Optimal cerium microalloying enhances SASS/Q235 weld corrosion and antibacterial performance

Summary

Super austenitic stainless steels (SASS) face challenges like galvanic corrosion and antibacterial performance when welded to carbon steel (Q235) in marine environments. This study demonstrates that adding 1.0 wt% cerium (Ce) to SASS refines the heat-affected zone (HAZ) grain structure (from 7 μm to 2 μm), suppresses detrimental σ-phase precipitation, and forms a dense oxide film. Electrochemical analyses confirmed this optimized composition increases charge-transfer resistance (to 2.2 × 10³ Ω cm²) and reduces passivation current density (to 0.12 μA/cm²), significantly enhancing corrosion resistance. Additionally, 1.0 wt% Ce disrupts sulfate-reducing bacteria (SRB) membranes, reducing survival to <1%. However, excess Ce (≥1.5 wt%) forms coarse CeO₂ particles, accelerating corrosion via porous films and micro-galvanic coupling. These findings provide a practical strategy for designing corrosion-resistant, antimicrobial welded joints in marine infrastructure.

Graphical abstract

Highlights

• •Ce element has a positive effect on the increase material’s surface potential
• •Adding 1.0 wt% La minimizes galvanic corrosion thus making it the optimal content
• •Ce increases the corrosion resistance of alloys to sulfate-reducing bacteria

Engineering; Materials science

Introduction

High concentrations of Cl⁻ and the activity of marine microorganisms make conventional austenitic stainless steels unsuitable for long-term, efficient service in harsh marine environments.^1,2,3 To address this, a series of super austenitic stainless steels (SASS) have been developed.^4,5,6,7 SASS, with its higher alloy content—particularly chromium (Cr) and molybdenum (Mo)—raises the pitting resistance equivalent number (PREN = Cr + 3.3Mo + 16N) to ≥45. This enhancement significantly improves corrosion resistance through the synergistic strengthening of Mo and nitrogen (N), making it an ideal material for offshore engineering, seawater desalination plants, and other corrosive environments.^8,9,10,11

However, SASS faces two challenges in the field of heterogeneous welding. First, the high alloy content leads to solidification segregation. Due to the elevated levels of alloy elements, SASS undergoes solute redistribution during solidification, forming liquid-phase segregation at the solidification interface, resulting in dendritic structures.^12,13,14 Studies have shown that high welding heat input causes grain coarsening in the heat-affected zone (HAZ).^15,16,17,18 Additionally, welding thermal cycles can cause the precipitation of metallic compounds along the γ/δ phase boundary, which severely affects the properties of the HAZ.^19,20,21 Second, when welded with conventional low-alloy steels, element diffusion at the Fe-Cr-Ni/Mo gradient interface leads to the formation of a brittle transition layer mainly composed of (Cr, Fe)₂₃C₆, which can cause intergranular fractures under seawater hydrothermal shock.^22,23,24 Research has shown that the high Cr content in SASS forms a concentration gradient with the Cr-poor region of low-carbon steel, promoting Cr migration to the interface. This, coupled with the increased pitting corrosion rate in chloride environments, results in HAZ sensitization in SASS.^25,26 Rare earth microalloying, particularly the addition of cerium (Ce), has been shown to effectively stabilize the austenitic matrix and suppress the precipitation of detrimental secondary phases, thereby improving both microstructural stability and corrosion resistance in SASS.^27,28,29 Cao et al. found that in high heat input welding, Ce refines the austenitic grains in the coarse-grain heat-affected zone (CGHAZ).³⁰ Barrie Mintz and others discovered that Ce combines with impurities like sulfur and phosphorus to form CeS or CeP, reducing the segregation of harmful elements (e.g., S, P) at grain boundaries and lowering the risk of intergranular embrittlement.³¹ Current research has primarily focused on optimizing Inconel-type welding materials for dissimilar steels. However, systematic understanding of rare earth microalloying’s effects on the microstructure and properties of dissimilar welded joints, particularly in the SASS/Q235 system, remains underdeveloped.

This study examines the dissimilar welding of stainless steel with varying cerium content and Q235 steel, with a focus on the microstructural evolution and corrosion resistance in the HAZ on the SASS side of the dissimilar weld. The addition of cerium and its interaction with the welding process significantly enhanced both corrosion resistance and antibacterial properties. The study highlights the role of Ce in the microstructural evolution, demonstrating that the optimal corrosion resistance and resistance to sulfate-reducing bacteria (SRB) occur at a cerium concentration of 1.0 wt%. This finding offers a scalable manufacturing strategy for producing corrosion-resistant welded materials for marine engineering applications.

Results

Microstructure analysis

Figure 1 illustrates the microstructural evolution characteristics of the fusion zone near the base metal side in the welded joint. The HAZ of the control sample without Ce addition (Figures 1.a₁–a₃) exhibits a typical epitaxial growth of coarse columnar crystals, with the dendrite arm spacing (DAS) calculated using Formula 7 as 7 μm ± 0.5 μm (Figure 1.a₃). With the addition of 0.5 wt% Ce (Figures 1.b₁–b₃), the dendritic crystals were refined, and the DAS decreased to 5 μm ± 0.5 μm (Figure 1.b₃), consistent with Zhang et al.,³² who observed similar refinement of the solidification structure by rare earth elements.³² When the Ce content was increased to 1.0 wt% (Figures 1.c₁–c₃), a columnar-to-equiaxed transition (CET) occurred in the HAZ, and discontinuous grain boundaries (black long lines) were observed. The average grain size was calculated as 10 μm ± 0.5 μm using Formula 8 (Figure 1.c₃). With 1.5 wt% Ce (Figures 1.d₁–d₃), the average grain size was further refined to 5 μm ± 0.5 μm (Figure 1.d₃). At 2.0 wt% Ce (Figures 1.e₁–e₃), the average grain size decreased to 2 μm ± 0.5 μm (Figure 1.e₃). Morphological analysis of the fusion line (left side of Figure 1) shows that the grain size of the weld metal also exhibited a trend of gradual refinement with increasing Ce content. This phenomenon can be attributed to changes in thermodynamic parameters during weld pool solidification. Due to extremely high cooling rates, steep temperature gradients form at the weld pool edges (on the substrate material side), and the solidification rate (R) is controlled by the direction of heat conduction, leading to an increase in the G/R ratio,^33,34 which results in epitaxial growth of columnar crystals. The introduction of cerium disrupts this solidification pattern through two mechanisms: (1) Ce-O-S composite nanoparticles act as heterogeneous nucleation sites, reducing the critical nucleation energy and promoting random nucleation of equiaxed crystals, and (2) the reduction in liquidus temperature caused by cerium significantly increases compositional undercooling at the solidification front.

Figure 1.

Microstructure of HAZ at different cerium contents: 0Ce (A₁, A₂, A₃), 0.5Ce (B₁, B₂, B₃), 1.0Ce (C₁, C₂, C₃), 1.5Ce (D₁, D₂, D₃), 2.0Ce (E₁, E₂, E₃)

Analysis of precipitated phases

Figure 2 demonstrates the microstructure and elemental distribution characteristics of the HAZ on the SASS side in dissimilar welded joints between SASS and Q235 steels with different Ce contents (0, 0.5, 1.0, 1.5, 2.0 wt%). The microstructure of the HAZ gradually refines with increasing Ce content (a₁-e₁), which is consistent with the previously discussed OM images. This indicates that Ce significantly inhibits grain coarsening caused by welding thermal cycling. Previous studies have shown that SASS is prone to precipitating σ-phase (Fe-Cr-Mo type intermetallic compounds) in the HAZ, and the aggregation of this brittle second phase leads to the formation of Cr-poor zones at grain boundaries, significantly deteriorating the material’s corrosion resistance.³⁵ However, the average size of the second phase in the HAZ increases with increasing Ce content (0 → 2.0 wt%) as shown by high-magnification SEM images (a₂-e₂) and corresponding energy dispersive spectrometry (EDS) facet distribution analyses. Notably, the EDS results indicate that Ce is enriched in the second-phase region, and the co-localization of elemental O suggests that Ce may form Ce-O compounds (e.g., CeO₂), which preferentially occupy grain boundaries. This inhibits σ-phase nucleation and reduces the total amount of deleterious precipitated phases. The diffuse distribution of Ce-O phases reduces the local electrochemical activity difference, thereby improving the pitting and intergranular corrosion resistance of the HAZ. To further clarify the composition of the second phase, X-ray diffraction (XRD) phase analysis and transmission electron microscopy (TEM) characterization were carried out.

Figure 2.

SEM and EDS of HAZ at different cerium contents: 0Ce (A₁, A₂, A₃), 0.5Ce (B₁, B₂, B₃), 1.0Ce (C₁, C₂, C₃), 1.5Ce (D₁, D₂, D₃), 2.0Ce (E₁, E₂, E₃)

The XRD spectra in Figure 3 show that the characteristic diffraction peaks of austenite (γ-phase) ((111), (200), and (220)) dominate in the HAZ of all the welded samples, indicating that the matrix structure in this region is predominantly austenitic. This is in agreement with the properties of super austenitic stainless steels, which inhibit martensitic phase transformation due to their high Ni and N content during the weld thermal cycle.³⁶ Notably, no significant characteristic CeO₂ diffraction peaks were detected at Ce additions ≤1.5 wt% (standard PDF #34–0394, main peaks at 2θ = 28.5°, 33.1°, and 47.5°). However, a broadening peak of weaker intensity was observed around 2θ = 47.5° when the Ce content was increased to 2.0 wt%. According to Scherrer’s formula (D = Kλ/βcosθ),³⁷ when the CeO₂ particle size is smaller than 5 nm, the diffraction peaks become significantly broadened due to the weakening of the coherence of reflections from the crystal surface. Consequently, their peak intensities may fall below the instrumental detection limit (<1 wt%). The strong diffraction peaks of the γ-Fe austenite (111) facet (2θ ≈ 43.5°) and the peaks of CeO₂ (111) facet (2θ = 28.5°), though not directly overlapping, may cause the weak peaks of low-content CeO₂ to be masked by the neighboring high-intensity austenite peaks (e.g., the (200) peak, 2θ ≈ 50.8°).³⁸ These results suggest that Ce may be diffusely distributed in the austenite matrix as ultrafine nanoparticles (<3 nm) or amorphous oxides at Ce contents ≤1.5 wt%. Such highly dispersed Ce-O phases can inhibit the precipitation of deleterious phases through the grain-boundary pinning effect, while also preventing the elevated local corrosion susceptibility caused by the aggregation of coarse second-phase particles.

Figure 3.

XRD of HAZ at different cerium contents

Figure 4 presents the TEM analysis results of the HAZ on the parent material side. Elemental surface distribution analysis using EDS clearly reveals the distribution and relative content of each element in the material. Research indicates that the σ phase has a tetragonal crystal structure, with experimentally measured unit cell parameters of a = 8.76 Å, c = 4.57 Å.³⁹ Due to differences in lattice geometry and symmetry across various crystal systems, the formulas for calculating interplanar spacings differ.

Figure 4.

TEM and EDS of HAZ at different cerium contents: 0Ce (a₁, a₂, a₃), 0.5Ce (b₁, b₂, b₃), 1.0Ce (c₁, c₂, c₃), 1.5Ce (d₁, d₂, d₃), 2.0Ce (e₁, e₂, e₃)

For the cubic crystal system, the interplanar spacing is calculated as:

$$d_{hkl}=\frac{a}{\sqrt{h^2+k^2{+l}^2}}$$ (Equation 1)

For the tetragonal crystal system, the interplanar spacing is:

$$d_{hkl}=\frac{1}{\sqrt{\frac{h^2+k^2}{a^2}+\frac{l^2}{c^2}}}$$ (Equation 2)

where h, k, and l are the Miller indices, and a is the lattice constant.

According to the EDP analysis in a₂, the interplanar spacing of the precipitated phase is d ≈ 0.371 nm, which is basically consistent with the value calculated using Formula 2 (d₁₁₁ of the σ phase ≈0.368 nm). Crystallographic analysis of CeO₂ indicates that this compound adopts a cubic fluoride structure, with a lattice constant a ≈ 0.541 nm, and the main exposed crystal planes are (111), (110), and (100).^40,41 The EDP data in Figures b and c show d ≈ 0.312 nm, which is consistent with the theoretical interplanar spacing of the (111) crystal plane calculated using Formula 1. The measured d ≈ 0.541 nm in Figures d and e corresponds to the characteristic spacing of the (100) crystal plane. Notably, the EDS mass ratio of the Ce-O system aligns with the stoichiometric ratio of CeO₂, while the Fe-Cr-Ni ratio corresponds to the stoichiometric ratio of the austenitic matrix. TEM-EDS analysis confirmed that the control sample without Ce addition primarily precipitated the σ phase, while Ce addition resulted in the formation of CeO₂ compounds in the system.

Analysis of electrochemical behavior

Figure 5 shows the comparison of electrochemical tests before and after seawater immersion (a-d before immersion, e-h after immersion) in the fusion zone (predominantly HAZ) close to the side of the parent material. The open-circuit potential (OCP) curves (a, e) show that the HAZ of the alloy containing 1.0% Ce shows the highest stabilization potential. It was found that Ce is preferentially enriched in the HAZ. This enrichment promotes a synergistic film-forming effect of the dense Cr₂O₃ passivation film and the Ce-O composite oxide.⁴² This structural feature greatly improves the chemical stability of the passivate film, enabling it to form an effective protective barrier in the early stages of corrosion.

Figure 5.

Electrochemical tests of the heat-affected zone before and after seawater immersion

(A–D) The OCP curve, Nyquist plot, Bode plot, and Tafel curve before immersion.

(E–H) The OCP curve, Nyquist plot, Bode plot, and Tafel curve after immersion.

(I) Comparison of corrosion potential before and after immersion.

(J) Comparison of self-corrosion current density before and after immersion.

(K) Corrosion mechanism.

Electrochemical impedance spectroscopy further verifies the effect of Ce on the corrosion resistance of the HAZ: the Nyquist plot (b) before immersion shows that the 1.0% Ce sample has the largest capacitive arc diameter, corresponding to a charge transfer resistance (Rct) of 2.2 × 10³ Ω cm², a 3-fold increase compared to the Ce-free sample. In the corresponding Bode plot (c), the 1.0% Ce sample exhibits a broad phase angle plateau of 70° in the mid-frequency region (10–100 Hz), while the impedance modulus reaches the order of 10³ in the low-frequency region (0.1 Hz), together indicating that an intact passivation film with high impedance characteristics is formed on the surface of the HAZ. After 12 days of seawater immersion, the HAZ maintains optimal protection: the 1.0% Ce sample retains the highest impedance value (8.5 × 10² Ω cm²) in the Nyquist plot (f), the phase angle in the mid-frequency region of the Bode plot (g) remains above 55°, and the low-frequency impedance modulus is two orders of magnitude higher than that of the blank sample. Microscopic mechanism studies show that Ce³⁺ effectively blocks the diffusion channel of Cl⁻ along the grain boundary by occupying the oxygen vacancies in the Cr₂O₃ lattice. Although prolonged immersion caused localized pitting of the passivation film, the corrosion rate of the 1.0% cerium sample in the HAZ was still lower than that of the 0% cerium sample because cerium enhances the repair capacity of the passivation film.⁴³ The diameter of the capacitive arc in the high-frequency region is positively correlated with the charge-transfer resistance (Rct), and the inductive loop in the low-frequency region may arise from the adsorption of localized corrosion products or residual inductive effects of the test system. It is inaccurate to equate the bilayer at the electrode interface (passivation film/electrolyte interface) to a pure capacitance. Therefore, a constant phase element (CPE) was used to characterize the capacitive response of the bilayer. The equivalent fitted impedance expression is shown in the following text:

$$Z_{CPE}=\frac{1}{{Q\left(j\omega \right)}^n}$$ (Equation 3)

$$Z_0=R_0+j\omega L$$ (Equation 4)

where Q denotes the amplitude parameter of the CPE, $j$ is the imaginary unit ($j$² = −1), $\omega$ is the angular frequency (ω = 2πf, where f is the frequency), and R₀ and L represent the equivalent resistance and equivalent inductance of the inductive impedance generation process.

By comparing the polarization curves before (d) and after (h) immersion, it is evident that passivation occurs in the anodic region, indicating that an oxide film forms on the material’s surface, hindering corrosion. Further analysis of the passivation current density (ipass) at different Ce contents revealed that the sample with 1.0 wt% Ce exhibited the lowest ipass value, indicating the densest oxide film and the best corrosion resistance. The corrosion potential (E_corr) and self-corrosion current density (I_corr) were calculated using the Tafel extrapolation method, as shown in (i and j). With the increase in Ce content, E_c was first positively shifted and then negatively shifted, peaking at 1.0 wt% with the highest thermodynamic stability. I_c was first reduced, then increased with the increase of Ce content, and finally decreased to the minimum value at 1.0 wt%, reflecting the slowest kinetic corrosion rate. The combination of high E_c (thermodynamically stable) and low I_c (kinetically slow) indicates that the material has both thermodynamic and kinetic corrosion resistance.⁴⁴ During electrochemical testing, the exposed metal substrate (anode) and the surrounding passivation film-covered area (cathode) form a microelectrode pair, initiating localized electrochemical corrosion. Oxidative dissolution of the metal occurs, releasing electrons, which are transferred through the substrate to the cathode region, driving the reduction reaction. Anodic dissolution leads to Cl⁻ enrichment in the micropores, forming an acidic microenvironment, which further accelerates metal dissolution, creating a vicious cycle. An appropriate amount of Ce can refine the metal matrix grain and reduce the active sites of the micro couple, while also promoting the formation of a dense oxide film that effectively blocks the penetration of erosion ions. However, excessive cerium weakens the repair capacity of the oxide film, thereby reducing its corrosion resistance. The specific mechanism is shown in Figure 5K.

Scanning Kelvin probe force microscope (SKPFM) revealed the direct effect of different Ce contents on the electrochemical activity and galvanic corrosion tendency of the micro-regions by measuring the difference between the surface voltametric potential and the morphology height. Figure 6 shows significant differences in the surface morphology and potential distribution of the HAZ region with changes in Ce content. When Ce (0 wt%) was not added, the surface potential difference was as high as 1.5 ± 0.1 μV, indicating a significant potential gradient within the material. This enhanced the galvanic coupling effect between the chromium-poor region and the surrounding substrate, leading to increased local corrosion susceptibility. When the Ce content was increased to 0.5 wt%, the potential difference decreased to 1 ± 0.1 μV, and the morphological height difference simultaneously decreased. This suggests that an appropriate amount of Ce effectively suppressed the potential difference between the chromium-poor region and the substrate by refining the grain boundary structure or stabilizing the passivation film, thereby reducing the driving force of galvanic coupling corrosion. At the optimal Ce content (1.0 wt%), the potential difference further decreased to 0.8 ± 0.1 μV, and the morphological height difference was only 22.7 nm. At this point, the material’s microstructural homogeneity was optimal, minimizing the electrochemical activity difference between the chromium-depleted area and the substrate, significantly weakening the galvanic coupling corrosion tendency. However, when the Ce content exceeded a critical value (1.0 wt%), both the potential difference and the topographic height difference increased. The potential difference rose to 1.5 ± 0.1 μV at 1.5 wt% Ce and further increased to 2 ± 0.1 μV at 2.0 wt% Ce. The introduction of excess Ce led to the precipitation of secondary phases, aggravated microstructural heterogeneity, and enlarged the potential difference between the Cr-poor and Ce-rich phases, thereby reactivating the galvanic coupling corrosion effect.

Figure 6.

Effect of different Ce contents on SKPFM in the HAZ region: a₁-c₁: 0Ce a₂-c₂: 0.5Ce a₃-c₃: 1.0Ce a₄-c₄: 1.5Ce a₅-c₅: 2.0Ce

Analysis of corrosion behavior

To assess the effect of different Ce contents on the galvanic coupling corrosion behavior of SASS/Q235 dissimilar joints, samples of welded joints containing 0–2.0 wt% Ce were immersed in natural seawater for 12 days. Macro-morphological observations (Figure 7A) showed no obvious corrosion products or large areas of spalling on the surface of all the samples, indicating that the passivation film effectively protected the substrate, in accordance with the requirements of the NORSOK M601 standard. However, microanalysis revealed a significant modulation of local galvanic coupling corrosion behavior by Ce content: corrosion behavior in the weld zone (Figure 7Ba₁-e₂) varied non-linearly with increasing Ce content. Without added Ce (0 wt%), the corrosion pits in the weld zone were dense and deep (Figure 7B-a₁), suggesting that the potential difference between the substrate and the passivation film coverage zone drove Cl⁻ enrichment, exacerbating anodic dissolution. When the Ce content was increased to 1.0 wt% (Figure 7B-c₁), the corrosion pits were the smallest in both density and size, indicating that moderate amounts of Ce significantly reduced the electrochemical activity difference between the anode (substrate) and cathode (passivation film area) by suppressing harmful phases at the grain boundaries (e.g., σ-phase) and optimizing the passivation film homogeneity. This weakened the galvanic coupling corrosion driving force. However, when the Ce content was ≥1.5 wt% (Figure 7B-e₁), the corrosion pit density and depth surged. Excess Ce led to the coarsening of CeO₂ particles, disrupting the continuity of the oxide film, increasing O²⁻ and Cl⁻ diffusion channels, and contributing to the enlargement of local micro-coupling active sites, which accelerated the galvanic coupling corrosion process. HAZ corrosion characteristics (Figure 7Ba₃-e₄) showed that the corrosion behavior of the HAZ on the SASS side was highly correlated with the weld zone. When Ce was not added (Figure 7B-a₃), the potential difference between the HAZ grain boundaries and the surrounding substrate was significant due to the formation of chromium-depleted zones, inducing intergranular corrosion and pitting (Figure 7B-a₄). After the addition of 1.0 wt% Ce (Figure 7B-c₃), Ce made the HAZ surface most uniform by purifying the grain boundaries and refining the grains (Figure 7B-c₄). CeO₂ nanoparticles were embedded in the passivation film, blocking the diffusion path of Cl⁻ along the grain boundaries and inhibiting micro-coupling corrosion. However, excess Ce (≥1.5 wt%) led to pronounced pitting and intergranular corrosion in the HAZ (Figure 7B-e₄) due to the formation of a heterogeneous interface between the coarsened CeO₂ particles and the substrate. This elevated the local potential gradient and reinitiated the galvanic coupling effect between chromium-poor zones and Ce-rich phases. The weight loss rate (Figure 8D) and standard deviation (Table 1) show that the corrosion rate of the 1.0 wt% Ce sample is the lowest, while the corrosion rate of the excess Ce (2.0 wt%) sample exceeds that of the Ce-free sample. This further confirms the threshold effect of Ce on galvanic coupling corrosion. Combined with the mechanistic model (Figure 7C), the moderate amount of Ce inhibited galvanic coupling corrosion through the following pathways: CeO₂ nanoparticles filled the defects of the Cr₂O₃ passivation film, reducing Cl⁻ permeability; Ce concentrated at the grain boundaries to inhibit the formation of chromium-depleted zones, reducing the anodic/cathodic activity difference. In contrast, excess Ce aggravated susceptibility to galvanic coupling corrosion by inducing the precipitation of heterogeneous phases and decreasing the repair capability ability of the passivation film.

Figure 7.

Corrosion morphology and mechanism after seawater immersion

(A) Macroscopic corrosion morphology.

(B) Microscopic corrosion morphology (0Ce (a₁-a₄), 0.5Ce (b₁-b₄), 1.0Ce (c₁-c₄), 1.5Ce (d₁-d₄), 2.0Ce (e₁-e₄).

(C) Corrosion mechanism.

Figure 8.

Three-dimensional morphology and corrosion rate after seawater immersion

(A) 3D morphology after seawater immersion.

(B) Working principle of CLSM.

(C) Surface roughness.

(D) Graph of weight loss rate.

Table 1.

Weight loss for different Ce contents

Alloys	Samples	M0/g	M1/g	S/cm2	T/h	g/(cm2·y)	Average	Deviation
0Ce	1#	24.037	23.887	20.8	288	2.50 × 10−5	2.43 × 10−5	5.35 × 10−7
	2#	23.930	23.788	20.8	288	2.37 × 10−5
	3#	23.896	23.751	20.8	288	2.42 × 10−5
0.5Ce	1#	25.864	25.560	20.8	288	5.08 × 10−5	5 × 10−5	5.89 × 10−7
	2#	25.618	25.323	20.8	288	4.93 × 10−5
	3#	25.719	25.421	20.8	288	4.98 × 10−5
1.0Ce	1#	25.639	25.523	20.8	288	1.94 × 10−5	1.97 × 10−5	2.08 × 10−7
	2#	25.592	25.474	20.8	288	1.97 × 10−5
	3#	25.491	25.372	20.8	288	1.99 × 10−5
1.5Ce	1#	25.809	25.512	20.8	288	4.96 × 10−5	4.93 × 10−5	4.97 × 10−7
	2#	25.787	25.496	20.8	288	4.86 × 10−5
	3#	25.589	25.291	20.8	288	4.97 × 10−5
2.0Ce	1#	25.684	25.385	20.8	288	4.99 × 10−5	4.99 × 10−5	1.43 × 10−7
	2#	25.614	25.312	20.8	288	5.04 × 10−5
	3#	25.480	25.185	20.8	288	4.93 × 10−5

The three-dimensional morphology of HAZ after seawater corrosion was further observed using confocal laser scanning microscopy (CLSM), as shown in Figure 8A. The working principle of CLSM is illustrated in Figure 8B, which provides high-resolution three-dimensional surface morphology information through laser scanning and optical detection technology. Combining this with the previous metallographic discussion, it is observed that the grains gradually refine with increasing Ce content, and the grain boundaries become more pronounced when the Ce content is greater than or equal to 1.5%. As seen in Figure 8A, in the HAZ (FZII) near the parent material, corrosion mainly manifests as pitting corrosion when the Ce content is less than 1.0%, while it shifts to intergranular corrosion when the Ce content exceeds 1.5%. This aligns with the earlier observation of grain refinement and the increased visibility of grain boundaries. By comparing the 3D morphology and height difference of the fusion zones (FZI and FZII) and the weld metal zone (WMZ) with different Ce contents in Figure 8A, it is evident that the surfaces are smoothest, with the smallest surface height difference, when the Ce content is 1.0%. Figure 8C further illustrates the trend of surface roughness in the fusion and weld metal zones with varying Ce concentrations, showing a decrease followed by an increase in surface roughness, with a minimum at 1.0% Ce. This suggests that the sample with 1.0% Ce exhibits the best corrosion resistance.

Analysis of corrosion resistance to SRB bacteria

Figure 9A illustrates the experimental procedure of immersing a dissimilar welded joint in a solution containing SRB, and Figure 9B shows the surface corrosion pattern after 12 days of immersion in the bacterial solution. Microscopic observation revealed differentiated corrosion characteristics in the weld zone and the HAZ: significant corrosion pits appeared on the surface of the control specimen, whereas the size of the pits was significantly reduced when the Ce addition was increased to 0.5 wt%. Notably, when the Ce content reached 1.0 wt%, no obvious corrosion morphology was detected on the surface of the specimens, indicating the best corrosion resistance. However, when the Ce addition exceeded 1.5 wt%, densely distributed macroscopic corrosion pits appeared on the surface of the material, suggesting that excess rare earth elements exacerbate localized corrosion. The effect of Ce addition on the antimicrobial properties of the materials followed a similar trend: as the Ce content increased, the SRB corrosion resistance of the specimens first improved and then weakened. The specimens with 1.0 wt% Ce exhibited the best corrosion resistance. This nonlinear trend results from the specific interaction mechanism between the rare earth element Ce and the SRB cell membrane/wall.^45,46 When Ce was added in moderate amounts (1.0 wt%), Ce³⁺ ions bound to the negatively charged bacterial membrane surface via electrostatic interaction, effectively enhancing membrane permeability and disrupting cellular structural integrity, thus inhibiting SRB metabolic activity (the survival rate decreased to 15%). However, excessive Ce (≥1.5 wt%) caused a significant increase in the material’s surface porosity, creating a porous structure that facilitated the penetration of H₂S, promoting the formation of thicker FeS corrosion products, and ultimately accelerating the corrosion rate. Figure 9C illustrates the regulatory mechanism of Ce content on the corrosion process of SRB through a schematic diagram.

Figure 9.

SRB corrosion process and mechanism diagram

(A) Flow chart of SRB bacterial solution immersion.

(B) Corrosion pattern.

(C) Corrosion schematic diagram.

To systematically evaluate the effect of rare earth element Ce addition on the corrosion resistance of welded joints against SRB, a fluorescent probe labeling method was used to characterize the living and dead bacterial populations on the surface of the specimens.⁴⁷ Based on the cellular structure of SRB as Gram-negative bacteria, Cy3 (red fluorescence) and FITC (green fluorescence) were used as dual-labeled probes, and in situ hybridization was carried out to detect SRB by specifically targeting the 16S rRNA sequence of SRB (see Figure 10B for the schematic diagram of the principle). Green fluorescence characterizes live bacteria with intact cell membranes, while red fluorescence indicates dead bacteria with damaged cell membranes.⁴⁸ The fluorescence microscopy images in Figure 10A show that the control group, without added Ce, exhibited a significant green fluorescence-dominant distribution (82.3 ± 0.5% of live bacteria) in the heat-affected zone (FZII) near the parent material, indicating that the base material had limited resistance to SRB. When the Ce content was increased to 0.5 wt%, the percentage of live bacteria decreased to 54.7 ± 0.5%. Notably, 1.0 wt% Ce demonstrated optimal bacterial inhibition, with the fluorescence signal in this region turning completely red (98.9 ± 0.5% of dead bacteria), and the live bacteria were below the detection limit. However, when the Ce content was further increased to 1.5 wt%, the percentage of live bacteria rebounded (12.4 ± 0.5%), indicating that the addition of excessive Ce may reduce antimicrobial efficacy. This concentration-dependent effect showed significant consistency in the heat-affected zone (FZI) of the Q235 steel and the weld metal zone (WMZ). The experimental data revealed a clear concentration threshold for the antimicrobial efficacy of elemental Ce: optimal inhibition was achieved at 1.0 wt% (98.2 ± 0.5% inhibition), while over-addition led to a decrease in antimicrobial performance (85.3 ± 0.5% inhibition at 1.5 wt%).

Figure 10.

SRB fluorescence test

(A) Fluorescent staining diagram.

(B) Principle of fluorescent staining.

Discussion

The present study is on the effect of Ce microalloying on the microstructure, corrosion resistance, and antimicrobial properties of the HAZ of SASS/Q235 dissimilar welded joints. The main conclusions are as follows.

• (1)Microstructural optimization: As the cerium content increases, the grains in the HAZ are refined, transforming from dendritic crystals to equiaxed crystals.
• (2)Improved corrosion resistance: The Rct of the cerium-containing sample with 1.0 wt% cerium was three times higher than that of the control group, while the passivation current density (ipass) was the lowest. This indicates that 1.0 wt% Ce facilitates the formation of a denser and more uniform oxide film. Electrochemical impedance spectroscopy (EIS) analysis after long-term immersion testing revealed that the material maintained a high impedance modulus, confirming the excellent long-term corrosion stability of the 1.0 wt% Ce alloy.
• (3)Antimicrobial synergistic effect: Under experimental conditions, the optimal cerium content (1.0 wt%) exhibited significant antibacterial effects, with a SRB survival rate of less than 1%. However, excessive cerium addition (≥1.5 wt%) resulted in the formation of porous oxide films, which accelerated the deposition of FeS corrosion products. This suggests that cerium’s antibacterial capacity has an optimal concentration range.
• (4)Optimal content: At 1.0 wt% cerium, grain refinement, passivation film integrity, and microbial resistance achieve a favorable balance. However, excessive cerium (>1.0 wt%) causes CeO₂ particle coarsening and reactivates galvanic corrosion between chromium-depleted zones and cerium-enriched phases, leading to a decline in material performance. This study identifies 1.0 wt% cerium as an industrially feasible addition level suitable for marine welding, providing a reference for applying laboratory research findings to marine infrastructure.

This study comprehensively investigates the regulatory effects of Ce microalloying on the microstructure, corrosion resistance, and antibacterial performance of the HAZ in dissimilar welded joints between SASS and Q235 carbon steel. The established optimization strategy and clearly defined cerium content threshold (1.0 wt%) provide a practical and scalable manufacturing solution for designing dissimilar steel marine engineering welded joints that combine high corrosion resistance and antibacterial performance. Future research could be extended to other welding processes (e.g., laser welding), dynamic marine environment simulations (considering tides and fluid dynamics), assessing the impact of more complex microbial communities (e.g., synergistic effects with iron-oxidizing bacteria), and systematically verifying the effects of cerium addition on the mechanical properties (fatigue strength, toughness) of welded joints to further optimize their industrial applications.

Limitations of the study

While this study demonstrates the significant benefits of 1.0 wt% Ce microalloying in enhancing the corrosion resistance and antibacterial properties of SASS/Q235 dissimilar welded joints, several limitations warrant consideration. First, the findings are based on tungsten inert gas (TIG) welding under specific parameters (0.57 kJ/mm heat input); performance under alternative welding methods (e.g., laser welding) or higher heat inputs, which may alter microstructural evolution and Ce distribution, remains unexplored. Second, corrosion and antibacterial tests were conducted in static natural seawater over 12 days, which may not fully replicate dynamic marine conditions (e.g., tidal fluctuations, hydrodynamic stress, or complex microbial consortia) encountered in long-term service. Third, the antibacterial evaluation focused exclusively on SRB; synergistic effects with other prevalent marine microorganisms (e.g., iron-oxidizing bacteria or algae) were not assessed. Additionally, industrial scalability was not addressed: vacuum induction melting for Ce incorporation may increase production costs, and the impact of Ce on mechanical properties (e.g., fatigue strength, toughness) in welded joints requires further validation. Future work should investigate these factors to optimize practical deployment.

Resource availability

Lead contact

Requests for further information and resources should be directed to and will be fulfilled by the lead contact, Quantong Jiang (jiangquantong@qdio.ac.cn).

Materials availability

This work did not produce fresh novel reagents. All resources utilized are commercially accessible or already published and cited appropriately.

Data and code availability

• •The data supporting the findings of this study can be obtained from the lead contact upon reasonable request.
• •No code or algorithm used for this study.
• •Any additional raw data or datasets created during this inquiry are available from the lead contact upon reasonable request.

Acknowledgments

This work was supported by Qingdao Key Technology Research and Development Project (no. 24-1-6-ghgg-2-hz), Wenhai Program of the S&T Fund of Shandong Province for Pilot National Laboratory for Marine Science and Technology (no. 2021WHZZB2301), and Hainan Provincial Joint Project of Sanya Yazhou Bay Science and Technology City (no. 2021CXLH0005).

Author contributions

X.L., formal analysis, investigation, and writing - original draft; Q.J., resources, methodology, and writing - review and editing; H.L., conceptualization and validation; D.L., resources and project administration; X.Z., visualization and supervision; N.L., conceptualization and supervision; J.D., investigation and supervision; B.H., funding acquisition.

Declaration of interests

The authors declare no competing interests.

STAR★Methods

Key resources table

REAGENT or RESOURCE	SOURCE	IDENTIFIER
Bacterial and virus strains
SRB strain	ATCC29579	N/A
Software and algorithms
Metallographic microscope	Carl Zeiss	Axio Lab A1
Scanning electron microscope	Hitachi	Regulus 8100
X-ray diffractometer	Bruker	Bruker D8 ADVANCE
Atomic force microscope	Oxford Instruments	MFP-3D
Electrochemical workstation	Princeton	PARSTAT 4000A
Fluorescence microscope	Leica	LEICA DM6 B
CLSM	Olympus	OLS5100 3D
ImageJ	N/A	https://imagej.nih.gov/ij/download.html- Version 1.53o

Experimental model and study participant details

This work did not need any unique experimental model.

Method details

After being cleaned and dried, the steel is placed in a vacuum induction furnace following evacuation, and a protective gas is introduced during melting. Once the steel is completely melted, the matrix is micro alloyed using a Ce–Fe intermediate alloy. Electromagnetic stirring ensures compositional homogeneity, and the molten steel is cast into ingots. After cooling, the ingots are demolded to obtain a 4 mm thick base material (BM). As gaseous elements, nitrogen (N), Sulfur (S), and oxygen (O) were quantified using combustion analysis. Phosphorus (P) was determined spectrophotometrically, while concentrations of iron (Fe), chromium (Cr), nickel (Ni), molybdenum (Mo), and the rare earth element cerium (Ce) were analyzed via inductively coupled plasma optical emission spectrometry (ICP-OES). The chemical composition of the alloy is presented in Table S1.

Tungsten Inert Gas (TIG) welding was used to join base metal (BM) steel and Q235 steel, using the following welding parameters: direct current positive connection (DCEN), voltage of 12 V, current of 90 A, welding speed of 80 mm/min, and heat input of 0.57 kJ/mm. ER-NiCrMo-3 welding wire, compliant with the AWS A5.14 standard, was selected. The welding process was divided into three passes: root pass, fill pass, and cover pass. During the root pass, the wire feed speed was set to 200 mm/min, and the interpass temperature was maintained at 25°C. For the fill and cover passes, the wire feed speed was increased to 400 mm/min, and the interpass temperature was adjusted to 100°C. Pre-welding preparation included a single-sided 30° bevel, a blunt edge of 1.5–2.0 mm, and a root gap of 2 mm. After welding, the specimen was cut into 40 mm × 20 mm × 2 mm samples, and the macrostructure of the welded joint is shown in Figure S1.

The samples were sequentially polished using 200#, 1200#, and 3000# sandpaper. After polishing, the surfaces were chemically etched for 10 s using a 3:1 mixture of hydrochloric acid and sulfuric acid (aqua regia), then dried with a hair dryer. The polished samples were characterized at multiple scales: microstructures and morphologies were observed using a Zeiss Axio Lab A1 optical microscope (OM) at 100×, 200×, and 500× magnifications. Microstructural analysis and elemental composition were performed using a Regulus 8100 scanning electron microscope (SEM) equipped with an OXFORD 550i energy dispersive spectrometer (EDS). The physical phase composition was determined using a Bruker D8-Advance X-ray diffractometer in the 10°–90° scanning range at a rate of 5°/min. Crystalline structures were examined with a Transmission Electron Microscope (TEM). Additionally, microstructure analysis was conducted using an MFP-3D Origin Atomic Force Microscope (AFM), and the surface potential distribution was measured with the Scanning Kelvin Probe Force Microscope (SKPFM) mode of the same AFM, within a scanning range of 10–20 μm.

Electrochemical performance was assessed using a Princeton P4000+ electrochemical workstation, which integrates open-circuit potential (OCP), electrochemical impedance spectroscopy (EIS), and dynamic potential polarization (Tafel) testing capabilities. The three-electrode system comprised a silver/silver chloride reference electrode, a 2 cm × 2 cm platinum counter electrode, and a working electrode with an effective surface area of 1 cm². Natural seawater from the Yellow Sea is used as the electrolyte. The experimental procedure was as follows: the heat-affected zone (HAZ) boundary of the sample was determined using a metal etching solution, followed by marking and cutting the sample into 10 mm × 10 mm × 2 mm specimens, which were placed on the surface of the working electrode for electrochemical testing. First, the open-circuit potential of the sample was measured for 1000 s, followed by EIS testing within a frequency range of 100 kHz to 0.01 Hz. Finally, Tafel scans were performed within a potential window of ±1 V around the open-circuit potential at a scan rate of 0.5 mV/s. To ensure the accuracy of the test results, three independent tests were conducted. To assess the long-term corrosion behavior, the samples were immersed in natural seawater for 12 days, with each content containing three parallel samples. The corrosion rate and standard deviation were calculated in accordance with ASTM G3121 and GB/T 27761-2011 standards, as follows:

$$V=\frac{\left(M_0-M_1\right)}{\left(\mathrm{S}·\mathrm{T}\right)}$$ (Equation 5)

where V is the weight loss rate, M₀ and M₁ represent the initial mass of the specimen and the mass after removal of corrosion products, respectively, S is the total surface area of the specimen (cm²), and T is the immersion time (days).

$$s=\sqrt{\frac{1}{n-1}{\sum }_{i=1}^n{\left(x_i-\overline{x}\right)}^2}$$ (Equation 6)

where $x_i$ is a single measurement, $\overline{x}$ is the mean value, and n is the number of experimental repetitions. The corrosion test was followed by simultaneous electrochemical retesting and characterization of the corrosion morphology.

Firstly, the medium for SRB was prepared, sterilized at 121°C under high temperature and pressure, and then transferred to a UV sterilization table. To maintain the anaerobic conditions required for SRB, nitrogen was used to remove oxygen from the medium for 30 min. The medium was then inoculated with SRB strain (ATCC 29579) and sealed using a double sealing system (plug-paraffin seal) to create an anaerobic culture environment. Pretreated metal specimens were fully submerged in the SRB bacterial solution (three parallel specimens were included in each setup) and placed in a thermostatic incubator at 30 ± 0.5°C for 12 days for microbial corrosion experiments(The concentration value of the SRB bacterial solution is 0.6). At the end of the experiment, the specimens were removed, and surface biofilm residues were gently rinsed with phosphate buffer solution (PBS, pH = 7.4). The specimens were then double-stained using the LIVE/DEAD BacLight Fluorescent Staining Kit (Thermo Fisher), and three fields of view were measured for each specimen. Bacterial survival was analyzed using ImageJ software. The corrosion morphology of the microorganisms was examined using scanning electron microscopy (SEM). The experimental setup and steps are shown schematically in Figure S2.

Quantification and statistical analysis

Figure 1: Images obtained via optical microscopy (OM) were analyzed using ImageJ software to measure dendrite spacing and grain size. Dendrite spacing was measured using the linear intercept method, where intercept lines were drawn on the image, and the number of dendrite arms intersecting each intercept line (n) was recorded. The dendrite spacing (λ) was then calculated using the following formula:

$$\lambda =\frac{L}{n}$$ (Equation 7)

where L is the intercept line length (μm), n is the number of intercept points, and each sample was measured 10 times, with the results averaged.

In accordance with ASTM E112 standards, the grain size was measured using the intercept line method. Test lines were drawn on the image, and the number of grain boundaries intersected by each test line was recorded. The average intersection length ($\overline{L}$) was calculated using the following formula:

$$\overline{L}=\frac{L_T}{P_T}$$ (Equation 8)

where $L_T$ is the total length of all test lines (μm), and $P_T$ is the total number of intersection points.

Table 1: The corrosion rate and standard deviation were calculated in accordance with ASTM G3121 and GB/T 27761-2011 standards, as follows:

$$V=\frac{\left(M_0-M_1\right)}{\left(\mathrm{S}·\mathrm{T}\right)}$$ (Equation 9)

where V is the weight loss rate, M₀ and M₁ represent the initial mass of the specimen and the mass after removal of corrosion products, respectively, S is the total surface area of the specimen (cm²), and T is the immersion time (days).

$$s=\sqrt{\frac{1}{n-1}{\sum }_{i=1}^n{\left(x_i-\overline{x}\right)}^2}$$ (Equation 10)

where $x_i$ is a single measurement, $\overline{x}$ is the mean value, and n is the number of experimental repetitions. The corrosion test was followed by simultaneous electrochemical retesting and characterization of the corrosion morphology.

Published: August 6, 2025