corrosion performance. The EHLA coatings exhibited dense, metallurgically bonded microstructures with minimal porosity. Microstructural analysis revealed Nb- and Mo-rich segregation in EHLA 625 and fine columnar dendritic morphology in all coatings. EHLA 625 developed a stable passive film with only a thin deposit of Mg-O-containing compounds, whereas EHLA 316L exhibited localised pitting and significant Si- and Mg-containing scale accumulation, especially in as-built conditions. Surface finishing reduced corrosion activity by minimising roughness and defect-driven localised attack. Critical pitting temperature (CPT) tests confirmed the superior localised corrosion resistance of EHLA 625 relative to EHLA 316L under laboratory conditions. While these results indicate promising corrosion and scaling resistance of EHLA coatings, further process optimisation and post-deposition thermal treatments might be required to achieve coating performance comparable to wrought alloys. The results indicate the potential of EHLA-fabricated coatings for producing corrosion and scaling resistance surfaces.

1. Introduction

Geothermal energy is a promising source of renewable and clean energy. However, the harsh conditions of geothermal environments pose significant challenges to the materials used in production equipment. Recent efforts to drill deeper geothermal wells to achieve up to 10 times increase in energy output compared with the conventional wells have resulted in a variety of failures due to the higher temperature, lower pH, supercritical state of the geothermal fluid, higher corrosivity, etc. [1, 2]. These geothermal fluids often contain high concentrations of solids, chloride ions, hydrogen sulphide (H₂S), and carbon dioxide (CO₂). Such conditions can be highly aggressive to carbon and low-alloy steels commonly used in geothermal wells [3, 4, 5, 6].

Various mitigation strategies have been explored to address the failures due to the high corrosion susceptibility of carbon steel in such environments. The use of corrosion-resistant alloys (CRAs) to mitigate corrosion has been widely considered [7, 8, 9]. However, their high cost makes full component fabrication economically unviable. Therefore, using CRA weld overlay or cladding is a more cost-effective alternative that has been successfully applied in other energy industries. Inconel 625 and 316LSS, known for good weldability and corrosion resistance, are among the candidates. CRA weld overlays offer good resistance to corrosion and erosion, while the backing carbon steel provides appropriate mechanical properties such as high strength.

Laser cladding is an advanced surface engineering technique for protecting underlying material using a deposition layer. In this process, a cladding material is deposited on a surface with thermal energy supplied by a laser beam [10, 11, 12]. While the wire-feed method is compatible with lasers, most claddings use powder due to better versatility [13, 14]. The overlay cladding is commonly used to enhance corrosion and tribological performance of the surface [15, 16, 17]. The Extreme High-speed Laser Application (EHLA) process which is a state-of-the-art technology based on laser cladding can operate at ultra-high processing speeds. In EHLA, the laser focuses above the surface of the substrate to melt the metal powder mid-air. On reaching the substrate surface, the molten droplet solidifies rapidly and forms a metallurgically bonded coating [18, 19]. EHLA can achieve deposition rates of up to 200 m/min and processing efficiencies approaching 500 cm²/min [20]. The high processing speed of this method allows large-scale fabrication to be carried out more economically, which is essential for its viability in the renewable energy sector. Another study compared AISI 431 stainless steel claddingdeposited by EHLA, and conventional laser processes and showed several benefits in using EHLA such as four times higher deposition rate and a more refined microstructure [21]. The fine microstructure was due to the ultra-high cooling rate achieved in EHLA technique.

Although CRA overlays improve corrosion performance of the substrate material for geothermal applications, their influence on scale deposition rate and associated issues are not yet well understood. A range of factors influence the tendency for scale formation, including temperature and concentration of species such as silica, calcium, carbonates, sulphates, and other heavy metal ions [22]. The cladding material and surface roughness are some additional parameters that influence scaling on clad surfaces.

This work examines the suitability of EHLA-produced corrosion-resistant alloy (CRA) overlays, namely Inconel 625 and 316 stainless steel, for service in high-silica geothermal environments. Particular attention is given to the role of surface condition (as-deposited versus as-polished) in governing corrosion behaviour and scaling tendencies under simulated geothermal exposure. While the microstructure and performance of EHLA-fabricated coating in the simulated geothermal brine was the focus of this study, the electrochemical results were also compared with those of conventional wrought alloys as a benchmark.

2. Materials and Methods

2.1. Materials and Laser Cladding Equipment

To compare the performance of conventionally manufactured (wrought) materials with those produced by the EHLA process, tests were carried out on samples fabricated using both methods. The materials examined in this study are as follows: precipitation hardened nickel alloy 718, Inconel^® 600, austenitic stainless steel 316L, and nickel alloy 625 (the last two were manufactured using EHLA). Alloy 718 and Inconel^® 600 were selected to provide an intermediate level of corrosion resistance between 316L and 625, although it is recognised that these alloys are not typically employed for the intended application. Both alloy 718 and Inconel^® 600 were purchased off-the-shelf from Goodfellow Cambridge Limited (Huntingdon, UK).

All claddings in this work were produced using the EHLA machine manufactured by Hornet Laser Cladding (Alblasserdam, The Netherlands) (Figure 1). The machine features 10-axis manipulation with a 6-axis robot, linear guide, and two spindles with capacity for parts up to 3 m long and 1 ton in weight. The laser specifications are a 11.5 kW laser with adjustable spot size for enhanced process flexibility.

For the EHLA-fabricated materials, trials were conducted to identify the best processing conditions that yield minimal porosity and defects (such as lack of fusion, micro-cracks, and inclusions) without compromising deposition efficiency. The manufacturing parameters investigated are spot size, laser power, speed, powder fed rate, feeder, and pitch. The detailed results of the trials can be found in reference [23]. The particle size ranged between 15 and 45 µm and 25 and 45 µm for 316L and nickel alloy 625, respectively. Carpenter Additive and AP&C were the source of the metal powders 316L and Inconel 625, respectively. The parameters used to produce the 625 and 316L samples are shown below in

Table 1. Process parameters employed during cladding deposition for the specimens used in the corrosion tests.

Material	Spot Size (mm)	Laser Power (kW)	Speed (m/min)	Powder Feed Rate (g/min)	Feeder (RPM)	Pitch %	Pitch (mm)
625	3	2	10	15	1.1	80	0.4
316L	3	2.5	10	25	4.6	70	0.48

. The coatings were manufactured under argon shielding gas.

The EHLA coating was deposited on commercially available carbon steel substrate purchased from Apollo Metals. The nominal composition of the carbon steel and the powders are presented in

Table 2. Nominal composition of the carbon steel (CS) and the two clad materials.

Element	316L wt%	Inconel 625 wt%	CS Substrate wt%
Fe	Balance	≤5	Balance
Cr	16.0–18.0	20.0–23.0	-
Ni	10.0–14.0	58.0 min	-
Mo	2.0–3.0	8.0–10.0	-
Nb	-	3.15–4.15	-
Al	-	≤0.40	-
Ti	-	≤0.4	-
Mn	≤2.0	≤0.50	0.8
Si	≤1.0	≤0.50	0.2
C	≤0.03	≤0.10	≤0.03

.

The assessments of the microstructure were carried out only for the EHLA manufactured layers, using samples prepared via conventional metallurgical procedures (i.e., sectioning from cladding, followed by grinding, polishing, and etching). The microstructure of the alloys studied was revealed by Kalling’s no. 2 reagent. Optical scanning electron microscopy (SEM) coupled with energy-dispersive X-ray spectroscopy (EDX) and electron backscatter diffraction (EBSD) were used for the assessment of the microstructure. The percentage of porosity in the materials was estimated by image analysis of the metallurgical sections using ImageJ 1.53k (National Institutes of Health, Bethesda, MA, USA) in as-polished condition (before etching).

2.2. Scaling and Corrosion Tests

2.2.1. Test Performed Simulating Geothermal Conditions

To assess materials’ degradation behaviour under representative geothermal conditions, corrosion and scaling tests were performed using a custom-built vessel with rotating paddles in the laboratory. Details of the test rig can be found in reference [24]. Test coupons fabricated from the alloys mentioned in Section 2.1 were mounted inside the vessel and instrumented for electrochemical measurements, allowing real-time monitoring of surface changes during exposure.

The simulated geothermal brine was formulated based on field analyses from Nesjavellir subcritical fluids [25]. The composition was selected to include corrosive species such as chloride and dissolved CO₂, as well as scaling agents including calcium, magnesium, and silica. Solution preparation was carried out in stages: NaOH (0.1 M, 2 g in 500 mL DI water) was used to dissolve 1.804 g of 325-mesh SiO₂ (solution 1), while 24.2 g of NaHCO₃ and 7.64 g of Na₂CO₃ were dissolved in 2 L of DI water (solution 2). After mixing the two solutions, CaCl₂·2H₂O (0.5388 g) and MgCl₂·6H₂O (0.7398 g) were added, and the final mixture was diluted to 4 L with DI water. All the chemicals used in this study were analytical reagent-grade. The as-mixed pH was measured to be 10.5. The subcritical fluid of Nesjavellir also contains H₂S, this was not included during these tests due to Health and Safety considerations. However, the solution was continuously purged and saturated with CO₂ during the 720 h test. The pH was dropped to 6.7 after purging with CO₂ at the beginning of the test and varied between 6.6 and 9.0 during the test.

Experiments were carried out at 70 °C, corresponding to the maximum temperature achievable by the test rig. While the subcritical fluid at Nesjavellir reaches higher temperatures, this testing regime was selected to allow early screening of materials, and the results should be considered as an indication for corrosion susceptibility at the tested temperature and to define preliminary operational limits. A 50 L reservoir was used to store the prepared solution, which was circulated through the test vessel using a peristaltic pump to ensure stable hydrodynamic and chemical conditions throughout the experiments. The fluid in the reservoir was continuously purged with CO₂ gas during the test.

2.2.2. Post-Tests Examination

Once the long-term tests simulating the geothermal conditions were completed, the samples were analysed visually, and images were taken to record observations. In the as-received post-testing condition, the surface of the samples was examined using light optical microscopy and scanning electron microscopy coupled with energy-dispersive X-ray spectroscopy (SEM-EDX). At this stage, EDX maps were taken from all the samples in an equal number of locations to quantify the nature (chemistry) and level of scaling in the samples.

Subsequently, cross-sectional metallographic specimens were prepared by sectioning through the regions of interest, mounting, grinding, and polishing to reveal the internal structure. These sections were examined to determine the depth and nature of corrosion damage, including pitting, intergranular attack, or scale formation.

2.2.3. Critical Pitting Temperature

In this study, the pitting corrosion resistance of the investigated EHLA materials was evaluated following a procedure based on ASTM G150 [26] using 35 × 35 mm specimens (prepared for electrochemical measurements ) in the as-built condition (the surface of the specimens was not ground). The procedure involved immersing the specimens in two different nitrogen deaerated solutions: (i) a solution containing 170 ppm of chloride ions and (ii) a standard ASTM G-150 solution (1 M of NaCl). The specimens were put under a controlled fixed potential while gradually increasing the solution temperature. The temperature at which current density rose above 0.1 mA/cm² was identified as the critical pitting temperature (CPT). Note: These tests were conducted in a slightly different solution to the geothermal brine, to have an idea of the influence of chlorides in the absence of the other species.

3. Results

3.1. Sample Manufacturing by EHLA

3.1.1. Inconel 625

The part manufactured to extract the coupons for the corrosion testing is shown in Figure 2a. The 10 mm thick disc was 150 mm in diameter. Metallographic sections and corrosion coupons were extracted from this flat disc (i.e., an example of the type of coupon used for the corrosion test is illustrated in orange in Figure 2a, the coupons for the metallurgical examination were extracted transverse to this). The as-built cladding surface was analysed using a 3D microscope and the roughness of the surface (Ra) was ~5.7 μm. The measured porosity percentage area was 0.12 ± 0.06 using ImageJ 1.53k software.

3.1.2. EHLA 316L

As with Alloy 625, a 150 mm-diameter steel disc was used as the substrate for depositing the 316L EHLA cladding. The roughness (Ra) of the as-built surface was ~7.3 μm. The measured porosity percentage area was 0.063 ± 0.04.

Note: For both 625 and 316L, the corrosion test coupons were carefully prepared to ensure that only the cladding surface was exposed, to avoid any influence of the steel base/substrate.

3.2. Microstructural Examinations

3.2.1. Inconel 625

The metallographic section extracted from the flat coupon was polished and etched appropriately. The revealed microstructure is shown in Figure 3. The microstructure exhibits a primarily elongated columnar dendritic structure growing perpendicular to the interface with the average thickness of 442 μm. The overlay exhibited a dense and uniform morphology, with no evidence of solidification cracking within the cladding material. However, some un-melted deposits were still present on the surface, indicating incomplete fusion in localised regions. High-magnification SEM imaging revealed distinct banding patterns consistent with elemental segregation, particularly of niobium and molybdenum, within the dendritic structure [27]. No ferrite formation was seen in the coating layer, with dilution of iron limited up to 20 µm into the coating.

3.2.2. 316L Stainless Steel

Figure 4 shows the revealed microstructure, which consists predominantly of elongated columnar dendrites aligned along the thickness of the cladding. The average measured thickness of the cladding layer is approximately 1240 μm. While the overlay appears generally dense, some linear and rounded features were observed. The linear indications are consistent with solidification cracks, whereas the rounded ones are likely due to porosity or non-metallic inclusions. SEM analysis of this sample showed a uniform microstructure with no signs of elemental segregation or a significant number of secondary phases. EDX analysis revealed iron diffusion from the substrate into the coating, observable only within the dilution-limited depth of 20 µm; however, austenitic and ferritic phases boundaries were clearly defined and, as observed, remained separated.

3.3. Flow Rig Test

Figure 5 shows a comparison of the coupons before and after testing, while representative images of the SEM examination after testing are presented in Figure 6. The surface of the coupons revealed varying degrees of surface alteration among the tested alloys. The 316L as-built coupon showed a yellowish hue on the surface after testing with visible signs of surface degradation. SEM examination revealed extensive surface coverage by deposits, with the underlying metal exposed only in isolated areas. Remnants of the original powder particles were still present, and EDX spectra indicated silica, oxygen, and magnesium scales (see Figure 6g); notably, the dark region at the base of right in Figure 6d corresponds to a silica-based scale.

The 316 ground coupon exhibited a noticeably darker appearance; however, the SEM image confirmed the absence of surface deterioration, with changes limited to scale formation. These scales were predominantly composed of silicon, oxygen, and magnesium, while other deposits rich in calcium and oxygen (with minor silicon content) were also observed.

Alloys 600 and 718 followed the same trend that was noticed in the 625 ground sample, in particular, the EDX spectra of the scales of alloy 600 identified magnesium, silicon, and oxygen as the principal elements. In 718, this type of scale was identified, but also with additional scales rich in calcium and oxygen, which agree with the white patches observed in Figure 5 after testing. The coupon of alloy 625 (ground) was the one that displayed less scaling.

For 625, the as-built sample showed no signs of surface degradation, although a significant layer of deposits was observed, primarily composed of magnesium and oxygen (this was the only sample that displayed significant amounts of this type of scale, see Figure 6h). In the ground condition, this alloy showed minimal surface change, with only a light layer of scale detected; EDX analysis identified magnesium, silicon, and oxygen as the principal elements in these scales, and no corrosion damage was evident.

Based on the elements identified in the EDX spectra of the scales, it could be speculated that the magnesium–calcium–oxygen, the magnesium–oxygen and the calcium–oxygen scales are carbonate scales. And although the exact nature is unknown, these could be compatible with magnesium containing calcite, magnesite, and calcite/aragonite, respectively. The silicon–oxygen scale is compatible with silica. Given that each of the scales changed from sample to sample, it is plausible to conclude that the substrate chemistry plays a role in the formation of these. Equally, another parameter that could have played a role in the scale formation is the evolution of solution chemistry over time. During the first 300 h of testing, the solution pH increased from approximately ≈6.6 to ≈9 and remained stable at this value up to the completion of the tests. Based on previous investigations by Scott et al. [24, 28], it is noted that silica precipitation tends to be higher at pH levels below 9, which could suggest that the silica scaling observed in the samples occurred during the first 300 h. The carbonate scaling is expected to start slightly later during the tests, as the pH was initially lower than 8. Therefore, the type of scale present in the samples is not only influenced by its initial surface chemistry but also due to the silica on the surface of the samples.

The open circuit potential (OCP) measurements of the six tested materials, 316L ground, 316L as-built, 625 ground, 625 as-built, 718, and 600, presented in Figure 7 revealed a similar electrochemical shift during the early stages of exposure. Within the first 200 h, all samples experienced a transition from more negative to more positive potentials, suggesting surface evolution likely linked to passive film formation or stabilisation. Among these, the 316L as-built specimen stands out due to the OCP values being consistently higher throughout the test, this is an unexpected result as more positive OCP tend to be associated with less reactive passive films. By the completion of the tests, 718, 625 as-built, and 600 show a very stable OCP, suggesting a steady electrochemical environment. However, for 316L ground and 625 ground, the trend indicates that the OCP is not stable yet, tending to decrease. The OCP of 718 although displaying similar OCP values to the other materials, exhibited greater fluctuations, pointing to possible changes in the surface of the sample, this seems to be related to deposit formation rather than localised events of substrate corrosion. However, it is interesting that 718 is only showing these fluctuations despite all the samples showing some level of surface deposition. Another alternative explanation could be electrical connection issues; however, the pattern of fluctuation is not compatible with this.

Figure 8 presents the cross-sectional view of the 316L as-built sample after testing. As expected for this alloy in a chloride-containing environment, the corrosion observed is localised, manifesting primarily as pitting. Figure 8a highlights the development of secondary pits within existing surface pits, suggesting that despite the flow conditions during testing, localised concentration of aggressive species is occurring. This may occur due to the scales observed in the sample or due to un-melted surface particles that shield specific regions. At higher magnification (Figure 8b), the pit bottom reveals a growth pattern that appears to follow the orientation of the dendritic microstructure. The pit propagation occurs in a toothed/serrated manner, indicating that the solidification structure of the as-built material influences the pit growth.

3.4. Critical Pitting Temperature

Table 3. Results of the ASTM G150 tests and modified G150 tests with lower chloride concentration.

Material	Environment	CPT Value
316L	1 M NaCl	7.7 °C
316L	170 ppm chloride	>80 °C
625	1 M NaCl	>80 °C
625	170 ppm chloride	>80 °C

shows a summary of the results of ASTM G150 tests. As expected, 625 displayed, in general, higher resistance to localised corrosion than 316L. Also, as expected, the environment containing the higher number of chlorides (1 M) proved to be more aggressive. It was interesting to observe that the CPT values of the EHLA 316L and EHLA 625 at 170 ppm chloride were >80 °C (this was the temperature limit of the CPT testing rig). In spite of this similarity in the CPT of the two EHLA coatings, different results were obtained in the tests simulating the geothermal brine conditions.

In addition to the ASTM G150 test results, Figure 9 presents the current density versus time curves recorded during the tests. Although the 625 alloy successfully passed, the data revealed that the current generated during exposure to 170 ppm chloride concentration brine was higher than that observed for 316L. In absence of any other knowledge of the material, this is unexpected given the well-established superior corrosion resistance of alloy 625.

4. Discussion

The as-built EHLA coatings exhibited substantially higher surface roughness than the ground samples, increasing the electrochemically active area and contributing to higher measured currents during testing. Although this difference introduces some variability when comparing the two conditions, it also reflects a practical consideration relevant to AM technologies, namely the possibility of deploying components without post-processing. For this reason, both as-built and ground surfaces were evaluated. However, roughness alone does not fully explain the differences observed. Prior studies on AM 316L have demonstrated that the effect of surface finish on corrosion behaviour is primarily governed by changes in surface chemistry and oxide composition rather than by topography itself. Teo et al. [29] showed that different post-processing routes modify the chemical state of the passive film, while Clark et al. [30] reported that variations in oxide chemistry, not roughness, control the electrochemical response of LB-PBF 316L. Thus, while increased roughness enhances the electrochemically active area, the dominant factor influencing corrosion performance is the stability and composition of the passive film, which is directly shaped by deposition-induced microsegregation and surface heterogeneity. These findings imply that for EHLA-produced alloys, the observed differences between as-built and ground conditions are likely linked to differences in passive film composition and surface heterogeneity introduced during deposition. Thus, the effect of surface condition must be interpreted as a combined effect of both morphology and surface chemistry.

To facilitate understanding of the experimental findings, potentiodynamic sweeps were carried out in the geothermal brine. The results are shown in Figure 10. Alloy 625 in the as-built condition exhibited the lowest OCP among the tested materials. This result is somewhat unexpected, as OCP typically reflects the thermodynamic tendency of a material to corrode. Among the alloys tested, 625 contains the least iron, which would suggest a higher OCP. However, the as-built 625 surface displayed several features detrimental to corrosion resistance, including high surface roughness, surface heterogeneities, and chemical segregation. These factors are likely to contribute to the unusually low OCP observed for this alloy. This indicates that the surface heterogeneity can override thermodynamic expectations from bulk composition, demonstrating that passive film chemistry and microstructural uniformity are the primary controls on corrosion behaviour. All the other alloys showed similar OCP values. In terms of i_corr, 316 as-built displayed the higher value, while all the alloys displayed relatively similar values. Analysis of the anodic branches reveals clearer distinctions in localised corrosion susceptibility. Excluding the anomalous response of as-built 625, EHLA 316L showed the lowest pitting potential, followed by alloy 600 and then ground 625, consistent with their respective alloy chemistries and microstructures. Alloy 718 displayed the lowest anodic current densities, reflecting its robust passive film stability. For the as-built 625, the apparent absence of a well-defined pitting potential, coupled with the presence of a sustained high-current plateau, indicates passive film instability likely linked to melt pool segregation and surface heterogeneity intrinsic to the as-built microstructure.

If the results of the different tests are compared, no one-to-one correlation can be identified. For 316L as-built, the OCP measurements were unexpectedly higher, despite this being the only sample that exhibited corrosion in the long-term test. This apparent contradiction can be explained by considering the chloride dependence of corrosion: modified ASTM G150 results show very low currents at low chloride concentrations, but a sharp increase in current at high chloride levels, indicating that the passive film remains largely stable under mild conditions but fails catastrophically once critical chloride thresholds are exceeded (it should be noted that other factors might play a role as the flow; also note the importance of long-term duration tests to assess materials corrosion performance). CO₂ also appears to assist in overcoming this threshold. The measured high i_corr, low pitting potential, and the presence of localised pits/crevice attack detected by the post-exposure surface examination in the long-term tests indicate that the passive film of 316L is unstable in this environment. This finding is consistent with the established understanding that elevated temperatures (>55–60 °C) and high chloride concentrations promote passive film breakdown and accelerate localised corrosion processes in austenitic stainless steels. Microscopic examination revealed subsurface pits, nucleation of pits inside primary pits, and crevice-like attack under deposits/corrosion products, etc., which are characteristic of chloride-induced pitting corrosion. The initiation of pitting likely occurred due to localised de-passivation when chloride ions locally penetrated the thin chromium oxide film, establishing an anodic site. The potential presence of dissolved CO₂ in the simulated geothermal brine would further exacerbate this behaviour, even without chlorides [31]. The synergistic interaction between chlorides and CO₂ at elevated temperature (~70 °C) creates an aggressive environment that possibly exceeds the materials’ resistance to corrosion in this condition [32]. Importantly, the relatively high pitting potential observed for 316L ground and conventionally manufactured 316L suggests that removal of surface defects through grinding could restore passivation capability, reduce microgalvanic sites, and increase resistance to localised attack. This highlights the practical role of surface engineering in enhancing the corrosion performance of as-built EHLA 316L. Regarding the pit morphology found in 316L, although no chemistry segregation or secondary phases were identified, the structure at the bottom of the pits suggested that the pit growth is governed by the microstructure; Voisin et al. [33] and Wang et al. [34] studied the microstructure of laser powder bead fusion (LPBF) 316L in detail. They found that in this alloy, there is a cellular structure within the grains creating partitioning of the elements. This observation is compatible with the morphology of pit growth presented in this paper. No pit growth rate was established during this study; however, if this cellular structure leads to potentially higher corrosion rates (to contrast this, Chao et al. [35] found that in samples of 316L fabricated using selective laser melting (SLM), the pitting potential was higher than in conventionally manufactured wrought 316L), then it implies that the qualification of AM material might require more detailed information that is generally used for wrought alloys. For example, the qualification of the materials might require macrographs, micrographs, and advance microscopic imaging.

Regarding 625, despite no pitting or localised corrosion events being identified during the long-term tests, this indicated some corrosion related information such as the increase in current during the ASTM G150 or the poor behaviour during the potentiodynamic scans. Thus, differentiating the 625 EHLA from their wrought counterpart, which is expected to show better corrosion indicators than any of the alloys presented in this study. The differences can be rationalised by examining the unique microstructural characteristics imparted by the AM process. Unlike wrought material, which typically exhibits a more homogeneous, often equiaxed, and well-controlled microstructure, AM alloys are shaped by rapid solidification, localised melting, and layer-by-layer deposition, factors that promote microstructural heterogeneity, elemental segregation, and defect formation, all of which influence corrosion performance.

One of the primary contributors to this increased corrosion susceptibility is the non-homogeneous nature of deposit formation, as observed in Figure 3d. For example, previous long-duration tests in sulfuric acid have revealed shallow but distinct selective attack along melt pool borders, indicating that these features create local electrochemical inhomogeneities that accelerate localised corrosion processes in LPBF 625 [36]. While wrought Alloy 625 generally forms equiaxed grains with more uniform chemistry, the layer-wise deposition in AM introduces melt pool boundaries and microsegregation that do not exist in conventionally processed material. Marchese et al. [37] have found that within the microsegregation regions, i.e., in the interdendritic regions, which are typically rich in elements such as niobium (Nb) and molybdenum (Mo), precipitates such as niobium carbides or gamma double prime (γ″) are present. While no specific Mo-rich precipitates were identified (during this research, it was found that the bands observed in Figure 3d are rich in niobium and molybdenum, no Cr enrichment was identified), it is well-established that the precipitates deplete the adjacent gamma-Ni matrix of protective elements, reducing the resistance to localised corrosion. Porosity and other manufacturing defects further amplify this effect. Although laser-based methods typically achieve relatively low porosity, any residual voids, can serve as initiation points for corrosion attack. In contrast, wrought alloys typically exhibit no discernible porosity, and precipitation of secondary phases typically occurs only after prolonged exposure or poor heat treatment. This understanding, which enables effective control of sensitisation in wrought alloys and has begun to be applied to conventional laser additive manufacturing (AM) processes [38], still needs to be developed for EHLA-fabricated claddings.

In addition to the parameters discussed in this study, one critical aspect that was not characterised is the passive layer. The condition and integrity of the passive layer formed during EHLA (or, in general, for AM parts), compared to ground parts, remains unknown [33]. A complete understanding of the corrosion behaviour would benefit from further characterisation of the material’s composition, thickness, and electrochemical properties. The effect of deposition parameters such as heat input, cooling rate, and thermal cycles on the corrosion performance of EHLA-fabricated materials is also important but unknown. Without this, the conclusions from the corrosion tests remain preliminary.

5. Conclusions

The following conclusions can be drawn from the work presented in the paper:

•

EHLA-fabricated coatings exhibited dense microstructures with excellent substrate bonding and minimal porosity, confirming the potential of the process to produce coatings suitable for corrosion-resistant applications.

•

Although the findings reported in this investigation are only indicative of the corrosion performance under the tested laboratory conditions and further investigation would be required to confirm equivalence in industrial environments, ground 316L EHLA shows performance comparable to wrought 316L.

•

EHLA alloy 625 showed better corrosion and scaling resistance than EHLA 316L in simulated geothermal brine due to its stable passive film and higher alloying content. However, in the ground condition, after removal of the as-deposited surface the differences in corrosion performance are not significant.

•

Localised Nb- and Mo-rich segregation in EHLA 625 had some influence on the corrosion resistance under the tested conditions. Surface conditions strongly influenced performance. Processes such as grinding reduced scale deposition and enhanced electrochemical stability. Post-deposition surface engineering is required to enhance the performance of the deposits. However, with the limited studies carried out in this project it is difficult to state with certainty if further process parameter optimisation could achieve this.

The results presented in this paper demonstrate that EHLA offers considerable potential as a technique for producing corrosion-resistant claddings suited to demanding geothermal conditions. Nonetheless, further work such as (i) targeted optimisation of EHLA processing parameters to reduce segregation and improve microstructural uniformity, (ii) evaluation of defined post-deposition treatments, and (iii) systematic long-term testing under conditions representative of geothermal service to assess scaling behaviour and durability are required to ascertain the potential of EHLA as a technique to produce effective corrosion-resistant coatings.