Introduction

Stainless steels resist corrosion primarily due to a natural passive film formed on the surface.1 However, the resistance of stainless steel to pitting corrosion is poor when the material is exposed to a chloride-containing environment.2 Exposure of stainless steel to hydrogen is known to increase corrosion susceptibility, so much work has been done to determine the effects of hydrogen on pitting.3,4,5,6,7 It was previously shown that hydrogen increases the surface activity of passive film and deteriorates passive film stability.5,6,7 In a recent review, Thomas et al.8 stated that hydrogen destabilizes the passive film by altering its chemical composition, thus changing the dielectric constant and the electronic properties of the passive film. However, the role of hydrogen in relation to the local electronic properties of passive film and pitting corrosion in stainless steel remains far from clarity.

It is known that corrosion failure of duplex stainless steels (DSS) often begins with pitting corrosion before subsequent stress corrosion cracking development in a chloride-containing environment.9,10,11,12,13,14 Yao et al.15 studied passive film behavior of hydrogen-charged 316 austenitic stainless steel and determined that hydrogen decreased the passive film breaking load, the Young’s modulus, and the fracture stress, and reduced the cohesive strength between the stainless steel matrix and the passive film. However, it is difficult to characterize the electrical behavior of hydrogen-containing passive film on duplex stainless steel using traditional macroscopic measuring methods, such as electrochemical impedance spectroscopy and photoelectrochemical analysis, because duplex stainless steel has a heterogeneous structure composed of ferrite and austenite grains. This leads to differences in semiconductor characteristics of the passive film formed on the different phase grains and grain boundaries. Fortunately, current-sensing atomic force microscopy (CSAFM) can detect the hydrogen-induced local electrical properties’ change of passive film on different areas of duplex stainless steel by acquiring the electrical conductivity.16,17,18 In previous research,18 we used CSAFM to characterize the semiconductor behavior of passive film on ferrite and austenite grains before and after hydrogen charging. However, the electrical behavior at grain boundaries was not investigated. In general, grain boundaries are more susceptible to pitting corrosion, especially in a hydrogen-containing environment. This is due to the grain boundaries containing more defects, making these areas preferentially gather hydrogen, which causes the subsequently formed passive film to have poor electrical properties.

Present work aims to use CSAFM current mapping and optical imaging to determine the performance of passive film formed on grain boundaries and different phase grains before and after hydrogen charging of duplex stainless steel. The more active areas of the passive film are identified as susceptible pitting initiation sites. This research furthers the understanding of the mechanisms of hydrogen-induced film breakdown and pitting corrosion in duplex stainless steel.

Results and discussion

Passive film conductivity

Fig. 1a illustrates the CSAFM map with 0.8 V applied tip bias of passive film formed before hydrogen charging. CSAFM measurements were conducted in the air with 328 nN tip normal load to obtain a clear current map image, prevent film fracture and limit electron tunneling effects.17,18 The ferrite and austenite phases can be identified in the CSAFM current maps based on our previous work.17,18,19 In Fig. 1a, the passive film on grain boundaries exhibits much higher conductivity (brighter color) than passive film on either the ferrite or austenite grains. This is because the grain boundaries act as hydrogen trap sites, thus affecting interfacial processes occurring at the material’s surface8 and creating unfavorable conditions for the development of stable, insulating passive film. Moreover, the passive film on austenite grains shows higher conductivity (brighter color) than on ferrite grains (darker color), which means that passive film on ferrite grains is more insulating (more stable) than on austenite grains. The current measured during CSAFM is negatively correlated with the corrosion resistance of the scanned area. It can be said that areas of the passive film with high CSAFM-measured current also have high pitting susceptibility. The low conductivity on the ferrite grains is attributed to enriched chromium oxides in the passive film formed in these areas.20 Current spikes in Fig. 1a are related to nanoscale defects present in the passive film. Figure 1b shows an optical image of pit initiation and propagation after hydrogen charging and FeCl3 exposure. More hydrogen is adsorbed in the austenite grains than in the ferrite grains due to hydrogen having a higher solubility and a lower diffusivity in the face-centered cubic austenite phase compared with the body-centered cubic ferrite phase.21,22 It can be concluded that hydrogen facilitates pitting corrosion initiation at grain boundaries and austenite grains, i.e. hydrogen makes the passive film less stable and more conductive.

Fig. 1
figure 1

a Current-sensing atomic force microscopy 3D current map of the air-formed passive film with high current sites at grain boundaries and the austenite phase; b Optical image showing pitting corrosion after hydrogen charging and FeCl3 exposure

It is generally accepted that pitting is a result of the localized breakdown of the passive film, which usually occurs at specific sites of low film stability.23 In our previous research,24 pitting of a hydrogen-charged specimen exposed to a 6% FeCl3 solution initially occurred at the grain boundaries or at the austenite grains, and then appeared at the ferrite grains. Therefore, areas with poorly insulating passive film, such as grain boundaries and austenite grains, can be considered more susceptible to pitting corrosion. Pitting nucleation occurs at grain boundaries because these areas gather hydrogen preferentially, which causes the passive film to easily break down. Meanwhile, the vast majority of hydrogen gathers in the austenite grains, which can accelerate the breakdown of the passive film and cause pitting nucleation in this area. However, under the same hydrogen charging conditions, less hydrogen gathers in the ferrite grains, so pitting nucleation appears later here compared to other areas.

Passive film current density

Combining the contact area calculated by the Hertzian elastic contact formula with the current output of the CSAFM, yields the average current density of passive film at the grain boundaries and austenite grains of 0.074 nA/nm2 and 0.003 nA/nm2, respectively. The maximum current density at the grain boundaries and austenite grains is 0.109 nA/nm2 and 0.077 nA/nm2, respectively. Current flow is negligible on the ferrite grains, corresponding to black regions in Fig. 1a. Thus, it can also be deduced that pitting corrosion will preferentially begin at the grain boundaries or in the austenite phase of 2507 duplex stainless steel, which matches our previous results.21,24,25

CSAFM was used to directly observe the conductivity of passive film formed on grain boundaries, austenite and ferrite grains prior to hydrogen charging. The passive film on austenite grains and grain boundaries had much higher conductivity than on ferrite grains. Hydrogen charging was seen to cause an increase in corrosion susceptibility due to poor passive film formation, especially at grain boundaries and at austenite grains. This resulted in pitting corrosion initially occurring in these areas before spreading to ferrite grains after hydrogen charging and FeCl3 solution exposure.

Methods

Sample preparation

The material used for the experiment was commercially available 2507 duplex stainless steel with the composition shown in Table 1, as reported previously.26 The specimen was machined into a square sheet with the dimensions of 10 × 10 × 0.7 mm. The specimen was prepared by first wet grinding with silicon carbide paper up to 2000 grit before mechanical polishing with 1.5 μm diamond paste. Final polishing with a solution of HNO3:H2O = 1:1 for 20 s at 1.2 V was performed to ensure that surface stresses were eliminated. The specimen was then cleaned in ethanol, dried with nitrogen gas flow, and held at 25 °C in the air for 24 h to develop a native passive film. This treatment was performed to achieve a very smooth surface suitable for the CSAFM measurements. Hydrogen charging was performed at room temperature in 0.5 M H2SO4 aqueous solution with 0.25 g/L As2O3 added to increase hydrogen adsorption. A platinum electrode was used as the anode and the specimens acted as the cathode. The specimen was charged from all sides at 2.5 mA/cm2 current density for 24 h. Then, the passive film was formed on the hydrogen-charged specimen with the same parameters as for the uncharged specimen.

Table 1 Chemical composition in wt.% of 2507 duplex stainless steel for ferrite (α) and austenite (γ) phases

Instrumentation

CSAFM measurements of the passive film before and after hydrogen charging were conducted in the air at 25 °C and 25% relative humidity using an Agilent 5500 AFM set to current-sensing mode and equipped with the DPE14/AIBS platinum-coated silicon tip (50 nm tip radius and 5.7 N/m force constant). The optimal tip normal load was determined to be 328 nN, which was found by bringing the CSAFM tip into contact with the sample surface. Then, the normal load was increased to a certain value and the I–V curve was recorded. After this, the tip was moved to a nearby unscanned area and the I–V curve was recorded with a higher tip normal load. Ten different normal loads were used. At the 328 nN tip normal load, a stable and reproducible I–V relationship was observed. The Hertzian equation neglects interfacial adhesion. Therefore, the error in the calculated contact area is larger at lower tip normal loads. At tip normal loads above 492 nN, the high tip pressure caused passive film deformation outside of the elastic range. This was determined by adding load and then observing if the measured current returned to original values after the additional load was removed. A tip bias was chosen based on the breakdown voltage of passive film and clarity of current maps at the 328 nN tip normal load.

The current density at the contact area

For current density calculations, it was necessary to determine the contact area of the CSAFM tip on the film surface. This was performed using the Hertzian elastic contact formula between a sphere and a flat plane,27,28 as shown in Eq. 1.

$$A = \pi a^2 = \pi \left( {\frac{{3PR}}{{4E}}} \right)^{\frac{2}{3}}.$$
(1)

The effective Young’s modulus, E, is calculated as shown in Eq. 2.

$$E = \left[ {\frac{{\left( {1 - {\mathrm{\nu }}_1^2} \right)}}{{E_1}} + \frac{{\left( {1 - {\mathrm{\nu }}_2^2} \right)}}{{E_2}}} \right]^{ - 1}.$$
(2)

In these equations, A is the contact area, a is the contact radius, and E is the effective Young’s modulus. R is the tip radius and P is the normal tip load, while ν1, ν2, and E1, E2 are the Poisson’s ratio and Young’s modulus of the tip and the specimen, respectively. For the materials used in this study, ν1 = 0.28, ν2 = 0.3, E1=169 GPa, and E2=200 GPa.