Corrosion Science 83 (2014) 111–123

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Corrosion Science

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In situ EC-AFM study of the effect of nanocrystals on the passivationand pit initiation in an Al-based metallic glass

http://dx.doi.org/10.1016/j.corsci.2014.02.0050010-938X/� 2014 Elsevier Ltd. All rights reserved.

⇑ Corresponding author.E-mail address: jqwang@imr.ac.cn (J.Q. Wang).

S.D. Zhang a, Z.W. Liu b, Z.M. Wang c, J.Q. Wang a,⇑a Shenyang National Laboratory for Materials Science, Institute of Metal Research, CAS, Shenyang 110016, PR Chinab Agilent Technologies, Wang Jing North Road No. 3 Chao Yang District, Beijing 100102, PR Chinac Shengli Engineering & Consulting Company, Shengli Petroleum Administrative Bureau, SINOPEC, Jinan Road 49, Dongying 257026, PR China

a r t i c l e i n f o

Article history:Received 15 October 2013Accepted 2 February 2014Available online 10 February 2014

Keywords:A. AluminiumA. GlassB. AFMB. XPSC. PassivityC. Pitting corrosion

a b s t r a c t

The effect of nanocrystals on pit initiation in metallic glasses is an unresolved issue. The passive film for-mation and pit initiation in the Al–Ni–Ce metallic glass were investigated using in situ electrochemicalatomic force microscope (EC-AFM). The a-Al nanophases were identified from the amorphous matrixbased upon the phase imaging in the tapping mode AFM. In the early stage of the passive film formation,the corrosion products Al(OH)3 formed on the a-Al nanoparticles due to the galvanic coupling. The cor-rosion products incorporated into the passive film changed the local structure and component of the pas-sive film, lowering its stability.

� 2014 Elsevier Ltd. All rights reserved.

1. Introduction

Al–TM–RE (TM: Transition Metal; RE: Rare Earth element)metallic glasses exhibit extremely good corrosion resistance [1,2]and have the potential application as corrosion resistant coatings[3]. One of the main mechanisms responsible for their improvedcorrosion resistance over crystalline counterparts is that they arechemically and structurally homogeneous [4]. Chemical and struc-tural heterogeneities in many crystalline alloys, such as stainlesssteel [5] and Al alloys [6,7], usually exert a crucial influence ontheir localised corrosion behaviour, especially pitting corrosion.Similarly, in Al-based metallic glasses, the precipitated crystallinephases in the amorphous matrix also have important influenceson their corrosion resistance [8–11]. Our previous work [8] sug-gests that the appearance of different types of nanocrystals in vary-ing degrees impact their pitting corrosion resistance. It is generallyrecognised that nanoscale crystals would not significantly changethe stable growth of the pits in metallic glasses [8,9,12]. Theymainly affect the pit initiation and its metastable propagationwhich is essentially related to the structure and stability of the pas-sive film. However, there is lack of direct evidence of how thenanocrystalline phases affect the passive film formation and pitinitiation in Al–TM–RE alloy systems and the nature of pit initia-tion at the nanometre scale is still unclear.

It is thus necessary to elucidate the detailed mechanisms of thepassive film formation for a better understanding of the metallicglass corrosion behaviour. In order to observe the earliest stagesof the effect of nanocrystals on passive film formation, novel sur-face measurement techniques are urgently necessary. In recentyears, the development of local surface measurement techniques,such as atomic force microscope (AFM), has enabled detailedinvestigations of different materials regarding morphology, surfacechemistry and electrochemical response. In situ electrochemicalAFM (EC-AFM) is a local probing technique that can provide directinformation of a sample surface exposed to a solution under ap-plied potentials. This opens up a new scope for understandingthe corrosion mechanism [13]. The ability to probe the local char-acteristics in real-space and real-time at the nanometre scale ofelectrode–electrolyte interfaces has advanced its applications invarious studies, such as the corrosion and inhibition of copper[14], localised corrosion and dissolution of Al alloys [15,16], pittingbehaviour of stainless steel [17–19] and the corrosion of Mg-basedbulk metallic glasses [20,21]. These AFM studies are related to theeffects of the micro- or submicron scale heterogeneities on the cor-rosion of the bulk alloys. But they could hardly identify the initialstages of localised corrosion which needs high-resolution AFMmeasurements. To date, no AFM study exists on the early stagesof the corrosion behaviour of nanocrystals embedded amorphousalloys, and possible reasons may be as follows. First, it is difficultto produce atomically smooth surfaces for these bulk alloys, whichis necessary for high-resolution AFM measurements. The second

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reason is that these nanocrystals usually have a size in the tens ofnanometres and they are hard to be verified by only using a topog-raphy image in contact mode AFM.

In addition to contact mode, tapping mode is another AFMimaging modes. Phase imaging is an extension of the regular tap-ping mode AFM (TM-AFM). Phase images of the local nanoscalevariations in the mechanical properties of samples such as hard-ness, friction and viscoelasticity [22] can be obtained simulta-neously with topography images during the same scan. Phaseimaging is based on detecting the phase lag of the output signalof the cantilever oscillation relative to the signal that drives thecantilever-tip oscillation. A more negative phase shift correspondsto a soft and viscoelastic surface due to greater energy lost tothe sample. Therefore, following the variations of the phase of thevibrating cantilever during the AFM scanning, one can image thedistribution of ‘‘hard’’ and ‘‘soft’’ portions of the material. Phaseimaging in TM-AFM has been widely used to characterise poly-meric materials [23] and biological materials [24]. But so far toour knowledge in situ AFM phase imaging has scarcely been re-ported in literatures to be used for the study of corrosion behaviourof alloys.

In our previous works, the correlation of nanoscale heterogene-ity with corrosion in Al-based metallic glasses has been studied [8],and ex situ AFM has been used to observe the nanoscale break-down phenomena [25,26]. In this paper, we have investigatedthe effect of precipitated a-Al nanocrystals on the formation ofpassive film and the pit initiation in the Al-based metallic glassby TM-AFM, aiming at clarifying the origin of the influence of thenanoscale crystalline phase on the corrosion behaviour of metallicglass. This work has some marked features. First, a good surfaceflatness is the precondition for obtaining a clear phase image inTM-AFM measurements, which is crucial for distinguishing thecrystalline phases with tens of nanometres from the amorphousmatrix. The amorphous ribbon sample is not as flat as the sputter-ing film sample. Special measures were taken for this purpose andfinally an atomically flat surface was achieved. Furthermore, byusing the TM-AFM, we have recognised the nanoscale crystallinephase from amorphous matrix for the first time. The size, shapeand distribution of the nanocrystalline a-Al have been successfullydetermined from the phase images. This made it possible to studythe effect of precipitated a-Al nanocrystals on the formation ofpassive film and the pit initiation in Al-based metallic glass.

2. Experimental

2.1. Material and sample preparation

The materials used in this study were Al88Ni8Ce4 amorphousribbons and partially nanocrystallised samples produced by an iso-thermal heat-treatment in evacuated quartz tubes at 180 �C for60 min. The amorphous ribbons were fabricated by induction melt-ing from the master alloy ingot in a silica crucible under argonatmosphere, and ejecting it onto a copper wheel rotating with aspeed of 20 m/s using a Bühler melt spinner (Hechingen, Ger-many). The master alloy ingot with nominal compositions of Al88-

Ni8Ce4, was prepared by melting appropriate amounts of elementalmetals (Al: 99.999%, Ni: 99.9%, Ce: 99%) under a Ti-gettered argonatmosphere on a water-cooled copper plate. Each ingot was flippedand remelted several times to avoid any compositional nonunifor-mity or segregation inside the ingot alloys. The ribbons, with athickness of about 50 lm and a width of about 3 mm, were cut intocoupons which have a length of 5 mm. The coupons used for elec-trochemical measurements and in situ EC-AFM experiments wereelectrically connected to an isolated wire and embedded in anepoxy resin leaving a cross-section (�3 mm � 50 lm, 0.001–0.002 cm2) exposed surface area. The samples were designed to

be discs of 20 mm in diameter and 5 mm in thickness in orderfor proper EC-AFM measurements.

2.2. Microstructure characterisation

The as-quenched ribbons and annealed samples were examinedby X-ray diffraction (XRD) using a Rigaku D/max 2400 diffractom-eter (Tokyo, Japan) with monochromated Cu Ka radiation(k = 0.1542 nm). The thermal stability of two types of samples wereanalysed in a Neztsch DSC 404C under purified flowing argon at aheating rate of 10 �C/min up to a temperature of 500 �C. An FEImodel Tecnai F30 transmission electron microscope (TEM) wasused to confirm the microstructure of the as-quenched and an-nealed samples. Thin foils for TEM observations were preparedby a standard twin-jet electropolishing method in a mixture of25 vol.% nitric acid and 75 vol.% methanol at –30 �C.

2.3. Electrochemical measurements

Electrochemical measurements of the as-quenched and an-nealed samples were conducted in 0.01 mol/L NaCl aqueous solu-tion using a potentiostat–galvanostat EG&G Princeton AppliedResearch model 2273. Solutions were made from reagent-gradechemicals mixed with deionised water. All electrochemical testswere performed in a conventional three-electrode cell, using alarge platinum plate as the counter electrode and a saturated cal-omel electrode (SCE) as the reference electrode. All the measure-ments were held at 30 ± 1 �C in a water bath and the potentialvalues mentioned in this paper were versus SCE. Prior to all elec-trochemical measurements, the samples were kept in the solutionuntil a stable open circuit potential (OCP) was attained. Themounted samples were mechanically ground with up to 2000 gritSiC paper, then rinsed in deionised water and dried in a streamof air prior to the experiments. Potentiodynamic and potentiostaticpolarisation measurements in the anodic regime were performedin 0.01 mol/L NaCl aqueous solution. The dynamic polarisationwas started from the cathodic regime (about �50 mV vs. OCP)and linearly swept in the anodic direction at a rate of 0.333 mV/still the current reached a limit of 10–5 A. Potentiostatic measure-ment at an anodic potential of �250 mVSCE was conducted in0.01 mol/L NaCl aqueous solution. Changes in the current signalwere recorded and collected at a data-sampling interval of 0.05 s.Electrochemical impedance measurements (EIS) were conductedat the OCP after the samples were kept in the solution at OCP for60 min and a stable state was attained. The EIS experiments usedan AC amplitude of 10 mV (rms) in a frequency range of 10 kHzto 5 mHz with ten points per decade. EIS results were calculatedbased on a suitable fitting of an equivalent circuit using Zview soft-ware. In order to obtain a high reliability of the results, some elec-trochemical measurements were repeated five times and theresults that can represent the mean value were selected.

2.4. In situ EC-AFM measurements

In situ EC-AFM observations were performed using an AgilentAFM series 5500 (Agilent Technologies). Imaging was carried outin the tapping mode (AAC mode) and a silicon cantilever with softtapping mode tip (NanoWorld Co., NCST; Typical size: T = 2.8 lm,W = 27 lm, L = 150 lm, H = 10–15 lm; Typical spring constant:C = 7.4 N/m, Typical resonance frequency: f = 160 kHz, radius ofcurvature of the tip: R < 8 nm) was used. The scanning rate was0.6–2.0 line/s.

The AFM liquid cell was made from chemical-resistant polycar-bonate and was 15 mm in diameter, which was pressed against thedisk-shaped sample surface with two screws and retaining clips.We noted that the sample surface must be flat and larger than

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the diameter of the cell. A viton O-ring gasket was used to avoidleakage. The measurements were carried out with a three-elec-trode setup. The working electrode was the sample on the bottomof the cell and the reference electrode was a silver wire with0.51 mm diameter which was located within the electrolyte butit did not contact the working electrode. The counter electrodewas a coiled Pt wire with 0.25 mm diameter and it encompassedas much of the inner rim of the liquid cell as possible. The cablehead from the AFM sample plate was connected to the Potentio-stat/Galvanostat (PARSTAT 2273). The Ag reference electrode wascalibrated with respect to the SCE in the same solution used forthe polarisation experiments. Typically, it had a potential of about�150 mV relative to the SCE. For consistency all potentials werequoted vs. SCE in this paper.

In this study, the nanocrystals embedded in Al-based metallicglasses were very fine, about 50 nm in diameter, and the differencein viscoelasticity between the two phases was tiny. Therefore, itwas a challenge to get a good phase image from which the nano-crystals could be identified from the amorphous matrix. In orderto detect the slightest surface changes, special care was taken forAFM sample preparation, which was also crucial for obtaining reli-able results. The samples were ground in sequence on 1000, 2000,3000 and 5000 grit SiC paper and then polished with 2.5 lm dia-mond paste, down to 0.5 lm diamond paste using deionised wateras the lubricant. They were subsequently ultrasonically cleanedtwo times in deionised water and air-dried. Finally, the sampleswere further flattened by argon-ion milling with Leica EMRES101 (Wetzlar, Germany) at an acceleration voltage of 3 kV witha beam current of 1 mA. The average corrugation (root-mean-square, RMS) of the surface as prepared was a few angstroms forareas of a few lm2. Any scratches resulting from the formermechanical polishing was removed.

The in situ AFM scanning was performed in an OCP condition, aswell as stepping the potential from OCP to pitting potential (Epit).All the AFM observations were carried out in a vibration isolationchamber at room temperature. The AFM images were processedby flattening in order to remove the background slope, and thebrightness and contrast were adjusted. The root-mean-square(RMS) roughness was evaluated by analysing the 2D height imageswith Picoimage software.

2.5. XPS characterisation

The surface films were analysed by X-ray photoelectron spec-troscopy (XPS) using an ESCALAB250 photoelectron spectrometerand a monochromatic Al Ka radiation (hm = 1486.6 eV). Quantita-tive information was calculated from the survey spectra usingstandard Scofield sensitivity factors [27]. The spot area for testingwas randomly selected at the flat surface with a diameter of0.5 mm. Binding energies were calibrated using the carbon con-tamination with C 1s peak value of 284.6 eV. The depth profilesof the alloy surface were characterised by in situ XPS ion beamsputtering with 2 kV, 2 lA argon ions at a sputtering rate of about0.2 nm/s. The samples for XPS analysis were flattened by argon-ionmilling with Leica EM RES101 (Wetzlar, Germany) after beingground with 2000-grit SiC paper. Then they were cleaned withdeionised water and immersed in 0.01 mol/L NaCl solution for1 h. Spectra were curve-fitted using commercial software.

Fig. 1. X-ray diffraction patterns of the as-quenched and annealed Al88Ni8Ce4

amorphous ribbon.

3. Results

3.1. Microstructure characterisation

The XRD patterns of the as-quenched and annealed Al88Ni8Ce4

samples are shown in Fig. 1. For the quenched sample, only a broad

hump at the angle of 2h = 38� without any diffraction peaks associ-ated with any crystalline phases can be seen. This indicates that itis a fully amorphous structure. For the annealed sample, five dif-fraction peaks are detected besides a weakened broad halo in thescanned range. The five diffraction peaks correspond to planes(111), (200), (220), (311) and (222) respectively and the crystal-line phase is identified as fcc-Al. This means that an amorphous-nanocrystalline composite with a-Al nanocrystalline phaseprecipitated in the amorphous matrix is formed. Fig. 2 presentsthe DSC curves of the as-quenched and annealed Al88Ni8Ce4 sam-ples at a heating rate of 10 �C/min. As shown in Fig. 2, for both sam-ples, three exothermic transformations occur on the DSC curves.The first wide exothermal reaction from 150 to 280 �C correspondsto a-Al nanocrystal formation, while the second and the thirdpeaks are ascribed to the formation of some intermetallic com-pounds [28]. However, from the enlargement of the selected DSCcurves, we can see that the onset transformation temperature ofthe first reaction for the annealed sample is about 50 �C higherthan that of the as-quenched sample and the peak position changeswith it. Two other peaks remain unchanged compared to the as-quenched sample. DSC methods can be used to evaluate the crys-tallinity of nanocomposites composed of the amorphous matrix.The volume fraction of Al crystals, Vf, is usually evaluated by theequation [29]: Vf = 1 � DHcryst/DHamor, DHcryst and DHamor are thetransformation enthalpy of a partially crystallised sample and afully amorphous sample, respectively. The volume fraction of thea-Al nanocrystal is about 17.5%. In order to determine the struc-ture of the samples further, both were observed and analysed byusing high-resolution transmission electron microscopy (HRTEM)and selected area electron diffraction (SAED). The results areshown in Fig. 3. For the as-quenched sample, as shown in Fig. 3a,the speckled appearance of the HRTEM image corresponds to theshort-range order structure at the atomic scale and the corre-sponding SAED pattern shows diffuse halo rings, further confirm-ing its fully amorphous structure. Dark-field TEM observation ofthe annealed sample is shown in Fig. 3b. The nanoparticles, witha diameter of about 20–80 nm, diffusely distributed in the amor-phous matrix are observed in this sample. Both the dark and brightspots (different crystal orientations relative to the electron beam)represent nanocrystals and the gray matrix between them is theamorphous structure. Fig. 3c presents the HRTEM micrograph ofthe nanocrystals embedded in the amorphous matrix. The HRTEMimage is taken along the h1 1 0 i direction of the nanocrystals. Theclosely packed (111) and ð1 1 1Þ planes are marked with twocrossing lines and the angle of the two planes is 109.5�. The(111) lattice interplanar spacing is determined to be 0.234 nm.

Fig. 2. Differential scanning calorimetry (DSC) continuous heating curves of the as-quenched and annealed Al88Ni8Ce4 amorphous ribbon, heating rate is 10 �C/min.Inset: an enlargement of the selected part of this figure.

Fig. 3. HRTEM and Dark-field TEM micrographs of as-quenched and annealedAl88Ni8Ce4 samples. (a) HRTEM image and selected area electronic diffraction(SAED, inset) of the as-quenched amorphous ribbon. (b) Dark-field TEM image andSAED for the annealed sample. (c) HRTEM image of the nanocrystals precipitated inthe amorphous matrix in Fig. 3b, the image is taken along the h1 1 0i direction of thenanocrystals.

Fig. 4. Potentiodynamic scans conducted on the as-quenched amorphous andannealed amorphous-nanocrystalline Al88Ni8Ce4 alloys in 0.01 M NaCl solution.

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This value matches exactly with the (111) plane of metallic Al,further confirming that the nanocrystal precipitated in theamorphous matrix is fcc-Al.

3.2. Potentiodynamic polarisation

Fig. 4 displays the potentiodynamic polarisation curves forthe as-quenched and the annealed samples in the test solution.The three dots along the polarisation curve in Fig. 4 represent thepotential values that have been applied to the sample during thein situ EC-AFM measurements. Under the applied anodic potentials,the two materials spontaneously enter the passive region and exhi-bit excellent passivation ability. The passive current density (ipass)for the as-quenched sample is basically constant at about 3.3 lA/

cm2, while the ipass for the annealed sample increases slightly withpotential in the passive region and the average value is about5.2 lA/cm2. The corrosion potentials (Ecorr) of both samples are verysimilar (about�0.50 VSCE), but the pitting potentials (Epit) are differ-ent. The Epit of the as-quenched amorphous sample is more positive(about 50 mV) than that of the annealed partially nanocrystallisedsample and the passive region is thus wider. The pitting resistanceof the partially devitrified alloy is inferior to that of the fully amor-phous precursor. The statistical analyses of the Epit of the as-quenched and the annealed samples are �0.110 ± 0.026 VSCE and�0.160 ± 0.039 VSCE, respectively. Pitting corrosion occurs in a seriesof consecutive steps [30,31], such as nucleation of embryonic pit,metastable pit state and stable growth of the pit. Stable growth ofthe pits on amorphous alloys mainly depends upon the matrix phaseand it is not significantly affected by the presence of nanocrystals[9,12]. The amorphous matrix of the annealed sample is composi-tionally and structurally the same as the as-quenched one exceptthe solute buildup at the nanocrystal/amorphous matrix interface[32]. Thus there would be no obvious differences for the growthkinetics of the pits between as-quenched and annealed samples,and the difference of pitting corrosion resistance should be relatedto pit initiation and its early growth. The reason is probably thatthe appearance of the nanocrystals destroys the uniform and stablepassive film formed on the chemically and structurally homoge-neous amorphous alloy. The stability of the passive film and thepit initiation in the nanocrystalline a-Al precipitated amorphousalloy will be discussed in detail below.

3.3. AFM characterisation of the nanocrystals in air

The TM-AFM measurements are performed on well-preparedannealed amorphous-nanocrystalline sample surfaces to map thedistribution of nanocrystals. This technique can provide valuableinformation concerning the sample difference of mechanical prop-erties, such as hardness, friction and viscoelasticity. Fig. 5 presentsthe simultaneous 1 lm by 1 lm TM-AFM images of the annealedAl88Ni8Ce4 amorphous-nanocrystalline sample in air. The topogra-phy and phase images were acquired simultaneously in eachexperiment, as described above. The AFM topography image(Fig. 5a) reveals that the surface of the sample is quite flat andno distinct height difference can be seen. Indeed, as can be noticedfrom the line profile in Fig. 5(b), the height difference between thepeak and the bottom is about 2.5 nm within scanning range. TheRMS roughness is 0.544 nm. However, much information can be

Fig. 5. Simultaneous 1 lm by 1 lm TM-AFM images of topography (a) and phase (b) for the annealed Al88Ni8Ce4 amorphous-nanocrystalline sample in air. The topographyline profile (c) and phase line profile (d) of the line as indicated in (a) and (b), respectively. 3D topography image (e) and phase image (f).

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acquired from the phase image (Fig. 5c). Some dark spots in differ-ent sizes are diffusely distributed in the phase image. The size ofthe spots are around 20–80 nm, which is generally consistent withthe TEM result in Fig. 3b. But the shapes of nanoparticles seem verydifferent. These differences in shape come from the methods used.TEM micrograph, a two-dimensional projection of the nanoparti-cles shows a ‘‘visual’’ shape whereas AFM shows a three-dimen-sional image in real space. Furthermore, the details of the darkspots can be seen from the phase line profile in Fig. 5d in whichthe phase shifts and size differences of the marked points are rep-resented. The marked particle in Fig. 5d has a size of 72 nm andexhibits a negative phase shift of about 30� compared to the sur-rounding matrix. The height difference between the same markedpoint is within 0.5 nm (Fig. 5b). The whole surface is very smooth,without marked lumps or indentations, which can be seen clearlyfrom the 3D topography image (Fig. 5e). It is evident from the 3Dphase image (Fig. 5f) that some round patches are distributed ran-domly in the sample.

3.4. In situ EC-AFM observation

After being identified by the TM-AFM in air, the annealed amor-phous-nanocrystalline sample was observed by using the TM-AFMin 0.01 mol/L NaCl solution. In situ EC-AFM measurements were

first taken at the open circuit potential in order to observe the ear-liest stages of the effect of nanocrystals on passive film formation.Fig. 6 shows AFM images with a 1 � 1 lm2 in-plane scan range ofthe annealed Al88Ni8Ce4 sample at OCP held for 3 min. As can beseen from the topography image (Fig. 6a and e), the sample surfaceis very flat with a RMS height variation of about 0.443 nm. The ex-tremely smooth surface makes it possible to investigate the nano-scale change in surface morphology during polarisation. We cansee clearly from the phase image (Fig. 6b and f) that the surfacepresents many small randomly distributed a-Al islands with anaverage diameter of about 20–60 nm, and an average peak to val-ley height of about 0.5 nm (shown in Fig. 6c). Some a-Al islandssurrounded by bright annular patches look like craters, whichcan be seen clearly in the local enlarged phase image in Fig. 6g.The nucleation of the passive film can be seen clearly from thephase image (Fig. 6b, f and g). More details about the a-Al islandmarked with green circle in Fig. 6a and b are exhibited in thetopography and phase line profiles represented by Fig. 6c and d.The a-Al island, about 55 nm in diameter, has a negative phaseshift with respect to the surrounding matrix. It is about 0.5 nmlower than the surrounding matrix, as shown in Fig. 6c. Fig. 6hshows the potential–time curve recorded during the EC-AFM mea-surements. We can see that the potential drops and then increasesas time goes on.

Fig. 6. In situ AFM images of the annealed Al88Ni8Ce4 amorphous-nanocrystalline sample at OCP hold for 3 min in 0.01 mol/L NaCl solution. Topography image (a) and phaseimage (b) and corresponding line profiles (c) and (d). 3D images of topography (e) and phase (f) and the local enlarged phase image of the marked circle in Fig. 6f (g).Potential–time curve of the sample during the EC-AFM measurement (h).

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The passive film are growing by stepping the potential to morepositive values up to E = �0.4 VSCE. This can also be regarded as theformation of anodic passive film. Fig. 7 depicts the in situ AFMimages of the evolution of a surface morphology and their corre-sponding line profiles after 6 min and 29 min, respectively. Tworepresentative oxide particles (A and B) are marked in the AFMimages in order to analyse the passive film formation process. Asthe upper and lower surfaces of the AFM sample are not absolutelyparallel, each scanning cause a small drift in position. The intervals

between Fig. 7a and b is about 23 min and several scans have beendone during this period of time, thus producing a relatively largedrift in position between Fig. 7a and b. It is clear that the oxide par-ticles are growing as the polarisation time increases from 6 min to29 min. The AFM amplitude and phase images (Fig. 7e�h) demon-strate clearly the changes of the passive film though the topogra-phy images look alike (Fig. 7a and b). From the amplitude andphase images (Fig. 7e–h), we can see clearly that the oxide particlesgrow large and round, and their surfaces become smooth as the

Fig. 7. In situ AFM images of the annealed Al88Ni8Ce4 amorphous-nanocrystalline sample in the growth stage of the passive film under anodic polarisation at �0.4 VSCE afterpassivation 6 min and 29 min in 0.01 mol/L NaCl solution. Topography images (a and b) and corresponding line profiles (c and d), amplitude images (e and f) and phaseimages (g and h).

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polarisation time increases. We can also see some small particlesmerging into larger particles. Detailed information of the individ-ual oxide particle can be expressed by the corresponding lineprofile. Fig. 7c and d show the corresponding line profiles of the

lines as indicated in Fig. 7a and b, respectively. For particle A, itis clear that the hollow of passive film grow from �14 nm to�12 nm while its bulges dissolve from 28 nm to 21 nm. Thecombined action of dissolution and growth gives rise to the

Fig. 8. In situ AFM images of the annealed Al88Ni8Ce4 amorphous-nanocrystalline sample in pit initiation stage under anodic polarisation at �0.25 VSCE after 7 min, 16 minand 25 min in 0.01 mol/L NaCl solution. Topography images (a, c and e) and phase images (b, d and f).

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smoothening of the surface, and consequently decreases the RMSroughness from 11.7 nm to 10.3 nm. Changes in the dimensionsof particle A with an increase in polarisation time are shown inFig. 7c and d. Such changes are clearly visualised from the corre-sponding line profiles and could represent the passive film growthrate. The increase in diameter of particle A is 120 nm (from 440 nmto 560 nm), and the decrease in height is 7 nm (from 28 nm to21 nm). The size of this particle diameter is about one order ofmagnitude higher than the size of the a-Al particles. It is probablythat the a-Al nanoparticles are oxidised very quickly at the begin-ning of the experiment and then they serve as the nucleation sitesfor the growth of the passive film. This scheme feels well-consis-tent with the smoothing of surface. The changes of particle B area 220 nm increase in diameter and a 16 nm decrease in height(The line profile is not shown in Fig. 7). Based on the area changerates of oxide islands A and B, the average lateral growth rate ofthe individual island of the passive film is estimated to be about100 nm2/s.

As the potential anodically shifts, the passive film becomes lo-cally instable and metastable pit events occur. Fig. 8 exhibits therepresentative AFM images of the evolution of a surface morphol-ogy under anodic polarisation at �0.25 VSCE after 7 min, 16 minand 25 min in 0.01 mol/L NaCl solution. This figure clearly showsthe dynamic process of the metastable pit initiating, passivatingand propagating. The metastable pit is marked with a green ellipsein each image. Around the pit, six islands of passive film (num-bered 1–6) are marked. The phase images can present more clearchanges of morphology than the topography images. As shown inFig. 8a and b, after the sample is polarised for 7 min, most probablya metastable pit forms at the weak point in the passive film. Theseweak points include the heterogeneities in oxide films (the a-Alnanocrystals in Fig. 6) and the boundary regions of the oxide filmislands (Fig. 7). The smooth bottom of the pit can be seen clearlyin the phase image. Fig. 8d reveals that islands 2 and 3 grow largerand that islands 1, 4, 5 and 6 rise around the pit after the sample ispolarised for 16 min. This causes the metastable pit to shrink and it

Fig. 9. The XPS spectrum of Al 2p for (a) the quenched and (b) the annealed Al88Ni8Ce4 alloys and (c) the schema (not to scale) of the relative components of photoelectronline Al 2p for both samples, the red rectangle represents the XPS detecting range. (For interpretation of the references to color in this figure legend, the reader is referred to theweb version of this article.)

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becomes shallow and thus repassivation occurs. However, island 2completely dissolves after the sample is polarised for a longer per-iod of time (Fig. 8f), which results in the growth of the metastablepit. Generally the anodic current peaks in potentiostatic or poten-tiodynamic polarisation curves are related to initiation of metasta-ble pits. But there is no anodic peak shown in the in Fig. 4 at�0.25 VSCE. The probably reason is that the transient current aretoo small to be detected by the electrochemical instruments. Infact, metastable pits can nucleate and even grow at potentials be-low the pitting potential [30] and sometimes the current is so tiny(nanoamperes even picoamperes) that it cannot be detected if notby advanced measurement instruments. This dynamic process ofmetastable pit can be observed by in situ AFM, demonstrating thatAFM is an effective tool for in situ corrosion study.

3.5. Chemical analysis of surface film

In order to provide more information about the components ofthe passive films, an XPS technique was exploited because of itscapability in giving the information on the chemical environmentof cations, namely aluminium in our case. The XPS results showthat the main component Al cations account for more than 90%of the chemical compositions of the passive film and the contentsof Ce cations and Ni cations are about 7% and 2%, respectively. Thecontents of Ce cations and Ni cations in the passive film show noobvious difference, thus the XPS spectrum of main component Alare analysed below. The Al 2p XPS spectrum of the as-quenchedand annealed Al88Ni8Ce4 samples immersed in 0.01 mol/L NaClsolution for 1 h are shown in Fig. 9a and b. An Ar ion sputteringwas performed before spectra acquisition. The Al 2p spectrum forboth samples contained three components, metallic state alumi-num (Al(0)), aluminium oxide (Al2O3), and aluminium hydroxide(Al(OH)3). The peaks of three components for both samples are lo-cated at 72.54 eV, 74.39 eV and 75.52 eV, respectively. The Al 2pphotoelectron peak positions, peak assignments, peak areas and

relative contents of three components are summarised in Table 1.In XPS experiment, the detecting area (X-ray beam cross section)and detecting depth are the same for both samples thus the rela-tive content of three components can be denoted by the area inXPS detecting longitudinal section. Fig. 9c exhibits the schema(not to scale) of the relative components of photoelectron line Al2p for both samples, the red rectangle represents the XPS detectingrange. The spectrum peaks of the oxide states correspond to signalsfrom the surface film while the peaks of metallic states are as-signed to the bulk alloy [33,34]. The peak position of the metallicstate aluminium (72.54 eV) for the annealed alloy is the same asthat of the quenched alloy. However, the overall peak position ofthe oxide state components has shifted to the high binding energydirection because the relative proportion of the hydroxide andoxide increases. Assuming that the oxide state components are inthe passive film, the relative components of the films can be deter-mined based on the XPS spectra. The percent of Al(OH)3 in the pas-sive film (Al2O3 and Al(OH)3) for the annealed alloy and thequenched alloy is 76.32% and 69.47%, respectively. This is to saythat the relative concentration of Al2O3 in the passive film is lowfor annealed alloy and the relative concentration of Al2O3 is18.2% and 20.0% (Table 1) for annealed and quenched alloy, respec-tively. There are two factors accounting for this phenomenon. Oneis the oxide film (Al2O3 film not including Al(OH)3 film) formed onthe amorphous-nanocrystalline alloy is thin (Fig. 9c II), and anotheris that the oxide film (Al2O3 film not including Al(OH)3 film) on theamorphous-nanocrystalline alloy is inhomogenous and thehydroxide is locally formed on the surface of Al particles (Fig. 9cIII). For the quenched and annealed alloys, they are composition-ally the same and the amorphous matrix of the annealed alloy isthus structurally and compositionally nearly identical to thequenched alloys except for local slight changes in composition pro-files around the nanocrystals [32]. Therefore, the oxide film formedon the amorphous matrix of the annealed alloy is most likely thesame as that of the quenched one. The first case (Fig. 9c II) that

Table 1Al 2p photoelectron peak positions, peak assignments, peak areas and relative content of chemical forms found in passive film of the quenched and annealed Al88Ni8Ce4 alloys.

Samples Peak position (eV) Peak assignment Peak areas Relative content (%)

Quenched 72.54 Al(0) 27682.29 34.474.39 Al2O3 16088.97 20.075.52 Al(OH)3 36606.14 45.6

Annealed 72.54 Al(0) 21228.88 23.274.39 Al2O3 16636.88 18.275.52 Al(OH)3 53628.56 58.6

Fig. 10. Electrochemical impedance behaviour of the as-quenched and annealedsamples at the OCP after 1 h exposure in 0.01 mol/L NaCl solution and correspond-ing fitting results of EIS data. (a) Bode impedance magnitude and phase angle plots,(b) Nyquist plots. (c) An equivalent circuit model used to fit the EIS data.

120 S.D. Zhang et al. / Corrosion Science 83 (2014) 111–123

the thin oxide film formed on the annealed alloy is unreasonable.Thus, the second case (Fig. 9c III) that the oxide film is inhomoge-neous is probably correct. This inhomogeneity is also observed inthe early stage of the passive film formation (Fig. 6).

3.6. Electrochemical impedance spectroscopy measurements

In order to evaluate the stability of the passive film for both theas-quenched and annealed samples, electrochemical impedancespectroscopy (EIS) was conducted after 1 h exposure at OCP andthe electrode achieving a steady state. The EIS complex planeimpedance plots are shown in Fig. 10. There is no significant differ-ence in electrochemical impedance Bode magnitude and phaseplots for both samples. The response of the Nyquist plane for bothsamples exhibits two time constants in the frequency range 105–5 � 10�3 Hz. Therefore, it can be modeled using an equivalent cir-cuit composed of two hierarchically distributed RC circuits, shownin Fig. 10c, with the two constant phase element (CPE) due to thecompressed nature of the semicircle in the Nyquist plot. The mean-ing of the model parameters is as follows: R0, the high frequencyresistance, corresponds to the solution resistance; CPE1 is a con-stant phase element associated with the double layer capacitanceand R1 with the charge transfer resistance; R2 and CPE2 are associ-ated with the resistance and capacitance of the passive layer. Thecurve fitting of the spectra obtained for both samples is shown inFig. 10. The data are accurately represented by this circuit model.The fitting parameters are summarised in Table 2. As seen fromFig. 10, the as-quenched sample has much higher R1 and R2 valuesthan that of the annealed sample, which indicates that the passivefilm formed on the amorphous alloy is more stable. Thus, for thefull amorphous alloy the migration of the metal ions throughthe passive film towards the electrolyte would be slowed down.The a-Al nanoparticles decreased the oxide film impedanceand the percentage of decrease in the impedance is about 10–12%.

4. Discussion

4.1. TM-AFM for identifying the nanometric a-Al particles

AFM has three differing modes of operation. These are the con-tact mode, non-contact mode and tapping mode. The advantages ofthe tapping mode are the elimination of a large part of lateralforces and it causes less damage to the sample surface. This benefitfacilitates high-resolution imaging. In TM-AFM, the complex tip-sample interactions cause changes in the amplitude and phase ofthe oscillating cantilever. The amplitude and phase images mayhighlight physical properties of the sample that are not readily dis-cernable in the topography image. The material properties varia-tions can be mapped by recording the phase shift between thedriving force and the tip oscillation. The relationship between thephase shift and tip-sample interactions developed by Magonovet al. [23] can be expressed as:

D/0 ¼p2� tan�1 k

Qr

� �� Qr

k¼ s

Qk¼ aE�

Qk

ð1Þ

where D/0 is the phase angle shift between the free and interactingcantilevers, r has been approximated by the surface stiffnesss = �aE�, � is a constant, a is the contact radius and E� is the effectivemodulus. Eq. (1) establishes a proportionality between the phaseshift and the reduced Young’s modulus and shows that a more neg-ative phase shift corresponds to a soft surface and hence appearsdarker in the phase image.

From the above discussion, it is evident that the phase imagecan highlight the local nanoscale variations in the mechanicalproperties of samples that cannot be achieved in the topographyimage. From the topography image (Fig. 5a) we can see that thesurface of the sample is quite flat and the height variations is

Table 2Equivalent electrical circuit fitting parameters of EIS data for Al88Ni8Ce4 alloy from the data set of Fig. 10 after 60 min at OCP in 0.01 M NaCl solution.

Samples CPE1-T (lF/cm2) CPE1-P R1 (X cm2) CPE2-T (lF/cm2) CPE2-P R2 (X cm2)

Quenched 12.128 0.92636 21,024 1222.4 0.9826 17,638Annealed 11.619 0.91744 18,634 1342.2 0.9784 16,011

S.D. Zhang et al. / Corrosion Science 83 (2014) 111–123 121

about 2.5 nm within the scanning range (Fig. 5b). However, asharp contrast picture is shown in the phase image (Fig. 5c),many dark spots are distributed randomly in the matrix. The darkspot has a negative phase shift of about 30� compared to the sur-rounding matrix (Fig. 5d). A negative phase shift means that theAFM tip sticks to the soft surface, confirming that the dark spothas a lower hardness and elastic modulus with respect to the sur-rounding matrix. The XRD analysis of this sample indicates thatthe sole crystalline phase a-Al exists in the amorphous matrix(Fig. 1). This is also verified by TEM observation (Fig. 3). The Vick-ers hardness of pure aluminium is 17 [35], while the hardness va-lue of amorphous Al88Ni8Ce4 alloy is 345 Hv [36], which issignificantly higher than that of pure aluminium. Hence, thesharp contrast in the phase image (Fig. 5c) should result frominherently different mechanical properties of the two phases.Combined with the results of XRD and TEM, it is therefore reason-able to infer that the round patches in the AFM phase image is a-Al nanocrystals.

It is noteworthy that the topography in the height image can re-sult in the contrast change of the phase image [37]. It is necessaryto reduce or eliminate all topography so that the phase contrast isonly contributed by the material properties. In our study, an atom-ically flat surface has been obtained by argon-ion milling and theRMS roughness is 0.544 nm. This value is comparable to the mag-netron sputtering film which has a RMS roughness of 0.3 nm [38].The topography line profile depicted in Fig. 5b shows that theheight variation is less than 2.5 nm. The sharp contrast in the phaseimage (Fig. 5c) demonstrates the nanoscale variation in the phaseshift. The phase line profile in Fig. 5d reveals that the discrepancyin the phase shift is as large as 30�. The comparison between thephase image and the topography image indicates that the markedregion with large phase shift does not correspond to a rough do-main. This indicates that the phase shift is independent of the sur-face topography. In this case, the effect of the topography on thephase shift can be ignored and the sharp contrast in the phase im-age mainly reflects intrinsic material characteristics. Therefore, theatomically flat surface benefits the identification of the nanometrica-Al particles in AFM observations.

4.2. Nucleation and growth of the passive film

Current–time transients have been successfully utilised tostudy the kinetics of passive film formation [39–42] and the earlystages of pit initiation and propagation [30,43]. In order to revealmore details about the formation of passive film, potentiostaticpolarisation measurements were performed. The widely used the-oretical model for electrochemical nucleation developed byScharifker and Hills [44] was used to analyse experimentalcurrent–time transients. This model can classify the experimentaltransients into two limiting nucleation mechanisms: instanta-neous nucleation and progressive nucleation. The equations ofthe theoretical models for instantaneous and progressive nucle-ation can be expressed as follows:

Instantaneous nucleation:

i2

i2m

¼ 1:9542t=tm

1� exp �1:2564t

tm

� �� �� �2

ð2Þ

Progressive nucleation:

i2

i2m

¼ 1:2544t=tm

1� exp �2:3367t

tm

� �2" #( )2

ð3Þ

where im and tm are the current density and the time of peak coor-dinates, respectively. Fig. 11a shows the experimental current tran-sient plots of two types of samples. Initially the current increasessharply as a response to the applied potential and then decreasesdue to the passive film nucleation. The plots of the experimentalcurrent transients in (i/im)2 vs. t/tm coordinates can follow eitherof the two limiting cases: instantaneous and progressive nucleation.Fig. 11b exhibits the experimental current transients plotted in a re-duced current–time coordinate for the as-quenched and annealedsamples. The curves associated with the two nucleation mechanismdescribed by Eqs. (2) and (3) are also included to illustrate theexperimental nucleation mechanism. It can be seen that in the caseof the passive film on the as-quenched sample, the decreasing por-tion of the experimental plot is closer to the instantaneous nucle-ation curve. This indicates that the passivation proceeds by meansof the instantaneous nucleation mechanism. That is to say, for theas-quenched amorphous sample, all the active sites are activatedrapidly and almost simultaneously. However, for the passive filmon the annealed sample, the decreasing portion of the experimentalplot agrees well with the theoretical progressive nucleation curve.This suggests that the nucleation mechanism complies with theprogressive nucleation. It also means that, for the formation processof the annealed amorphous-nanocrystalline sample, active sites areactivated gradually during polarisation. It is thus clear that the pre-cipitation of the nanocrystalline a-Al in the amorphous matrix haschanged its nucleation mechanism of the passive film from instan-taneous to progressive nucleation. In general, the amorphous alloyhas a higher chemical reactivity due to the metastable nature[45,46]. The amorphous alloy has the ability to form supersaturatedsolid solutions in one or more corrosion-resistant elements. Theseelements, such as TM and RE elements in Al-based metallic glasses,possess a high affinity to oxygen and a high chemical stability in theoxidised state. Therefore, the amorphous alloys containing suffi-cient concentrations of corrosion-resistant solute elements arecharacterised by a rapid formation of uniform corrosion-resistantpassive film. In the nanocrystals embedded amorphous alloy, theactive sites are not activated rapidly and simultaneously. Thus,the occurrence of nanocrystals changes the nucleation mechanismof the passive film.

The nucleation and growth of the passive film can be observedby using in situ EC-AFM. The nucleation of the passive film occursafter the sample is immersed in NaCl solution at the OCP (Fig. 6b, fand g). The bright annular lumps form on a-Al islands at the a-Al/amorphous matrix interface, leaving a cavity on each a-Al island(Fig. 6g). The bright annular lump has a different phase shift com-pared to the surrounding matrix, indicating that a new substance isproduced. The potential’s change over time can give interestinginformation about the evolution of corrosion or passivationphenomena. From the potential–time curve recorded during theEC-AFM measurements (Fig. 6h), we can see that the potentialdecreases and then increases. This classic potential–time curvereflects an attack followed by passivation of the alloy [47]. The cor-rosion potential of a-Al is about 300 mV lower than that of the

Fig. 11. Potentiostatic experiment results of the as-quenched and annealedsamples at an anodic potential of �250 mVSCE conducted in 0.01 mol/L NaClaqueous solution. (a) current–time curves. (b) Theoretical and experimentaldimensionless plots (i/im)2 vs. t/tm plots for the passive film on both samples.

122 S.D. Zhang et al. / Corrosion Science 83 (2014) 111–123

amorphous matrix [8], and the less noble a-Al phase acts as an an-ode and is attacked by galvanic corrosion. From the phase images(Fig. 6b and g), we have observed a new substance forming onthe a-Al particles at the a-Al/amorphous matrix interface. TheXPS results demonstrate that the percent of Al(OH)3 in the passivefilm for the annealed sample is higher than that of the as-quenchedsample (Fig. 9c). According to our previous discussion in XPS anal-ysis section, the oxide film on the annealed amorphous-nanocrys-talline alloy is inhomogenous and most probably the hydroxide islocally formed on the surface of Al particles (Fig. 9c III). This inho-mogeneity could be observed in the early stage of the passive filmformation in Fig. 6. The increased hydroxide content in the passivefilm of the annealed alloy can be mainly attributed to the corrosionproducts which are formed at the early stage of the passive filmformation as discussed in EC-AFM observations (the annular lumpsaround the a-Al islands in Fig. 6). Therefore, it can be asserted thatthe annular patches surrounding the a-Al islands are the corrosionproducts Al(OH)3 which have a different characteristic from theoxide film formed on the amorphous matrix. The hydroxides inthe passive film are of a lower density than the oxides [47], whichdecrease the compactness of the passive film. These heterogene-ities weaken the protective properties of the passive film andmay be the weak points preferentially attacked by Cl� during ano-dic polarisation.

5. Conclusions

In situ EC-AFM was successfully performed on the a-Al nano-crystals precipitated Al-based metallic glass to study the passive

film formation and pit initiation during the potential scan in chlo-ride-containing solutions. Tapping mode AFM was exploited andphase imaging allowed us to identify the nanocrystalline a-Al fromthe amorphous matrix. Phase imaging combined with the atomi-cally flat surface made it possible to follow the very early stagesof the passive film formation process. The corrosion productsAl(OH)3 are formed because of the galvanic coupling between thea-Al particle (anodic site) and the surrounding amorphous matrix(cathodic site) during the earliest stages of passive film formation.These corrosion products were incorporated into the passive film,which changed the local structure and the components of the pas-sive film and lowered its stability. Therefore, the nanometric a-Alprecipitated metallic glass has a low corrosion resistance than thatof fully metallic glass.

Acknowledgements

This work was supported by the National Natural Science Foun-dation of China (Nos. 51131006 and 51071151). The authors wouldlike to thank Dr. S.J. Wang for the assistance in TEM observationsand Dr. H.W. Yang for interesting discussions. The authors are alsograteful to David English for the modification of English.

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