Effect of Nano-sized Cu-Rich Phase on Microbiological Corrosion Behavior of Cu and Ni-Added Steel for Oil Country Tubular Goods

Abstract

The Cu-bearing antibacterial low-alloyed steel pipe is a new strategy to mitigate microbiologically influenced corrosion (MIC) for oil and gas industry. It can effectively alleviate the occurrence of bacterial corrosion, but cannot avoid the MIC totally. To enhance the resistance to MIC of Cu-bearing low-alloyed steel, the MIC behavior of a Cu and Ni-added experimental steel with different nano-sized Cu-rich precipitates was investigated. Results showed that the synergistic effect of antibacterial ability of nano-sized Cu-rich precipitates and protective Cu–Ni enrichment layer formed on the steel surface contributed to the good MIC resistance. Although the Cu-rich precipitates possessed antibacterial actively, they also increased the surface potential difference simultaneously, resulting in promoting MIC. The extremely fine and dispersed Cu-rich precipitates with high-density was the preferred microstructure to achieve better MIC resistance for the steel.

1. Introduction

Nowadays, fossil fuels still account for a significant proportion of all energy sources, among which oil and gas are the main energies in the world today. With the continuous exploitation of oil fields, most of them are already in the middle or later stage. Accordingly, the water content of the produced fluids from oil wells is constantly increasing. After filtering, the water from the produced fluids can be used for the secondary oil recovery by waterflooding. It not only saves a large amount of clean water but also expels crude oil, and prevents environmental pollution from sewage discharge at the same time. However, the reinjection of water is usually a habitat for microorganisms, which can cause microbiologically influenced corrosion (MIC) to oil and gas pipelines during their growth, reproduction, and metabolism processes.¹⁾ In recent years, the pipeline failure cases due to MIC are increasing.^2,3,4,5,6) Except for literature reports, on-site cases have also occurred. For example, in a newly drilled oil field in Southwestern China, multiple wells were perforated due to MIC of reinjected water in just six months. Therefore, the attention to MIC of pipelines caused by the reinjected water from oilfield has been paid.

To solve the MIC trouble of pipelines used in oil/gas exploitation and transportation, various mitigation countermeasures such as physical and chemical approach,⁷⁾ protective coatings,^8,9) biological competition,^10,11) and antibacterial steel pipe have been developed.^{12,13,14,15,16,17,18,19,20,21,22)} Among them, antibacterial steel pipe is a new strategy to mitigate MIC from the aspect of material itself, having the characteristics of long-term durability.^12,13,14) By proper Cu alloying design to the traditional pipeline steels, aiming at continuous release of Cu ions to kill the bacteria and inhibit the formation of bacterial biofilm, a creative strategy for improving the MIC resistance of pipeline steels has been proposed.¹⁵⁾ Yu et al.^16,17) reported that a Cu-added L245NCS steel with 0.6 wt.% Cu content exhibited strong antibacterial activity against sulfate reducing bacteria (SRB). These results were not only confirmed under the laboratory condition but also in the shale gas field environment. More importantly, the oil country tubular goods (OCTGs) manufactured by Cu-added L245NCS steel has achieved good MIC resistance when applied to a certain oil field in Western China. Huang and Li et al.^20,21) proved that a new Cu-added X70 grade anti-bacterial corrosion high frequency welding (HFW) pipe could also effectively inhibit the SRB corrosion, and ensure the safe service of pipes.

Practice has proved that the Cu-added OCTGs could effectively alleviate the bacterial corrosion. However, they did not avoid MIC totally. In our previous research, it was found that although the depth and density of pits induced by SRB was reduced significantly compared to the Cu-free steel, there was still shallow pitting corrosion on the surface of Cu-bearing steel.^14,19) What’s more, the strong or weak to MIC resistance of Cu-added steel was not only related to the Cu content, but also closely related to the existing form of Cu in the steel.²³⁾ It is believed that the presence of Cu in the form of precipitation in steel was more beneficial to reduce the harm of MIC.¹⁴⁾ Thus, it is reasonable to believe that the morphology characters, including size and number density, of Cu-rich precipitates in steel will also have a significant impact on its MIC resistance. This is because there are often micro-areas with uneven physical or chemical properties between the precipitates and the matrix, where provide the initiation sites for the corrosion. The different morphology characters of Cu-rich phase in a same Cu-added steel may lead to a difference in MIC resistance. Therefore, the work of the correlation between the nano-sized Cu-rich phases and MIC behavior is of great significance for optimizing the MIC resistance of Cu-added low-alloyed steel.

In general, the Cu-bearing low-alloyed steel requires the addition of appropriate amount of Ni to avoid “hot shortness” during hot working.^24,25) In this work, a Cu-added OCTG-purpose steel with 1.35 wt.% Cu and 1.05 wt.% Ni was fabricated. To obtain different sizes and quantity densities of nano-sized Cu-rich phases in the steel, different tempering temperatures were selected. The influence of the changes in Cu-rich phases generated after different tempering temperatures on the MIC resistance of OCTG steel was studied, and the possible mechanism of Cu and Ni improving the MIC resistance of the steel was proposed. This study can provide a guidance for the optimization of MIC resistance of Cu-added OCTG-purpose steel.

2. Material and Methods

2.1. Materials

The Cu-added OCTG-purpose steel with 1.35 wt.% Cu and 1.05 wt.% Ni was melted in a 180 kg vacuum induction melting furnace. For comparison, a Cu-free with 1.05 wt.% Ni steel and an X65 steel were used. Their compositions are listed in Table 1. The Cu-added and Cu-free steel ingots were forged into blocks at first, and then austenitized at 1200°C for 2 hours. Seven hot rolling steps were performed at a similar temperature to reduce the thickness of steels from 90 mm to 12 mm. After final rolling, a fast cooling rate was achieved by water spray. To investigate the effect of morphology characters of Cu-rich phase on the MIC behavior, the Cu-added hot-rolled steel was tempered at 500, 550 and 600°C for 2 hours, respectively. For convenience, the experimental materials were named as 500Cu, 550Cu, 600Cu, Cu-free and X65 steel.

Table 1. Chemical compositions of the experimental steels (wt.%).

Steels	C	Mn	Cu	Ni	Nb	Ti	Si	Al	P	S	Fe
Cu‑added steel	0.029	0.59	1.35	1.05	0.019	0.017	0.19	0.024	0.009	0.0021	Bal.
Cu-free steel	0.028	0.59	<0.01	1.05	0.018	0.017	0.18	0.026	0.008	0.0020	Bal.
X65 steel	0.060	1.64	<0.01	0.09	0.040	0.013	0.14	0.028	0.010	0.0010	Bal.

2.2. Methods

2.2.1. Bacteria and Solution

The SRB seed used in this study was isolated from the produced water in an oil field in the southwest of China. The bacterial culture medium was API RP-38, which was the same as the previous work.¹⁹⁾ The pH of the medium was adjusted to 7.0–7.2 by 5 mol/L NaOH. After that, a high-purity nitrogen was injected into the medium for 2 hours to obtain an oxygen-free environment. Then the SRB-inoculated API RP-38 medium was stored in a 4°C icebox for standby. To simplify the work, the corrosion solution used in the study was near-neutral, which contains of 0.122 g/L KCl, 0.483 g/L NaHCO₃, 0.181 g/L CaCl₂·2H₂O and 0.131 g/L MgSO₄·7H₂O. After the solution preparation, it was deoxygenated with mixed gas including 5%CO₂+95%N₂ for 2 hours, and then sterilized by autoclaving at 121°C for 20 min. Prior to tests, the SRB strains were activated in an incubator for 12 hours. After that, 5% of SRB-inoculated API RP-38 medium mixed with 95% of sterilized corrosion solution was used.

2.2.2. Weight Loss Test

Weight loss was measured for the specimens after immersion for 14 and 28 days by a balance with accuracy of 0.0001 g. Before the test, three specimens for each group were ground to 800# with sandpapers, cleaned in absolute ethanol and dried in a hot air stream. After immersion, attached biofilms and corrosion products were removed from surface of the specimens using a de-rusting solution containing 500 mL hydrochloric acid (HCl) + 500 mL deionized water + 20 g hexamethylenetetramine, which were then cleaned with absolute ethanol. The corrosion rate was calculated by Eq. (1) as follows:

  

$$v=87\mathrm{\hspace{0.17em} }600\times \frac{W_0-W_t}{\rho \text{St}}$$(1)

where v is the corrosion rate in mm/a; W₀−W_t is the weight loss in g; S is the exposed surface area in cm²; t is the immersion time in hours; and ρ is the steel density 7.85 g/cm³ in the present study.

2.2.3. Biofilm and Corrosion Product Analysis

The biofilm and corrosion product morphologies on surfaces of different steels were examined by confocal laser scanning microscope (CLSM) and scanning electron microscopy (SEM), respectively. The same experimental methods as the previous study¹⁹⁾ were used, including biofilm staining test, biofilm thickness, elements distribution on the surface and cross-section of the corrosion products, corroded surfaces observation, as well as pit morphologies.

The live/dead BacLight bacterial viability kit was used to differentiate the dead and live cells. Before being stained, the surface of the sample was gently washed with phosphate buffer saline (PBS) to remove the loosely adhered bacteria. Then the dye was extracted and dropped onto the sample in a dark environment for 20 min. After that, when observed under a CLSM, the live cells show a green fluorescence, while complete dead cells show red fluorescence. Finally, the 15 NIS Viewer software was used to reconstruct three-dimensional images, and at least 3 images were used to perform quantitative analysis by ImageJ software.²⁶⁾

Besides, X-ray diffraction (XRD, Rigaku D/Max 2500PC) was used to analyze the crystal structure of corrosion products. The XRD was performed with Cu–Kα radiation (λ = 0.154 nm) operating at 45 kV and 200 mA. The XRD patterns were collected within the 2θ range of 10° to 90° at a scanning speed of 10°/min.

2.2.4. Surface Potential Measurement

Scanning Kelvin probe force microscopy (SKPFM) was conducted using atomic force microscope (AFM), and mappings were performed to evaluate the surface potential difference between Cu-rich phase and the matrix.²⁶⁾ The samples tested by SKPFM were Cu-added steel samples that had undergone the aging treatment at 500, 550, and 600°C for 2 hours, respectively. To observe the nano-sized Cu-rich phases clearly, the samples were carefully mechanically ground to #2000 with sand papers, polished, and etched in a 2% nital solution. Then the surface potential was measured in the dual-scan mode. For the first line scan, topography images were obtained using the tapping mode, while Volta potential data were measured in the second line scan. During the operation, all of the scans were carried out at 25°C and humidity of 40% to 65% with a scan frequency rate of 0.3 Hz.

3. Results

3.1. Nano-sized Cu-rich Phases

Figure 1 shows the micrographs of nano-sized Cu-rich phases in the Cu-added steel after tempering. As the tempering temperature increased, the size of Cu-rich phase increased but the number density decreased. According to the results of statistics (Fig. 2), the average size of Cu-rich phase tempered at 500, 550 and 600°C are 5.4, 7.7 and 10.1 nm, and the number density are 8.02×10¹⁴, 5.68×10¹⁴, 2.49×10¹⁴ No. /m², respectively. As expected, the size and number density of Cu-rich phases in the Cu-added steel changed with the tempering temperature.

Fig. 1. Morphologies of Cu precipitates and the corresponding TEM-mapping images in Cu-added steel tempered at different temperatures for 2h: (a) 500°C; (b) mapping images of (a); (c) 550°C; (d) 600°C. (Online version in color.)

Fig. 2. Size distributions and number density of Cu precipitates in Cu-added steel tempered at different temperatures for 2 h: (a) 500°C; (b) 550°C; (c) 600°C. (Online version in color.)

3.2. Corrosion Rate

Corrosion rate was obtained from weight loss based on the totally exposed surface area and time of immersion. Figure 3 shows the corrosion rate of different experimental steels after immersion for 14 and 28 days. During the 14 days immersion, 600Cu and X65 steels had the highest corrosion rate, which was approximately the same. The corrosion rate of Cu-free steel was the second, 550Cu steel was the third, and 500Cu steel was the last. After immersion for 28 days, the corrosion rates of Cu-added steel with different tempering temperatures were lower than those of Cu-free and X65 steels. The lower corrosion rate of Cu-added steel demonstrated a better corrosion resistance in the bacterial environment. For the Cu-added steel, its corrosion rate increased with increasing tempering temperature, with the lowest corrosion rate at tempering temperature of 500°C. Besides, the average corrosion rates of all the experimental steels were higher during the first 14 days than the 28 days. This indicates that the accelerated corrosion mainly occurred within the first 14 days.

Fig. 3. Average corrosion rate of different experimental steels after exposed in SRB-inoculated medium for 14 and 28 days. (Online version in color.)

3.3. Biofilm and Corrosion Product

Figures 4(A)–4(E) show the live/dead stained biofilms measured by CLSM. The difference in biofilm thickness among the five steel samples was small within 14 days immersion. The maximum biofilm thickness was 28 μm for X65 steel, and the minimum one was 16 μm for 600Cu steel. However, there was a significant difference in the number of live/dead cells on the steel surfaces. Large green areas (live cells) and nearly no red areas (dead cells) were found on the X65 and Cu-free steels surfaces, while there were only approximately half of dead cells were found on the Cu-added steel surfaces. Figure 5 shows the statistical results of the live/dead cell number on the steel surfaces. The number of dead cells on the Cu-added steel surfaces (1.8~5.0×10⁵ cells/cm²) was much more than those on X65 steel (6.0×10² cells/cm²) and Cu-free steel (4.6×10² cells/cm²). The much more dead cells on the Cu-added steel surfaces demonstrated a better antibacterial property. For the Cu-added steel, the number of dead cells was in the same order of magnitude (10⁵ cells/cm²), showing no much difference. This indicates that the variations in size and number density of Cu-rich phase after tempering at different temperatures on the antibacterial property were negligible.

Fig. 4. Live/dead stained morphologies of X65 (A), Cu-free (B), 500Cu (C), 550Cu (D) and 600Cu (E) steels after 14 days of immersion. Morphologies of corrosion products on surfaces of X65 (a), Cu-free (b), 500Cu (c), 550Cu (d) and 600Cu (e) steels after 28 days of immersion. (Online version in color.)

Fig. 5. Live/dead sessile cells counts on the surfaces of different steels after exposed in SRB-inoculated medium for 14 days. (Online version in color.)

Figures 4(a)–4(e) show the corrosion products measured by SEM after 28 days of immersion in the SRB-inoculated medium. A large amount of biofilm and white corrosion products covered on the surfaces of all the experimental steels. Worm-shaped bacteria were widely attached to the products on the surfaces of X65 and Cu-free steels, while it was not obvious on different Cu-added steel samples. EDS detection was carried out on the positions marked by red circle in Figs. 4(a)–4(e). The results in Table 2 show that types of the elements were basically same but Cu and Ni. The Cu-added steel showed the existence of Cu and Ni in the corrosion products, but not much difference in the concentration was found.

Table 2. EDS results of different positions on sample surfaces after exposed in SRB inoculated medium for 28 days (wt.%).

Position	C	O	Si	Na	Fe	Mn	P	S	Cu	Ni
1	0.62	5.86	0.34	0.84	84.11	1.28	–	6.94	–	–
2	0.47	4.81	–	0.69	84.58	0.42	–	8.23	–	0.79
3	0.62	8.30	0.36	–	78.63	0.49	0.31	8.81	1.47	1.02
4	0.62	4.72	0.26	–	85.80	0.47	–	5.89	1.43	0.81
5	0.59	3.52	0.30	–	85.97	0.51	–	7.14	1.20	0.77

XRD was used to identify compositions of the corrosion products on different steel surfaces. The corrosion products of all the steels were mainly composed of iron sulfide. As shown in Fig. 6, except for the Fe in matrix, Fe was mainly in the form of FeS. However, no Cu and Ni were found in the corrosion products.

Fig. 6. XRD patterns of the corrosion products on different steel surfaces after 28 days exposure in the SRB inoculated medium. (Online version in color.)

To further confirm the role of Cu and Ni in steels, X65 and Cu-free as well as 500Cu steels were chosen to conduct the SEM mapping along the cross-section of corrosion products, and the results are shown in Fig. 7. The corrosion product layer was located between the matrix and epoxy, as remarked by the yellow dotted lines. The distribution of those important elements as Cu, Ni, S and Fe in corrosion products is clearly present. S was found to widely distribute in the corrosion product layer on the three steels, indicating that the main corrosion products were composed of sulfides. Unlike X65 steel, it should be noticed that the segregation of Ni could be seen in the corrosion product on Cu-free steel. Besides, pit was coincidentally found, corresponding to the yellow and blue areas in Fig. 7(b) showing high contents of S and Ni. It is indicated that the S and Ni compounds might form near the matrix. Except for the Ni-rich layer in corrosion product, a continuous and Cu-enhanced Cu/Ni-rich layer was observed on 500Cu steel. The thin Cu/Ni-rich layer was just nest to the matrix, as shown in Fig. 7(c). It is evident that it could prevent the further corrosion.

Fig. 7. Cross-sectional morphologies and EDS mapping results of different steels after 28 days immersion in the SRB-inoculated medium: (a) X65 steel; (b) Cu-free steel; (c) 500Cu steel. (Online version in color.)

EPMA with higher resolution was used to verify the reliability of the above results. Figure 8 shows the elements distribution along the cross-section of corrosion products on Cu-free steel and 550Cu steel. As same as the results of SEM analysis, the Ni-rich (Fig. 8(a)) and Cu/Ni-rich (Fig. 8(b)) but poor Fe layers located at the interface between the matrix and the biofilm/corrosion product were obviously observed. Another point should be noticed that the layer was not only quite rich in Cu and Ni but also rich in S on 550Cu steel.

Fig. 8. EPMA mapping results of the cross-section of corrosion products on Cu-free steel (a) and 550Cu steel (b) after 28 days immersion in the SRB inoculated medium. (Online version in color.)

3.4. Pitting Corrosion

After removing the biofilm and corrosion products, the corroded surfaces were observed by SEM. Figures 9(A)–9(E) show the images of surface morphologies of five steel samples after 28 days immersion. It should be noticed that uniform corrosion almost did not occur on all steels because the original scratches were still clear. However, there were large amounts of aggressive pits appeared on the surfaces of all steels. It was extremely terrible for both X65 and Cu-free steels. Pits on two steels were large-sized and clustered in groups. Although the addition of Cu did not make the steel to avoid pitting corrosion, it was still alleviated compared with the Cu-free steels.

Fig. 9. Surface corrosion morphologies (A–E) and pit equivalent diameters (a–e) of different steels after immersion for 28 days: (A, a) X65, (B, b) Cu-free, (C, c) 500Cu, (D, d) 550Cu and (E, e) 600Cu. (Online version in color.)

To evaluate the severity of pitting corrosion on different steels, pit depths and equivalent diameters on two random fields were also counted.¹⁹⁾ The results are shown in Figs. 9(a)–9(e) and Fig. 10. The maximum equivalent diameter of pits on X65 and Cu-free steels could reach to 10 μm, which was almost triple that of Cu-added steels (~4 μm). It should be noticed that more than 80% of the pits on surfaces of Cu-added steels had an equivalent diameter less than 2 μm. The pitting profiles imaged by CLSM (Fig. 10) showed that the deepest pits were 25.0 μm, 22.4 μm, 7.8 μm, 10.2 μm and 8.9 μm for X65, Cu-free, 500Cu, 550Cu and 600Cu steels, respectively. Obviously, the deepest pit on those Cu-free steels was much deeper than that on the Cu-added steel. Thus in terms of pitting corrosion resistance, the Cu-added steel performed better than Cu-free steels. For the Cu-added steel, 500Cu showed the best corrosion resistance in the presence of SRB.

Fig. 10. 3D images of pit morphologies on the surfaces of different steels after immersion for 28 days: (A, a) X65, (B, b) Cu-free, (C, c) 500Cu, (D, d) 550Cu and (E, e) 600Cu. (Online version in color.)

4. Discussion

4.1. Corrosion Against SRB of Cu/Ni-added Steel

SRB can lead to severe pitting corrosion of OCTGs. In the present work, MIC behaviors of Cu/Ni-added OCTG-purpose steel with different nano-sized Cu-rich precipitates were investigated in the presence of SRB. From the observation of corrosion morphologies shown in Fig. 9, it can be noticed that all steels suffered from pitting corrosions. Although the corrosion was not avoided totally by addition of Cu and Ni in the steel, it seemed that the pitting corrosion occurred more easily and severely on Cu-free steels. Cu is a desirable element in the antibacterial metal materials, which has been confirmed from many studies.^27,28,29,30) The antibacterial mechanism was attributed to the “contact killing” of Cu which possesses a high toxicity to bacteria.^{26,31,32,33,34,35)} However, the Cu-bearing stainless steels could not well mitigate MIC against SRB. Liu et al.³⁶⁾ claimed that the Cu precipitates on the steel would react with sulfides produced by SRB metabolizing and form the copper sulfide on the steel surface, thus, losing its toxicity to SRB. The result is different from our present and previous studies.^12,13,14,19) Yu et al.^16,17) and Huang et al.²⁰⁾ also reported that the Cu-bearing carbon steel exhibited strong antibacterial activity against SRB. This is also confirmed recently by study of Fan et al.²²⁾ Their results showed that MIC resistance of the Cu-bearing pipeline steel was better than those of X70 and Cr-bearing steels. Although the direct reason on difference of corrosion inhibition effect was not found, the lower solid solubility of Cu in ferritic steel led to form more and finer Cu-rich phases in the matrix may be one of the reasons.

Besides, the role of Cu in promoting the formation of corrosion product layers cannot be excluded in inhibiting bacterial activity of the Cu-added steel.³⁷⁾ Results presented by Kim et al.³⁸⁾ revealed that the sulfide layer formed on the surface of Cu and Ni-bearing steel showed lower corrosion rate and higher resistance to hydrogen induced cracking than the steel containing only smaller amount of Ni. Another research conducted by Choi et al.³⁹⁾ also showed that Cu and Cr compounds would be formed in the corrosion product layer on the weathering steel, which acted as a factor for the corrosion resistance in aqueous solutions. In the present study, a continuous and dense Cu- and Ni-rich layer was also observed (Figs. 7(c) and 8(b)). For the Cu-free steel (Ni-added steel only), the Fe and Ni would react with sulfides produced by SRB metabolizing and form FeS and NiS. Due to the solubility product constant (Ksp) of NiS is smaller than that of FeS,⁴⁰⁾ NiS is more likely to precipitate before FeS. So that the Ni-rich layer accumulated near the substrate (Fig. 7(b)). In addition, the kinetics of water exchange rate of Ni²⁺ is two orders of magnitude lower than that of Fe²⁺,⁴¹⁾ and the amount of Fe²⁺ released from the matrix is much more than that of Ni²⁺. Therefore, Ni²⁺ tended to attached onto FeS to form (Fe, Ni)S (Fig. 8(a)). This result was confirmed by Mansor et al.⁴²⁾ They found that as increasing the ratio of [Fe²⁺]aq/[Ni²⁺]aq, Ni-sulfides became rarer, and finally Ni-rich mackinawite (FeS) became the primary component when the ratio of [Fe²⁺]aq/[Ni²⁺]aq exceeded five. So, in the present study, only FeS and no NiS were found in the XRD analysis (Fig. 6). More importantly, the doping of Ni in FeS could also promote the growth of FeS and improve its thermodynamic stability,^43,44) which is also beneficial for the corrosion resistance. Thus, the corrosion resistance of Ni-added steel (Cu-free steel) performed better than X65 steel (Fig. 3). As mentioned above, a high content of sulfur was detected by EDS for all steels in the presence of SRB (Table 2), indicating that the main corrosion products were composed of sulfides. In addition to FeS and NiS for the Cu- and Ni-added steel, Cu₂S would also be formed.³⁶⁾ Due to the smallest Ksp of Cu₂S among the three sulfides in the present study,⁴⁰⁾ Cu would firstly form Cu₂S, so the Cu-Ni-rich layer tightly bonded to substrate was observed in Fig. 7(c). Kolthoff et al.⁴⁵⁾ and Baba et al.⁴⁰⁾ reported that Cu₂S could promote the FeS co-precipitation, forming a denser corrosion product layer on the surface of the substrate. The denser the corrosion product layer, the stronger its ability to suppress corrosion.⁴⁶⁾ Thus, the Cu-added steel showed good corrosion resistance in the presence of SRB.

Based on the experimental results and analyses, a proposed corrosion mechanism is schematically illustrated in Fig. 11. In the early stage, the biofilms would be formed on the surfaces of the steels. To maintain the energy for reproduction, sessile cells underneath the biofilms would obtain electrons, accordingly the pitting corrosion would appear. For the Cu-added steel, the nano-sized Cu precipitates on surface of the steel matrix would reduce the number of live bacteria by “contact killing” (Figs. 4 and 5). During the late stage, the steel surface accumulated a large amount of corrosion products (Fig. 4). The FeS layer on the X65 steel could not effectively prevent the matrix from further corrosion (Fig. 11(a)). Although the Ni-rich layer would be formed on the Cu-free steel (Fig. 11(b)), it was still unable to resist the attack of SRB. Only the Cu/Ni-rich layer formed on the Cu-added steel could alleviate the pitting corrosion induced by SRB. But it still did not avoid it completely (Fig. 11(c)).

Fig. 11. Schematic illustration on corrosion mechanism of different steels immersed in SRB-inoculated medium: (a) X65, (b) Cu-free, (c) Cu-added. (Online version in color.)

4.2. Effect of Nano-sized Cu-rich Phases on SRB Corrosion Resistance

According to morphologies of the pitting corrosion shown above, pitting corrosion should mostly occur on the Cu-added steel with negligible uniform corrosion. The 500Cu steel showed the minimum pit depth and pit density. The main reason for such best corrosion resistance to SRB was attributed to the difference of nano-sized Cu precipitates (Fig. 1), resulting in two different corrosion results. One is the antibacterial property, and the other is the corrosion potential. In general, the finer the Cu-rich precipitates, the larger the number density of precipitates, and the greater the probability of contacting and killing bacteria. But the dead cell number was not much difference for the Cu-added steel with different tempering treatments (Fig. 5). Namely, the antibacterial ability of three different nano-sized Cu-rich precipitates was almost similar. Thus, the antibacterial property of 500Cu steel did not enhance its SRB corrosion resistance.

Precipitates in steel are usually the original location of corrosion. They often have different potentials with the steel matrix, thus easily forming a corrosion couple and then accelerating the corrosion process. But would the abiotic factor affect the biological corrosion? To clarify the best corrosion resistance against SRB for 500Cu steel, the difference in surface potential between the Cu-rich phase and the matrix was measured by SKPFM. Figure 12 shows the topographical and surface potential images as well as the corresponding linear profiles of surface potentials of differently treated Cu-added steels. It can be can found that the difference in surface potential between the Cu-rich phase and the matrix was ~3 mV, ~5 mV and ~10 mV for 500Cu, 550Cu and 600Cu steels, respectively. Namely, the potential difference gradually increased with increasing the tempering temperature. This indicates that larger size with low density Cu-rich phases would enhance the surface potential difference. While the higher surface potential difference implied a higher charge transfer potential.⁴⁷⁾ The metabolism of SRB is mediated by the transfer of electrons.^48,49,50) Therefore, a higher charge transfer potential can undoubtedly make this metabolism process easier. It is reasonable to believe that the higher surface potential difference would promote the MIC process. Accordingly, the 600Cu steel had the largest corrosion rate and pit density. In short, the nano-sized Cu-rich phase is the prerequisite for safeguarding the antibacterial activity. Meanwhile, the surface potential difference caused by Cu-rich phase and matrix could also promote the MIC. That is, it is very important to minimize the surface potential difference as much as possible but not impair the antibacterial property. Therefore, the uniformly high-density distributed Cu-rich precipitates with the smallest size of 500Cu steel possessed the best MIC resistance against SRB.

Fig. 12. Topographical (a, d, g), surface potential images (b, e, h) and cross-sectional profiles of surface potential (c, f, i) generated along the red line in the topographical images of Cu-added steels: (a, b, c) 500Cu, (d, e, f) 550Cu, (g, h, i) 600Cu. The brighter the image, the higher the position/potential on the sample surface. Red arrows in topographical images indicate the Cu-rich phase. (Online version in color.)

5. Conclusion

In this study, the MIC behavior of a Cu and Ni-added OCTG-purpose steel with different nano-sized Cu-rich precipitates was investigated. Although the addition of Cu and Ni still did not make the steel to avoid MIC completely, the results showed that the synergistic effect of antibacterial ability of nano-sized Cu-rich precipitates and protective Cu-Ni enrichment layer formed on the steel surface played a very important role in improving the MIC resistance. Besides, the different morphology characters of Cu-rich phase in the Cu-added steel resulted in the difference of MIC resistance. The Cu-rich precipitates possessed antibacterial actively, however, they also increased the surface potential difference simultaneously, resulting in promoting MIC. Therefore, the 500Cu steel with extremely fine and dispersed as well as high-density Cu-rich precipitates achieved the best MIC resistance. This study provides valuable guidance for the optimization of MIC resistance of Cu-bearing low-alloyed steel.

Acknowledgments

This work was supported financially by the National Natural Science Foundation of China (Grant No. 52201093), Project of Liaoning Marine Economic Development (Development of high strength pipeline steel for submarine oil and gas transmission) and State Key Laboratory of Metal Material for Marine Equipment and Application Funding (Grant No. SKLMEA-K202205).

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