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Article

Effects of Cathodic Hydrogen Charging on the Mechanical Properties and Fracture Behaviour of Wire Arc Additively Manufactured AA2319

by
Tinashe Mazarire
*,
Alexander Galloway
and
Athanasios Toumpis
Department of Mechanical & Aerospace Engineering, University of Strathclyde, 75 Montrose St, Glasgow G1 1XJ, UK
*
Author to whom correspondence should be addressed.
Metals 2026, 16(6), 647; https://doi.org/10.3390/met16060647
Submission received: 11 May 2026 / Revised: 3 June 2026 / Accepted: 10 June 2026 / Published: 12 June 2026
(This article belongs to the Special Issue Advanced Metallic Materials for Hydrogen Production, Storage and Transportation: Design, Degradation and Reliability)
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Abstract

The growing adoption of wire arc additive manufacturing (WAAM) requires an understanding of how WAAM-fabricated aluminium alloys respond to environmental factors that may degrade mechanical performance. This study investigates the effects of cathodic charging on the mechanical properties and fracture behaviour of WAAM AA2319 aluminium alloy. Cathodic charging was conducted in an electrolyte containing 3.5 wt.% NaCl and 3 g/L ammonium thiocyanate using different applied current densities. The resulting changes in mechanical performance were assessed through uniaxial tensile and Charpy impact toughness tests. The results demonstrate that cathodic charging led to a progressive reduction in ductility with increasing current density. Elongation decreased by up to approximately 45% relative to the uncharged condition, while ultimate tensile strength and yield strength were marginally affected. Charpy impact testing revealed a corresponding reduction in impact toughness of approximately 40% following hydrogen charging. Fractographic analysis showed a transition from ductile fracture dominated by microvoid coalescence in the uncharged material, to a mixed ductile–brittle fracture in hydrogen-charged specimens, characterised by shallow dimples and quasi-cleavage features. The observed changes in mechanical behaviour and fracture morphology suggest that cathodic charging promoted hydrogen-assisted mechanical degradation, with features consistent with hydrogen-enhanced localised plasticity (HELP) and hydrogen-enhanced decohesion (HEDE).

1. Introduction

Aluminium alloys are widely used in the electrical, biomedical, transport, and defence sectors due to their excellent electrical conductivity, high strength-to-weight ratio, and good corrosion resistance properties [1,2,3]. Among these, aluminium alloy (AA2319), an Al-Cu alloy, is of particular interest for structural applications requiring high strength, including cryogenic tanks [4,5,6] and aircraft fuselage [7,8]. This alloy also exhibits good weldability, which has led to extensive research into its application in wire arc additive manufacturing (WAAM) [9,10,11]. WAAM is an additive manufacturing technique that utilises the principle of fusion welding to deposit material [12]. This fabrication method offers higher deposition rates compared with other additive manufacturing techniques such as powder-based processes [13], while also reducing material waste and environmental impact relative to conventional subtractive manufacturing methods [14,15], making it a highly sustainable and resource-efficient fabrication process.
Despite these advantages, the structural integrity and environmental compatibility of WAAM-fabricated aluminium alloys remain an area of active research [16]. The layer-by-layer deposition process and associated thermal cycles result in heterogeneous microstructures, residual stresses and process-induced defects such as porosity and lack of fusion, all of which can influence the mechanical performance [17]. As the adoption of WAAM continues to increase, it is essential to understand how WAAM-fabricated aluminium components respond to environmental factors that may degrade their mechanical properties in service.
One such factor is hydrogen exposure. Hydrogen is becoming an increasingly recognised energy carrier in the transition to a net-zero future [18], and additively manufactured components are being considered for hydrogen infrastructure and structural applications [19] in which exposure to hydrogen may occur during processing or service [20]. During WAAM processing, hydrogen may be introduced through contaminated wire feedstock, shielding gas impurities, or atmospheric moisture [21]. In aluminium alloys, this is important because hydrogen has limited solubility in solid aluminium and may contribute to porosity during solidification [22]. In service, hydrogen uptake may arise from corrosion reactions, electrochemical environments, or surface interactions with hydrogen [23,24]. Consequently, it is crucial to assess hydrogen embrittlement in additively manufactured aluminium alloys to determine their compatibility within hydrogen environments.
Hydrogen embrittlement (HE) refers to the degradation of mechanical properties of a material due to the presence of hydrogen within the material often leading to premature failure [25,26]. Extensive research has been conducted on the susceptibility of wrought aluminium alloys to HE [27,28,29,30,31,32] and several mechanisms have been proposed to account for this behaviour. The primary HE mechanisms in aluminium alloys include hydrogen-enhanced localised plasticity (HELP), hydrogen-enhanced decohesion (HEDE) and the synergistic action of HELP and HEDE [33,34,35]. In the HELP mechanism, hydrogen facilitates plastic deformation by increasing dislocation mobility [36], often leading to a ductile fracture characterised by microvoid coalescence (MVC). In contrast, the HEDE mechanism involves a reduction in lattice cohesion due to hydrogen-induced weakening of atomic bonds or hydrogen segregation between matrix atoms, thereby decreasing the atomic binding strength [37] and promoting brittle, cleavage-like fracture. The synergistic effect of HELP and HEDE occurs when both mechanisms are active, with their dominance depending on the hydrogen concentration in the specimen [38,39].
HELP and HEDE have both been reported in aluminium alloys [40], although their relative contribution depends on factors such as hydrogen concentration and strain rate. For example, Baltacioglu et al. [33] investigated the effect of strain rate on the hydrogen embrittlement susceptibility of AA7075 and reported a synergistic contribution of HELP and HEDE at higher strain rates (1 s−1), with MVC attributed to HELP and quasi-cleavage features and deep cracks attributed to HEDE. Similarly, Hou et al. [41] reported that HELP dominates at relatively low hydrogen concentrations in 7xxx aluminium alloys, whereas HEDE becomes more prominent at higher hydrogen concentrations.
Cathodic hydrogen charging is a widely used method to introduce hydrogen into specimens and assess their susceptibility to hydrogen-assisted degradation [42]. The method involves cathodic polarisation in an electrolyte, often containing recombination poisons to promote hydrogen uptake [43]. During cathodic charging, hydrogen is generated at the specimen surface through electrochemical reactions, after which atomic hydrogen may adsorb onto the surface, absorb into the material, and diffuse into the bulk of the material. Hydrogen uptake during cathodic hydrogen charging depends on several factors including electrolyte composition, ccd, charging duration and the microstructure of the investigated material [44]. This method provides a useful approach for evaluating changes in mechanical response and fracture behaviour after hydrogen exposure.
Most existing studies on hydrogen embrittlement in additively manufactured materials have focused on alloys such as austenitic stainless steel [45,46], Inconel 718 [47,48] and Ti-6Al-4V [49,50,51,52]. There is currently a lack of published work addressing hydrogen embrittlement in WAAM-fabricated AA2319. Given the growing interest in utilising WAAM for hydrogen-related components, this represents a gap in current knowledge, and a lack of understanding of how hydrogen affects the mechanical properties of WAAM AA2319.
The present study addresses this gap by investigating the effects of cathodic hydrogen charging on the mechanical properties and hydrogen embrittlement behaviour of WAAM-fabricated AA2319. Specimens were subjected to electrochemical cathodic charging, followed by uniaxial tensile testing and Charpy impact testing to quantify changes in mechanical properties and fracture resistance under both quasi-static and impact loading conditions. Fracture surface analysis was then conducted to compare fracture morphology and damage characteristics between charged and uncharged specimens. The aim of this work is to assess how cathodic hydrogen charging influences the mechanical performance and fracture characteristics of WAAM AA2319, and to evaluate the observed degradation in relation to hydrogen-assisted embrittlement mechanisms.

2. Materials and Methods

ER2319 wire was used to fabricate the WAAM component. The wire had a diameter of 1.2 mm, and its chemical composition, as provided by the supplier (WB Alloys, Glasgow, UK), is presented in Table 1.
The bulk material used in this study was manufactured using WAAM based on cold metal transfer (CMT) technology. The deposition cell comprised an ABB IRB 2400L robot arm (Zurich, Switzerland) controlled by an ABB IRC5 control and a teach pendant, integrated with a Fronius CMT welding torch and wire feeding system. Figure 1 depicts the WAAM equipment in operation during the fabrication of the bulk material used in this study.
High-purity argon (99.97%) was used as the shielding gas during the deposition process, supplied at a constant flow rate of 20 L/min. The WAAM parameters employed to fabricate the bulk material are summarised in Table 2.
Following deposition, the as-deposited material was solution treated at 535 °C for 2 h, polymer-quenched and artificially aged at 191 °C for 20 h to achieve a T6 temper condition. Samples for microstructural analysis and mechanical testing were harvested from the bulk material through waterjet cutting. The sample extraction location and orientation are illustrated in Figure 2.

2.1. Metallographic Preparation and Characterisation

Samples for microstructural analysis were prepared using standard metallographic procedures. Samples were progressively ground using SiC papers (Struers ApS, Ballerup, Denmark) from 550 grit to 2400 grit, followed by final polishing with a 0.01 μm oxide polishing suspension to obtain a mirror finish. Scanning electron microscope (SEM) with backscattered electron (BSE) imaging was used for high-resolution microstructural analysis of the polished samples, while energy-dispersive X-ray spectroscopy (EDS) was used to determine chemical composition of intermetallic phases. The polished samples were then electro-etched using Barker’s solution (20 mL HBF4, 80 mL H2O) in order to reveal the microstructure and grain boundaries. Optical microscopy (OM) was the performed using an Olympus GX-51 light optical microscope (Tokyo, Japan) for colour imaging.

2.2. Mechanical Testing

Vickers microhardness testing was conducted on polished samples using a load of 200 g and a dwell time of 10 s. Both uniaxial tensile and Charpy V-notch specimens were machined from the extracted blocks to the required geometry. The tensile specimens had a cylindrical geometry as shown in Figure 2b, with dimensions in accordance with ASTM E8 [53], and were tested following the same standard. The Charpy specimens had dimensions shown in Figure 2c, conforming to ASTM E23 [54], and were tested in accordance with the standard. Triplicate tensile and Charpy tests were conducted for each test condition.

2.3. Hydrogen Charging and Testing

Hydrogen was introduced into both tensile and Charpy specimens via electrochemical charging. The electrolyte consisted of 3.5 wt% NaCl and 3 g/L ammonium thiocyanate (NH4SCN), the latter serving as a hydrogen recombination poison to enhance hydrogen uptake [29,43]. Each specimen functioned as the working electrode, while a platinum wire was used as the counter electrode. The cathodic charging setup is depicted in Figure 3. Cathodic charging was carried out at room temperature for 24 h at current densities ranging from 2 mA/cm2 to 8 mA/cm2.
Following charging, the tensile specimens were lightly abraded with fine grit paper to remove surface impurities and for a clean surface for strain measurement using a video extensometer. All tensile tests were initiated within 5 min of charging and were performed on a 50 kN Instron universal testing machine a strain rate of 1.6 × 10−2 s−1.
HE was quantified based on tensile test results by calculating the hydrogen embrittlement index (HEI) using the following equation:
H E I = φ u − φ c φ u
where φ u represents the parameter for uncharged specimen, and φ c represents the corresponding parameter for the hydrogen-charged specimen.
A subset of Charpy specimens was hydrogen-charged at a current density of 8 mA/cm2 for 24 h and subsequently tested to assess the effect of hydrogen on impact toughness. Charpy impact tests on hydrogen-charged specimens were also initiated within 5 min after hydrogen charging. Following impact testing, some samples were lightly abraded and subsequently subjected to macrohardness testing using a 5 kg load and a dwell time of 10 s on the surface that had been exposed to the electrolyte during hydrogen charging.

2.4. Fractography

Fractographic analysis was carried out on both tensile and Charpy specimens after mechanical testing to assess the influence of hydrogen on fracture behaviour. A Hitachi S3700-N tungsten-filament SEM (Tokyo, Japan) operating at 15 kV was used to characterise fracture surface morphology. EDS was conducted at selected locations on the fracture surfaces to characterise and determine local elemental composition.

3. Results and Discussion

3.1. Microstructural Characterisation

The polished surface of the WAAM AA2319 is shown in Figure 4a, which reveals porosity within the material. This porosity primarily arises from the difference in hydrogen solubility between the molten and solid aluminium. During deposition, hydrogen originating from the atmosphere and feedstock readily dissolves into the molten pool due to its high solubility in molten aluminium [22]. As the material solidifies, the solubility of hydrogen decreases and the excess hydrogen precipitates in the form of gas bubbles. During rapid solidification of the molten aluminium, these bubbles do not have sufficient time to escape and are therefore trapped within the solidifying matrix, resulting in the formation of porosity [21]. The volume fraction of porosity was measured (33.6 mm2 across nine deposited layers) and calculated as 0.2%. The average pore size was measured as 20 μm, while the maximum pore size was 47 μm. Figure 4b shows a histogram of the number of pores vs. pore diameter. The results illustrates that most of the pores are less than 25 μm.
The microstructure of the WAAM-produced specimen (Figure 5) is complex due to the layer-by-layer deposition process, which induces changing thermal cycles across each deposited layer [55,56]. Each deposited layer, located between two fusion lines as illustrated in Figure 5a, can be divided into interzone (ITZ) and inner zone (INZ). The INZ can be further subdivided into the lower inner zone (L-INZ) and the upper inner zone (U-INZ), with the L-INZ corresponding to the region adjacent to the ITZ and the U-INZ representing the top of the deposited layer (see Figure 5b).
The microstructure within each deposited layer comprises fine equiaxed, columnar, and equiaxed grains. Fine equiaxed grains are observed at the top of the fusion line within the ITZ, formed due to the rapid solidification of the deposited layer [57]. The average grain size in this region was 23 μm. Following this region is the L-INZ, characterised by columnar grains which grow in the cooling direction [58,59], with an average grain size of 57 μm. The upper portion of the deposited layer (U-INZ) is remelted and mixed with subsequent deposition, effectively acting as the heat-affected zone. Heat accumulation from subsequent layers, along with thermal influence from preceding layers, creates favourable conditions for equiaxed grain growth in this region [60], where the average grain size was 62 μm.
The SEM image of the polished surface (Figure 6) reveals the presence of second-phase particles distributed both along the grain boundaries and within the grains. EDS analysis of these precipitates (Figure 6b) indicates that they contain aluminium and copper, corresponding to Al-Cu eutectic phases [61]. The formation, segregation and evolution of eutectic phases during the WAAM process have been extensively discussed elsewhere [62,63].
For the alloy investigated in this study, which was subjected to a T6 heat treatment condition, solution heat treatment at 535 °C promotes the dissolution of second-phase particles formed during WAAM into the aluminium matrix [10]. Subsequent quenching retains a supersaturated solid solution, while artificial ageing leads to the precipitation of second-phase precipitates along grain boundaries and inside grains, leading to the observed microstructure. The microstructure of the WAAM-fabricated AA2319 examined in this study contains heterogeneous grain structures, precipitates, and porosity, all of which can influence hydrogen transport and trapping behaviour [64]. Consequently, these microstructural features may contribute to local variations in hydrogen-assisted degradation and embrittlement susceptibility.

3.2. Mechanical Properties

3.2.1. Hardness Results

The microhardness results are presented in Figure 7a. The measured hardness values ranged from 136 HV to 160 HV, with an average of 148 HV. This variation in hardness across the sample corresponds to differences in microstructural features between the zones formed during the WAAM fabrication process. As shown in Figure 7a, regions containing fine equiaxed grains exhibited higher hardness values, whereas zones with larger or columnar grains displayed comparatively lower hardness.
Finer grains have a greater number of grain boundaries per unit area compared with coarser grains. During indentation, each grain boundary acts as a barrier to dislocation motion [65], thus impeding plastic deformation [66]. Consequently, regions with finer grains have higher hardness due to higher grain boundary density within the indented area. In contrast, coarser grains contain fewer grain boundaries to restrict dislocation motion, resulting in lower hardness values. This behaviour can also be interpreted using the Hall–Petch relationship, which describes the relationship between a material’s yield strength (and consequently hardness) and its grain size [67]:
σ y = σ 0 + k d − 1 2
where σy is the yield strength, d is the grain size, σ0 and k are constants. σ0 is the friction stress required to move dislocations and k is the Hall–Petch slope.
As the grain size decreases, the material’s hardness increases and vice versa. The fine grains had an average grain size of 23 μm, whereas the grain size of the coarse grains was 62 μm. Hence, there is higher hardness in the fine grain zone and lower hardness in the coarse grain zone. The hardness variation observed across the deposited layers reflects the influence of the non-uniform thermal cycles during the WAAM process, which governs grain size and structure. Overall, the hardness values obtained for the studied material fall within the range reported in the literature for the WAAM-fabricated AA2319 in T6 condition (143 HV–150 HV [10,68]).
The macrohardness results shown in Figure 7b indicate that the surface hardness of the material increased by approximately 14% following hydrogen charging at 8 mA/cm2. El-Amoush [69], similarly observed increased hardness in hydrogen-charged AA7075, which was attributed to hydrogen-induced increase in dislocation density. Comparable trends have also been reported in steels, where hydrogen charging leads to an increase in hardness due to hydrogen–dislocation interactions [70,71]. The increase in hardness observed in the present study suggests a corresponding increase in the brittleness of the material following hydrogen charging.

3.2.2. Tensile Test Results

Figure 8 presents the tensile test results for uncharged and hydrogen-charged WAAM AA2319 specimens. The results show that the uncharged specimen exhibits superior mechanical properties compared to all charged specimens.
An increase in charging current density (ccd) leads to an overall reduction in elongation, indicating a progressive loss of ductility with increasing hydrogen charging. The mean percentage elongation decreases from 4.2% ± 0.3 in the uncharged condition to 3.1% ± 0.2 at 2 mA/cm2 and 3.3% ± 0.1 at 4 mA/cm2, followed by a pronounced reduction to 2.3% ± 0.2 at 6 mA/cm2 and 2.5% ± 0.1 at 8 mA/cm2. The reported elongation values correspond to the arithmetic mean of three specimen per test condition, with the associated ± values representing the standard deviation. At lower current densities, the reduction in ductility is less pronounced, whereas at higher current densities there is greater loss of ductility.
Moderate variation in elongation is observed between intermediate charging conditions—for example, an 8.7% increase in elongation between the 6 mA/cm2 and 8 mA/cm2 conditions; however, the absolute difference in mean values between the two (2.5% and 2.3%) is only 0.2% which is comparable or in the same order of magnitude as the standard deviation (~0.1–0.3%). The microstructure of the WAAM material is inherently inhomogeneous, as shown in Section 3.2.1, and this is expected to contribute to scatter in mechanical behaviour even between nominally identical specimens. In addition, hydrogen concentration in the specimens is governed by hydrogen diffusion which is significantly influenced by the microstructure and microstructural defects such as grain boundaries and porosity which all act as hydrogen traps [72]. Consequently, these differences can lead to variations in hydrogen concentration and embrittlement severity [73]. The observed difference between the intermediate specimens is therefore attributed to microstructural and hydrogen-related variability rather than to a statistically meaningful change associated with current density. Nevertheless, the overall trend demonstrates a progressive reduction in elongation with increasing current density.
The HEI values for UTS, yield strength and elongation are illustrated in Figure 9. The effect of hydrogen embrittlement is most pronounced for elongation, with a decrease of approximately 45% calculated between the uncharged specimens and the specimens charged at 6 mA/cm2, indicating a substantial loss of ductility due to hydrogen charging. In contrast, the HEI values for yield and UTS illustrates minor changes as depicted in Figure 9a,b. This indicates that for WAAM-fabricated AA2319, cathodic hydrogen charging primarily affects ductility rather than strength, a trend consistent with previously reported findings in the literature [30,33] for other materials.

3.2.3. Charpy Impact Results

The Charpy impact results in Figure 10 show that the uncharged specimens have an average impact energy of 5.3 J, whereas the specimens charged at 8 mA/cm2 exhibit a reduced average impact energy of approximately 3.3 J. This corresponds to a reduction of approximately 40% in impact toughness following hydrogen charging. The percentage decrease is equivalent to the HEI value calculated for the specimen charged at 8 mA/cm2. These results demonstrate that hydrogen charging significantly degrades the toughness of WAAM AA2319, leading to a more brittle response. This trend is consistent with the tensile results presented in the Section 3.2.2, which also showed a reduction in ductility following hydrogen charging. Similar reductions in toughness after cathodic hydrogen charging have also been reported in the literature [34,74].

3.3. Fractography

Figure 11 reveals the fracture surface of the uncharged tensile specimen, together with the magnified image of the fractured region. The fracture surface displays a microvoid coalescence (MVC) morphology characterised by a high density of dimples. The MVC process initiates at inclusions (highlighted in Figure 11b), where voids nucleate, grow, and coalesce, leading to the observed ductile fracture mode [33]. The observed fracture is similar to the fracture behaviour of WAAM AA2319 reported in previous studies [58,75]. Additionally, porosity is evident on the fracture surface, which originates from the WAAM process as previously discussed in Section 3.1.
The tensile fracture surfaces of the uncharged and hydrogen-charged tensile specimens are presented in Figure 12. Higher magnification SEM images (×1000) acquired from the middle section of the fracture surfaces shown in Figure 12 are depicted in Figure 13.
In the cathodically charged specimens, increasing the ccd leads to a progressive transition from deep dimples to shallow dimples and, ultimately, quasi-cleavage fracture as illustrated in (Figure 13b–e). After cathodic charging at 2 mA/cm2 and 4 mA/cm2, the fracture surfaces still showed dimpled morphology; however, the dimples became shallower and quasi-cleavage fractures appeared in the specimen charged at 4 mA/cm2. The presents of shallow dimples suggest that cathodic charging altered the local plastic deformation behaviour during fracture. For cathodically hydrogen-charged specimens, such features are associated with the HELP mechanism [33,41], where hydrogen increases dislocation motion thus enhancing localised plasticity. However, because hydrogen content was not directly quantified in this study, the association with HELP is hypothesised rather than directly confirmed.
At higher ccd, the fracture surfaces showed a greater proportion of quasi-cleavage features and decohesion cracks, as indicated by the orange and yellow arrows in Figure 13. These features suggest a more pronounced loss of fracture resistance and are consistent with a greater contribution from HEDE, where hydrogen reduces cohesive strength at crack tips, grain boundaries, or particle matric interfaces. The presence of both shallow dimples and quasi-cleavage regions in the charged specimens suggest that the fracture was not governed by a single mechanism but involved the interaction of localised plasticity and decohesion.
The observed fracture behaviour is consistent with the hydrogen embrittlement mechanism described by Djokic et al. [39] in which synergistic effects of HEDE and HELP differ depending on hydrogen concentration in the material. In the present study, hydrogen uptake was not quantified; nevertheless, the predominance of shallow dimples at lower ccd is consistent with a greater contribution from HELP, whereas quasi-cleavage and decohesion cracks at higher ccd suggests a stronger contribution from HEDE. Similar HELP and HEDE mechanisms have also been reported in cathodically hydrogen-charged aluminium alloys [33,41,74].
Overall, the fractographic observations and tensile results indicate a progressive transition from ductile to brittle behaviour with increasing ccd. Based on these results and the previous literature, this behaviour is interpreted as evidence of cathodic charging-induced hydrogen-assisted mechanical degradation in WAAM AA2319, which became more pronounced as the ccd increased.

3.3.1. Charpy Fractography

Figure 14 shows the fracture surfaces of the uncharged and 8 mA/cm2 hydrogen-charged Charpy specimens. The side view image (Figure 14a) highlights pronounced differences in crack propagation between the two conditions. The uncharged specimen exhibits a non-rigid, deformed fracture appearance indicating plastic deformation during failure whereas, the hydrogen-charged Charpy specimen shows a straight crack propagation path indicative of a brittle fracture. This behaviour is also evident in the fracture surface macrographs shown in Figure 14b.
The fracture surface of the uncharged specimen depicts a bowed profile indicating plastic deformation during fracture. In contrast, the hydrogen-charged specimen retained a geometrically regular square profile, indicating limited plastic deformation and fast crack propagation.
The SEM image of the uncharged specimen (Figure 14c) shows a high density of dimples associated with ductile fracture through microvoid nucleation, growth, and coalescence. A similar fracture morphology was observed for the uncharged tensile specimen, Figure 11b. The hydrogen-charged specimen (Figure 14d) displays a fracture surface with very shallow dimples and quasi-cleavage features, indicating a transition towards a more brittle fracture behaviour. This transition from ductile to a semi-brittle fracture accounts for the reduction in impact toughness following hydrogen charging. In ductile fracture, a substantial amount of energy is absorbed through plastic deformation, whereas in brittle fracture lower energy is absorbed due to limited plasticity.

3.3.2. Surface Corrosion

Corrosion products were observed on the fracture surfaces of some hydrogen-charged specimens, near the edges of the samples, as shown in Figure 15. EDS analysis of these regions revealed a high oxygen content, consistent with the formation of aluminium hydroxide [28]. Additionally, chlorine from the charging electrolyte was detected, which is a known contributor to pitting corrosion in aluminium alloys [76]. Figure 15 illustrates the EDS results at the highlighted spot.

4. Conclusions

The effects of cathodic hydrogen charging on the mechanical properties and fracture behaviour of WAAM-fabricated AA2319 aluminium alloy were investigated using uniaxial tensile testing, Charpy testing and fracture surface analysis. The key findings are as follows:
•
Cathodic hydrogen charging reduced the ductility and impact toughness of WAAM AA2319, indicating susceptibility to hydrogen-assisted mechanical property degradation under the charging conditions used in this study.
•
Increasing cathodic current density resulted in progressive reduction in elongation of up to approximately 45% relative to the uncharged condition. Charpy impact toughness also decreased by approximately 40% following cathodic charging. In contrast, the UTS and yield strength were marginally affected, indicating that cathodic charging primarily degraded ductility and fracture resistance without substantially altering the macroscopic strength of the alloy.
•
Fractography analysis showed a transition from ductile fracture characterised by microvoid coalescence in the uncharged specimens to a mixed ductile—brittle fracture in charged specimens. The charged specimens exhibited shallow dimples, quasi-cleavage features, and decohesion cracks, with these features becoming more pronounced at higher cathodic current densities.
•
The fracture morphology suggests that HELP was more dominant at lower cathodic current densities leading to reduced elongation. At higher ccd, the increased quasi-cleavage features and decohesion cracks suggests a greater fracture contribution from HEDE.
Overall, the results show that cathodic hydrogen charging degrades the ductility and fracture resistance of WAAM-fabricated AA2319 under both quasi-static and dynamic loading conditions. The combined mechanical and fractographic evidence indicates HELP mechanism at lower ccd and HEDE being dominant at higher ccd.

Author Contributions

Conceptualization, T.M.; methodology, T.M., A.G. and A.T.; investigation, T.M.; resources, A.G. and A.T.; writing—original draft preparation, T.M.; writing—review and editing, T.M., A.G. and A.T.; visualisation, T.M.; supervision, A.G. and A.T.; All authors have read and agreed to the published version of the manuscript.

Funding

This research received no external funding.

Data Availability Statement

The original contributions presented in this study are included in the article. Further enquiries can be directed to the corresponding author.

Conflicts of Interest

The authors declare no conflicts of interest.

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Metals 16 00647 g001
Figure 1. WAAM robotic cell in operation.
Figure 1. WAAM robotic cell in operation.
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Figure 2. (a) Final component; (b) schematic of samples as extracted from deposited material; (c) tensile specimen and dimensions (mm); (d) Charpy V-notch specimen and dimensions (mm).
Figure 2. (a) Final component; (b) schematic of samples as extracted from deposited material; (c) tensile specimen and dimensions (mm); (d) Charpy V-notch specimen and dimensions (mm).
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Metals 16 00647 g003
Figure 3. Electrochemical hydrogen charging setup.
Figure 3. Electrochemical hydrogen charging setup.
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Figure 4. (a) Porosity distribution in WAAM AA2319 under OM; (b) histogram of pore size distribution.
Figure 4. (a) Porosity distribution in WAAM AA2319 under OM; (b) histogram of pore size distribution.
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Metals 16 00647 g005
Figure 5. WAAM AA2319 microstructure (a) etched surface; (b) magnification of (a).
Figure 5. WAAM AA2319 microstructure (a) etched surface; (b) magnification of (a).
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Figure 6. (a) BSE image of polished sample; (b) EDS results of highlighted particles.
Figure 6. (a) BSE image of polished sample; (b) EDS results of highlighted particles.
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Figure 7. Hardness results showing: (a) microhardness of the polished sample; (b) surface macrohardness of the uncharged and hydrogen-charged specimen (8 mA/cm2).
Figure 7. Hardness results showing: (a) microhardness of the polished sample; (b) surface macrohardness of the uncharged and hydrogen-charged specimen (8 mA/cm2).
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Figure 8. Tensile testing results of WAAM AA2319 with and without hydrogen charging: (a) force displacement graph; (b) engineering stress–strain and total elongation.
Figure 8. Tensile testing results of WAAM AA2319 with and without hydrogen charging: (a) force displacement graph; (b) engineering stress–strain and total elongation.
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Metals 16 00647 g009
Figure 9. HEI results of different mechanical parameters (a) yield strength; (b) UTS; (c) elongation.
Figure 9. HEI results of different mechanical parameters (a) yield strength; (b) UTS; (c) elongation.
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Metals 16 00647 g010
Figure 10. Charpy impact results of uncharged and hydrogen-charged specimens.
Figure 10. Charpy impact results of uncharged and hydrogen-charged specimens.
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Metals 16 00647 g011
Figure 11. SEM images of fracture surface on uncharged specimen (a) ×200, overall fracture surface (b) ×1.10k, showing dimple fracture and inclusions.
Figure 11. SEM images of fracture surface on uncharged specimen (a) ×200, overall fracture surface (b) ×1.10k, showing dimple fracture and inclusions.
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Metals 16 00647 g012
Figure 12. Tensile fracture surfaces of (a) uncharged; (b) 2 mA/cm2; (c) 4 mA/cm2; (d) 6 mA/cm2; (e) 8 mA/cm2 specimens.
Figure 12. Tensile fracture surfaces of (a) uncharged; (b) 2 mA/cm2; (c) 4 mA/cm2; (d) 6 mA/cm2; (e) 8 mA/cm2 specimens.
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Metals 16 00647 g013
Figure 13. SEM fracture surface images of uncharged and cathodically charged tensile specimens: (a) uncharged; (b) 2 mA/cm2; (c) 4 mA/cm2; (d) 6 mA/cm2; (e,e1) 8 mA/cm2. White arrows indicate porosity; orange arrows show quasi-cleavage fracture and yellow arrows mark decohesion cracks. Orange dashed lines highlight sections of shallow MVC.
Figure 13. SEM fracture surface images of uncharged and cathodically charged tensile specimens: (a) uncharged; (b) 2 mA/cm2; (c) 4 mA/cm2; (d) 6 mA/cm2; (e,e1) 8 mA/cm2. White arrows indicate porosity; orange arrows show quasi-cleavage fracture and yellow arrows mark decohesion cracks. Orange dashed lines highlight sections of shallow MVC.
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Figure 14. Charpy specimens after impact tests (a) side view, (b) macro image of fracture surfaces, (c,d) fracture surface of uncharged and hydrogen-charged specimens.
Figure 14. Charpy specimens after impact tests (a) side view, (b) macro image of fracture surfaces, (c,d) fracture surface of uncharged and hydrogen-charged specimens.
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Metals 16 00647 g015
Figure 15. EDS spectrum of the corrosion products on 6 mA/cm2 specimen.
Figure 15. EDS spectrum of the corrosion products on 6 mA/cm2 specimen.
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Table 1. Chemical composition of ER2319 wire.
Table 1. Chemical composition of ER2319 wire.
ElementCuMnFeTiZrVSiAl
wt.%5.660.270.220.130.110.100.07Balance
Table 2. WAAM process parameters.
Table 2. WAAM process parameters.
ParameterValue
Average current72
Average voltage18
Average heat input (kJ/mm)0.14
Travel speed (mm/s)10
Wire feed speed (m/min) 6.5
Standoff distance (mm)15
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Mazarire, T.; Galloway, A.; Toumpis, A. Effects of Cathodic Hydrogen Charging on the Mechanical Properties and Fracture Behaviour of Wire Arc Additively Manufactured AA2319. Metals 2026, 16, 647. https://doi.org/10.3390/met16060647

AMA Style

Mazarire T, Galloway A, Toumpis A. Effects of Cathodic Hydrogen Charging on the Mechanical Properties and Fracture Behaviour of Wire Arc Additively Manufactured AA2319. Metals. 2026; 16(6):647. https://doi.org/10.3390/met16060647

Chicago/Turabian Style

Mazarire, Tinashe, Alexander Galloway, and Athanasios Toumpis. 2026. "Effects of Cathodic Hydrogen Charging on the Mechanical Properties and Fracture Behaviour of Wire Arc Additively Manufactured AA2319" Metals 16, no. 6: 647. https://doi.org/10.3390/met16060647

APA Style

Mazarire, T., Galloway, A., & Toumpis, A. (2026). Effects of Cathodic Hydrogen Charging on the Mechanical Properties and Fracture Behaviour of Wire Arc Additively Manufactured AA2319. Metals, 16(6), 647. https://doi.org/10.3390/met16060647

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