The role of Cu-based intermetallics on the pitting corrosion behavior of Sn–Cu, Ti–Cu and Al–Cu alloys

Electrochimica Acta 77 (2012) 189–197

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The role of Cu-based intermetallics on the pitting corrosion behavior of Sn–Cu, Ti–Cu and Al–Cu alloys Wislei R. Osório a,b,∗ , Celia M. Freire b , Rubens Caram b , Amauri Garcia b a b

School of Applied Sciences/FCA, University of Campinas, UNICAMP, Campus Limeira, 1300 Pedro Zaccaria St., Jd. Sta Luiza, 13484-350 Limeira, SP, Brazil Department of Materials Engineering, University of Campinas, UNICAMP, P.O. Box 6122, 13083-970 Campinas, SP, Brazil

a r t i c l e

i n f o

Article history: Received 26 February 2012 Received in revised form 27 May 2012 Accepted 30 May 2012 Available online 7 June 2012 Keywords: Intermetallics Pitting corrosion Lamellar microstructure Sn–Cu, Ti–Cu, Al–Cu alloys

a b s t r a c t The aim of this study is to evaluate the effect of three different intermetallics; Cu6 Sn5 , Ti2 Cu and Al2 Cu, on pitting corrosion of Sn–2.8 wt.% Cu, Ti–5 wt.% Cu and Al–5 wt.% Cu as-cast alloys, respectively. Chillcast and centrifuged-cast samples were subjected to potentiodynamic polarization corrosion tests in a sodium chloride solution at 25 ◦ C. It was found that the intermetallic particles play an important role in the process of pitting corrosion. The results have shown that the pitting potentials (EPit ) of all three alloys experimentally examined have been displaced about 10 mV (SCE) toward the nobler-side potential when compared with the corresponding corrosion potentials (ECorr ). It was also found that the lamellar morphology formed by alternated solvent-rich/intermetallics phases, which characterizes the microstructure of any alloy examined, provides an enveloping effect giving rise to a protective barrier against corrosion. Other microstructural features associated with the ratio of anode/cathode areas for each alloy examined are also discussed. © 2012 Elsevier Ltd. All rights reserved.

1. Introduction Although the metallurgical and micromechanical aspects of factors controlling microstructure, unsoundness, strength and ductility of as cast alloys can be considered complex, it is well known that the solidification processing variables play an extremely important role. Besides, it is also known that both the macrostructural and microstructural morphologies influence the corrosion behavior. It was found that the improvement in the electrochemical corrosion resistance depends on the cooling rate imposed during solidification, which affects the scale of cellular and dendritic microstructures and the solute redistribution, and on the electrochemical behavior of solute and solvent of each alloy [1–7]. Recently, intermetallic compounds have attracted much attention mainly due to their potential technological applications as high-temperature materials. Moreover, microstructures combing a refined matrix and homogeneously distributed intermetallic particles may result in high hardness and, as a consequence, high mechanical strength. By examining the equilibrium phase diagrams of a number of important binary and multicomponent metallic systems, it can be seen that these diagrams predict the

∗ Corresponding author at: School of Applied Sciences/FCA, University of Campinas, UNICAMP, Campus Limeira, 1300, Pedro Zaccaria St., Jd. Sta Luiza, 13484-350 Limeira, SP, Brazil. Tel.: +55 19 3521 3320; fax: +55 19 3289 3722. E-mail address: [email protected] (W.R. Osório). 0013-4686/$ – see front matter © 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.electacta.2012.05.106

formation of intermetallic compounds for a wide range of alloys compositions. It is also expected that the intermetallic particles may have an important role on the alloy electrochemical behavior. For instance, a higher susceptibility to pitting corrosion can eventually be induced by the presence of intermetallic particles in the microstructure. In particular, alloys which are applied in the manufacture of reliable solder joints in microelectronics, aerospace (turbine) and automotive components [5,8–11], are good examples of practical situations where a better understanding of the influence of such intermetallics on the corrosion mechanisms can be used to improve operational conditions in order to manufacture components with appropriate final properties. Sn–Cu alloys are becoming interesting lead-free solder alternatives for the classical Sn–Pb solder alloy, since Pb is considered a poisonous metal with damage effects for human health [1]. New legislations worldwide demonstrate much concern about the use of lead in the manufacture of industrial parts [1,12,13]. Despite this potential of Sn–Cu alloys, few studies highlighting the effects of the nature and scale of the phases forming their microstructure on both mechanical and corrosion resistances can be found in the literature, which usually is much more concerned with the addition of a third element to the alloy composition [12,13]. In recent studies with a Sn–Cu solder alloy [3], it was reported that when a high cooling rate (of about 15 ◦ C/s) is applied, finely and homogeneously distributed Cu6 Sn5 intermetallic particles associated with fine inter-branch spacing (±3 ␮m) can be obtained. This resulting microstructural arrangement proved to have a deleterious effect on the

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electrochemical corrosion parameters [3]. On the other hand, when such alloy was more slowly cooled, a coarser inter-branch spacing (±10 ␮m) was produced and as direct consequence, a better electrochemical corrosion resistance was attained [3]. This resulting electrochemical behavior was shown to be intimately associated with the different growth natures of the Sn-rich (non-faceted) and the intermetallic (faceted) phases, which provoked certain strain in the atomic level along the boundaries between these phases favoring electrochemical corrosion. Furthermore, intermetallic particles having finer inter-branch spacings were shown to provide more accentuated pitting corrosion due to the higher number of galvanic couples that are formed in the microstructure [3]. Industrial applications of Ti–Cu alloys are reported in the literature since the 1950s. In the last 20 years, dental titanium casting technology has made considerable progress. It has been reported that Ti–Cu alloys may show reasonable mechanical strength associated with good formability [14]. It is important to remark that Ti–Cu alloys attain yield strengths, which are considerably higher than that of c.p. Ti and of most of the dental casting Co–Cr alloys [14]. It is known that melting and casting of Ti alloys are not simple tasks as these processes are associated with several complexities, including high melting temperature, affinity with interstitial elements and considerable activity with mould materials. The use of Cu as an alloying element can be an alternative method permitting these aforementioned difficulties to be minimized as well as to improve the mechanical behavior of c.p. Ti. Electrochemical studies will be very useful to establish correlations between the resulting Ti alloy microstructure and the kinetics of electrochemical corrosion and oxide film formation. In a recent investigation [15], the general electrochemical corrosion resistances of three Ti–Cu alloys were evaluated. It was concluded that with the increase in Cu content of the Ti–Cu alloy, the volumetric fraction of Ti2 Cu intermetallic particles also increases. Although the increase in Cu content significantly decreased the electrochemical performance of Ti–Cu alloys, it was evidenced that heat treatment can provide better electrochemical corrosion behavior when compared with that of a centrifuged cast sample [4]. Copper is added to aluminum alloys to increase their strength, hardness, fatigue, creep resistances and machinability. It is well known that Cu generally reduces the resistance to general corrosion and, in specific compositions and material conditions, the stress corrosion susceptibility [16]. Al–Cu hypoeutectic alloys (commercially classified as 2xxx series) have microstructures formed mainly by an Al-rich matrix and a eutectic mixture constituted by the Alrich phase and Al2 Cu intermetallic particles. This intermetallic has a cathodic behavior with respect to the Al-rich phase [5,16,17]. Over the years, a number of studies have been carried out in order to assess the effect of specific intermetallic particles and individual alloying additions upon the corrosion resistance of Al alloys [5,17,18]. Osório et al. [5] reported that Al2 Cu intermetallic particles are more susceptible to corrosion action than pure aluminum. Besides, it was also reported that in the as-cast condition, the Al2 Cu particles were enveloped by the Al-rich phase in the alternated eutectic mixture, which behaved as a protection against corrosion of the intermetallics. In this sense, smaller dendritic spacings and hence smaller eutectic interphase spacings have proved to provide a more extensive distribution of the “protective barrier”. The present study aims to contribute to the understanding of the effects of Cu-based intermetallics on the corrosion resistance of alloys representative of three different important metallic systems, i.e. Sn–2.8 wt.% Cu, Ti–5 wt.% Cu and Al–5 wt.% Cu alloys. The microstructures of these chemical compositions will be analyzed in the as-cast condition and after the corrosion tests in order to evaluate the role of the intermetallics with respect to the alloy matrix.

2. Experimental procedure Commercially pure (c.p.) metals were used to prepare the Sn–2.8 wt.% Cu, Ti–5 wt.% Cu and Al–5 wt.% Cu alloys samples: (i) Sn (99.993 wt.%) and Cu (99.991 wt.%); (ii) Ti (99.86 wt.%) and Cu (99.99 wt.%); and (iii) Al (99.72 wt.%) and Cu (99.93 wt.%). The impurities detected were commonly lower than 0.007 wt.% and less than 100 ppm, as previously reported [3,5,15]. Because of the complexities usually associated with Ti-castings, the Ti–Cu alloy sample was melted in an arc-melting furnace with a non-consumable tungsten electrode and water-cooled copper hearth under ultra-pure argon atmosphere [15,19]. In contrast, for the Sn–Cu and Al–Cu alloys a controlled atmosphere was not necessary, and these alloys were melted (and solidified) in a directionally solidification apparatus in which heat is extracted only through a water-cooled bottom, promoting vertical upward directional solidification [5–7]. After arc-melted, a homogenized Ti–Cu alloy sample (of about 30 g) was centrifuged cast (rotation of about 1000 rpm) by using a permanent copper mold, as detailed in a previous study [15]. Specimens for metallographic examination were extracted from longitudinal sections of the samples, ground by using silicon carbide papers up to 1200 mesh, polished and etched to reveal the microstructure. Microstructural characterization was performed by using a scanning electron microscope (SEM, Jeol JXA 840A). The working electrodes for the corrosion tests consisted of Sn–Cu, Ti–Cu and Al–Cu alloys samples which were positioned at the glass corrosion cell kit, leaving a circular area in contact with a naturally aerated and stagnant 500 (±10) cm3 of sodium chloride solution at 25 ◦ C having a pH between 6.5 and 7.5. Solution molarities (concentrations) were: 0.5 M NaCl solution for both Sn–Cu and Al–Cu alloys and 0.15 M NaCl solution for the Ti–Cu alloy. The 0.15 M and 0.5 M NaCl solutions were chosen in order to simulate human physiological fluids reactions with Ti–Cu alloys and general corrosion behavior of Sn–Cu and Al–Cu alloys in aggressive environment, respectively, as prescribed by ASTM standards G3 and G31 [20,21]. The samples were further ground to a 1200 grit SiC finish, followed by distilled water washing and air drying before electrochemical measurements. The tests began after an initial delay of 30 min for the sample to reach a steady-state condition. Potentiodynamic polarization curves were obtained by stepping the potential at a scan rate of 0.1667 mV s−1 from −250/+250 mV (SCE) at open-circuit. Using an automatic data acquisition system, these polarization curves were plotted and both corrosion rate and potential were estimated by Tafel plots using both anodic and cathodic branches. Duplicate tests were carried out.

3. Results and discussion 3.1. Resulting as-cast microstructures Fig. 1 shows the corresponding binary phase diagrams of Sn–Cu, Ti–Cu and Al–Cu systems [22]. Black arrows are indicating the intermetallic chemical compositions of each binary alloy system. Although the Sn–Cu alloy system provides basically two distinct intermetallic compounds, i.e. ␧-Cu3 Sn with about 60 wt.% Cu (75 at.% Cu) and ␩-␩ -Cu6 Sn5 with about 40 wt.% Cu (±55 at.% Cu), only the Cu6 Sn5 intermetallics was examined in the present study, since it predominates due to the composition of the Sn–Cu alloy. The Cu6 Sn5 intermetallics, can be intimately correlated with the cooling rate applied during casting, as previously reported [3,23–25], and this intermetallics is randomly distributed in the Sn-rich matrix as “H” [23] or “M-shaped” particles [24].

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Fig. 2. Typical SEM micrographs of the Sn–2.8Cu, Ti–5Cu and Al–5Cu alloys which are respectively constituted by: (a) Sn-rich matrix and Cu6 Sn5 intermetallics evidenced by white arrows; (b) ␣-Ti phase combined with Ti2 Cu intermetallics, and (c) Al-rich dendritic matrix involved by an interdendritic mixture.

Fig. 1. Phase diagrams of Sn–Cu, Ti–Cu and Al–Cu alloys. Redrawn from Ref. [21].

The Ti–Cu phase diagram with the Ti2 Cu intermetallic compound (40 wt.% Cu) indicated by the black arrow is shown in Fig. 1(b). Although five other intermetallic compounds can also be formed, as predicted by the phase diagram, only Ti2 Cu was examined in this study since the XRD patterns only exhibited this phase, which can be associated with the alloy Cu content. Considering this Ti–5 wt.% Cu alloy, the intermetallic Ti2 Cu is commonly combined with the ␣-Ti rich phase in the alloy microstructure, which was formed from the ␤-Ti decomposition by the eutectoid transformation [15]. A martensitic morphology prevails since the rapid cooling during casting has limited the eutectoid coupled growth. A previous study [26] has reported a microstructural arrangement of Ti2 Cu and the Ti-rich phase in a lamellar morphology.

The black arrow in Fig. 1(c) indicates the composition of the intermetallic Al2 Cu (␪-phase with ±54 wt.% Cu). The microstructure of the examined Al–5 wt.% Cu alloy has a eutectic mixture formed by the ␪-phase associated with the Al-rich phase characterizing a lamellar alternation of each phase. These three foreshown phase diagrams evidence that the three examined intermetallics are Cu-base compounds, which permit comparison among their corresponding polarization curves, as well as the corresponding analyzes concerning their reactions with respect to the matrix phase (Sn, Ti and Al-rich matrices). SEM micrographs of the Ti–Cu centrifuged cast and chill-cast Sn–Cu and Al–Cu alloys samples are shown in Fig. 2. The ascast microstructure of the Sn–2.8 wt.% Cu alloy is constituted by a Sn-rich matrix (␣: solid solution of Cu in Sn) and Cu6 Sn5 intermetallics randomly distributed in the Sn-rich matrix, as shown in Fig. 2(a). Dark regions indicated with arrows are intermetallic particles, which are commonly characterized by “H or M-shaped” Sn–Cu intermetallics [24,25], while the light regions are formed by the

W.R. Osório et al. / Electrochimica Acta 77 (2012) 189–197

Sn-rich matrix. This Sn–Cu alloy sample corresponds to a position in the casting of about 8 mm from the water-cooled bottom (experimental details can be found in a previous study [3]). In this sense, fine Cu6 Sn5 particles (corresponding to an inter-branch spacing, , of about 3 ␮m) were found close to the casting cooled surface. The corresponding cooling rate (dT/dt) operative during solidification at this location in casting was about 15 (±2) ◦ C/s [3]. Fig. 2(b) evidences a typical SEM micrograph of the Ti–5% Cu alloy (hypoeutectoid) sample, which is constituted by a ␣-Ti phase combined with Ti2 Cu intermetallics, which were formed from the ␤-Ti decomposition by the eutectoid transformation. As expected, rapid cooling of ␤ grains limits the eutectoid coupled growth and gives rise to a martensitic pattern (lamellar morphology of ␣Ti + Ti2 Cu). Although lamellar Ti2 Cu particles have been observed in the examined Ti–Cu alloy sample, this intermetallic may also occur as spherical precipitates [15,26]. From these microstructural observations, it can be said that the Ti2 Cu particles have a lamellar morphology, from 1 to 2 ␮m thick and formed under a cooling rate of about 10 (±2) ◦ C/s [15,26]. Fig. 2(c) shows the micrograph of the Al–Cu alloy corresponding to a sample obtained at about 10 mm from the water-cooled bottom of the casting. The hypoeutectic Al–5 wt.% Cu alloy is constituted by an Al-rich dendritic matrix (␣-phase) involved by an interdendritic mixture (eutectic) of ␣ phase and Al2 Cu intermetallic particles following a lamellar alternation of each, as shown in Fig. 2(c). The secondary dendrite arm spacing (2 ) is about 15 ␮m and the Al2 Cu lamellae are between 1 and 2 ␮m thick. These lamellae have the ␣-Al rich phase with about 5.65 wt.% Cu and the Al2 Cu particles with 53.5 wt.% Cu, forming the eutectic mixture [5]. From these observations, it is interesting to remark that for the two alloys containing X2 Cu intermetallic compounds, i.e. Ti2 Cu and Al2 Cu, the intermetallic is immersed in a lamellar mixture alternating X-rich phase/intermetallics. The X-rich phases (e.g. Al: −850 mV, SCE) from the electrochemical point of view are less noble than the corresponding intermetallics (e.g. Al2 Cu: −695 mV, SCE). In contrast, the Sn–Cu alloy exhibits an abrupt variation in the Cu content at the Cu6 Sn5 /Sn-rich phase interface [3]. However, the Cu6 Sn5 intermetallic compound is also considered nobler than the Sn-rich phase. In order to evaluate the microstructural area ratio between the intermetallic compound and the corresponding matrix, selected images of the microstructures of each alloy were converted into binary images and 10 measurements were carried out using the software ImageJ® . These area ratios will be very useful to the discussions concerning the potentiodynamic polarization curves. The following area ratios between the Sn, Ti and Al-rich matrix and the corresponding intermetallics have been obtained: Cu6 Sn5 /Sn (1:9), Ti2 Cu/Ti (1:2) and Al2 Cu/Al (1:4). From these results, it can be seen that the alloys characterized by lamellar arrangements present higher intermetallics/solvent-rich phase area ratios. It is important to remark that a higher area ratio is associated with a more homogenous distribution of the lamellar intermetallics in the matrix. 3.2. Polarization curves and microstructure arrays Fig. 3 shows potentiodynamic polarization curves (from −1000 mV to 0.0 mV, SCE) for the three examined Sn–Cu, Ti–Cu and Al–Cu alloys samples, which were carried out in a stagnant and naturally aerated sodium chloride solution at room temperature 25 (±3) ◦ C. The corrosion current densities (i) were obtained by Tafel extrapolation considering both the cathodic and anodic branches from each potential interval of the polarization curves and are shown in Fig. 4. In order to minimize the effects of distortion in Tafel slopes and current densities, small scan rates were used, as suggested in the literature [27].

0,00 Potential ( E ) / Volts (SCE)

192

Cu-based intermetallics Sn-2.8Cu

-0,25

Ti-5Cu Al-5Cu

-0,50

-0,75

-1,00 -10 -9 10 10

-8

10

-7

10

-6

10

-5

10

-4

-3

10

Current density ( i ) / A cm

10

-2

10

-1

10

-2

Fig. 3. Experimental potentiodynamic polarization curves carried out in a stagnant and naturally aerated sodium chloride solution at room temperature 25 (±3) ◦ C for the three examined Sn–Cu, Ti–Cu and Al–Cu alloys samples.

It is important to remark that the aim of the present investigation is not focused on the comparison among potentiodynamic polarization curves, but on the effects of each intermetallics with respect to the matrix and the corresponding pitting behavior. By analyzing the three potentiodynamic polarization curves between 0 and −1000 mV, SCE, it can be said that the Sn–Cu alloy shows a transient passivity phenomenon initiated at about −588 (±2) mV (SCE) which seems to indicate corrosion of the Sn-rich phase, formation of tin oxide (Sn II oxide) and possibly initiation of mechanisms of precipitation and dissolution of a number of corrosion products, e.g. SnCl2 , Sn(OH)4 , SnO, SnCl−3 and SnCl6 −2 , as previously reported [28,29]. Considering the Ti–Cu alloy sample, it is observed a partial stabilization in current density in the range 5–7 × 10−8 A cm−2 , which is associated with the range −420 and −450 mV, SCE. Slight potential turbulences between −315 and −140 mV (SCE) are also observed. These potential breakdowns can be associated with passive oxide film forming and dissolving. It is important to remark that the Ti oxide formation become more stable near +500 mV (SCE) [15]. A primary passive current density of about 5 ␮A cm−2 at about +500 mV (SCE) has been reported [15]. The potentiodynamic polarization curve of the Al–Cu alloy sample is also shown in Fig. 3. It can be seen that the anodic current density increases rapidly of about 2 decades (from 2 × 10−6 A cm−2 to 2 × 10−4 A cm−2 ) at −675 mV to −640 mV, SCE. This Al–Cu alloy is the only alloy showing a plateau of cathodic current density, occurring between 7 and 5 ␮A cm−2 which initiated from −936 mV to −700 mV, SCE. It seems that this plateau can be associated with a limiting cathodic current density for dissolved gases reactions, as reported by Birbilis and Buchheit [17,18,30]. Fig. 4 shows the corresponding corrosion current densities (i), corrosion potentials (ECorr ), pitting potentials (EPit ) and differences (E) between ECorr and EPit for each examined alloy. These values are also shown in Table 1 associated with anode/cathode area ratios and estimated galvanic potentials. Fig. 4(a) shows the potentiodynamic polarization curve of the chill-cast Sn–Cu alloy sample from potentials of −750 mV to −650 mV (SCE). It is clearly evidenced a slightly partial stabilization in the current density of about 2 × 10−8 A cm−2 at −715 mV (SCE) and a range of perturbations in the potential in the range −707 mV to −702 mV (SCE), which characterizes pitting initiation. Evidently this perturbation is associated with the galvanic couple Sn//Cu, i.e. Sn-rich and Cu6 Sn5 intermetallics regions. A couple ␣-Sn//Cu6 Sn5 seems to produce a galvanic potential driving-force of about 220 mV (SCE) since the estimated corrosion potentials of ␣-Sn and Cu6 Sn5 are about −520 mV and 300 mV

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Table 1 Characteristics and effect of microstructure array of three examined Sn–Cu, Ti–Cu and Al–Cu alloys on the corrosion behavior. Characteristics

Alloys Sn–2.8Cu

Ti–5Cu

Al–5Cu

Microstructure

Sn-rich matrix (␣ solid solution Cu in Sn) + ␩-phase Cu6 Sn5 (H-shaped) N/Aa

␣-Ti matrix (martensite) + Ti2 Cu (lamellae) Decrease

␣-Al matrix (␣ solid solution Cu in Al) + Al2 Cu (lamellae, eutectic mixture) Decrease

Decrease

Increase

Increase

1:9 −715 mV −707 mV 0.15 ␮A cm−2 Localized deformation (faceted/non faceted phases) ␣-Sn//Cu6 Sn5 520 mV//300 mV  = 220 mV

1:2 −515 mV −507 mV 0.10 ␮A cm−2 Ti2 Cu enveloping protection

1:4 −678 mV −670 mV 4.2 ␮A cm−2 Al2 Cu enveloping protection

␣-Ti//Ti2 Cu 20 mV//100 mV  = 80 mV

␣-Al//Al2 Cu 690 mV//530 mV  = 160 mV

Increased alloy Cu content (effect on corrosion resistance) Fine microstructure (effect on corrosion resistance) Cathode/anode area ratio ECorr b EPit i Pitting corrosion (driving-force) Galvanic couple (corrosion potential)

a b

Not available. Values are SCE referenced.

(SCE), respectively, as shown in Table 1. This Sn–2.8 wt.% Cu alloy exhibits a corrosion current density of 0.15 (±0.01) ␮A cm−2 with ECorr about −717 (±2) mV (SCE). As reported by Zaid et al. [31], a pitting potential can be coincident with or near to the corrosion potential when the anodic Tafel slopes are practically zero. In this context, from the polarization curves shown in Fig. 4, the Sn–Cu alloy evidences an EPit displaced about 10 mV (SCE) toward the nobler-side potential than its corresponding ECorr . Fig. 4(b) shows the corrosion current density of 102 (±8) × 10−9 A cm−2 with ECorr and EPit about −515 (±2) mV and −507 (±2) mV (SCE), respectively for the Ti–5 wt.% Cu alloy sample. Pitting seems to be initially provoked by the galvanic couple Ti//Ti2 Cu, i.e. Ti-rich phase and lamellae of the corresponding intermetallics. In this sense, the Ti-rich phase has a corrosion potential of about −20 mV (SCE) while that of the Ti2 Cu intermetallics is about −100 mV (SCE), forming consequently the lowest galvanic couple (80 mV) of all examined alloys. Fig. 4(c) shows the potentiodynamic polarization curve of the chill-cast Al–5 wt.% Cu alloy which depicts the highest current density (of about 4 ␮A cm−2 ) with ECorr and EPit about −679 mV and −670 mV (SCE), respectively. Its corresponding galvanic couple potential is about 160 mV (SCE), the ␣-Al phase (in the eutectic mixture) has a corrosion potential about −690 mV (SCE) and the Al2 Cu intermetallics about −530 mV (SCE). These corrosion potentials are similar to those reported by Birbilis and Buchheit [18] for a 0.6 M NaCl solution. Besides, it is important to point out that EPit is ever displaced toward the noble-side potential when compared with ECorr , as also reported by Birbilis and Buchheit [18]. For instance, these authors indicated an ECorr of −696 mV (SCE) and EPit of −652 mV (SCE) for Al2 Cu intermetallics in a 0.6 M NaCl solution. Evidently that these experimental results of ECorr , EPit and current densities should not been considered as fixed references, since both the electrolyte characteristics (e.g. concentration, type of electrolyte and pH) and the microstructural array (e.g. dendritic spacing, second and intermetallic phases) have to be considered [3,5,15]. Besides, other electrochemical parameters (obtained from Bode and Nyquist plots) can also be used in order to give support to the discussions and to permit the electrochemical behavior of the three alloys examined to be better understood. For example, when as-cast Al–Cu alloys are considered, it has been reported that the corrosion resistance increases with refined microstructural arrays, i.e. with lower dendritic arm spacings [5]. A lower corrosion resistance has been reported for coarser dendritic structures of an Al–4.5 wt.% Cu. This has been attributed to a more homogeneous distribution of Al2 Cu intermetallics in the

Al-rich phase (both in the eutectic mixture and the dendritic Al-rich matrix) [5]. In this sense, when high cooling rates are imposed during casting, finer dendritic arrangements will be produced. In such case, the eutectic mixture (which includes the Al2 Cu particles) will be more extensively distributed along the finer interdendritic arms. Such intermetallics (when distributed in a fine dendritic arrangement) “envelope” the Al-rich phase which acts as a protection against corrosion (protective barrier against corrosion). It has also been reported [5] that as the alloy Cu content increases, a higher susceptibility to corrosion action in a NaCl solution is detected. This was attributed to the increase in the fraction of Al2 Cu particles [32]. In contrast, tests carried out in a H2 SO4 solution have evidenced similar corrosion rates for three alloys with different Cu contents, i.e. 5, 10 and 15 wt.% Cu [32]. The effect of alloy Cu content associated with the intermetallics has also been observed for Ti–Cu alloys. It has also been concluded that with the increase in alloy Cu content, the volumetric fraction of Ti2 Cu also increases, i.e., with lower alloy Cu content a more homogeneous and smoother oxide film is provided [15]. Considering the aforementioned assertions, some correlations between microstructure (␣: solvent-rich phases and intermetallics) and corrosion behavior (i.e. ECorr , EPit , i) for the three examined Sn–Cu, Ti–Cu and Al–Cu alloys have been made. Table 1 synthesizes microstructural conditions of ␣-rich phases and the corresponding intermetallics, as well the effects of fine microstructure array associated with anode/cathode area ratios, galvanic couples and Cu content on the corrosion resistance in a sodium chloride solution. ECorr , EPit and corrosion current densities, used to discuss the mechanism of general and pitting corrosion, respectively, are also shown in Table 1. It can be clearly observed that both the Al–Cu alloy and the Ti–Cu alloy present the ␣-phase (solventrich) and the corresponding X2 Cu intermetallic particles alternated in a lamellar microstructure, and that the corrosion resistance increases for a refined microstructure and decreases with the increase in the alloy Cu content. Although the intermetallics and the ␣-phase lamellae constitute galvanic couples, finer microstructures provide an “enveloping effect” between these phases, which minimize the galvanic effects improving the corrosion resistance. These Ti–Cu and Al–Cu alloys also have cathode/anode area ratios (1:2 and 1:4), which are considerably higher than that of the Sn–Cu alloy (1:9). However, it is not ever true that both Ti–Cu and Al–Cu alloys exhibit only cooperative growth mechanisms resulting in a lamellar morphology. This condition is strongly dependent on thermodynamics and chemical characteristics, such as alloy Cu content,

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Fig. 4. Experimental potentiodynamic polarization curves evidencing corrosion current densities (i), corrosion potentials (ECorr ), pitting potentials (EPit ) and differences (E) between ECorr and EPit for: (a) chill-cast Sn–2.8 wt.% Cu alloy, (b) centrifuged cast Ti–5 wt.% Cu alloy, and (c) chill-cast Al–5 wt.% Cu alloy.

cooling rate (in case of castings or soldering applications), heat treatment conditions and some possible interactions between these phases with other phases or other elements constituting the alloy composition. For instance, it was reported [15] that heattreated samples of a hypereutectoid Ti–15 wt.% Cu alloy gave rise to a eutectoid transformation when slowly cooled. This has resulted in a morphology considerably different of the typical martensite, with Ti2 Cu spheroids dispersed in the eutectoid microstructure [15]. SEM images after the corrosion tests are shown in Fig. 5. Typical images of the Sn-rich phase and of Cu6 Sn5 particles (Sn–2.8 wt.% Cu alloy) are shown in Fig. 5(a) and (b). It can be seen a corroded microstructure basically constituted by 4 elements: i. preserved

Cu6 Sn5 intermetallics (mixture: H-shaped + fine hexagonal); ii. thin Sn oxide layers in the Sn-rich matrix; iii. slight cavities from where the Cu6 Sn5 particles detached from the microstructural arrangement; and iv. eutectic mixture (lamellae of Sn + Cu6 Sn5 ) located inside the aforementioned cavities. Fig. 6 shows high magnification SEM images (post-corroded) of the examined Sn–Cu alloy evidencing primary Cu6 Sn5 intermetallic particles (H-shaped) and pits with fiber-like Cu6 Sn5 intermetallics (of about 0.5 ␮m thick) forming the eutectic mixture (Sn-rich + Cu6 Sn5 ). It is known that during solidification, the Snrich phase grows from the liquid in a non-faceted manner (growth with a rough interface) while the Cu6 Sn5 solidifies with interfaces that are smooth [3,23,33]. Recently, Gourlay et al. [34] have reported the growth of faceted primary Cu6 Sn5 intermetallics and non-faceted Sn-rich dendritic matrix. As a result, the boundaries (interfaces) between these two phases are subjected to a localized strain (will not be perfectly conformed) and will be more susceptible to corrosion than the “H-shaped” Cu6 Sn5 intermetallics (nobler areas). This situation will accelerate the corrosion of the Sn-rich phase surrounding the intermetallics provoking the detachment of these particles and forming as a result both small spheroid-like pits and elongated cavities, as shown in Fig. 5(a) and (b). Table 1 indicates that this examined Sn–Cu alloy has a galvanic couple between ␣-Sn//Cu6 Sn5 of about 220 mV (SCE) which seems to be a driving-force which is sufficient to promote pitting in the Snrich matrix. Another interesting observation is concerned to the presence of the eutectic mixture inside these cavities which was not previously observed by using BSEM (back scattered electron microscopy). C¸adirli et al. [35] have also reported the formation of Cu6 Sn5 intermetallics (Sn–39.61 wt.% Cu) and a eutectic mixture in a hypereutectic Sn–3 wt.% Cu alloy. Panchenko et al. [36] have used BF (bright field) and CP (cross polarized) contrast in order to characterize Sn–Cu solder alloys. They reported a resulting microstructure consisting of Sn-rich cells and dendrites with small Cu5 Sn6 particles into the interdendritic regions [36]. They have observed that only a Sn–3 wt.% Cu alloy has provided large Cu5 Sn6 intermetallics randomly distributed in the Sn-matrix lattice which are similar to those called as “H-shaped” examined in the present investigation. It was noticed that other dilute Cu content Sn–Cu alloys, e.g., Sn–1.5Cu, Sn–0.9Cu, Sn–0.5Cu and Sn–0.25Cu solder alloys had no evidence of this “H-shaped” morphology [36]. From these observations, it seems that after the detachment of the H-shaped intermetallics (with about 40 wt.% Cu content) from the microstructure, thus exposing the eutectic mixture (with about 0.9 wt.% Cu), the driving-force to galvanic corrosion decreased. Fig. 5(c) shows a typical SEM image of a Ti–5 wt.% Cu alloy. Although it has been reported [4,37–39] that a very protective Ti oxide film can be formed on the surface of Ti–Cu alloy samples, some defects on this oxide can be observed, as shown in Fig. 5(c). These defects can be attributed to slight potential turbulences observed in the potentiodynamic anodic polarization plot which indicate a tendency of formation of more irregular or porous oxide layer [4,35–37]. From the point of view of the bone-implant interface (applications in both orthopedic and dental prostheses), some minimal roughness is required in order to privilege both bone ingrowth and bone retention [15,40]. On the other hand, from the electrochemical point of view, the Ti-rich regions (␣-Ti) are expected to be nobler regions when compared with the Cu-rich regions, as shown in Table 1. Again, it seems that the lamellae of Ti2 Cu (nobler region) “envelopes” the Ti-rich phase (less noble region) providing corrosion protection. From these aforementioned observations about the present experimental impedance parameters and the resulting microstructures, it can be concluded that the eutectoid transformation (␣-Ti + Ti2 Cu) has an important role on the corrosion behavior of the Ti–5 wt.% Cu alloy.

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Fig. 5. Typical SEM images after corrosion tests for: (a) Sn–2.8 wt.% Cu alloy, (b) Ti–5 wt.% Cu, and (c) Al–5 wt.% Cu alloy.

By examining the images shown in Fig. 5, it can be seen that the Sn–Cu and Ti–Cu alloys evidence more compact oxide layers when compared with that of the Al–Cu alloy. This seems to be associated with the lowest corrosion current density, as shown in Table 1. However, considering the pits areas, it seems that the Sn–Cu alloy has more corroded areas, the Al–Cu alloy is in an intermediate position and the Ti–Cu alloy exhibited less corroded areas, as evidenced in Fig. 5. This seems to be associated with the corresponding values of galvanic corrosion potential and EPit for each alloy, as shown in Table 1. Fig. 5(d) and (e) evidences pits in the microstructure of the Al–Cu alloy caused by the difference in pit corrosion potentials between the aluminum-rich phase and the intermetallic particles. It seems that these pits are formed on the surroundings of the intermetallics that finally detached from the alloy microstructure. Although the Al–Cu alloy has lower incidence of pit areas than the Sn–Cu alloy, its corresponding pits are deeper and slimmer than those observed for the Sn–Cu alloy. In this context, if these intermetallic particles are distributed in a fine dendritic array, they will “envelope” the Al-rich lamellae (in the eutectic mixture) providing a more extensive distribution of the “protective barrier”. It can be said that the cathode intermetallic area is considerably

lower than or similar to the anode area, which induces slower corrosion. Considering the resulting microstructural arrays, pitting corrosion potentials (galvanic couples) and anode/cathode area ratios, it can be concluded that the Cu content and the cooling rate during casting, which determines the scale of the microstructure arrays, are important parameters that should be considered in the manufacture of components with a view to improving their pitting corrosion resistances. From these experimental observations, it can be said that a finer microstructure array improves the corrosion resistance. This can be easily concluded when both Ti–Cu and Al–Cu alloys are considered, because of the aforementioned enveloping protection effect. However, this is not directly true when the examined Sn–Cu alloy is considered. Although it has been previously reported [3] that coarser microstructures in a Sn–2.8 wt.% Cu solder alloy provided better electrochemical corrosion resistance than finer ones, it can be supposed that if a higher number of fine “H-shaped” Cu5 Sn6 intermetallic particles could be produced, a better pitting corrosion resistance could also be attained. This is based in the increase of the cathode/anode (c/a) area ratio which was shown to be 1:9 in the present study. For instance, if this present c/a area ratio could be increased three times (1:3) (i.e. finer

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used processing condition, it is suggested the use of operational conditions conducive to refined microstructure arrays. Although the intermetallics and the ␣-phase lamellae constitute galvanic couples, fine microstructures provide an “enveloping effect” modifying the cathode/anode area ratio between these phases which minimize the galvanic effects improving the corrosion resistance. However, when a hypereutectic alloy is considered, as it is the case of the Sn–2.8Cu alloy, due to primary intermetallics formation, a coarse microstructure will probably be associated with a better corrosion resistance. 2. The pitting potential (EPit ) of all three examined alloys was displaced about 10 mV (SCE) toward the nobler-side potential when compared with the corresponding corrosion potential (ECorr ). 3. The microstructure of the Sn–Cu alloy is formed by a eutectic mixture constituted by the Sn-rich phase and thin (±0.5 ␮m thick) fiber-like Cu6 Sn5 associated with primary faceted Cu6 Sn5 intermetallics (H-shaped). The interfaces between the H-shaped intermetallics and the Sn-rich matrix are characterized by a localized strain caused by the different growth mechanisms of each phase, thus increasing corrosion. However, when the primary Cu6 Sn5 intermetallic particles were detached during the corrosion process, cavities containing a lamellar microstructure (eutectic mixture) could be observed, which seem to be the responsible for the reduction in corrosion after the detachment of these intermetallics. Acknowledgements The authors acknowledge the financial support provided by CNPq (The Brazilian Research Council), FAEPEX-UNICAMP and FAPESP (The Scientific Research Foundation of the State of São Paulo, Brazil). References Fig. 6. SEM micrographs of the Sn–Cu alloy evidencing the primary Cu6 Sn5 intermetallic particles (H-shaped) and pits with eutectic mixture (lamellae of Sn-rich phase + Cu6 Sn5 ).

and more homogeneously distributed intermetallic particles in the alloy matrix), that could probably displace the pitting corrosion potential toward the nobler-side reducing pitting and increasing the enveloping effect. This is certainly expected when a hypoeutectic Sn–Cu alloy is considered, since only a eutectic mixture of fine needles or fiber-like Cu6 Sn5 particles and Sn-rich phases are attained. Another supposition could also be made concerning the coarsening of microstructures of both Al–Cu and Ti–Cu alloys. If the present cathode/anode area ratios could be reduced (1:8 and 1:4, respectively), it is expected that the coarser lamellae would decrease the “enveloping protection” considering the c/a area ratio and thus decrease the corrosion resistance. 4. Conclusions Based on the present experimental results, which include the microstructural array imposed by the cooling rate during casting (which defines the anode/cathode area ratio and the galvanic couple) associated with potentiodynamic polarization curves, the following conclusions can be drawn: 1. Cu-base intermetallic particles, such as Cu6 Sn5 , Ti2 Cu and Al2 Cu have important roles on the resulting pitting corrosion responses of Sn–2.8 wt.% Cu, Ti–5 wt.% Cu, and Al–5 wt.% Cu alloys, respectively. Since the Cu-base intermetallic compounds are always present in the microstructures of these alloys, no matter the

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