Influence of filler wires on hardness and coating dilution in high-speed laser-welded Al–Si-coated 22MnB5 steel

Abstract

22MnB5 steel blanks (Usibor^®1500) are widely used in automotive crash-resistant components because of their high strength. However, laser welding of Al–Si-coated 22MnB5 is challenging because coating dilution in the weld pool promotes soft δ-ferrite phase formation, reducing weld strength. An industrial solution is laser-ablation, while filler wire addition provides an alternative means to modify fusion-zone chemistry, with or without ablation. In this study, micro-structural development and mechanical performance of laser-welded Al-Si-coated 22MnB5 were investigated using ER4130 and ER420 filler wires. Energy-dispersive spectroscopy showed reduced bulk aluminum-(Al) content in welds with both filler wires relative to as-received welded hot-stamped welds, correlating with reduced δ-ferrite formation and increased martensite content. Higher chromium in ER420 further enhances hardenability, while carbon addition on wires further promoted martensite formation. These findings offer insights for Cr–C-based filler wire selection and advancing laser welding strategies for high-strength structural applications.

Introduction

The automotive industry increasingly relies on high-strength-to-weight ratio materials like advanced high-strength steels (AHSSs) for structural and safety-critical components.^1–4 AHSSs, such as Al–Si-coated 22MnB5 (also known as Usibor^®1500 and hereafter referred to as 22MnB5), are valued for their strength and ductility.^5–8 During hot stamping, steel is heated to 920 °C for 6 min and die-quenched for 12 s. Al–Si coating protects steel against oxidation and de-carburisation during austenitisation process, preserving carbon and properties.^9,10

Laser welding is often utilised to join 22MnB5 sheets due to stringent demands for high joint quality, which requires high energy efficiency and deep penetration, making lasers ideal for welding various thicknesses of sheets. ¹¹ Laser welding of 22MnB5 steel is commonly performed in the unquenched (as-received) condition prior to hot stamping; however, several studies have also reported laser welding in the quenched (hot stamped) condition, which avoids post-weld heat treatment but can introduce susceptibility to heat-affected zone (HAZ) softening. ¹² However, a major challenge is the Al–Si coating, which melts inside the weld pool, raising aluminium (Al) content in the fusion zone.^13,14 Aluminium stabilises ferrite, which promotes δ-ferrite formation, creating a dual-phase structure of ferrite in the martensite matrix.^14,15 Initially, it was demonstrated that dilution of Al–Si coating into the molten weld pool led to the formation of Fe–Al intermetallic compounds along the fusion zone boundary, as reported by Kim et al. ¹⁶ However, later Saha et al. ¹⁷ confirmed that dilution of Al–Si coating into the molten pool results in the formation of δ-ferrite along the fusion boundary (FB) due to the ingress of Al, which stabilises ferrite. Rapid solidification in laser welding limits coating diffusion, leaving inhomogeneous fusion zones with white ferrite bands in martensite.^16–18 These soft bands act as failure sites during post-processing.^17,18

A study by Khan et al. ¹⁵ showed that adding carbon in the form of colloidal graphite to the coated steel surface can help suppress the influence on the Al inside the weld pool, leading to an increase in the amount of martensitic micro-structures. To mitigate the adverse effects of Al–Si coating, filler wires have been used as a strategy to modify weld pool chemistry, thereby enhancing mechanical properties and reducing defect formation.^19,20 Use of carbon-steel filler wires or addition of austenite stabilising elements has been shown to reduce δ-ferrite formation in the fusion zone and enhance the overall joint strength.^19,21 Recent studies have claimed that filler wire welding can raise strength to ∼1550 MPa with ˃3% elongation.^20,21

Chromium, a strong solid-solution strengthener and hardenability element, further influences weld micro-structure and mechanical properties. ²² Similarly, carbon also enhances hardenability, martensite formation and dilutes aluminium in the weld, reducing δ-ferrite formation. ¹⁵ This study examines the effects of using different filler wires, specifically commercial ER4130 high-strength tool steel and ER420 high-chromium stainless steel, as well as the addition of carbon, on the micro-structure and mechanical properties of joints, while also achieving a good weld geometry using a welding speed of 8 m min⁻¹, with the filler wire. Results of this study assess the feasibility of using Cr-based filler wires along with the influence of added carbon during the welding of 22MnB5. ²³

Materials and experimental methods

Materials

In this study, 22MnB5 (1 mm thick) along with Al–Si coating was used as a base material, cut into 200 × 80 mm² specimens by shear cutting. Composition of base material is (0.24C–1.20Mn–0.21Cr–0.28Si–0.05Cu). Commercial ER4130 high-strength tool steel (0.30C–0.42Mn–0.95Cr–0.22Si–0.15Cu–0.19Mo) and high-chromium stainless steel (0.31C–0.33Mn–13.7Cr–0.55Si–0.29Ni–0.03Mo) ER420, 0.9 mm diameter, were used as filler wires. To raise the carbon content by ∼0.5 wt-%, graphite paste (isopropyl-based) was applied to filler wires, yielding ER4130 + C (0.80C–0.42Mn–0.95Cr–0.22Si–0.15Cu–0.19Mo) and ER420 + C (0.81C–0.33Mn–13.7Cr–0.55Si–0.29Ni–0.03Mo). Weight gain before and after carbon coating confirmed the addition of carbon.

Welding process and hot stamping

Laser welding was performed using a high-power fibre laser (IPG-YLS 8000) paired with a Fanuc robot. The laser wavelength was 1069.7 nm, with a focal length of 300 mm and a default spot size of 0.3 mm. A positive de-focus of 10.8 mm was used to achieve a 0.7 mm spot size, as the laser follows a Gaussian beam power distribution. Filler wire was fed using a Lincoln Electric power supply. Power used was 4.5 kW for welding with filler wire and 3.75 kW for autogenous welding. A higher heat input is required while welding with filler wire to provide additional energy to melt the welding wire. These specific power values are used after multiple experiment trials with power values to minimise surface defects, such as under-cut and reinforcement. Argon shielding gas (99.9%) was supplied at 40 L min⁻¹ with a cross-jet to remove the plume. No alcohol cleaning was applied, reflecting industrial practice.

Filler wire (fed at 2.75 m min⁻¹) was delivered at 30° through a copper guide positioned 15 mm ahead of the pool. Welding speed was 8 m min⁻¹; the laser beam was tilted to 8° as illustrated in Figure 1(a) and (b). Nominal heat input was calculated to be 28.1 J mm⁻¹ for the autogenous weld, and 33.7 J mm⁻¹ when filler wire is used. The sheet edges were shear cut, such that at the time of the butt joint, there was a ∼0.05 mm gap verified by a feeler gauge.

Figure 1.

(a–d) Schematic of laser and sheet setup. (e) Coating/base micro-structure in as-received, (f) hot-stamped 22MnB5 and (g) tensile specimen dimensions.

After welding, the samples were then hot stamped by heating to 920 °C for 6 min, and cooling to room temperature between two flat dies. Figure 1(e) and (f) compares the Al–Si coating in both its as-received and hot-stamped states. In the as-received condition, the coating (35–40 µm) comprises an Al–Si layer and a (Fe–Al) intermetallic coating layer, as illustrated in Figure 1(e). The as-received steel was ferrite-pearlite, transformed into martensite after hot stamping (Figure 1(f)).

Characterisation of the welds

Welded sheets were cut into 10 mm × 15 mm specimens by water jet, mounted in bakelite, ground to 1200 grit SiC and polished up to 1 µm DiaPro Nap B1 abrasive media. Specimens were etched with a nital solution (5% nitric acid (HNO₃), 95% methanol (CH₃OH)) for 4–5 s to reveal the micro-structure. A digital microscope (Keyence VHX-7000 series) and scanning electron microscope (SEM; Zeiss FE-Leo) were used for micro-structural and phase analysis of the hot-stamped welds. Energy-dispersive spectroscopy (EDS) was also performed to examine the elemental distribution in the fusion zone. The volume fraction of δ-ferrite was quantified through image analysis using Clemex Vision Lite software.

To evaluate the mechanical properties of the welded joints, the hardness of the weld metal, HAZ and base material was measured using a Clemex CMT (v.8.0.197) automated Vickers tester, with an indentation load of 200 or 50 gf, and indent spacing as per ASTM E384. The indentation was performed with two loads. As the load size decreases, the indent diagonal also reduces, leading to closer and narrower spacing between the loads. This improves the ability to capture hardness variations with better mapping compared to 200 gf. The tensile properties of the welded joints were assessed as per ASTM E8 standards using an MTS-810 machine at a displacement rate of 1 mm min⁻¹, and a digital image correlation (DIC) system was utilised to analyse the strain with a 7-pixel step and a 29-pixel subset size. Results were analysed using an average of five replicated tests, and the dimensions of the tensile specimens are presented in Figure 1(g).

Results and discussion

Weld geometry

Figure 2 shows cross-sectional weld geometries for all conditions, both as-received and hot stamped, at 8 m min⁻¹. Figure 2(a) to (e) shows the formation of Y-shaped welds without significant defects, such as lack of fusion or porosity. All welds remained within a ±10% of base material thickness in reinforcement and under-cut, meeting EN 10359:2023 standards. The average fusion zone width in the as-received condition ranged from 0.9 mm in autogenous to 1.1 mm with ER420 + C, while in the hot-stamped condition, it varied from 0.9 mm in as-received welded hot-stamped welds (ARWHS) to 1.4 mm in ER4130 + C, the widest. These differences indicate that both thermal cycles and filler wire chemistry influence laser fusion behaviour.

Figure 2.

Cross-sections of laser-welded 22MnB5: (a–e) as-received – autogenous, ER4130, ER4130 + C, ER420, ER420 + C; (f–j) hot stamped – ARWHS, ER4130, ER4130 + C, ER420, ER420 + C. ARWHS: as-received welded hot-stamped weld.

Following hot stamping, the excess weld reinforcement was flattened by quenching the die, resulting in a flattened final weld surface, ensuring dimensional accuracy in stamped components for proper assembly fit and preserving weld integrity.

Elemental distribution

Elemental composition of the fusion zone in the hot-stamped condition was examined by area-scan EDS at the top, middle and bottom regions (Figure 3). The average wt-% values are given in Table 1, with elemental distribution maps of Cr and Al in Figure 3. A trend is observed in the average Al wt-% distribution, with the ARWHS showing the highest average Al content ∼3.13 wt-%. This is attributed to the dilution of the Al–Si coating from the top and bottom in the absence of filler wires. This high level of average Al content inside the weld enhances δ-ferrite formation by broadening the body-centred cubic phase region and raising the A3 transformation temperature.^24,25 Filler-wire welds (ER4130, ER420 and their carbon-added variants) show a consistently lower Al content of ∼1.1–1.2 wt-%, compared with ARWHS welds. A higher laser power used during filler wire welding will cause more vapourisation of the coating, which may result in less Al content inside the weld. Also, the filler wire acts as a liquid-bridge transition, increasing molten-pool flow and significantly increasing molten-pool convection and expanding the molten volume, making a wider fusion zone but with lower Al content. The same trend in reduced Al content with a widening the fusion zone when using filler wire weld is observed by Xu et al. ²⁰ As the average Al content of the weld drops, δ-ferrite formation is suppressed and the martensite phase dominates. A similar observation by Khan et al. was reported in their study, where a carbon coating increased the heat absorptivity and heat input, resulting in a lower average Al content inside the weld pool compared to ARWHS. ¹⁵ To quantify further, the EDS maps in Figure 3 confirmed lower and more dispersed Al in filler-wire welds.

Figure 3.

EDS of fusion zones welds: (a) ARWHS, (b) ER4130, (c) ER4130 + C, (d) ER420, (e) ER420 + C, showing backscattered electron SEM (left) and Cr/Al maps (right). ARWHS: as-received welded hot-stamped weld.

Table 1.

Average elemental composition in wt-% norm.

Welding condition	Chromium (wt-%)	Iron (wt-%)	Nickel (wt-%)	Molybdenum (wt-%)	Aluminium (wt-%)	Silicon (wt-%)	Manganese (wt-%)
ARWHS	0.2	94.83	0.02	0.01	3.13	0.58	1.21
ER4130	0.2	97.43	0.07	0.06	1.11	0.36	0.82
ER4130 + C	0.2	97.31	0.07	0.08	1.18	0.36	0.85
ER420	1.8	95.25	0.11	0.04	1.22	0.36	0.88
ER420 + C	2.1	95.23	0.14	0.03	1.23	0.38	0.88

ARWHS: as-received welded hot-stamped weld.

The role of chromium is especially significant in welds made with ER420 and ER420 + C filler wires, where the average Cr content inside the fusion zone (FZ) is ∼1.9–2.1 wt-%. The EDS maps in Figure 3 show the Cr distribution across the weld zone, confirming a higher Cr wt-% than the base material. Chromium increases hardenability by delaying diffusional transformation kinetics from austenite to ferrite in Cr-alloyed steels and shifts the time-temperature-transformation curves to the right. ²² This promotes more martensite formation upon cooling. Cr does not interact directly with Al but suppresses its ferrite-promoting effect by enabling faster martensitic transformation kinetics. This behaviour is consistent with previous findings. ²²

The carbon addition through graphite paste, as in the ER4130 + C and ER420 + C joints, further enhances hardenability. As shown by Luo et al., ²⁶ the addition of carbon in Mn–B steels favours martensite formation during rapid cooling in 38MnB5Nb compared to 22MnB5. The addition of carbon suppresses the martensite start temperature (M_s) and promotes martensite formation during rapid solidification.

Chromium and carbon both delay diffusional transformations, increasing hardenability, suppressing δ-ferrite and facilitating a more martensite structure.^26,27 In ER420 wire-fed joints, chromium impedes austenite-to-ferrite diffusion and, together with higher carbon, promotes martensite formation and further limits δ-ferrite.^28,29

Micro-structures

SEM images in Figure 4(a) to (e) show the micro-structural transformation at the FB among the various welding conditions. The main phases identified were martensite (M) and δ-ferrite (F), with their distribution along the weld primarily influenced by the filler wire composition, average Al wt-% and carbon addition. Image analysis was used to determine the volume fraction of δ-ferrite phase present in the weld affected by composition change, and the results are shown in Figure 4(f) to (j). When no filler wire is used, it is found that the FZ had a ferrite volume of about 65(±5.30)%, as shown by the blue phase in Figure 4(f). When analysing the welds made with filler wire and carbon addition, it was observed that the ferrite phase concentration consistently decreased as the Cr wt-% increased, along with the carbon content. The ferrite volume fraction was reduced to about 2.6(±1.9)% when welded with ER420 with carbon addition. A similar trend of decreasing ferrite fraction was observed by Khan et al., where a 1.5 mm thin sheet in ARWHS condition had 40% ferrite fraction when there was approximately 1.06 wt-% Al inside the FZ, and 2% ferrite fraction when carbon was added to the FZ through colloidal graphene. ¹⁵

Figure 4.

SEM images (a–e) of fusion boundary (FB) and base material under different welds: (a) ARWHS, (b) ER4130, (c) ER4130 + C, (d) ER420, (e) ER420 + C (M: martensite, F: δ-ferrite) Clemex phase maps (f–j) show ferrite (blue). ARWHS: as-received welded hot-stamped weld.

Across the FB, ARWHS welds and those joined with ER4130 had a significant δ-ferrite region, while ER420-based welds contained a higher fraction of martensite. The slight decrease in δ-ferrite content confirms that the selection of filler wire and the element composition of the weld induce phase transformation. The transition across the FB shows that ER420 with carbon has the most considerable martensite phase fraction.

Micro-hardness

The hardness variation across the welded joints after the hot-stamping condition is shown in Figure 5. For a broader view of hardness distribution, the weld cross-sections were subjected to a systematic hardness measurement across the fusion zone and base material using a 200 gf indentation load. The resulting hardness graph, shown in Figure 5(a), visualises the average fusion zone hardness distribution for each welding condition. The ARWHS weld shows the most pronounced softening in the fusion zone with an average hardness of 370 HV, due to higher ∼3.13 wt-% Al average content in the weld and a related high number of δ-ferrite-rich softened regions. Compared to other welds with filler material, they show improved fusion zone hardness, indicating the influence of alloying elements in reducing δ-ferrite formation.

Figure 5.

(a) Average micro-hardness across weld zones in the hot-stamped joints. (b–e) Hardness maps at 50 gf for ER4130, ER4130 + C, ER420 and ER420 + C.

Among the filler wires, ER4130 welds exhibited the lowest average fusion zone hardness (450 HV) due to high fractions of δ-ferrite in the weld metal. With the addition of carbon to ER4130 wire, due to its hardenability, it promotes the formation of martensite and suppresses the δ-ferrite phase fraction, increasing the average fusion zone hardness to 470 HV. Similarly, the ER420 filler wire resulted in a higher average fusion zone hardness (480 HV), due to the Cr atoms acting as both a solid-solution strengthener and increasing hardenability. This promotes martensite formation while reducing the δ-ferrite fraction. The ER420 with a carbon addition shows the highest average fusion zone hardness (550 HV), indicating a predominantly martensite micro-structure with minimal δ-ferrite fraction. Such a high hardness is due to the formation of high-carbon martensite, resulting from increased carbon content and the combined effect of Cr and C on hardenability.

To study the micro-structural hardness variation across the weld, several micro-hardness indentations using 50 gf were performed across the FZ, including measurements on visually identified phases such as δ-ferrite and martensite. These localised indentations helped differentiate the hardness response of each phase, with selected results shown in Figure 5(b) to (e), highlighting the contrast between softer ferrite regions and harder martensite areas. When ER420 filler wire is used, carbon addition results in a more martensitic phase in the fusion zone with increased hardness. These features, along with the effects of over-hardened martensite or grain boundary carbide segregation, can influence local toughness even as overall hardness rises. As illustrated in Figure 5(e), hardness mapping with a lower indent load reflects these patterns, showing a hardness distribution in butt welds with higher carbon content. The greatest retention of hardness occurs in the ER420 wire-fed weld containing carbon, indicating enhanced hardenability of the fusion zone. However, substantial variation in hardness remains a factor that should be minimised, as it may contribute to crack initiation in the fusion area and impact quasi-static tensile or dynamic Charpy testing outcomes.

Tensile strength

A comparison with the published results further highlights the influence of coating removal on weld performance. Lin et al. ³⁰ showed that de-coated 22MnB5 welds contain only ∼0.5 wt-% Al in the fusion zone, resulting in fully martensitic micro-structures after hot stamping and significantly improving mechanical behaviour compared with the coated welds. They observed that welds on sheets with the coating removed achieved an ultimate tensile strength (UTS) of ∼1428 MPa and ∼3% elongation in the hot-stamped state. However, coated welds exhibited a pronounced decrease in strength (∼1124 MPa) and ductility due to δ-ferrite formation. These results demonstrate the mechanical benefit of reducing Al dilution – whether by coating removal or by altering weld pool chemistry through filler wire-based approaches used in this study. Figure 6(a) to (c) summarises the tensile properties of the hot-stamped specimens of all welding conditions, presenting UTS and total elongation. Figures 6(d) to (i) and 7(f) to (k) further illustrate local elongation at failure and fracture location, emphasising how filler wire selection and carbon addition affect mechanical performance.

Figure 6.

UTS (MPa) and elongation (%) for each weld (a–c), with DIC strain maps (d–f). UTS: ultimate tensile strength; DIC: digital image correlation.

Figure 7.

(a–e) Tensile specimens after elongation to 70% peak load: ARWHS, ER4130, ER4130 + C, ER420, ER420 + C. (f–k) Fractured specimens: base material, ARWHS, ER4130, ER4130 + C, ER420, ER420 + C. ARWHS: as-received welded hot-stamped weld.

The base material, 22MnB5, achieves a UTS of 1512 MPa and a total elongation of 4%, demonstrating a fully martensitic structure following hot stamping and setting a benchmark for comparison. The ARWHS welds exhibited the lowest UTS of 612 MPa due to significant δ-ferrite formation caused by Al retention from the Al–Si coating. As discussed earlier, the average Al content in ARWHS welds is ∼3.13 wt-%, leading to a high δ-ferrite fraction of ∼65(±5.30%), which contributed to the reduction in UTS. In contrast, welding with ER4130 showed an increase in UTS to 1201 MPa due to reduced average Al content in the weld and increased martensite formation relative to the δ-ferrite phase. With the addition of carbon (ER4130 + C), the tensile strength of joints produced using ER4130 showed a slight increase in nominal fracture stress to 1278 MPa, as carbon increases the hardenability and further suppresses δ-ferrite formation. When welded using the ER420, containing high chromium content (13.7 wt-% in filler wire), the highest UTS (1401 MPa) was achieved due to the higher martensite content and reduced δ-ferrite. However, ER420 + C welds had a slightly decreased tensile strength (1357 MPa). Although carbon addition increases hardness due to high-carbon martensite, it can also introduce local embrittlement due to carbide formation or Cr–C interactions, as seen by the hardness maps, which show regions greater than 550 HV.

The location of the fracture in the specimen consistently occurred at the weld interface, regardless of the welding conditions or type of filler wire used. ARWHS welds fractured within the fusion zone, resulting in the lowest elongation, which can be explained by high quantities of δ-ferrite due to a high amount of average Al content (3.13 wt-%) in the fusion area. This δ-ferrite type is usually coarse, irregular and poorly bonded to nearby phases. Instead of improving ductility, it serves as a site for crack initiation, leading to cleavage-type fractures and inadequate strain accommodation.^17,26 Additionally, the presence of a small under-cut in the weld geometry may have further contributed to pre-mature failure. Conversely, welds using ER4130 exhibited greater strength and improved ductility than ARWHS welds. This is due to a reduction of the average Al content (1.1–1.2 wt-%) inside the FZ, which suppresses δ-ferrite further and facilitates more martensite structure.

Adding carbon to ER4130 increases elongation slightly to 1.33%, compared to 1.00% in the ER4130 weld. This enhancement results from a small amount of carbon, suppressing further δ-ferrite. Consequently, a slightly more martensite matrix enables a more stable strain distribution and delays δ-ferrite initiation, leading to a slight increase in elongation, even with an overall rise in hardness and strength. On the other hand, when the same amount of carbon was added to the ER420, the mechanical response was different. The ER420 alone led to the highest UTS of 1401 MPa and a total elongation of 1.98% due to the optimised alloy composition. The high chromium content in ER420 increases hardenability and solid solution strengthening, which produces a slightly more martensite structure when welding with ER4130. The absence of excessive carbon limits high-carbon martensite, carbide precipitation and local embrittlement. Reduction of δ-ferrite fraction in ER420 welds, along with alloying element properties of chromium, contributes to strain distribution during tensile loading and allows the weld to preserve ductility while retaining high strength as compared to other welds. In contrast, ER420 + C exhibits a decrease in UTS to 1357 MPa and reduced elongation to 1.73%. The existing levels of chromium and carbon in ER420 wire, when further supplemented by additional carbon, can also lead to carbide formation, increased high-carbon martensite hardness and localised embrittlement. These factors, combined with the presence of under-cut, can contribute to micro-structural instability near the FB, affecting the weld's strain tolerance even though the martensite structure in the fusion zone remains higher than in other welds.

Figure 6(d) to (i) shows the DIC strain maps recorded just before fracture, showing the localised strain distribution across each sample under tensile load. The base material presents a characteristic V-shaped strain profile with a maximum local engineering strain of 17.5%, indicating uniform plastic deformation. This behaviour results from the martensitic micro-structure formed during hot stamping, offering a balance between strength and ductility, and is not affected by residual stresses or weld-induced heterogeneities.

In contrast, the ARWHS shows a much lower peak local strain of 1.25%, mainly along the fusion zone. The early strain localisation and pre-mature failure are likely caused by the heterogeneous micro-structure, where the high Al content from the Al–Si coating dilution encourages δ-ferrite formation. This phase compromises local mechanical strength and disrupts uniform strain distribution. Additionally, the presence of an under-cut along the weld FB further reduces the weld's strain tolerance and contributes to pre-mature failure. Welds made with ER4130 and ER4130 + C display a moderate local strain of ∼2.19%. These weld joints show a partially transformed micro-structure with reduced δ-ferrite fraction. Although these structures are harder than ARWHS, they still limit full plastic deformation. Adding carbon slightly improves phase uniformity, hardenability, and slightly enhances strain distribution due to a further decrease in the δ-ferrite fraction inside the weld. But still, the δ-ferrite and the presence of under-cut made the weld susceptible to failure at the interface.

Using ER420 results in a maximum local strain of about 2.81%, indicating further reduction of δ-ferrite fraction and an increase in martensite phase inside the FZ. This micro-structure enhances localised deformation and delays fracture compared to other welds. However, adding carbon to ER420 keeps the peak strain at 2.81%, suggesting that excess carbon, while promoting martensite formation, may cause grain boundary embrittlement through carbide precipitation and carbon segregation. The presence of under-cut, residual δ-ferrite and embrittlement still causes the fracture at the fusion zone. Consequently, the bulk strain capacity of the weld metal is not fully realised, and overall ductility decreases despite more martensitic structure and less δ-ferrite.

The samples with different filler wire compositions were pulled apart using approximately 70% of the peak load of the specimen, as shown in Figure 6(c), to evaluate the initiation of failure in the specimen before the failure of the sample occurred, as shown in Figure 7(a) to (e). At 70% peak load, the ARWHS sample (Figure 7(a)) does not exhibit any indication of necking or distributed plastic deformation along the gauge length. This pre-mature damage correlates with the low local strain (∼1.25%) as seen in Figure 6(e), showing limited strain accommodation. The brittle-like behaviour is due to a high fraction of δ-ferrite in the fusion zone. Additionally, the under-cut, along with the weld interface, likely serves as a stress concentrator, showing early crack initiation under loading. The ER4130 weld shows localised fracture due to heterogeneous phase distribution, alongside the influence of under-cut (Figure 7(b)). As the carbon percentage increases in ER4130, Figure 7(c) shows no necking and much plastic deformation. The ER420-welded specimens remain largely intact, with no major necking occurring in the fusion zone, as shown in Figure 7(d). This behaviour is consistent with the local strain (∼2.81%) seen in the DIC figure and reduced δ-ferrite content, suggesting that the weld metal is harder and slightly more ductile but capable of withstanding higher stress without pre-mature failure. However, as the carbon percentage increases in the ER420, the weld material shows improved strength and hardness, though it also exhibits reduced ductility and increased brittleness, as illustrated in Figure 6(c). ER420 + C compared to ER420 has a lower average total elongation. Although ER420 + C has higher hardness, the ER420 weld without carbon demonstrates greater toughness due to a higher combination of strength and ductility compared to the ER420 + C weld. This shows a balance compromised can be achieved with the increased carbon-content filler wire. Figure 7(f) to (k) illustrates these points with optical images of the fractured specimens, where fracture sites match the regions of highest strain localisation observed in the DIC maps.

Conclusions

This study examines the effect of Cr-based filler wire composition and carbon addition on the elemental distribution, micro-structural transformation, hardness and mechanical properties of laser-welded 22MnB5. The following conclusions are drawn:

Welding at 8 m min⁻¹ with optimised settings produced welds within ±10% tolerance, meeting the EN 10 359:2023 industry standard.

The addition of filler wires containing chromium decreased δ-ferrite formation and enhanced martensite transformation across the weld pool. The volume fraction of δ-ferrite decreased while the volume fraction of martensite increased with an increase in Cr% in filler wire from ER4130 (0.9%) to ER420 (13.7%), and the addition of C via carbon paste.

The average weld micro-hardness increased with an increase in Cr% in filler wire from ER4130 (0.9%) to ER420 (13.7%) and the addition of C via carbon paste. The highest average micro-hardness of 550 HV in the fusion zone was achieved with ER420 wire coated with carbon paste.

The use of carbon paste may have resulted in the formation of carbide regions and areas of high-hardness martensite, sometimes causing local hardness measurements above 600 HV.

An increase in martensitic content within the micro-structure is associated with higher hardness and tensile strength. The base material demonstrates a UTS of 1512 MPa, while ER420 reaches a UTS of 1401 MPa, which is approximately 93% of the base material's strength. Adding carbon to ER420 slightly lowers the UTS to 1357 MPa (about 90%) due to increased weld embrittlement.

Fractography analysis indicated that failure continued to occur at the FB despite changes in weld hardness. This was attributed to the presence of δ-ferrite at the FB, and the resulting failure mode did not satisfy the requirements of EN 10359:2023.

The results indicate that controlling fusion zone mechanical properties through filler wire chemistry is complex, as none of the approaches tested have yet met industry requirements for tensile performance, as per EN 10359:2023, which requires fracture to occur in the base metal (or as a ductile weld failure meeting minimum strength). The addition of chromium and carbon has been observed to increase hardness and strength in hot-stamped laser-welded specimens by reducing δ-ferrite formation; however, excessive carbon may result in embrittlement at grain boundaries, introducing new weaknesses in the weld. These results indicate the importance of determining the minimum and maximum effective ranges for each element incorporated into the wire. Additionally, completely eliminating the ablation process may not be practical, and it may be necessary to use both methods together to obtain the required properties.

Moreover, since hardness measurements and quasi-static tensile tests do not comprehensively reflect the dynamic crashworthiness or ductility required for safety-critical automotive components fabricated with hot-stamped laser welds, future work will incorporate Charpy testing following the methodology established by ArcelorMittal Global R&D.^31,32