Protective coating on Mg AZ91D alloy – The

Corrosion Science 49 (2007) 1777–1798

www.elsevier.com/locate/corsci

Protective coating on Mg AZ91D alloy – Theeffect of electroless nickel (EN) bath stabilizers

on corrosion behaviour of Ni–P deposit

Woo-Jae Cheong a,*, Ben L. Luan a,b, David W. Shoesmith a

a Department of Chemistry, University of Western Ontario, Chemistry Building, 1151 Richmond St. London,

Ont., Canada N6A 5B7b Integrated Manufacturing Technologies Institute, the National Research Council Canada, London, Ont.,

Canada N6G 4X8

Received 1 November 2005; accepted 21 August 2006Available online 25 January 2007

Abstract

The effects of the electroless nickel (EN) bath stabilizers; thiourea (TU) and maleic acid (MA), onthe corrosion of the EN coatings on Cu and Mg substrates were investigated in an aggressive oxy-genated chloride environment.

Corrosion was limited to the external EN layer, and thus, the intermediate Cu layer as well as MgAZ91D substrate remained intact. On the EN surface, two different types of corrosion morphologieswere observed. Specimens prepared with TU showed preferential attack along particle boundaries,while pitting corrosion was observed for specimens prepared either without a stabilizer or withMA. The latter deposits showed higher corrosion resistances due, in part, to the formation of aP-enriched layer on the surface. The poorer corrosion performance of EN specimens prepared withTU was attributed to the presence of the particle boundaries and the trace impurities adsorbed on,or included in, them. The key impurity was sulphur incorporated during formation of deposits inTU-stabilized baths.

It was clearly shown that the corrosion behaviour of the EN specimens was predominantly influ-enced by the morphological and surface compositional aspects of the deposits.Crown Copyright � 2006 Published by Elsevier Ltd. All rights reserved.

0010-938X/$ - see front matter Crown Copyright � 2006 Published by Elsevier Ltd. All rights reserved.

doi:10.1016/j.corsci.2006.08.025

* Corresponding author. Tel.: +1 519 661 2111x86248; fax: +1 519 661 3022.E-mail addresses: [email protected], [email protected] (W.-J. Cheong).

mailto:[email protected]

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1778 W.-J. Cheong et al. / Corrosion Science 49 (2007) 1777–1798

Keywords: A. Magnesium AZ91D alloy; Metal coating; B. EIS; SEM; C. Corrosion electrochemical behaviour

1. Introduction

Magnesium (Mg) and its alloys are plated for various reasons, but primarily to improvetheir corrosion resistance. While Mg is relatively stable under alkaline conditions due tothe rapid formation of the hydroxide, Mg(OH)2, it is subject to severe attack in acidicand chloride environments. The metal is also prone to galvanic attack when in contactwith another metal [1]. Various methods for applying protective coatings, such as electro-chemical plating, chemical conversion coating, hydride coating, anodizing, gas phasedeposition, and organic/polymer coatings have been tried [2]. Since it requires a minimumcapital investment, chemical conversion coating has been extensively investigated. How-ever, the use of CrO3 or chromate solutions for the conversion process is harmful to theenvironment. Moreover, a conversion coating can only provide a suitable base for anadherent organic coating [3].

Electroless deposition can also provide a protective coating for common metal sur-faces [1,4,5]. In general, an electroless nickel (EN) deposit is less porous and more uni-form than an equivalent electrolytically plated deposit, and thus provides bettercorrosion and wear protection [6]. However, electroless deposition has the inherent prob-lem of sudden bath decomposition, which can be aggravated in the case of reactive sub-strate plating (i.e. Mg alloys), since the dissolved metal ions enhance decomposition. Thisresults in both an increase in costs and the production of environmentally hazardous pol-lutants due to large waste generation [5]. To overcome this problem, EN baths com-monly contain small amounts of stabilizers to prevent the homogeneous reactions thattrigger decomposition [5]. Previously, the stability of EN baths were evaluated with anaccelerated stability test in which the bath decomposition time was measured after theaddition of the optimum concentration of PdCl2 (40 mg/L) solution [8]. The EN bathstability was increased significantly with the addition of either thiourea (TU) or maleicacid (MA). The bath life increased from 1.3 ± 0.5 min to >264 min for an increase inTU concentration from 0 mg/L to 2.0 mg/L and from 0.7 ± 0.1 min to 101 ± 9 minfor MA concentrations from 0 to 3000 mg/L [8]. While enhancing bath stability, stabi-lizers may also influence deposit morphology [7,8], composition [8–11], and microstruc-ture [8]. These changes would be expected to affect the corrosion behaviour of the EN(Ni–P) coating.

The corrosion behavior of protective Ni–P coatings has been extensively investigatedand is considered to be dependent on many processes and features. These include thefollowing:

1. The process used to apply the Ni–P coating (i.e. electrolytic, electroless, and meltingand quenching processes).

2. The composition of the Ni–P coating, especially its P content [12–14].3. The homogeneity of elemental distribution, especially P distribution in the coating sur-

face [12], and throughout the deposit [15].4. The surface morphology of the deposit: i.e. whether it is smooth, nodular, or possesses a

cauliflower-like morphology [16].

W.-J. Cheong et al. / Corrosion Science 49 (2007) 1777–1798 1779

5. The microstructure of the deposit: i.e. whether it is crystalline, microcrystalline, amor-phous, or composed of mixed phases [17].

6. The exposure environment, especially the pH [12].7. The presence of impurity inclusions and their distribution in the deposit; e.g., Pb or S

inclusions [12,18].8. The nature of the substrate under the deposit [19].9. The presence of surface defects such as microscopic pores or pits [18].

Although the corrosion behaviour of protective Ni–P deposits have been extensivelystudied, only limited information is available for Ni–P deposits prepared in the presenceof the common bath stabilizers, TU and MA. In this paper, we have studied the corrosionof such deposits over longer time periods in an aggressive chloride environment. Ni–Pcoatings on both copper substrates and on magnesium AZ91D alloy substrates (with anintermediate copper layer produced by a novel Cu immersion process) have been studied.

2. Experimental

2.1. Preparation of Ni–P deposits

Ni–P binary alloy deposits with P contents ranging from 6% to 9% (wt.%) were preparedon a dumbbell shaped copper (99.9%) substrate in a plating bath containing 30 g/LNiSO4 Æ 5H2O, 20 g/L NaH2PO2 Æ xH2O, 20 g/L CH3COONa, and various concentrationsof the stabilizers, thiourea (TU) or maleic acid (MA). Six different Ni–P alloy specimenswere prepared with different TU (0.5 mg/L, 0.9 mg/L, and 2.0 mg/L) and maleic acid(180 mg/L, 1000 mg/L, and 3000 mg/L) concentrations. A reference specimen (RE) wasprepared without any stabilizer. These specimens were coded with the name of the stabilizerfollowed by their concentration. For example, the EN specimen prepared in the solutioncontaining 0.5 mg/L of TU is coded as TU0.5. When prepared on a Mg AZ91D substratethe EN specimen is designated as TU0.5 on Mg. During plating, the EN bath was slowlyagitated to minimize local acidity at the substrate surface. The deposition conditionsemployed were a pH of 4.5, a temperature of 85(±1) �C, a deposition time of �6.0 h,and a specimen surface area to volume ratio of 7.8 cm2/L. A weight gain analysis was doneon a known surface area of specimen to determine the average thickness of the coating. Toensure the Cu substrate or the Cu intermediate layer were not exposed during subsequentcorrosion tests, the thickness of the coating was made >50 lm, as suggested by Innes [20].

For the Mg AZ91D alloy substrate, an intermediate Cu layer was required to facilitatedeposition of the Ni–P coating. This intermediate coating was deposited using a novel Cuimmersion process in a bath containing 0.67 M CuSO45H2O and 3.3 M HF followed by asurface sealing process.

2.2. Electrochemical measurements

The electrolyte used was a 5% NaCl solution prepared with reagent grade NaCl anddeionized water (resistivity >18 MX) from Millipore (Milli-Q academic) system. A stan-dard three-compartment cell was used with a saturated calomel electrode (SCE) and a2 · 3 cm2 platinum electrode as a reference and counter electrode, respectively. Beforeintroduction of the specimen to the corrosion cell, the solution was purged with Argon


gas for >30 min. The specimen is only partially immersed in the electrolyte solution toavoid crevices. Before measurement of the open circuit potential, the electrode was cathod-ically polarized at �0.8 VSCE to remove any air-formed oxide on the surface of the Ni–Pcoating. Corrosion potential (Ecorr) and intermittent electrochemical impedance spectros-copy (EIS) measurements were then made, first in deaerated (argon) (50 ± 1 h), and thenin oxygenated (145 ± 1 h) environments.

EIS measurements were performed with a Solartron 1255 Frequency Response Ana-lyzer coupled to a 1480 Multistat. The EIS measurements were obtained by applying asinusoidal perturbation of ±5 mV, and measuring the modulus of impedance and thephase shift over the frequency range from 100 kHz to 1 mHz. The electrode potentialfor each sample was held at its open circuit value during EIS measurements. Curve fittingof the EIS data was performed using Zview software, which uses the Levenberg–Marqu-ardt non-linear least squares method. This method is based on the complex non-linearleast squares (CNLS) program written by Macdonald [21].

2.3. Surface analysis

Surface analysis was performed before and after corrosion experiments. The surfacemorphology of the Ni–P deposits was investigated using scanning electron microscopy(Hitachi S3500N electron microscope). Energy dispersive X-ray spectrometry (IncaEnergy400, Oxford Instrument Ltd.) was used to estimate the Ni and P contents (wt.%). Theentire 1.2 · 1.2 mm2 area was scanned to enhance the reliability of the measurements ofaverage P content.

3. Results and discussion

3.1. Selection of Ni–P coated specimens

The P content of the electroless deposits is shown as a function of stabilizer concentra-tion in Fig. 1. The stabilizer concentrations used to produce specimens for corrosion test-ing were selected so that the difference in P content was insignificant (�±0.5 wt.%) overthe range studied. The only exception was the deposit prepared using TU2.0 that contains�1.0 wt.% less P than generally observed. Thus, with this exception, differences in the cor-rosion behaviour are unlikely to be attributable to differences in the P content of thedeposits.

3.2. Surface layer composition (XPS)

In a previous publication [22], X-ray photoelectron spectroscopy (XPS) was used toexamine the composition of films formed on similar EN deposits after corrosion underthe same conditions as those used here. Two distinct corrosion films were observed. Anenrichment of elemental phosphorus compared to nickel was observed on the surface ofEN deposits produced with either no stabilizer or maleic acid. By contrast, extensive sur-face oxidation was observed on EN deposits prepared with TU. Trace impurities of sulfurwere blamed for the extensive build-up of corrosion products on EN deposits formed withTU, whereas a P enriched chemical passivation layer was formed on deposits preparedwithout any stabilizer and with MA.

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Fig. 1. Phosphorus content in a Ni–P binary alloy produced by EN deposition with different types and amountsof the stabilizer.


3.3. Surface morphology and composition (SEM/EDS)

‘As deposited’ surfaces of the EN specimens are shown in Fig. 2(A1)–(G1). The surfacescan be categorized into two types; (a) surfaces with particle boundaries and (b) smoothsurfaces. Deposits formed in an unstabilized bath or in a bath containing larger than1 g/L MA exhibited smooth surfaces with a number of small defects (Fig. 2(A1), (F1)and (G1)). For surfaces with particle boundaries, slightly different structures wereobserved depending on the type of stabilizer and its concentration. The MA180 specimenhad regular small nodular features �100 nm in diameter (Fig. 2(E1)). The TU0.5 specimenshowed non-uniformly ordered particle boundaries (Fig. 2(B1)) while very clear uniformlyordered boundaries were produced with TU0.9 (Fig. 2(C1)). A cauliflower like surfacewith many boundaries was observed on the TU2.0 specimen (Fig. 2(D1)).

After immersion in 5% NaCl solution for 195 h (Ar-purged (50 h) and O2-purged(145 h)), two distinctly different corrosion morphologies were observed (Fig. 2(A2)–(G2)). For specimens with particle boundaries, preferential attack occurred along theboundaries (Fig. 2(B2)–(E2)), and the severity of attack increased with an increase inthe concentration of TU. The blurred SEM micrograph (Fig. 2(D2)) observed for theTU2.0 specimen may indicate the presence of a thick corrosion/oxide layer on this surface.By contrast, evenly distributed small pits with a high number density were formed on thesurfaces of smooth deposits (Fig. 2(A2), (F2), and (G2)).

These differences in morphology of corrosion attack were correlated with the twodistinctly different surface layer compositions determined by XPS [22]. EDS analysis con-firmed the higher O content and lower P content (compared to the bulk P content of�6 wt.%) obtained on the corroded TU2.0 specimen prepared on Mg AZ91D substrate(Fig. 3). For the corroded RE and MA3k specimens prepared on a Mg AZ91D substrate,the P and Ni contents were comparable to the bulk values (Fig. 3). In addition, an SEM

Fig. 2. The SEM micrograms of Ni–P deposits surface prepared with various concentrations of TU and MA: (A)no stabilizer; (B) 0.5 mg/L TU; (C) 0.9 mg/L TU; (D) 2.0 mg/L TU; (E) 180 mg/L MA; (F)1.0 g/L MA; (G) 3.0g/L MA. The SEM micrograms of each specimen were obtained (1) before and (2) after corrosion experimentsin 5% NaCl purged with Ar (48 h) and O2 (145 h).


Fig. 2 (continued)


investigation of the specimens (Fig. 4(A)–(E)) after the 50 h Ar purged period indicatedthat the pitting and intergranular attack initiated during the Ar-purged period.

3.4. Corrosion potential (Ecorr) measurements and electrochemical impedance spectroscopy

(EIS)

3.4.1. General corrosion behaviour and equivalent circuit modelsThe corrosion potentials (Ecorr) of different specimens were monitored in deaerated (Ar)

(50 ± 1 h) and then oxygenated (145 ± 1 h) 5% NaCl solution (Fig. 5). In addition, EISmeasurements were acquired periodically for every specimen (Fig. 6). For all specimens,the absolute impedances at low frequencies (10�3–10�2 Hz) were much higher for

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TU2.0 on Mg 8.05 3.54 88.41 MA3k on Mg 0.51 6.30 93.19

Fig. 3. EDS spectrum of EN deposits prepared with different stabilizers (RE – no stabilizer, TU – thiourea, andMA – maleic acid) on an Mg AZ91D alloy substrate (with a Cu intermediate later) after immersion in Ar-purged(50 h) and O2-purged (145 h) 5% NaCl solution. The table shows the elemental content of each specimen.


Ar-purged conditions than for O2-purged conditions (Fig. 6(A2)–(G2)). The increase inthe Ecorr (Fig. 5) observed on switching from Ar to O2-purged conditions, can be attrib-uted to the increased polarization of the corrosion process by oxygen reduction leadingto an increase in corrosion rate. In addition, the introduction of oxygen leads to less noisein the Ecorr. The noise amplitude was ±5 mV for Ar-purged conditions, but only ±0.1 mVfor the O2-purged condition. The Ecorr of the uncoated Mg AZ91D alloy substrate wasextremely negative (�1.5 VSCE, Fig. 5), consistent with its instability in water. In addition,the introduction of oxygen did not change Ecorr, indicating that the oxygen reduction reac-tion does not influence the already rapid corrosion of the Mg AZ91D alloy (Fig. 5) in 5%NaCl solution.

Inspection of the EIS spectra suggests three different equivalent circuit models arerequired to fit the results (Fig. 7). A one time constant (s) model (Fig. 7(a)) is appropriatefor smooth deposits (RE, MA1k or MA3k) corroding under Ar-purged conditions(Fig. 6(A), (F), and (G); plots in black). Such a circuit suggests a uniform corrosionprocess on smooth deposits. The development of small pits in the latter stages of theAr-purged period required the fitting of the spectra with a 2 s-circuit (Fig. 7(b)), consistentwith a corrosion process occurring predominantly at the base of pits. A 2 s-circuit modelwas also required to fit the spectra measured on specimens with particle boundaries whenpreferential corrosion along particle boundaries occurred (Fig. 6(B)–(E)). Lastly, a 2 s cir-cuit with an added Warburg circuit element (Fig. 7(c)) was required to fit the spectra gen-erated from specimens prepared with >0.9 mg/L TU after prolonged exposure tooxygenated solution (Fig. 6(C) and (D); plot in olive). Such a model is consistent with a

Fig. 4. The SEM micrograms of Ni–P deposits prepared with various concentrations of TU and A: (A) nostabilizer; (B) 0.5 mg/L TU; (C) 2.0 mg/L TU; (D) 1.0 g/L MA; (E) 3.0 g/L MA. The SEM micrograms wereobtained after immersion in 5% NaCl purged with Ar (48 h).


corrosion process occurring preferentially at the base of deep particle boundary channelsin which O2 transport is partially rate controlling.

3.4.2. Ar-purged condition

For the deposit prepared in an unstabilized bath, Ecorr rapidly established a value of�400 mV to �425 mV followed by a slow rise (Fig. 5). A deposit grown in the presenceof a low TU concentration (TU0.5) exhibited similar behaviour, but as the TU concentra-tion was increased, an Ecorr value in the range �450 mV to �500 mV was established moreslowly and no subsequent gradual increase was observed (Fig. 5). The MA180 specimenbehaved similarly to these last specimens. Inspection of Fig. 2 shows this latter behaviour

Fig. 5. Corrosion potential of deposits formed in differently stabilized baths exposed to 5% NaCl solution.


is consistently observed with deposits comprising fine and nodular features (Fig. 2(C1),(D1), and (E1)). The Ecorr behaviour in the MA1k specimen was similar to that observedfor the RE specimen. For the MA3k specimen, Ecorr was initially very low but eventuallythe Ecorr rose as if undergoing a transition (Fig. 5).

A comparison of values after the Ar-purging period shows Ecorr shifted to less noblevalues with an increase in TU concentration from 0 mg/L to 2.0 mg/L, Fig. 8. By contrast,Ecorr values for the MA1k and MA3k specimens (Fig. 8) were closer to the Ecorr for the REspecimen. The onset of pitting during the Ar-purging period forced the use of a 2 s-circuitto fit the impedance results. Fig. 8 also shows the total surface resistances determined fromthe low frequency limit of the impedance spectra. The resistances (Rt) clearly correlatewith the Ecorr values, with specimens with Ecorr < �0.45 V possessing lower resistance thanthose with Ecorr P �0.40 V. This indicates that the positive shift in Ecorr under Ar-purgedconditions leads to an enhanced corrosion resistance, suggesting the formation of a corro-sion resistant surface film/layer. The increases in Ecorr and total resistance for the RE andthe MA3k specimens can be attributed to the gradual inhibition of the anodic reaction dueto the build-up of a P-enriched layer containing PO3�

4 , H2PO�2 , and P0, as indicated byXPS [22]. It is likely this build-up is due to the selective dissolution of Ni at the beginningof the immersion period [17].

The formation of a hypophosphite film has also been reported on both electroless Ni–P(12.8% and 8.0%) alloys under natural corrosion conditions in 3.5% NaCl at pH of 10 [4]and on electroless Ni–P (13.2%) deposits in a 0.2 N HCl solution [23]. The oxidation ofphosphorus to P1+ was considered to occur as follows [24]:

Pþ 2H2O ¼ H2PO�2 þ 2Hþ þ e� ð1ÞThe P1+ species can then be further oxidized to a P5+ species [24]

H2PO�2 þ 2H2O ¼ H2PO�4 þ 4Hþ þ 3e� ð2Þand stabilized on the metal surface by an Ni2+ cation.


By comparison, XPS analysis showed a low P/Ni ratio on the TU2.0 specimen [22], andthe low total resistance indicated that the anodic reaction (i.e. Ni! Ni2+ + 2e�) was notinhibited in this case. The TU0.5 specimen exhibits a similar high resistance to the unsta-bilized system, Fig. 8, but in this case XPS showed the inhibiting surface possessed a highO/Ni ratio and only a small P/Ni ratio. The high resolution Ni 2p3/2 peak indicated mostof nickel species on the surface were bound to hydroxide [22]. The MA1k specimen exhib-ited a particularly high Rt value suggesting the presence of an air-formed oxide, incom-pletely removed by the cathodic cleaning step.

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Fig. 6. Bode plots recorded after various periods of corrosion in 5% NaCl solution with Ar (50 h) or O2 (145 h)purging for deposits produced from differently stabilized EN baths: (A) RE; (B) TU0.5; (C) TU0.9; (D) TU2.0;(E) MA180; (F) MA1k; (G) MA3k.

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Fig. 6 (continued)


Fig. 7. Equivalent circuit models for electroless Ni–P deposits exposed in 5% NaCl solution. (a) Uniformcorrosion film on the Ni–P surface, (b) corrosion film with pores of low aspect ratio and (c) corrosion film withpores of high aspect ratio.

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Fig. 8. Corrosion potential and total resistance (from EIS measurements) for deposits produced with differentamounts of stabilizer: MA – maleic acid; TU – thiourea; Ref – no stabilizer; 50 h immersion in Ar-purged 5%NaCl.


3.4.3. O2-purged condition

The onset of O2 purging resulted in a shift of Ecorr to more positive values for all thespecimens (Fig. 5), as expected. After the initial increase, Ecorr varied for individual spec-imens. Ecorr for the MA3.0, MA1.0 and RE specimens rapidly stabilized (Fig. 5), while the


TU specimens relaxed to lower values (Fig. 5). This relaxation was most marked for theTU0.9, TU2.0 and MA180 specimens, and suggests an accelerated corrosion reaction,which is consistent with the significantly lower Rt values obtained for these specimens,Fig. 9.

The behaviour of the three specimens prepared in TU makes an instructive series. Therelaxation in Ecorr is much less marked for TU0.5 compared to TU0.9 and TU2.0 (Fig. 5),and so is the decrease of Rt with time (Fig. 9). Inspection of Fig. 2(B)–(D) shows that theextent of intergranular attack is also lower for the TU0.5 specimen. However, the decrease

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Fig. 9. Total resistances from EIS measurements of the deposits produced from differently stabilized baths aftervarious periods of corrosion in 5% NaCl; 0–50 ± 1 h Ar-purged; >50 h oxygenated: (a) RE and TU specimens;(b) RE and MA specimens.


of Rt with time was much greater for the TU0.9 specimen than that for the TU2.0 speci-men. This may be attributed to the fact that the well defined particle boundaries present onthe TU0.9 specimen (Fig. 2(C)) are prone to the intergranular attack. Also, the decrease inRt for TU2.0 on Mg in the oxygenated solution was much less than that prepared on thecopper substrate (Fig. 9(a)). This difference might be attributed to the high roughness ofthe Mg alloy surface and the electrochemically produced Cu intermediate layer, leading tothe production of a deposit with less well defined particle boundaries. By comparison,noticeable differences were not observed for RE and MA specimens (Fig. 9(b)).

The EIS spectra for the TU specimens evolve considerably as Ecorr relaxes during theO2-purging period. This is best appreciated from the phase angle plots in Fig. 6(B1)–(D1). A second time constant evolves with time at low frequencies and for the TU0.9and TU2.0 specimens, a Warburg impedance is required to accurately fit the spectra. Thissuggests partial control of the corrosion by O2 transport within the deepening intergran-ular sites (Fig. 6(C1) and (D1)). By contrast, the RE, MA1k, and MA3k specimens do notobviously develop a second time constant during the O2-purging period, although MA180shows a slight tendency to do so. This is consistent with the SEM micrograms obtained atthe end of the Ar-purged (Fig. 4) and the O2-purged (Fig. 2) periods which both show agenerally uniform distribution of numerous pits. The slight differences might be due to thepresence of the patches of surface film on the surface as indicated at the end of the Ar-purged period.

XPS analyses showed these differences coincided with the formation of two distinctlydifferent corrosion films/layers after the oxygenated period [22]. Higher P/Ni ratios(1–2) were observed on smooth specimens (RE, MA1k, and MA3k) than on the nodularTU specimens (<0.5). The P/Ni ratios after O2-purging were much higher than those mea-sured on sputter-cleaned surfaces for the RE, MA1k, and MA3k specimens. By contrastthe TU0.5 and TU2.0 specimens had similar P/Ni ratio as those of the sputter-cleaned sur-face [22]. In general, corrosion of smooth deposits lead to an increase in P/Ni ratio whilethose specimens exhibiting intergranular corrosion experienced little change in this ratio[22].

The composition of corroded smooth surfaces was mainly P0 along with PO3�4 , H2PO�2

and only a minor amount of nickel oxide or hydroxide. By comparison, the intergranularlyattacked TU specimens had a high O content mostly in the form of nickel hydroxide. Thenature of these surface layers influenced the corrosion rates of the specimens. The Penriched surface layers of the RE and MA specimens lead to the higher Rt values observed,relative to those for the intergranularly attacked TU2.0 specimens (Fig. 9). However, athick generally uniform hydroxide layer also provides some corrosion resistance as inthe case of the TU0.5 specimen and the TU2.0 specimen prepared on the Mg AZ91D sub-strate when the particle boundaries are not as well defined or subsequently corroded(Fig. 2(B1) and Fig. 9(a)).

The use of the two time constant circuit shown in Fig. 7(b) allows the determination ofthe resistance associated with pits and intergranularly corroded sites. The evenly distrib-uted preformed defects resulted in pit formation on the surface of the RE deposit,Fig. 4(A), and the high pore resistance (Rpo > 10,000 X cm2) (Fig. 10(a)) suggests a largeaspect ratio. The pore resistance increased slightly during the Ar-purged period as theP-rich surface layer started to form on the surface. However, for more aggressive O2-purg-ing conditions, Rpo initially decreased by an order of a magnitude, suggesting opening upof the pits. As the pits propagated, the aspect ratio became larger resulting in the observed

0 20 40 60 80 100 120 140 160 180 200

10

100

1000

10000

Rpo

,Ωcm

2 Rpo - RE Rpo - TU0.5 Rpo - TU0.9 Rpo - TU2.0 Rpo - RE on Mg Rpo - TU2.0 on Mg

10

100

1000

10000

Rpo

,Ωcm

2

Rpo - RE Rpo - MA180 Rpo - MA1k Rpo - MA3k Rpo - RE on Mg Rpo - MA3k on Mg

Time, h

0 20 40 60 80 100 120 140 160 180 200

Time, h

Fig. 10. Pore resistances of the deposits produced from differently stabilized baths after various periods ofcorrosion in 5% NaCl; 0–50 ± 1 h Ar-purged; >50 h oxygenated: (a) RE and TU specimens; (b) RE and MAspecimens.


recovery in Rpo (Fig. 10(a)). The Rpo value for the RE on Mg specimen did not show thedecrease and subsequent increase, but the Rpo was similar to that of the RE specimen afterlong immersion in the oxygenated solution.

For the TU specimens (including the TU2.0 on Mg specimen), a similar drop in Rpo

occurred on switching from Ar to O2, except for the TU0.5 specimen. This can be attrib-uted to the enlarged corrosion surface area with the introduction of O2. For the TU0.5specimen, the particle boundaries are not clearly defined (Fig. 2(B2)). Also, the high sur-face oxide content prevented the corrosion from penetrating to the particle boundaries.


Thus, the Rpo was much smaller compared to the others (Fig. 10(a)). After the introduc-tion of oxygen, the corrosion reaction started to occur along the particle boundaries as inthe case of the other TU specimens, and the Rpo values comparable to those of the otherTU specimens were observed through the long oxygen purged period (Fig. 10(a)).

For the MA specimens, two different behaviours were observed. A drop in Rpo wasobserved for the MA180 and MA3k specimens but not for the MA1k and the MA3kon Mg specimens on switching from Ar to O2 (Fig. 10(b)). This drop can be attributedto the increase in the number densities and/or opening up of pits, in the more aggressivecorrosion environments, Fig. 2(E) and (G). For the MA1k specimen, the high Rpo value(�20,000 X cm2) under Ar-purging did not change on switching to O2. This is consistentwith the presence of very few pits on this specimen, Fig. 2(F2). XPS results indicated thisspecimen is protected by a thick oxide layer under Ar-purged conditions and a strong sur-face enrichment of P after O2-purging.

Steady and high film resistances (Rfilm � 10,000 X cm2) were observed on all MAspecimens (Fig. 11(b)) while significantly lower film resistances were observed on mostTU specimens (Fig. 11(a)). The high values of Rfilm can be attributed to passivation byP-enriched surface layers as indicated by XPS [22]. The film resistances for the TU samplesdecreased as intergranular corrosion became severe with oxygen purging. The noticeablylower Rfilmobserved for the TU0.9 specimen (compared to the TU2.0 specimen) can beattributed to the much larger particle boundaries for TU0.9 compared to TU2.0. Thisresults in the accumulation of a thinner layer of corrosion products on the particles.

3.5. Significant factors governing the corrosion characteristics of the protective Ni–P coating

3.5.1. Microstructure vs. microscopic defects

The microstructure of the deposit as characterized previously [8] indicated that the ENdeposits are polycrystalline with mixed amorphous and nanocrystalline phases. The depos-its with a higher degree of amorphicity or smaller grain sizes (i.e. the MA3k specimen)exhibited slightly poorer corrosion characteristics than those with ordered nanocrystals(i.e. the RE specimen). This result is not consistent with the literature; amorphous Ni–Pcoatings passivate in acidic and neutral solution, but the presence of crystalline structuresincreased corrosion rates dramatically [12,25]. This conclusion was drawn knowing thatthe deposit only differs in microstructure [12] by applying heat treatment, or microstruc-ture and P content [25] by adjusting electroplating bath composition. In the present cor-rosion study, for a deposit with smooth morphology, the dominant factor for thecorrosion performance appeared to be the pre-formed microscopic defects/preferentiallyattacked sites since the morphologies observed after etching and corrosion were similar;i.e. both exhibited evenly distributed small pit formation (Fig. 2(A2) and Fig. 12).

3.5.2. Morphology

Specimens with a smooth morphology (i.e. the RE, MA1k and MA3k specimens) weremuch more corrosion resistant than those with a nodular morphology, except for theMA180 specimen. The variation in P content between specimens was insignificant and,therefore, not a factor in determining the differences in corrosion behaviour. However,the corrosion resistance decreased at the high TU concentration, consistent with publishedobservations [18]. Our study of deposit microstructure [8] showed TU0.9 to be more amor-phous than the RE specimen. Clearly amorphicity is not critically important in determin-

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Rfil

m,Ω

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Time, h

Rfilm - RE Rfilm - TU0.5 Rfilm - TU0.9 Rfilm - TU2.0 Rfilm - RE on Mg Rfilm - TU2.0 on Mg

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100000

Rfil

m,Ω

cm2

Time, h

Rfilm - RE Rfilm - MA180 Rfilm - MA1k Rfilm - MA3k Rfilm - RE on Mg Rfilm - MA3k on Mg

Fig. 11. Film resistances of the deposits produced from differently stabilized baths after various periods ofcorrosion in 5% NaCl; 0–50 ± 1 h Ar-purged; >50 h oxygenated: (a) RE and TU specimens; (b) RE and MAspecimens.


ing corrosion behaviour since the corrosion rate for TU0.9 is greater than that of the REspecimen. It appears that the presence and density of particle boundaries are more impor-tant. As the particle boundaries become more clearly defined (Fig. 2(A2), (B2), and (C2)),the corrosion resistance decreases, due to intergranular corrosion along these boundaries.However, the MA180 specimen dose not fit this pattern, since despite the presence of fineparticle boundaries, the total resistance value was high and compatible with that of theMA3k specimen.

Fig. 12. SEM microgram of a polished Ni–P deposit surface after etching with dilute HNO3 (5%) for 5 s.


3.5.3. Surface composition and impurity

The higher corrosion resistance observed for the RE specimen and the MA specimencan be attributed to the formation of a P-enriched layer shown to be present by XPSand mostly in the P0 valence state. Such a layer has been claimed to be a chemical passiv-ation layer against corrosion [12,17,26], the accumulation of elemental P on the alloy sur-face acting as a barrier against dissolution. On the contrary, the more rapidly corrodingTU specimens possessed a low P/Ni ratio but a high O/Ni ratio (>2.0).

In addition, trace amounts of the sulfur were detected by the XPS on TU specimenbefore and after corrosion [9–11,22] and its presence was confirmed by high resolutionTOF-SIMS [22]. The S peaks were detected at a binding energy of �167–168 eV corre-sponding to an oxidized form of sulfur, most likely SO2�

4 [27]. This form of sulfur was pro-posed by Eklund as the dominant dissolved sulfur species originating from the corrosionof sulfide inclusions in stainless steel (SS) [28]. The acidity formed via this corrosionprocess

MnSþ 4H2O!Mn2þ þ SO2�4 þ 8Hþ þ 8e� ð3Þ

was considered to activate the surface.The stability domain of adsorbed sulfur extends beyond the usually predicted range of

stability of metal sulfides; i.e., adsorbed sulfur can exist under conditions in which no bulksulfide is stable [29]. The presence of the adsorbed sulfur has been shown to increase thedissolution rate of nickel based alloys, a weakening the metal–metal bond induced byadsorbed sulfur. This lowers the activation energy barrier for the passage of metal atomsfrom the surface to the solution [29]. Furthermore, this influence can be localized if thesulfur is adsorbed on specific sites, where it can be more tightly bonded [29]. For TU spec-imens, sulfur adsorption may be concentrated on the particle boundaries (Fig. 2(A2)–(D2)), where the higher adsorption density of TU during the EN deposition is expected


due to the higher diffusive flux. Adsorbed sulfur is also thought to retard the growth ofpassive films by blocking the sites utilized in oxide growth by OH� adsorption. Anodicsegregation of sulfur caused by the selective dissolution of nickel could lead to large sur-face concentrations of adsorbed sulphur on the surface [29]. These sites could be contin-uously exposed to further corrosion. Such processes could explain the large differencesin total resistance between surfaces with particle boundaries (TU0.9, TU2.0 and MA180).The significantly lower resistance values for TU0.9 and TU2.0 compared to MA180 couldbe attributed to the presence of S on particle boundaries.

In addition to the influence of S impurities, local acidity could be enhanced at particleboundaries by Ni2+ hydrolysis to NiOH+, Ni(OH)2, and Ni2OH3+ [30]

Ni-P

Ni dissolution is blocked

PNi

Ni2+

Ni-P

P-rich

5 % OH-

O2

Ni dissolution enhanced by SadsSads

S2-+ 4H2O

SO42-

Ni2+

8H+

Ni

O2OH-

Ni2++ H2O

H+

Ni(OH)+

O2OH

Ni(OH)2

Sads

O2

P + 4H2OPO4

3-

Ni2+

8H+

Ni-P5 % NaCl Solution

P oxidation

Sulfide oxidation

Ni2+ hydrolysis

Ni

Ni

Fig. 13. Schematic models for (a) intergranular corrosion of Ni–P deposits prepared with TU and (b) chemicalpassivation of Ni–P deposits prepared with no stabilizer or with MA during corrosion in oxygenated neutral 5%NaCl.


Ni2þ þH2O ¼ NiOHþ þHþ ð4Þ2Ni2þ þH2O ¼ Ni2OH3þ þHþ ð5Þ

The schematic in Fig. 13(a) attempts to summarize the reactions described as occurringon an EN surface prepared with TU present. Ni2+ hydrolysis, combined with S and P oxi-dation, lowers the pH at the particle boundaries and increases the solubility of variousNi2+ hydrolysis products such as NiOH+ and Ni2OH3+ [30]. Thus, these Ni2+ hydrolysisproducts readily diffuse out of particle boundaries and precipitate to form a thick oxidelayer on the outer surfaces as shown by XPS analyses. Due to the acidity within the cor-roding particle boundaries, oxide/hydroxide formation is prevented and corrosion conti-nues. On the contrary, in the absence of the particle boundaries and sulfur impurities,the general preferential dissolution of nickel produces a P-rich chemically passive layerwhich hinders further dissolution of nickel as illustrated schematically in Fig. 13(b).

4. Conclusions

The effects of the EN bath stabilizers, TU and MA, on the corrosion of the EN coatingson copper and magnesium substrates (with an intermediate copper layer produced by anovel Cu immersion process) were investigated in an aggressive oxygenated chlorideenvironment.

The dissolution or the corrosion attack of the AZ91D Mg alloy and Cu substrates werenot observed throughout the corrosion experiments, with corrosion limited to the externalEN layer.

Two different types of corrosion morphology were observed. Specimens prepared withTU, and the MA180 specimen, showed preferential attack along particle boundaries whilethose prepared with no stabilizer and MA showed a pitting-type corrosion on the smoothsurfaces. Higher corrosion resistances were observed for smooth deposits, which is attrib-uted to the formation of a P-enriched layer on the surface. The poorer corrosion perfor-mance of EN specimens prepared with TU is attributed to the presence of the particleboundaries and the trace impurities adsorbed on, or included in, them.

It was clearly shown that the corrosion behaviour of the EN specimens studied was notgoverned by the microstructural aspects of the EN deposit but was predominantly influ-enced by morphological and surface compositional aspects (i.e. surface impurities).

Acknowledgements

This research is funded by the Natural Science and Engineering Research Council ofCanada (NSERC). The authors gratefully acknowledge the general technical help of JohnNagata (IMTI-NRC) and help of Mike Meinert (IMTI-NRC) and Brad Kobe (SSW) inthe acquisition of SEM micrographs. The authors would like to thank Dr. Zack Qin(UWO) for the design of the corrosion cell and Dr. Jamie Noel (UWO) for the help inshowing problems related to electrochemical instrumentation.

References

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