To meet the fast-growing demand of alloys serving in harsh environments, great efforts have been made to design alloys combining enhanced mechanical properties and improved corrosion resistance.1, 2 One strategy to achieve this goal is to modify the alloy surface using surface engineering techniques, such as carburization. The deleterious precipitation of intergranular Cr carbides during carburization of stainless steels by standard methods can be prevented by low temperature carburization (LTC),3, 4, 5 which allows the diffusion of carbon atoms into the alloy matrix resulting in a high concentration, without precipitation of carbides. Generally, carbon atoms are incorporated in the octahedral sites of face-centered cubic (FCC) structures, forming a subsurface zone of concentrated interstitial carbon, with carbon fraction up to 0.15 and the mean depth on the order of 10 mm.6 Even low concentrations of interstitial carbon can improve the corrosion properties of iron,7, 8 but the benefits are small relative to those imparted by higher concentrations.
Concentrated interstitial carbon was widely reported to improve the corrosion resistance of stainless steel,3, 4, 5, 9, 10, 11, 12, 13, 14, 15, 16, 17, 18 Ni-based superalloys,19 and high entropy alloys.20 Most previous studies focused on the changes in passive film properties.4, 10, 11, 12, 13, 16, 18, 20, 21 However, no consensus regarding the underlying mechanism has been reached. Some considered that the interstitial carbon increased passivity, including the growth kinetics and stability of the passive film,4, 10, 12, 21 while some others espoused the totally opposite opinion.13, 16 Furthermore, in some studies the passive current density of carburized samples, a key parameter to evaluate the passivity, was found to be lower than or similar to untreated samples,3, 4, 10, 14 while others have reported higher passive current densities for carburized samples.11, 15, 17 Clearly, the change in passive film properties is not a satisfactory explanation for the greatly improved corrosion resistance, and the effect of interstitial carbon needs to be further explored.
Generally, the evolution of a corrosion pit occurs in three stages: pit nucleation, metastable pit growth and stable pit growth. According to a recently proposed unified framework for pitting corrosion,22, 23, 24, 25, 26 there are two critical steps that will influence the overall pitting process: (1) pit initiation associated with the breakdown of the passive film; (2) the transition from metastable pit to stable pit growth. The corrosion resistance of an alloy can be attributed to either the resistance to passive film breakdown, which is determined by the passive film properties, or the resistance to the transition of a pit from metastable to stable growth, or both.
To better understand the key role of concentrated interstitial carbon in enhancing the pitting corrosion resistance, its effect on pit stability must be understood. According to the unified pitting framework,22, 23, 24, 25, 26 pit stability can be evaluated by the parameter idiss,max, which is the pit dissolution current density in the salt-film-free state and is also the maximum dissolution current density that a pit can achieve under given conditions. A larger idiss,max corresponds to a greater ability to stabilize a pit, thus promoting the transition of a pit from metastable to stable pit growth. In contrast, a lower idiss,max represents a lower pit stability thus a higher pitting corrosion resistance. A detailed description of pit growth stability is described in Supplementary Materials. In this study, we demonstrate that the superior pitting resistance of carburized alloy is not derived from the changes in passive film properties, but should be attributed to the significantly reduced idiss,max because of the strong covalent bonds formed between interstitial carbon and its neighboring metal atoms, making it more energetically difficult to dissociate metal atoms for metal dissolution.
Origin of Superior Pitting Resistance: Passive Film Breakdown vs. Pit Growth Stability
To exclude effects of the untreated substrate, fully carburized AISI 316 stainless steel (hereafter referred to as SS316) foils (50 mm in thickness) and wires (51 mm in diameter) were prepared by LTC treatment. Details of the samples are given in the Supplementary Materials. Figure 1a shows typical cyclic potentiodynamic polarization (CPP) curves of carburized and untreated foil samples tested in 0.6 M NaCl at room temperature. Although the carburized SS316 sample exhibited a lower breakdown potential compared with the untreated sample and the carburized sample was polarized to high potentials, no severe corrosion attack was observed on the carburized sample after the tests (Figure 1b). The light corrosion attack likely originated from the transpassive dissolution at high potentials, rather than pitting corrosion, which is in agreement with the small positive hysteresis observed during the reverse scan. In contrast, the curve of the untreated sample exhibits a large positive hysteresis (Figure 1b), and large penetrating pits were observed after the test (Figure 1c). Therefore, in spite of a lower breakdown potential, the pitting resistance of the carburized sample is actually significantly higher than that of the untreated sample. On the other hand, the carburized SS316 also exhibits higher passive current density with substantial noise, probably associated with metastable pitting events. This indicates that the passive film formed on the carburized SS316 is less protective than that of the untreated SS316. Additionally, in the region between the breakdown and transpassive dissolution, the carburized SS316 exhibited pseudo passivity with the nominal current density of about 10-4 A cm-2, which is about one order of magnitude higher than the passive current density before breakdown. The breakdown associated with the increase in current density at the potential of about 0.28 VSCE is really a transition from passivity to another condition, referred to here as pseudo passivity. Note that this pseudo passivity is not a true passive state for stainless steel as the current density is significantly higher than that in the passive regime. Furthermore, different than the untreated sample, the carburized SS316 did not fully repassivate during the reverse scan but again exhibited pseudo passivity, indicating that the repassivation ability of the carburized sample is much lower than that of the untreated sample. Based on the CPP results, it can be seen that the carburized SS316 forms a weak passive film, and its resistance to passive film breakdown and the ability to repassivate after the breakdown is actually low. Therefore, the superior pitting resistance of carburized SS316 is not derived from improvements in passive film properties.
To investigate the pit growth stability, untreated and fully carburized SS316 wires were made into one-dimensional (1D) artificial pit electrodes,26 which are able to create a single pit, allowing pit growth in one direction (Figure 2a). Figure 2b shows CPP curves for 1D electrodes in 0.6 M NaCl solution at room temperature. The absence of metal at the pit mouth after the tests, as indicated by the black round area in the inserted optical images, shows that both samples experienced considerable dissolution and 1D pits were successfully created. The electrochemical behavior of the carburized SS316 1D electrode is similar to that of the carburized foil sample (Figure 1a), exhibiting passivity, a lower breakdown potential and a current peak, followed by pseudo passivity and transpassivity. The behavior after reversal of scanning is also the same: a small hysteresis followed by pseudo passivity. The decrease of current density to the pseudo passive level indicates that the 1D pit could not transition into stable growth even though it experienced high rate dissolution in the transpassive region. The untreated SS316 1D electrode exhibited different behavior than the untreated foil sample; breakdown was not observed until at high potentials associated with transpassive dissolution. After a period of reverse scan that followed back along the upward scan, the transpassive dissolution transitioned into stable 1D pit growth under diffusion control, as indicated by the increase in current density to a high stable value. The slight decrease of the current during the following downward scanning is typical of diffusion-controlled 1D pit growth caused by the effect of increasing pit depth on the diffusion-limited current density. Additionally, the 1D pit of untreated SS316 exhibited clear repassivation at low potentials associated with a sharp decrease in current. CPP measurements of 1D electrodes were also performed in 1 M HCl (Figure 2c). In such a harsh environment, the untreated SS316 is still able to passivate, as indicated by the clear active-passive transition followed by a broad passive region. Stable diffusion limited 1D pit growth was also observed in the reverse scan. In contrast, the carburized SS316 did not fully passivate and only exhibited pseudo passivity with current density on the order of 10-3 A cm-2. However, even in this harsh bulk environment, the carburized SS316 1D pit still did not develop into to stable 1D pit growth, as indicated by the small hysteresis during the reverse scan. Based on CPP tests on 1D electrodes, it can be seen that, although the carburized SS316 exhibited a much lower resistance to passive film breakdown, its resistance to the transition of a pit to stable growth is extremely high.
To quantitatively evaluate the pit growth stability, the metal dissolution rate in local pit environment, i.e. idiss,max, must be determined. According to the unified pitting framework,26 idiss,max must be measured from a growing pit under charge transfer control, which corresponds to the salt-film-free condition. CPP tests show that the carburized SS316 exhibits pseudo passivity at potentials below the transpassive region (Figure 2b). Therefore, idiss,max of carburized samples can only be measured within the transpassive region where active dissolution is achieved. Figure 2d shows the potentiostatic polarization of a carburized SS316 1D electrode initiated at 1.2 VSCE. The continuous increase of current density is generated by the increase of pit solution concentration at the pit surface, and it indicates that the 1D pit growth remains under charge transfer control.26 Therefore, the measured pit current density is the idiss,max of carburized SS316 at 1.2 VSCE,[*] and it is on the order of 10-1 A cm-2. After stepping to 0.7 VSCE, the current density plunged to the pseudo passive level. In contrast, the untreated SS316 1D pit transitioned to diffusion-controlled pit growth immediately after the final breakdown, as indicated by the continuous decrease of the current density (Figure 2e). After stepping to 0.7 VSCE, the diffusion-controlled pit growth was still maintained, which, again, indicates a much higher pit stability compared with the carburized sample. The diffusion-controlled pit growth indicates the presence of salt film on the pit surface, thus idiss,max of the untreated sample cannot be obtained directly from potentiostatic polarization measurements. The approach using a 1D electrode to determine idiss,max for alloys with relatively high pit stability has been developed previously.26 The idiss,max of untreated SS316 at 1.2 VSCE was determined and is on the order of 1010 A cm-2 (details described in Method), which is about 11 orders of magnitude higher than that of the carburized sample. According to the unified pitting framework,22, 23, 24, 25, 26 the extremely reduced idiss,max will remarkably decrease the growth stability of a pit formed on carburized SS316, preventing it from transitioning into stable growth, thus significantly improving the pitting corrosion resistance of the carburized alloy.
Effect of Interstitial Carbon on Bond Strength of an Alloy System
According to a conceptual model developed by Marcus to explain the effects of certain alloying elements on localized corrosion,27 the metal–metal bond strength is a fundamental property that will significantly influence the passivity and metal dissolution. Generally, the addition of dissolution moderators such as Mo, Nb, and W can increase the bond strength and make it energetically more difficult to separate metal atoms from the substrate. However, when a considerable amount of non-metallic elements is incorporated into the alloy without the formation of precipitates, the effect of non-metallic elements on the bond strength must be considered to interpret the behavior of passivity and metal dissolution (thus pit growth stability).
The effect of interstitial carbon on the bond strength of the FeCr:C alloy system can be evaluated by bond energies, which can be calculated using a recently developed bond-order bond energy (BOBE) model based on density-functional theory (DFT) calculations.28 Figure 3 shows the fitting of total bond energies (Ebond,tot) of a 32-metal-atom FeCr:C supercell (FCC structure) with various compositions. The extracted metal–metal (M–M) and metal–carbon (M–C) bond energies are listed in Table 1 (detailed approach described in the Method section). A more negative bond energy value indicates a stronger bond. The calculated bond energies for Fe–C and Cr–C bonds are stronger than Fe–Fe and Cr–Cr bonds by a factor of 1.7 and 2.0, and stronger than Fe–Cr bonds by a factor of 1.8 and 1.6, respectively. Because the fitting is based on a variety of different FeCr:C compositions (28 in total), including those with and without interstitial carbon, the excellent fit quality suggests that the presence of interstitial carbon does not change the M–M bond strength in an appreciable way. Therefore, the formation of the added M–C bonds with significantly higher strength than the M–M bonds should be the crucial factor responsible for the decreased passivity and metal dissolution rate. Note that the bulk bond energies obtained here should be different from surface bond energies due to the different bonding situation. However, previous DFT work29 has shown that the common approximation of direct proportionality between bulk and surface bonds27, 30 is often justified. Our results thus suggest that if a metal atom on the surface is bonded to a neighboring interstitial carbon, it will have stronger bonding, making it more difficult for this atom to dissolve or break its bonds to participate in oxide formation.
To get further insight into the nature of the M–C bond, the charge density difference between the DFT-optimized ground state distribution in the solid and the superposition of the electron density of the unbound atoms was calculated for austenitic Fe with and without interstitial carbon. The redistribution of electrons in the solid allows for analysis of the types of bonds that form. Looking at pure Fe in Figures 4a and b, the electron distribution is found to be nearly uniform in the region between the metal ions, indicating that the valence electrons are more or less completely delocalized from their cores, which is consistent with the classical model of metallic bonding. Figures 4c and d show the charge density difference in the Fe32:C1 supercell. High electron accumulation can be found between the interstitial carbon and its neighboring Fe atoms, forming what is typically referred to as a bond charge in covalent solids.31 These bond charges are typical of the bonding in group IV elemental semiconductors32 and look equivalent to that of Fe–C bond obtained here.
Next, the density of states (DOS),33 which is the number of electron states per unit volume at a given energy, was calculated to further investigate the nature of the Fe–C bonds. The local density of states (LDOS),34 which is the density of states projected on the atomic orbitals of each atom, shows that, in the energy region from -7.5 to -5.0 eV and the region around -12.5 eV, new states (peaks) appear in the LDOS of the Fe atoms neighboring carbon, whose energy ranges overlap with peaks of the interstitial carbon LDOS (Figure 5a). Overlapping LDOS curves typically indicates the formation of a bond with hybridized orbitals found in covalent bonds and is expected for interstitial carbon in an octahedral site bonding to transition metal atoms, mainly involving σ bonds between the C 2s, 2p and Fe 4s orbitals, and π bonds between C 2p and Fe 3p, while the Fe 3d orbitals should be mostly non-bonding to carbon.35, 36, 37 Screening should prevent such bonds from forming with Fe atoms outside of the carbon nearest-neighbor shell, which is indeed already the case for the second nearest Fe atom, where the LDOS has no peak and looks more or less identical to that in pure iron (Figure 5a). This is also in agreement with the charge density difference plot in Figure 4d, where the charge distribution around Fe atoms only differs from that of pure Fe on the nearest Fe neighbors of the interstitial carbon. These findings are also in agreement with the calculated bond energy results and explain why interstitial carbon does not affect the M–M bond strength in the alloy.
Since the above qualitative comparisons suggest the formation of covalent bonding between Fe and interstitial carbon, pCOHP analysis was then performed to identify covalent bonding, nonbonding, and antibonding contributions for the electrons represented in the DOS. Positive values of -pCOHP indicate bonding states whereas negative values indicate anti-bonding states. Figure 5b shows the -pCOHP plot of an interstitial carbon atom and one of its neighboring Fe atoms. All bonding states are found below the Fermi level (EF = 0 eV) and all antibonding states are found above the Fermi level, confirming covalent bonding characteristic between interstitial carbon and its neighboring Fe atoms.38 Figure 5c compares the LDOS of interstitial carbon and Fe at the first neighbor site with the total -pCOHP and orbital-resolved -pCOHP results between interstitial carbon and neighboring Fe atoms. We found the dominating bonding states around –12.5 eV to involve C 2s and Fe 4s orbitals (Figure 5c-iii) as expected from molecular orbital theory,36, 37 with a smaller contribution of C 2s – Fe 3d bonding (Figure 5c-iv), which is not observed in molecules.36, 37 The bonding states in the energy range between –7 and –4 eV are contributed by σ bonding between the C 2p orbitals and Fe 4s orbitals (Figure 5c-v) as expected by molecular orbital theory, whose geometry/symmetry is mostly responsible for the octahedral bonding configuration, as well as weaker interactions between C 2p orbitals and Fe 3d and 3p orbitals (Figures 5c-vi and 5c-vi; the integral over the -pCOHP curve is a measure for the related bond strength). In addition, all the high LDOS values between –5 and 0 eV, which are from the Fe 3d orbitals, has -pCOHP values of zero and does not form covalent bonds. The Fe 3d orbitals are thus the main origin of the homogeneous charge density responsible for metallic bonding, as shown in Figure 4.
In addition to the covalent bond charges, the entire region of the interstitial carbon shows an increased charge density, accumulated from its Fe neighbors, which is similar to the situation in polar binary semiconductors such as SiC.39 Such charge transfer is expected when considering the different electronegativities of Fe (1.83) and carbon (2.55), for which Pauling’s theory40 predicts 12% ionic character. To quantify the transfer, Bader charge analysis41 was performed on the interstitial carbon and its six neighboring Fe atoms. The Bader charge of the interstitial carbon is -1.1, while that of neighboring Fe atoms is 0.12 with the discrepancy spilling to second-nearest neighbors with charge values around 0.03. Therefore, Fe–C bonds are of polar covalent nature, where the shared electrons in the bond charge are closer to the interstitial carbon atom. Such covalent bonds are also expected formed between interstitial carbon and other transition metal atoms, such as Cr (electronegativity 1.66), Ni (electronegativity 1.91), etc.
The formation of a covalent M–C bond is also supported by the Hall effect measurements on carburized SS316 foils, where the conduction electron concentration (Ne) was found to decrease with increasing concentration of interstitial carbon.42 In light of the present findings, the decreased Ne of carburized SS316 should be attributed to the formation of covalent M–C bonds, where a portion of the mobile Fe electrons become localized in the covalent M–C bonds. For the fully carburized foil with an average atomic fraction of 0.1 interstitial carbon, the Ne drops by 29%,42 which approximates to the value of 33% obtained by theoretical analysis on covalent bonds. Detailed analysis is presented in the Supplementary Materials. As a result, both the electrical and thermal conductivities are decreased in the carburized alloy.42 Additionally, XPS results on the carburized alloys, including both martensitic and austenitic steels,8, 43 show that the electron binding energy of Fe atoms for the sample with high interstitial carbon is higher than that of the sample without interstitial carbon, which can also be attributed to the formation of covalent M–C bonds.
Summarizing the above analysis, there is strong and widespread evidence that interstitial carbon forms covalent bonds with its nearest neighbor metal atoms. It is also found that the effect of interstitial carbon is rather local and already becomes negligible for second nearest neighbors. If the added strong bonds are the dominant cause for increased pitting resistance and decreased passive film formation, this suggests that the concentration of interstitial carbon needs to be high for a significant effect, since only then the cumulative bond strength of the alloy system will be greatly increased. In that case, the covalent bonds formed by interstitial carbon can make it more energetically difficult to dissociate metal atoms for oxide formation or metal dissolution. The formation of M–C covalent bonds can thus well explain the low passivity and extremely low idiss,max (thus low pit growth stability) of carburized SS316 observed in this work. The fundamental understanding accomplished in this work provides a new strategy for designing highly corrosion resistant alloys, namely alloying concentrated interstitial atoms that form strong covalent/ionic bonds with the matrix atoms. Suitable alloying elements include but not limited to carbon; concentrated nitrogen and boron interstitials are also expected to have similar effects.
[*] Potential drop of the solution for carburized SS316 1D pit is calculated to be very small, only 0.034 V, thus it can be ignored. Therefore, the applied potential (1.2 VSCE) can be considered as the potential at pit surface.