Beneficial effects of CeO2 addition on microstructure and corrosion behavior of electrodeposited Ni nanocrystalline coatings

Surface & Coatings Technology 235 (2013) 433–446

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Beneficial effects of CeO2 addition on microstructure and corrosion behavior of electrodeposited Ni nanocrystalline coatings Xiaowei Zhou, Yifu Shen ⁎ College of Materials Science and Technology, Nanjing University of Aeronautics and Astronautics, Nanjing 210016, PR China

a r t i c l e

i n f o

Article history: Received 18 May 2013 Accepted in revised form 31 July 2013 Available online 8 August 2013 Keywords: Nanocrystalline coatings Corrosion behavior Microstructure CeO2 adsorption

a b s t r a c t This paper presents an experimental study of the effects of CeO2 addition on the microstructure and corrosion behavior of nanocrystalline Ni coatings. A novel method of ultrasound-assisted pulsed electrodeposition was used to prepare the coatings from a modified Watts-nickel electrolyte. Surface morphologies and microstructural evolution of the coatings before and after aging treatments were characterized using E-SEM and TEM observations. Furthermore, electrochemical techniques attached with the XRD and XPS analysis were applied to study their corrosion behavior in a 5 wt.% HCl aqueous solution. The underlying mechanisms of corrosion inhibition resulted from the Ce-rich passive layer were discussed in detail. Experimental results indicated that the existence of well-distributed CeO2 phase in the coatings promoted dispersion strengthening through a proper aging treatment at 650 °C for 4 h, thereby achieving an effective bonding link between interfacial boundaries for denser microstructure. According to the measured results from polarization curves and the AC-impedance measurements, it exhibited superior corrosion resistance of the Ni-CeO2 coatings as compared to pure nickel. During the long-term static immersion tests in an acid 5 wt.% HCl aqueous solution, a small amount of Ce3+/4+ ions was released from the dissolution–precipitation of a Ce-rich phase and served as the corrosion inhibitors. Furthermore, some insoluble corrosion products resulted from the process of electrochemical reactions were favorable for precluding the corroded surface from the corrosive attack and diffusion behavior by Cl− ions. © 2013 Elsevier B.V. All rights reserved.

1. Introduction In order to meet the ever-increasing demands for high-performance coatings applied in the harsh conditions, many attempts (e.g., plasma spraying [1], laser cladding [2], surface mechanical alloying [3], etc) have been carried out for surface modification during the past decades. Among them, the traditional electroplating technique, which is considered as one of the most effective methods for preparing nanocrystalline coatings, has been extensively employed to prepare Ni nanocomposite coatings used in a wide range of engineering applications in aerospace engines, petroleum pipelines, gas-turbines, medical industries, etc. [4,5]. Well documented that the properties of such composite coatings are closely correlated with the degree of reinforcing phase. More recently, nanoparticle reinforced Ni-based coatings have attracted significant interest and received intensive study of their potential applications in the surface modification of corrosion protection and thermal barrier coatings [6]. However, due to the mismatch with the coefficients of thermal expansion between Ni grain and its reinforcing phase, they are prone to take place cracking boundaries, thermal micro cracks, coarsening growth or inter-crystalline oxidation at the coating surface where subjected to a severe environment [7]. Based on above shortages,

⁎ Corresponding author. Tel.: +86 25 84895940; fax: +86 25 84896170. E-mail address: [email protected] (Y. Shen). 0257-8972/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.surfcoat.2013.07.070

we should search an available reinforcing phase for the Ni-based composite coatings. Since the beneficial effects of rare-earth (RE) elements or RE-oxides were primarily applied in the field of surface modification, various researchers have devoted them to the surface coatings for the improvements in micro-hardness, oxidation and corrosion resistance [8]. For example, it was reported that, in the case of laser-clad nickel-based alloy coatings, the addition of nano-sized CeO2 and La2O3 particles managed their textural evolution and corrosion resistance [9]. Besides, an experimental study by Aruna et al. [10] manifested that the deposited Ni/ceria coatings showed super corrosion resistance and microhardness as compared to pure nickel, which was in well accordance with the results of our recent studies [11]. Balathandan et al. [12] evaluated the Ni-based coatings incorporated with micron- or nano-sized CeO2 particles, indicative of excellent oxidation and corrosion resistance as relative to Ni-ZrO2, Ni-La2O3 and pure Ni. On the same basis, Antill et al. [13] illustrated that the introduction of CeO2 phase in the coatings was contributed to modify the plastic deformation of Ce-rich precipitates through accommodating different coefficients of thermal expansion between Ni-based coatings and their oxide scales. We are perceived from the arguments that large volume fraction of grain boundaries existing in nanocrystalline materials can provide much growing space for reinforcing nanoparticles to make them precipitated along enlarged boundaries through a proper aging treatment. Unfortunately, they are no clear works carried out for this research of the

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beneficial effects of CeO2 nanoparticles on making an effective dispersion strengthening for better interfacial adhesion between coarsegrained Ni grains. In addition, there are also less papers reported on the corrosion behavior of Ni-CeO2 coatings immersed in a 5 wt.% HCl aqueous solution. Hence the objective of this paper is to study the effect of CeO2 addition on managing the textural evolution and corrosion behavior of nanocrystalline Ni-CeO2 coatings by means of microstructural characterizations and electrochemical measurements. 2. Experimental procedures 2.1. Specimen preparation A commercial nickel plate (purity ≥ 99.5%) with dimensions of 65 mm × 20 mm × 2 mm was used as the anode, and a low-carbon steel specimen with dimensions of 50 mm × 10 mm × 0.5 mm was employed as the cathode. Detail processes of surface pre-treatment for the electrodes were reported elsewhere [11]. Bath composition and its operating conditions used for electrochemical deposition of the coatings were summarized in Table 1. The electrolyte used in this study was prepared from reagent grade chemical and distilled water. The pH value, between 2.3 and 2.5, was adjusted by adding 0.5 wt.% dilute sulfuric acid. An experimental apparatus of ultrasound-assisted pulsed electrodeposition was performed as previously described [14]. Pure nickel and the Ni-CeO2 coatings were fabricated from a modified Wattsnickel electrolyte without or with 15 g L−1 CeO2 addition using a novel method of ultrasound-assisted pulsed electrodeposition. In order to ensure the uniform distribution of CeO2 solid particles within the electrolyte before deposition, the as-prepared electrolyte was subject to magnetic stirring (150 rpm) for 2 h, followed by ultrasonic vibration (45 kHz and 80 W) for 60 min. Moreover, this electrolyte was magnetically stirred at a rational rate of 100 rpm to promote that the nanoparticles be well transported to the cathode/electrolyte interface during the process of deposition. Typical coatings with a nominal thickness of ~ 25 μm were obtained at an average current density of 1.2 A dm− 2 for 60 min. Afterwards, the as-deposited coatings were conducted in a muffle furnace for isothermal aging treatments at 650 °C for 4 h, exposed to dry air. 2.2. Reinforcing phase Cerium oxide (n-CeO2), as one of the most powerful oxide particles, has been widely used in the surface modification for the improvements in mechanical properties of the coatings, which is closely associated with its special characteristic of both strong adsorption and high catalytic activity [15]. In this experiment, high purity of CeO2 particles (Alfa Aesar, 99.95%) with a nominal particle size of ~30 nm was added into the bath and served as the reinforcing phase for composite coatings. Table 1 Composition of the electrolytic bath and its operating conditions. Bath composition and operating conditions for composite coatings 350 g L−1 40 g L−1 50 g L−1 15–20 g L−1 0.5–1 g L−1 0.02 g L−1 40 ± 2 °C 2.3–2.5 120–160 rpm 45 kHz/80 W 1.0–1.2 A dm−2 Double-pulsed wave

Nickel sulfate Nickel chloride Boric acid CeO2 nanoparticles Surfactant, R-SO2-M Saccharine Temperature pH Mechanical stirring Ultrasonic field Peak current density Current mode Duty cycle (%)

On time (ms)

Off time (ms)

Pulse frequency (Hz)

25

10

30

1500

Prior to electrodeposition, the as-preformed aqueous emulsion composed of some surfactants (e.g., saccharin sodium (C7H4O3NSNa), sodium dodecyl benzene sulfonate (C12H25 C6H4NaO3S), etc) were used to make 12–15 g L−1 CeO2 particles pre-treated into a suspending liquid for the co-deposition with Ni coatings. As concluded from Ref. [14], the proper amount of CeO2 addition in the bath should be less than 20 g L−1, or it will be probable for those with coagulation problems. With the aid of strong acoustic streaming and fluidal attack generated by ultrasonic field, those nanoparticles could be well transported into the electrical double layer, hence leading into an irreversible adsorption onto the coating surface. With the presence of CeO2 addition, defective regions existed in the coatings were strongly adsorbed by nano-sized CeO2 particles, as confirmed from Fig. 1a. Moreover, in Fig. 1b, it exhibited nano-crystal structure of Ni grains for the Ni-CeO2 coatings attached with well-distributed CeO2 phase through the coatings. The cobblestone-like CeO2 nanoparticles, as presented in Fig. 1c, were favorable for the effects of both nanoparticle reinforcement and dispersion strengthening. 2.3. Surface characterizations and electrochemical measurements Surface morphologies and micro-structural evolution of investigated coatings were characterized by an environmental scanning electron microscopy (E-SEM, Philips/S-4800) performed with EDAX analyzer and transmission electron microscope (TEM, Hitachi/JEOL-2000FX). The determination of corrosion products was carried out by an X-ray diffractometer (XRD,/D8 Focus Bruker) at 40 kV and 30 mA using the Cu Kα (λ = 1.5406 Å) radiation. In order to clarify element distribution of corrosion products formed on the corroded surface, X-ray photoelectron spectra (XPS) measurements were employed and conducted at Thermo ESCALAB 250 using non-monochromatic Al Ka (1486.6 eV) radiation. The electron energy analyzer of XPS analysis was operated at pass energy of 70 eV for survey scans and of 30 eV for high resolution scans. In addition, an ultrasonic generator (KQ-100VDB) was employed for the simulation of ultrasonic field at 45 kHz and 80 W. Electrochemical measurements were performed at a CHI660C electrochemical analyzer in a 5 wt.% HCl acid aqueous solution at 25 °C. A standard three-compartment cell was used with a saturated calomel electrode (SCE) as the reference electrode, a platinum electrode as the counter electrode and the as-prepared specimen with an exposed surface area of 1 cm2 as the working electrode, respectively. Electrochemical impedance spectroscopy (EIS) measurements were conducted at the open circuit potential (OCP) with the AC amplitude of 5 mV and an applied frequency ranging from100 kHz to 0.01 Hz. To further investigate the electrochemical characteristics of the exposed coatings measured in this system, an available method of electrical equivalent circuits (EECs) were proposed from ZsimpWin software to simulate experimental data. For potentiodynamic polarization curves, they were recorded at a sweep rate of 20 mV min−1, starting from a moment when the OCP curve reached a steady-state condition. All of the electrochemical measurements were reproduced at least three times in order to confirm the reproducibility of experimental results. 3. Results and discussion 3.1. CeO2 adsorption CeO2 nanoparticles have attracted extensive interest because of their strong adsorption capacity and high catalytic activity. When added into the bath, they will preferentially adsorb onto the crystal defects and be acted as the catalytic sites for more crystal nucleus. As a result, not only the process of electron charge transfer between Ni2+ ions and [H] atoms hinders, but also the free energy of cathode surface decreases, which leads to a significant increase of the nucleation sites for crystal nucleus of Ni grains. As long as the rate of nucleation and grain growth exceeds dynamic re-crystallization, crystal growth will therefore get suppressed

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435

(a)

200 nm

(b)

(c) CeO2 particles

500 nm

200 nm

Fig. 1. E-SEM images of (a) CeO2 adsorption, (b) well-distributed CeO2 nanoparticles through the coatings, and (c) high magnification of (b) showing the cobblestone-like CeO2 feature.

with the consequent effect of grain refinement [15]. Based on the above results, it can predict that the introduction of CeO2 addition into the bath will affect the electrocrystallization of Ni grains. The predictive model, as illustrated in Fig. 2, exhibited CeO2 nanoparticles adsorbed along interfacial gaps or grain boundaries. While it should be noted that the degree of transfer process was not only associated with the pH value (~2.5) of the bath, but also depended upon the effects of acoustic streaming and temporary cavitation generated from ultrasonic field. During the process of deposition, small Ce4+ ions were released from CeO2 phase carrying with CeO2 particles for nucleation sites. In that case, they could be attracted as catalytic centers to provide more nucleation sites for fine-grained growth. The existence of Ce4+ ions or nano-sized CeO2 particles in the bath was performed with strong adsorption towards Ni grains, as well as blocked crystal defects for vacancy condensation. Due to their strong adsorptive capacity and high catalytic activity, they would be acted as the superficial modifier and complex reagent of Ni2+ cations to reduce the interfacial free energy for promoting more dynamic re-crystallization occurring at the electrode surface. Meanwhile inert CeO2 phase is typical of high resistance (~015 Ω cm2) that can be regarded as the dielectric phase in the coatings to take the shielding effect on preventing the coarsening growth of Ni grains [16].

The formation of cracking interfaces would be therefore deteriorated into the short-circuit diffusion paths, resulting in the conductive pathways for aggravating anion migration and cation transportation upon the subjection to long-term aging treatments in air. Due to the existence of uniform-distributed CeO2 phase and be served as a softening phase to reduce internal stresses, it exhibited denser structure together with a refiner grain size (~160 nm) of Ni grains in Fig. 3d, which can be well explained by: 1) generating more crystal nucleus for dynamic recrystallization that achieved by adding CeO2 nanoparticles into the bath; 2) blocking void defects for vacancy condensation by means of their high adsorption towards defect regions; and 3) making the precipitated Ce-rich phase along interfacial boundaries to create an effective

+ + + 4+ + Ce + + + +

Ni Impules

3.2. Surface morphologies As described from Fig. 3a, lots of micro-pores or cracking boundaries were observed at the coating surface of pure nickel. The clusters of voids defects were probably due to the out gases of H2 or O2 that escaped from the hydrogen evolution reactions near the electrode surface in the bath. Nevertheless, well-dispersed CeO2 nanoparticles in the coatings would make a significant effect on filling the interfacial boundaries for structural integrity. In Fig. 3b, it exhibited denser structure with a finergrained size of ~105 nm of Ni grains for the Ni-CeO2 coatings as compared to that of ~200 nm for pure nickel samples. After aged at 650 °C for 4 h, it displayed the cracking boundaries attached with coarsening growth of Ni grains (~320 nm) for pure nickel, as shown in Fig. 3c.

pH~2.5

CeO 2

+ + + 4+ + Ce + + + +

Small transfer

U

Ultrasonic field

Fig. 2. Schematic diagram of catalytic co-deposition obtained from strong adsorption by Ce4+ ions.

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(b)

(a)

200 nm

200 nm

(d)

(c)

Ni

CeO2 200 nm

(e)

200 nm

(f)

Micro-pores

20 µm

25 µm

Fig. 3. E-SEM images of the coating surfaces for pure nickel: (a) as-deposited, (c) aged at 650 °C; the Ni-CeO2 coatings: (b) as-deposited, (d) aged at 650 °C; and the oxide scales of (e) pure nickel and (f) the Ni-CeO2 coatings, respectively.

bonding link for dispersion strengthening through a proper aging treatment [17,18]. After aged at 650 °C for 4 h in air, it displayed the spalling behavior of oxide scales for pure nickel in Fig. 3e, most probably due to the void defects or thermal micro-cracks existed in the coating surface to accelerate the process of internal diffusion and external oxidation. While for the Ni-CeO2 samples, an undamaged surface of oxide scales was observed in Fig. 3f, which was attributed to the formation of the precipitated Ce-rich phase that could be acted as a passive layer covered on the coating surface to reduce the exposed area available for reducing external oxidation and internal diffusion by active [O] atoms. Moreover, the continuous Ce-rich oxide scale was likely to take an interstitial diffusion on combining the high-angel oxide boundaries for vacancy condensation, thereby promoting a self-repaired effect on completing the spalling of oxide scales [19]. 3.3. XRD analysis Fig. 4 presents X-ray diffraction patterns for pure nickel and the NiCeO2 coatings before and after aged at 650 °C, in which both of them exhibited predominant orientations of Ni (200) and (111) over Ni (220), (311) and (222). Meanwhile some phase of CeO2 (331), (420) and (511) was observed in the Ni-CeO2 coatings, revealing that inert CeO2 phase was well distributed into the coating surface. In addition,

it appeared as the increasing intensities of Ni (220) and (222) as relative to the decreasing trend of Ni (111) and (200), indicative of the effect of grain refinement that was achieved by adding CeO2 phase. The change of preferred orientations and grain refinement was actually attributed to the following reasons: 1) the existence of well-distributed CeO2 nanoparticle or its Ce4+ ions in the bath could partially adsorb onto the surface of Ni2+ and Ni[B(OH)4]+ cations for promoting a complex catalyst, and then the adsorbed cations were attracted as catalytic sites of crystal nucleus for more dynamic electrocrystallization; and 2) inert CeO2 phase existed in the coatings would make a shielding effect surrounding the growth centers to modify the directions of crystal growth and also prevent the coarsening growth of Ni grains [20,21]. During aged at 650 °C for 4 h, an intermetallic compound of NiCe2O4 (044) phase was detected at the 2θ diffraction of 42°, and its formation process of sintering reactions can be expressed as [22]: Ni + 2CeO2 → NiO + Ce2O3 → NiCe2O4 (ΔG = −225.2 kJ mol−1; T = 900 K). In order to estimate the effect of CeO2 addition on the crystallite size of Ni grains, the classical Scherrer equation was used according to the XRD line-boarding patterns of predominated orientations on Ni (111), (222) and (200). The calculated results showed a finer-grained size of Ni grains with ~125 (as-deposited) and 185 nm (aged at 650 °C) for the Ni-CeO2 coatings as relative to ~172 and 254 nm for pure nickel. While it should be noted that the results of grain size for Ni grain were a litter higher than the observed ones (as seen in Fig. 3), which

X. Zhou, Y. Shen / Surface & Coatings Technology 235 (2013) 433–446

(511)

(331) (311) (420)

(220) (200)

(222)

NiO NiCe2O4

CeO2

(111)

(111)

Intensity /a.u.

(044)

Ni

Ni-CeO2, 650 C Ni-CeO2, deposited Pure Nickel, 650 C Pure Nickel, deposited

20

40

60

80

100

2 /degree Fig. 4. XRD patterns of pure nickel and the Ni-CeO2 coatings before and after aged at 650 °C.

was probably attributed to the fact that we did not consider the affecting effect resulted from strain broadening or micro-strains in the diffraction peaks' integral widths. 3.4. Micro-structural evolution Fig. 5a displays a typical bright-field TEM image of coarse-grained Ni grains for pure nickel. Without CeO2 addition, interfacial gaps were developed from the relaxation of internal stresses or thermal cracks after aged at 650 °C, as concluded from Fig. 5b. Nevertheless, high volume fraction of interfacial boundaries existing in nanocrystalline coatings could not only be expected as numerous of growing space for

(a)

437

embedding CeO2-reinforced phase for nanoparticle reinforcement, but also considered as the inter-diffused channels for the precipitated Cerich phase with the consequent improvement of solution treatment and dispersion strengthening through a proper aging treatment [23]. As deduced from Fig. 5c, the process of aging treatment at 650 °C for 4 h was served as a driving force for well-distributed CeO2 phase to make precipitated along enlarged boundaries, hence leading to the hindrance of sliding dislocations or boundaries for better interfacial adhesion. It was evident from Fig. 5d that it displayed the precipitated Cerich phase along the boundaries, well bonded with cracking interfaces to create an effective blocking effect on combining the cracking interfaces for vacancy condensation. Additionally, the precipitation of Cerich phase could reduce thermal stresses that released from coarsening growth or textural evolution, which was mainly contributed to its great plastic deformation to modify their different coefficients of thermal expansion. According to each SAED pattern (seen the top-right corners), crystal texture of Ni grains in pure nickel was transformed from amorphous (as-deposited) state into micro-crystal structure (aged at 650 °C); while for the case of Ni-CeO2 specimen, it was typical of nano-crystal texture for Ni grains that was attributed to the fact that the existence of CeO2 addition promoted an effective dispersion strengthening along the interfacial boundaries, thereby achieving a self-repaired effect on completing those cracking interfaces and the suppression of coarsening growth for denser structure. As proposed from Refs. [24–26], the formation of micro cracks or residual stresses in the coating surface was mainly generated from enlarged boundaries or impurities phase, but rarely suffered from the precipitated RE-rich phase. Up to now, many studies of strengthening mechanisms produced from RE-precipitated phase have attracted extensive interest, which was involved in the beneficial effects on the stabilization of micro structural evolution through effective interactions with the extended defects of dislocations and grain boundaries [27,28]. Such reports have spawned an available proposal of the survey research that focused on the suppressing grain growth in nanocrystalline materials. In

(b)

Cracking boundaries

Coarse Ni grain

100 nm

200 nm

(c)

(d)

Ni

Precipitated Ce-rich phase Nano-CeO2

CeO2

50 nm

50 nm

Fig. 5. TEM images attached with the SAED patterns for pure nickel: (a) as-deposited, (b) aged at 650 °C; and the Ni- CeO2 coatings: (c) as-deposited, and (d) aged at 650 °C.

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general, there are two typical methods referred to the improvement of micro-structural evolution. One is the kinetic approach that involves in a pinning force to trap the sliding grain boundaries in place thereby halting grain growth, which can be accomplished by means of the pinning link with either secondary phases or the effects of solute drag. While for the other one, it is a thermodynamic approach that is used to reduce the driving force for the sliding movement of grain boundaries. This approach has actually relied on the stabilizing process of solute atoms or precipitated phase for reducing the free energy penalty of grain boundaries in nanocrystalline coatings (significant grain boundaries) [18,26]. In order to better understanding the effect of CeO2-dispersed phase on governing micro-structural evolution, a diagrammatic sketch involved in the bonding link on completing interfaces or boundaries is illustrated in Fig. 6. It is well known that there is a high volume fraction of grain boundaries existing in the nanocrystalline materials so that they can be regarded as the growing space for embedding CeO2 nanoparticles to make dispersion strengthening and particle reinforcement. In this study, nano-sized CeO2 particles are supposed to adsorb onto Ni grains and make them precipitated along vacant sites for vacancy condensation, thereby leading into a pegging effect on the suppression of cracking interfaces for structural integrity (seen Fig. 5d). In that case, nanosized CeO2 particles can provide extra nucleation sites for inducing heterogeneous nucleation and more re-crystallization of Ni grains. Well-distributed CeO2 phase, which can be acted as a softening phase in the coatings, will perform with super plastic deformation leading into the consequent modification of thermal stresses that released from textural evolution or coarsening growth of Ni grains. Besides, their vicinal atoms are considered as spherical sites for the direction of Ni growth to reduce thermal residual stresses because of its high adsorptive capacity and low electronegativity of the precipitated Ce-rich phase [28,29]. During the process of aging treatments, large numbers of grain boundaries can be supposed as the inter-diffused channels for the formation of the precipitated Ce-rich phase, which leads to an effective pinning force on bonding the interfacial boundaries for better interfacial adhesion. The existence of CeO2 nanoparticles in the coatings can be acted as a softening phase to make a buffering effect on the decreasing of internal stresses. In addition, the effect of triple junction (TJ) drag generated from the precipitated Ce-rich phase will promote an effective bonding link between the sliding boundaries or interfaces for structural integrity, which effect is commonly referred to as “Zener pinning” [18]. The as-described TJ drag is intrinsically favorable for constructing a TJ obstacle on the adjacent interfaces to prevent the movement of sliding boundaries and coarsening growth to some extent. Meanwhile the solute-drag effect resulted from the Ce-rich phase or its ions is closely correlated with the diffusion activation energies of grain boundaries, hence leading into the reduction of the driving force for the mobility of enlarged boundaries. According to predictive modeling of Fig. 6d and as-mentioned results from our study, it can predict that the significant grain boundaries in nanocrystalline materials can be utilized to manage the structural evolution. In particular, the effective bonding effect produced from CeO2 nanoparticles are contributed for the precipitation of well-distributed Ce-rich phase along interfacial boundaries for denser microstructure. 3.5. Electrochemical measurements 3.5.1. Open circuit potential (OCP) The OCP vs. time curves, as depicted in Fig. 7, were recorded for pure nickel and the Ni-CeO2 coatings in a 5 wt.% HCl solution at 25 °C. All of the OCP curves reached a steady-state condition at a short period, indicative of the stability of measured system. As for pure nickel sample, it displayed the negative potential attached with an unstable trend at the early stage, revealing anodic dissolution and localized corrosion occurred at the exposed surface [30]. With the CeO2 addition, a positive potential by more than ~150 mV was observed for the Ni-CeO2 coatings

as relative to pure nickel, correlating well with above results that welldistributed CeO2 particles on the coating surface were served as the dielectric phase to prevent the special adsorption of Cl− ions in the corroding solution. Additionally, inert CeO2 phase is of a positive potential (+1.61 V) so that it is favorable for an increase in hydrogen evolution over-potential to reduce the cathodic reactions occurred at the exposed surface. After aged for 4 h at 650 °C, the precipitated Ce-rich oxide scales formed on the coating surface were acted as a passive layer to reduce the surface defects, thereby decreasing the active sites for the initiation of localized corrosion [31]. 3.5.2. Polarization curves For comparison of their corrosion behavior, potentiodynamic polarization (Tafel) curves were employed and measured in a 5 wt.% HCl aqueous solution. They were polarized from −800 to +600 mV (vs. SCE) at a scan rate of 20 mV min−1. As described from Tafel curves in Fig. 8, an overlap of anodic polarization curves was observed together with a slight transition of self-passivation behavior, revealing super corrosion resistance of the as-prepared coatings. Electrochemical parameters such as corrosion potential (Ecorr), corrosion current density (Icorr) and anodic/ cathodic Tafel constants (βa and βb), were derived using Tafel extrapolation from polarization curves. Moreover, polarization resistance (Rp) was determined from the typical Stern–Geary equation [32]. Based on the results in Table 2, both of Ecorr and Icorr values positively shifted from −0.38 V and 2.22 × 10−6 A cm−2 of the as-deposited NiCeO2 coatings into −0.43 V and 4.47 × 10−6 A cm−2 of pure nickel. The slightly increased corrosion resistance was actually associated with the addition of CeO2 phase that adsorbed along crystal defects in the coating surface for denser coatings, hence leading into the reduction of specific adsorption by Cl− ions. After aged for 2 h at 650 °C, it showed a higher Ecorr value of −0.29 V together with a lower Icorr value 2.2 × 10−7 A cm−2 for the Ni-CeO2 coatings as compared to −0.38 V and 5.12 × 10−6 A cm−2 for pure nickel. It was found that the corrosion current density of as-deposited Ni-CeO2 coating was about one order of magnitude lower than that of pure nickel, but slightly higher than that of the Ni-CeO2 coating after aged at 650 °C. Compared to pure nickel, 100 mV more positive potential can decrease the driving force of electrochemical reactions for the Ni-CeO2 coatings to some extent. During the process of Tafel measurements, the aggressive Cl− ions not only permeated into the Ni-based coatings to aggravate the corrosive attack, but also accelerated the forms of localized corrosion occurred at the corroded surface. With the presence of CeO2 addition, diffusion behavior by Cl− ions was much weaker than that of those without CeO2 addition, which was most probably due to inert CeO2 phase in the coatings that would be acted as the dielectric phase to make a shielding effect on decreasing the specific adsorption of Cl− ions along the exposed area or pitting holes. As a result, pure nickel exhibited an obvious pitting behavior in an acid HCl solution; while for the case of the Ni-CeO2 sample, it was immunized from the Cl- bearing solution, showing a slightly selfpassivated behavior at the anodic transition. The shift to a lower current value (Icorr) by almost one order of magnitude of the Ni-CeO2 coating (aged at 650 °C) than that of pure nickel can be explained by the following reasons [33–35]: (i) Inert CeO2 phase (more than 1015 Ω cm2) is regarded as dielectric phase that can make a shielding effect on limiting the electrical charge transport and ions exchange to prevent the diffusion behavior by Cl− ions occurred at the corroded surface; (ii) Similar lattice type, face-centered cubic (fcc), of CeO2 phase and Ni grain is favorable for denser microstructure, thereby leading into the reduction of pitting holes and localized corrosion that induced by crystal defects or surface micro-pores; (iii) During the process of aging treatment, the continuous Ce-rich oxide scales acted as a passive layer formed on the coating surface are contributed to reduce surface defects. In addition, the precipitation of the Ce-rich phase is also useful for dispersion

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439

Coarse Ni grain

Ni grain The precipitated Ce-rich phase

C 50 @6 n o i t ddi 2a CeO

Nanocrystalline Ni grain

t hou Wit Wi th Ce O2

Relaxtion of internal stresses add itio n@

650 C

Note: A high volume fraction of grain boundaries are expected as the adsorption space for embedding the CeO 2-reinforced phase.

Bonding link produced from the precipitated Ce-rich phase

(b)

200 nm CeO2 adsorption

(a)

(c)

CeO2

100 nm Cracking boundaries in pure nickel

50 nm Bonding link on boundaries

Fig. 6. Predictive modeling of a bonding link on combining interfacial boundaries: (a) the defects of cracking boundaries existed in pure nickel, (b) the adsorption of CeO2 phase along interfaces, and (c) a bonding link produced from the precipitated Ce-rich phase.

strengthening by means of internal diffusion, hence resulting in the decreasing of the active sites for the initiation of localized corrosion. 3.5.3. EIS diagrams The AC-impedance (EIS) measurements were used to further study their corrosion behavior in a 5 wt.% HCl aqueous solution, which can offer deep insights into electrochemical characteristics of the measured system. In Fig. 9a and b, they presented the tested results of EIS diagrams for investigated specimens, including Nyquist plots (frequency

vs. impedance modulus |Z|) and Bode plots (frequency vs. phase angle). It was noted from Nyquist plots that only one difference among EIS diagrams tested for specimens was the size of capacitive loops. As a general rule, the larger size of the capacitive loop is, the better the corrosion resistance of a coating will present. Compared to pure nickel, a larger size of the capacitive loop was observed in the entire frequency range for investigated coatings reinforced with CeO2 phase. The detail size of the capacitive loop was about 5 orders of magnitude larger than that of pure nickel, indicative of improved corrosion resistance for the Ni-CeO2 coatings. Such results were coincided well with above

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0.0

Potential /VSCE

-0.1

-0.2

-0.3 Pure nickel (As-deposited) Pure nickel (650 C) Ni-CeO2 (As-deposited) Ni-CeO2 (650 C)

-0.4

-0.5 0

600

1200

1800

2400

3000

3600

Time /sec Fig. 7. Variation of open circuit potential (OCP) with time for pure nickel and the Ni- CeO2 coatings immersed in a 5 wt.% HCl solution at 25 °C.

observations of polarization curves. As mentioned above, pinholes or micro cracks that produced from void defects and stress relaxation were prone to bring about the initiation of localized corrosion occurred at the corroded surface. Nevertheless, when the inert CeO2 phase was embedded into the coatings, it would be acted as the dielectric phase to make an effective shielding effect for reducing charge transfer and the specific adsorption of Cl− ions [36]. Moreover, the Ce-rich oxide scale, which was characteristic of a super hydrophobic surface, could be served as the passive layer covered on the coating surface to decrease the defective regions for a denser surface [37,38]. As a result, larger impedance module was measured for the Ni-CeO2 coatings (aged at 650 °C) as relative to both of pure nickel and the Ni-CeO2 coatings (as-deposited), which was mainly contributed to the precipitation of the Ce-rich phase along enlarged boundaries to promote an effective bonding link on completing the cracking interfaces or pinholes for denser coating surface. As for pure nickel samples, the diffusion behavior by Cl− ions was observed because of severe localized corrosion occurred at the corroded surface, hence resulting in Warburg impedance. As apparent in Fig. 9b, Bode plots exhibited a larger capacitive arc for Ni-CeO2 coatings attached with the largest phase angle close to ~−80° over a wide range of middle frequencies (103–100 Hz). While for pure nickel sample, it displayed the largest phase angle of ~−60° over a

0.6 Pure nickel (As-deposited) Pure nickel (650 C) Ni-CeO 2 (As-deposited) Ni-CeO 2 (650 C)

0.4

E (V vs. SCE)

0.2 0.0 -0.2 -0.4 -0.6 -0.8 -9

-8

-7

-6

-5

-4

-3

-2

-1

0

1

Log (i, A/cm2) Fig. 8. Potentiodynamic polarization curves recorded for pure nickel and the Ni-CeO2 coatings.

narrow frequency range (103–101 Hz). As depicted from the result of Bode plots, it displayed the log-linear relationship between impedance value (Z) and applied frequencies (f), indicative of an integrated exposed surface and a predominantly capacitive response measured in the corroding solution. On the other hand, two time constants were observed for the present system of the Ni-CeO2 coatings, revealing two different types of electrochemical characteristics occurred at the interfaces of the coatings/corrosive solution and the coatings/the passive layer. As described before, a continuous Ce-rich oxide scale formed on the coating surface was interpreted as the passive layer covered on the exposed area leading into the decreasing of localized corrosion, instead of electrochemical dissolution or diffusion behavior by Cl− ions for pure nickel sample where subjected into a 5 wt.% HCl solution. Therefore, an increased value of absolute impedance was tested at the low-frequencies (LF) region for Ni-CeO2 coatings as against pure nickel, most probably due to this barrier layer enrich of a Ce-rich phase to self-repair the void defects, micro-cracks or the spalling behavior of oxide scales for a denser surface of the coatings. Considering similar EIS results of the Ni-CeO2 coatings reported by Aruna et al. and Gray et al. [10,34], electrochemical interfaces in this system could be defined as: Ni-based coatings/the passive layer (e.g., Cerich oxide scales)/the corroding solution. The absolute impedance |Z| recorded at the low-frequencies (LF) region revealed intrinsic properties of measured specimens, corresponding to the resistance of charge transfer. While an impedance value measured at the high-frequencies (HF) range only responded to the surface characteristics of corroded area. According to the above arguments, two typical electrical equivalent circuits (EECs) were employed and used to fit experimental data, as shown in Fig. 10, where Rs corresponds to the solution resistance, Qdl is a constant phase element (CPE) which represents the capacitance of the electric double layer at HF region, Rct denotes the charge-transfer resistance, the values of Qpf and Rpf are the capacitance and resistance occurred at the passive layer/corroding solution interface happened at LF range. In Fig. 10a, an EEC model of Rs(Qdl(RctW)) was illustrated for pure nickel, including one time constant of the coating's resistance (Rct) and the electrode double-layer capacitance (Qdl), in parallel with a Warburg (W) resistance resulted from the diffusion behavior by Cl− ions. Taking into account the precipitation of Ce-rich phase served as the passive layer on the coating surface, the EEC model composed of two time constants was introduced to fit the experimental data. Fig. 10b presents the fitting model of Rs(Qdl(Rct(QpfRpf))) for the Ni-CeO2 coatings. It was contained of two time constants, for instance, one time constant of (QdlRct) occurred at HF range (~103 Hz) was closely related to the intrinsic properties of the as-prepared coatings; the other time constant (QpfRpf), around LF region (~10−1 Hz), was well correlated with the Ce-rich passive layer and governed by the chargetransfer process. In view of non-uniform current distribution caused by the roughness or local defects of electrode surface, it will bring about the frequency dispersion effect. So a constant phase-angle element (CPE) was introduced to replace an ideal capacitance element for better fitting procedures in this study. The impedance of a CPE can be expressed as: ZCPE = [Y0(jω)n]−1, where Y0 is a constant of the CPE which represents the passive-layer capacitance, j is an imaginary number, ω is the angular frequency (ω = 2πf rad s−1), and the value of n is an adjustable empirical exponent for the nature of the CPE, being n = 1 for a purely capacitive behavior associated with a perfectly smooth surface and n ~ 0.5 for a mass transfer process occurred at the porous surface. The simulated EIS results, as listed in Table 3, were obtained from the fitting procedures according to the ZsimpWin software. The fitting quality was evaluated using the statistical method of chi-squared (χ2) tests. The fitting results were of the order of 10−3, implying a good agreement between experimental data and simulated results. As calculated from the given EIS models, Rct value of pure nickel was ~0.45 (as-deposited) and 1.02 kΩ cm2 (aged at 650 °C). While for the case of Ni-CeO2 coatings, they were about 2.89 and 7.49 kΩ cm2, respectively. An increasing

X. Zhou, Y. Shen / Surface & Coatings Technology 235 (2013) 433–446

441

Table 2 Derived results from the potentiodynamic polarization measurements. Type of coatings

Ecorr (mV vs. SCE)

lg (Icorr, A/cm2)

bc, mV/dec

ba, mV/dec

Rp, kΩ cm2

Pure nickel,deposited Pure nickel, 650 °C Ni-CeO2, deposited Ni-CeO2, 650 °C

−0.43 −0.48 −0.38 −0.29

−4.47 −5.12 −2.22 −0.22

−5.58 −5.89 −6.23 −6.74

5.63 7.87 7.22 6.74

10.78 06.98 16.88 57.24

value of Rct was attributed to the well-distributed CeO2 phase in the coating surface that acted as the passive layer with a shielding effect on limiting the diffusion behavior by Cl− ions. According to the result of higher Rpf and lower Qpf values measured at LF regions, it reflected super corrosion resistance of the Ni-CeO2 coatings which was probably due to the fact that micro-pores acted as the conductive pathways for Cl− ions were well blocked by insoluble corrosion products (as described below). In particular, a continuous Ce-rich passive layer covered on the coating surface compelled the corroding solution in contact with either dielectric CeO2 phase or the Ce-rich oxide scales, but rarely exposed to the Ni-based coatings. Due to the formation of a Ce-rich passive layer with the typical hydrophobic property and higher electric resistance, the process of charge transport and Cl− ion migration had got effectively inhibited. Moreover, the existence of Ce4+ or Ni2+ cations in the corroding solution were likely to combine with the OH− anions, thereby generating some insoluble products formed on the exposed area and also reducing the active sites for localized corrosion. In order to clearly study the effect of the Ce-rich passive layer on the capacitive characteristics of the passive layer/corroding solution interface, especially with respect to the Qpf capacitance. The value of capacitance (Qpf) can be defined by: εr ε0 Ac d

(a)

10000 Pure nickel (As-deposited) Pure nickel (650 C) Ni-CeO2 (As-deposited) Ni-CeO2 (650 C) Simulated

8000

ð1Þ

Z'' /Ω cm2

Q pf ¼

passive layer covered on the coating surface for inhibiting the process of ions exchange and charge migration. As a result, in Fig. 11d, it displayed small corrosion products without any pitting holes for the Ni-CeO2 coatings (aged at 650 °C). During long-term static immersion tests in an acid 5 wt.% HCl aqueous solution, Ce3+ ions reduced from CeO2 phase were favorable for generating some insoluble corrosion products (e.g., CeCl3, Ce(OH)3) covered on the corroded surface. The corrosion products, as determined by the RD analysis, were shown in Fig. 11e and f. As proposed from Refs. [30,35], the formation of such corrosion products was contributed to make the plugging effect on filling the pitting holes for an integrated surface of exposed area, as well as to effectively reduce electrochemical dissolution of Ni metal or diffusion migration of Ni2+ ions, thereby decreasing the initiation of localized corrosion or electrochemical dissolution occurred at the corroded surface.

where ε0 and εr are the dielectric constant of the passive layer and the constant of permittivity of free space (about 8.85 × 10−12 F/m), d is the thickness of the passive layer and Ac is the exposed area. As conducted form Eq. (1), the value of Qpf is well correlated with the dielectric constant of the passive layer. The dielectric constant, εr, of the passive layer will get increased with the existence of inert CeO2 phase due to its stable physical and chemical characteristics, high resistance (~1015 Ω cm2) and low electronegativity [39].

1200 800 400

6000

0 0

400

800

1200

1600

4000 2000 0 0

2000

3.6. Corrosion morphologies and their product analysis

4000

6000

8000

10000

Z' /Ω cm2

Phase angel (degree)

(b) 80

Log (|Z| /Ω cm2)

The corroded surfaces of these specimens were observed after being immersed in a 5 wt.% HCl solution for 10 days at 25 °C exposed to air, and their corrosion products were determined by XRD analysis. As noted from Fig. 11a, it exhibited selective corrosion attached with some pitting holes on the corroded surface of pure nickel, most probably because of pinholes or micro cracks that acted as the active sites for the initiation of localized corrosion. After aged for 2 h at 650 °C, lots of pitting holes were consequently generated from the release of internal stresses or coarsening growth, hence resulting in the short-circuit diffusion channels at the corroded surface to suffer from the corrosive attack by aggressive Cl− ions. Worse still, severe localized corrosion was induced from surface defects, showing an overall corrosion around the corroded surface for pure nickel in Fig. 11c. The results obtained in the present work were well consistent with our previous study [40]. With the presence of CeO2 addition, uniform corrosion surface was detected for the Ni-CeO2 coatings, as shown in Fig. 11b, probably due to inert CeO2 phase that adsorbed onto the coating surface and acted as the dielectric phase with a hydrophobic surface to prevent the diffusion behavior by Cl− ions. After aged at 650 °C, the precipitated Ce-rich phase produced from well-distributed CeO2 phase was beneficial to form the

1600

4

60 40 20 0 Pure nickel (As-deposited) Pure nickel (650 C) Ni-CeO2 (As-deposited) Ni-CeO2 (650 C)

3 2 1 -2

-1

0

1

2

3

4

5

Log (Frequency /HZ) Fig. 9. Electrochemical impedance plots recorded for investigated specimens: (a) Nyquist plots and (b) Bode plots. Symbol: experimental data; line: simulated data.

442

X. Zhou, Y. Shen / Surface & Coatings Technology 235 (2013) 433–446

(b) Ni-CeO2 coatings

(a) Pure Nickel

0.2 µm

0.2 µm Surface Review: Lots of pinholes

Surface Review: Denser structure by addition of CeO2 phase

Qdl CPE1

CPE1

Qdl

Rs

RE

WE

Rs

RE

WE

W

Rct

Rct

EEC simulation: Rs (Qdl (Rct W))

Qpf CPE2 Rpf

EEC simulation: Rs (Qdl (Rct (Qpf Rpf )))

Substrat Coating A passive layer

Diffusion behavior by Cl "Warburg"

Cl

Cl

H H

Rpf H

Rs

Rct

Rct

Cl

Cl

Rs Cl

H Cl

Cl

HCl solution

Qpf

Qdl

Qdl

Rs – Solution resistance

Qdl – The double layer capacitance occurred at HF region

R ct– The charge transfer resistance of coatings

W – Warburg impedance, corresponding to charge transfer in LF region

Rpf – The resistance of a passive layer

Q pf– The capacitance produced from the Ce-rich passive layer

Fig. 10. Equivalent electrical circuits used to simulate experimental data of impedance spectra for (a) pure nickel and (b) the Ni-CeO2 coatings.

3.7. XPS studies and formation mechanisms of corrosion inhibition In order to accurately examine their corrosion products formed on the corroded surface, XPS measurements were carried out using nonmonochromatic Al Ka (1486.6 eV) radiation. All of the binding energies

were referenced to C1s (284.6 eV) for compensating the effects of surface charging. Fig. 12(a-1) and (b-1) shows the full XPS spectra for corrosion products of the Ni-CeO2 coatings (aged at 650 °C) before and after static immersion tests in a 5 wt.% HCl solution, several peaks of Ni, Cl, O, C and Ce elements were identified according to a standard

Table 3 Fitting results of impedance spectra from the proposed equivalent circuit simulation. Type of coatings

Rs (Ω cm2)

Rct (Ω cm2)

CPE1 −1

Qdl − Y0 (Ω Pure Ni, deposited Pure Ni, 650 °C Ni-CeO2, deposited Ni-CeO2, 650 °C

8.07 7.59 8.01 8.12

1.82E−6 1.08E−6 2.08E−6 2.93E−6

n

−2

s cm

)

−1

Qpf − Y0 (Ω

nd 0.89 0.75 0.90 0.95

CPE2

0.45E+3 1.02E+3 2.89E+3 7.49E+3

– – 10.45E−6 15.76E−6

n

−2

s cm

)

Rpf (Ω cm2)

W (Ω cm2 s0.5)

χ2 (×10−3)

– – 8.56E+3 9.88E+3

10.52 03.44 – –

3.55 6.88 0.59 0.42

nf – – 0.92 0.95

X. Zhou, Y. Shen / Surface & Coatings Technology 235 (2013) 433–446

(a)

443

(b)

2 m

(c)

5 m

(d)

2 m

(e)

5 m

(f)

NiCl2

Ni(OH)2 Ce(OH)3

80

100

Intensity /a.u.

Ni(OH)2

NiCl2 CeCl3 CeO2

650 C

650 C

Deposited

Deposited

20

40

60

80

100

2 /degree

20

40

60

2 /degree

Fig. 11. E-SEM images of the corroded surfaces for pure nickel before (a) and after (b) aged at 650 °C; the Ni-CeO2 coatings before (c) and (d) after aged at 650 °C; and the determination of their corrosion products by XRD analysis listed in Fig. 10 (e–f).

XPS database, in which the presence of C element was from a contaminant hydrocarbon layer in the XPS instrument. From quantification of XPS intensities for these elements of the coatings before immersion tests, they were composited of 3.22 at.% Cl, 21.82 at.% O, 48.52 at.% Ni, 18.32 at.% C and 8.12 at.% Ce, respectively. To clarify the valence states of Ce and Ni, XPS high resolution scans of Ni 2p3 and Ce 3d spectrum were estimated and shown in Fig. 12 (a-2) and (a-3). Meanwhile, their characteristic peaks were de-convoluted using the Gaussian–Lorentzian line-shapes together with a Shirley-type background. Two fitting peaks of Ni 2p3 located at binding energies of 858.12 and 863.86 eV were closely related to the NiO phase. In addition, the satellite peaks separated from the Ce 3d spectrum at 881.95, 899.02 and 885.98, 904.13 eV were characteristic of CeO2 and Ce2O3 phase. The measured results of XPS analysis revealed that the occurrence of a Ce-rich oxide phase acted as the barrier layer covered on the coating surface. After long-term static immersion tests, the corrosion products formed on the exposed surface contained 25.88 at.% Cl, 29.01 at.% O, 28.07 at.% Ni, 12.45 at.% C and 4.59 at.% Ce. To study the valence states

of Ce and Ni, XPS high resolution scans of Ni 2p3 and Ce 3d spectrum were exactly scanned and shown in Fig. 12 (b-2) and (b-3). The 2p3 spectrum was fitted with three main peaks with the respective binding energies at 851.89, 855.54 and 861.15 eV, indicative of the existence of Ni2+ as Ni(OH)2, NiCl2 and NiOOH phase. Besides, curve fitting of Ce3+ and Ce4+ peaks conducted at 887.8, 884.5 and 905.8 eV was separated from Ce 3d spectrum, showing the formation of ionization state (Ce3+, CeCl3) and oxidation state (Ce4+, CeO2) in the corrosion products [41,42]. According to the above analysis, the XPS results confirmed that the corrosion products formed on the corroded surface were enriched with the Ce-rich phase. It is well known that the existence of Ce3+ ions and inert CeO2 phase is typical of strong adsorption and high catalytic activity, which will lead to a shielding effect on limiting the special adsorption of Cl− ions. Furthermore, due to the formation of insoluble corrosion products covered on the corroded surfaces, it can predict that the process of charge transport and ion exchange surrounding the pitting holes has been inhibited, as well as to prevent the diffusion behavior by Cl− or H+ ions and against further corrosion [43].

X. Zhou, Y. Shen / Surface & Coatings Technology 235 (2013) 433–446

20000

2 µm 40000 20000

Full XPS spectrum

Full XPS spectrum 0

0 0

200

400

600

800

1000

0

200

Binding Energy /eV

400

600

800

1000

Binding Energy /eV

32000

48000

(a-2)

(b-2)

Ni 2p3

Ni 2p3

45000

31000

NiO

30000

XPS intensity /a.u.

XPS intensity /a.u.

Ce 3d

O 1s

60000

C 1s

O 1s

C 1s

Cl 2p

40000

80000

Cl 2p

Ce 3d

2 µm 60000

XPS intensity /a.u.

100000

80000

Ni 3p Ni 3s

XPS intensity /a.u.

100000

(b-1) Ni 2p3

120000

(a-1) Ni 2p3

120000

Ni 3p Ni 3s

444

29000 28000

NiCl2

42000

NiOOH

Ni(OH)2 39000 36000 33000

27000

30000 26000 848 850 852 854 856 858 860 862 864 866 868 870

849

852

Binding Energy /eV

(a-3)

861

(b-3)

Ce 3d

864

867

Ce 3d

36000

CeCl 3

Ce 2O3

Ce2O3

38000

CeO 2 36000

CeO 2 34000 32000 30000

XPS intensity /a.u.

XPS intensity /a.u.

858

38000

42000 40000

855

Binding Energy /eV

34000

CeO 2

CeCl 3

32000 30000 28000

28000 26000 876 880 884 888 892 896 900 904 908 912 916

Binding Energy /eV

876

882

888

894

900

906

912

918

Binding Energy /eV

Fig. 12. XPS analysis of the Ni-CeO2 coatings (aged at 650 °C) before (a) and after (b) long-term static immersion tests: full XPS survey spectra and high-resolution XPS scans of Ni2p3 and Ce 3d, respectively.

In order to better understand the mechanism of corrosion inhibition, a schematic diagram was illustrated in Fig. 13, which involves in the fact that a small amount of Ce3+ ions was released from the Ce-rich oxide scales or CeO2 phase. As mentioned above, these ions are provided with strong adsorption capacity so that they can be expected as

corrosion inhibitors by means of producing some insoluble corrosion products on the corroded surfaces. Meanwhile they will also take an important role on plugging the pitting holes for the decreasing of shortcircuit diffusion paths for Cl− ion migration and the charge-transfer process. During long-term static immersing tests, the existence of Cl− ions

X. Zhou, Y. Shen / Surface & Coatings Technology 235 (2013) 433–446

in the corroding solution will make specific adsorption towards the exposed area to aggravate surface defects, eventually resulting in the occurrence of localized corrosion. Nevertheless, both of inert CeO2 phase and small Ce3+ ions could be served as the catalytic nucleus for complex cations and therefore, generating the formation of insoluble corrosion products on the corroded surfaces. As indicated from Eq. (2), in a 5 wt.% HCl aqueous solution, the HCl is likely to be divided into H+ and Cl− ions in order to provide enough H+ ions for the Ce-rich oxidation product of Ce2O3 phase to occur in the following reaction: Ce2 O3 þ 6HCl→2CeCl3 þ 3H 2 OðpHe2:3–2:5Þ:

ð2Þ

When nickel metal is immersed into an acid aqueous solution, surface ionization will take place due to the electric charge difference at the solid/liquid interface. According to Eqs. (3) and (4), they are typical of electrochemical reactions of nickel coatings occurred in the acid 5 wt.% HCl corroding solution. For example, nickel metal dissolves into an acid HCl solution in the form of positively charged Ni2+ ions, where 2þ

Ni−2e→Ni :

ð3Þ

by means of combining with H+ ions and the reduction reaction of oxygen, the amount of H+ ions in the corroding solution will decrease with increasing immersion time leading into the increase of pH value. Meanwhile active Cl− ions are prone to drill into Ni-based coatings resulting in the diffusion behavior for the initiation of localized corrosion. In order to maintain the electrical neutrality of pitting holes (significant Cl− ions and free electrons with negative charge), the free Ni2+ ions with positive charge are likely be attracted by the negatively charged Cl− or OH− ions to make the accumulation of corrosion products filling on the pitting holes. As described above, small Ce3+ ions released from the precipitated Ce-rich phase were favorable for promoting the traditional corrosion products (e.g., Ni(OH)2, NiCl2) turned into some insoluble corrosion products (e.g., CeCl3, Ce(OH)3) or a complex compound of (Ni, Ce)(OH)5 [44]. As a result, the exposed area is well covered by such insoluble corrosion products, thereby achieving an effective shielding effect on precluding the diffusion behavior and electrochemical reactions by H+ or Cl− ions by means of blocking the short circuit diffusion paths. During the long-term static immersion tests, several corrosion reactions can be expressed as [45,46]for pure nickel: 2Ni

Without the presence of air or oxygen, hydrogen ions will be reduced by the excess of electrons at the cathode surface to take the reaction of hydrogen evolution: þ

−

2H þ 2e →H2 :

ð4Þ

While it should be noted that in air where oxygen concentration is available, oxygen-derived areas could be regarded as the oxygen-rich area so that corrosive reactions of Eqs. (5) and (6) are closely dependent upon the process of oxygen reduction and hydrogen evolution reaction. The occurrence of oxygen reduction is indeed supposed to take place at the air-solution interface in a dilute 5 wt.% HCl aqueous solution, fully aerated the exposed area to drive the corrosion of deeper, airdeprived areas. As a result, the resulting concentration cells or oxygen deprivation will bring about electrochemical reactions that occurred at the anodic areas of localized corrosion. Common forms of such corrosive attack are commonly referred to as “oxygen concentration-cell corrosion”. With the presence of oxygen or air, two possible reactions of oxygen reduction that occurred at the corroded surface are given by: þ

O2 þ 4H þ 4e→2H2 O −

O2 þ 4e þ 2H2 O→4OH

ð5Þ −

ð6Þ

In a slightly acid solution (pH ~ 2.5–2.8), both of hydrogen evolution (Eq. 4) and oxygen reduction (Eqs. 5, 6) probably take place at the corroded surface. However, due to the formation of hydroxyl ions or water

HCl solution Covered by insoluble corrosion products

Oxide layer Ni coating Iron substrate

2þ

−

−

þ 2OH þ 2Cl →NiðOHÞ2 þ NiCl2 ;

2þ

−

3þ

-

ð8Þ

4. Conclusions The beneficial effects of CeO2 addition on phase transformation and micro structural evolution of nanocrystalline Ni coatings were evaluated using XRD, E-SEM (EDX) and TEM analysis. It exhibited a finer-grained size of ~105 nm for the as-deposited Ni-CeO2 coatings as compared to that of ~200 nm for pure nickel. High density of grain boundaries or interfaces were expected as adsorption space for embedding CeO2 nanoparticles to make the coexistence with Ni grains, as well as acted as diffused channels for CeO2-precipitated phase to promote an effective bonding link on completing the cracking interfaces for denser microstructure through aging treatments at 650 °C. It was well evident from the corrosion morphologies by E-SEM observations that a slight uniform corrosion was observed for the NiCeO2 coatings instead of severe localized corrosion for pure nickel when measured in a 5 wt.% HCl solution. The results of EIS diagrams and polarization curves showed super corrosion resistance to the NiCeO2 coatings as against pure nickel. Moreover, the determination of corrosion products by XRD and XPS analysis reveled that both of two type coatings were mainly composed of NiCl2 and Ni(OH)2 phase,

3+

Ce

-

OH

+

Ni+ H+

−

2Ni þ Ce þ 5OH þ 2Cl →NiðOHÞ2 þ CeðOHÞ3 þ NiCl2 →ðNi; CeÞðOHÞ5 þ NiCl2 :

O2

Cl

ð7Þ

while for the Ni-CeO2 coatings:

+

H

445

(Ni, Ce)(OH)5

Localized corrosion

+ + + 3+ + Ce + + + +

Intermetallic inclusions

CeO2

Fig. 13. Schematic representation of corrosion inhibition mechanism involved in the precipitation of some insoluble corrosion products that resulted from Ce3+ ions or its hydroxide.

446

X. Zhou, Y. Shen / Surface & Coatings Technology 235 (2013) 433–446

except some insoluble products of CeCl3 and CeO2 phase generated on the corroded surface of the Ni-CeO2 coatings. During the process of long-term static immersion test in an acid 5 wt.% HCl solution, a small amount of Ce3+ ions was released from the precipitated Ce-rich phase, thereby accumulating some insoluble products (e.g., CeCl3, Ce(OH)3 and CeO2) on the corroded surface to prevent the diffusion behavior by Cl− ions and localized corrosion. The novel method of ultrasound-assisted pulsed electrodeposition can be expected as an available approach to fabricating nanocrystalline coatings. Moreover, additional studies of dynamic simulation should be performed to further study the effect of a Ce-precipitated phase along interfacial boundaries on managing textural evolution for achieving micro-structural integrity. Acknowledgments

[12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29]

The authors acknowledge the financial support from the Fundamental Research Funds for Central Universities and Jiangsu Innovation Program for Graduate Education (no. CXLX12_0151). We would like to express our sincere gratitude to Prof. H.M. Jin group for E-SEM and TEM observations.

[30]

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