Introduction

In a corrosive environment, the inclusions in low-alloy steel preferentially experienced pitting corrosion and extended to uniform corrosion1,2. With the progress of smelting technology, corrosion-resistant elements, such as Cu, Ni, Cr, and Sb, were added, and weathering steel with good corrosion resistance was obtained3. The corrosion behaviors of corrosion-resistant elements on weathering steel have been studied extensively due to the characteristics of the rust layer. Corrosion-resistant elements had important effects on the stability of the rust layer of low alloy steel. In this process, corrosion-resistant elements acted alone or in coordination. Zhang et al.4 suggested that the addition of Cu accelerated the formation of corrosion products and increased the corrosion current density, and Cu improved the long-term protection of the rust layer by forming CuFeO2. There was an interaction between Cu and Sb5,6. Wu et al.7 added Cu and Sb with different contents into 3Ni steel and found that there was a synergistic effect between 0.5% Cu and 0.2% Sb, which improved the corrosion resistance of 3Ni steel. The inner rust layer was rich in Cu, Sb, and Ni, and its compactness, phase composition, element distribution, and electrochemical activity were significantly optimized. Zhang et al.8 considered that Sb increased the consumption of H+ and decreased the electrochemical activity of steel during the corrosion process. The process of Sb depletion reduced the local acidification of the steel surface and accelerated the transformation from γ-FeOOH to Fe3O4 in the rust layer. Sb2O3 made the rust layer more uniform and denser; thus, the layer had more effective resistance to corrosion ions9. Sun et al.10 suggested that the addition of Cr in steel promoted the transformation of α-FeOOH. Cr-doped α-(Fe,Cr)OOH and FeCr2O4 were ~10-nm spherical particles, which improved the compactness of the rust layer. Cr formed Cr2O3 and Cr(OH)3 with good corrosion resistance and effectively refined the rust particles. The diffusion resistance levels of corrosive substances were increased. Dong et al.11,12 found that Cr(OH)3 produced by Cr hydrolysis in low-alloy steels inhibited the growth of FeOOH in rust and led to a denser rust layer. Furthermore, the presence of Cl− decreased the formation of NiFeO4 and Cr(OH)3, decreased the nucleation site of FeOOH, and promoted the formation of large-sized FeOOH, increasing the porosity in the corrosion product film. Jiang et al.13 considered that the enrichment of Ni in corrosion products inhibited the oxidation reaction of corrosion products, especially the formation of amorphous materials. Sun et al.14 suggested that Ni and Cu were enriched in the inner rust layer of NiCu steel and synergistically enhanced the compactness of the rust layer. In addition, Zhang et al.15 found that Cu, Ni, and Mo promoted the agglomeration process and improved the performance of the rust layer in NaCl solutions with low concentrations; with high concentrations, the rust layer was loose.

The corrosion-resistant elements Cu, Ni, Cr, and Sb played important roles in improving the rust layer characteristics of low-alloy steel; however, the role of corrosion-resistant elements in the local corrosion behavior of low-alloy steel was not clear. Among the many studies on the local corrosion mechanism of inclusions, only some researchers have observed the corrosion characteristics of corrosion-resistant elements in inclusions. Wang et al.16 studied the local corrosion behaviors of (Ca, Mg, Al)–Ox–Sy inclusions in EH36 steel and found that the corrosion products covered the surfaces of the inclusions and contained Cu2S. Liu et al.17 observed the enrichment of Cr and Cu in the solution products around the Al2O3 inclusion, suggesting that it promoted the formation of a dense rust layer and inhibited the cathodic corrosion reaction. Zhou et al.18 studied the effects of Cu on the dissolution of MnS in stainless steel. Cu could replace Mn in the surface layer of MnS to form insoluble copper sulfide compounds, and Cu could stably dissolve the harmful sulfur forms released from MnS; this phenomenon was helpful for reducing the adverse effect of manganese sulfide on the corrosion resistance of stainless steel. Jeon et al.19 suggested that the addition of Cu to the matrix increased the activity of Cr, promoted the precipitation of Cr-containing inclusions, and decreased pitting resistance. Ma et al.20 added rare-earth elements to 2205 duplex stainless steel and found that Cr could form Cr2O3–RE2O3 and (Ca, Mg, Cr, RE)xOy inclusions in stainless steel. Williams21 found the halo of CrxSy in MnS inclusions.

Localized corrosion behavior induced by inclusions in low-alloyed steels. Low-alloyed steels improved the electrochemical properties of the material as well as the rust layer transformation characteristics through the addition of corrosion-resistant elements, which played an important role in enhancing the corrosion resistance of low-alloyed steels. However, the influence of the characteristic state of corrosion-resistant elements in the initial corrosion and its distribution in the matrix on the initial corrosion behavior of low-alloy steels was not clear. In this article, the distribution of corrosion-resistant elements in the matrix was studied, and the distribution characteristics of the elements in the inclusion region were investigated. The local corrosion behaviors of Cu and Sb in the inclusion region and the corrosion characteristics of Cr in the pearlite region were investigated, and the segregation behaviors of Cu and Sb in inclusion defects and their effects on local corrosion were analyzed.

Results

Characteristics of segregation at grain boundaries

The distribution characteristics of Cu and Sb in low-alloy steel affected the corrosion behaviors of Cu and Sb at the initial corrosion stage. The crystal defects in low-alloy steel included grain boundaries, phase boundaries, subgrain boundaries, and dislocations, which had high-energy levels. When the elements dissolved in the alloy, they interacted with crystal defects. Corrosion-resistant elements greatly influenced the intracrystalline distortion energy. Solute atoms that were larger or smaller than the matrix atoms migrated from the crystal to the grain boundaries, phase boundaries and dislocations, creating grain boundary segregation. The main reason for grain boundary segregation was the elastic interaction between the solute atom and Fe atom, and the segregation process was spontaneous. The contents of Cu, Ni, Cr, and Sb in Q500 were very small, and they were enriched in the defect area because of the interactions with crystal defects, greatly influencing the microstructures and properties. Precipitation of Cu had a pinning effect on dislocations and grain boundaries, which could effectively prevent recovery and recrystallization and significantly affected grain growth inhibition, thus increasing the strengths of metals22. However, the segregation of elements at the grain boundary reduced the surface energy of the grain boundary, and the impact toughness of the steel decreased significantly23.

The studied material is the low-alloy steel Q500, the composition of which is shown in Table 1, and its organization is shown in Fig. 1. The segregation of corrosion-resistant Cu, Ni, Cr, and Sb at the grain boundary of low-alloy steel Q500 is shown in Fig. 2. During the cooling process of supercooled austenite, Cr and Mn formed CrxCy and MnxCy precipitates with carbon. The elastic interactions between solute atoms and matrix atoms resulted in the segregation of Mn, Ni, Cr, and Sb at the grain boundaries. Cu precipitated near the grain boundary. There was competition among the elements of grain boundary segregation, and the segregation rate was affected. In addition, some of the elements had promoting effects on cosegregation. The nucleation of Cu in the segregation process was more difficult than that of carbide, and the precipitated phase Cu was the result of the mutual promotion of Cu, Ni, and Sb in the segregation process. Some studies showed that the contents of Cu precipitates and Ni in Sb-bearing steel were higher24. Ni promoted the enrichment of Cu-rich precipitates in high-strength low-carbon steels25,26.

Table 1 The composition of low-alloy steel Q500.
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Fig. 1
figure 1

Microstructure of pearlite and ferrite of low-alloy steel Q500.

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Fig. 2: Segregation of corrosion-resistant elements at grain boundaries.
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a SAED analysis of adjacent grains in the grain boundary. b EDS analysis of the grain boundary with corrosion-resistant elements.

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The behaviors of corrosion-resistant elements in the grain are shown in Fig. 3. During the transformation from austenite to pearlite, the diffusion coefficient of carbide-forming element Cr in austenite was much smaller than that of C. The diffusion and segregation of Cr and Mn in the Fe3C hindered the formation of eutectoid carbides, reducing the pearlite cementite. The high-resolution transmission (HRTEM) analysis of the cementite in Fig. 3a revealed that the lamellar phase was (Cr,Fe)7C3, while Mn was present around it in the form of segregation. Figure 3b shows that Cu precipitated in the ferrite grain boundary region, while granular CrxCy and MnxCy precipitated obviously in the ferrite grain. During the undercooling of austenite, C was captured by Cr and Mn to form carbides, and the formation of carbides reduced the diffusion coefficient of carbon in austenite and delayed the transformation of proeutectoid ferrite. The supercooled austenite region eventually formed ferrite.

Fig. 3: Segregation of corrosion-resistant elements in grain and cementite.
figure 3

a SAED analysis of adjacent grains in the pearlite. b EDS analysis of cementite and grain internal with corrosion-resistant elements.

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Characteristics of the structure

The segregation of elements at grain boundaries had an important effect on the corrosion characteristics of the materials in grain boundaries and the intragrain and pearlite regions. The microstructure of low-alloy steel Q500 was analyzed by EBSD, and the results are shown in Fig. 4. The IPF diagram in Fig. 4b shows that the grains of the material were ferrite, and the ferrite grains were randomly distributed along an equiaxed thermodynamic orientation. Figure 4a shows the BC diagram in the EBSD, where pearlite was present in the material; however, the content of Fe3C was relatively low, phase recognition was not obtained in Fig. 4c, and the pattern quality in the pearlite region was ambiguous. Combined with the KAM diagram in Fig. 4d, there were defects or high internal stresses in the pearlite region, and the stress distributions in the grain boundary and intragrain region were uniform. On this basis, the Volta potential characteristics of the grain boundary and pearlite region were studied by SKPFM. The results are shown in Fig. 5. According to the Volta potential characteristics in Fig. 5b, the cementite lamellae in the pearlite region had a high Volta potential and could be used as the cathode phase during corrosion. The ferrite in pearlite was preferentially corroded as the anode phase. The cementite as the cathode phase was the result of the interaction of Fe3C and element segregation, and there was stress concentration in the pearlite region, which promoted the corrosion process. The grain boundary region had a higher Volta potential relative to the grain boundary, and the intragranular ferrite dissolved during pearlite corrosion.

Fig. 4: EBSD patterns of microstructure of low-alloy steel Q500.
figure 4

a BC diagram, b IPF diagram, c Phase diagram, and d KAM diagram.

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Fig. 5: SKPFM characteristics of grain boundaries and pearlite regions.
figure 5

a Topography diagram and b Volta potential diagram.

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Segregation characteristics of elements in inclusions

In addition to the crystal defects, the corrosion-resistant elements Cu and Sb exhibited segregation in the boundary region between the inclusions and the matrix. The corrosion of low-alloy steel originated from the inclusion region. The study on the corrosion behavior of inclusions was helpful for studying the effects of corrosion-resistant elements Cu and Sb on local corrosion from corrosion initiation.

Different Cu-containing inclusions were selected, as shown in Fig. 6. The precipitation behaviors of Cu at the interfaces of inclusions were analyzed, and the segregation distributions of different alloy elements at the boundary defects of Cu-containing inclusions were observed. The cluster structures of alloy elements and boundary inclusions around Cu precipitates were obtained. The results of the Cu-containing inclusion test of Fig. 6a are shown in Fig. 7a. There were Cu-rich precipitates in the boundary of MnS, and CrxCy segregation occurred between the boundaries of Cu precipitates and the crystal. Cu interacted with Ni and Sb, and the segregation of Ni and Sb existed in the boundary region of Cu precipitates. The distribution of alloying elements was analyzed at the interface between Cu precipitates and the matrix. Figure 7b shows that Ni and Sb were enriched at the interface.

Fig. 6: Segregation characteristics of Cu in inclusions.
figure 6

a The precipitation behaviors of Cu at the TiN–MnS–MgO–Al2O3 inclusion. b The precipitation behaviors of Cu at the MnS–MgO–Al2O3 inclusion.

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Fig. 7: The element distribution characteristics of Cu-containing inclusion and Cu precipitation interface.
figure 7

a Distribution characteristics of corrosion-resistant elements at the boundary of inclusion and matrix. b Relative content of elements at the interface between the Cu precipitation phase and the matrix.

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The segregation characteristics of Ni and Sb at the precipitate boundary of Cu were analyzed, as shown in Fig. 8a, and the results are shown in Fig. 8b. Cu interacted with both Ni and Sb; however, there were segregation regions of Ni in the matrix, while the segregation behavior of Sb was always the same as that of Cu. There was no segregation region of Sb in the matrix, proving that Sb was more easily trapped by Cu. Myers and Follstaedt found that Sb was trapped by Cu precipitation27. The corrosion-resistant element Cu easily precipitated in the boundary region of the inclusion. The precipitation of Cu in the inclusion region was beneficial for reducing the precipitation of Cu at the grain boundary and within the grain and for improving the material properties. Sb was trapped by Cu and segregated at the boundary of the Cu precipitate phase and matrix. Some corrosion-resistant elements of Ni showed segregation behavior during the process, mainly within and at grain boundaries. Part of the Ni is segregated in the Cu precipitates, mainly within and at the grain boundaries. Cr either formed CrxCy in the grain boundary or segregated in the cementite.

Fig. 8: The characteristics of Ni and Sb capture by Cu in Cu-containing inclusions.
figure 8

a Distribution characteristics of corrosion-resistant elements at the boundary of inclusion and matrix. b Elemental content of Ni and Sb in different Cu precipitation phases and matrix interfaces.

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Local corrosion initiation behavior of Cu-containing inclusions

The precipitation behavior of Cu existed in the inclusion region. Figure 9 shows the distribution characteristics of Cu in all kinds of inclusions. The local corrosion of Cu-containing inclusions were studied by CaS/CaO–TiN–MgO–Al2O3 inclusions.

Fig. 9: Various inclusions containing Cu precipitates.
figure 9

a The precipitation behaviors of Cu at the CaS/CaO–MgO–Al2O3 inclusion. b The precipitation behaviors of Cu at the TiN–TiS2–MnS–MgO–Al2O3 inclusion. c Cu exists as a precipitated phase at the boundary of the compound inclusions.

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The corrosion initiation characteristics of Cu-containing inclusions were observed. CaS/CaO–TiN–MgO–Al2O3 inclusions were immersed in pure water for 10 s, and the results are shown in Fig. 10a. In the process of corrosion initiation, the matrix in the inclusion boundary region was partially dissolved; however, the TiN of the inclusion was not dissolved. Microcracks appeared at the interfaces between the inclusions and the matrix, and CaS partially dissolved and participated in the initiation of corrosion of the inclusions. After the surface of the inclusion was thinned, the observation of secondary polishing was carried out; the results showed that the corrosion characteristic products of the inclusion, as shown in Fig. 10b, were obtained. Cu was oxidized as an anode to form copper oxide during the local corrosion of inclusions.

Fig. 10: Corrosion morphology and products of CaS/CaO–MgO–Al2O3 inclusion with 10 s water immersion.
figure 10

a Before thinning and b after thinning.

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The result of the inclusion SKPFM test after immersion is shown in Fig. 11. At the early stage of inclusion corrosion, the anodic dissolving phase was the matrix, and this process was accompanied by the dissolution of part of CaS. The lower Volta potential values at the contact points between the inclusions and the matrix indicated that the local corrosion process was still expanding in the matrix region of the boundary of the inclusions. The dissolution of CaS occurred, and the oxide that formed in the inclusion region existed in the inclusion products.

Fig. 11: The morphology and Volta potential of the CaS/CaO-MgO-Al2O3 inclusion.
figure 11

a Dissolution profile; b the surface topography; c substance distribution; and d the surface potential.

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After the material analysis and SKPFM testing of Cu-containing inclusions as shown in Fig. 12a and b, the corrosion states of the inclusions were analyzed by immersion tests in 0.1 mol/L NaCl solution at different periods. The results are shown in Fig. 12c. The local corrosion behaviors of CaS/CaO–TiN–MgO–Al2O3 inclusions occurred, and the corrosion of inclusions happened preferentially in the gaps between the inclusions and the matrix. In the corrosion process, the corrosion preferentially occurred in the matrix structure defects, the corrosion pits gradually magnified, and the corrosion products gradually accumulated.

Fig. 12: The characteristics and in-situ corrosion behavior of CaS/CaO–MgO–Al2O3 Cu-containing inclusion.
figure 12

a SKPFM analysis of the CaS/CaO–MgO–Al2O3 inclusion. b EDS analysis of the CaS/CaO–MgO–Al2O3 inclusion. c Corrosion morphology of the CaS/CaO–MgO–Al2O3 inclusion after different times of immersion.

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The corrosion morphologies and products of Cu-containing inclusions at 1800 and 3600 s were analyzed, and the results are shown in Fig. 13. The inclusion region (I) and matrix boundary region (II) after 1800 s of immersion are shown in Fig. 13a. The inclusions were partially corroded and dissolved, and the oxide products of copper were found at the boundaries of the inclusions. After the 3600 s immersion test, the corrosion products gradually changed from cluster to lamellar. The results of EDS and Raman analysis showed that there were 265, 450, 624, and 1120 cm−1 peaks in the inclusion region. The peak positions of 265 and 1120 cm−1 corresponded to γ-Fe2O3 and α-FeOOH, respectively28,29, 450 cm−1 corresponded to Sb2O330,31,32, and 624 cm−1 corresponded to Cu2O. There were 518, 624, 1054, and 1335 cm−1 peaks in the Raman spectra of the corrosion products at the substrate corrosion boundary. The 1054 and 1335 cm−1 peaks corresponded to γ-FeOOH and Fe(OH)3, respectively, and 518 and 624 cm−1 both corresponded to Cu2O33,34. Cu and Sb contributed to the process of local corrosion of inclusions and produced corresponding oxides. The distribution characteristics of the elements within the inclusions were analyzed by TOF-SIMS, and the results are shown in Fig. 13b. CaS and CaO completely dissolved, and MgO and Al2O3 remained in the CaS/CaO–MgO–Al2O3 inclusion. Cu in the boundary region of the inclusion was retained as Cu2O in the corrosion products.

Fig. 13: Corrosion morphology and product analysis of Cu-containing inclusion with 0.1 mol/L NaCl immersion.
figure 13

a EDS and micro-Raman analysis of the inclusion after different times of immersion. b TOF-SIMS analysis of the CaS/CaO–MgO–Al2O3 inclusion after 3600 s immersion.

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Local corrosion development behavior of Cu-containing inclusions

The effects of the segregation of Cu, Ni, and Sb at the grain boundary and the segregation of Cr in the cementite lamellar region on the initial corrosion were analyzed. After 1800 and 3600 s immersion tests in 0.1 mol/l NaCl solution, the inclusions and boundary corrosion morphologies are shown in Fig. 14a and b. The local corrosion behavior originated in the inclusion region, and the grain boundary was extended during the corrosion process, during which the grain dissolved. Because of the segregation of Cu, Ni, Cr, and Sb at the grain boundary, the grain boundary acted as the cathode phase during the dissolution process. In addition, the ferrite around the cementite lamellae in the pearlite region dissolved, and the cementite lamellae remained, which was consistent with the results in Fig. 5.

Fig. 14: The corrosion morphology and distribution of corrosion-resistant elements in matrix and pearlite region of inclusion boundary.
figure 14

a EPMA analysis after 1800 s immersion and b EPMA analysis after 3600 s immersion.

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The distributions of Cu, Ni, Cr, and Sb in Fig. 14b show that Cu and Sb were enriched in the inclusion region. In combination with the test results shown in Fig. 13, Cu and Sb were distributed as products of Cu2O and Sb2O3. Cr and Mn were enriched in the cementite lamellae, which was consistent with the results shown in Fig. 3, and some Cr was enriched in the inclusions. In addition, no Ni enrichment was observed in the inclusion and matrix, which indicated that Cu was easier to precipitate and had a stronger Sb trapping effect. The Raman analyses of Cu, Cr, and Sb in the inclusion and matrix are shown in Fig. 15b. There were 624, 1054, and 1320 cm−1 peaks in the inclusions. The peak positions of 1054 and 1320 cm−1 corresponded to γ-FeOOH and Fe3O4, respectively, and the peak position of 624 cm−1 corresponded to Cu2O. There were 616, 696, and 884 cm−1 peaks in the cementite lamellar region. The peak positions of 616 and 696 cm−1 corresponded to Fe3O4 and Fe(OH)3, respectively, and the peak position of 884 cm−1 corresponded to CrO335,36,37,38. The corrosion products are analyzed by XPS analysis and the results are shown in Fig. 15c. Some of the results are consistent with Raman test results, Cu exists in the corrosion products as Cu2O and Cr exists in the corrosion products as CrO3. In addition, the oxidation product of Cr is also partly Cr2O3, while Sb exists in the corrosion product as Sb2O3. During the corrosion dissolution of the inclusion and matrix, Cu, Cr, and Sb contributed to the corrosion reaction and produced corresponding oxidation products. Cr existed in the form of Cr2O3 and CrO3 in the cementite lamella, and CrO3 had good stability in the dry environment. Sb existed in the corrosion product as Sb2O3 and Cu were oxidized with Cu2O.

Fig. 15: Corrosion morphology and corrosion products analysis of inclusion and pearlite region.
figure 15

a Microscopic corrosion morphology of the CaS/CaO–MgO–Al2O3 inclusion after 3600 s immersion. b EDS analysis and micro Raman analysis of the localized corrosion area after immersion. c XPS analysis of the yellow boxed area.

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Discussion

The segregation of Cu, Cr, Ni, and Sb in the grain boundary defect region of steel was the result of elastic interaction between solute atoms and Fe atoms, and the segregation process was spontaneous39. Cr segregated in the crystalline or cementite lamellar regions. The segregation of elements had a great effect on the microstructures and properties of low-alloy steel. Cu precipitation had a pinning effect on dislocations or grain boundaries, which increased the strength of the material. However, the grain boundary surface energy and the impact toughness were reduced40,41. Cu precipitation existed in the inclusion region, which significantly affected the properties and corrosion resistance of the material. The precipitation of Cu in the inclusion region was beneficial for reducing grain boundary segregation and improving the comprehensive properties of the material42. The precipitation of Cu in the inclusion region had a strong ability to capture Sb and a certain ability to capture Ni, and the Ni in the steel was still dominated by grain boundary and intragrain segregation43,44.

The distribution characteristics of Cu, Ni, Cr, and Sb affected the initial corrosion behavior of the steel. In the corrosive environment, local corrosion occurred preferentially in the inclusions. The segregation of Cu and Sb existed in the inclusion, and Cu and Sb participated in the dissolution reaction of the inclusion. The local corrosion and expansion processes of CaS/CaO–MgO–Al2O3 inclusions containing Cu and Sb are illustrated in Fig. 16. Corrosion was initiated in the CaS/CaO–MgO–Al2O3 inclusion with crack defects, and the matrix around the inclusion dissolved to form pits. In the corrosion dissolution stage of the matrix, CaS dissolution occurred, and part of Cu was oxidized to copper oxide. MgO–Al2O3 within the inclusions dissolved or shed gradually16,45.

Fig. 16: Schematic diagram of local corrosion of Cu-containing inclusion.
figure 16

a Characteristics of the distribution of alloying elements in the steel matrix; b localized corrosion occurs preferentially in the area of inclusions; c Cu and Sb are involved in the localized corrosion of inclusions; and d localized corrosion development and Cr in the cementite is involved in the corrosion process.

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The electrochemical reaction occurred in the local corrosion area, and the dissolution reaction of the matrix serving as the anode and the electrochemical dissolution reaction of CaS occurred. Cu and Sb were oxidized to Cu2O and Sb2O3, respectively. The electrochemical reaction equations were shown in Eqs. (1) and (2).

Anodic reaction:

$${{{\mathrm{Fe}}}} - {{{\mathrm{2e}}}}^ - \to {{{\mathrm{Fe}}}}^{2 + }$$
(1)

Cathode reaction:

$${{{\mathrm{O}}}}_2 + {{{\mathrm{H}}}}_2{{{\mathrm{O}}}} + 4{{{\mathrm{e}}}}^ - \to 4{{{\mathrm{OH}}}}^ -$$
(2)

In the process of anodic dissolution, the matrix was gradually transformed from thermodynamically unstable Fe(OH)2, Fe(OH)3, and γ-FeOOH to thermodynamically stable Fe3O4, α-FeOOH, and γ-Fe2O346.

$${{{\mathrm{Fe}}}}^{2 + } + 2{{{\mathrm{OH}}}}^ - \to {{{\mathrm{Fe}}}}\left( {{{{\mathrm{OH}}}}} \right)_2$$
(3)
$${{{\mathrm{Fe}}}}\left( {{{{\mathrm{OH}}}}} \right)_2 + 2{{{\mathrm{H}}}}_2{{{\mathrm{O}}}} + {{{\mathrm{O}}}}_2 \to 4{{{\mathrm{Fe}}}}\left( {{{{\mathrm{OH}}}}} \right)_3$$
(4)
$$6{{{\mathrm{Fe}}}}\left( {{{{\mathrm{OH}}}}} \right)_2 + {{{\mathrm{O}}}}_2 \to 2{{{\mathrm{Fe}}}}_3{{{\mathrm{O}}}}_4 + 6{{{\mathrm{H}}}}_2{{{\mathrm{O}}}}$$
(5)
$$6{{{\mathrm{Fe}}}}\left( {{{{\mathrm{OH}}}}} \right)_2 + {{{\mathrm{OH}}}}^ - \to \gamma - {{{\mathrm{FeOOH}}}} + 6{{{\mathrm{H}}}}_2{{{\mathrm{O + e}}}}^ -$$
(6)
$$\gamma - {{{\mathrm{FeOOH}}}} \to \alpha - {{{\mathrm{FeOOH}}}} \to \gamma - {{{\mathrm{Fe}}}}_2{{{\mathrm{O}}}}_3$$
(7)

CaS was electrochemically dissolved at the edge of the inclusion and gradually dissolved.

$$2{{{\mathrm{CaS}}}} + 6{{{\mathrm{OH}}}}^ - - 8{{{\mathrm{e}}}}^ - \to 2{{{\mathrm{Ca}}}}^{2 + } + {{{\mathrm{S}}}}_2{{{\mathrm{O}}}}_3^{2 - } + 3{{{\mathrm{H}}}}_2{{{\mathrm{O}}}}$$
(8)

Cu and Sb were involved in the reaction and oxidized to corresponding oxides34.

$${{{\mathrm{Cu}}}} + {{{\mathrm{OH}}}}^ - \to {{{\mathrm{CuOH}}}} + {{{\mathrm{e}}}}^ -$$
(9)
$$2{{{\mathrm{CuOH}}}} \to {{{\mathrm{Cu}}}}_2{{{\mathrm{O}}}} + {{{\mathrm{H}}}}_2{{{\mathrm{O}}}}$$
(10)
$$4{{{\mathrm{Sb}}}} + 6{{{\mathrm{OH}}}}^ - \to 2{{{\mathrm{Sb}}}}_2{{{\mathrm{O}}}}_3 + 3{{{\mathrm{H}}}}_2{{{\mathrm{O}}}}$$
(11)

The inclusion corroded the grain boundary of the matrix, and the interior of the crystal was preferentially corroded47. The ferrite crystals in the inclusion region exhibited intragranular corrosion propagation and epitaxy. The ferrite in the pearlite tissue acted as the cathode phase, and the ferrite within the pearlite preferentially dissolved and corroded48. Cr was enriched in the cementite lamellae, formed (Cr,Fe)7C3 and contributed to the corrosion process, resulting in CrO3. Sun et al.10 suggested that the product CrO3 would transition to Cr2O3.

$${{{\mathrm{Cr}}}} + 6{{{\mathrm{OH}}}}^ - \to {{{\mathrm{CrO}}}}_3 + 3{{{\mathrm{H}}}}_2{{{\mathrm{O}}}}$$
(12)
$${{{\mathrm{CrO}}}}_3 \to {{{\mathrm{Cr}}}}_2{{{\mathrm{O}}}}_3$$
(13)

Methods

Specimen preparation

The samples were cut into 10 mm*10 mm*2 mm and finely ground using 400#, 800#, 1200#, 2000#, and 3000# sandpaper in turn. The ground samples were mechanically polished with 0.5 μm metallographic polishing agent, and the surface of the samples was polished to a mirror surface without scratches, and finally, the surface of the samples was rinsed with deionized water and alcohol in turn. And the clean samples were dried and stored.

Materials characterization

The polished pearlite steel Q500 was etched with 4% alcohol nitrate, after rinsing and blowing dry with deionized water and alcohol, and its microstructure was observed by a VHX-2000 ultradepth of the field microscope system. The polished samples were finely polished with a suspension polishing solution and then further characterized by electron backscatter diffraction (EBSD) to analyze the structure of the matrix.

The microstructure and morphology of the matrix were characterized by a JEM-F200 field emission transmission electron microscope (FE-TEM). The selected region was analyzed by selected area electron diffraction (SAED), and the element distributions in the grain boundary and intragrain region were analyzed by energy dispersive spectrometry (EDS). The TEM samples were prepared by electrolytic double-spray technique with the solution of 90 vol% acetic acid and 10 vol% perchloric acid. The working voltage was 20 V and the temperature was 243 K.

The inclusions containing Cu on the surface of the material were observed by a GEMINISEM500 field emission scanning electron microscope (FE-SEM), and the element distribution of the inclusion was analyzed by energy-dispersive spectrometry (EDS). The microscope Vickers hardness tester was used to locate the target inclusions. The Cu-containing inclusions were cut by a focused ion beam (FIB), the distributions of elements at the boundary of Cu-containing inclusions were analyzed by CAMECA three-dimensional atomic probes (3DAP); the segregation distribution characteristics of different alloy elements at the boundary defects of inclusions were obtained. The cluster structures of alloy elements and boundary inclusions around Cu precipitates were obtained.

The local corrosion behaviors were observed by FE-SEM, and the e-local corrosion products were analyzed by EDS. ION-TOF GmbH time-of-flight secondary ion mass spectrometry (TOF-SIMS) was used to analyze the distribution of Cu at the inclusion boundary, determine the inclusions and analyze the local corrosion and internal dissolution characteristics of the inclusions. The size of the sputtering area was 15 µm*15 µm, the sputter ion beam was O2+, the potential was 2 keV, the incidence was 45°, and the sputter velocity was 1.52 nm/s.

The grain boundary, pearlite structure, and surface potential characteristics between the inclusions and matrix of low-alloy steel were measured by MULTIMODE8 scanning Kelvin probe force microscopy (SKPFM). The Kelvin probe model was PFQNE-AL. The SKPFM experiment adopted a single-channel method, and the scanning frequency was 0.5 Hz. The surface topography and the corresponding surface potential were measured simultaneously.

Corrosion testing

Local corrosion initiation and expansion of the inclusion and matrix were observed in situ. The local corrosion characteristics of the inclusions were observed after 10 s of immersion in H2O, and the inclusions were polished and thinned with 3000 # SiC sandpaper. Local corrosion and corrosion propagation of inclusions were investigated by immersion tests in 0.1 mol/L NaCl solution at 25 °C. The soaking times were 0, 60, 1800, and 3600 s. A VK-X250 confocal laser scanning microscope (CLSM) was used to analyze the depths and morphologies of the inclusions. EPMA-1720H electron probe microanalysis (EPMA) was used to study the distribution characteristics of corrosion-resistant elements in the local corrosion region, and the products were analyzed by a LabRAM HR Evolution high-resolution confocal micro-Raman spectrometer. The laser wavelength was 532 nm. In addition, the local corrosion areas of the inclusions were analyzed by X-ray photoelectron spectroscopy (XPS) for the present patterns of Fe, Sb, Cr, and Cu elements. The X-ray source was an Al target with a spot diameter of 15 µm and a power of 150 W. The operating voltage and current were 14.8 kV and 1.6 A, respectively.