Impact of CeO₂ incorporation in electrodeposited Ni-Fe on structure and physical properties of multifunctional nanocomposites

3.1 Effect of current density

With a view to study the effect of current density on coating, the electrodeposition was carried out in its range, 1.0–5.0 A dm⁻². Bright, adherent, fine grained and smooth deposits were obtained at a current density of 3.0 A dm⁻² but the quality of the coatings deteriorated at a higher current density. Variations in the content of iron and co-deposited CeO₂ particles with current density are shown in Fig. 1. The Fe content in the coating decreases from 33 wt% to 19 wt% when the current density increases from 1.0 A dm⁻² to 5.0 A dm⁻². A marginal decrease in the iron content at higher current densities indicates the system to come under diffusion control.

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Figure 1

Variation of (a) CeO₂ and (b) iron content in the nanocomposite with current density.

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The content of CeO₂ increases sharply in the initial stages with increasing current density and decreases thereafter achieving the maximum at a current density of 2.0 A dm⁻². A gradual decrease in the CeO₂ particle content is observed at higher current densities. The observed trend of variation of particle content in the coatings can be examined in light of Guglielmi’s two step adsorption model (Guglielmi, 1972): physical (loose) adsorption of the particles followed by a field-assisted strong adsorption. It is believed that due to strong adsorptive capacity of CeO₂ particles, nickel ions get adsorbed on their surface (Zhou and Shen, 2013). Once the particles are adsorbed on the cathode surface, metal begins building around the particles slowly, encapsulating and incorporating the particles. The strong adsorption can be promoted by the higher current density and consequently the particle incorporation in the coatings initially increases with increasing current density. The highest content of CeO₂ particles in the coating at current density 2.0 A dm⁻² is due to the saturation in adsorption on cathode surface. Observed decrease in the CeO₂ content above current density 2.0 A dm⁻² is likely due to enhanced reduction of metal ions with current density while the CeO₂ particles are transported by mechanical agitation from bulk to the cathode under diffusion control. Thus the process is dominated by metal deposition leading to the observed decrease in particle content of the coating. Kasturibai and Kalaignan (2014) reported that at higher current densities nickel content in the coatings decreases due to increase in hydrogen evolution which causes a decrease in nickel electrodeposition efficiency. Similar results have also been observed in our earlier investigations (Chaudhari and Singh, 2014; Singh and Singh, 2011).

3.2 XRD analysis

The effect of CeO₂ particle incorporation on the growth of Ni-Fe alloy crystals was studied by X-ray diffraction. The XRD patterns of CeO₂ nanopowder and Ni-Fe/CeO₂ nanocomposite prepared at various current densities are illustrated in Figs. 2 and 3, respectively, and results are listed in Table 1. The XRD pattern of CeO₂ powder reveals typical intense lines appearing at 2θ = 28.59 (1 1 1), 33.13 (2 0 0), 47.55 (2 20 ), 56.43 (3 1 1), 59.20 (2 2 2), 69.53 (4 0 0), 76.83 (3 3 1), 79.30 (4 2 0) and 88.50 (4 2 2) corresponding to the fcc structure with a = b = c = 5.403 Å (ICDD: 65-5923). The XRD patterns of the as deposited composite indicate that at lower current densities, the coatings are found to exhibit (1 1 1) preferred orientation while a shift in orientation from (1 1 1) to (2 2 0) crystallographic plane is observed for the coatings obtained at a higher current density. Homogeneously distributed CeO₂ particles play an important role in the crystal growth and the observed change in the texture of the coating can be attributed to the following reasons: (i) inert CeO₂ particles (or Ce⁴⁺ ions) that are partially adsorbed on the surface of the cathode during the electrodeposition process give rise to a shielding effect surrounding the growth center and thus alter the direction of crystal growth; (ii) Ni²⁺, Fe²⁺, Ni[B(OH)₄]⁺ ions are supposed to get adsorbed on the surface of the CeO₂ particles and these cations along with CeO₂ particles are attracted toward the cathode surface which shield the growth center from the incoming metal ions preventing their further nucleation and (iii) with increasing current density more hydrogen is likely to be evolved and the cathode surface got modified resulting in the growth of the coating along the plane with less atomic hydrogen coverage i.e. (2 2 0) plane (Chaudhari and Singh, 2014). Qu et al. (2006) reported that the atomic hydrogen or the combined action of both molecular hydrogen and nickel hydroxide adsorbed onto the cathode surface leads to (2 2 0) preferred orientation. These results are in contradiction to our earlier investigation with In₂O₃ particles (Chaudhari and Singh, 2015) where preferred orientation of the grain growth was observed along (1 1 1) plane at all the current density ranges due to more conducting nature of the In₂O₃ particles. However nanocomposites show similarities with the XRD pattern of the N-Fe alloy although, there is less shifting in preferred orientation toward (2 2 0) plane due to incorporation of CeO₂ particles (Chaudhari and Singh, 2015). Reflections corresponding to CeO₂ particles are feeble due to low amount of incorporation; nevertheless, it can be identified at 2θ angles 80.1 (4 2 0) and 28.59 (1 1 1) in each XRD pattern which substantiates the incorporation of CeO₂ particles in Ni-Fe alloy matrix. The Ni-Fe phase gives lattice parameter between 3.522 and 3.547 Å corresponding to fcc Ni-Fe alloy. It is worth to mention that although crystallite size of the deposits does not show much variation at higher current densities due to almost constant iron content, the CeO₂ particles have grain refinement effect on alloy matrix resulting in smaller crystallite size (7–15 nm) than the Ni-Fe alloy (16–18 nm) (Chaudhari and Singh, 2015). However, the coating obtained at 1.0 A dm⁻² has small crystallite size which seems to be due to synergistic effect of CeO₂ particle incorporation and higher iron content which has a profound grain refining tendency (Cheung et al., 1995). It is observed from Table 1 that the calculated strains of all the deposit are almost negligible (0.0033–0.0095).

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Figure 2

XRD pattern of CeO₂ particles.

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Figure 3

XRD pattern of as deposited Ni-Fe/CeO₂ nanocomposites prepared at current density: (a) 1.0 A dm⁻², (b) 2.0 A dm⁻², (c) 3.0 A dm^−2, (d) 4.0 A dm⁻² and (e) 5.0 A dm⁻².

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Fig. 4 depicts the diffraction patterns of the samples annealed at 800 °C for 1 h and related data are reported in Table 2. CeO₂ reflections can be distinguished in all these XRD patterns. It is evident from the figure that as a result of change in the crystalline structure after annealing there is preferred growth of the deposit along (2 2 0) plane although there is not any phase change. The peaks emerged remarkably sharp due to recrystallization and grain growth. These micro-structural changes can be easily recognized by SEM observations. The lattice constant of the coating was found slightly higher (3.555–3.560 Å) than that of the as deposited specimen and crystallite size calculations shows that annealed deposits have larger crystallite size (20–45 nm). This increase in the crystallite size may be attributed to the recrystallization and grain growth. However, coatings obtained at 3.0 A dm⁻² show large crystallite size probably due to more particle content than the coatings obtained at 1.0 and 5.0 A dm⁻² resulting into more grain growth along (2 2 0) plane. Kotan et al. also reported inconsistent result for the determination of grain size of Ni-Fe alloy by the Scherrer equation at higher annealing temperatures (Kotan et al., 2012). Observed higher value of the lattice constant may be due to stress relief after annealing and probable formation of a little Ni₃Fe inter-metallic compound in nanocrystalline Ni-Fe alloy because of faster grain boundary diffusion in comparison to conventional alloys (Li and Ebrahimi, 2003). The calculated strain of the annealed specimen was found lower (0.0009–0.0029) in comparison to the as deposited coatings due to grain boundary relaxation and release in the stress.

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Figure 4

XRD pattern of annealed (800 °C for 1 h) Ni-Fe/CeO₂ nanocomposites prepared at current density: (a) 1.0 A dm⁻², (b) 3.0 A dm⁻² and (c) 5.0 A dm⁻².

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3.3 SEM study

Morphological examination of the coatings was done by field emission scanning electron microscope (FESEM) and the representative images from surface and cross-section of the samples (Fig. 5) indicate that the morphological and structural characteristics of the alloy matrix are strongly influenced by the presence of CeO₂ particles. Well-dispersed CeO₂ nanoparticles in the coatings significantly contribute to the filling the interfacial boundaries for structural integrity and improvements of surface densification and mechanical properties. At lower current density (1.0–2.0 A dm⁻²) the deposits show a fine grained regular surface with nodular microstructure (Fig. 5a–d) and the morphology is changed to hemispherical grain structure with uniform grain distribution at higher current densities (Fig. 5e–h). This indicates that co-deposited CeO₂ nanoparticles are much uniformly distributed in the alloy matrix with an increase in the current density. The size of CeO₂ particles is so small that it is difficult to clearly differentiate between alloy matrix and the particles although CeO₂ particle rich regions can be identified as slightly bright spheres throughout the coating and somewhere protruding from the coatings which is also evidenced by EDAX analysis. The cross sectional examination of the coating was done to affirm the incorporation and uniform distribution of the particles in the film (Fig. 5i and j). The thickness of the coating, estimated by cross sectional SEM, was found around 20 μm. Surface morphology of the coating after annealing at 800 °C (prepared at current density 3.0 A dm⁻²) is shown in Fig. 5k and l. These images displayed growth of the larger grains (1–4 μm) with closely packed arrangement in the structure at the cost of the smaller grains. The grain boundaries and triple junctions become distinct. The CeO₂ particles can also be identified between these grains due to their segregation after annealing. These results are well supported by the results of XRD analysis.

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Figure 5

SEM images of Ni-Fe/CeO₂ nanocomposites at different current densities-(a) & (b) 1.0 A dm⁻², (c) & (d) 2.0 A dm⁻², (e) & (f) 3.0 A dm⁻², (g) & (h) 4.0 A dm⁻², (i) & (j) cross sectional 3.0 A dm⁻², (k) & (l) 3.0 A dm⁻² (after annealing at 800 °C for 1 h).

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Figure 5

SEM images of Ni-Fe/CeO₂ nanocomposites at different current densities-(a) & (b) 1.0 A dm⁻², (c) & (d) 2.0 A dm⁻², (e) & (f) 3.0 A dm⁻², (g) & (h) 4.0 A dm⁻², (i) & (j) cross sectional 3.0 A dm⁻², (k) & (l) 3.0 A dm⁻² (after annealing at 800 °C for 1 h).

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3.4 AFM study

For more detailed study on surface morphology, roughness and grain analysis of the nanocomposite coating and Ni-Fe alloy, atomic force microscopy (AFM) was employed. It is evident from the AFM images (Fig. 6a and c) that surface morphology of the alloy is more granular and nodular which becomes more compact and smooth after the incorporation of CeO₂ particles (Fig. 6d and f). The grain size analysis of the coatings revealed that the alloy film is covered with large globular grains in size ranging over 34–180 nm but the average grain diameter is 105 nm and most of the grains have their size between 50 and 90 nm (Fig. 6b). However, both the grain size (20–70 nm) and the average grain size (56 nm) become much less for nanocomposite coating, obviously, due to incorporation of CeO₂ particles which might have contributed to different morphology and granular growth of the film (Fig. 6e). To obtain a more quantitative description, the topography of the surface was expressed in terms of roughness factor i.e. root mean square roughness (RMS) and average roughness (ARF). Roughness analysis of the coatings showed that Ni-Fe/CeO₂ nanocomposite coating has lower RMS (3.35 nm) and ARF (2.55 nm) compared to Ni-Fe alloy (4.90 nm and 3.84 nm, respectively). These observations indicate that both the coatings have very low roughness but the composite coating is smoother than the alloy.

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Figure 6

AFM images of Ni-Fe alloy (a: 2D morphology, b: grain analysis & c: 3D morphology) and Ni-Fe/CeO₂ nanocomposites (d: 2D morphology, e: grain analysis & f: 3D morphology).

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3.5 TEM study

To study the structural evolution and crystal transformation of the nanocomposites, bright field TEM characterization was performed along with selected area electron diffraction (SAED) patterns. Fig. 7(a) shows a typical bright field TEM image of a fine grained Ni-Fe/CeO₂ nanocomposite. The mottled contrast indicates that the microstructure of Ni-Fe alloy matrix is composed of crystalline structure with distinct grain boundaries and the average grain size seems to be ∼12 nm (±2 nm). The figure demonstrates that the nanocomposite comprises homogeneously distributed grains without dislocations and defects, which are usually caused by hydrogen occlusion in aqueous medium. CeO₂ particles can also be identified in the micrographs which appear to possess rough surface, and the average particle size is found ∼40 nm. The electron diffraction patterns recorded from the different regions of the coatings can be indexed for fcc Ni-Fe alloy matrix together with fcc ceria (Fig. 7b). The SAED pattern shows the characteristic of fine grained, polycrystalline and randomly oriented material with lattice parameter 3.55 Ǻ (±0.01 Å). The grain size and lattice parameter obtained are consistent with the finding from the XRD data.

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Figure 7

Bright field TEM images showing fine distribution of grains (a) and selected area electron diffraction (SAED) pattern (b) of Ni-Fe/CeO₂ nanocomposite prepared at 2.0 A dm⁻².

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3.6 Microhardness

The variation in microhardness of the coatings obtained at different current densities is shown in Fig. 8. The Ni-Fe/CeO₂ coatings exhibited significantly improved microhardness (up to 900 HV) compared to parent Ni-Fe alloy (410–660 HV) (Chaudhari and Singh, 2014) because of the decrease in the grain size and the increase in deposits compactness. The microhardness of the nanocomposite coating increased initially with increasing current density due to the incorporation of more CeO₂ particles. An increase in microhardness can be associated with the strengthening effect of the nanoparticles by two kinds of mechanisms namely, solid solution strengthening and dispersion strengthening. Formation of solid solution is substantiated by the XRD results, and the improved microhardness of the nanocomposite coating than the Ni-Fe alloy is due to incorporation of CeO₂ particles in the metal matrix. Large improvement in the microhardness is mainly attributed to: (i) formation of a dense structure of the coating resulting its refined microstructure when particles are embedded by suppressing the electro-crystallization and grain growth of metal grains and it will increase the load carrying capacity and the resistance for plastic deformation; (ii) uniform dispersion of CeO₂ phase into the coatings is acting as a softening phase to reduce thermal stresses for the inhibition of coarsening growth which promotes effective bonding interactions with cracking interfaces, thus hindering the sliding dislocations or defective grain boundaries (Zhou and Shen, 2014). With further increase in current density the microhardness of the coatings is almost constant. At higher current densities metal hydroxides may be formed near the cathode and get incorporated in the coating resulting in higher microhardness value (Suzuki et al., 2001). It is also important to note that deposits prepared at higher current densities have almost same grain size and very little change in the iron content resulting nominal variation in the microhardness value. There is only nominal variation in the microhardness of the coatings obtained at different loads i.e. 10 g and 25 g for 5 s due to very little variation in the applied load. The heat treated Ni-Fe/CeO₂ nanocomposite coatings showed a marked decrease in microhardness (225–320 HV) compared to as deposited coatings. This decrease in the microhardness may be attributed to release of stress and grain coarsening due to more mobility of atoms at elevated temperature which favors grain growth resulting in deterioration of the composite microhardness. These results are in accordance with SEM observations and XRD results.

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Figure 8

Variation of microhardness with current density at applied load (a) 10 g for 5 s and (b) 25 g for 5 s.

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3.7 Electrical resistivity

The electrical resistivity of the prepared nanocomposites is dependent on the complex interplay between several factors, including microstructure, grain size, composition and chemical nature. Variation in apparent electrical resistivity of the nanocomposites is illustrated in Fig. 9. It can be seen in the figure that with increase in current density there is a slight decrease in electrical resistivity of the composites and then, it increases with further increase in deposition current density. The effect of particle incorporation on the electrical resistivity can be easily recognized because the coatings show lower electrical resistivity than the electrodeposited Ni-Fe alloy (2.6 × 10⁻⁸–6.97 × 10⁻⁸ Ω cm) (Chaudhari and Singh, 2015), and deposits with maximum CeO₂ particle content (2.0 A dm⁻²) show the minimum electrical resistivity. CeO₂ is an example of mixed electronic-ionic conductor whose conductivity depends on the impurities and oxygen activity in the ambient atmosphere. Electronic conductivity of the deposits can be promoted by hopping between Ce⁺³ and Ce⁺⁴ valence states, while with decrease in size <100 nm, there is a decrease in enthalpy of oxygen vacancy formation resulting in an increase of over 4 orders of magnitude in the conductivity of nanocrystalline CeO₂ (Aruna et al., 2006). Hartmann et al. reported that it is possible to change in nature of nano CeO₂ from mixed ionic electronic conductor to a predominantly electronic conductor with a high total conductivity due to space charge overlap with the grain interior (Hartmann et al., 2013). The increase in electrical resistivity of nanocomposites with increasing deposition current density is ascribed to the grain refining effect of current density which causes an increase in the grain-boundaries and in turn offering increased resistance (Tripathi et al., 2015) which is contradictory to our results. Pores in the samples act as electrical resistances and entail an increase in the total intra-granular void space. It is evident from the surface observations by SEM and AFM that the coatings have a smooth surface with very little porosity resulting into low resistivity. The electrical resistivity of the coatings is also dependent on the microstructure and being almost dislocation and defect free nanostructures (as observed by TEM), the electron scattering is greatly reduced and consequently the electrical resistivity. However, with further increase in current density, electrical resistivity of the coating increases due to grain boundary segregation resulting from diffusion controlled metal deposition. Thus the electrical resistivity of the coatings is altered mainly due to morphological and microstructural changes in the metal matrix after incorporation of CeO₂ nanoparticles.

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Figure 9

Variation in apparent electrical resistivity at different current densities.

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3.8 Magnetic study

The magnetic field dependence of magnetic moment (M-H curve) for the Ni-Fe/CeO₂ nanocomposite (prepared at current density 3.0 A dm⁻²) at applied magnetic field up to 5000 Oe at 300 K is presented in the Fig. 10. Magnetization measurements show that the coating has low coercivity, very low remnant magnetization and high saturation magnetization than the Ni-Fe alloy (Table 3), which indicate the effect of particle incorporation on the magnetic properties. Very low value of coercivity indicates that the film has soft magnetic properties. The shape of the loop, saturation of magnetization at H about 1–2 tesla and high magnetic susceptibility can be recognized as ferromagnetic ordered state of Ni-Fe alloy. The change in values may be ascribed to the change in microstructure, grain size and nature of CeO₂ particles. It is interesting to notice that the nanocomposites have remarkably high magnetic saturation than the alloy showing an improved magnetization due to decrease in crystallite size which results an easier rotation of domain walls and also minimizes the distance between two grains, consequently the inter-granular interactions become more important. It has been also suggested that this change in magnetic behavior results when the valence of cerium changes from +4 to +3 generating an unpaired spin in the cerium f orbital (Zhang et al., 2014). Oxygen vacancies are created during exchange which can magnetically couple with Ce⁺³ ions causing long range ferromagnetic ordering. The Ni and Fe metal ions can also promote the formations of oxygen vacancies, particularly at the surface of nanoparticles to maintain charge neutrality. Recently it has been reported (Swatsitang et al., 2016) that the transition metal doping in CeO₂ creates oxygen vacancies and magnetic ordering which enhances room temperature ferromagnetic behavior. Beside the usual magnetic parameters, the M-T measurements i.e. Field Cooling (FC) and Zero Field Cooling (ZFC) were also performed from 2 K to 310 K at 500 Oe applied field (Fig. 11). It is noticed from the curves that the splitting between FC and ZFC is at least above 310 K which is also a characteristic feature of ferromagnetic ordered state of Ni-Fe alloy. FC curve has higher magnetic moment than ZFC in the entire temperature range and with decreasing temperature both the curves show downwards trend of magnetic moment. An indication of a magnetic transition near 10 K is also observed.

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Figure 10

Hysteresis loop of Ni-Fe/CeO₂ nanocomposite at 300 K.

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Figure 11

M-T (FC and ZFC) magnetization curves obtained at 500 Oe applied field.

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3.9 Polarization behavior and corrosion study

The cathodic polarization curves of Ni-Fe alloy and Ni-Fe/CeO₂ nanocomposite are given in Fig. 12 to investigate the peculiarities of their deposition from ethylene glycol bath. Ni–Fe alloy deposition seems to be sluggish compared to nanocomposite deposition as the current associated with nanocomposite deposition increases sharply than that for the Ni–Fe alloy deposition. The incorporation of CeO₂ particles has a remarkable influence on deposition potential and the curves show that it is more negative for the alloy than that for the nanocomposite indicating thereby that the alloy deposition is associated with considerable polarization. Commencement of deposition of the nanocomposite at nobler potential is much likely due to the availability of more active site on the surface of CeO₂ particles because of improved conductivity caused by hopping of cerium between 3+ and 4+ oxidation states (Aruna et al., 2006), which act as catalytic center to provide more nucleation sites. When they are added to the bath, due to adsorption of nickel ions on their surface their movement toward cathode increases and they will preferentially adsorb on the crystal defects and act as catalytic site for more crystal nucleation (Zhou and Shen, 2013). Thus, it is believed that the CeO₂ particles do not block the cathode surface rather they promote the reduction of metal ions by providing extra sites for deposition. Such depolarization caused by ceramic particles has been observed earlier also (Chaudhari and Singh, 2015; Ibrahim et al., 2013). However, Argirusis et al. (Argirusis et al., 2008) reported that addition of CeO₂ particles to the aqueous plating bath led to deposition process slowdown for Ni-CeO₂ composites due to enhancement of cathodic side reactions.

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Figure 12

Polarization curves for Ni-Fe alloy (a) and Ni-Fe/CeO₂ nanocomposite prepared at 3.0 A dm⁻² (b) deposition in ethylene glycol bath (Conditions of the deposition: Nickel sulfamate 100 g/L, ferrous sulphate 4.0 g/L, boric acid 30 g/L, CeO₂ particle 2.0 g/L, stirring rate 750 rpm, and temperature 34 ± 0.5 °C).

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Polarization curves of as deposited Ni-Fe alloy (Ni 79.6 wt%, Fe 20.4 wt%) and Ni-Fe/CeO₂ nanocomposite in 3.5 g/L NaCl solution at 25 °C are illustrated in Fig. 13. It is observed that the corrosion potential (Ecorr = 80 mV) of Ni-Fe/CeO₂ nanocomposite coatings is more positive with lower corrosion current density (Icorr = 0.095 mA/cm²) compared to Ni-Fe alloy (Ecorr = 40 mV and Icorr = 0.21 mA/cm², respectively). Thus, the corrosion resistance of the Ni-Fe/CeO₂ nanocomposite coating is better than that of the alloy. It can be argued that when nanosized CeO₂ particles are incorporated in the alloy matrix the corrosion process is influenced due to surface modifications of the coating. A plausible explanation can be offered as follows: (i) nano-sized inert CeO₂ phase acts as physical barriers in the initiation and propagation of corrosion process by filling existing crevices, gaps and micro holes on the surface of the coating, most probably, which acts as dielectric phase by impeding the electrical charge transport and ion exchange in turn it reduces the surface area of alloy exposed to corrosive media; (ii) the fine grain structure arising out of co-electrodeposition of CeO₂ nanoparticles leads the deposits to more compactness and thus promotes the corrosion resistance; (iii) both the matrix and particles have similar lattices (fcc) that produce denser microstructure and cause reduction in crystal defects leading to improved corrosion resistance; and (iv) it is also suggested that the incorporation of CeO₂ nanoparticles can influence the kinetics of both the anodic dissolution and cathodic hydrogen evolution reaction by producing many micro corrosion cells which facilitate homogeneous corrosion. The improvement in corrosion resistance of composite coatings on incorporation of CeO₂ particles in nickel matrix has been reported earlier by other workers (Qu et al., 2006) due to distortion in corrosion path and fine grained structure, which is in accordance with our results.

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Figure 13

Corrosion behavior of Ni–Fe alloy (a) and Ni–Fe/CeO₂ nanocomposite (b) coatings in 3.5% NaCl solution (prepared at 3.0 A dm⁻²).

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