Abstract
Corrosion degrades metals through complex microscopic processes that traditional methods often fail to capture. In situ transmission electron microscopy (TEM) bridges this gap by providing real-time, atomic- and nanoscale imaging of oxidation and corrosion dynamics. Gas-phase TEM uncovers critical oxidation mechanisms like oxide nucleation and alloying effects, while liquid-phase TEM tracks corrosion initiation. Coupled with emerging techniques like ultrafast direct electron detection cameras, these advancements promise unprecedented resolution in understanding material degradation.
Introduction
Oxidation and corrosion are fundamental phenomena that profoundly impact the performance and longevity of metals and alloys across various applications. These processes can lead to significant degradation of material properties, compromising the structural integrity of essential components in industries such as aerospace, automotive, energy, and infrastructure. Oxide formation due to oxidation can alter mechanical characteristics, including strength and ductility, leading to failure of metallic components exposed to aggressive environments. Economically, the consequences of corrosion are staggering, with costs running into billions of dollars annually due to maintenance, repairs, and lost productivity1. A thorough understanding of the mechanisms behind oxidation enables engineers to select suitable materials and coatings tailored to specific environmental conditions, thus enhancing both durability and reliability. Additionally, with growing emphasis on sustainability and environmental protection, addressing oxidation also plays a vital role in preventing environmental contamination, particularly in systems that transport hazardous materials. Ongoing research into oxidation phenomena drives innovation, leading to the development of new alloys and advanced coatings, such as high-entropy alloys2,3,4,5, which offer improved resistance to these material degradation processes. Therefore, achieving a comprehensive understanding of oxidation is indispensable not only for enhancing existing technologies but also for advancing new applications that demand high-performance materials in challenging environments across a wide range of industries. This knowledge is crucial for optimizing material properties, extending service life, and ensuring reliability in critical applications, ultimately driving progress in fields such as aerospace, energy, and manufacturing.
Despite the seeming simplicity of the reaction between metal and oxygen to form an oxide, oxidation is a complex process that involves multiple hierarchical length and time scales. It typically begins with dissociative chemisorption of molecular oxygen onto the metal surface, followed by the nucleation, growth and coalescence of oxide islands. This sequence ultimately leads to the formation of a continuous, macroscopically thick oxide layer that thickens at a steady rate. When a metal is exposed to an oxygen-containing atmosphere, a series of structural changes occur, including the initial formation of oxygen chemisorbed adlayers, subsurface diffusion of oxygen, and the eventual transformation of metal crystal lattice to a bulk oxide phase. These transformations are influenced not only by the intrinsic properties of the metals—such as crystal orientation, surface structure, defects, and alloy composition—but also by external environmental factors like the type of oxidizing gases, pressure and temperature. To adequately address the fundamental mechanisms underlying the initiation of an oxidation process, it is essential to achieve an atomic-level understanding of surface structures. For decades, surface science techniques have been utilized to investigate surface structures in ultrahigh vacuum (UHV), typically involving careful dosing of small amounts of oxygen. While this approach has yielded valuable insights into the mechanisms of oxygen chemisorption—such as resolving the atomic structures of chemisorbed layers on single crystals under low oxygen coverage6—it often falls short when translating findings to real-world oxidation processes7. The pressure and materials gaps between idealized conditions of surface science experiments and the practical applications of engineering metals and alloys under high temperature and pressure conditions can lead to significant differences in surface structure and composition8,9,10,11,12.
On the other hand, studies of both low and high-temperature oxidation under practically relevant conditions have traditionally focused on the steady-state growth of thermodynamically stable, macroscopically thick oxide layer during the later stages of metal oxidation. A classic method for measuring oxide growth kinetics is thermogravimetric analysis (TGA), which assesses weight gain and oxygen consumption. While TGA can provide valuable insights into the dominant migrating defects—particularly when measurements are conducted as a function of oxygen partial pressure—it does not capture the structural dynamics within the oxide overlayer and at the oxide/metal interface. Consequently, established oxidation theories, such as the Cabrera-Mott13 and Wagner14 theories, often rely on the simplifying assumption of a uniformly growing oxide film. Furthermore, TGA lacks the sufficient accuracy and sensitivity needed to investigate the initial formation of oxides. As a result, the transient early stages of metal oxidation—ranging from the oxide nucleation and initial growth to the establishment of a steady-state oxide scale—remain poorly understood and represent a critical area for further research.
This review highlights the pivotal role of in situ environmental transmission electron microscopy (TEM), encompassing both gas-phase and liquid-phase techniques, in bridging critical knowledge gaps across different stages of metal oxidation and corrosion. We begin by discussing the fundamental principles underpinning the development of environmental TEM for in situ studies of these processes. Subsequently, we present several case studies demonstrating the unique capabilities of environmental TEM in revealing atomic- and nanoscale mechanisms underlying metal oxidation and corrosion, offering insights unattainable through conventional techniques.
In situ environmental transmission electron microscopy
TEM has demonstrated exceptional versatility in studying materials with atomic precision. Recent advancements have further enhanced its capabilities, particularly through the integration of environmental features that facilitate temperature-, pressure-, and time-resolved investigations15,16,17,18. This allows in situ TEM imaging with flowing gases or liquids in the sample area while simultaneously probing atomic structural evolution from the outermost surface layer to the bulk and buried interfaces. Such capabilities provide critical structural and chemical insights into oxidation mechanisms at the nanometer scale and below, using time- and temperature-resolved high-resolution imaging, electron diffraction, and electron energy-loss spectroscopy (EELS) under controlled surface and environment conditions.
In situ studies of surface oxidation and corrosion using environmental (E)-TEM present unique challenges due to the requirements for a vacuum better than 10−6Torr for electron gun operation. To facilitate oxidation experiments under gaseous environments within TEM, reactive gases must be confined to the sample region while preserving the high vacuum conditions required for both the column and gun. Two primary designs have been developed for gas confinement: (1) thin window cells (Fig. 1a) and (2) differential pumping systems (Fig. 1b). The thin window design involves constructing a sealed sample holder that isolates the gas cell from the column vacuum, allowing gases to oxidize the sample positioned between two windows. These window materials must be electron-transparent and sufficiently robust to withstand the surrounding vacuum. While it is theoretically possible to achieve gas pressures up to one atmosphere without compromising the electron gun, using thicker window materials to support such pressures can adversely affect the instrument performance. Additionally, the presence of window materials can introduce extra structural information into the imaging and diffraction patterns of the sample. Consequently, materials with weak electron scattering properties, such as amorphous carbon and SiNx thin films, are typically preferred for the windows. The second design for gas confinement utilizes multi-stage small apertures placed above and below the samples, enabling differential pumping through these apertures within the objective pole-piece. This configuration allows gases to be directly introduced into the sample region while maintaining a significantly lower pressure near the electron gun. In this differential pumping system, the gas pressure achievable in the sample region is influenced by both the leak rate through the apertures and the pumping speed in their vicinity. This interplay determines how effectively the microscope can maintain the desired gas pressure while ensuring the vacuum conditions necessary for the electron gun operation.
Fig. 1: In situ gas- and liquid-phase TEM setup.
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a Thin-window cell specimen holder. b Differentially pumped TEM system.
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Both designs—thin-window environmental cells and differential pumping systems—offer distinct advantages and disadvantages. Thin-window cells, while convenient, can reduce image contrast, resolution, and analytical capabilities such as energy-dispersive X-ray spectroscopy (EDS) and electron energy-loss spectroscopy (EELS) due to the combined thickness of the windows and sample. Traditionally, these cells were also limited by their inability to withstand significant heating, making them unsuitable for high-temperature oxidation experiments. However, recent advancements, including the integration of microelectromechanical system (MEMS)-based heaters, have expanded their capabilities, allowing thin-window cells to operate at temperatures ranging from 800 to 1000 °C19,20. Additionally, because the gas is confined within the sample holder, in situ oxidation experiments using thin-window cells can be conducted with commercial or custom-built holders, requiring no modifications to the TEM and supporting gas pressures of up to 1 atmosphere21. In contrast, differential pumping systems require reconfiguring the TEM and typically limit gas pressure in the sample region to around 20 Torr to avoid adverse effects on the electron gun and TEM column. Despite this limitation, differential pumping offers several advantages, including superior resolution, the ability to heat samples to higher temperatures, and greater sample tilting flexibility. This design also supports analytical tools such as EELS, enabling simultaneous acquisition of structural and chemical information during oxidation.
The thin-window design (Fig. 1a) also enables encapsulation of a thin liquid layer between two electron-transparent membranes, with variations in the thickness and type of materials for viewing windows, spacer, and metal contacts connected to external electrical terminals22,23,24,25,26,27,28. This configuration allows real-time observation of electrochemical reactions, such as metal dissolution, passivation, and oxide formation, under applied potentials or in the presence of corrosive liquids. However, liquid-cell TEM faces significant challenges. Maintaining a stable and thin liquid layer (<1 µm) is essential to minimize electron scattering and preserve imaging resolution, while preventing beam-induced artifacts like electrolyte radiolysis or unintended reactions that can obscure or alter natural corrosion processes. Additionally, ensuring the chemical compatibility of window materials with corrosive environments is critical to avoid unwanted reactions with the liquid or metal samples. Effective thermal management and control of beam-induced heating are also necessary to prevent the unintentional acceleration of corrosion dynamics.
To achieve atomic-level imaging resolution, graphene liquid cells have been introduced, wherein liquid is encapsulated between graphene layers29,30. Compared to microfabricated liquid cells with silicon nitride windows typically tens of nanometers thick, the single-atom thickness and low atomic number of graphene enable reduced electron scattering and improved imaging resolution. However, graphene liquid cells present unique limitations. The liquid solution remains stagnant without mass flow between the graphene layers. Furthermore, functions available in microfabricated cells, such as biasing via integrated electrodes, are challenging to implement in graphene-based systems, restricting their versatility for certain electrochemical experiments.
The advancements in aberration correction for electron lenses, coupled with the integration of the environmental cell designs, significantly enhance dynamic in situ oxidation and corrosion experiments. Aberration-corrected environmental TEM offers the ability to capture a full range of spatial frequencies without the spherical aberration (Cs) issues that can cause image contrast reversals and minima in the contrast transfer function (CTF) typical of uncorrected TEM31,32. This correction also minimizes electron delocalization effects in TEM images, which are particularly critical at discontinuities—such as structural defects and surfaces—that critically influence material behavior and chemical activity during oxidation processes. Moreover, the development of the aberration-corrected environmental scanning TEM (STEM) further extends the capabilities for in situ atomic resolution studies33,34,35. It incorporates full analytical tools like energy-dispersive X-ray spectroscopy (EDS) and EELS, along with unrestricted high-angle annular dark field (HAADF) imaging. This allows for detailed observations of dynamic changes in surface, interface and bulk of materials during in situ experiments, while continuously flowing gases at elevated temperatures36.
Oxidation studies using gas-phase TEM
E-TEM enables direct visualization of oxygen-adsorption-induced reconstructions on metal surfaces, kinetics of oxide nucleation and growth, and morphological evolution of oxide islands in gas environments37,38,39,40. This approach yields critical insights into active sites, transient states, and the kinetics and thermodynamics of oxidation processes at local regions38,39,40,41,42,43,44,45,46,47,48,49,50,51,52,53,54,55, which are often difficult to capture using conventional ensemble techniques. The early stages of oxidation present a particularly complex interplay of processes, such as oxide island nucleation and the growth of individual islands. Conventional ensemble-averaging methods typically measure overall reaction kinetics, making it challenging to distinguish between these processes. In contrast, E-TEM is uniquely suited to address this limitation by simultaneously tracking oxide island nucleation events and the growth rates of individual islands in real time39,40,47,52,53,54,55,56,57,58,59,60. This ability to capture structural evolution at the atomic and nanoscale under reaction conditions is crucial for understanding early-stage oxidation dynamics, effectively bridging the gap between surface science experiments and bulk material characterization. Due to its unique capabilities, E-TEM has gained significant interest for studying gas-surface reactions. Table 1 summarizes recent oxidation studies conducted via gas-phase TEM on metals and alloys.
Table 1 Representative oxidation studies using gas-phase TEM
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Transition from oxygen-induced surface restructuring to monolayer oxide formation
Surface oxidation typically progresses through a series of stages, beginning with oxygen adsorption-induced surface restructuring, followed by oxide nucleation, growth, and ultimately the development of a continuous oxide layer61,62,63,64. While considerable research has focused on characterizing the atomic structure of oxygen-induced surface reconstructions under ultrahigh vacuum conditions6,65,66,67, directly observing the atomic processes that drive the transformation from a two-dimensional oxygen-chemisorbed layer to a continuous oxide layer during cumulative oxygen uptake remains a challenge. This difficulty largely stems from experimental limitations in monitoring the spatiotemporal evolution of local atomic structures under reactive O2 gas and high-temperature conditions. However, the advent of environmental TEM has opened new avenues for bridging this knowledge gap, allowing for atomic-scale imaging of reaction dynamics at gas-solid interfaces.
Figure 2 illustrates in situ high-resolution (HR) TEM observations of Cu(110) oxidation. A monoatomic Cu2O step divides two terraces: the lower right terrace displays a well-developed c(6×2)-O reconstruction, induced by oxygen chemisorption on pristine Cu(110), while the upper left terrace features a single-atom-thick Cu2O overlayer (highlighted in yellow). Scanning tunneling microscopy (STM) imaging of O₂-exposed Cu(110)65,66 and HRTEM simulations68 further confirm the c(6×2)-O structure. In situ HRTEM imaging (Fig. 2a–e) shows that Cu2O grows via lateral propagation of the Cu2O step toward the c(6×2)-O terrace, proceeding in a stop-and-go manner. The step movement intermittently pauses for 1 to 4 s, as depicted in the distance-time plot (Fig. 2f). A zoomed HRTEM image (inset, Fig. 2a) overlaid with the atomic structure of the Cu2O growth front reveals the mechanism: the Cu2O step temporarily halts at isolated Cu atom columns within the c(6×2)-O reconstruction before advancing to the next isolated Cu column. During each step, half the unit cell length of the c(6×2)-O structure is transformed into a monoatomic Cu2O layer. This process of periodic pauses and half-unit-cell propagation results in the stop-and-go migration of the Cu2O growth front toward the c(6×2)-O region, providing insights into the atomic-scale dynamics of Cu2O growth.
Fig. 2: Intermittent transformation of c(6×2)-O to Cu2O on Cu(110).
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a–e Time-sequence HRTEM images illustrating stop-and-go step growth of a monoatomic Cu2O layer at 623 K and 1 × 10−3 Pa O2. The Cu2O growth front pauses intermittently (1 to 4 s) at isolated Cu columns within the c(6×2)-O terrace. Inset in (a): zoomed-in HRTEM image with the overlaid atomic structure of the Cu2O growth front. The yellow dashed circle represents a Cu vacancy beneath an isolated Cu column in the c(6×2)-O reconstruction. Cyan, yellow, gray, and red balls represent substrate Cu, Cu in Cu2O, Cu in c(6×2)-O, and O, respectively. Inset in (e): image intensity profile along the yellow dashed line, showing the periodic undulations of the Cu2O layer. Pseudo color is applied to the Cu2O layer for clarity. f Distance–time plot from in situ HRTEM imaging, showing the stop-and-go movement of the Cu2O growth front. Adapted from68 Copyright @2024 American Physical Society.
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The resulting Cu2O overlayer exhibits periodic hill-and-valley morphology. The intensity profile along the interfacial Cu layer (yellow dashed line in Fig. 2e) shows weaker intensity in valley regions, attributed to the self-assembly of atomic vacancies. A density functional theory (DFT) optimized structure model (inset, Fig. 2a) reveals that isolated Cu atoms in the c(6×2)-O reconstruction bond with oxygen in the Cu2O overlayer. This bonding causes the upward relaxation of isolated Cu atoms and the formation of Cu vacancies beneath the Cu2O growth front (inset in Fig. 2a). The transformation from c(6×2)-O to Cu2O involves oxygen adsorption at the growth front, prompting nearby Cu atoms to migrate toward the created vacancies. These vacancies are then injected into the Cu2O/Cu interface. DFT analysis indicates that these injected vacancies preferentially self-assemble into nanocavities at the Cu2O/Cu interface, causing local downward relaxation of Cu and O atoms in the Cu2O overlayer68. This interfacial self-assembly of vacancies drives the hill-and-valley undulations observed in the Cu2O overlayer.
In situ TEM imaging comparing the oxidation of Cu(110) and Cu85Au15(110)68 reveals that injected vacancies can be directed to deeper regions within the alloy. This reduces metal/oxide interfacial cavitation and improves oxide layer adhesion. The established oxide breakdown model69 describes the evolution of the oxide/metal interface through stages of wrinkling, voiding, buckling, cracking, and eventual spalling. In situ TEM provides atomistic insights into the origins of interfacial wrinkling and voiding caused by vacancy injection and aggregation. Additionally, it demonstrates the potential to enhance oxide scale stability by alloying to manipulating the fate of atomic vacancies injected during oxide growth.
Grain boundary-driven unidirectional oxide growth
The spontaneous formation of oxide whiskers during metal oxidation is a well-documented phenomenon that remains incompletely understood. Oxide whiskering typically occurs during the later stages of oxidation, alongside the growth of a thick oxide scale. Studies suggest that whisker formation arises from the accumulation and relaxation of compressive stresses generated by interfacial reactions between different oxide phases in multilayer scales. While whiskering is observed across various metal systems, Cu oxidation is a well-known example70,71,72,73,74,75,76,77. As shown in Fig. 3a, oxidized Cu exhibits a three-layer structure: a 3-µm-thick Cu₂O layer above the Cu substrate, an intermediate CuO layer (600 nm thick), and a top layer of CuO whiskers. The formation of CuO whiskers is attributed to high compressive stresses within the oxide layer, driven by the significant molar volume differences between Cu and its oxides71,78,79. Both Cu₂O and CuO are cation-deficient p-type oxides, with growth primarily governed by the outward diffusion of cations71,80,81,82,83,84,85. New oxide forms at the gas/oxide interface, enabling growth into free space. This unconstrained surface growth, facilitated by outward Cu cation diffusion, typically prevents stress development at the gas/oxide interface.
Fig. 3: Unidirectional CuO growth during Cu oxidation.
figure 3
a Cross-sectional SEM image illustrating the layered morphology of oxidized Cu at 673 K and ~2 × 104Pa O2 for 2 h. b TEM image of a CuO whisker with a bicrystal grain boundary along its axis. c HRTEM image revealing the atomic structure of the bicrystal grain boundary. d–l In situ HRTEM imaging of CuO tip growth at 673 K under 0.5 Pa O2. d Step-terrace configuration at the CuO whisker tip with the bicrystal boundary marked by green dashed lines. e Schematic illustrating Cu ion transport along the grain boundary to the tip. f Step-flow growth of atomic planes on the left terrace, starting from the bicrystal boundary corner, with the green dashed line indicating the initial configuration from (d). g Locally accelerated oxide growth (protrusion) at the grain boundary, marked by yellow arrows. h–l Repeated nucleation and step-flow growth of CuO atomic layers on the right terrace of the whisker tip. Adapted from91 Copyright @2019 Wiley.
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An interfacial stress mechanism has been proposed to explain oxide whiskering during metal oxidation86,87. For Cu, the formation of CuO whiskers is linked to reactions at the CuO/Cu₂O interface, where Cu₂O grows at the expense of CuO through a continuous oxide-forming reaction. The differing molar volumes of Cu₂O (77.83 ų) and CuO (81.16 ų) create compressive stresses at the interface. These stresses drive the outward diffusion of Cu cations along grain boundaries in the CuO layer, facilitating whisker formation.
However, the mechanism driving unidirectional oxide whisker growth during metal oxidation has been a longstanding topic of debate. The anisotropic shape of whiskers arises from differential axial and radial growth. Axial growth occurs at the whisker tip, while radial growth happens along its sidewalls. Axial growth requires metal ions to be transported from the whisker base to its tip, either through surface diffusion along the whisker’s exterior or via internal pathways. However, surface diffusion is unlikely to drive axial elongation, as metal ions on the whisker surface are more prone to reacting with the abundant oxygen in the surrounding environment, leading to new oxide formation at the base and promoting radial growth instead. Given that whiskering occurs at intermediate temperatures (e.g., below 500 °C for Cu), internal diffusion must differ from conventional lattice diffusion.
TEM characterization reveals that oxide whiskers typically exhibit a bicrystal boundary along their axial direction86,87,88,89,90, which provides an alternative diffusion pathway for metal ions. Recent in situ TEM imaging has confirmed that these bicrystal boundaries guide unidirectional CuO whisker growth during Cu oxidation91. As shown in Fig. 3b, c, the bicrystal boundary extends along the whisker’s axis. Time-sequence HRTEM images (Fig. 3d–l) detail the growth process, showing the nucleation and lateral propagation of atomic layers at the whisker tip. At the tip, the bicrystal boundary forms a terrace-step structure, with one grain elevated by five atomic layers (Fig. 3d). Faster oxide formation occurs at the grain boundary emergence, creating a small protrusion (highlighted by yellow arrows in Fig. 3g). This localized oxide growth indicates significant Cu ion flux through the bicrystal boundary, where diffusing Cu ions react with adsorbed oxygen, forming oxide islands at the tip.
The oxidation of metals and alloys, such as Cu86,92, Fe93,94,95,96, Ni97, Zn98,99,100, Mo101,102, and Cu-Zn103, leads to the formation of diverse oxide nanostructures, including nanowires, nanblades, nanobelts, and nanoplates. The common presence of bicrystal boundaries in these oxidation-induced structures underscore the importance of the bicrystal grain boundary diffusion mechanism. This mechanism plays a unique role in shaping growth morphologies by influencing the boundary structure, enabling either one-dimensional crystal growth through anisotropic diffusion or two-dimensional growth via isotropic diffusion along the grain boundary plane. Variation in oxide growth morphologies—1D nanowires86[,87[,104 versus 2D nanobelts, nanobelts, and nanoplatelets93,94,96—can be traced to the atomic structure of the bicrystal grain boundary. In the case of 1D whisker growth, the bicrystal grain boundaries exhibit an anisotropic structure, allowing for highly unidirectional diffusion. Conversely, an isotropic bicrystal grain boundary structure promotes isotropic diffusion across the bicrystal plane, resulting wider, thinner platelet growth.
Atomic insights into CO2-promoted oxidation of chromia-forming alloys
CO2, a common byproduct of petrochemical fuel combustion, poses significant corrosion risks to critical alloy components in energy and power systems. For instance, structural alloys in next-generation power cycles using supercritical CO2 as the working fluid, offering improved conversion efficiency over traditional steam systems, are vulnerable to high-temperature corrosion105,106,107. Similarly, CO2 gas-cooled nuclear reactors depend on various steels as construction materials, and future oxy-combustion power plants, designed for efficient CO2 capture, will require alloys resistant to flue gases with high CO2 concentrations. In these applications, CO2 is highly corrosive, causing substantial oxidation of steel components. Chromia-forming steels, which perform well in air, experience breakaway oxidation in CO2 environments63,108,109. This indicates that CO2 is significantly more corrosive to chromia-forming alloys than O2 or air110,111,112,113. The accelerated corrosion in CO2 is linked to rapid chromia growth, which generates stresses within the chromia scale, causing local failures and depleting Cr in the subsurface zone needed for sustained Cr2O3 growth110,112,114.
Current understanding of CO₂-induced oxidation is based on macroscopic or mesoscopic measurements of weight or thickness changes during steady-state oxide growth. However, these methods lack the resolution to uncover the atomic-scale mechanisms driving oxide formation. A detailed microscopic understanding is essential for selecting and designing high-temperature alloys for future energy systems. NiCr, a key chromia-forming alloy, plays a critical role in high-temperature applications, and gaining insights into its oxidation mechanisms is crucial for managing oxide scale growth and preventing alloy degradation.
Environmental TEM offers an ideal platform for observing Cr2O3 growth in CO2 atmospheres48. Fig. 4a shows in situ bright-field (BF) TEM images of the initial-stage oxidation of Ni–25 wt%Cr in CO2, revealing the nucleation and growth of a Cr2O3 island. A similar pattern is observed in O2 under the same conditions (Fig. 4b). Tracking the Cr2O3 island coverage allows for direct determination of the oxidation rate. Fig. 4c shows the area evolution of Cr2O3 islands, with a growth rate of ~215 nm²/s in CO2, much faster than the ~17 nm²/s in O2. Fig. 4d presents real-time HRTEM images of dynamic void formation and migration in Cr2O3 during NiCr oxidation in CO2, with vacancy clusters exhibiting migration, aggregation, and disaggregation. In contrast, NiCr oxidation in O2 (Fig. 4e) results in a more compact Cr2O3 structure without noticeable vacancy clusters. HAADF-STEM imaging confirms a highly porous Cr2O3 structure from oxidation in CO2, compared to the compact structure in O2 (Fig. 4f).
Fig. 4: In situ TEM observations of Cr2O3 growth during NiCr oxidation at 773 K and 5 Pa CO2 and O2.
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a Time-elapse TEM images of Cr2O3 island growth in CO2. b Time-sequence TEM images of Cr2O3 island growth in O2. c Plots of oxide island area versus oxidation time, revealing faster oxidation in CO2 than in O2 gas. d In situ HRTEM images of vacancy cluster dynamics during Cr2O3 growth in CO2. e Time-sequence HRTEM images of Cr2O3 growth in O2. f, g HAADF images of Cr2O3 formed in CO2 and O2, respectively. Adapted from48 Copyright @2024 Elsevier.
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The issue of carbon uptake during oxidation in high-temperature CO2 is debated63,115,116. In situ EELS analysis detects C in the Cr2O3 layer48, suggesting CO2 decomposes into C and O on the Cr2O3 surface, with C atoms dissolving into the Cr2O3 lattice. In situ TEM shows that oxidation in CO2 leads to faster oxide growth and a more porous structure, with O acting as an oxidizer and C infiltrating the lattice. DFT modeling provides insights into how dissolved C promotes vacancy formation and migration, resulting in a defective oxide structure that enhances through-lattice diffusion, driving Cr2O3 growth in CO248.
The oxidation atmosphere significantly influences the generation and dynamics of atomic defects in the oxide scale. Complete CO2 decomposition leads to C atom dissolution into the oxide lattice, increasing atomic vacancy concentration and resulting in a defective oxide layer that weakens its self-protection properties. Experimentally revealing these processes is challenging but crucial for controlling oxidation behavior. Steering CO2 dissociation toward partial decomposition (CO2 → CO + O) instead of complete decomposition (CO2 → C + 2O) could reduce C dissolution into the Cr2O3 lattice, enhancing the oxide scale’s self-protection. This could be achieved by designing multicomponent alloys that favor partial CO2 decomposition and prevent C penetration into the oxide.
Atomistic mechanisms behind surface passivation of aluminum
Surface passivation is crucial for various processes, including corrosion resistance, catalysis, electrochemistry, and electronic device fabrication. One widely referenced model for low-temperature passivating oxide film formation is the Cabrera-Mott model13, which describes rapid oxide growth followed by slower growth to a self-limiting thickness117,118,119,120,121,122,123,124,125. This behavior is attributed to a self-generated electric field that enhances ionic transport during the initial phase, but this effect diminishes as the oxide thickens. Despite its widespread use, the atomic-level mechanisms behind this self-limiting growth remain unresolved, primarily due to challenges in directly investigating atomic exchanges between the surface and subsurface regions. Surface defects further complicate this investigation, making it difficult to probe structural evolution at the atomic scale. ETEM provides a solution by offering atomic-scale dynamic insights into both surface and subsurface regions under controlled atmospheres, making it ideal for studying passivation processes, especially with materials like aluminum (Al), which is widely used for its corrosion resistance.
In situ TEM observations of Al surfaces in O2 show that the passivation process occurs in two stages41. The first stage is characterized by intralayer disordering, while the second stage is dominated by interlayer disordering, leading to the complete amorphization of the Al2O3 layer. In contrast to the formation of an amorphous Al2O3 film in O2, in situ HRTEM imaging of Al surfaces exposed to H2O vapor reveals the development of a bilayer film consisting of an upper Al(OH)3 layer and inner layer of amorphous Al2O342. Fig. 5a–g present in situ HRTEM images capturing the dynamic transformation of the Al lattice into Al oxides as clean Al(111) is exposed to H2O vapor. Following H2O exposure, certain areas of the outermost layer exhibit weakened image contrast (marked by solid white lines in Fig. 5b, c), which is attributed to the extraction of Al atoms from the topmost layer due to the significant exothermic reaction associated with dissociative adsorption of H2O molecules. This process leads to the formation of atomic vacancies in the topmost layer, as shown by the simulated HRTEM image (inset in Fig. 5b). The H2O adsorption-induced extraction of surface atoms aligns with the results from surface science experiments and atomistic simulations, which indicate that the dissociative adsorption of gas molecules on metallic surfaces typically results in one-atomic-layer-deep pits and adatoms removed from the metal surface65,126,127. The exfoliation of the outermost Al results in the hydroxylation of the two inner atomic layers, causing the interplanar spacing to expand from 2.3 Å in the pristine Al(111) lattice (Fig. 5a) to ~ 3.2 Å in the region marked by the two small triangles in Fig. 5d. This increased spacing closely corresponds to the interplanar spacing of Al(OH)3(112)128,129.
Fig. 5: H2O-vapor induced passivation of Al surfaces.
figure 5
a–g Time-sequence HRTEM images showing the growth of the Al(OH)3/Al2O3 bilayer film on Al(111) at 298 K and ~5 × 10−3Pa H2O. The solid white lines indicate weakened lattice contrast due to Al atom extraction from the surface during H2O adsorption. In (g), the white line overlays the Al2O3/Al(111) interface profile at t = 65.4 s from (d). Insets in (b, g) show simulated HRTEM images of the Al lattice with vacancies (b) and the Al(OH)3 structure (g). h Schematic of the Al(OH)3/Al2O3 bilayer structure. i Time dependence of Al(OH)3/Al2O3 bilayer thickness (green) and displacement of the Al2O3/Al(111) interface (black). j–o Time-sequence HRTEM images of Al(OH)3/Al2O3 bilayer growth on a stepped Al surface at 298 K and ~5 × 10−5 mbar H2O. The white dotted rectangles in (j–l) highlight the disordering of the stepped facet before Al(OH)3 formation. Dashed yellow and solid blue lines in (d–g, m–o) indicate the Al(OH)3/Al2O3 and Al2O3/Al(111) interfaces, respectively, with pseudo yellow and blue colors representing the Al(OH)3 and Al2O3 layers. Adapted from42 Copyright @2023 The Authors. Published by AAAS.
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Continued H2O exposure transforms the crystalline Al lattice in the subsurface into an amorphous Al2O3-like oxide (Fig. 5d–g), while the top Al(OH)3 layer remains crystalline, as shown by the simulated HRTEM image in Fig. 5g. These observations confirm the growth of an Al(OH)3/Al2O3 bilayer, supported by ambient-pressure X-ray photoelectron spectroscopy measurements42. The Al2O3/Al(111) interface is atomically rough and moves inward, while the Al(OH)3 layer maintains a constant thickness of ~5.0 Å. The Al(OH)3/Al2O3 bilayer configuration is shown in Fig. 5h, with thickness evolution in Fig. 5i, where rapid initial growth is followed by slower thickening, consistent with the Cabrera–Mott model’s self-limiting growth behavior. A similar bilayer oxide film was also observed when the Al surface was immersed in liquid H2O42.
The formation of the Al(OH)3/Al2O3 bilayer film is observed on stepped Al surfaces. Figure 5j–o shows HRTEM images of hydroxylation at the intersection of a flat (100) and a stepped facet. H2O adsorption abstracts Al atoms from the (100) facet, causing contrast loss (Fig. 5j, k) and hydroxylating the top two atomic layers, expanding the lattice spacing from 2.0 Å to 2.9 Å. In contrast, the stepped facet undergoes disordering and transitions to crystalline Al(OH)3 (Fig. 5l). Prolonged exposure forms and thickens amorphous Al2O3 beneath the Al(OH)3 layer (Fig. 5m–o). The Al(OH)3 layer remains ~5.0 Å thick, while the Al2O3 layer grows to a limiting thickness. This bilayer oxide formation is consistent with the behavior observed when an amorphous Al2O3 overlayer is exposed to H2O vapor, confirming the intrinsic formation of a crystalline Al(OH)3 overlayer through H2O-induced surface hydroxylation42.
The atomic origins of the Al(OH)3/Al2O3 bilayer formation in H2O are elucidated through DFT modeling42. H2O molecules dissociate on pristine Al(111) into OH- and H+, hydrogenating the surface and forming an Al(OH)3 overlayer. DFT calculations show that a perfect, stoichiometric Al(OH)3 surface does not react with dissociative H2O adsorption. Based on in situ TEM observations, Al2O3 growth at the Al2O3/Al interface requires O ion transport from the Al(OH)3 layer, creating O vacancies. These vacancies promote H2O dissociation, providing atomic O that diffuses inward for Al2O3 growth and producing H2, as confirmed by residual gas analyzer measurements during Al(111) exposure to H2O vapor42.
The widespread use of Al is due to its ability to form a passivation layer, typically thought to be amorphous. However, in situ TEM imaging reveals that the passivating layer in H2O consists of a crystalline Al(OH)3 outer layer and an amorphous Al2O3 inner layer. This Al(OH)3/Al2O3 bilayer differs in microstructure from the single amorphous Al2O3 layer formed in dry O2 and alters passivation mechanisms. In H2O, the growth mechanism shifts from a combination of inward O diffusion and outward Al diffusion in O241 to predominantly inward O diffusion, forming O vacancies and excess H in the Al(OH)3 overlayer. This promotes H2O dissociation and H2 production, with implications for clean-energy applications, such as the use of Al in vehicular H2 storage130,131,132. The findings also provide insights into surface passivation dynamics for other metals in humid conditions.
Corrosion studies using liquid-phase TEM
Liquid-phase TEM has provided exciting but challenging opportunities to enhance the mechanistic understanding of corrosion phenomena by enabling real-time, nanoscale observation of dynamic processes in aqueous environments133,134,135,136,137. Studies have revealed key phenomena such as the initiation and propagation of pitting corrosion, the formation and breakdown of passive oxide layers, and the role of localized electrochemical reactions in driving metal dissolution135. Liquid-cell TEM revealed localized pitting behavior of Al thin films in saltwater environments, showing that pitting events could be directly observed and correlated with ion-irradiated regions of Au138,139. Liquid-cell TEM also allows for monitoring the formation of nanoscale corrosion products, such oxides, and their impact on corrosion propagation140. These findings have deepened our understanding of the atomic and nanoscale processes governing corrosion, offering new perspectives for designing corrosion-resistant materials. Table 2 summarizes recent corrosion studies conducted via liquid-phase TEM on metals, alloys, and oxides.
Table 2 Representative corrosion studies using liquid-phase TEM
Full size table
Dealloying-driven local corrosion: from surface initiation to in-depth propagation
A representative example is the study of the de-alloy process in intermetallic phases within aluminum alloys exposed to aqueous environments. Using liquid-cell HAADF-STEM imaging, Kosari directly observed localized corrosion initiated by the dealloying of intermetallic Al2CuMg particles141. Fig. 6a illustrates the time-resolved morphological evolution (top view) of an Al2CuMg particle in 0.01 M NaCl solution. At 0 min, a bright Al2CuMg particle with two distinct grains is visible within the alloy matrix, separated by a grain boundary (indicated by a white arrow). By 5 min, the particle’s center darkens, signaling the onset of material removal due to corrosion. At 8 min, the corroded region expands radially, forming a porous structure, while partial dissolution of the adjacent matrix begins (white arrow). At 11 min, corrosion extends further, with granular contrast variations appearing throughout the particle. By 16 min, trench-like dissolution around the particle becomes apparent, accompanied by the corrosion of a nanosized Al2CuMg dispersoid, leaving behind a nano-pit. At 19 min, trenching completely disconnects the particle from the matrix. By 21 min, bright Cu clusters form along the matrix edges and on dispersoids (red arrow), while the particle size reduces due to Cu dissolution. Finally, at 28 min, some Cu clusters near the trench rim disappear, making the advanced stages of the dealloying process.
Fig. 6: In situ HAADF-STEM imaging of localized corrosion initiated at intermetallic Al2CuMg particles in 0.01 M NaCl solution (pH 6.5).
figure 6
a Time-resolved top-view images illustrating the morphological evolution of an Al2CuMg particle and the surrounding Al matrix during corrosion. b Time-resolved cross-sectional images revealing the role of a deep Al2CuMg particle in driving corrosion propagation. Adapted from141 Copyright @2020 The Authors. Published by Elsevier Ltd.
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Figure 6b presents cross-sectional HAADF-STEM images illustrating the in situ corrosion progression of an Al2CuMg particle within the alloy matrix exposed to 0.01 M NaCl solution. At 20 min, distinct contrast variations appear in the upper portion of the particle, which is directly exposed to the electrolyte. Corrosion initiates at the surface, dissolving the surrounding alloy matrix and propagating deeper along the particle/matrix interface, leading to extensive matrix degradation. By 25 min, corrosion advances further into both the matrix and interfacial regions of the particle, although the middle portion of the particle remains less affected. The trench surrounding the particle widens and deepens as the process continues. At 31 min, the particle becomes fully disconnected from the alloy matrix. Rapid dealloying at this stage leads to a breakdown crack near the top of the particle, highlighted by the yellow arrow. These in situ STEM observations, combining both top- and cross-sectional views, provide a comprehensive visualization of the corrosion process, from its initiation at the surface to its deep propagation. The findings underscore the critical role of dealloying in intermetallic particles as the primary driver of localized degradation in aluminum alloys.
Atomic-scale visualization of corrosion in Sn nanocrystals
Liquid-cell TEM is often limited by poor resolution, primarily due to electron scattering from the SiN window and the liquid layer. While liquid layers up to several micrometer thick can be imaged, the resolution is significantly compromised142. One effective way to mitigate electron scattering from the window is to use alternative membrane materials, such as graphene29,30. Recent studies have shown that using amorphous carbon film as the membrane material can also achieve atomic resolution141. Fig. 7a illustrates the corrosion behavior of a Sn nanocrystal in a relatively thick aqueous solution of salt chlorides (e.g., SnCl4, NiCl2, and BeCl2). The entire particle dissolved within seconds, exhibiting a smooth and rounded solid/liquid interface during the etching process. As the primary Sn nanocrystal dissolved, small Sn nanoparticles formed in the vicinity, indicating a rapid increase in Sn ion concentration, reaching an oversaturation level due to the dissolution of the primary nanocrystal. In contrast, distinct corrosion behavior was observed for a Sn nanocrystal in a thin liquid layer (tens of nanometers). The Sn nanocrystal surfaces became rough during etching (Fig. 7b–e), displaying “ocean wave-like” motions. Fig. 7e shows in situ HRTEM images that atomically revealed a single crystalline structure at the tips of the surface curvature in the liquid. The complete dissolution of the Sn nanocrystal takes a significantly longer time compared to the thick liquid layer (Fig. 7a). Additionally, a dense solution of Sn ions, resulting from the etching process, moves away from the nanocrystal surfaces (Fig. 7b). This dense layer of Sn ion products exhibits distinct contrast than the surrounding solution, making it easily distinguishable. It was suggested that reduced ion diffusion—both in terms of Sn ions diffusing away from the nanocrystal and the transport of oxidative species to the nanocrystal surface—plays a key role in the slower corrosion kinetics and the distinct solid-liquid interfaces observed in the thin liquid layer.
Fig. 7: Corrosion of Sn nanocrystals in an aqueous salt chloride solution.
figure 7
a Sequential TEM images illustrating the etching process of a Sn nanocrystal in a thick liquid layer under an electron dose rate of 7470 e− Å−2 s−1. b Sequential HRTEM images capturing the corrosion of a Sn nanocrystal in a thin liquid region at a higher electron dose rate of 12000 e− Å−2 s−1. c Contours of the Sn nanocrystal surfaces at the initial stage (0.0–13.8 s) and (d) at a later stage (20.0–33.0 s) during the corrosion process shown in (b). The colors represent the time sequence, with blue indicating the initial time and red representing the later time. e Enlarged view of representative images from panel (b), highlighting key stages of the corrosion process. Adapted from191 Copyright @2024 American Chemical Society.
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Summary and perspectives
The oxidation and corrosion processes of metals and alloys are complex phenomena involving a series of interrelated events at the atomic and nanoscale. Understanding these processes requires exploring several key aspects, including complex surface and interface dynamics, point and extended defect dynamics, diffusion mechanisms, phase evolution, the interplay between these factors, and their dependence on environmental variables such as temperature, pressure, and gases. Environmental TEM has emerged as a transformative tool for investigating oxidation dynamics at the atomic and nanoscale. By combining the high-resolution imaging and diffraction capabilities of TEM with a gaseous environment, this advanced technique allows for observing oxidation processes in situ under controlled environmental conditions, mimicking the real-world scenarios in which the metallic materials are used.
One of the most significant advantages of E-TEM is its ability to capture transient states and dynamic processes that would otherwise remain hidden in traditional bulk studies. The technique has revealed how environmental factors, such as humidity and temperature, influence oxidation pathways, leading to the discovery of previously unknown phenomena, such as the role of surface and interface defects in oxide nucleation38,40,53,55, interface migration143,144 or the influence of alloying elements on oxide nucleation and growth kinetics37,145,146,147. Recent studies utilizing aberration-corrected E-TEM have provided critical insights into the atomic-scale mechanisms underlying oxidation phenomena. For instance, the oxygen-adsorption-induced surface restructuring and subsequent oxide nucleation and growth can now be monitored as they occur68, allowing for a more comprehensive understanding of the initial stages of oxidation beyond the traditional UHV-based surface science experiments. Defects such as atomic vacancies significantly influence lattice diffusion and thus oxidation kinetics. Extended defects, such as dislocations and grain boundaries, can also serve as pathways for mass transport and thus oxide growth morphologies. The interaction between point defects and extended defects can further influence the kinetics of oxide growth. E-TEM is capable of probing atomically the effects of these structure defects on the oxide nucleation and growth as well as their dependence on oxidizing gases such as O2, H2O, CO2, etc.
Oxidation is inherently complex, due to the presence of multiple interface reactions. The key interfaces involved include the gas/oxide interface, oxide/metal interface, and the interface between different oxide phases because many metals exhibit multiple oxidation states and form various stable oxides and suboxides. As oxidation progresses, multiple oxide phases can form, each with distinct structural and chemical properties. The interfaces between these different oxide phases are critical for the overall stability and growth kinetics of the oxide layer. Phase transformations can occur due to changes in temperature, pressure, or local chemical environment, leading to the formation of new oxide structures. The interactions at these phase boundaries can influence defect generation and migration, as well as the diffusion pathways available for both cations and anions. For example, the transition from a less stable oxide phase to a more stable one can lead to a reduction in the number of vacancies, altering the diffusion dynamics and potentially impacting the rate of oxidation.
Directly investigating interfacial transformations during oxidation is inherently challenging, as it requires atomic-level resolution of the interface region. The insulating properties of oxides hinder the application of many surface-sensitive techniques that rely on the detection of charged particles, such as electrons and ions. Additionally, ensemble techniques like in situ X-ray diffraction struggle to differentiate between various interfacial reaction fronts due to the spatial and temporal overlap of signals from the heterogeneous mixture of bulk materials and interface boundaries. In contrast, E-TEM is not subject to these limitations. It excels at precisely resolving structural inhomogeneities and interfacial reaction dynamics at the atomic scale. Moreover, E-TEM’s capability for temperature-, pressure-, and time-resolved imaging allows for investigating how interfacial reaction dynamics are influenced by factors such as temperature, oxidizing gas, its pressure, and alloy composition. This comprehensive approach enables a deeper understanding of the complex mechanisms governing oxidation processes.
Although E-TEM is a powerful tool for studying oxidation processes at the atomic level, several limitations must be carefully considered when interpreting results from E-TEM such as the thinness of samples, temperature and pressure constraints and electron irradiation effects. One significant limitation of E-TEM is the requirement for sample thickness to be less than 100 nm to achieve sufficient electron transparency. The oxidation kinetics and mechanisms observed in thin foils may differ from those in bulk materials due to the high surface-to-volume ratio in thin foils, which can lead to enhanced reactivity and different oxidation rate compared to bulk materials. Although E-TEM can operate under various gas pressures, achieving the conditions found in real-world environments (such as high-pressure conditions) can be difficult. Additionally, variations in gas pressure during experiments can affect oxidation kinetics and mechanisms.
The electron beam used in gas-phase TEM can have significant effects such as enhancing dissociative ionization and excitation of gas molecules148, thus higher chemical activity for oxidation, radiation damage such as the formation of atomic defects149 to affect true oxidation dynamics, and irradiation-induced phase transformations in the oxide150,151, complicating the interpretation of results. In liquid-phase TEM experiments, the electron beam can induce various artifacts152,153, such as heating and radiolysis, leading to side reactions like the decomposition of H2O into aqueous electrons, H*, and OH*. To minimize these effects, the best approach is to use high beam energy to increase the mean free path and reduce scattering events, alongside a low electron dose rate to further decrease total scattering152,153. Recent advancements in direct electron cameras allow for (S)/TEM imaging at much lower electron doses, effectively minimizing the unwanted beam-induced events while maintaining high-quality imaging.
Looking ahead, several future perspectives can be envisioned for the application of E-TEM in the study of oxidation dynamics. First, advancements in detector technology, such as direct-electron cameras that capture 3500 image frames per second, combined with enhanced imaging algorithms154,155, are set to significantly improve the temporal resolution of E-TEM. These developments will facilitate the observation of faster oxidation processes and transient events occurring on the microsecond timescale, providing unprecedented insights into the rapid dynamics of oxidation phenomena. This enhanced capability will allow for the development of more accurate kinetic models of oxidation and enable real-time observation of the onset of oxidation, leading to a deeper understanding of the critical factors that govern these processes.
Second, the integration of machine learning and artificial intelligence with E-TEM data analysis holds great promise for accelerating materials discovery156,157,158,159. By analyzing vast datasets generated by E-TEM experiments, machine learning algorithms can identify patterns and correlations that may not be apparent to human observers. This approach can lead to the rapid identification of novel materials with optimized oxidation resistance or the prediction of long-term performance in corrosive environments. Additionally, expanding the capabilities of E-TEM to include multiple gas environments will be essential for studying the complex interactions that occur during oxidation. The ability to introduce various gases, vapors, or even electrolytes during imaging could provide insights into how different environmental conditions affect oxidation mechanisms. This would be particularly beneficial for developing materials for renewable energy applications, such as catalysts for fuel cells or batteries, where understanding the interplay between oxidation and reduction reactions is crucial.
Data availability
Raw data was not used in the preparation of this review article. The data presented are publicly available and can be found in the references section of the paper.
References
Li, X. G., Zhang, D. W., Liu, Z. Y., Du, C. W. & Dong, C. F. Materials science: share corrosion data. Nature 527, 441 (2015).
CASPubMedGoogle Scholar
Anber, E. A. et al. Oxidation resistance of Al-containing refractory high-entropy alloys. Scr. Mater. 244, 115997 (2024).
CASGoogle Scholar
Kumar, P. et al. Recent progress in oxidation behavior of high-entropy alloys: a review. APL Mater. 10, 120701 (2022).
CASGoogle Scholar
Gwalani, B. et al. Mechanistic understanding of speciated oxide growth in high entropy alloys. Nat. Commun. 15, 5026 (2024).
CASPubMedPubMed CentralGoogle Scholar
Gorr, B. et al. Current status of research on the oxidation behavior of refractory high entropy alloys. Adv. Eng. Mater. 23, 2001047 (2021).
CASGoogle Scholar
Besenbacher, F. & Norskov, J. K. Oxygen chemisorption on metal surfaces: general trends for Cu, Ni and Ag. Prog. Surf. Sci. 44, 5–66 (1993).
CASGoogle Scholar
Ertl, G. Reactions at surfaces: from atoms to complexity. Angew. Chem. Int. Ed. 47, 3524–3535 (2008).
CASGoogle Scholar
Stampfl, C., Ganduglia-Pirovano, M. V., Reuter, K. & Scheffler, M. Catalysis and corrosion: the theoretical surface-science context. Surf. Sci. 500, 368–394 (2002).
CASGoogle Scholar
Stampfl, C., Soon, A., Piccinin, S., Shi, H. Q. & Zhang, H. Bridging the temperature and pressure gaps: close-packed transition metal surfaces in an oxidizing environment. J. Phys.: Condens. Matter 20, 184021 (2008).
Google Scholar
Reuter, K., Stampfl, C. & Scheffler, M. Handbook of Materials Modeling, Part A: Methods (Springer, 2005).
Freund, H.-J. et al. Bridging the pressure and materials gaps between catalysis and surface science: clean and modified oxide surfaces. Top. Catal. 15, 201–209 (2001).
CASGoogle Scholar
Ferrer, S., Ackermann, M. D. & Lundgren, E. In situ investigations of chemical reactions on surfaces by X-ray diffraction at atmospheric pressures. MRS Bull. 32, 1010–1014 (2007).
CASGoogle Scholar
Cabrera, N. & Mott, N. F. Theory of the oxidation of metals. Rep. Prog. Phys. 12, 163–184 (1949).
CASGoogle Scholar
Wagner, C. Beitrag zur Theorie des Anlaufvorgangs. Z. Phys. Chem. B21, 25–41 (1933).
CASGoogle Scholar
Sharma, R. An environmental transmission electron microscope for in situ synthesis and characterization of nanomaterials. J. Mater. Res. 20, 1695–1707 (2005).
CASGoogle Scholar
Sharma, R. Kinetic measurements from in situ TEM observations. Microsc. Res. Tech. 72, 144–152 (2009).
CASPubMedGoogle Scholar
Boyes, E. D. & Gai, P. L. Environmental high resolution electron microscopy and applications to chemical science. Ultramicroscopy 67, 219–232 (1997).
CASGoogle Scholar
Gai, P. L. et al. Atomic-resolution environmental transmission electron microscopy for probing gas–solid reactions in heterogeneous catalysis. MRS Bull. 32, 1044–1050 (2007).
CASGoogle Scholar
Yaguchi, T., Gabriel, M. L. S., Hashimoto, A. & Howe, J. Y. In-situ TEM study from the perspective of holders. Microscopy 73, 117–132 (2023).
Google Scholar
Allard, L. F. et al. Novel MEMS-based gas-cell/heating specimen holder provides advanced imaging capabilities for in situ reaction studies. Microsc. Microanal. 18, 656–666 (2012).
CASPubMedGoogle Scholar
Unocic, K. A. et al. Introducing and controlling water vapor in closed-cell in situ electron microscopy gas reactions. Microsc. Microanal. 26, 229–239 (2020).
CASPubMedGoogle Scholar
Grogan, J. M. & Bau, H. H. The nanoaquarium: a platform for in situ transmission electron microscopy in liquid media. J. Microelectromech. Syst. 19, 885–894 (2010).
CASGoogle Scholar
Zheng, H. et al. Observation of single colloidal platinum nanocrystal growth trajectories. Science 324, 1309–1312 (2009).
CASPubMedGoogle Scholar
Xin, H. L. & Zheng, H. M. In situ observation of oscillatory growth of bismuth nanoparticles. Nano Lett. 12, 1470–1474 (2012).
CASPubMedGoogle Scholar
Liquid Cell Electron Microscopy (Cambridge University Press, 2016).
Kosari, A. et al. Challenges to electrochemical evaluation of nanometric sandwiched thin specimens using liquid cells designed for application in liquid-phase TEM corrosion studies. Corros. Sci. 192, 109864 (2021).
CASGoogle Scholar
Hwang, S., Chen, X. B., Zhou, G. W. & Su, D. In-situ transmission electron microscopy on energy-related catalysis. Adv. Energy Materials 10, 1902105 (2019).
Yang, R. et al. Fabrication of liquid cell for in situ transmission electron microscopy of electrochemical processes. Nat. Protoc. 18, 555–578 (2023).
CASPubMedGoogle Scholar
Yuk, J. M. et al. High-resolution EM of colloidal nanocrystal growth using graphene liquid cells. Science 336, 61–64 (2012).
CASPubMedGoogle Scholar
Yuk, J. M. et al. Real-time observation of water-soluble mineral precipitation in aqueous solution by in situ high-resolution electron microscopy. ACS Nano 10, 88–92 (2016).
CASPubMedGoogle Scholar
Haider, M. et al. Electron microscopy image enhanced. Nature 392, 768–769 (1998).
CASGoogle Scholar
Batson, P. E., Dellby, N. & Krivanek, O. L. Sub-ångstrom resolution using aberration corrected electron optics. Nature 418, 617–620 (2002).
CASPubMedGoogle Scholar
Boyes, E. D., LaGrow, A. P., Ward, M. R., Martin, T. E. & Gai, P. L. Visualizing single atom dynamics in heterogeneous catalysis using analytical in situ environmental scanning transmission electron microscop. Philos. Trans. R: Soc. A 378, 20190605 (2020).
Google Scholar
Boyes, E. D. & Gai, P. L. Visualizing reacting single atoms in chemical reactions: advancing the frontiers of materials research. MRS Bull. 40, 600–609 (2015).
CASGoogle Scholar
Gai, P. L. & Boyes, E. D. Advances in atomic resolution in situ environmental transmission electron microscopy and 1Å aberration corrected in situ electron microscopy. Microsc. Res. Tech. 72, 153–164 (2009).
CASPubMedGoogle Scholar
LaGrow, A. P., Ward, M. R., Lloyd, D. C., Gai, P. L. & Boyes, E. D. Visualizing the Cu/Cu2O interface transition in nanoparticles with environmental scanning transmission electron microscopy. J. Am. Chem. Soc. 139, 179–185 (2017).
CASPubMedGoogle Scholar
Zhou, G. W. et al. Composition effects on the early-stage oxidation kinetics of (001)Cu-Au alloys. J. Appl. Phys. 101, 033521 (2007).
Google Scholar
Zhou, G. W. et al. In situ atomic-scale visualization of oxide islanding during oxidation of Cu surfaces. Chem. Comm. 49, 10862–10864 (2013).
CASPubMedGoogle Scholar
Zhou, G. W. & Yang, J. C. In situ UHV-TEM investigation of the kinetics of initial stages of oxidation on the roughened Cu(110) surface. Surf. Sci. 559, 100–110 (2004).
CASGoogle Scholar
Zhou, G. W. & Yang, J. C. Initial oxidation kinetics of Cu(100), (110), and (111) thin films investigated by in situ ultra-high vacuum transmission electron microscopy. J. Mater. Res. 20, 1684–1694 (2005).
CASGoogle Scholar
Chen, X. B. et al. Passive oxide film growth observed on the atomic scale. Adv. Mater. Interfaces 9, 2102487 (2022).
CASGoogle Scholar
Chen, X. B. et al. Atomic mechanisms of water-vapor induced surface passivation. Sci. Adv. 9, eadh5565 (2023).
CASPubMedPubMed CentralGoogle Scholar
Li, L. et al. Surface-step-induced oscillatory oxide growth. Phys. Rev. Lett. 113, 136104 (2014).
PubMedGoogle Scholar
Luo, L. L. et al. Atomic origins of water-vapour-promoted alloy oxidation. Nat. Mater. 17, 514–518 (2018).
CASPubMedGoogle Scholar
Luo, L. L. et al. In situ atomic scale visualization of surface kinetics driven dynamics of oxide growth on a Ni–Cr surface. Chem. Comm. 52, 3300–3303 (2016).
CASPubMedGoogle Scholar
Zhou, G. W. et al. Step-edge induced oxide growth during the oxidation of Cu surfaces. Phys. Rev. Lett. 109, 235502 (2012).
PubMedGoogle Scholar
Zhou, G. W. & Yang, J. C. Formation of Quasi-one-dimensional Cu2O structures by in situ oxidation of Cu(100). Phys. Rev. Lett. 89, 106101 (2002).
PubMedGoogle Scholar
Zhu, D. D. et al. Atomic origin of CO2-promoted oxidation dynamics of chromia-forming alloys. Acta Materialia 264, 119578 (2024).
CASGoogle Scholar
Gulec, A. et al. Early stage of oxidation of Mo3Si by in situ environmental transmission electron microscopy. Corrosion 74, 288–294 (2018).
CASGoogle Scholar
Wang, C.-M., Schreiber, D. K., Olszta, M. J., Baer, D. R. & Bruemmer, S. M. Direct in situ TEM observation of modification of oxidation by the injected vacancies for Ni–4Al alloy using a microfabricated nanopost. ACS Appl. Mater. Interfaces 7, 17272–17277 (2015).
CASPubMedGoogle Scholar
Zhu, D. D. et al. One-dimensional γ-Al2O3 growth from the oxidation of NiAl. Corr. Sci. 216, 111069 (2023).
CASGoogle Scholar
Zhou, G. W., Wang, L. & Yang, J. C. Effect of surface topology on the formation of oxide islands on Cu surfaces. J. Appl. Phys. 97, 063509 (2005).
Google Scholar
Zhou, G. W. & Yang, J. C. Initial oxidation kinetics of copper (110) film investigated by in situ UHV-TEM. Surf. Sci. 531, 359–367 (2003).
CASGoogle Scholar
Zhou, G. W. & Yang, J. C. Temperature effect on the Cu2O oxide morphology created by oxidation of Cu(001) as Investigated by in situ UHV TEM. Appl. Surf. Sci. 210, 165–170 (2003).
CASGoogle Scholar
Zhou, G. W. & Yang, J. C. Temperature effects on the growth of oixde islands on Cu(110). Appl. Surf. Sci. 222, 357–364 (2004).
CASGoogle Scholar
Zhou, G. W., Chen, X. D., Gallagher, D. & Yang, J. C. Percolating oxide film growth during Cu(111) oxidation. Appl. Phys. Lett. 93, 123104 (2008).
Google Scholar
Yang, J. C., Bharadwaj, M. D., Zhou, G. W. & Tropia, L. Surface kinetics of copper oxidation investigated by in-situ UHV transmission electron microscopy. Microsc. Microanalysis 7, 486–493 (2001).
CASGoogle Scholar
Yang, J. C., Evan, D. & Tropia, L. From nucleation to coalescence of Cu2O islands during in situ oxidation of Cu(001). Appl. Phys. Lett. 81, 241–243 (2002).
CASGoogle Scholar
Yang, J. C., Kolasa, B., Gibson, J. M. & Yeadon, M. Self-limiting oxidation of copper. Appl. Phys. Lett. 73, 2841–2843 (1998).
CASGoogle Scholar
Yang, J. C., Yeadon, M., Kolasa, B. & Gibson, J. M. The homogeneous nucleation mechanism of Cu2O on Cu(001). Scr. Mater. 38, 1237–1241 (1998).
CASGoogle Scholar
Birks, N., Meier, G. H. & Pettit, F. S. Introduction to the High Temperature Oxidation of Metals 2nd edn. (Cambridge University Press, 2006).
Gattinoni, C. & Michaelides, A. Atomistic details of oxide surfaces and surface oxidation: the example of copper and its oxides. Surf. Sci. Rep. 70, 424–447 (2015).
CASGoogle Scholar
Young, D. J. High Temperature Oxidation and Corrosion of Metals. 2nd ed.; Elsevier: (2016).
Zhou, G. et al. Revealing atomic-to-nanoscale oxidation mechanisms of metallic materials. MRS Bull. 48, 852–863 (2023).
CASGoogle Scholar
Liu, Q. Q., Li, L., Cai, N., Saidi, W. A. & Zhou, G. W. Oxygen chemisorption-induced surface phase transitions on Cu(110). Surf. Sci. 627, 75–84 (2014).
CASGoogle Scholar
Wu, D. X. et al. Revealing an intermediate Cu–O/OH superstructure on Cu(110). J. Phys. Chem. Lett. 13, 2396–2403 (2022).
CASPubMedGoogle Scholar
Besenbacher, F. Scanning tunnelling microscopy studies of metal surfaces. Rep. Prog. Phys. 59, 1737–1802 (1996).
CASGoogle Scholar
Wu, D. X., Sun, X. H., Yang, J. C. & Zhou, G. W. Intermittent oxidation kinetics and metal/oxide interfacial undulation. Phys. Rev. B 110, 085402 (2024).
CASGoogle Scholar
Suo, Z. Wrinkling of the oxide scale on an aluminum-containing alloy at high temperatures. J. Mech. Phys. Solids 43, 829–846 (1995).
CASGoogle Scholar
Goncalves, A. M., Campos, L. C., Ferlauto, A. S. & Lacerda, R. G. On the growth and electrical characterization of CuO nanowires by thermal oxidation. J. Appl. Phys. 106, 034303 (2009).
Google Scholar
Raynaud, G. M. & Rapp, R. A. In situ observation of whiskers, pyramids and pits during the high-temperature oxidation of metals. Oxid. Met. 21, 89–102 (1984).
CASGoogle Scholar
Kumar, A., Srivastava, A. K., Tiwari, P. & Nandedkar, R. V. The effect of growth parameters on the aspect ratio and number density of CuO nanorods. J. Phys.: Condens. Matter 16, 8531–8543 (2004).
CASGoogle Scholar
Chen, J. T. et al. CuO nanowires synthesized by thermal oxidation route. J. Alloy. Compd. 454, 268–273 (2008).
CASGoogle Scholar
Xu, C. H., Woo, C. H. & Shi, S. Q. Formation of CuO nanowires on Cu foil. Chem. Phys. Lett. 399, 62–66 (2004).
CASGoogle Scholar
Kaur, M. et al. Growth and branching of CuO nanowires by thermal oxidation of copper. J. Cryst. Growth 289, 670–675 (2006).
CASGoogle Scholar
Zhang, K., Rossl, C., Tenalleau, C., Alphonse, P. & Chane-Ching, J. Y. Synthesis of large-area and aligned copper oxide nanowires from copper thin film on silicon substrate. Nanotechnology 19, 275607 (2007).
Google Scholar
Chopra, N., Hu, B. & Hinds, B. Selective growth and kinetic study of copper oxide nanowires from patterned thin-film multilayer structures. J. Mater. Res. 22, 2691–2699 (2007).
CASGoogle Scholar
Appleby, W. K. & Tylecote, R. F. Stresses during the gaseous oxidation of metals. Corros. Sci. 10, 325 (1970).
Morin, F. Copious whisker growth on copper scale. J. Mater. Sci. Lett. 2, 383–384 (1983).
CASGoogle Scholar
Evans, H. E. Stress effects in high temperature oxidation of metals. Int. J. Mater. Res. 40, 1–40 (1995).
CASGoogle Scholar
Rapp, R. A. The high temperature oxidatoin of metals forming cation-diffusing scales. Metall. Trans. B 15B, 195–212 (1983).
Google Scholar
Zhu, Y., Mimura, K. & Isshiki, M. Oxidation mechanism of Cu2O to CuO at 600–1050 °C. Oxid. Met. 62, 207–222 (2004).
CASGoogle Scholar
Zhu, Y. F., Mimura, K. & Isshiki, M. Oxidation mechanism of Copper at 623-1073K. Mater. Trans. 43, 2173–2176 (2002).
CASGoogle Scholar
Grzesik, Z. & Migdalska, M. On the mechanism of Cu2O oxidation at high temperature. Defect Diffus. Forum 289-292, 429–436 (2009).
CASGoogle Scholar
Zhu, Y. F., Mimura, K., Lim, J. W., Isshiki, M. & Jiang, Q. Brief review of oxidation kinetics of copper at 350 °C to 1050 °C. Metall. Mater. Trans. A 37A, 1231–1237 (2006).
CASGoogle Scholar
Yuan, L., Wang, Y. Q., Mema, R. & Zhou, G. W. Driving force and growth mechanism for spontaneous oxide nanowire formation during the thermal oxidation of metals. Acta Materialia 59, 2491–2500 (2011).
CASGoogle Scholar
Yuan, L. et al. The origin of hematite nanowire growth during the thermal oxidation of iron. Mater. Sci. Eng. B 177, 327–336 (2012).
CASGoogle Scholar
Wang, R. M., Chen, Y. F., Fu, Y. Y., Zhang, H. & Kisielowski, C. Bicrystalline hematite nanowires. J. Phys. Chem. B 109, 12245–12249 (2005).
CASPubMedGoogle Scholar
Yuan, L., Wang, C., Cai, R. S., Wang, Y. Q. & Zhou, G. W. Spontaneous ZnO nanowire formation during the oxidation of Cu-Zn alloy. J. Appl. Phys. 114, 023512 (2013).
Google Scholar
Yuan, L. & Zhou, G. W. Enhanced CuO nanowire formation by thermal oxidation of roughened copper. J. Electrochem. Soc. 159, C205–C209 (2012).
CASGoogle Scholar
Sun, X. H. et al. Atomic-scale mechanism of unidirectional oxide growth. Adv. Funct. Mater. 30, 1906504 (2019).
Zou, G. et al. Well-aligned arrays of CuO nanoplatelets. J. Phys. Chem. B 110, 1632–1637 (2006).
CASPubMedGoogle Scholar
Yuan, L., Jiang, Q. K., Wang, J. B. & Zhou, G. W. The growth of hematite nanobelts and nanowires-tune the shape via oxygen gas pressure. J. Mater. Res. 27, 1014–1021 (2012).
CASGoogle Scholar
Yuan, L. et al. Morphological transformation of hematite nanostructures during oxidation of iron. Nanoscale 5, 7581–7588 (2013).
CASPubMedGoogle Scholar
Wang, Y. Q. et al. Coincidence-site-lattice twist boundaries in bicrystalline α-Fe2O3 nanoblades. J. Phys. Chem. C 118, 5796–5801 (2014).
CASGoogle Scholar
Zhu, W. H. et al. In situ atomic-scale probing of the reduction dynamics of two-dimensional Fe2O3 nanostructures. ACS Nano 11, 656–664 (2017).
CASPubMedGoogle Scholar
Dufour, L. & Morin, F. The growth of oxide platelets on nickel in pure oxygen. II. Surface analyses and growth mechanism. Oxid. Met. 39, 137–154 (1993).
CASGoogle Scholar
Kong, Y. C., Yu, D. P., Zhang, B., Fang, W. & Feng, S. Q. Ultraviolet-emitting ZnO nanowires synthesized by a physical vapor deposition approach. Appl. Phys. Lett. 78, 407–409 (2001).
CASGoogle Scholar
Xing, Y. J. et al. Optical properties of the ZnO nanotubes synthesized via vapor phase growth. Appl. Phys. Lett. 83, 1689–1691 (2003).
CASGoogle Scholar
Yuan, L., Wang, C., Cai, R., Wang, Y. & Zhou, G. Temperature-dependent growth mechanism and microstructure of ZnO nanostructures grown from the thermal oxidation of zinc. J. Cryst. Growth 390, 101–108 (2014).
CASGoogle Scholar
Zhou, J. et al. Large‐area nanowire arrays of molybdenum and molybdenum oxides: synthesis and field emission properties. Adv. Mater. 15, 1835–1840 (2003).
CASGoogle Scholar
Zheng, L., Xu, Y., Jin, D. & Xie, Y. Novel metastable hexagonal MoO3 nanobelts: synthesis, photochromic, and electrochromic properties. Chem. Mater. 21, 5681–5690 (2009).
CASGoogle Scholar
Yuan, L., Wang, C., Cai, R., Wang, Y. & Zhou, G. Spontaneous ZnO nanowire formation during oxidation of Cu-Zn alloy. J. Appl. Phys. 114, 023512 (2013).
Google Scholar
Yuan, L. & Zhou, G. W. Enhanced CuO nanowire formation by thermal oxidation of roughened copper. J. Electrochem. Soc. 159, C205 (2012).
CASGoogle Scholar
Holcomb, G. R., Carney, C. & Doğan, Ö. N. Oxidation of alloys for energy applications in supercritical CO2 and H2. O. Corr. Sci. 109, 22–35 (2016).
CASGoogle Scholar
Oleksak, R. P., Carney, C., Teeter, L. & Doğan, Ö. N. Oxidation behavior of welded Fe-based and Ni-based alloys in supercritical CO2. Oxid. Met. 123–139 (2022).
Li, K., Zeng, Y. & Luo, J.-L. Corrosion of SS310 and Alloy 740 in high temperature supercritical CO2 with impurities H2O and O2. Corr. Sci. 184, 109350 (2021).
CASGoogle Scholar
Nguyen, T. D., Zhang, J. & Young, D. J. Effects of cerium and manganese on corrosion of Fe–Cr and Fe–Cr–Ni alloys in Ar–20CO2 gas at 818 °C. Corr. Sci. 76, 231–242 (2013).
CASGoogle Scholar
Gheno, T., Monceau, D. & Young, D. J. Mechanism of breakaway oxidation of Fe–Cr and Fe–Cr–Ni alloys in dry and wet carbon dioxide. Corr. Sci. 64, 222–233 (2012).
CASGoogle Scholar
Young, D. J., Nguyen, T. D., Felfer, P., Zhang, J. & Cairney, J. M. Penetration of protective chromia scales by carbon. Scr. Mater. 77, 29–32 (2014).
CASGoogle Scholar
Zheng, X. G. & Young, D. J. High-temperature corrosion of Cr2O3-forming alloys in CO-CO2-N2 atmospheres. Oxid. Met. 42, 163–190 (1994).
CASGoogle Scholar
Nguyen, T. D., Xie, Y., Ding, S., Zhang, J. & Young, D. J. Oxidation Behavior of Ni–Cr Alloys in CO2 at 700 °C. Oxid. Met. 87, 605–616 (2017).
CASGoogle Scholar
Jiang, C., Xie, Y., Kong, C., Zhang, J. & Young, D. J. Corrosion behaviour of Ni-Cr alloys in mixed oxidising gases at 650 °C. Corr. Sci. 174, 108801 (2020).
CASGoogle Scholar
Young, D. J. & Zhang, J. Corrosion by Hot CO2 Gases. Electrochem. Soc. Interface 30, 73 (2021).
CASGoogle Scholar
Pint, B. A., Pillai, R., Lance, M. J. & Keiser, J. R. Effect of pressure and thermal cycling on long-term oxidation in CO2 and supercritical CO2. Oxid. Met. 94, 505–526 (2020).
Google Scholar
Oleksak, R. P. et al. Molecular-scale investigation of the oxidation behavior of chromia-forming alloys in high-temperature CO2. npj Mater. Degrad. 5, 46 (2021).
CASGoogle Scholar
Reichel, F., Jeurgens, L. P. H. & Mittemeijer, E. J. The effect of substrate orientation on the kinetics of ultra-thin oxide film growth on Al single crystals. Acta Materialia 56, 2897–2907 (2008).
CASGoogle Scholar
Jeurgens, L. P. H., Sloof, W. G., Tichelaar, F. D. & Mittemeijer, E. J. Thermodynamic stability of amorphous oxide films on metals: application to aluminum oxide films on aluminum substrates. Phys. Rev. B 62, 4707 (2000).
CASGoogle Scholar
Liu, Q. Q., Tong, X. & Zhou, G. W. H2O dissociation induced aluminum oxide growth on oxidized Al(111) surfaces. Langmuir 31, 13117–13126 (2015).
CASPubMedGoogle Scholar
Cai, N., Liu, Q. Q., Tong, X. & Zhou, G. W. X-ray photoelectron spectroscopic study of the passivation of NiAl(100) surface by water vapor. Langmuir 30, 774–783 (2014).
CASPubMedGoogle Scholar
Cai, N., Qin, H. L., Xiao, T. & Zhou, G. W. Growth of ultrathin amorphous alumina films during the oxidation of NiAl(100). Surf. Sci. 618, 20–26 (2013).
CASGoogle Scholar
Cai, N., Zhou, G. W., Muller, K. & Starr, D. E. Tuning the limiting-thickness of a thin oxide layer on Al (111) with oxygen gas pressure. Phys. Rev. Lett. 107, 035502 (2011).
PubMedGoogle Scholar
Cai, N., Zhou, G. W., Muller, K. & Starr, D. E. Effect of oxygen gas pressure on the kinetics of alumina film growth during the oxidation of Al(111) at room temperature. Phys. Rev. B 84, 125445 (2011).
Google Scholar
Cai, N., Zhou, G. W., Muller, K. & Starr, D. E. Temperature and pressure dependent Mott potentials and their influence on self-limiting oxide film growth. Appl. Phys. Lett. 101, 171605 (2012).
Google Scholar
Cai, N., Zhou, G. W., Muller, K. & Starr, D. E. Comparative study of the passivation of Al(111) by molecular oxygen and water vapor. J. Phys. Chem. C 117, 172–178 (2013).
CASGoogle Scholar
Zhu, Y. G. et al. Atomistic mechanisms of the initial oxidation of stepped Cu3Au(100). Phys. Rev. B 105, 075422 (2022).
CASGoogle Scholar
Brune, H. et al. Interaction of oxygen with Al(lll) studied by scanning tunneling microscopy. J. Chem. Phys. 99, 2128–2148 (1993).
CASGoogle Scholar
Frenzel, J., Oliveira, A. F., Duarte, H. A., Heine, T. & Seifert, G. Structural and electronic properties of bulk gibbsite and gibbsite surfaces. Z. Anorg. Allg. Chem. 631, 1267–1271 (2005).
CASGoogle Scholar
Fleming, S., Rohl, A., Lee, M. Y., Gale, J. & Parkinson, G. Atomistic modelling of gibbsite: surface structure and morphology. J. Cryst. Growth 209, 159–166 (2000).
Xiao, F., Yang, R. & Liu, Z. Active aluminum composites and their hydrogen generation via hydrolysis reaction: a review. Int. J. Hydrog. Energy 47, 365–386 (2022).
CASGoogle Scholar
Chen, X. et al. Popcorn-like aluminum-based powders for instant low-temperature water vapor hydrogen generation. Mater. Today Energy 19, 100602 (2021).
CASGoogle Scholar
Elitzur, S., Rosenband, V. & Gany, A. Study of hydrogen production and storage based on aluminum-water reaction. Int. J. Hydrog. Energy 39, 6328–6334 (2014).
CASGoogle Scholar
Ross, F. M. Opportunities and challenges in liquid cell electron microscopy. Science 350, aaa9886 (2015).
PubMedGoogle Scholar
Mirsaidov, U., Patterson, J. P. & Zheng, H. Liquid phase transmission electron microscopy for imaging of nanoscale processes in solution. MRS Bull. 45, 704–712 (2020).
Google Scholar
Unocic, R. R., Jungjohann, K. L., Mehdi, B. L., Browning, N. D. & Wang, C. In situ electrochemical scanning/transmission electron microscopy of electrode–electrolyte interfaces. MRS Bull. 45, 738–745 (2020).
Google Scholar
Kosari, A. et al. Application of in situ liquid cell transmission electron microscopy in corrosion studies: a critical review of challenges and achievements. Corrosion 76, 4–17 (2019).
Google Scholar
Hattar, K. & Unocic, R. R. in Recent Developments in Analytical Techniques for Corrosion Research (ed. Toor, I. U.) 121–150 (Springer International Publishing, 2022).
Chee, S. W., Duquette, D. J., Ross, F. M. & Hull, R. Metastable structures in Al thin films before the onset of corrosion pitting as observed using liquid cell transmission electron microscopy. Microsc. Microanal. 20, 462–468 (2014).
CASPubMedGoogle Scholar
Chee, S. W. et al. Studying localized corrosion using liquid cell transmission electron microscopy. Chem. Commun. 51, 168–171 (2015).
CASGoogle Scholar
Hayden, S. C. et al. Genesis of nanogalvanic corrosion revealed in pearlitic steel. Nano Lett. 22, 7087–7093 (2022).
CASPubMedPubMed CentralGoogle Scholar
Kosari, A. et al. In-situ nanoscopic observations of dealloying-driven local corrosion from surface initiation to in-depth propagation. Corros. Sci. 177, 108912 (2020).
CASGoogle Scholar
Jonge, N. D., Peckys, D. B., Kremers, G. J. & Piston, D. W. Electron microscopy of whole cells in liquid with nanometer resolution. Proc. Natl Acad. Sci. 106, 2159–2164 (2009).
PubMedPubMed CentralGoogle Scholar
Sun, X. H. et al. Dislocation-induced stop-and-go kinetics of interfacial transformations. Nature 607, 708–713 (2022).
CASPubMedGoogle Scholar
Sun, X. et al. Atomic dynamics of multi-interfacial migration and transformations. Small 23, 2305746 (2023).
Google Scholar
Zhou, G. W. et al. Cu2O island shape transition during Cu-Au oxidation. Phys. Rev. Lett. 96, 226108 (2006).
PubMedGoogle Scholar
Luo, L. L. et al. Comparative study of the alloying effect on the initial oxidation of Cu-Au(100) and Cu-Pt(100). Appl. Phys. Lett. 104, 121601 (2014).
Google Scholar
Luo, L. L., Kang, Y. H., Yang, J. C. & Zhou, G. W. Effect of Au composition on the orientations of oxide nuclei during the oxidation of Cu-Au alloys. J. Appl. Phys. 111, 083533 (2012).
Google Scholar
Cosby, P. C. Electron-impact dissociation of oxygen. J. Chem. Phys. 98, 9560–9569 (1993).
CASGoogle Scholar
Jiang, N. Electron irradiation effects in transmission electron microscopy: random displacements and collective migrations. Micron 171, 103482 (2023).
CASPubMedGoogle Scholar
Zhu, Q. et al. Electron-irradiation induced unconventional phase transition of β-Ga2O3 epitaxial single-crystal thin film observed by in-situ TEM. J. Mater. Res. Technol. 30, 2397–2405 (2024).
CASGoogle Scholar
Jiang, N. Electron beam damage in oxides: a review. Rep. Prog. Phys. 79, 016501 (2016).
PubMedGoogle Scholar
Egerton, R. F., Li, P. & Malac, M. Radiation damage in the TEM and SEM. Micron 35, 399–409 (2004).
CASPubMedGoogle Scholar
Pu, S., Gong, C. & Robertson, A. W. Liquid cell transmission electron microscopy and its applications. R. Soc. Open Sci. 7, 191204 (2020).
CASPubMedPubMed CentralGoogle Scholar
Ishida, T. et al. Development of silicon-on-insulator direct electron detector with analog memories in pixels for sub-microsecond imaging. Microscopy 73, 511–516 (2024).
Faruqi, A. R. & McMullan, G. Direct imaging detectors for electron microscopy. Nucl. Instrum. Methods Phys. Res. A: Accelerators Spectrometers Detect. Associated Equip. 878, 180–190 (2018).
CASGoogle Scholar
Zheng, H., Lu, X. & He, K. In situ transmission electron microscopy and artificial intelligence enabled data analytics for energy materials. J. Energy Chem. 68, 454–493 (2022).
CASGoogle Scholar
Botifoll, M., Pinto-Huguet, I. & Arbiol, J. Machine learning in electron microscopy for advanced nanocharacterization: current developments, available tools and future outlook. Nanoscale Horiz. 7, 1427–1477 (2022).
CASPubMedGoogle Scholar
Kalinin, S. V. et al. Machine learning for automated experimentation in scanning transmission electron microscopy. npj Comput. Mater. 9, 227 (2023).
Google Scholar
Treder, K. P., Huang, C., Kim, J. S. & Kirkland, A. I. Applications of deep learning in electron microscopy. Microscopy 71, i100–i115 (2022).
CASPubMedGoogle Scholar
Zhou, G. W. Stress-driven formation of terraced hollow oxide nanorods during metal oxidation. J. Appl. Phys. 105, 104302 (2009).
Google Scholar
Zhou, G. W., Slaughter, W. S. & Yang, J. C. Terraced hollow oxide pyramids. Phys. Rev. Lett. 94, 246101 (2005).
Google Scholar
Luo, L. L., Kang, Y. H., Liu, Z. Y., Yang, J. C. & Zhou, G. W. Dependence of degree of orientation of copper oxide nuclei on oxygen pressure during initial stages of copper oxidation. Phys. Rev. B 83, 155418 (2011).
Google Scholar
Luo, L. L., Kang, Y. H., Yang, J. C. & Zhou, G. W. Influence of the surface morphology on the early stages of Cu oxidation. Appl. Surf. Sci. 259, 791–798 (2012).
CASGoogle Scholar
Yang, J. C. & Zhou, G. In situ ultra-high vacuum transmission electron microscopy studies of the transient oxidation stage of Cu and Cu alloy thin films. Micron 43, 1195–1210 (2012).
CASPubMedGoogle Scholar
Zou, L. F. et al. Segregation induced order-disorder transition in Cu(Au) surface alloys. Acta Materialia 154, 220–227 (2018).
CASGoogle Scholar
Nguyen, L. et al. Atomic-scale insights into the oxidation of aluminum. ACS Appl. Mater. Interfaces 10, 2230–2235 (2018).
CASPubMedGoogle Scholar
Li, S. et al. Selective atomic sieving across metal/oxide interface for super-oxidation resistance. Nat. Commun. 15, 6149 (2024).
CASPubMedPubMed CentralGoogle Scholar
Liu, S. et al. Microstructure and oxidation of NiCr alloys studied by analytical in situ environmental TEM. Corros. Sci. 224, 111525 (2023).
CASGoogle Scholar
Luo, L. et al. Deciphering atomistic mechanisms of the gas-solid interfacial reaction during alloy oxidation. Sci. Adv. 6, eaay8491 (2020).
CASPubMedPubMed CentralGoogle Scholar
Luo, L. L. et al. In-situ transmission electron microscopy study of surface oxidation for Ni–10Cr and Ni–20Cr alloys. Scr. Materialia 114, 129–132 (2016).
CASGoogle Scholar
Togaru, M., Sainju, R., Zhang, L., Jiang, W. & Zhu, Y. Direct observation of tungsten oxidation studied by in situ environmental TEM. Mater. Charact. 174, 111016 (2021).
CASGoogle Scholar
Unocic, K. A., Shin, D., R.R. Unocic, R. R. & Allard, L. F. NiAl oxidation reaction processes studied in situ using MEMS-based closed-cell gas reaction transmission electron microscopy. Oxid. Met. 88, 495–508 (2017).
CASGoogle Scholar
Yu, X.-X. et al. In situ observations of early stage oxidation of Ni-Cr and Ni-Cr-Mo alloys. Corrosion 74, 939–946 (2018).
CASGoogle Scholar
Qu, B. & van Benthem, K. In-situ anisotropic growth of nickel oxide nanostructures through layer-by-layer metal oxidation. Scr. Materialia 214, 114660 (2022).
CASGoogle Scholar
Gulec, A. et al. Early stage of oxidation of Mo3Si by in situ environmental transmission electron microscopy. Corrosion 74, 288–294 (2017).
Google Scholar
Li, Y. et al. In situ tem studies of the oxidation of Li dendrites at high temperatures. Adv. Funct. Mater. 32, 2203233 (2022).
CASGoogle Scholar
Burke, M. G., Bertali, G., Prestat, E., Scenini, F. & Haigh, S. J. The application of in situ analytical transmission electron microscopy to the study of preferential intergranular oxidation in Alloy 600. Ultramicroscopy 176, 46–51 (2017).
CASPubMedGoogle Scholar
Xin, H. L. et al. Revealing the atomic restructuring of Pt–Co nanoparticles. Nano Lett. 14, 3203–3207 (2014).
CASPubMedGoogle Scholar
Dai, S. et al. Revealing surface elemental composition and dynamic processes involved in facet-dependent oxidation of Pt3Co nanoparticles via in situ transmission electron microscopy. Nano Lett. 17, 4683–4688 (2017).
CASPubMedGoogle Scholar
Song, B. et al. In situ oxidation studies of high-entropy alloy nanoparticles. ACS Nano 14, 15131–15143 (2020).
CASPubMedGoogle Scholar
Wang, C.-M. et al. In-Situ TEM visualization of vacancy injection and chemical partition during oxidation of Ni-Cr nanoparticles. Sci. Rep. 4, 3683 (2014).
PubMedPubMed CentralGoogle Scholar
Gao, J., Yan, J., Zhao, B., Zhang, Z. & Yu, Q. In situ observation of temperature-dependent atomistic and mesoscale oxidation mechanisms of aluminum nanoparticles. Nano Res. 13, 183–187 (2020).
CASGoogle Scholar
Zhu, B. et al. Reshaping of metal nanoparticles under reaction conditions. Angew. Chem. Int. Ed. 59, 2171–2180 (2020).
CASGoogle Scholar
Hayden, S. C. et al. Localized corrosion of low-carbon steel at the nanoscale. npj Mater. Degrad. 3, 17 (2019).
Google Scholar
Tian, M., Choundraj, J. D., Voisin, T., Wang, Y. M. & Kacher, J. Discovering the nanoscale origins of localized corrosion in additive manufactured stainless steel 316L by liquid cell transmission electron microscopy. Corros. Sci. 208, 110659 (2022).
CASGoogle Scholar
Kosari, A. et al. Nanoscopic and in-situ cross-sectional observations of Li-based conversion coating formation using liquid-phase TEM. npj Mater. Degrad. 5, 40 (2021).
CASGoogle Scholar
Gross, D. et al. Observations of Corrosion in Nanocrystalline Fe Thin Films. In Processing, Properties, and Design of Advanced Ceramics and Composites II 327–338 (2018)
Key, J. W. et al. Investigating local oxidation processes in Fe thin films in a water vapor environment by in situ liquid cell TEM. Ultramicroscopy 209, 112842 (2020).
CASPubMedGoogle Scholar
Agrawal, S., Sk, M. H., Langford, R. M. & Clarke, S. M. In-situ iron corrosion in brine using TEM. MRS Adv. 8, 376–380 (2023).
CASGoogle Scholar
Wang, X., Hung, T.-F., Chen, F.-R. & Wang, W.-X. In situ tracking of crystal-surface-dependent Cu2O nanoparticle dissolution in an aqueous environment. Environ. Sci. Technol. 57, 1006–1016 (2023).
CASPubMedGoogle Scholar
Peng, X. et al. Unveiling corrosion pathways of Sn nanocrystals through high-resolution liquid cell electron microscopy. Nano Lett. 24, 1168–1175 (2024).
CASPubMedPubMed CentralGoogle Scholar
Shi, F. et al. Strain-induced corrosion kinetics at nanoscale are revealed in liquid: enabling control of corrosion dynamics of electrocatalysis. Chem 6, 2257–2271 (2020).
CASGoogle Scholar
Li, X. et al. Dissolution behavior of isolated and aggregated hematite particles revealed by in situ liquid cell transmission electron microscopy. Environ. Sci. Technol. 53, 2416–2425 (2019).
CASPubMedGoogle Scholar
Hu, H. et al. Corrosion dynamics of carbon-supported platinum electrocatalysts with metal–carbon interactions revealed by in situ liquid transmission electron microscopy. Nano Lett. 24, 2157–2164 (2024).
CASPubMedGoogle Scholar
Liu, P. et al. Dealloying kinetics of AgAu nanoparticles by in situ liquid-cell scanning transmission electron microscopy. Nano Lett. 20, 1944–1951 (2020).
CASPubMedGoogle Scholar
Vigil, S. A. & Moreno-Hernandez, I. A. Dissolution heterogeneity observed in anisotropic ruthenium dioxide nanocrystals via liquid-phase transmission electron microscopy. J. Am. Chem. Soc. 146, 11133–11140 (2024).
CASPubMedGoogle Scholar
Chen, L. et al. Imaging the kinetics of anisotropic dissolution of bimetallic core–shell nanocubes using graphene liquid cells. Nat. Commun. 11, 3041 (2020).
CASPubMedPubMed CentralGoogle Scholar
Yang, Y. et al. Elucidating cathodic corrosion mechanisms with operando electrochemical transmission electron microscopy. J. Am. Chem. Soc. 144, 15698–15708 (2022).
CASPubMedGoogle Scholar
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Acknowledgements
This work was supported by the U.S. Department of Energy (DOE), Office of Basic Energy Sciences, Division of Materials Sciences and Engineering under Award No. DE-SC0001135. This research used resources from the Center for Functional Nanomaterials, which is a U.S. DOE Office of Science Facility, at Brookhaven National Laboratory under Contract No. DE-SC0012704.
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Department of Mechanical Engineering & Materials Science and Engineering Program, State University of New York at Binghamton, Binghamton, NY, 13902, USA
Zhikang Zhou, Xiaobo Chen, Dongxiang Wu & Guangwen Zhou
School of Materials Science and Engineering, Xiangtan University, 411105, Xiangtan, China
Dingding Zhu, Jianmin Chen & Canying Cai
School of Chemical Sciences, University of Chinese Academy of Sciences, 101408, Beijing, China
Xianhu Sun
Center for Functional Nanomaterials, Brookhaven National Laboratory, Upton, NY, 11973, USA
Meng Li & Judith C. Yang
Department of Chemical and Petroleum Engineering, University of Pittsburgh, Pittsburgh, PA, 15261, USA
Judith C. Yang
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Z.Z.: data collection, analysis, and manuscript preparation; X.C., D.W., D.Z., J.C., X.S.: data collection, analysis; M.L., C.C., J.Y.: analytic tools; G.Z.: project design, manuscript writing.
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Zhou, Z., Chen, X., Wu, D. et al. In situ electron microscopy: atomic-scale dynamics of metal oxidation and corrosion. npj Mater Degrad 9, 28 (2025). https://doi.org/10.1038/s41529-025-00568-9
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DOI:https://doi.org/10.1038/s41529-025-00568-9
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