Localized corrosion at nm-scale hardening precipitates in Al-Cu-Li alloys

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Localized Corrosion at nm-Scale Hardening Precipitates in Al-Cu-Li Alloys Yakun Zhu , Jonathan D. Poplawsky , Sirui Li , Raymond R. Unocic , Leslie G. Bland , Christopher D. Taylor , Jenifer S. (Warner) Locke , Emmanuelle A. Marquis , Gerald S. Frankel PII: DOI: Reference:

S1359-6454(20)30183-X https://doi.org/10.1016/j.actamat.2020.03.006 AM 15893

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Acta Materialia

Received date: Revised date: Accepted date:

7 December 2019 30 January 2020 4 March 2020

Please cite this article as: Yakun Zhu , Jonathan D. Poplawsky , Sirui Li , Raymond R. Unocic , Leslie G. Bland , Christopher D. Taylor , Jenifer S. (Warner) Locke , Emmanuelle A. Marquis , Gerald S. Frankel , Localized Corrosion at nm-Scale Hardening Precipitates in Al-Cu-Li Alloys, Acta Materialia (2020), doi: https://doi.org/10.1016/j.actamat.2020.03.006

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Localized Corrosion at nm-Scale Hardening Precipitates in Al-Cu-Li Alloys Yakun Zhu1,2, Jonathan D. Poplawsky3, Sirui Li1, Raymond R. Unocic3, Leslie G. Bland 1, Christopher D. Taylor1, Jenifer S. (Warner) Locke1, Emmanuelle A. Marquis2, Gerald S. Frankel1*1 1 2

Fontana Corrosion Center, The Ohio State University, Columbus, OH 43210

Department of Materials Science and Engineering, The University of Michigan, Ann Arbor, MI 48109 3

Center for Nanophase Materials Sciences, Oak Ridge National Laboratory, Oak Ridge, TN 37830

Graphical abstract

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* Corresponding author: Gerald S. Frankel, email: [email protected]

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Abstract The localized corrosion of Li-containing nm hardening precipitates in the 3rd generation of AlCu-Li alloys was investigated based on a quasi in situ approach by sequentially exposing the material to NaCl solution and characterizing the structural, chemical, and electrochemical evolution at atomic scale using electron microscopy, spectroscopy, 3D tomography, electrochemical measurements, and DFT calculations. Localized corrosion of Al7.5Cu4Li (TB phase) initiated along {001} family of planes through the dealloying of Al and Li due to a low surface work function. Cu was enriched along the Cu (110) // TB (011) // Al (100) orientations on and around corroded TB precipitates. No strong galvanic interactions were observed at the TB and Al matrix interface due to the formation of a Li-C-O rich passivation layer during electrolyte exposure. Similarities and differences between TB and other common Al-Cu-Li precipitates (Al2CuLi, Al6CuLi3, and Al3Li) with respect to corrosion are discussed. The reported corrosion mechanism can assist in the assessment of the localized corrosion susceptibility of precipitationhardened Al alloys and assist in the design of new alloys. Keywords: Al-Cu-Li hardening precipitate, atomic/nm-resolution analysis, localized corrosion, passivation, electrochemical methods 1. Introduction The 3rd generation of lightweight Al-Cu-Li alloys, such as AA2x99 and AA2060, are designed to provide low density, high specific strength and specific modulus, and excellent fatigue-resistance performance, which makes these alloys ideal for structural materials for aerospace applications [1, 2]. However, the alloys still remain susceptible to forms of localized corrosion associated with these intermetallic phases such as pitting, trenching, and microgalvanic corrosion in corrosive environments. The phases of interest include nm to sub-μm precipitates, 10 to 100 nm scale dispersoids (e.g. Zn, Mn, or Cr containing), and large 10 to 100 μm scale intermetallic particles (IMPs) [3-5]. Therefore, it is necessary to evaluate the corrosion behavior of these intermetallic phases upon exposure to aggressive environments to better understand the corrosion mechanism and to devise better corrosion mitigation strategies. Large constituent IMPs e.g. Fe-containing particles can be detrimental to both mechanical properties and corrosion resistance [1, 3, 5]. Recent studies have addressed their role in the localized corrosion of Al alloys e.g. AA2024-T3, AA2060-T8, AA2070-T8, and AA7075-T651 2

[3, 4, 6-14]. For example, the micro-galvanic interaction between the IMPs and the matrix plays a dominant role in the development of trenches around the IMPs. Dealloying occurs in the compounds themselves by the corrosion of relatively active elements e.g. Al or Mg, which increases the nobility of the IMPs [5, 11]. The hydrolysis of Al3+ ions in the trench generates H+ ions, resulting in an acidic environment that supports hydrogen evolution (HE) [5]. The nm-scale precipitates, including Al7.5Cu4Li (TB), Al2CuLi (T1), Al6CuLI3 (T2), Al3Li (), AlLi (), Al2Cu (), Al2CuMg (S), and MgZn2 [15, 16], are of importance because, unlike the large IMPs, they play a critical role in matrix hardening. However, a detailed understanding of the very early stage of corrosion, in particular of Li-containing precipitates and their interactions with the matrix, is limited [17-19]. Whether nm scale precipitates generate the above corrosion phenomena and how they proceed have never been reported. Accurate

corrosion

assessment

of

hardening precipitates

through

electrochemical

measurements, however, is very challenging due to the nm size. Studies of the corrosion of hardening precipitates have been performed, including relatively high-resolution examination of surfaces, [20, 21] but these works were not at the atomic/nm level. The presence of Li further increases the difficulty in capturing the chemistry of localized corrosion phenomena due to the very active electrochemical nature of Li. Moreover, separate electrochemical processes both at these precipitates and in the surrounding matrix, create complex localized environments with steep gradients. Therefore, techniques that provide high spatial resolution and high precision in chemical and structure analysis are necessary to probe nm-scale phenomena associated with localized corrosion of these hardening precipitates. This study, for the first time, addresses the corrosion behavior of nm-scale Al-Cu-Li hardening precipitates, primarily TB precipitate, at the early stage of corrosion and the interaction with, and the propagation into, the Al matrix in corrosive environments at atomic resolution. The nanostructure and electrochemical evolution of TB precipitate during electrolyte exposure was investigated using a quasi in-situ scanning transmission electron microscopy (STEM) on thin foils in combination with electrochemical measurements. Atom probe tomography (APT) was adopted to explore the atomic level elemental distribution. A mechanism is proposed to describe the changes of nm structural, chemical, and electrochemical properties of the TB hardening precipitate in the corrosion process, which can assist materials degradation assessment and 3

provide input for the design of new alloys. Similarities and differences between TB and other common Al-Cu-Li precipitates (Al2CuLi, Al6CuLi3, and Al3Li) with respect to corrosion are discussed. 2. Experimental 2.1. Materials A model alloy with the nominal composition Al-3Cu-1.5Li (wt%) was purchased from ACI Metals. Composition analysis of the precipitate-free region in the matrix based on APT gives 1.2 wt% Li, 2.9 wt% Cu, and 96 wt% Al. The as-received alloy was solution treated at 515 C for 24 h followed by aging at 180 C for two weeks to produce nm-scale precipitates. This model alloy was then mechanically polished to a 1 m surface finish and cut into standard 3-mm diameter disc TEM samples with thickness of around 100 m. The disc samples were mechanically polished to a 0.25 m surface finish in diamond paste and then electropolished in 75% methanol + 25% HNO3 at 30 V and -35 C until they were electron-transparent. Ar ion milling was then applied to the sample surface to eliminate the oxidation layer formed during electropolishing. In the corrosion exposure test, the as-prepared 3-mm disc samples were immersed in 0.1 M NaCl aqueous solution and removed after 2.5 min, 15 min, 30 min, 120 min, and 15 h, rinsed with water, and air dried in the lab environment. 2.2. Tomography examination Topographic maps of the surfaces of 3-mm disc samples before and after corrosion exposure were collected at room temperature (~21 C) in air using a Veeco MultiMode 8 atomic force microscope (AFM) with a Nanoscope V controller. The scanning tip (Bruker OSCM-PT-R3) was coated with platinum on the tip apex. The scan frequency was set at 0.25 Hz with 512 pixels per scanning line. The images were post-analyzed using the software NanoScope Analysis 1.5. 2.3. Structure and chemical examination For atomic resolution scanning transmission electron microscope (STEM) imaging of the electro-polished surface near perforation holes of 3-mm discs before and after sequential exposure to 0.1 M NaCl, a CS aberration-corrected FEI Titan TEM S/TEM operated at 300 kV

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was used. Plasma cleaning was performed on the sample before STEM imaging of the thin region in as-prepared 3-mm disc samples. An FEI Helios 650 Nanolab Dual Beam FIB-SEM was used to lift out thin foil cross sections of the corroded 3-mm disc samples for STEM imaging. The accelerating voltage initially was set at 30 kV for cross-sectioning and trenching, 5 kV for thinning, and 2 kV for final cleaning. A Cs aberration-corrected JEOL 2100 AEM was used for STEM imaging and EDS mapping of the cross-sections of the corroded samples. Nano beam diffraction (NBD) was conducted using JEOL 2010F analytical electron microscope operated at 200 kV. Bright field (BF), dark field (DF), high angle annular dark field (HAADF) images, and diffraction patterns were captured and analyzed using Digital Micrograph. Before performing APT needle preparation liftouts, the corroded samples were coated with 100 nm of Ni using a South Bay Technologies Ion Beam Deposition System. Traditional liftout and sharpening methods using a Si microtip array were used for APT sample preparation [22]. A FEI Nova 200 dual beam FIB-SEM, operated at 30 kV for trenching and sharpening and 2 kV for final cleaning, was used for the APT needle fabrication. After FIB sample preparation, APT analyses were performed using a CAMECA LEAP 4000X HR. The LEAP was operated in laserpulse mode at 25 K, a 30-50 pJ laser pulse energy, a pulse repetition rate of 200 kHz to allow for all ions to be collected in the time window for the mass spectrum, and detection rate of 1 ion per 200 pulses. The resulting data were reconstructed and analyzed using CAMECA’s IVAS software. 2.4. Electrochemical measurements Bulk model alloys with a same composition as the representative precipitates, Al7.5Cu4Li, Al2CuLi, Al6CuLi3, and Al3Li, were purchased from ACI Metals. Potentiodynamic polarization measurements were performed on the surfaces of model alloys, pure Al (99.999%), Al-3Cu1.5Li alloy, and AA2070 alloy in 0.1 M NaCl using the same microcell setup in a previous relevant work [7]. Measurements were also carried out on the Al-3Cu-1.5Li model alloy matrix and 99.999% pure Al for comparison. A platinum wire was used as counter electrode, and a saturated calomel electrode (SCE) connected to the microcell through a salt bridge was used as reference electrode. A Gamry potentiostat coupled with a microcell setup was used for polarization tests with a scanning rate of 5 mV/s after an open circuit potential (OCP) period 5

ranging between 0 – 120 min. The potential scanning range was from -0.2 V vs. OCP towards the anodic direction until at least 0.2 V vs. OCP to guarantee a sharp increase of corrosion current density. Each test was repeated at least twice to confirm reproducibility. Separate cathodic polarization was carried out for TB from 0.01 to -0.8 V vs. OCP for the evaluation of galvanic current density at the corrosion potentials of pure Al, Al-3Cu-1.5Li, and AA2070. 2.5. Crystal structure and DFT calculations CrystalMaker was used to build the crystal structures used in this study, and Vienna Ab-initio Simulation Package (VASP) was used for density functional theory (DFT) calculations. The relevant crystallographic information files (CIF) were selected from American Mineralogist Crystal Structure Database and the Springer Materials database. The CIF of TB was based on Xray powder diffraction data as documented in the Springer Materials database [23]. To maintain the Al:Cu:Li stoichiometry ratio of 7.5:4:1, partial occupancy at the (0.25, 0.25, 0.25) site for Al was 0.88 and Li 0.12, and at the (0, 0, 0) site was Cu 0.96 and Al 0.04. Because VASP does not yet have official support for Virtual Crystal Approximation (VCA), we circumvented the partial occupancy by utilizing the supercell approach, which is documented in the software Supercell [24]. The codes can generate combinatorial structures of vacancy or substitution defects in crystalline materials. In the case of TB phase, the total number of generated 1x2x2 supercells is 35960. We selected one of the output structures using the random selection algorithm built into software. The 1x2x2 supercell was then cut into low index surfaces of (001), (110), and (111). Density Functional Theory (DFT) implemented in VASP [25] was used to calculate the (001), (110), and (111) surface work function (WF) of TB phase. Calculations were carried out using the

Projector-Augmented-Wave

(PAW)

pseudopotentials

and

a

generalized

gradient

approximation (GGA) exchange correlation functional [26]. The valence electron eigenfunctions were expanded in a plane-wave basis set with an energy cutoff of 550 eV and a second order Methfessel-Paxton smearing width 0.2 eV was used throughout the calculations. The Monkhorst-Pack k-point was sampled using 1x6x6 mesh size. The convergence criterion for the self-consistent electronic step was 10-4 eV, and the structures were relaxed until the forces acting on the atoms were less than 0.03 eV/A. 3. Results 6

3.1. Nanostructure, chemical and electrochemical properties of precipitates The chemical

composition,

precipitate shape, topography,

crystal

structure,

and

electrochemical polarization behavior of the heat treated (solutionized and aged) Al-3Cu-1.5Li alloy were analyzed using APT, AFM, HAADF-STEM, and the microcell technique, and the data shown are in Fig. 1. Li is enriched at the interfaces of the matrix and precipitate, and the precipitates show a higher Li content than the matrix (Fig. 1a). Clear enrichment of Li is evident in the interfacial region, measured at 10-15 at%, and in the precipitate, about 8 at% (Fig. 1b). The composition ratio of Al, Cu, and Li defines the precipitate as TB phase based on the analyses of different APT needles. Similar in nature to micron-scale IMPs [6], these nm-scale TB precipitates stick out of the surface on the AFM topography map (Fig. 1c). The crystal structure of the TB (100) plane [27] and Al (110) [28] plane is displayed on the magnified STEM images in the insets and precisely overlaps the lattice found in the STEM-HAADF image (Fig. 1d). The inset FFT analysis indicates a coherent correlation between the TB (100) and the Al (110) matrix, or equivalently between TB {100} [29] and Al {110} [30] since TB and Al both have cubic structure. Note that in the crystal structure of TB, Al atoms share the same atomic site with Cu and with Li atoms at the ratio of 0.04 and 0.88, respectively. The potentiodynamic polarization curve of a TB bulk analog in 0.1 M NaCl measured by the electrochemical microcell technique [7] is plotted in Fig. 1e along with those of a solution treated Al-3Cu-1.5Li alloy matrix, AA 2070-T8, and pure Al. The contents of Al, Cu, and Li in the Al3Cu-1.5Li model alloy matrix are a very close representation of these three elements in Al-Cu-Li base alloys, such as AA2050, AA2060, AA2070, AA2099, and 2A97 [6, 13, 31-33]. The corrosion potential (Ecorr) of TB is about 0.25 V and 0.4 V (deviation within 0.03 V) more noble than the Ecorr of Al matrix in Al-Cu-Li alloys and pure Al, respectively. The corrosion current density of TB is around 1x10-5 A/cm2 at its Ecorr. However, TB can support strong cathodic current densities at the corrosion potentials of 2070-T8, Al-3Cu-1.5Li, and pure Al, intersecting the cathodic polarization curve of TB at currents of around -5.3x10-5 A/cm2, -8.8x10-5 A/cm2, and -2.1x10-4 A/cm2, respectively. These values suggest a potentially strong detrimental effect of microgalvanic coupling. Above its Ecorr, TB exhibited no clear evidence of passivity, and the current density rapidly reached 0.1-1 mA/cm2, a signature current density range for the breakdown region of Al alloys [34, 35]. The two alloy materials have passivation regions and the

7

breakdown potentials are around -0.6 V vs. SCE. (For additional information see Fig. S1 and Video S1 in the supplementary information.) 3.2. Quasi in-situ examination of corrosion behavior TB morphology examined by high resolution (HR)-STEM imaging before and after sequential exposure and the separate potentiodynamic curves measured on bulk TB samples are shown in Fig. 2. Because the material is composed of the elements Al, Cu, and Li, the atoms with bright contrast in HADDF images indicate Cu-occupied atomic sites. STEM-HAADF images of TB plates formed in the Al matrix with a TB (100) // Al (110) crystallographic correlation are shown in Fig. 2a1. Some visible material on the edge of the TB precipitate near the hole indicates that the oxide layer covering the TB precipitates is thinner than that on the matrix. TB had no passivation region and was more noble than the Al-3Cu-1.5Li alloy matrix which exhibited a passivation region (Fig. 2a2). After the whole disc sample was exposed to 0.1 M NaCl for 2.5 min, dried and re-imaged under STEM (Fig. 2b1), contrast is visible on the top surface of the TB (100) plane as well as into the depth of TB along certain (010) and (001) planes as a result of the loss of Cu atoms (bright spots in the image and indicated by the red lines in Fig. 2b1). The TB lattice close to the edge of the hole is even sharper and clearly shows much thinner amorphous oxide coverage on TB (highlighted by the dashed red circle in Fig. 2b1) than on the matrix and interfacial region, which can explain the localized corrosion preferentially occurring in TB {001} planes including (100) and (010) than at the interface. No evidence of galvanic interaction was observed at the interface between Al matrix and TB precipitate as little contrast variation appeared in this region. However, galvanic corrosion would be expected at that interface due to the 0.2-0.3 V difference in their Ecorr values (Fig. 2b2). After 30 min exposure to 0.1 M NaCl, more Cu atoms are missing due to further corrosion in the TB precipitate along the TB (001) and (010) planes into the depth. Even in a single plane of either (001) or (010), the HAADF image in Fig. 2c1 exhibits a different contrast, indicating varying corrosion extent in the precipitates. A lattice structure with a different crystallographic orientation appeared to the right of the TB precipitate and on the edge of hole as shown in the green dashed region due to corrosion of TB (Fig. 2c1). The corrosion potential of TB decreased,

8

and TB started to exhibit behavior consistent with passivation after 30 min hold at the OCP (Fig. 2c2). After 90 more min exposure (a total of 120 min) and re-examination with STEM-HAADF imaging, the missing Cu atoms tended to become continuous along the TB (001) and (010) planes (Fig. 2d1). A slight expansion of the bright contrast along the red arrows appeared due to the enrichment of Cu atoms at the interface. However, there was still no evidence of galvanic interaction was observed at the interface. The passivity of TB improved according to the polarization curve (Fig. 2d2). Additionally, the corrosion potential of the matrix increased, approached to that of TB, and the passivation region disappeared as the exposure time increased (from 0 to 120 min), indicating that corrosion not related to galvanic driving force of the precipitate occurred in the matrix. 3.3. Post-corrosion cross-sectioning and 3D analysis A schematic representation of a cross-section thin foil lift out from a traditional TEM disc sample with perforation following the above 2-h immersion tests is shown in Fig. 3a, and the HAADF image of the thin foil bearing TB phase shows the top surface covered by corrosion products and deposited Pt in Fig. 3b. The white dashed line in Fig. 3c indicates the interface of the subsurface matrix and corrosion products. Localized corrosion in TB did occur along the (001) planes in the precipitate, as indicated by the red dotted lines for the loss of Cu atoms (Fig. 3d). It was ruled out that the loss of Cu atoms was the change of foil thickness or stacking faults as they could not produce the missing of Cu atoms in columns similar to the feature of missing Cu atoms in Fig. 2b1-d1 and the rearrangement of these Cu atoms in Fig. 3e. As the top of the precipitate was corroded, Al and Li atoms were dissolved, and pure Cu remained and covered the very top of the precipitate (Fig. 3e). The dashed red lines in Fig. 3e indicate the interface of TB and Cu sites. There is a variation of corrosion in the depth along TB (001) planes, which is supported by the unattacked TB in the dashed hexagonal region and the rearranged pure Cu structure in the dashed rectangle area. (Additional information is in Fig. S2-3) APT analysis was carried out on needles lifted out from the cross section of the corroded alloy matrix bearing TB precipitates after 2 h electrolyte exposure, and the results are shown in Fig. 4. The needles were fabricated starting from the top of the corroded surface into the depth direction of the matrix with the needle pointing towards the surface. Compared to the matrix composition, 9

O, C and Li were enriched in the top layer of corroded structure (Fig. 4a). Directly beneath this layer is the Cu rich layer across the corroded matrix and the TB precipitates as revealed by the Cu atom maps. The isosurfaces with a 90 rotation in Fig. 4b show two crossed TB precipitate plates with a plate along the plane and the other into the plane. According to the isoconcentration surface of Li (Fig. 4b) and the concentration profiles (Fig. 4c), Al and Li were depleted to a certain extent in the top part of the precipitate where Cu was enriched, and Li, O, and C were enriched in the corrosion product layer. The enrichment of C is believed to be from not only beam contamination (if there is any) but also other sources as will be discussed below. Li, C, and O peaks are observed long the concentration profile in Figs. 4b and d which are perpendicular to the precipitate/matrix interface and their concentrations are higher than the uncorroded matrix. Increased Cu concentration was also seen in the both sides of the interface in Fig. 4d. These observations indicate that the Li-enriched interfacial region was corroded to some extent, and a Li-O-C rich layer formed at the interface. This is believed to function as a passivation film at the interface, which hinders the propagation of galvanic corrosion towards the matrix on the sides. Note that longer exposure times may be necessary to further understand the corrosion of the T B precipitate, the role of the existing of Li-O-C layer and its role in corrosion initiation and propagation. (Additional information is in Fig. S4 and Video S2-3.) After a relatively long-term exposure (15 h) of a 3-mm disc sample, preferential aggregation of NaCl on TB precipitates was observed based on surface examination (Fig. 5a and Fig. S5). Foils crossing random precipitates were lifted out from the corroded top surface in Fig. 5a for HR-STEM examination of the cross section. Cu rearrangement and Cu rich sites were confirmed at the top and in the surroundings of the TB precipitates following the orientation correlation of Cu (110)//TB (011) // Al (100) (Figs. 5b-5e). The rearranged atomic structure and the precipitate are different materials in composition and structure. If they were the same type, say TB, the Cu enrichment region located on the top and the sides of TB precipitate would never keep an intact atomic structure without corrosion attack. Alternatively, if they both were Cu crystal, TB would not get corroded. However, the observation is that the Cu enrichment region does have an intact structure, but the TB precipitate beneath the Cu region was corroded as confirmed by the structure and elemental distribution in the top surface examination in Fig. 2 and cross section examination in Figs. 3, 4, 5f and 5g. During the corrosion process, the loss of Cu atom columns initiating along (001) TB planes expanded across the precipitates such that the loss of Cu atom 10

columns is even observed along the (111) TB planes (Fig. 5f). As a result of this, the loss of Cu atoms seems to be nearly random when viewing the precipitate from the [011] TB surface normal. In addition, the EDS map of Cu does show increased Cu counts at the Cu rearrangement regions on the top and the sides of the corroded TB precipitate, further verifying that the Cu rearrangement region and the precipitate represent two different materials. NBD was performed in the Li, C, and O rich corrosion product layer as indicated by the bright spot in Fig. 5d, and the diffraction pattern is given in Fig. 5e. However, precise indexing of the diffraction was not successful. The electrochemical measurements (Fig. 2), the APT analysis, and the cross section analysis, demonstrate that such a layer of corrosion products functioned as a protective barrier against strong galvanic interactions between the relatively noble precipitates and the matrix in the interfacial region, thus preventing corrosion propagation from the inner region of the precipitates towards the matrix. The surface of TB was passivated by the formation of this Li-O-C layer and its corrosion potential neared that of the Al matrix as the exposure time increased (Fig. 2). It is known that the metallic surfaces can be passivated by the precipitation of corrosion products in the process of corrosion [36]. The corrosion product compound Li2CO3, similar to this Li-C-O rich layer with respect to elemental constituents, was indeed reported to form a uniform surface coverage contributing to corrosion resistance [37]. The EDS maps (Fig. 5g) provide additional information regarding the corrosion propagation, corrosion product formation, and NaCl aggregation on the precipitates. NaCl preferentially aggregated and adsorbed on the TB precipitates based on the EDS map of Cl and penetrated into the depth of TB in the subsurface (Figs. 5a and 5g), which is an indication of the role Cl- ions played in the corrosion of the oxide covered surface [38, 39]. The interaction of Cl- with oxidecovered Al has been well documented and generally follows the following process [38, 40, 41]: Cl- adsorption and incorporation into oxide film, oxide thinning through Cl- replacing O2- and/or OH-, and localized damage to the substrate Al metal. As corrosion proceeded, the aggregated clusters of NaCl of around 100 nm scale on top of the TB precipitates provided a source of Clion for the continued localized corrosion underneath the surface. The corrosion attack stopped somewhere in the depth direction (less than 100 nm) as indicated by an unattacked TB precipitate with intact structure and the absence of C, O, Li, and Cu enrichment further beneath the corroded subsurface matrix (Figs. 5b, S5-6).

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4. Discussion A schematic representation of the proposed corrosion mechanism of TB precipitate and the matrix is given in Fig. 6. The oxides in the cross section view represent the oxides covering the top of the precipitates with varying thicknesses. The first issue to be addressed is why the corrosion attack initiated in the precipitates rather than at the interface or matrix. As has already been observed (Figs. 2b1, 2c1, and 2d1), the oxide layer covering the precipitate is thinner than the oxides on the matrix. As such, the as-prepared surface of TB is not as strongly passivated as the matrix and perhaps is not passivated at all. The polarization curve of as-prepared TB surface in Fig. 2a2 showed no strong passivation region, which supports this view and explains why corrosion initiated in the TB precipitate, even though TB possesses a higher Ecorr. Previous research reported that corrosion can initiate in intermetallic compounds due to thin or defective oxide coverage, which supports the preferential corrosion initiation in the TB phase [14, 42]. However, following the depassivation and corrosion initiation in the matrix after sufficient time of exposure e.g. 2 h, it is most likely that matrix corroded faster than TB because the TB surface passivated and matrix did not (Fig. 2d2), which is further supported by the protrusion of TB after 15 h exposure in the cross section view in Fig. 3b and c. After the preexisting oxide film covering the TB precipitate was penetrated, corrosion attack initiated along the {001} planes, including (100) and (010), of the TB precipitate (Figs. 2 and 3). This is a result of the electrochemical reactivity (or nobility) of the atoms in the particular crystallographic planes. Al (with a standard electrode potential of -1.662 V vs. SHE) and Li (3.050 V vs SHE) are electrochemically active due to their negative electrochemical potential relative to Cu (+0.342 V vs. SHE) [43] and were therefore preferentially dissolved. As can be seen in schematics Figs. 6a and 6b, all the atomic sites in each of the {001} TB planes were either occupied by Cu0.96Al0.04 or by Al0.88Li0.12 (Figs. 6a and 6b). For the latter case where Al0.88Li0.12 occupied the {001} TB planes, the atoms in these planes were preferentially corroded. The Cu atoms (or Cu0.96Al0.04) in the rest of the {001} TB planes could not remain in their original positions due to lack of the support from Al and/or Li atoms in the TB lattice structure. However， when the corrosion got close to the matrix/precipitate interface, e.g. the columns of atoms on the outer most side of the TB precipitate (Figs. 2c1 and d1), attack of every two columns along {001} planes was not observed due to the enrichment of Li and the formation of Li-C-O layer, as

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confirmed by APT analysis. Note that the atomic resolution DF STEM images of the precipitates, e.g. Figs. 2, 3, and 5, can only show Cu0.96Al0.04 atomic sites because of the Z (atomic number)contrast. However, this does not mean that corrosion initiated at the Cu0.96Al0.04 atomic sites instead of at the Al0.88Li0.12 atomic sites. As can be seen in Figs. 6c and 6d, the preferential dissolution of Al and Li atoms in {001} TB planes leads to the isolated {001} planes completely composed of Cu0.96Al0.04 atoms. Therefore, these atoms tended to leave their original positions and rearrange in the materials due to the lack of support from Al/Li atoms (Figs. 6c and 6d). The matrix/precipitate interface could physically support the accumulation of Cu atoms due to very limited corrosion attack. The WF has been used to characterize the surfaces of metallic materials in terms of electrochemical reactivity and has been found to closely relate to the OCP of exposed surfaces [44, 45]. Thus, the understanding of corrosion initiation at the TB phase can be enhanced by comparing calculated WF values of planes. The calculated WFs of {001}, {110}, and {111} surfaces in the TB phase are 3.54 eV, 3.95 eV and 3.99 eV, respectively (Fig. 6e). The {001} surface of TB has the lowest WF of all the three surfaces, suggesting that it should have a lower OCP, which correlates with the {001} surface initially undergoing active dissolution prior to that of the other two surfaces. The correlation of WF values to corrosion potential of the nm-scale TB precipitate is based on the notion that relatively low WF values indicate a low energy barrier to trigger electrochemical reactions. The DFT calculations of WF were carried out assuming an ideal surface of TB precipitate in vacuum environment, excluding the variation of oxide film coverage, the interface between precipitate and matrix, and the corrosion attacking directions. Therefore, the WF values do not differentiate between (001) and (010) in {001} family planes. If the preexisting oxide film, Li enriched interface, or the perforation hole close to the edge of TB were considered, slight differences in (001) and (010) might be seen as the corrosion rate would be kinetically different. However, those effects were not included in the DFT calculations of WF. The corrosion initiation along {001} family planes in TB, as confirmed by the WF calculation and atomic arrangement, include the thermodynamically equivalent (001) and (010) planes as labeled in Fig. 2. The {001} planes composed by Al and Li atoms are relatively easy to get corroded, lose electrons, and thus Cu atom columns next to Al/Li atom columns lost support, 13

migrated, and weaken the constrast. One column in every two columns along (001) planes got corroded and lost constrast until close to the interface, which could be treated as a corrosioninduced ‘periodic structure’ in (100) plane along [010] direction. However, it is not expected the two planes (001) and (010) following corrosion show exactly the same level of contrast at the atomic level imaging, due to the variation of sample thickness, the pre-existing oxide film, the defect in the film, and Li enriched interface near both the ends of (010) planes/directions. For example, the region near

the edge of the perforation hole is darker on average than the region further away from the hole because the brightness in the DF image depends on both the Z number and the total amount of materials. Li enriched interface can effectively hinder corrosion at the both ends of (010) planes (or [010] direction) and prevent the formation of a continuous corroded structure as those along (001) planes. Another possible explanation for the stronger contrast along (001) planes than (010) planes, as indicated by the displacement of more Cu atoms from their original positions, could be the differences in corrosion kinetics associated with the directions from which electrolyte attacked the precipitates. Corrosion of (001) planes occurred in three directions close to the hole: top, bottom of the thin sample, and the side close to the hole, whereas (010) planes only got attacked from the top and the bottom of the thin sample. On the other hand, the limited amount of Al atoms, which constituted only 0.04 at% of these atomic sites, might also experience preferential dissolution, further contributing to the rearrangement of Cu atoms. Further, corrosion is a kinetic process and even though other planes have relatively higher WFs, corrosion still occurred in these planes but just at relatively lower rates that, given enough time, can lead to noticeable corrosion attack. Therefore, as the corrosion attack continued, corrosion could transition to other planes, e.g. random rearrangement of Cu atoms and corrosion in (111) TB planes, and eventually consume the entire TB precipitate (Figs. 5e, 5f, and 6d). As a result of this, pure Cu sites and/or Cu rich layers, with orientation relationship of Cu (110) // TB (011) // Al (100), were observed on the top and in the surroundings of the corroded TB precipitate. The Ecorr difference between TB and the matrix is about 0.25 - 0.3 V and TB can support a maximum cathodic current density close to 1x10-4 A/cm2. However, the absence of strong galvanic corrosion observed at the TB and matrix interface must be explained. Corrosion product layer primarily composed of Li-C-O layer has been observed covering the top of the corroded TB precipitates and also distributed along the interface into the depth direction (Fig. 4 and as 14

suggested by the schematic plot in Fig. 6d). As corrosion proceeded on the surface of TB, the surface was passivated due to the formation of an Li-C-O layer, and its corrosion potential neared that of the Al matrix over the increased exposure time (Fig. 2). Such a layer functions as a protective barrier against strong galvanic interactions between the relatively noble precipitates and the matrix in the interfacial region. It is known that the metal/alloy surface can be passivated by the precipitation of corrosion products [36]. The compound Li2CO3 was indeed reported to form a uniform surface coverage contributing to corrosion resistance, as determined by X-ray photoelectron spectroscopy depth profiling [37]. However, the composition of Li in the matrix is not as high as at the interface between precipitate and matrix, which is the major difference between the interfacial region and the matrix: 10-20 at % Li in the interface (and corroded TB surface) vs. 3.0-3.5 at% Li in the matrix. The composition of Li is thus believed to play a major role in passivating the TB surface and interface to form the protective corrosion product layer of Li-C-O. Even though TB contains about 8 at% Li, it did not passivate successfully in 0.1 M NaCl during the early stages of electrolyte exposure but did passivate at later stages with the formation of the Li-rich corrosion product layer. Given this, there is a threshold value of Li composition for successful passivation, which is expected to be above 8 at%. Such a threshold value explains why no strong galvanic corrosion phenomena was observed after short-term exposure between the Cu sites on top of the TB precipitate and the matrix at the interface (Figs. 2d1 and 3e) where Li was above 8 at%. On the other hand, contrast in the matrix near sites of Cu rearrangement started to appear after relatively long period exposure (Fig. 5b, 6c, 6d), indicating Cu sites reached a critical size with sufficient driving force to overcome the protection of Li2CO3 and/or Li-C-O rich layer and initiate galvanic attack in the nearby matrix. The analysis results of the corroded structure and Na/Cl elemental mapping indicate that Clplayed a role in the corrosion initiation, development, and propagation in the precipitates where Cl- accumulated (Fig. 5a and g). It is particularly interesting that Cl- ions penetrated into the depth direction of the precipitates as the corrosion path developed, which is an indication of the role Cl- ions played in the initiation of corrosion of the oxide-covered surface [38, 39]. Clinteracting with oxide-covered Al has been reported [38, 40, 41] and, further based on the current study, generally follows the process: Cl- adsorption and incorporation into oxide film, movement toward metal/alloy-oxide interface, oxide thinning through Cl- replacing O2- and/or OH-, and 15

localized damage to the substrate Al metal. Since Al is a major composition in TB phase, the above mechanism of Cl- interacting with Al applies to both the matrix and the precipitate TB here. On the other hand, relevant work [7, 46] has demonstrated a decrease in corrosion potential and pitting potential with the increased Cl- concentration for intermetallic compounds, which means the intermetallic compounds would become more thermodynamically susceptible to corrosion once Cl- reaches the intermetallic/oxide interface. As corrosion proceeded, the aggregated clusters of NaCl of size of around 100 nm scale on top of the TB precipitate did not serve as a barrier against corrosion. These clusters provided a source of Cl- ion for the continued localized corrosion underneath. As the anion of the strong acid HCl, chloride solutions provide an environment where many metal cations such as Al3+ exhibit considerable solubility [47]. Several other Al-Cu-Li nm precipitates common to Al-Cu-Li alloys, including Al2CuLi (T1), Al6CuLi3 (T2), and Al3Li (), are electrochemically more active than TB precipitates based on the electrochemical polarization measurements (Fig. 7). Even though the breakdown potentials are similar, their OCP values are lower than that of TB and the matrix due to higher Li content. These precipitates and the Al-3Cu-1.5Li matrix all exhibit a passive region up to a breakdown potential of around -0.6 V vs. SCE, and their nobilities from high to low follow the order of TB>Al-3Cu-1.5Li>T2>=T1> (Fig. 7a). If galvanic corrosion of the matrix could occur between any Li-containing precipitate and Al matrix, the TB precipitate is the most likely phase around which such galvanic corrosion of matrix can be observed. Other precipitates are all more active than the matrix and would be less likely to cause galvanic dissolution of matrix. Practically, the galvanic interaction between an active matrix and noble precipitates/particles is more of a concern as during this process the matrix gets corroded and degraded. However, the finding in the present work is that, even though TB is the most noble phase among these precipitates, it still was corroded in the coupling with matrix. Therefore, it is expected that all the other precipitates would also be anodic to the matrix and be preferentially corroded and follow a corrosion path similar to that of TB precipitates. On the other hand, the dominant cathodic reaction is oxygen reduction reaction (ORR) for TB and hydrogen evolution reaction (HER) for the other precipitates (Fig. 7b). Major HER can initiate at the potential value of around -1.2 V vs. SCE, and the cathodic branches of T1, T2, and  fall below that value. 5. Conclusions 16

In summary, this study reported localized corrosion of Li-containing nm-sized hardening precipitates, in particular TB, in Al-Cu-Li alloys based on a quasi in situ approach by sequentially exposing a sample to 0.1 M NaCl solution and characterizing the structural, chemical, and electrochemical evolution. Plate-shaped TB precipitates formed in the Al-3Cu-1.5Li matrix with a crystallographic relationship of TB (001) // Al (110) and TB (011) // Al (100). TB was electrochemically more noble than the Al matrix by 0.2-0.4 V. Localized corrosion initiated along TB {001} planes by the dealloying of Al and Li due to their relatively low WF. This corrosion process was accompanied by the enrichment of Cu in the orientation correlation of Cu (110) // TB (011) // Al (100) on and around corroded TB precipitates. No strong galvanic interactions were observed at the interface between cathodic TB and the anodic Al-3Cu-1.5Li matrix, due to Li enrichment at the interface and in the corrosion product on TB where an Li-C-O rich passivation layer formed during electrolyte exposure. The threshold value of Li composition for successful passivation is above 8 at%. Cl- ions penetrated into the depth direction of the precipitates as the corrosion path developed. The methodology and results of the current work can be extended to assess corrosion associated with other intermetallic phases and other Al alloy systems and can provide inputs for the design of new Al alloys with improved corrosion resistance and high strength. Acknowledgements This work is partially based on research sponsored by Office of Naval Research under agreement #N00014-14-2-0002 through a consortium of LIFT. Dr. Ken Smith at UTRC, Prof. John Allison and Dr. Kai Sun at the University of Michigan are acknowledged for helpful discussions throughout the entire work. APT and part of the STEM imaging were conducted at the Center for Nanophase Materials Sciences (CNMS) at Oak Ridge National Lab, which is a U.S. DOE Office of Science User Facility. DFT calculations were performed on the Ohio Supercomputer Center (OSC). Part of this work were supported by UT-Battelle, LLC under Contract No. DE-AC05-00OR22725 with the U.S. Department of Energy. The Department of Energy will provide public access to these results of federally sponsored research in accordance with the DOE Public Access Plan (http://energy.gov/downloads/doe-public-access-plan). The U.S. Government is authorized to reproduce and distribute reprints for Governmental purposes notwithstanding any copyright notation thereon. Any opinions, findings, and conclusions or

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recommendations expressed in this material are those of the authors and do not necessarily reflect the views of the U.S. Government. Declaration of interests The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

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Fig. 1 Characterization of the TB phase in Alloy Al-3Cu-1.5Li. a, Atom probe tomography maps showing the distribution of Al, Cu, and Li atoms as blue, orange, and pink dots, respectively. b, Proximity histogram concentration profiles of elements based on the top 5 at% Li isoconcentration surface shown in the inset. c, AFM topography map of the TB phase in the matrix. d, HAADF image of the TB in the Al matrix. TB (001) plane and Al (110) plane in the inset STEM-HAADF images overlap with their simulated crystal structure. e, polarization curves of 99.999% pure Al, Al-3Cu-1.5Li matrix, and TB as determined by the microcell technique.

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Fig. 2 Quasi in-situ HR-STEM DF imaging examination of (001) plane of T B phase in (110) plane of Al matrix before and after exposure to 0.1 M NaCl and electrochemical measurements. The exposed disc contained TB precipitates in the electron-transparent region near the perforation hole. a1, as-electro-polished and ion-milled surface; the lattice planes are confirmed by the overlapped crystal structures with the imaged lattice structures. b1, c1, & d1, the same region after the sequential exposure time of 2.5, 30, and 120 min, respectively. a2-d2: micro-cell polarization measurements of the T B and model alloy Al-3Cu-1.5Li on the as-polished surfaces and surfaces of bulk cast ingots after 2.5, 30, and 120 min exposure at OCP, respectively. Dotted blue lines indicate OCP of Al-3Cu1.5Li, OCP of TB, and pitting potential of TB.

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Fig. 3 Thin foil lift out and HAADF images of the cross section of a corroded T B precipitate imbedded in an Al3Cu-1.5Li matrix exposed to 0.1 M NaCl for 2 h. a, schematic representation of thin foil lift out from a corroded traditional 3-mm TEM disc sample. b, HAADF image of thin foil bearing TB phase at low magnification showing top corroded surface covered by deposited Pt. c, the cross section of the TB (011) plane imbedded in the Al-Cu-Li (100) matrix. The whited-dotted curve indicates the interface between the matrix and corrosion products. The white arrow is normal to the top surface and points towards the deposited Pt with bright contrast. d, HR examination of the corroded precipitate and the matrix. Red-dotted lines indicate the missing columns of Cu atoms after electrolyte exposure. e, HR examination of the red-rectangle top region of the corroded TB in a. Red-dotted lines in c indicate the TB precipitate and Cu metal interface. Lattice of Cu viewed from the plane normal [110]. Simulated T B and Cu crystal structure sitting on the imaged T B lattice and Cu lattice in b and c, respectively.

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Fig. 4 Example of APT analysis of a needle lifted out from the cross section of the corroded alloy matrix bearing T B precipitates after 2 h electrolyte exposure in 0.1 M NaCl. The needle sample was fabricated from the very top of the corroded surface into the depth direction of the matrix. a, Atom probe maps of all major ions, including Al, Cu, Li, O, and C. b, Atom maps based on an isosurface of 5 at% Li overlapped with all the other major ions showing two crossed TB precipitates. The map on the left was rotated 90 relative to the map on the right to show the crossing of two TB precipitates. c, One dimensional (1D) concentration profiles of all major ions across the T B precipitates from bottom to the top surface in b. d, 1D concentration profile of all major ions across the corroded matrix surface and TB phase showing Li, O, C, and Cu rich at the precipitate/matrix interface.

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Fig. 5 Surface and cross section analysis of a 3-mm Al-3Cu-1.5Li disc sample exposed to 0.1 M NaCl for 15 h. a, BSE image showing NaCl clusters adsorbed on T B precipitates on corroded sample surface. b, DF and c, BF STEM images of the cross section right beneath the coated Ni covering the surface. The red arrow is normal to the top surface and points towards the Ni coating. The white arrows indicate the contrast between Cu rich sites and matrix around them. Beneath the surface Ni coating is the corrosion product and then the corroded T B precipitates and the matrix. d, f, DF lattice imaging of the corroded structure, corrosion product, Cu sites and matrix of the squared regions in b. e, diffraction pattern of Li-C-O rich layer in the white dotted area in c. g, EDS maps showing elemental distribution of the black squared region in c. Scale bars in g represent 20 nm.

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Fig. 6 Schematic representation of the corrosion processes in specific planes in TB and matrix and the mechanism. a and b, top view and cross section of the as-prepared TB surface respectively in Al matrix based on the imaged precipitates in Figs. 1, and 3 and 5 (uncorroded region). c and d, top view and cross section of the T B surface respectively in Al matrix after electrolyte exposure. e, DFT calculated work function of the (001), (011), and (111) planes, which are equivalent to {001}, {011}, and {111} planes, of TB phase. The arrow in e points towards the WF increase. The oxides in b represent the oxides covering the top of the precipitates with varying thicknesses. The T B precipitate and Al matrix follow orientation relationships of T B (100) // Al (110) in the top view and of T B (011) // Al (100) in the cross-sectional view. Red lines indicate (001) planes of the T B precipitate and black lines represent the TB (111) planes.

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Fig. 7 Polarization evaluation of corrosion properties of common Al-Cu-Li precipitates as comparison to alloy matrix Al-3Cu-1.5Li. a, potentiodynamic polarization scans. b, cathodic polarization scan for T B phase.

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