Properties of electrodeposited functional Ni–Fe/AlN nanocomposite coatings

2.1 Electrodeposition

Nickel sulphamate tetrahydrate (Fluka-Aldrich), ferrous sulphate heptahydrate (Fisher-Scientific India), boric acid (Qualigens Fine Chemicals India), AlN nanopowder (Sigma–Aldrich, particle size <100 nm) and N,N-dimethylformamide (Fluka-Aldrich) were used in the bath. All chemicals were of analytical grade and used as procured.

The bath composition and electrolysis conditions are given in Table 1. The required amount of electrolytes (nickel sulphamate 24.0 g, ferrous sulphate 4.0 g and boric acid 5.0 g) was mixed together with 100 mL of N,N-dimethylformamide (DMF) in the form of slurry which was then transferred to the electrolysis cell and the volume was made up to 200 mL by adding DMF. After dissolution of the electrolytes, 5 g of the AlN nanoparticles was added to the electrolyte solution. Before electrolysis the solution was stirred continuously with the help of DC regulated mechanical stirrer for 24 h at 45 ± 0.5 °C to get uniform suspension.

The electrolysis cell has an assembly of two parallel pure nickel anodes (2 cm × 2 cm × 0.5 cm) on either side of a centrally placed commercial copper (Cu-99.28% & Ni-0.72%) plate (2 cm × 1 cm × 0.2 cm) cathode. Electrodeposition was carried out using a thermostated water bath (Haake K15 with Haake DC30 control, Netherlands) at 45 ± 0.5 °C employing DC current density in the range of 0.5–5.0 A dm^–2 for 20–90 min. The details of the electrolysis cell assembly, substrate preparation and working procedures are described elsewhere (Tripathi et al., 2015a).

2.2 Characterization

The phase composition and crystallographic structure of the as-deposited films were characterized by X-ray diffraction (XRD, Bruker D8 ADVANCE ECO) using Cu Kα radiation of wavelength 1.5418 Å. X-ray data were collected at a scan rate of 2°/min. The average crystallite sizes of the coatings were calculated from the diffraction peak widths (Scherrer equation), according to the equation:

Morphology of the deposits was examined by Scanning Electron Microscope (SEM, Philips XL20) and Atomic Force Microscope (AFM, NT-MDT SOLVER Next AFM/STM) in semi-contact mode. The compositional analysis was done spectrophotometrically (Tripathi et al., 2013) and it was further substantiated by Energy Dispersive X-ray Analysis (EDAX, JEOL 840A). In EDAX analysis usually 5–10 spots on the sample surface are selected randomly and the values are averaged. The microstructure of the composite coatings was examined under Transmission Electron Microscope (TEM, TECNAI G² FEI) at an accelerating voltage of 200 kV.

The hardness of the nanocomposite coatings was evaluated using Vickers diamond indentation tester (Shimadzu HMV-2) applying 25 g f (0.0980665 N) load for 10 s. All the indentations were made by applying the load perpendicular to the smooth, free surface of the coatings. A minimum of five measurements was taken for a sample to obtain a statistical average value for microhardness.

The magnetic measurements on the nanocomposites were done using fine powder (scraped from the substrate) on a Vibrating Sample Magnetometer (VSM, Lakeshore 735). The M–H data were recorded at 300 K while M–T was recorded between temperatures 500 K and 78 K at 1000 Oe applied magnetic field.

Thermal study of the samples was performed on DSC822^e Mettler Toledo equipment in the temperature range of 30–900 °C at a heating rate of 3 °C min⁻¹. Electrical resistivity of the samples was measured by the four-probe method with a Keithley Nanovoltmeter 2182A and Keithley Source-meter 2400 at 27 °C.

Polarization studies for nickel–iron alloy and composite deposition were carried out at 45 °C using a potentiostat (POS73 Wenking) by step-wise increase of 20 mV potential under mechanical agitation. Cu plate was used as working electrode and Pt foil as counter electrode. The potentials are reported with respect to Saturated Calomel reference electrode (SCE).

3.1 Effect of current density

Out of several deposition parameters such as temperature, electrolyte concentration, particle bath load, and agitation current density is the most important one. It has pronounced effect on the particle incorporation, microstructure, morphology, crystallographic structure, and mechanical, electrical and magnetic properties of the deposits. The effect of applied current density on AlN particle incorporation in the deposits was investigated and is represented in Fig. 1. The figure depicts that initially on increasing current density the wt% incorporation of AlN particle in the deposits increases gradually, reaches the maximum, 23 wt% at 3.0 A dm⁻² and then decreases to the limiting value symptotically.

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Figure 1

Effect of current density on wt.% incorporation of AlN particle in the deposits. (AlN particles 5 g/L.)

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It is conceived from the results that the particle incorporation follows Guglielmi’s model (Guglielmi, 1972). Accordingly the particles suspended in the electrolyte solution get surrounded by the cations. Initially, when current is applied, the cation laden particles present in the vicinity of the cathode get adsorbed on the surface and then simultaneous reduction of cations on the bare cathode and on the surface of the particles results in the strong embrace of the particles in the metal film being deposited. As the current density is further increased, the particle incorporation is also increased, reaching the maximum at 3.0 A dm^–2. A further increase in the current density enhanced reduction of the metal cations while the particle transport from bulk to the cathode becomes diffusion controlled. As a result, the flux of particles compared to the flux of metal ions decreases, consequently the particle content of the deposit decreases.

3.2 Phase structure (XRD analysis)

Fig. 2 shows the X-ray diffraction patterns of AlN nanopowder (Fig. 2a) and the Ni–Fe/AlN composite coatings (Fig 2b–f) prepared at various current densities. The XRD pattern of AlN powder reveals typical broad intense lines appearing at 2θ = 33.22 (1 0 0), 36.01 (0 0 2), 37.93 (1 0 1), 49.80 (1 0 2), 59.37 (2 −1 0), 66.01 (1 0 3), 69.75 (2 0 0), 71.43 (2 −1 2), 72.65 (2 0 1), 76.38 (0 0 4), 81.09 (2 0 2) and 85.87 (1 0 4) corresponding to the hexagonal structure with a = b = 3.109 Å and c = 4.988 Å [JCPDF 65-1902]. On the other hand the XRD patterns of the composite coatings show typical peaks of the substrate, Ni–Fe alloy and few feeble peaks of the AlN. The peaks of matrix can be indexed to (1 1 1), (2 0 0) and (2 2 0) crystal planes of FCC type Ni–Fe alloy.

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Figure 2

XRD patterns of AlN powder (a), Ni–Fe/AlN composite prepared at current density 1.0 A dm^–2 (b), 2.0 A dm^–2 (c), 3.0 A dm^–2 (d), 4.0 A dm^–2 (e) and 5.0 A dm^–2 (f). (AlN particles 5 g/L.)

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The crystallite size of the as deposited samples was calculated from the Scherrer’s equation with (1 1 1) peaks in the XRD patterns (Fig. 2b–f) after subtracting the instrumental broadening (Table 2) which was found around 9–15 nm. The maximum microstrain was found in the composite prepared at current density 3.0 A dm^–2 while the minimum value was obtained for that at current density 1.0 A dm^–2. Observed maximum value of microstrain can be ascribed to the maximum incorporation of the AlN particles causing more deformation of the matrix grains. The results are consistent with our earlier report (Tripathi et al., 2015a,b).

3.3 Morphology

3.3.1

3.3.1 SEM

Fig. 3 represents typical fine granular and nearly smooth surface morphology of the Ni–Fe/AlN deposits obtained at current density 3.0 A dm^–2. The cluster of AlN particles can be seen easily on the surface. There are some holes on the surface which might have formed due to detachment of incompletely adhered AlN clusters.

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Figure 3

SEM morphology of Ni–Fe/AlN composite coating deposited at current density 3.0 A dm^–2. (AlN particles 5 g/L.)

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The cross-sectional SEM image of the coating is shown in Fig. 4. The matrix grains are clearly visible in the deposits (Fig. 4a). There is good adherence of the coatings on the substrate surface. The coating thickness was estimated by cross-sectional SEM and was found between 20 μm and 40 μm. It can be seen from Fig. 4b that the larger grains are of uniform size with semi-spherical or oval shape. The size of these grains varies between 4 μm and 9 μm. Each larger grain also consists of many tiny grains with size less than 0.5 μm and these grains are closely packed.

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Figure 4

Cross-sectional SEM of Ni–Fe/AlN nanocomposite on the substrate (a) and the coating (b) prepared at 3.0 A dm^–2 from a bath having 5 g L^–1 AlN nanoparticles.

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3.3.2

3.3.2 AFM

Fig. 5 shows the AFM image of Ni–Fe/AlN composite coating obtained at current density 2.0 A dm^–2 (Fig. 5a) and 3.0 A dm^–2 (Fig. 5b) from a particle bath load of 5 g L^–1. A compact nodular morphology of the surface can be examined in the figure. The nodules are cluster of a large number of grains and the grains consist of the aggregation of many small crystallites, which is consistent with the results of XRD and SEM. The grain analysis of the surface (10 μm × 10 μm) showed that the clusters are of various sizes ranging between 150 nm and 1105 nm. The grain size obtained from the grain analysis of surface (1 μm × 1 μm) for fine grains, ranged between 9 nm and 34 nm. Roughness analyses of the surface revealed an average surface roughness (R_a) 25.525 nm and the root-mean square roughness (R_q) 33.758 nm.

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Figure 5

AFM image of Ni–Fe/TiN composite prepared at 2.0 A dm^–2 (a) and 3.0 A dm^–2 (b) from a bath with load of 5 g L^–1 of AlN nanoparticles.

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3.4 Microstructure (TEM)

A typical microstructure and corresponding SAED pattern of the as prepared Ni–Fe/AlN electrodeposit are shown in Fig. 6. The bright field image demonstrates fine granular structure of the deposits. The average grain size seems to be ∼14 ± 2 nm which is consistent with the crystallite size calculated from the XRD data. The deposits are dense and compact without dislocations, twins and other defects except grain boundaries (GB).

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Figure 6

Microstructure (TEM) of Ni–Fe/AlN deposit along with SAD pattern. (3.0 A dm^–2; AlN particles 5 g/L.)

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The SAED pattern (inset of Fig. 6) of the deposit consists of continuous rings in a specific manner, which confirms polycrystalline FCC crystallographic structure of the coatings. The rings observed in the SAED pattern are assigned as to correspond to (1 1 1), (2 0 0), (2 2 0), (3 1 1) and (3 3 1) planes of Ni–Fe alloy.

3.5 Microhardness

Microhardness of the composite deposits obtained at different deposition current density is shown in Fig. 7. It can be seen that the microhardness increases with increasing current density up to 4.0 A dm^–2 and decreases thereafter. Among all the deposition variables, current density has the most pronounced effect on the microhardness of the deposits. As can be corroborated from Table 2, the crystallite size of the deposits is not influenced by the applied current density, but as illustrated in Fig. 1, the particle content differs greatly. The initial increase in microhardness is due to the incorporation of AlN particles up to current density 3.0 A dm^–2. It is intuitive that the maximum hardness achieved at 4.0 A dm^–2 is presumably attributed to the synergistic effects of the incorporated particles and to the contribution of grain boundaries due to size confinement effect of current density.

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Figure 7

Variation of microhardness with current density. (AlN particles 5 g/L.)

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Present results can be visualized in light of several reported strengthening mechanisms in MMC based on Al, such as Orowan strengthening, grain boundary strengthening and hardening due to the difference in thermal expansion coefficients between phase constituents (Ferguson et al., 2014; Schultz et al., 2011).

When the grain size becomes less than a certain value in the nano range (less than 50 nm), the grains become free from dislocation, and therefore, the dislocation mediated plastic deformation becomes hindered (Meyers et al., 2006; Kumar et al., 2003). At this nano-size level the contribution of grain boundaries to the microhardness dominates. Grain boundary mediated plastic deformation processes take place through grain boundary sliding, coble creep and grain rotation (Conrad, 2003). Keeping the above facts in view it can be considered that the incorporated AlN particles prevent grain boundary sliding and thus results into increase in the microhardness.

From the crystallite size calculations it is intuitive to arrive that the crystallite size of the matrix is much smaller than the AlN particles; therefore, in the composite coatings, all the AlN nanoparticles cannot be considered embedded entirely inside the matrix grains; instead they could be surrounded by several grains; consequently some of the matrix grain boundaries are replaced with particle–matrix interfaces. At these interfaces the AlN nanoparticles can strengthen the grain boundaries and hinder the plastic deformation (Bakhit and Akbari, 2012).

3.6 Magnetic property

The field-dependent magnetization (M–H hysteresis) curve of the Ni–Fe/AlN deposit (prepared at 2.0 A dm^–2 from a bath load of 5 g L^–1 of the AlN particles) at room temperature is represented in Fig. 8. Typical hysteresis is clearly observed which indicates ferromagnetic behaviour of the deposited composite. The composite has low coercivity (H_c = 28 G), very low magnetic remanence (M_r = 2.22 emu/g) and high saturation magnetization (M_s = 62.02 emu/g). Earlier we have reported the room temperature values H_c, M_r and M_s for the Ni–Fe alloy deposit (Tripathi et al., 2015b) as 61.2 G, 3.07 emu/g and 51.2 emu/g, respectively. In the present study, the values of H_c and M_r for Ni–Fe/AlN deposits are found comparatively very low. The difference in the values can be thought of due to the microstructure of the composite deposit (single domain grain size resulting in strong particle–particle interaction) and the nature of AlN particles. The presence of the hysteresis loop and small coercivity indicates the soft magnetic nature of the deposit but from the shape of the loop a marked signature of superparamagnetism can be recognized. The M_r/M_s value for the deposit is 0.035 which also supports the above inference (Couderchon and Tiers, 1982). The saturation of magnetic moment at very high fields (above 2 kG) reflects the random magnetic anisotropy (Chudnovsky and Serota, 1982) existing in the materials with crystallite size below 20 nm (Moustaïde et al., 1999). These results agree with our earlier observations (Tripathi et al., 2015b).

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Figure 8

M–H curve for NFAN-2 measured at 300 K. (AlN particles 5 g/L.)

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In order to substantiate the results of M–H hysteresis M–T measurements were also performed. The variation of magnetization with temperature at 0.1 T applied field is illustrated in Fig. 9 wherein the Curie temperature (T_C) was identified and was found to be ∼540 K while in the case of Ni–Fe/TiN the value of T_C was found to be 590 K (Tripathi et al., 2015b). The bifurcation of FC and ZFC curves occurs at a temperature 297 K and they never meet below this temperature. The ZFC curve at lower temperatures goes down and above this temperature the curves coincide. Bifurcation between ZFC and FC curves is generally a characteristic feature of superparamagnetism, resulting from competing ferro- and antiferromagnetic pair interactions between the components. More specifically, below the bifurcation temperature irreversibility sets in and below the ZFC-peak temperature, the system passes to a mixed magnetic phase. It has been reported that the FC and ZFC curves start bifurcating at higher temperatures with the decrease in inter-particle separation (Sato et al., 2000; Morup et al., 1995). Above type of bifurcation is an impression of magnetic metastability as a consequence of blocked magnetic moments in superparamagnetic particles below T_B (blocking temperature). From these results it is apparent that the magnetic behaviour is superparamagnetism at the temperature above 270 K. There is a gradual increase in FC magnetization with decreasing temperature. The FC/ZFC measurements show typical superparamagnetic behaviour and the ZFC curve yields a cusp at 270 K, characteristic blocking temperature.

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Figure 9

M–T (FC & ZFC) curve of NFAN-2. (AlN particles 5 g/L.)

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Present results of magnetic measurements can be discussed for magnetic materials, in terms of domain walls which play an important role. Upon application of a magnetic field, initially domain walls will tend to move reversibly, so that the domains that are magnetized with a component in the direction of the field grow at the expense of domains magnetized in opposite directions. Continuing to apply an increasing field the domain walls begin to move from their original positions irreversibly into new positions. Usually in this state the increase in magnetization with field is steeper and reaches the maximum and finally in the saturated magnetic state. Thus, coercivity in such materials generally increases with decreasing particle size or, in fully dense materials, with decreasing grain size. The nonmagnetic second phase particles and the grain boundaries are ought to have pinning effect on domain wall motion in nanocrystalline materials resulting in an increase in the coercivity. Soft magnetic materials usually possess high permeability and low coercive force and in general, their magnetic behaviour is governed by domain wall pinning at heterogeneities such as grain boundaries, surfaces, precipitates and inclusions.

3.7 Electrical resistivity

Electrical resistivity is an important physical property of the coatings which is responsible for many applications. However, the studies on the resistivity measurements of electrodeposited nanocomposite coatings are rarely reported. Therefore, the measurement was done for the Ni–Fe and Ni–Fe/AlN composite coatings. Fig. 10 depicts the measured apparent electrical resistivity of the composite coatings deposited at different current densities. The apparent electrical resistivity of the deposits was found to be of the order 10^–8 Ω cm. It can be seen, in the figure, that the electrical resistivity value of the coatings does not show a regular trend of variation with current density; in general, more or less it decreases with an increase in the deposition current density except at 5.0 A dm^–2. A correlation of electrical resistivity (Fig. 10) of the composite deposits with the particulate content (Fig. 1) and the crystallite size (Table 2) indicates that the particulate content in the deposits does not seem to affect the resistivity of the coatings, as very little variation in crystallite size of the coatings i.e. only few nm is observed; however, it is known that below a critical size limit (below 20 nm) any alteration in crystallite size poses drastic change in the properties of the nanomaterials (Cao, 2004). Therefore, explanation for the observed trend in the electrical resistivity of the composite can be alternatively looked into their crystallite size. The higher value of resistivity at initial deposition current density can be due to the larger crystallite size of the deposit where the grain boundaries offer increased resistance. As the deposition current density increases, the crystallite size decreases. Once the crystallite size becomes much below the critical limit, though the density of grain boundaries increases yet the crystallites/grains offer hopping conductance resulting in declined resistivity. The highest value of electrical resistivity offered by the deposits at the highest deposition current density seems due to the grain boundary segregation resulting from the diffusion controlled metal deposition. It is worth noting that the resistivity of the Cu-substrate and the Ni–Fe alloy (3.0 A dm^–2) was measured to be 2.5 × 10^–7 Ω cm and 4.97 × 10^–8 Ω cm, respectively. Thus a comparison of the apparent electrical resistivity between bulk Cu, Ni–Fe alloy and the Ni–Fe/AlN nanocomposites depicts that the nanocomposites offer much less resistivity than the bulk Cu and Ni–Fe alloy (Chaudhari and Singh, 2015; Tripathi et al., 2015a,c). Furthermore, there is almost marginal decrease in the electrical resistivity of the Ni–Fe matrix on inclusion of the nano-AlN particles. Thus it can be conceived that there is negligible effect of AlN particle incorporation in the Ni–Fe alloy matrix.

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Figure 10

Apparent electrical resistivity of composites deposited at different current densities. (AlN particles 5 g/L.)

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Since resistivity has been measured at room temperature, therefore, the surface scattering may be ignored which is a temperature dependent phenomenon. In the present investigation the coatings have been prepared under such conditions that the presence of residual stress and dislocation is expected to be greatly reduced. Again, since the bath and the particles were fairly stable, the occlusion of any decomposition product in the composite or at the interfaces can be safely ignored. Such interfaces are known to decrease electrical resistivity (Singh and Singh, 2014).

The effect of size of nanostructures on electrical resistivity (conductivity) is complex (Cao, 2004) and based on surface scattering, quantized conduction, ballistic conduction, coulomb charging, tunnelling, size of band gap and change in microstructure. Moreover, in near absence of the literature pertaining to the resistivity of MMNCs, it is difficult to extend a definite mechanism to explain the observed results.

3.8 Electrochemical study

3.8.1

3.8.1 Polarization

In order to determine deposition potential of Ni–Fe alloy and the Ni–Fe/AlN nanocomposite, cathodic polarization studies were performed. Fig. 11 shows cathodic polarization curves for Ni–Fe alloy and Ni–Fe/AlN composite deposition. The curves show that the deposition potential of Ni–Fe alloy is ∼600 mV cathodic (vs. SCE) while in the presence of AlN particles it shifts towards less cathodic direction by ∼50 mV. From the polarization studies it is arrived that on addition of AlN nanoparticles to the Ni–Fe plating solution the current density of the Ni–Fe alloy deposition increases. A similar depolarization was observed in the presence of ceramic particles by several workers (Ibrahim et al., 2013; Tripathi et al., 2015a; Benea, 2009). However, such behaviour is not the same with other composite electrodeposition systems where addition of ceramic particles to the plating bath led to higher polarization of the deposition process (Vaezi et al., 2008; Shi et al., 2006). Therefore, from the present observed results it is enunciated that the AlN nanoparticles have facilitating/promoting effect on the Ni–Fe alloy deposition. It is speculated that when the particles are added to the plating bath, the metal cations are adsorbed on their surfaces and thus they increase the ionic transport towards cathode; consequently, the deposition potential shifts towards positive potentials and the associated current density increases due to the increased rate of metal deposition (Tuaweri and Wilcox, 2006). Secondly, the insulating AlN particles cover the cathode surface and thus the active cathode surface area decreases which can increase the associated current density.

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Figure 11

Cathodic polarization curves for Ni–Fe alloy and Ni–Fe/AlN composite deposition. (AlN particles 5 g/L.)

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3.8.2

3.8.2 Corrosion

To compare the corrosion resistance of the Ni–Fe alloy and Ni–Fe/AlN nanocomposite cathodic and anodic polarization studies were performed. The polarization curves for Ni–Fe alloy and Ni–Fe/AlN composite coatings (both prepared at 3.0 A dm^–2) in 3.5% NaCl solution are given in Fig. 12. The alloy and the composite both exhibited similar polarization curves with active and poor passive characteristic. The corrosion potential (E_corr) and corrosion current (I_corr) were determined extrapolating cathodic and anodic curves. The E_corr for Ni–Fe alloy was found to be −600 mV (SCE) and that for the Ni–Fe/AlN coating was −720 mV (SCE). Both the alloy and the composite show the tendency of passivation which might be due to the formation of a salt film on the surface. The I_corr for Ni–Fe coating was measured to be ∼0.029 mA/cm² while that for the Ni–Fe/AlN nanocomposite coating was ∼0.015 mA/cm². These current values indicate a marked decrease in corrosion of Ni–Fe/AlN nanocomposite coating compared to the Ni–Fe alloy coating. The improvement in corrosion resistance on incorporation of ceramic particles has been reported earlier by several workers (Zhu et al., 2011; Vaezi et al., 2008; Aruna et al., 2006; Li et al., 2005; Shi et al., 2006) which is in accordance with our results. On the other hand a reduction in corrosion resistance has also been reported in the presence of ceramic particles in the plating solution (Tripathi et al., 2015b; Lekka et al., 2009). It emerges out from the above discussions that the observed improved corrosion resistance of the Ni–Fe/AlN nanocomposite coating over Ni–Fe alloy, in the present investigation, can be due to the positive effect of the incorporated AlN nanoparticles which being insulating in nature act as inert physical barrier to the initiation and development of defects. Additionally, the particles reduce the corrosion active surface area resulting in increased corrosion resistance of the nanocomposite coating. Moreover, the electrical resistivity of the Ni–Fe/AlN nanocomposite has been found to be lower than that of the Ni–Fe alloy which implies a more tendency towards corrosion. However, the results are unlike, which can be most likely caused by change in local microstructure around the occluded particles and their inertness. A comparison of Ni–Fe/AlN and Ni–Fe/TiN (Tripathi et al., 2015b) nanocomposite deposits shows that Ni–Fe/AlN nanocomposites are more resistant to corrosion than both the Ni–Fe alloy and Ni–Fe/TiN composite coatings.

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Figure 12

Corrosion behaviour of Ni–Fe alloy and Ni–Fe/AlN composite coatings. (3.0 A dm^–2; AlN particles 5 g/L.)

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