Advances in Corrosion of High-Temperature Materials: Interfacial Migration and Alloy Design Strategies

Abstract

High-temperature structural materials face severe degradation challenges due to oxidation and corrosion, leading to reduced long-term stability and performance. This review comprehensively examines the interfacial migration mechanisms of reactive elements (REs) such as Ti, Al, and Cr in Ni/Fe-based alloys, emphasizing their role in forming and stabilizing protective oxide layers. We discuss how these oxide layers impede ion migration and mitigate environmental degradation. Key findings highlight the importance of selective oxidation, oxide layer healing, and the integration of novel alloying elements to enhance resistance under ultra-supercritical conditions. Advanced insights into grain boundary engineering, alloy design strategies, and quantum approaches to understanding charge transport at passive interfaces are also presented. These findings provide a foundation for developing next-generation high-temperature alloys with improved degradation resistance tailored to withstand extreme environmental conditions.

1. Introduction

In metallurgy, pure metals have limited applications, and alloying is critical for achieving desired material properties [1]. Often, metals are encouraged to be formed into alloys to obtain better and desired properties of interest. The development of human civilization has been closely linked to the utilization and alloying of metals such as copper, tin, and gold, enabling the creation of materials with superior properties, such as those seen in the Bronze Age. The alloying of Cu and tin yielded a metal stronger than Cu itself and initiated the era of the Bronze Age. Years later, the intimate mixing of Fe in activated charcoal resulted in a much stronger alloy [2], and the ancestor of steel was born. It was only due to the invention of steam engines that hastened the bulk production of steel. Indeed, introducing alloying elements provides the desired property but also brings unavoidable issues. Highly complex microstructures eventually degenerate during in-service conditions, with grain boundary (GB) segregation [3] during high-temperature operations being among the critical issues.

In the modern era of civilization, it would have been hard to imagine life without alloys, as different alloy grades have made their way into every nook of human life. Conventional alloys perform satisfactorily to a temperature range of ~873–888 K and pressure up to 25 MPa, but, as soon as temperature enters the ultra-supercritical and advanced ultra-supercritical stage [4,5,6,7,8], the performance of traditional high-temperature steels starts to degrade drastically. The decrease in performance and properties can be attributed to the affinity of alloys and the alloying elements toward oxidation that corroded at higher temperatures [9]. Corrosion initiation is driven by a range of intriguing factors, including stress-induced corrosion [10], crevice corrosion [11], pitting [12,13], erosion, hot corrosion [14], and erosion-corrosion [15,16,17].

The word “CORROSION” is derived from the Latin word “CORRODERE”, which means ‘gnaw away’. Chemically speaking, it can also be defined as the physiochemical interaction between a metal and its environment, which results in changes in the properties of the metal that may often lead to the function impairment of the material, the environment, or the engineering system of which it forms a part. All elements, alloys, and compounds have a love-at-first-sight relationship with oxygen. Their propensity toward oxygen is even better at higher temperatures. Furthermore, structural and functional alloys of technological interest are inherently unstable; they become involved in oxidative processes and tend to revert to their mineral form, which is considered the most stable form of elements. At higher temperatures, the diffusion of atoms is erratic, and the reactive elements (REs) have a chance to migrate to the surface and look for oxygen-reactive sites. The change in the ratios of alloying elements due to the migration from the lattice sites creates vacancies in the matrix, weakening the bond and, as a result of oxidation, altering the near-surface microstructure, which makes the material prone to environmental degrading factors. The addition of Al or Cr to the alloys forms a stable oxide, Al₂O₃ or Cr₂O₃, respectively, acting as a barrier toward the further oxidation of the alloys in harsh oxidative environments [18]. The product of corrosion, the thermodynamic rate, and the corrosion mechanism is a strong function of variables like environment, pressure, and temperature [19]. The spontaneous growth of an oxide layer on the surface kinetically governs this mechanism. The oxide layer stability is a matter of great concern in relation to the behavior of metals, especially in corrosive [20] and at ultra-high temperatures (1100 °C) and pressure (200 bar) [21,22]. The detachment of this oxide layer at very high temperatures ends the material’s life. Oxidation remains the terrifying and undefeated warrior in the war of alloys against corrosion resistance at high temperatures. Almost all alloys are oxidized above a specific temperature, leading to scaling, an average percentage weight loss (APWL) [23] of material, and a degradation in properties.

2. The Big Question

The big question is as follows: Why do all materials (metals and alloys) corrode? Is there any benefit to undergoing corrosion? What are the protective measures? In addition, it is essential that we answer the key issues that arise from interfacial migration, particularly in the context of high-temperature materials. Can new alloy design strategies alleviate these issues? Specifically, could the use of alloying elements, chemical inhibitors [24,25], certain amino acids, and grain boundary engineering (GBE) [26,27,28,29,30,31,32] provide corrosion resistance even at very high temperatures?

This review examines material degradation issues at high temperatures, where corrosion is a fatal disease that weakens the material and degrades its performance. We also examine the corrosion mechanism and protective measures that will allow alloying at high temperatures. This review also attempts to understand the quantum charge transfer through the passive layers formed by the various oxide formers.

The fact that very few metals, such as gold (Au) and platinum (Pt), remain idle towards even the most potent reactive agents like aqua regia [33], whilst other materials show a strong affinity toward oxidation, leading to corrosion at high temperatures is a vindication that metals and alloys have a strong affinity for oxygen. Their affinity towards oxygen is justified by the fact that their atomic orbitals are incomplete. Metals exist either in their pure metallic form with a zero oxidation state or can be seen interacting with compounds that lower their Gibbs free energy or fill their vacant orbital and exhibit a variable positive oxidation state. Transition metals, in particular, often have varying oxidation states [34] due to an incomplete d-orbital. In a real-world scenario, most metallic elements that are in contact with oxygen form compounds, indicating a strong desire to achieve greater stability in their oxidized forms. For this apparent reason, to procure a metal in its pure metallic state from one of its compounds/ores, it is indispensable to put in energy. The vice versa scenario is also true: when a metal is exposed to its environment, it tends to emancipate this stored energy via the very profound processes of corrosion. This phenomenon of corrosion can be explained as an analogy to precisely what exactly happens with an object suspended at a point above the ground, which is equivalent to the metallic state with zero oxidation states. When allowed to fall freely to reach a stable state, it returns to a position of minimum energy on the ground, which is equivalent to the metal’s oxidized state. In the oxidized states, metals are usually stable.

3. Corrosion Mechanism

To better understand the protective method of corrosion, one needs to know the reaction mechanism of corrosion. The chemistry that governs the corrosion mechanism is stimulated by redox reactions (a chemical reaction in which one element becomes reduced, and, at the other end, another element becomes oxidized, hence the name: reduction–oxidation). The prerequisites of such reactions are the co-existence of an element that is reduced (the oxidizing agent) and another that is oxidized (the metal). Thus, the overall reaction can be split into two partial reactions: reduction and oxidation, as shown in Figure 1.

At the cathode, we have the following:

(Oxidant)/Mⁿ⁺ + ne⁻ → (Reductant)/M (Reduction Process)

(1)

At the anode, we have the following:

(Reductant)/M → (Oxidant)/Mⁿ⁺ + ne⁻ (Oxidation Process)

(2)

In oxidation, the metal loses electrons, and the zone in which this happens is known as the anode (Figure 1). On the reduction side, the oxidizing agent captures the electrons that the metal has shed during oxidation, and the zone in which this happens is the cathode. The heart of the corrosion protection methods lies in controlling the thermodynamical rate at which these reactions proceed. Conversely, the corrosion rate can be slowed down by the increasing the resistance of the oxide layer formed on the surface of metals and alloys and improving its strength, particularly in aqueous and high-temperature environments. The other method that can be deployed to curtail the corrosion mechanism is eliminating the possibility of redox reactions. If possible, conditions can be created to have only anodic or cathodic reactions, but these conditions are ideal and seem more or less a herculean task.

4. Why Does Metal Diffuse Outwards and Oxygen Inward in the Alloy?

Metals in the alloy are appreciably smaller than their non-metallic counterparts. Their cationic nature makes them slightly more mobile than their peers and, therefore, helps in faster diffusion at the interface of an alloy/material. The metal ions move on the surface and have fewer nearest neighbors, or coordination numbers [35], and, as a result, have dangling bonds exposed to the surface. In bulk, electrons move inside the material. Dangling bonds create surface instability and elevate energy states. These bonds must find ways to compensate for the apparent deficiencies by whatever chemical or physical means are available to them. Suppose surface atoms do exist in a state of higher energy than the fully coordinated volume atoms. In that case, the metallic ions at the surface will have a higher propensity towards the reaction than the bulk. This possibility is additionally backed by the fact that metals have more stored energy and want to stabilize them by forming oxides and hydroxides, as they have a negative Gibbs free energy [36,37,38]. Oxidation occurs first at the metal–environment interface, resulting in the formation of a metal scale that acts as a barrier to restrict further oxidation. To sustain the oxidation process, either oxygen must diffuse inwards through the scale to the underlying metal or metal must diffuse outwards through the scale to the surface. Both transfers occur, but the outward diffusion of metal is generally much more rapid than the inward diffusion of oxygen (Figure 2) since the metal ions are appreciably smaller than the oxygen anion and move with much higher mobility.

The creation of vacancies at the metal–oxide interface in the high-temperature atmosphere is a key factor for the detachment of the protective oxide layer, leading to significant material loss. Vacancies grow in size after the considerable loss of material and often can be seen in many systems and are known by many names, such as spalling failure in Fe-Cr-Al systems [39] and thermally grown oxide and bond-coat in thermal barrier coating (TBC) systems [40,41,42], and can often be seen by various characterization techniques, such as SEM and TEM. In the case of materials deployed in high-temperature in-service conditions, the protective oxide layer degenerates and falls off, and it is now evident that the interface has failed completely.

5. Quantum Approach in Understanding the Electron Transfer at Passive Interfaces

Electron transfer (ET) is an ideal framework for developing the theory of corrosion science and understanding corrosion control parameters due to its simplicity in reaction dynamics [43]. It is often critical to understand the corrosion mechanism in layers passivated by corrosion-protecting elements and materials, such as Cr, Al, Ti, Mo, W, etc. Several attempts using first-principles methods to predict reaction rates have proven futile. However, the seminal work of Gurney [44] on quantum electron transport (QET) theory was the first unassailable attempt to elucidate the corrosion mechanism and highlight the significance of Tafel’s law (an exponential dependence of charge transfer to the applied potential). Even after 100 years, Tafel’s law remains highly relevant, evidenced by its widespread use in electrochemistry [45], corrosion science [46], and energy devices [47]. In addition to Tafel’s law, inverse Tafel’s law is also a powerful tool for understanding the dynamics of oxide layer formation on the metal surface. The formation of this oxide layer further resists corrosion by forming a passivation layer on the metal surface during its exposition to high-temperature and environmental conditions. The Tafel equation must be introduced to understand the passive layer’s electron transfer and corrosion mechanism. Tafel formulated his famous equation connecting the overpotential (η) and the current density (j) by the equation given below:

η = a + b log (j)

(3)

where a and b are constants [48] and bears the relationship as follows a = b ln (j_o) and b = RT/zαF (R: universal gas constant, T: temperature in Kelvin, F: Faraday’s constant, z: charge number, j_o: exchange current density, and α: transfer coefficient). The parameters a and b are now commonly known as Tafel’s slope. The nomenclature of an equation in the name of its inventor is quite common, but having a slope in one’s name is rare and remarkable.

This seminal equation was initially applied to the charge transfer reactions occurring under controlled conditions and had nothing to do with the corrosion mechanism whatsoever. Over a period of time, people used this equation in the entire field, wherever electrochemistry was applied. To quote a few lines from the special issue on the centenary of the first publication of Tafel’s equation, Tafel’s equation is not only a remarkable achievement but a remarkable equation itself [49].

According to the QET, transferring an electron or a charge between two states of energy is only feasible when both energy levels are nearly equal. This transfer can be made possible by supplying energy to the molecule. Marcus [50,51,52,53] describes only a slight possibility of electronic orbital overlap in the activated complex, and, often, the electron transfer mechanism should proceed via a “slight overlap” mechanism. Greater detail of the Marcus mode is discussed elsewhere [54]. According to the classical electron transport theory, the amount of current flowing (i) in any redox reaction where QET is occurring is given by the following:

$$i \alpha \iint n\left(E,V\right)N\left(E,d\right)P\left(E,d\right)dEdx$$

(4)

n(E, V): electron density with energy E at the applied voltage V, N(E, d): acceptor species number having energy E at a distance d from the metal surface, and P(E, d): the probability of tunnelling electron from the metal to the acceptor.

The following assumptions have been undertaken to formulate the above equation: (1) the electric field inside the passive layer bears a linear relationship with the applied potential, (2) defects inside the matrix are constant, (3) the oxide film is homogenous and consists of a single phase, and (4) the thickness of the passive layer is independent of the applied potentials. As developed by Schmickler and the group [55], this theory is based on ideal conditions, which are hard to meet in superalloys and materials of interest. Metals, alloys, and materials of high-temperature interest often violate these ideal assumptions. Hence, better formalism has to be adopted to understand the current flow or charge transfer via passive layers and control the menacing corrosion. In a practical scenario, the electric field intensity (ε) is a function of distance (l) and varies with the applied potential (V), as given by the following:

ε = V/l

(5)

This simple expression explains the decrease in ε when the passive layers thicken over time. The thick passive layers resist the further migration of cations from the bulk through the lattice to seek oxygen from the surface. Verwey introduced a temperature-dependent model known as the high field model (HFM) to explain the electron transfer through the passive layer to match the real-world scenario and formulated an non-linear equation:

$$i=A \exp \left(B\mathsf{\epsilon }\right)$$

(6)

i is the current density, and A and B are temperature-dependent constants.

This model assumes that the charge motion within the oxide layers acting as a potential barrier is the rate-determining step. For sufficiently high fields and cations, the conduction current is given by the following:

$$i=2aCv\exp [-\left(W-zFa\epsilon \right)/RT]$$

(7)

a: half-jump distance, C: ion concentration, v: vibration frequency of charge, W: activation energy, z: ionic charge, F: Faraday’s constant; and R and T: universal gas and temperature in Kelvin, respectively.

The outcome of this equation fails to satisfactorily explain the experimental findings of the steady-state achievements of the film thickness or current, and, once again, this theory was discarded, leading to the proposal of Macdonald’s point defect model (PDM). This theory assumes that, in steady-state, passive films, ε is independent of the position of the film in the matrix. This finding has its roots in a high ε > 2 × 10⁶ V m⁻¹, which is sufficiently high and close to the dielectric strength, enabling the creation of electron-hole pair sufficient enough to buffer the medium against any rise or fall in the potential gradient [56]. PDM, like HFM, also suggested a decrease in the current density due to the increased film thickness at the interface but no change in the field strength. From an electronic standpoint, the passive films are analogous to a highly doped metal/n⁺-i-p⁺/solution junction. Depending on the relative concentrations of the donor oxygen/anion and acceptor/cation vacancies, the width of the passive layer increases with the increase in applied potential, maintaining a constant potential gradient. This model resembles the resonance tunnel diode (RTD) [57] and exhibits negative temperature coefficient characteristics. This model closely approximates the experimental condition, such that, at the higher temperature, the passive layers become weaker due to the tendency of the locked charges to conduct electricity, demonstrating the negative temperature coefficient characteristics. In reality, passivation helps gradually increase the oxidation resistance until a certain temperature range. The layer breaks via a spalling failure mechanism at a specific threshold temperature (depending upon individual alloys). The resonance tunneling of charges or electrons proves to be more promising than direct tunneling. The quantum mechanistic approach says there is always a possibility and indeed a probability of an electron tunneling through any potential, but large enough to confine it and restrict it from crossing; it is only a matter of time. Schmickler and Ulstrup [58] postulated a two-electron coherent tunneling theory for metal oxide films. This two-electron transfer mechanism and RTD can conjointly help to understand the diffusion of the charge through passive layers.

Another aspect of the quantum theory addresses the barrier layer as a defective structure characterized by a high density of point defects. The oxygen being a donor and metal cations acting as an acceptor resembles an electrolytic cell capable of redox reaction feasibility in superalloys with Cr, Ni, and Fe as cations, and diffused or ambient oxygen forming a complete cell.

6. Metallic Additions in Alloys to Scale up the Corrosion Resistance

The addition of alloying elements either during alloy formations or the coating of corrosion-resistant elements [59] on precious alloys has been a trend toward the protection of and increase in the longevity of high-performance material. Certain alloys are tailor-made to operate above 923 K and a pressure of 25 MPa or higher and are commonly known as superalloys. The operating temperature of these alloys is raised to increase the energy efficiency and scale back CO₂ emissions in one go in nuclear and thermal power plants, with turbine blades experiencing massive thrust, jet engines to enhance performances, chemical industries for their corrosion resistance, photochemical industries, and tools and dies for extremely hot working sections of metals. All these domains require corrosion-resistance elements to be inherently alloyed during the alloy formation. To meet these challenges, several classes of materials are explored; however, only Ni-based superalloys have shown the potential to meet these criteria [60].

Superalloys are multi-alloyed elements; hence, their reaction with oxygen is often very complex [61]. However, these alloys are developed to achieve resistance to oxidation by utilizing the concept of selective oxidation [62]. Eventually, the oxidation processes in superalloys are not as complex as one might initially presume. The selective oxidation phenomenon in superalloys, which enables oxidation resistance, involves oxidizing only a specific element in the alloy and relying on its oxide layer for protection. For the feasibility of this process, the oxide layer formed must perfectly cover the entire surface of the alloy, as it serves as the medium through which the diffusion of reactants (i.e., the selectively oxidized element and oxygen) occurs at relatively slow rates. However, there are very few elements capable of meeting these stringent conditions. The only elements capable of satisfying these conditions are silicon, aluminum, titanium, and chromium. The presence of silicon at a sufficient level allows selective oxidation conflicts with the high melting point requirements of superalloys; therefore, silicon is discarded in favor of selective oxidation for oxidation resistance. The available techniques to induce selective oxidation at lower concentrations are reported in the literature [63,64]. Singh et al. report the presence of silicon intermetallic compounds in high-temperature superalloy 617, but its origin in the matrix and its functionality are still under investigation [6]. Superalloys with Ni, Co, and Fe utilize the elements Cr or Al for selective oxidation and are often referred to as Cr₂O₃ and Al₂O₃ formers and act as a guard against oxidation degradation [6]. The successful selective oxidation process depends upon the phases that need to be selectively oxidized, and they must have the elements to be selectively oxidized at a concentration sufficient enough to form scales of Cr₂O₃ and Al₂O₃ to form on alloys, and phases that do not contain these are preferentially attacked at high-temperatures. Nair et al. [65] have reported an increase in the corrosion resistance of the Al_0.1CoCrFeNi high-entropy alloy (HEA). HEAs are a special class of alloys that have equiatomic combinations of 5 or more elements [66]. These HEAs are finding wider applications in almost all energy conversion/storage systems due to their tailor-made physical properties to suit a particular task [67,68]. These HEAs offer a wider space for chemical combinations from the periodic table, allowing the entropic contribution to overcome the enthalpic contribution and stabilizing the solid solution [69]. These HEAs are better-suited to corrosion and offer greater resistance by bringing passive film stability [70]. A study primarily composed of aluminum (Al), cobalt (Co), chromium (Cr), iron (Fe), and nickel (Ni), commonly referred to as Al(Co)CrFeNi HEAs, has been carried out by the Ulrike Hecht group [71]. The corrosion activity of this material is investigated in a 3.5% NaCl solution. The results indicate that the minor addition of Mo can dramatically influence the corrosion resistance of this alloy by restricting pit formation [72]. It was also found that the face-centered cubic (FCC) phase in equimolar Al(Co)CrFeNi is more resistant to corrosion, demonstrating a more noble corrosion potential. In addition, a lower corrosion current density and corrosion rate than the body-centered cubic (BCC) phase counterpart is an additional benefit. Although the exact mechanism of the corrosion prevention of Mo in the presence of Cr in steels is still unestablished, it is believed that both Mo and Cr form a barrier against Cl⁻ ions [73]. Mo has the ability to form Mo-oxo-chiro, a stable and soluble complex, along with insoluble chlorides and oxide chlorides on the surface of the steel, thereby preventing corrosion. Critically, the ratio of Al/Cr plays another dominant role in corrosion inhibition in this alloy. Corrosion resistance in high-temperature alloys can be further improved by the addition of W due to their improved mechanical properties. Another study on the role of HEA coatings (HEACs) using different compositions of Mo in FeNiCoCrMo_x (x = 0, 0.15, 0.20, 0.25) on 316-grade stainless steel (SS) brought some interesting results [74]. The microhardness increased in the order of 70.1% (Mo0), 77.0% (Mo0.15), 84.9% (Mo0.20), and 90.5% (Mo0.25), as shown in Figure 3a. The average friction coefficient of Mo0.25 is among the lowest, indicating that it has the best wear resistance (Figure 3b). The wear rates decreased dramatically in the order Mo0.15 (11.1%), Mo0.20 (27.8%), Mo0.25 (38.9%), with respect to Mo0, as shown in Figure 3c. The results of the electrochemical tests indicated that HEACs exhibited intergranular corrosion as their predominant type of corrosion, and all HEACs demonstrated superior corrosion resistance compared to both 304 SS and 316 SS. Notably, among the HEAC samples, Mo0.20 HEACs exhibited the highest level of corrosion resistance, which can be attributed to the beneficial influence of MoO₃ on the formation of a protective Cr₂O₃ passivation film. The increase in hardness for Mo0.25 is due to the strengthening of solid solutions offered by Mo. Rodriguez et al. investigated the corrosion effect of an Mo addition to CoCrFeNi₂ and CoCrFeNi₂Mo_0.25 under 3.5 wt % NaCl [73]. The cathodic current density, as obtained by the polarization curve, shows a straight line, indicating control over electron transfer. The presence of Mo in the alloy makes a passivation layer and offers a high electron transfer resistance. Mo also offers transpassivity, achieved through oxidation at higher potentials [75]. The presence of Cr in the alloys undoubtedly offers the first guard against corrosion, but the addition of Mo provides stability to the protective layer prepared by the Cr. In addition to this, Mo also interacts with S in the alloys and provides global repairment to the local weak spots [76]. The presence of Mo in the alloys also administers the small potential difference between the re-passivation potentials (E_rep) and breakdown potential (E_bre), offering resistance to corrosive environments and protecting the material from pitting corrosion in the NaCl solution.

As an alloying element, Al always plays a dominant positive role in corrosion inhibition in high-temperature alloys. Al is thus regarded as another vital metal addition that significantly changes the corrosion properties of the alloys and has a wider application in energy materials [77,78]. In the presence of oxygen, Al tends to combine with it to form alumina. Once there is a sufficient amount of Al, alumina acts as a protective barrier, preventing the permeation and penetration of carbon. However, when the Al content is relatively low, a situation arises where an excess segregation of the Al-rich and Cr-rich phases occurs. This excessive segregation ultimately leads to a deterioration in corrosion resistance. Thus, its study become essential in order to understand the protection mechanism. Recently, the role of an Al addition in a series of Al_x(TiZrHfNb)_100−x (x = 0, 3, 5, 7, 12 at%) refractory HEA (RHEA) was investigated [79]. The addition of Al significantly improves the mechanical properties of the alloys. It was found that the addition of Al brings a more compact BCC structure to the alloys. The yield strength (YS) results indicate that YS-Al bears a linear relationship (Figure 3d). YS for Al-0% is 331.3 MPa, while, for Al-12%, it jumps by 3.3 times (1097.8 MPa). The results of the wear rate and Vickers hardness of RHEA with a varying composition of Al is shown in Figure 3e. It is obvious that, with the increase in Al content, the hardness values increase and wear resistance improves, reflecting the lower wear rate. The gradual increase in hardness is due to the significant localized lattice distortion arising from the atomic size difference between the various components, which brings improved solution hardening. It is evident that the atomic radii of Al are similar to those of Nb and Ti, but they are significantly smaller than those of Hf and Zr. The δ values for Al-0% to Al-12% alloys are 5.733, 5.794, 5.830, 5.863, and 5.929, respectively, and are less than 6.6, indicating that the current BCC phase is essentially a disordered solid solution alloy. Eventually, the increasing δ values marks a trend in the increase in hardness, and the strong bonds between Al and the constituent transition metals also contribute significantly to the hardening effect in compliance with a previous report [80]. Therefore, the excellent hardness can be attributed to both the strong lattice distortion and the robust bonds between Al and the alloy constituents. Similarly, to reveal the oxidation behavior of alloys, normally a power law is used ($∆m=k_1{\mathrm{t}}^{\mathrm{n}}$); here, Δm is the mass gain/area, k₁ is the oxidation rate constant, t is the exposure time (s), and n is the time exponents. The oxidation data for this alloy are shown in Figure 3f. The n values for the different Al doping wt % are Al-5% (0.91), Al-7% (0.85), and Al-12% (0.78). As n values lie between/above 0.5 (n < 0.5: parabolic) and are less than/approaching close to 1 (n = 1: linear), this indicates that the oxidation behavior of this alloy is in between parabolic and linear. Al-12% displays a minimum oxidation rate, reflecting the wear mechanism to be oxidative at elevated temperatures for this alloy.

Figure 3. (a) Microhardness profile for 316L-grade SS and various Mo composition in HEACs. (b) Obtained average friction coefficient. (c) Histogram plots showing the specific wear rates. Reproduced with permission from [74], Elsevier, 2022. (d) The linear fit between YS and Al concentration for Al_x(TiZrHfNb)_100−x. (e) Profile showing wear rate and hardness variation with respect to Al concentration. (f) Logarithm plot of mass change against time (s) for Al-5/7/12 wt %. Reproduced with permission from [79], Elsevier, 2022.

Figure 3. (a) Microhardness profile for 316L-grade SS and various Mo composition in HEACs. (b) Obtained average friction coefficient. (c) Histogram plots showing the specific wear rates. Reproduced with permission from [74], Elsevier, 2022. (d) The linear fit between YS and Al concentration for Al_x(TiZrHfNb)_100−x. (e) Profile showing wear rate and hardness variation with respect to Al concentration. (f) Logarithm plot of mass change against time (s) for Al-5/7/12 wt %. Reproduced with permission from [79], Elsevier, 2022.

Kim’s group [81] studied the effect of a Cu addition in a virgin Al ingot alloy comprising varying wt % of Cu in an artificial acid rain environment containing 200 ppm of Cl ions. The results obtained by time-of-flight secondary ion mass spectrometry (ToF-SIMS) with low Cu concentrations (0.005, 0.01, and 0.03 wt % Cu) reveal that only 0.03 wt % Cu with Al gives reliable information. For the other compositions, the intensity is either very low (0.01 wt % Cu) or it was difficult to detect (0.005 wt % Cu). The detection of Cu on the Al surface is indicated by the intermetallic Al₂Cu phase and has been supported by various research groups as well [82]. It is believed that the higher Cu concentration causes more precipitation of the Al₂Cu phase [83] and increases the cathodic current density owing to its higher electron transfer in 0.03 wt % Cu. The high cathodic current density of 0.03 wt % Cu is a clear indication that pitting corrosion will be dominant in the matrix, and Al with this composition of Cu has the least corrosion resistance. Thus, care must be taken while alloying an addition of Cu to high-temperature material, and the wt % must be controlled around 0.01 wt %. Yan et al. [84] studied the effect of alloying elements on corrosion behavior in Co–Al–W material. Adding the elements, for instance, Cr, Si, and Al, to this alloy improved the corrosion resistance of this alloy. The order in which these oxides resist oxidation is given in decreasing order: 10Cr > 20Fe > 1-Si > 2Ta > Base > 20Ni > 2Ti > 2Mo > 2V > 6Ni–4V. The possible mechanism of oxidation protection is the formation of three types of oxides of CoO at high temperatures. Firstly, we have the formation of a stable cubic rock type of CoO above 1173 K with a lattice constant of 4.26 Å [85], and a spinel-type Co₃O₄ with a lattice constant of 8.15 Å [86], while, in other cases, a thermodynamically unstable oxide Co₂O₃ also forms under controlled conditions [87]. The Co₃O₄ oxide was the first to be formed on the outer surface of the alloy; the inward diffusion of oxygen forms it, and the rich Co content in the alloy promotes the upward diffusion of Co, resulting in the formation of a stable Co₃O₄. The decrease in the partial pressure of oxygen and the depletion of Co from the base metal initiated the formation of a second defensive layer composed of Al₂O₃. The continuous formation of Al₂O₃ depletes the percentage of Al locally and helps in initiating the precipitations of Co₃W. This reduction in major alloying elements promotes the porous-type layered structures, and, hence, coarsening is avoided or impeded, and, therefore, only fine-grained structures are present. The presence of fine-grained structures helps in strengthening the material at high temperatures. A substantial reduction in the partial pressure of oxygen induces the formation of CrO₂ but this oxide becomes thermodynamically unstable.

The presence of Cr in the alloys acts as an oxygen getter and further inhibits oxygen diffusion into alloys. Mo and W also offer selective oxidation by acting as an oxygen getter, which causes the diffusion of Al in alloys to decrease considerably. In this way, the presence of Cr, Al, Mo, and W [88], along with the presence of some oxide particles, helps in promoting the selective oxidations process and preserves a considerable number of precious alloys operating at very high temperatures.

7. Role of Texture Control and Surface Energy in Corrosion Inhibition

The microstructural texture in alloys plays a significant role in determining the physical properties of the alloys during their service conditions. Surface textures in metal and alloys are introduced during the mechanical deformation or material processing, for instance, forging, rolling, and drawing, and these significantly affect their physical properties, including the wear and friction characteristics. The corrosion resistance of the material exhibits a strong relationship with the crystallographic orientation and crystallite interfaces at the weld junctions [89,90]. Thus, it becomes imperative to understand precisely the role of texture in corrosion control. There is yet another important factor that plays a significant role in controlling the corrosion of alloys at high temperatures—surface energy. It is established that the activation energy required for the dissolution of the densely packed surface is relatively higher than that of the loosely packed surface. However, the opposite rule governs the surface energy, where a densely packed surface has a lower surface energy (SE) than that of a loosely packed surface [89]. Benefitting from these two established facts, the high-temperature alloys are designed to have a low surface energy to obstruct corrosion for a longer time. It becomes even more important when materials are joined together via welding. Often, a stress gradient is generated at the interface of the base and weld metal, which, later, during service conditions, becomes an active site for corrosion-induced failure. A primitive study on steel found that the carbon-steel texture changes after corrosive and wear resistance tests [91]. The orientation density function (ODF) of the alloy showed different trends under the same applied stress of 9.6 N to corrosive and wear tests in corrosive environments. It was noticed that the Goss and brass texture in corrosion wear situations completely disappeared contrary to the dry wear test environments. The difference in the ODF is ascribed to the reduced formation of shear/friction forces in the presence of NaCl, which acts as a lubricant. It was further found that weight loss in the case of corrosive tests was much lower than that of dry wear corrosion tests, the reason being, again, the lubricating effect of NaCl. In another study, it was found that sulfide stress corrosion cracking (SSCC) and strain rate tests bear a close relationship with the surface texture. Here, it is necessary to mention that there are certain favorable planes for the dislocation motion during crack initiation and crack propagation, arising due to the SSCC/hydrogen-induced cracking (HIC) often encountered in high-temperature alloys. Thus, it is required that the easy motion of these favorable planes should be restricted during the crack initiations. Failure arising from the SSCC includes two steps: crack initiation and crack propagation, dependent upon the crystallographic texture [92]. The crack propagations and crack initiation follow the following orders of planes (100) > (110) > (111) and (110) > (100) > (111), respectively [93]. In addition to this, the surface normal to the transverse direction (TD) having a family of planes (552) has more corrosion-resistant potential and the lowest corrosion rate compared to the other two directions, namely, the rolling and normal direction abbreviated as (RD) and (ND), respectively. The improved corrosion resistance is shared partially with grain boundaries (GBs) and texture effects. The importance of texture was once again established from the study carried out on steels alloyed with Sb and Cu in extremely corrosive environments of pH 0.3, having a % vol fraction of H₂SO₄ and HCl at 16.9 and 0.35, respectively [94]. A comparison of hot- and cold-rolled steels showed that the preferential orientation of planes with {111} along with {101} and {001} remain dominant, as opposed to {001}-dominated planes, along with {101} and {111} in the cold-rolled steels. The low corrosion resistance of cold-rolled steels along the direction {100} is ascribed to its high surface energy, thus favoring the faster dissolution of atoms.

It should be understood that cold rolling during alloy processing brings deformation twins and dislocation arrays inside the alloys, which are known as one-dimensional crystal defects. The presence of crystal defects inside the matrix allows the faster diffusion of Cr/Mo and easier carbide nucleation assisted by a low free-energy barrier. These easier diffusions and nucleation act as different phase formations than the matrix and introduce strain inside the material matrix. During external stress, these sites act as preferential sites for crack initiation, leading to material catastrophic failure.

The other factor that significantly controls the corrosion resistance of a material is surface energy, and the phenomenon that needs to be addressed is hydrophobicity. Hydrophobicity does not allow water to stay for longer on metal surfaces by reducing the metal–water contact area, thus protecting the material from corrosion. For corrosion protection in high temperatures, the surface must have minimum contact with the corrosive environments.

In a study carried out on 316 SS, a superhydrophobic (SHP) surface was achieved through a novel one-step simultaneous H₂O₂ and acid (HF) etching, along with perfluorooctyl trichlorosilane (PFOS) modification, which significantly improved the corrosion properties of the material [95]. This study accounted for several factors, such as varying the molar ratio of H₂O₂/HF, reaction time, and mass ratio of PFOS/HF, that could influence the hydrophobic surface. As a result, a maximum water contact angle (WCA) of 161.78° and a minimum slip angle (SA) of 1.94° were achieved (Figure 4a). Although a downward and upward trend were observed with an increase in peeling time for WCA and SA, respectively, the results were much better than those of several other techniques reported earlier [96]. To reveal its corrosion properties, polarization curves were obtained against bare SS (Figure 4b). The result of the corrosion inhibition rate for SHP-SS was much higher, at 83.5%, against some of the most complicated techniques reported earlier [97]. The corrosion inhibition parameters such as the corrosion potential (E_corr: V), the current (i_corr: A cm⁻²), and the corrosion rate (V_corr: mm/A) for bare SS/SHP-SS are −0.2265/−0.1088, 2.400 × 10⁻⁶/3.965 × 10⁻⁷, and 2.805 × 10⁻²/4.652 × 10⁻³, respectively. These values indicate that the corrosion resistance of SHP-SS has improved remarkably. The excellent SHP 316 SS is attributed to the creation of a micro-nano structure achieved through acid etching, coupled with the successful grafting of PFOS onto the metal surface with the assistance of H₂O₂ (Figure 4c). Furthermore, this hydrophobicity of SS was maintained even after 3 months of continuous exposure to the ambient air. Another study found that using a 515 nm pulsed laser with a short duration of 10 picoseconds, followed by treatment with a 0.01 mol/L solution of stearic acid, produced a superhydrophobic coating on AISI 420 SS, achieving a contact angle of up to 163° [98]. Although this method initially produced corrosion-tolerant surfaces, extended immersion in a 0.5 M NaCl solution for 30 days made it difficult to confirm consistent corrosion resistance [99]. To further improve the corrosion resistance, Wang et al. [100] applied a modified treatment technique in 1095 carbon steel with a 1064 nm laser and a 20 ns pulse duration. To increase the treatment speed, they used a relatively large distance of 600 µm between laser passes, with a laser spot size of around 50 µm. This process created high-edged ditches on the surface. Three types of surface textures were then hydrophobized with perfluorooctylsilane, all achieving superhydrophobic states with varying contact angles. The coating with the highest contact angle of ~160° demonstrated the best corrosion resistance, showing a corrosion current of 2.9 × 10⁻⁸ A/cm² compared to 5.4 × 10⁻⁷ A/cm² for untreated steel.

Corrosion protection through the hydrophobic surfaces can be understood in terms of trapping air inside their structures due to their hierarchical structures that allow them to put a restriction on the corrosive elements/ions to directly interact with the alloy surfaces [101]. Another aspect of understanding the corrosion protection of hydrophobic surfaces is the development of the Laplace pressure built on superhydrophobic surfaces with contact angles > 150°.

A recent survey on hydrophobic techniques for protecting commercial alloys highlighted that many studies claiming enhanced protection often omit negative results, or present them only in comparison to successful outcomes [99]. This mindset goes against the true spirit of scientific progress, where both positive and negative results are valuable. Such bias is unfortunate, as it prevents the community from learning and leads to repeated mistakes. Although a detailed study dedicated to hydrophobic protection is not the aim of this review, we summarize (

Table 1. Summary chart of superhydrophobic coatings.

Alloy	Treatment Method	Hydrophobic Agent	Contact Angle	Corrosion Inhibition Efficiency (CIE), %	Coating Corrosion Current (Ic), A/cm2	Bare Steel Corrosion Current (I₀), A/cm2	Key Observations	Ref.
316SS	H2O2, HF etching + PFOS modification	PFOS	161.78	83.5	3.965 × 10−7	2.4 × 10−6	Achieved high corrosion resistance and maintained hydrophobicity for 3 months in ambient.	[95]
AISI 420SS	515 nm pulsed laser + stearic acid	Stearic acid	163				Initial corrosion tolerance; reduced after 30 days in 0.5 M NaCl.	[98]
1095 Carbon steel	1064 nm laser with 20 ns pulse	Perfluorooctylsilane	160	94.6	2.9 × 10−8	5.4 × 10−7	Achieved superhydrophobicity with high corrosion resistance.	[100]

) quickly some of the notable findings and request readers to read other reviews in this direction [99].

The role of crystal planes in corrosion control cannot be ignored in alloys operating at high temperatures. Recently, a study conducted upon DD5 (a Ni-based single crystal) on different crystal planes indicate that the corrosion rate follows the order (011) < (001) < (111), with the passive films growing faster and denser on the planes having higher surface energy [102]. The enhanced performance of the (011) plane is ascribed to the presence of Cr₂O₃ and enrichment of MoO₃ in the passive layer that improves the corrosion resistance. Before moving further, it is essential to mention here that, in metal and alloys, the corrosion behavior depends on the orientation of the crystal planes, with a simple and straightforward chemistry. However, this could be more complex in Ni-based superalloys, particularly with the passive nature of films and, often, dozens of alloying elements. Furthermore, the passive behavior varying with different planes make the corrosion analysis even more complex, as has been reported in superalloy 738 [103]. This complex behavior is due to various proportions of Cr₂O₃, TiO₂, and Ni(OH)₂ passive films on differently oriented surfaces inside the alloy. Coming back to the results and discussions on the DD5 superalloy with its three crystal planes (Figure 4d), the polarization curves indicate that the current densities of all the planes rapidly increase, initially reflecting the active dissolution of the surfaces (Figure 4e). However, the current gradually declined as the potential went up due to the formation of passive films. A smaller current (inset) value reflects a better corrosion property of (011). An additional polarization plot of current vs. testing time (s) exhibits a smoother (Figure 4f and inset) current profile for the (001) and (011) planes compared to (111). Here, the current spikes show the pitting corrosion event occurs on the planes due to the localized passive layer breakdown and re-passivation of the active sites on the alloy surfaces. Evidently, the (111) surface with more current spikes is more prone to corrosion than the other two. It is conclusive now that the barrier layer growth is more facile on planes with the highest surface energy and lowest atomic packing density, while the closed packed plane (111) has the lowest protection ability. The interfacial energy of the plane can also be tuned by external doping such as boron, aluminum, silicon, and so on in trace amounts. In a recent study, boron substitution was made in FCC Fe (001)/Cr₂₃C₆ (001) to alter the interfacial strength of the precipitate and the matrix (“Boron substitution induced FCC Fe/Cr₂₃C₆ interfacial strengthening: An ab initio study”). DFT calculations reveal that boron substitution has a significant impact on the interfacial energy and the atomic arrangements at the interface by bringing increased covalent bonding. The increased covalency ensures more ordered structures at the interface, thereby making the diffusion process less favorable energetically. Eventually, precipitates like NbC, Cr₂₃C₆, other phases like sigma and μ-phases, and other microstructural evolutions are greatly affected, which play a dominant role during creep activities. A reduction in the interfacial energy by (~0.09 J m⁻²) after boron doping (0.29 J m⁻²) at the interface of the matrix and precipitate play a dominant role in creep retardation and improving interfacial strengthening. Several sites were selected, and the results obtained in Figure 4g clearly indicate that interface energy optimization is crucial for alloy design and optimizations. The intentional addition of boron at A-terminated Cr₂₃C₆ sites significantly reduces the interfacial energy from 0.38 to 0.29 J m⁻². To make a clear difference in interfacial energy, it is necessary that we consider local arrangements. The side view of the most symmetrical (001) interface between the FCC Fe and Cr₂₃C₆ sites is shown here (Figure 4h). Fe atoms have a covalent bonding with C, and the mostly metallic ones with Fe/Cr are projected onto the Cr-C sublattice. This mixed interface allows a significant charge transfer of Fe to C atoms. With B-doping, short-range interface strengthening is manifested, and strong covalent bonding is established between Fe-B, thus stabilizing the interface disordering. Although this study brings several insights on boron-doping to stabilize the interface, an extension to this study considering different layers (second, third, fourth, and so on) could be an exciting field and can bring a wealth of novel information.

Figure 4. (a) Water contact angle (WCA) and slip angle (SA) of SS and SHP-SS. (b) Polarization curve for SS (raw material) and SHP-SS (SS5; here, 5 indicates sample number with H₂O₂/HF ratio in 0.2/1, reaction time (50 min), and mass ratio of PFOS/HF in 1/100). (c) SEM image of SS5. Reproduced with permission from [95], Elsevier, 2022. (d) Schematic illustration of DD5 superalloy with different crystallographic planes. (e) Potentiodynamic polarization curve in 3.5 wt % NaCl solution. (f) A potentiostatic polarization curve obtained in 3.5 wt % NaCl solution for DD5 superalloy with different planes. Reproduced with permission from [102], Elsevier, 2022. (g) Various chemical energies for FCC Fe/Cr₂₃C₆ interface with and without B-substitution of carbon atoms. Here, A1–A4 and B1–B3 signify the boron atom replacing carbon atoms in the first, second, third, and fourth layers of Cr₂₃C₆ from A- and B-termination, respectively. (h) Side view of highly symmetrical (001) interface of FCC Fe/Cr₂₃C₆. Reproduced with permission from [104], Elsevier, 2023.

Figure 4. (a) Water contact angle (WCA) and slip angle (SA) of SS and SHP-SS. (b) Polarization curve for SS (raw material) and SHP-SS (SS5; here, 5 indicates sample number with H₂O₂/HF ratio in 0.2/1, reaction time (50 min), and mass ratio of PFOS/HF in 1/100). (c) SEM image of SS5. Reproduced with permission from [95], Elsevier, 2022. (d) Schematic illustration of DD5 superalloy with different crystallographic planes. (e) Potentiodynamic polarization curve in 3.5 wt % NaCl solution. (f) A potentiostatic polarization curve obtained in 3.5 wt % NaCl solution for DD5 superalloy with different planes. Reproduced with permission from [102], Elsevier, 2022. (g) Various chemical energies for FCC Fe/Cr₂₃C₆ interface with and without B-substitution of carbon atoms. Here, A1–A4 and B1–B3 signify the boron atom replacing carbon atoms in the first, second, third, and fourth layers of Cr₂₃C₆ from A- and B-termination, respectively. (h) Side view of highly symmetrical (001) interface of FCC Fe/Cr₂₃C₆. Reproduced with permission from [104], Elsevier, 2023.

8. Atomistic Study of Corrosion

In this section, detailed insights into the corrosion mechanism are presented, bringing the atomic level considerations into focus. A clear picture of corrosion and its initiation in its very early stage of triggering will not only help us in understanding the corrosion propagation over time but can also assist in bringing about delayed corrosion by applying pre-measures and controls to metal and alloys surfaces. Often, metal surfaces are coated with thin layers of protective oxide layers or may be found coated with an inorganic/organic coated layer that continually interacts with the aggressive/corrosive environments [105]. This section will focus on the interaction of the surfaces with the aqueous environments. To know more detailed aspects of the corrosion mechanism, its modelling, and various methods, the readers are requested to consult refs. [105,106,107].

When metal surfaces are exposed to aqueous corrosive environments, the polarized metal environments cause water dissociations, leading to the formation of OH⁻ ions. These sporadic oxide ions start vehemently reacting with metal atoms, and, in the process, the metal/water interface is modified completely, resulting in the formation of the passive anodic oxide layer. Here, we will discuss the interactions of these OH⁻ ions with respect to the Ni surface as most of the high-temperature alloys are made up of Ni-based alloys or at least contain Ni as a dominant element. On the Ni, the oxide surface also has a faceted structure similar to those obtained on Cu [108] and Ag surfaces. This faceted structure indicates that there exists a slightly tilted epitaxy between the preventive oxide layers of the NiO and Ni (111) lattice. The tilt between the Ni (111) and NiO surface has been measured to be about 3.3° [109]. Interestingly, the NiO (111) [1,2,3,4,5,6,7,8,9,10] grows with an antiparallel epitaxy on the Ni (111) [-110] substrate. A hexagonal-terrace-type lattice of 0.3 ± 0.02 nm is grown on the interface of the NiO and Ni substrate [110,111]. The density functional theory (DFT) confirms the preferential orientation on the interface is due to stability achievements by the adsorption of one monolayer of OH⁻ ions/groups [112,113] and is ascribed to the presence of dissociated water molecules at the interface, which favors the nucleation and growth mechanism. The major hurdles in impeding corrosion lie in the weakening of passivating films grown on the substrate. Even when the substrate is a well-grown single-crystal, the passivating films turn towards having a polycrystalline nature with plenty of grains and grain boundaries. From atomic force microscopy (AFM) and scanning tunnelling microscopy (STM) measurements, it is confirmed that, to impede corrosion completely or delay corrosion, the thickness of the grown passive layer should be around ~2.0 nm [110,114,115]. These GBs separating the oxide grains play a significant role in the breakdown of passive layers and initiate localized corrosion [114,116,117] because GB themselves are considered to be 2D defects [118]. To further reveal the mechanism of corrosion inhibition in Ni-based alloys, a study was conducted considering Cr and other impurities present in the alloy [119]. In this study, the Ni (111) surface is considered with an aim to reveal its resistance to dissolution and the segregation behavior of Cr in the presence of other co-doped elements. The research findings reveal that impurities S, P, O, and H tend to be preferentially trapped near the surface, while Cr exhibits a uniform distribution within the Ni crystal, influencing the segregation behavior of impurities S and P towards the surface and causing impurities N and O to shift towards the subsurface (Figure 5a). The study also observes the formation of near-surface Cr nitrides, possibly Cr₂N, and highlights the beneficial impact of introducing Cr on the structural stability of the Ni (111) surface, safeguarding it against corrosion in the presence of impurities. The investigation offers valuable microscopic insights into the creation of a Cr-depleted zone, a phenomenon associated with the local corrosion of the Ni alloy surface. Furthermore, the theoretical calculations provide explanations that different calculation methods bring different surface energy values but are within the error limit. Due to the lower surface energy value (γ_s) of Ni (111), there are significant improvements in surface stability due to the solid solution attributes of Cr. While the surface stability of the impurities elements follows N > O > S > P (Figure 5b), Cr significantly lowered the O stability over the Ni (111) surface, increased the instability of S and P, and enhanced the permeating on the Ni surface. Additionally, a discussion on segregation energy reveals that several impurities behave differently when it comes to GB segregation. Figure 5c showcases various amounts of segregation energy of different impurities (a positive and negative segregation energy means an impurity likely occupies the Ni surface and the GB, respectively). In summary, Cr enhances the Ni (111) surface structural stability and protects it from corrosion in the presence of impurities.

9. Corrosion and Sensitization Control by Grain Boundary Engineering

GBs are essentially two-dimensional (2D) lattice defects or interfaces where two neighboring grains contact and join each other on the atomistic scale, which generally exists in almost all types of polycrystalline materials, whether metallic or non-metallic and irrespective of the crystal structure (cubic or non-cubic) [120]. After the pioneering work by Watanabe [121] and Randle [122] on GBE, a resurgence of interest in the field of the old peculiar question of corrosion was instigated to find the solution near and around the atomic scale and GB. It will not be a metaphor to say that no other aspect affects the bulk properties of material more than GBs do. Of the many features influenced by GBE, for instance, intergranular brittleness [123,124], sensitization [125], oxidation-induced intergranular brittleness [126], segregation-induced brittleness [127], abnormal grain growth, heterogeneous microstructure [128,129], and corrosion control, we discuss only the phenomenon of corrosion control by GBE here.

To explain the importance of GBE on the materials’ corrosion protection, various research groups have undertaken significant efforts [130,131,132]. The alloys that are widely deployed in the high-temperature zone in marine environments, such as naval ships, pressure vessels, and aquatic hulls, are more prone to degrading due to corrosion. Thus, significant efforts must be taken to strengthen the GBs by judiciously selecting the alloying elements. The Al-Mg 5xxx series of alloys are frequently used in marine environments owing to their excellent high strength, weldability, and favorable corrosion resistance. In general, the strengthening in these alloys is commonly achieved via a solid solution, oxide dispersion, and work hardening mechanism [133,134]. The susceptibility to sensitization and stress-induced corrosion are the certain drawbacks to these alloys. During sensitization, Mg atoms diffuse toward the grain boundary (GB) and form a β-phase Al₃Mg₂. The formation of the β phase proceeds via the following path:

Solid solution → Guinier-Preston (GP) zones → β″ → β′ → β

(8)

where β″ and β′ are metastable phases. The intergranular β phase corrodes first in almost all environmental conditions, leading to stress corrosion and intergranular corrosion cracking [135,136]. Reports in this direction have tried to demonstrate that the precipitation of β-phase kinetics is governed by many factors, such as mechanical processing [137,138], chemical composition [139,140], thermo-mechanical treatments [141,142], and high-temperature exposure [143], and very little has been carried out by grain boundary engineering [144,145]. To completely describe a GB, five variables are needed [146]: one variable defines a misorientation angle, two variables are engaged in representing the misorientation axis, and the other remaining two variables characterize the GB plane orientation [125]. The literature has tried to investigate which of these five parameters plays a significant role in the sensitization and corrosion mechanism [147,148,149], but this remains a ground of open debate. For instance, Homer et al. [147] have stressed the GB plane orientation and GB misorientation as essential factors in determining the properties of a polycrystalline material; another effort by Davenport et al. [148] emphasized that the degree of precipitation and acid attack susceptibility for a boundary is related to the crystallographic misorientation. In contrast, Kaigorodova et al. [149] in his work explained that precipitation at the GB existed with a low-angle GB misorientation (5–10°) rather than high-angle GBs, contradicting the earlier prevalent idea that β precipitation was very much suited to high-angle boundaries [150]. This contradictory viewpoint is a pathway for future research in this direction. For our interest, we approach the GB and coincident site lattice (CSL) boundaries to improve the sensitization and corrosion properties.

Brandon’s study [151] investigated the boundaries in polycrystalline material and came up with the finding that not all boundaries are exact CSL boundaries, so it is customary to divide boundaries into special GBs and general boundaries. CSLs have a preferred notation (∑), with the usual meaning being to define the degree of coincidence sites at a GB. The prime motive behind introducing the concept of “grain boundary design and control” is to improve the bulk properties of polycrystalline materials by augmenting the percentage of ‘special’ GBs. General GBs can be tailored to adopt the CSL boundary, also termed as special boundaries. However, not all boundaries can achieve CSL. A twist in the tale of CSL is that not all CSL interfaces have an improved influence on bulk properties. Special boundaries, with low ∑ boundaries, have only been successful in improving the properties at bulk in comparison to their high-∑ interfaces. The ∑ value denotes the fractional number of special boundaries inside the crystal/matrix.

10. Effect of Grain Misorientation

To study the effect of GB parameters on sensitization and corrosion, the Al-Mg alloy is etched with 10% H₃PO₄, and the SEM micrograph is shown in Figure 6a,b [125]. The GBs are numbered from GB1 to GB56, as shown in Figure 6b. High angles (>15°) are marked by black, and low angles (≤15°) by yellow lines denoting the boundaries, and a rectangular mark is given to ∑13b GB. To further categorize the 56 GB surface on the etching behavior, we divide them into three major groups, namely, (a) fully etched boundaries for a continuous attack, (2) partly etched surface nomenclature for a discontinuous attack, and (3) no etch boundaries for no attack. The length % of attack/non-attack for different misorientation angles is shown in Figure 6c. When the misorientation angle is <10°, all the GBs show very good immunity toward attack. With the misorientation angle >15°, the % etched length far exceeds 95%. GBs with a misorientation angle between 10–15° occupy 70% of the etched length. It is now evident how high the misorientation angle plays in the precipitation and etching behavior of alloys. It is also quite clear that more precipitates form in high-angle GBs as compared to low-angle GBs [148]. Al₃Mg₂ precipitates formed at the high-angle GBs are anodic to the Al matrix [135], and, hence, high-angle GBs are severely attached.

The impact of GB misorientation can be explained with the assistance of the Gibb’s free energy, and the change in the Gibbs free energy (∆G) can be expressed as follows:

$$∆\mathrm{G}={∆\mathrm{G}}_{\mathrm{s}}+{∆\mathrm{G}}_{\mathrm{c}}+{∆\mathrm{G}}_{\mathsf{\varphi }}$$

(9)

where ∆G_s is the surface free-energy term, ∆G_c is the strain energy term, and ∆G_ϕ is the chemical free-energy term.

Thermodynamically, a system with lower energy responds slowly to the activation energy and is stable. Similarly, low-angle GBs are low-energy areas and are slow to activate for atomic diffusion during the formation of the precipitate, and, as a result, low-angle GBs favor less the formation of precipitates along their boundaries. An exception to this rule does exist, although most of the low-angle GBs showed an immunity to attack, and high angles are prone to attack. There are GBs, such as GB9 and GB17, with misorientation angles > 15° that are less attacked, and some with misorientation angles ≤ 15° are severely attacked, such as GB24 and GB43. There is a special GB with ∑13b GB (GB35); despite having a misorientation angle of 27.8°, it showed a high reluctance to attack. Therefore, at this junction, we can affirm that the misorientation angle is not the only parameter, but there are undoubtedly other parameters that affect the GB precipitation, corrosion, and sensitization behaviors in alloys.

Effect of GB Plane Orientations

The literature published in this direction compels us to think that precipitation behavior may bear a relationship with the GB orientation plane [147,152]. To understand the influence of the GB plane orientation, the top-surface electron backscatter diffraction (EBSD) orientation maps and the GB traces [3,153] were analyzed. The same H₃PO₄ etched surfaces were further considered to understand the influence of GB plane orientations on precipitation. The GB plane distributions in standard triangles are shown in Figure 6d–g. The plane orientations for fully and partially etched boundaries are uniformly distributed in the standard triangles. Interestingly, the plane orientations of non-etched boundaries are predominantly aligned near [6] orientations. This raises a pertinent question: Why do these specific orientations resist etching? To address this, we examine the orientations of the etched boundaries to understand what makes them more susceptible to attack. The other planes close to {100} have facilitated the nucleation and growth of the β phase as they were the habit planes [154]. Habit planes are certain planes on which certain phenomena such as twinning [155,156,157], dislocation loops [158], or certain transformations [159] (FCC → hexagonal close packing (HCP)) are favored (here, we do not discuss much about the habit plane, as it falls beyond the capacity of this review). When the GB orientation planes are close to the habit plane of a particular variant, copious nucleation and growth occur in that particular variant. From the preceding discussion, it is now obvious that the planes are not habit planes, and, hence, these planes are reluctant to attack.

Coincidence site lattice (CSL) boundaries are special because they have a given fraction of atoms in the GB plane, which are nearly coincident with both the lattices which are generally separated by the GB. ∑13b showed a good immunity toward attack which indicates that this special boundary offers some special features that resist acid attacks and prevent corrosion. For an in-depth analysis of the Al-Mg alloy, the focused ion beam (FIB) and TEM cross-sectional area are considered, as shown in Figure 7. The arrow in Figure 7a,b indicates the {111} growth direction of the sputtered film, and columnar grains of ~200 nm size are clearly seen. There are few nanotwins [160] clearly visible inside the columnar grains and contain ∑3 as a special GB, the reflection of which is shown by selected area electron diffraction (SAED) in Figure 7b inset. The columnar grains are very fine; hence, it does not serve many purposes to quantify them into a size distribution profile chart. Therefore, a t-EBSD is shown in Figure 7c; the color denotes the plane orientation. What is conclusive from Figure 7c is the variation of the thickness of precipitates across the different GBs. GB1 is a Σ21a special GB which corresponds to the θ of 18.7°, GB2 is Σ37c with a θ of 49.1°, and GB3 is Σ39a with a θ of 31.3°. It is observed that the β-precipitation at GB1 is much thinner than that at GB2 and GB3. This is expected since GB1 has a low-angle GB misorientation. Moving to GB3 and GB5 (θ of 47.9°), precipitates at this GB are clearly seen as expected, owing to their high-angle GB. Although GB4 has a misorientation angle of 39.7°, the precipitates are much thinner on account of having a low Σ7, which belongs to the special GB.

At this junction, a clear picture can be presented from the above discussion on the influence of GB parameters on stress corrosion and the sensitization behavior of materials. The sensitization and stress corrosion are not only influenced by the grain misorientation but also influenced by the GB plane orientation. The prediction dependent upon the grain misorientation is straightforward only when the misorientation angle is less than 10°, where all the GBs show an excellent immunity toward attack. The situation becomes unpredictable when conditions arise where GBs are being severely attacked despite having a low-angle misorientation and in other instances where GBs are less attacked although they have a misorientation angle even > 15°. There is also no set pattern when considering the GB plane orientation to predict the corrosion behavior of the materials. Anomalous behavior does exist in this case, as well; for instance, precipitation is thinner in the GBs with a low-angle misorientation and also in GBs with a high-angle misorientation; there are certain exceptions to having a thinner precipitate zone even though they have a larger misorientation angle. In certain other cases, it is also found that a few peculiar planes, often known as habit planes, are the most favorable planes for the nucleation and growth of precipitates, favoring their formation on their planes. The situation is more complicated than expected if only the misorientation angles or GB plane orientation are taken into consideration. But the prediction becomes quite simple if CSL (Σ) is taken into consideration for the prediction of precipitates, sensitization, and corrosion activity in materials.

11. The Emergence of New Material as a Guard Against Corrosion

Mg alloys are often employed in the aerospace and automotive industries owing to their lower weight imbibed with high strength [161,162,163]. The low oxidation resistance restricts its applicability to other domains [164]. The Mg alloys are often prone to surface degradation at elevated temperatures during metal forming, welding, and heat treatments [165]. The low oxidation resistance of Mg alloy is ascribed to its low Pilling–Bedworth (PB) ratio of 0.81 and the compactness of the MgO layer [166,167]. Like another oxide layer, MgO does provide some initial protection over certain incubation periods, but the migration of Mg from the lattice to search for oxygen increases internal stresses, internal cracking, and de-bonding. This surge of migration causes the ignition and rapid oxidation to Mg alloys [168]. Alloying with Be has been shown to escalate the oxidation resistance of the Mg and Al alloys. It is believed that alloying with Be in Mg alloys lowers the inclusion of impurities in the Mg melt and favors the formation of the more stable BeO. Be is one of the elements that has a higher affinity towards oxygen than Mg, and this affinity toward oxygen is a prime factor in lowering down the oxidation rate, increasing the ignition temperature, and corrosion resistance of the alloys [169]. Zhao et al. [170], in his work, reported that alloying with Be improves the ignition temperature to 1033 K. Several reports also published directly an increase in the incubation temperature of Mg alloying in the ppm addition of Be [171,172,173]. The mechanism by which Be improves oxidation resistance is still unclear, but greater emphasis has been given to the formation of the more stable BeO for the increased oxidation resistance of Be-doped alloys. The following reactions of Be with Mg alloys are understood in light of the Gibbs free energy change ∆G⁰:

Be(l) + MgO(s) → BeO(s) + Mg(l)

(10)

The change in Gibbs free energy is given by the following:

$$∆G={∆G}^{0 }+RT ln{\displaystyle \frac{{\alpha }_{BeO}{\alpha }_{Mg}}{{\alpha }_{Mgo}{\alpha }_{Be}}}={∆G}^{0 }+RT ln{\displaystyle \frac{{\alpha }_{Mg}}{{\alpha }_{Be}}}$$

(11)

ΔG: change in free energy; α: activity factor; and R and T are universal gas and temperature in Kelvin, respectively.

Applying the above free energy equation, Zeng and co-workers [173] arrived at an interesting conclusion and reported that the formation of BeO in the presence of MgO is only possible when the Be concentration exceeds 0.88 wt % at 923 K. This conclusion agrees well with the findings of Inoue et al. [174], who calculated the PB ratio of CaO/Mg-10 atm. %10Al-5 atm. %Ca and BeO on AZ91 to be 1.17 and 0.62, respectively. Although the BeO layer is formed, it is insufficient to provide a protective layer to protect the Mg substrate from oxidation due to its low PB ratio. The segregation of Be²⁺ ions along with MgO on and along the GBs inhibits the migration of Mg²⁺ ions through the GBs and looks for reactive oxygen species, resulting in a lowering of the oxidation rate. The partly substituted Mg²⁺ by Be²⁺ (shown in Figure 8a) has many added advantages, as (1) the partial substitution of Mg²⁺ by Be²⁺ causes the lattice distortion of FCC-bearing Mg alloys and causes an increase in solid-solution hardening and impedes dislocation motions and prevents grain boundary sliding at high temperatures; and (2) the higher hardness values significantly play a major role in reducing grain boundary cracking during stress-induced corrosion. In recent decades, Ce and Nd have commonly been used to provide better oxidation resistance to Mg alloys. A schematic representation of Ce and Nd alloying with Mg alloys has been shown in Figure 8b and Figure 8c, respectively.

The idea is exactly similar in the case of Ce and Nd as well. The formation of their respective oxides offers a passive protective film on the metal surface, and its segregation along with MgO on and near the GBs inhibits the upward diffusion of cations to look for reactive oxygen. This inhibition of migration toward capturing the reactive oxygen is termed a corrosion control mechanism. The corrosion control via the addition of Al either during the alloy design or coating a layer of aluminum oxide is primitive and is well-supported through the literature as well, but corrosion protection via yttrium has recently been reported [176,177]. In a study on the influence of the rare earth (RE) metal yttria on Mg-Y alloys, it was found that adding a certain percentage of Y in this alloy significantly boosted the corrosion behavior. The synthesized alloy, denoted as Mg-mY (1 ≤ m ≤ 5), was heated in an environment of SF₆ at 793 K for 16 h. The corrosion resistance of this alloy immersed in 3.5 wt.% NaCl solutions showed no sign of appreciable degradation in structure due to the absence of a scattered Mg₂₄Y₅ structure. This second phase is detrimental when it disperses evenly inside the crystal lattice structure, which causes an impurity inside the crystal. This impurity inside the crystal is detrimental to corrosion resistance. The electrochemical impedance spectra measurement indicated that the corrosion resistance of the alloy increased to 6945 Ω cm⁻² with a corresponding current of 4.01 µA cm⁻² even at a time duration of 16 h with m = 5. When the immersion time in the NaCl solution was further increased, the corrosion resistance value started to decline at 5223 and 4731 Ω cm⁻², with a corresponding increase in current with a value of 4.99 and 5.5 µA cm⁻², respectively. The reason for the decline in the corrosion resistance value is attributed to the dispersion of the Mg₂₄Y₅ structure in the crystal, which acts as an impurity. Recently, a single crystal of the Ni-based superalloy coated with a Si-doped Pt-modified aluminide (denoted as PtSiAl) showed excellent oxidation/corrosion resistance even at a temperature of 1373 K for 320 h [178]. The mass gain in the case of PtAl and PtSiAl coatings for the first 30 h shows a surprising behavior with a greater increase in the mass gain of PtSiAl over PtAl. However, as time progresses, the gain in the mass of PtSiAl (1.26 mg cm⁻²) is much lower than that of the PtAl (1.37 mg cm⁻²), with both eventually settling down to a saturated value. This interesting behavior will be discussed a little later, but, before that, we highlight the diffusion pathway of the major elements. In both coating cases, a thick layer of Al₂O₃ is formed. At high temperatures and with a prolonged aging time, there is a gradual increase in the Al₂O₃ layer, which induces growth stress with a phase transition from the γ-Al₂O₃-stable α-Al₂O₃. At a threshold value, which depends upon the alloy, when the total internal stress is too high, the layers fail to resist any further growth of the oxide layer, and it begins to crack and peel off from the surface. This spallation and these cracks offer a pathway for the rapid diffusion of the oxidative elements and initiate the degradation of the protective mechanism in the alloys. At this stage, if the Al content is high enough to sustain the growth stress, another α-Al₂O₃ is readily formed, and the protection schemes continue. Otherwise, in the absence of Al, another oxide takes over the protective Al₂O₃ layer. The formation of mixed oxides is generally porous in nature (and, thus, the mixed oxides are not protective in nature), which, eventually, instigates the degradation of the substrate and the material loses its efficacy. Now, we return to our previous discussions on the protection benefits of the Si-doped coating in relation to the undoped PtAl coating. The area fraction of γ′-Ni₃Al in the outer layer of PtAl and PtSiAl is 12.9 and 11.7%, respectively, indicating that more of the β-(Ni, Pt)Al has been consumed in PtAl than PtSiAl. It is not only the outward diffusion of the Al that is important to the formation of the protective layer, but the mutual diffusion of the coating and substrate is also vital in deciding the overall protection of the alloy at high temperatures [179]. The possibility of Ni and W diffusing outward from the substrate and Al diffusing inward creates the possibility of β-phase formation which grows in size over time, eventually leading to the formation of M₂₃C₆ and the µ-phase. These phases are detrimental to the corrosion protection of the alloys at high temperatures as they induce cracks just beneath the top layer and act as an initiation site for crack propagation, thus weakening the mechanical properties and corrosion resistance. Turning toward the beneficial roles of Si is acknowledged by the percentage of available Si content in the outer (0.7 at. %) and the inner layer (2–4 at. %). The lower at % of Si in the outer layer indicates the dissolution of Si, thus retarding the formation of PtAl₂ and voids. The Si addition in the alloy promotes the upward migration of Ni, but the vacancies created are duly filled by Si, which benefits the alloy in promoting the formation of silicides due to the affinity of Si with refractory elements, especially the Cr [180]. The formation of silicides impedes the migration of refractory elements on the surface, impeding the detrimental phase formation on the surface, thus assisting the corrosion protection of the alloys at high temperatures.

To provide a concise summary of the literature on emerging materials for corrosion protection, we have included

Table 2. A quick summary on emerging materials for corrosion protection.

Base Material	Alloying Element(s)	Corrosion Resistance Improvement	Mechanism	Ref.
Mg alloy	Be	Increased ignition temperature to 1033 K; enhanced oxidation resistance	Formation of stable BeO; reduction in inclusion impurities; Be2⁺ segregation inhibits Mg2⁺ migration	[169,170,171,172,173]
Mg alloy	Ce, Nd	Improved oxidation resistance and passive film stability	Formation of CeO2/Nd2O3 layers at GBs; inhibition of Mg2⁺ cation migration	[168,175]
Mg-Y alloys	Y (1–5 wt.%)	Corrosion resistance increased to 6945 Ω·cm2; reduced current density (4.01 µA·cm2)	Absence of Mg24Y5 phase; suppression of structural impurities	[176,177]
Ni-based superalloys	Si (in PtSiAl coating)	Enhanced oxidation resistance at 1373 K for 320 h	Formation of silicides; retardation of voids and phase transitions; stabilization of Al2O3 layer	[178,179,180]

, which highlights the key alloying elements, mechanisms, and their respective contributions to enhancing corrosion resistance. The table summarizes the findings from a variety of studies to offer readers a comparative understanding of the strategies employed to improve the oxidation and corrosion resistance of metals and alloys.

12. Alloy Design Strategy in a Nutshell

Alloy design strategies are a multifaceted domain requiring a comprehensive understanding of material design from various perspectives. The first step in strengthening corrosion resistance for high-temperature applications is a thorough understanding of the corrosion mechanisms involved. It is envisaged that corrosion mechanisms at low and high temperatures operate through different processes [181]. For instance, many materials oxidize during oxidation at high temperatures, forming an oxide layer that can either protect the material or, if unstable, lead to further degradation. In certain environments, materials may also undergo carburization or sulfidation, which can weaken the structure and accelerate corrosion. Additionally, interfacial migration, such as the movement of atoms or vacancies at grain boundaries, can lead to localized corrosion [182] and structural weaknesses.

Once the corrosion mechanism is established, the second step is the alloying addition. Alloying elements play a crucial role in enhancing corrosion resistance, particularly in high-temperature environments. Chromium (Cr) is a key element that promotes the formation of a stable chromium oxide (Cr₂O₃) layer, which acts as a protective barrier against further oxidation and other corrosive processes [183,184]. Similarly, aluminum (Al) forms a protective oxide layer (Al₂O₃), which is especially effective in resisting both oxidation and sulfidation. In addition, Al forms a major precipitate of the type Ni₃(Al), popularly known as γ′ precipitates in Ni-based superalloys [184]. This precipitate is generally responsible for increasing the rupture strength of the alloys, particularly at high temperatures [185]. Nickel (Ni) is often used in superalloys due to its ability to resist oxidation and corrosion at high temperatures, while also stabilizing the austenitic phase and facilitating the formation of protective oxide layers. Cerium is another trace dopant in Ni-based alloys, acting as a deoxidizing agent that protects these commercial alloys [186]. Silicon (Si) in trace amounts contributes to corrosion resistance by forming a silicon oxide (SiO₂) layer, which is highly beneficial in aggressive environments. Finally, molybdenum (Mo) enhances the resistance to carburization and strengthens the overall alloy structure at elevated temperatures. These elements work synergistically to improve the performance and longevity of alloys in challenging conditions.

Strategic alloy design could be another strategy essential for developing materials that can withstand the harsh conditions of high-temperature environments. A critical aspect of this strategy is phase stability—ensuring that the alloy maintains a stable microstructure at elevated temperatures [187]. This stability prevents undesirable phase transformations that compromise the material’s mechanical properties and corrosion resistance. Another key approach is elemental balancing, which involves carefully optimizing the composition of alloying elements. This balance is crucial for achieving the right combination of corrosion resistance and mechanical strength, as different elements contribute to the formation of protective oxide layers while others enhance the alloy’s structural integrity. Additionally, the incorporation of RE elements plays a significant role in improving the adhesion of these protective oxide layers. RE elements such as yttrium or cerium can anchor oxide scales more firmly to the surface of the alloy, reducing the likelihood of spallation and, thus, extending the material’s service life in corrosive environments. Together, these strategies provide a comprehensive approach towards designing alloys that perform reliably under extreme conditions.

Another design strategy includes protective coatings and surface treatments. These are another crucial technique in enhancing the corrosion resistance of alloys, particularly in high-temperature environments [188]. Oxidation-resistant coatings involve the application of protective layers, such as aluminide or chromide coatings, to the surface of the alloy. These coatings work by forming a stable and adherent barrier that resists the penetration of oxygen and other corrosive agents. The aluminide coating, for example, forms a layer of aluminum oxide (Al₂O₃) on the surface, which is highly resistant to further oxidation. Similarly, chromide coatings create a protective chromium oxide (Cr₂O₃) layer. These oxide layers are particularly effective in high-temperature environments, where they prevent the underlying alloy from oxidizing, thereby extending its lifespan and maintaining its structural integrity.

Another recent design strategy that is gaining much attention is diffusion treatment [189]. This is another effective method with which to enhance corrosion resistance in high-temperature alloys. Techniques like pack cementation involve diffusing elements such as aluminum or chromium into the surface layer of the alloy. During this process, the alloy is exposed to a powdered mixture containing the desired element (e.g., Al/Cr) at high temperatures. This causes the element to diffuse into the alloy’s surface, forming a rich layer of that element on the surface. This diffused layer then reacts with oxygen to form a dense and protective oxide barrier, similar to the effect of direct coatings. The diffusion barrier created by this process is highly effective in preventing not only oxidation but also other forms of high-temperature corrosion, such as carburization and sulfidation. By strengthening the surface of the alloy, these treatments significantly enhance the material’s durability and resistance to harsh environmental conditions, making them essential for applications in industries like aerospace, power generation, and chemical processing.

In addition to these, GBE is another modern technique used to obtain the desired grain size for the optimal performance of the alloys engaged in high-temperature alloys (for more detail, refer to Section 9).

We present the following table to summarize the impact of metallic additions on corrosion resistance. This summary (

Table 3. A quick summary of metallic additions to enhance corrosion resistance.

Metallic Addition	Alloy Type	Environment	Outcome	Ref.
Cr	Ni-based	High-temperature oxidations	Formation of protective Cr2O3 layer; enhances oxidation resistance.	[73]
Al	Ni-based	High-temperature	Formation of Al2O3 oxide layer, providing a dense protective barrier.	[6]
Mo	CoCrFeNi HEA	3.5% NaCl solution	Formation of stable Mo-oxides; reduces pitting.	[74]
Cu	Al alloy	Acid rain (200 ppm Cl−)	Formation of Al2Cu phase; with high cathodic current density.	[81]
W	High-temperature alloy	High-temperature alloy	Selective oxidation; improves mechanical stability.	[88]
Co	Co-Al-W alloy	High-temperature air	Formation of Co3O4 and Al2O3 for dual-layer protection.	[87]
Y	Mg-Y alloy	3.5% NaCl solution	Formation of stable Y oxides; prevents Mg2+ migration.	[176,177]
Si	Ni-based superalloy (PtSiAl)	High-temperature air	Formation of silicides; controls phase formation.	[179]
Be	My alloys	High-temperature oxidizing environments	Forms BeO, which stabilizes MgO layer, and reduces oxidation rate and internal stress, enhancing oxidation resistance.	[170]
Re	Ni-based superalloys	NaCl solution	Forms passivation layer, and enhances corrosion resistance by reducing current density and promoting impedance.	[190]
Ti	Ni-based superalloys	High-temperature air	Contributes to passive oxide layer formation, enhancing oxidation resistance at elevated temperatures.	[64]
Ce	Ni alloys	High-temperature oxidizing environments	Acts as a deoxidizing agent.	[186]
Nd	Mg alloys	High-temperature oxidizing environments	Forms Nd2O3 layer, enhancing corrosion resistance by stabilizing the passive film.	[175]

) provides readers with a concise overview of the recent progress in this area, offering a foundation for understanding the current advancements and inspiring new research directions.

13. Current Challenges in High-Temperature Corrosion and Material Design

Having discussed various aspects of challenges in high-temperature alloy design and their protective measures, we come to the standpoint of looking at and re-exploring the available long-standing challenges in corrosion science. Recently, Mortazavi and co-workers explained the formation of a messy nanostructure, which helps to offer a protective layer on the high-temperature oxide-dispersion-strengthened alloy [191]. Despite several advancements in alloy design, corrosion protection has been an open ground that has yet to be covered successfully. A recent study demonstrated that corrosion resistance in high-temperature alloys, particularly Ni-based alloys, could be significantly improved by the addition of 5 wt % rhenium (Re) [190]. In a 3.5 wt % NaCl solution, the corrosion studies showed that Re can significantly form a passivation layer which significantly reduces the current density, promotes higher impedance arc, and enhances corrosion resistance. The corrosion morphologies of bare Ni-based and Re-doped alloys after polarization are revealed under an optical microscope (Figure 9a,b). This image clearly reveals that continuous pitting corrosion and several pits are interconnected, revealing that a corrosion attack has severely damaged this bare alloy. While Re-doped (Figure 9b) surfaces show somewhat identical pitting corrosion, they are uniformly distributed, and the depth of the pitting diameter is significantly reduced. The addition of Re assists in reducing the pitting corrosion kinetics, thereby reflecting enhanced corrosion resistance against oxidations at high temperatures. Potential polarization studies further reveal that the Re-doped alloy has a relatively lower i_corr (3.847 × 10⁻⁶ A cm⁻²) current than bare Ni (2.005 × 10⁻⁵ A cm⁻²), and, hence, a better corrosion oxidation resistance (Figure 9c). This also implies that the Re-doped alloy brings better corrosion-protective passivation films to the alloy surface. To further reveal the depth of corrosion penetration, surfaces were examined under a super-depth optical microscope (Figure 9d,e). It could be seen that the maximum variation in the surface of the Re-doped alloy is very small (~11 μm) against the 37.92 μm in undoped Ni. Such a small variation in the surface indicates a better corrosion-inhibition ability of Re-doping. These better corrosion resistances of this alloy are attributed to Re-doping. The presence of a uniform distribution of Re in the alloy results in the formation of a surface layer enriched with Nb elements. This enrichment hinders the diffusion of Mn, leading to a more continuous and dense oxide layer. As a consequence, the high-temperature oxidation resistance of the alloy is significantly improved.

Although several novel elements have been used for corrosion inhibition in high-temperature alloys, they have a limitation, particularly in the time duration fighting against corrosion. Technically, the lack of long-term data is a serious challenge in designing high-temperature alloys against corrosion. Many of these components are designed to last for many years. However, it is difficult to obtain long-term data on the corrosion behavior of materials, as this requires long-term exposure to high temperatures and corrosive environments. This lack of data makes it difficult to assess the long-term reliability of materials in high-temperature corrosion applications. On the other hand, the complex behavior of the corrosion mechanism makes it highly difficult to predict the nature of corrosion in the real-world environment, and, moreover, several of these mechanisms can further interact with each other, making corrosion studies further complex, particularly at high temperatures. Coating is another domain where less ground has been covered, and more in-depth studies are awaited to bring novel insights into corrosion in high-temperature alloys. Thus, materials for extreme environmental [192] conditions need to be further studied thoroughly to develop better alloys for tomorrow’s applications to realize a sustainable society.

14. Market Potentials of Protective Systems Used in High-Temperature Materials

The corrosion phenomenon occurring during high-temperature operations is more severe than the corrosion occurring in day-to-day life, where graphene and many other protective coatings can easily assist in building a protective layer, whereas it has little or no significant role to play in high-temperature corrosion resistance. Often, the corrosion at high-temperature is aggravated by the presence of high temperatures, moisture, oxygen, and many oxidizing elements that make the corrosion phenomenon very complex, and, thus, the protection system succumbs to these conditions. Even very effective graphene finds itself helpless in high-temperature corrosion protection due to its inherent defects. Thus, corrosion protection at high temperatures imposes a significant challenge on material scientists and invites them to open the ground to designing materials resistant enough to high temperatures. In a report published by Precedence Research, Asia Pacific alone dominated the corrosion market, with a record share of 52% and an estimated market of 5.76 billion in 2023. The market is projected to grow at a compound annual growth rate (CAGR) of 3.95% from 2024 to 2033 [193]. These statistics are quite alarming and emphasize the urgent need for improved protective systems, especially in high-temperature environments where conventional coatings like graphene fall short due to inherent defects and limitations. Recent reports by National Association of Corrosion Engineers (NACE) and several other technical reports estimate that the cost of corrosion protection amounts to 3.5% of any country’s global Gross Domestic Product (GDP) in addition to irreplaceable human loss due to catastrophic failure [194,195]. For instance, the annual cost of corrosion is as high as $276 billion in China, and India alone invests $110 billion per year, which is nearly 3.5 GDP of these countries. These staggering figures highlight the market potential for innovative protective systems designed specifically for high-temperature applications, presenting a lucrative opportunity for material scientists to develop new solutions that can better resist the complex and aggressive conditions of high-temperature operations.

Corrosion is a ubiquitous issue and a serious threat to the structural integrity of the components. Serious steps have been raised by the governments from all the countries to tackle this menacing threat. This can be realized from the fact that corrosion data have now been made available and are shared with all to benefit from the greater understanding that results.

15. Diverse Ideas in Research Do Wonder in Science

Chemical diversity in corrosion science has a lot to offer to material protection. With the sincere efforts of researchers, academicians, and industry partners in collaboration, the past few years have witnessed an unprecedented development in the corrosion protection of materials. We have seen a sublime transition in corrosion protection techniques from the early days of the inherent addition of elements like Cr, Al, W, Mo, Co, and Si to protect precious alloys, to a few sophisticated techniques in the modern day. Welcoming alloying elements to unite against corrosion protection is an old tradition, yet the most effective method available so far. Alloying elements bring the desired properties along with the complex situations of phase separations, creep growth, and grain boundary degradation over the service life. A trade-off must often be made to quench our desire in the venture of perfect corrosion-free material.

No technology is perfect; certain drawbacks do exist in all, and the same is true with corrosion protection and all other materials engaged in corrosion control. Novel materials, theoretical tools, and corrosion-resistant alloys must be discovered that can provide corrosion resistance at high temperatures for a longer duration.

Chemical diversity in corrosion science has significantly advanced material protection strategies. The collaborative efforts of researchers, academicians, and industry partners have driven the unprecedented developments in corrosion protection over the past few years. The evolution of corrosion protection techniques has been remarkable, transitioning from traditional methods like the addition of alloying elements (e.g., Cr, Al, W, Mo, Co, and Si) to more sophisticated modern techniques [196]. While alloying elements have been fundamental in enhancing corrosion resistance, they also bring challenges such as phase separations, creep growth, and grain boundary degradation over time. This necessitates a trade-off between achieving desired properties and mitigating long-term material degradation. However, current research is expanding beyond traditional alloying and exploring diverse and innovative approaches. Some of the emerging ideas and technologies in the field include the following:

Nanocoatings and Thin Films: Utilizing nanotechnology, researchers are developing ultra-thin coatings that provide enhanced corrosion resistance by creating more uniform and defect-free protective layers. These include nanocomposite coatings and self-healing coatings that can repair minor damage autonomously [197].

High-Entropy Alloys (HEAs): HEAs are being explored for their exceptional corrosion resistance and mechanical properties at high temperatures. By combining multiple principal elements in near-equimolar ratios, HEAs create complex microstructures that offer superior resistance to oxidation and corrosion [198].

Ceramic and Glass-based Coatings: Non-metallic coatings like ceramics and glass offer high-temperature stability and corrosion resistance. Techniques such as sol–gel processing and thermal spraying are being refined to create robust protective layers that can withstand extreme environments [199].

Electrochemical Surface Treatments: Advanced electrochemical methods, such as anodizing and plasma electrolytic oxidation, are being developed to modify surface properties and improve the corrosion resistance of materials. These techniques allow for the formation of protective oxide layers tailored to specific operating conditions [200].

Additive Manufacturing (3D Printing): Additive manufacturing enables the precise fabrication of complex geometries and tailored microstructures, which can be optimized for corrosion resistance. This approach allows for the incorporation of corrosion-resistant materials directly into the manufacturing process, reducing the need for subsequent protective treatments [201].

Computational Modelling and Machine Learning: The integration of computational tools and machine-learning algorithms is revolutionizing the design of corrosion-resistant materials. These technologies enable the prediction of material behavior under various conditions, accelerating the discovery and optimization of new alloys and coatings [202].

Green and Sustainable Corrosion Inhibitors: The development of environmentally friendly corrosion inhibitors derived from natural sources is gaining attention [203]. These inhibitors are designed to minimize environmental impact while providing effective protection in diverse industrial applications. The ongoing exploration of these diverse ideas reflects the dynamic nature of corrosion science and its commitment to overcoming the limitations of current technologies. By embracing novel materials, advanced theoretical tools, and innovative approaches, the field is steadily advancing towards more effective and sustainable solutions for corrosion resistance, particularly in high-temperature environments.

16. Outlook and Conclusions

In conclusion, a transparent and critical appraisal has been presented on the efforts of different materials in combating corrosion mechanisms. The quest to find the perfect material should continue along two key prospects: First, improvements in the alloying elements should be made to push the performance of the materials toward high-temperature applications. Another prospect is the development of models and systems to theoretically calculate the corrosion mechanism and offer a hybrid system. In such systems, the judicious selection of fewer materials can ensure not only the high-temperature performance capability but also the resistance to various corrosive environments. How well such systems are employed depends upon how accurately systems had been designed taking into consideration various corrosion-causing parameters. Within the next decade, we may expect the venture to discover a perfect material will come to see the light of the Sun and its applications in power plants and various high-temperature applications.