Next Article in Journal
Microwave Bow-Tie Diodes on Bases of 2D Semiconductor Structures
Previous Article in Journal
Atomistic Simulation Studies of Na4SiO4
 
 
Search for Articles:
Title / Keyword
Author / Affiliation / Email
Journal
Article Type
 
 
Advanced Search
 
Section
Special Issue
Volume
Issue
Number
Page
 
Journals
Crystals
Volume 14
Issue 8
10.3390/cryst14080719
crystals-logo
Submit to this Journal Review for this Journal Propose a Special Issue
► Article Menu

Article Menu

share Share announcement Help format_quote Cite question_answer Discuss in SciProfiles thumb_up
...
Endorse textsms
...
Comment
first_page
settings
Order Article Reprints
Font Type:
Arial Georgia Verdana
Font Size:
Aa Aa Aa
Line Spacing:
Column Width:
Background:
Article

Early Stages of Crack Nucleation Mechanism in Fe39Mn20Co20Cr15Si5Al1 High-Entropy Alloy during Stress Corrosion Cracking Phenomenon: Pit Initiation and Growth

by
Pranshul Varshney
and
Nilesh Kumar
*
Department of Metallurgical and Materials Engineering, The University of Alabama, Tuscaloosa, AL 35487, USA
*
Author to whom correspondence should be addressed.
Crystals 2024, 14(8), 719; https://doi.org/10.3390/cryst14080719
Submission received: 15 July 2024 / Revised: 2 August 2024 / Accepted: 7 August 2024 / Published: 11 August 2024
(This article belongs to the Special Issue Preparation and Applications of High-Entropy Materials)
Browse Figures
Versions Notes

Abstract

:
This study investigated the susceptible sites for pit nucleation in a transformation-induced plasticity (TRIP) Fe39Mn20Co20Cr15Si5Al1 (at.%) high-entropy alloy (HEA) in 3.5 wt.% NaCl solution. The investigation involved a constant-load stress corrosion cracking (SCC) experiment. The SCC testing was interrupted at different pre-determined time intervals to characterize the specimen surface using a scanning electron microscope (SEM), electron backscattered diffraction (EBSD), and a three-dimensional optical stereomicroscope. The EBSD results revealed pit nucleation at the susceptible γ–ε interphase and ε–ε interlath/plate boundaries. The three-dimensional profile and SEM results indicated an increase in pit depth with no change in pit diameter on the surface of the specimen as the experiment progressed over time. This study highlights the importance of microstructural features and mechanical loading in the corrosion behavior of TRIP HEAs, providing insights into the mechanisms of pit nucleation and growth under aggressive environmental conditions.

1. Introduction

Over time, as research on stress corrosion cracking (SCC) has progressed, several authors have formulated various definitions of SCC, suggesting that it is a phenomenon caused by the combined effect of applied load (external or internal) and a corrosive environment, resulting in premature failure of a susceptible material [1,2]. Cracks due to SCC can nucleate and propagate rapidly, resulting in sudden failure and catastrophic accidents. For instance, the “Silver Bridge” connecting Ohio and West Virginia collapsed in 1967, which was attributed to SCC [3]. The passive film is an important factor that plays a vital role during SCC. Since the passive film acts as a barrier between the metal and corrosive ions, it kinetically inhibits the electrochemical process (reduces reaction kinetics) at the metal–electrolyte interface, thus resisting corrosion [4]. According to the slip-step dissolution model, one of several micro-mechanisms proposed for SCC, the passive film ruptures under load, exposing the underlying material to corrosive ions, leading to anodic dissolution and causing pit nucleation and growth. Some of these pits eventually transition to cracks. Repassivation of the crack surface temporarily arrests the crack propagation, which resumes upon rupture of the passive film, continuing cyclically until the component fails [5,6,7].
Some widely accepted techniques to understand the SCC characteristics of materials include the slow strain rate tensile (SSRT) test, constant load tensile test, and compact tension test [8]. The constant load experiment involves applying a pre-determined load and recording the time to fracture in the given environment. As the stress (or load) level increases, the time to failure decreases, allowing the determination of the threshold stress at which SCC occurs [8]. This type of statically loaded experiment also evaluates the influence of metallurgical (e.g., chemical composition, grain size, precipitates, and secondary phase) and environmental factors (e.g., temperature and pH) on SCC.
The current micro-mechanistic understanding regarding SCC is based on experimental observations on conventional alloys consisting of one principal alloying element, and the other elements added in minor concentrations to achieve the desired characteristics. A study was conducted on austenitic steels (Type 304, 310, and 316) in a boiling saturated magnesium chloride solution to understand SCC behavior using the constant load method; transgranular cracking was observed, and the slip-step dissolution model was identified as the operating mechanism [9]. Forty and Humble [10], Engseth and Scully [11], and several other authors have investigated the SCC behavior occurring in a wide range of metals and alloys, e.g., aluminum alloys [12,13,14,15,16,17], magnesium alloys [18,19], steels [20,21,22,23], titanium alloys [24,25,26,27], and Ni-base superalloys [28,29,30]. Several authors studied the pit nucleation stages. For instance, Klapper et al. [31] studied the nucleation stages of transgranular SCC in 15.1 wt.% Cr–14.9 wt.% Mn stainless steel in a chloride environment at 88 °C using in-situ electrochemical and optical methods under constant load conditions. They reported that with an increase in load, the electrochemical activity significantly increased, forming several stable pits and resulting in specimen failure. The failure occurred from the base of the larger pits, suggesting the influence of load on pit formation and its transition to a crack during SCC. Calabrese et al. [32] investigated the evolution of damage during SCC on 17-4PH stainless steel using electrochemical noise (EN) and acoustic emission (AE) techniques. They reported that using EN and AE techniques, the four stages of damage (Stage I—incubation time before pitting corrosion; Stage II— stable pits grow to a size that cracks nucleated and subsequent stages of repassivation; Stage III—quiescence phase (progressive increase of the plastic zone at the crack tip during crack propagation); and Stage IV—catastrophic failure) can be characterized.
In the past two decades, multi-principal element alloys (MPEAs), also termed high-entropy alloys (HEAs) or complex concentrated alloys (CCAs), developed by Yeh et al. [33] and Cantor et al. [34], have seen intense research activities. HEAs are based on a novel approach to alloy design, which requires five or more principal alloying elements. A majority of these alloys exhibit simple crystal structures like face-centered cubic (FCC), body-centered cubic (BCC), and hexagonal close-packed (HCP) structures [35,36,37]. According to the literature, some of the newly designed alloys have excellent mechanical properties [38,39,40]. For instance, at room temperature, the two HEAs, AlCoCr2FeMn0.5Ni (at.%: 2 and 0.5 in the subscript denote Cr is 2 times and Mn is 0.5 times of other elements Al, Co, or Ni) and Al0.2Co1.5CrFeNi1.5Ti, exhibited very high hardness (more than the Inconel 718 superalloy). Although the hardness decreased with increasing temperature, it remained higher than Inconel 718 at similar temperatures [38,39]. The non-equiatomic HEA AlCoCrFeNiTi0.5 exhibited a high yield strength of 2.26 GPa and a ductility of 23% [41].
In addition to excellent mechanical properties, some HEAs have also demonstrated excellent corrosion behavior. Some of the HEAs outperformed SS304 in terms of corrosion resistance. For instance, the single-phase FCC CoCrFeNi (at.%) HEA exhibited better corrosion resistance compared to SS304L due to its high Cr and Ni content [42]. Although understanding SCC is essential, researchers have performed very few studies on HEAs. Ayyagari et al. [43] observed a rise in anodic current density under constant load conditions for Al0.1CoCrFeNi HEA, which was attributed to a change in alloy behavior under load. Varshney et al. [44,45] reported a decrease in ductility when tested transformation-induced plasticity (TRIP) Fe39Mn20Co20Cr15Si5Al1 (at.%) HEA in a 3.5 wt.% NaCl solution at room temperature. The decrease in ductility was attributed to SCC. In this work, Varshney et al. [46] noted the presence of corrosion pits at specific microstructural features of the alloy.
In the present work, the initial stages of pit nucleation in a TRIP Fe39Mn20Co20Cr15Si5Al1 (at.%) HEA in a 3.5 wt.% NaCl solution at room temperature were investigated. As aforementioned, SCC initiates from the fracture of a passive film, resulting in pit nucleation and, eventually, crack nucleation at such sites. This work focused on identifying the susceptible sites in the studied HEA that cause pit nucleation and understanding the initial stages of pit nucleation and growth. According to the published literature and the authors’ prior work on this alloy, it has two phases: ε—martensitic phase (HCP crystal structure); and γ—austenite phase (FCC structure) [47]. The studied alloy possesses several important characteristics that makes it an important HEA for studying the pit-to-crack transition [47,48,49,50,51]: (a) stability of the high-temperature and high-pressure phase (ε) at room temperature, (b) high strength due to the high fraction of the ε phase, (c) excellent ductility, and (d) excellent uniform corrosion resistance. Note that in conventional high-Mn steel, the ε phase is brittle. These characteristics position this alloy as a promising candidate for load-bearing applications in corrosive environments. Therefore, it becomes imperative to understand the initial stages of pit nucleation and identify microstructural features acting as favorable pit nucleation sites. A constant load SCC technique was employed to understand the initial stages of pit nucleation. The constant load SCC experiment was interrupted at specific time intervals. The research utilized advanced characterization techniques, including scanning electron microscopy (SEM), electron backscattered diffraction (EBSD), and a three-dimensional optical stereomicroscope, to study the early stages of pit nucleation.

2. Experimental Procedure

2.1. Material Processing and Microstructural Characterization of the Alloy (As-Received)

The vacuum arc melting method was utilized to synthesize the Fe39Mn20Co20Cr15Si5Al1 (at.%) HEA studied in the present work. It was synthesized by Sophisticated Alloys, PA, USA. The HEA was synthesized and cast in a 10″ × 8″ × 2″ water-cooled Cu mold. The HEA synthesis process involved the addition of master alloys in bulk form in the required composition andof the creation ofa vacuum in the chamber to remove the oxygen from the casting chamber. Due to the difficulty in getting rid of all the oxygen, multiple Ar backfilling and vacuum cycles were performed to ensure the removal of oxygen from the casting chamber. The vacuum level was 300 μm, and Ar was backfilled to 1 atm before melting. The alloy was cast and homogenized for 2 h at a temperature of 900 °C in an argon atmosphere. The homogenization temperature used in the process was determined from the pseudo-binary phase diagram, which was estimated using Thermo-Calc software. The temperature was chosen to achieve complete homogenization and prevent excessive grain growth. Following homogenization, the HEA underwent water quenching and was then warm-rolled at 500 °C until the desired thickness was achieved. The rolling temperature was selected to prevent recrystallization and enhance the austenite to martensite phase transformation. For further characterization, the HEA plate was sectioned into thin sheets with a thickness ranging from 1.2 to 1.4 mm using wire electrical discharge machining (EDM) provided by Mitsubishi.
A slow-speed diamond precision saw was employed to cut a specimen with dimensions of 10.0 (L) × 5.0 (W) × 1.2 (T) mm3. The specimen was mechanically ground and polished with grit papers ranging from 400 to 1200 and final polishing with 0.05 µm (colloidal silica solution). EBSD in SEM was employed to characterize the alloy for phase fractions. The EBSD experiment was conducted using Thermo Fisher Scientific, USA (Apreo S) instruments. An EBSD of the as-received specimen was performed at a magnification of 200X, with a step size of 1.0 µm, and a scanned area of 536 × 425 µm2. For EBSD in SEM, the accelerating voltage and beam currents were 20 kV and 1.6 nA, respectively. The experimental procedures involved in estimating the average grain size and the phases present are discussed in our previous publications [44,45,47].

2.2. Uniaxial Tensile Test

Since the constant load SCC test required the knowledge of the yield strength of the alloy to ensure the applied stress remained within the elastic regime (to avoid changes in microstructure due to plastic deformation), a flat dog-bone-shaped tensile specimen was used to conduct the tensile test from which the yield strength of the alloy was determined. The specimen was machined out using a wire-EDM instrument. Figure 1a illustrates the geometry of the tensile specimen, which was used to perform the uniaxial tensile test. The rolling direction in the tensile specimen was parallel to the tensile loading axis. The tensile specimen was mechanically ground and polished to achieve a final thickness of ~1.0 mm and a surface roughness of 0.05 µm using a colloidal silica solution. A uniaxial tensile testing instrument (ADMET) was utilized to perform the tensile test at room temperature in an air environment. The strain rate used during the tensile test was 10−6 s−1. The corresponding constant crosshead velocity during the tensile test was 0.0232 mm/s.

2.3. Constant Load SCC Experiment

Another flat dog-bone-shaped tensile specimen with a mirror-polished surface condition was used for the constant load experiment. To expose a pre-defined area of the specimen (gauge section) to the corrosive environment and prevent the formation of a galvanic couple between the specimen and the specimen fixture, the specimen was covered with chemical-resistant tape, as shown in Figure 1b. The exposed area had dimensions of 5.73 mm (l) and 1.83 mm (w). Figure 1c shows the constant load testing instrument and the salt solution chamber mounted on the lower arm of the crosshead. A magnified view of the salt solution chamber is shown in Figure 1d. The specimen was securely placed on the fixture and a load of 578 N (192 MPa—~71% of yield strength) was applied. Subsequently, a 3.5 wt.% NaCl salt solution was poured into the chamber. Initially, the specimen was kept under load in the 3.5 wt.% NaCl salt solution for 30 min (0.5 h), and then the test was interrupted to observe any changes that might have taken place during the experiment. The specimen was taken out, cleaned with acetone, and subjected to microstructural characterization (experimental details discussed in Section 2.4). Following the microstructural characterization, the specimen was cleaned again with acetone, covered with chemical-resistant tape, and subjected to the same load in the instrument. Multiple cycles of constant load experiments, interruption, microstructural characterization, and more constant load experiments with increasing time intervals (between two subsequent interruptions) were performed to understand the early stage of pit nucleation and growth at the microscopic scale.

2.4. Microstructural Characterization: Before and after Constant Load SCC Experiment

The flat dog-bone-shaped tensile specimen was subjected to microstructural characterization before starting the constant load experiment. The initial microstructural characterization was conducted using SEM to compare any changes in the specimen after each interruption in the constant load experiment. After the first interruption, which occurred at 0.5 h, the specimen was cleaned with acetone and subjected to microstructural characterization. SEM microscopy was performed, but no changes were observed on the specimen surface. Therefore, the constant load experiment was resumed without any other microstructural characterization. The experiment was interrupted again after 1.5 h (cumulative time), and microstructural characterization was performed using SEM, EBSD, and a three-dimensional digital stereomicroscope across the observed pits on the surface. The EBSD across the pits was performed at a magnification of 800×. The scanned area was 131 × 105 µm2 with a step size of 0.25 µm. Using a three-dimensional profilometer, multiple images at different pit depths were taken and then stitched together to obtain three-dimensional optical micrographs of the pits. These micrographs were used to determine the depth of the pit at each interruption. After microstructural characterization across the pits, the constant load experiment was resumed. Multiple cycles of interruption were conducted at cumulative times of 3.5 h, 5.5 h, 10 h, 20 h, 40 h, 98 h, and 194 h.

3. Results

3.1. Microstructural Characterization of the As-Received Alloy

According to the XRD result reported in the previously published work, the alloy possessed two phases, the ε martensitic phase, which has an HCP structure, and the γ austenite phase, which has an FCC structure [44,45,47]. The EBSD results are shown in Figure 2. The EBSD of the as-received specimen shows that the alloy has a higher fraction of the ε martensitic phase (97%), with the remaining portion beingγ austenite phase. According to the literature, the ε martensitic phase contributes to the strength [48]. The average grain size for the γ austenite phase was estimated to be 5.0 ± 3.0 µm. Note that the microstructural elements quantified using EBSD are dependent on parameters used during EBSD data acquisition and analysis. The average grain size of the ε phase was measured through optical microscopy. It showed the alloy consisted of coarse grains of the ε phase. The average grain size for the ε martensitic phase was estimated to be 220 ± 11 µm. The XRD and optical micrographs can be referred to in our published papers [44,45,47].
Figure 2c shows the misorientation angle distribution (MAD) plot of the two phases together. The plot shows two peaks at ~60° and ~70° misorientation angles. Figure 2d,e show the MAD plots of the γ austenite and ε martensitic phases, respectively. For the γ austenite phase, two peaks can be noted. One, at very low misorientation angles (<10°), indicates a high degree of plastic deformation, and the other, at ~60°, which is attributed to Σ3 twin boundaries. For the ε martensitic phase, two peaks were noted, one at a ~70°–75° misorientation angle, which is attributed to interlaths/plate boundaries, and the other at low misorientation angles indicating low-angle grain boundaries (LAGBs). The phase fraction of LAGBs in the γ austenite phase is ~60%, which is in contrast to the ε martensitic phase that has ~8% of LAGBs.

3.2. Early Stages of Crack Formation: Pit Initiation and Growth

3.2.1. Constant Load Selection

To investigate the initial stages of pit nucleation, an interrupted constant load experiment was performed, as mentioned in Section 2.3. To estimate the load that needs to be applied to the specimen during a constant load experiment, a uniaxial tensile test was performed at a strain rate of 10−6 s−1 in an air environment. Figure 3 shows the engineering stress vs. engineering strain plot, marked with the yield strength (270 MPa) and ~71% of yield strength (192 MPa). The ultimate tensile strength of the HEA is 656 MPa, and ductility is ~25%. The applied stress selected is in the elastic regime so that the surface condition does not change significantly due to macroscopic plastic deformation. The applied stress is also high enough to cause pit nucleation in a reasonable time frame.

3.2.2. Microstructural Characterization during Interruption

Figure 4 shows the SEM micrographs captured at different times during the interruption. Within the first 0.5 h of starting the experiment, no pits were detected. After 1.5 h, several pits emerged on the specimen’s surface, indicating the fracture of the passive film at certain locations, initiating anodic dissolution and resulting in pit nucleation. The film fracture may be attributed to localized plastic deformation [52]. The pits formed on the surface were a few microns in diameter after 1.5 h. Up to 5.5 h, no new pits were observed to form, but the depth of the pits changed (discussed in the later part), suggesting the occurrence of electrochemical activity or anodic dissolution (at a relatively slower rate) on the specimen’s surface.
At time t = 10 h, additional pits nucleated on the surface, and no new pits formed until t = 40 h. After 98 h, more pits nucleated, and some of them grew to approximately hundreds of microns in size. It is possible that as the time interval between interruptions increased during the constant load experiment, the localized acidic conditions within the pit were maintained for a longer duration before the subsequent interruption, resulting in pit growth compared to pits that nucleated within the shorter time intervals. Figure 4 demonstrates that the pit density increased with time, and as the time between interruptions extended, both the density and size of the formed pits increased. Note that due to lower magnification, except a few, the majority of the smaller micron-sized pits are not visible in the SEM macrographs included in Figure 4.
Figure 5 shows high-magnification SEM micrographs capturing the formation of pits on the surface at different time intervals. Figure 5a shows the pit that was observed at time t = 1.5 h after starting the constant load test. The pits shown in Figure 5b–d were observed after 20 h, 40 h, and 98 h, respectively. The pit shown in Figure 5d has a diameter of ~80 µm, and it is relatively larger than the other pits formed. Figure 6 shows a plot of pit area fraction (%) vs. time (h) for the pit shown in Figure 5a,b. The plot indicated that the pit size remained unchanged throughout the experiment. Positions 1 and 3, labeled on the plot, represent the times at which the two pits nucleated, while positions 2 and 4 indicate the termination time of the constant load experiment. The SEM micrographs corresponding to the labeled numbers on the plot are shown on the top and right sides of the plot. The SEM micrographs pictorially show that the pit size remained constant with time. However, it is probable that, instead of the size, the depth of the pits increased during the course of the experiment.
To validate this argument, three-dimensional digital profile microscopy was conducted on the pits shown in Figure 5a,b. Figure 7 shows a plot of the pit depth (µm) vs. time (h) across the two pits. The plot reveals a continuous increase in depth with immersion times for both pits. The plot is labeled with numbers 1 to 4, where 1 and 3 denote the pit nucleation time, and 2 and 4 show the termination time of the experiment. The three-dimensional optical micrographs corresponding to the labeled numbers on the plot are shown on the top and right sides of the plot. It is worth noting that these pits formed at boundaries or triple junctions, suggesting that boundaries are susceptible to corrosion ion attack.
To study the role of microstructure on the pitting susceptibility of the alloy, EBSD was performed on the pits shown in Figure 5a,b. Figure 8a,d show the phase maps overlaid with high-angle grain boundaries, HAGBs (>15°). A straight line was drawn across the boundaries adjacent to the pit to learn the misorientation angle of the boundaries susceptible to corrosion attack. A digitally magnified view of this area is shown in Figure 8b and Figure 8e, respectively. The drawn line intersects with boundaries labeled as numbers 1, 2, and 3 for pit 1 (Figure 8b) and number 1 for pit 2 (Figure 8e). Figure 8c,f show the point-to-point misorientation variation along the line and are also labeled with the corresponding numbers. For pit 1, the misorientation angles measured across the intersected boundaries were 60° (numbers 1 and 2) and ~75° (number 3), suggesting the presence of a γ–ε interphase and ε–ε interlath/plate boundaries, respectively [50,51]. Regarding pit 2, the point-to-point misorientation angle at the intersected boundary was ~75° (number 1), suggesting the presence of ε–ε interlath/plate boundaries.
The presence of two boundaries (γ–ε interphase boundary and ε–ε interlath/plate boundaries) across the pits shows that the pits likely nucleated along these susceptible boundaries, which could eventually serve as crack nucleation sites under the influence of load. The proximity of the pit to the γ austenite phase (Figure 8b) and the presence of the γ austenite phase within the pit (Figure 8e) suggest that the attack might have started from the interphase regions or the γ austenite phase. The TRIP alloys transform from the γ austenite phase (metastable state) to the ε martensite phase (relatively lower energy state). Corrosive ions typically attack high-energy sites [53]. Furthermore, the presence of an applied load locally increases the energy and catalyzes the corrosion attack at these susceptible sites.

4. Discussion

The present work is focused on investigating the initial stages of pit nucleation in a TRIP Fe39Mn20Co20Cr15Si5Al1 HEA. The HEA has two phases, i.e., ε - martensitic phase having an HCP structure and γ austenite phase with an FCC structure with a high fraction of the ε (97%) phase. The microstructural analysis revealed the presence of two types of boundaries in the HEA, i.e., γ–ε interphase and ε–ε interlath/plate boundaries at misorientation angles of 60° and 70–75°, respectively. The misorientation variation analysis across the pits revealed that γ–ε and ε–ε boundaries are the susceptible sites for pit nucleation. The three-dimensional micrographs taken across the pits at different times show an increase in depth.

Pit Nucleation

When the constant load experiment was started, no pits nucleated for the first 0.5 h. However, after 1.5 h of starting the experiment, four micron-sized pits were observed. According to the literature, pitting involves three kinetic stages: pit nucleation, metastable pit propagation, and stable pit propagation [54]. When a pit nucleates, it can either proceed towards the metastable pit propagation stage or cease to exist. This depends on the local potential (whether it is close to the pitting potential or not) and the repassivation rate (relatively lower than the anodic dissolution rate) [54]. After forming a metastable pit, it can again either develop into a stable pit or cease to exist. This depends on several factors that need to be fulfilled for a pit to stabilize and grow.
Some of these important environmental conditions are as follows: the presence of a pit cover [31,55], the prevention of diffusion of corrosive ions/cations from/to the bulk solution [55], the formation of a local acidic solution within the pit (decrease in pH) [56], and a change in the potential within the pit [56]. When such conditions are met, the metastable pit stabilizes and grows. The pit cover restricts the diffusion of cations to the bulk solution, and cations undergo a hydrolysis reaction with water, which reduces the pH within the pit because of the increase in H+ concentration. This causes a decrease in the repassivation rate and allows for the further dissolution of the material [57]. If the pit cover fractures before achieving the critical pit size, the local acidic solution gets diluted, and the pit growth stops due to the repassivation. However, if a pit achieves a critical size (diameter or depth), then even without the pit cover, the pit continues to grow but at a reduced rate [55,57,58]. Probably, within 0.5 h of starting the experiment, the necessary conditions did not exist and, due to this, the pit did not nucleate. It is also probable that the pits were covered with pit covers and, consequently, were not observed under SEM at the end of 0.5 h. However, within 0.5–1.5 h of starting the experiment, either the conditions were fulfilled and caused stable pit formation or the pit covers ruptured, revealing the nucleated pits. In the later stages of the experiment, several more pits were observed at multiple locations, shown in Figure 4, suggesting the fulfillment of important conditions for pit growth.
Pit nucleation also depends on the susceptible site. According to the literature, we know that corrosive ions generally attack sites that have relatively high energy [53]. The presence of a load enhances the susceptibility of a site to corrosion ion attack. From the interatomic distance vs. energy plot, we know that a change in interatomic distance due to the application of load will increase the energy of the system from its minimum or equilibrium energy level [59]. The increase in energy affects the interaction of atoms (especially on the surface) with the corrosive ions. The increase in energy is not uniform in polycrystalline, multiphase materials. This possibly increases the electrochemical activity, especially at the most susceptible sites present in the material.
Feng et al. [60] studied the corrosion behavior of low-nickel stainless steel in a groundwater solution under elastic and plastic strain conditions. They observed an increase in grain boundary energy and electrochemical activity with an increase in elastic strain. Lu et al. [52] mentioned that elastic load induces microplastic deformation in the material, generating more active sites and increasing anodic current density. The studied alloy consists of two types of boundaries, i.e., γ–ε interphase boundary and ε–ε interlath/plate boundaries, and microstructural characterization using SEM and EBSD revealed that the pit nucleation sites are γ–ε interphase boundary and ε–ε interlath/plate boundaries, shown in Figure 8. In the present study, under the influence of load, the increase in energy was possibly more at these two susceptible boundaries. Also, preferential deformation of the γ phase might cause the development of anodic coupling between the γ and ε phases, causing film fracture or pitting at these locations preferentially. This means that γ–ε interphase boundary and ε–ε interlath/plate boundaries could be the susceptible sites for pit nucleation, and they might eventually become crack nucleation sites.
Note that the diameter of some of the pits nucleated early on did not change, as shown in Figure 6. After several interruptions, we observed an increase in pit depth without a change in pit diameter, as shown in Figure 7. This generally happens when a pit achieves a stable size; then, instead of diameter, depth increases. Note that the increase in pit depth was not significantly high with every interruption as compared to the initial depth observed. This could be related to the absence of a pit cover. The SEM micrographs (Figure 5) show that the pits do not have a pit cover. According to the literature, for pit growth, a pit cover is needed to maintain a local acidic environment until the pit achieves a critical size for self-sufficient growth [57,61]. After achieving a critical pit size, the pit will continue to grow without its cover. It can be speculated that the interruption of the test did not maintain the aggressive anolyte, and it became diluted into the bulk solution [57]. Additionally, when the experiment was resumed, the pit cover was not present to maintain the acidic environment. This possibly affected the growth rate of the pits (depth). The presence of a pit cover might cause a relatively faster pit growth rate.
Additionally, when the experiment was interrupted after a relatively long time interval (t = 98 h), a relatively large pit (~80 µm) was observed on the specimen’s surface, as shown in Figure 5d. Possibly, because of the longer time duration between interruptions, the incubation time for pit nucleation and growth was sufficient for a pit of such size to form on the specimen’s surface [31]. However, when the experiment was resumed and interrupted at t = 194 h (after 96 h), the diameter of the same pit (Figure 5d) did not change, but the depth increased (SEM micrograph not shown here). This suggests that the local environment plays an important role in deciding the pit nucleation and its growth. Furthermore, the increase in pit depth could be related to the rupture and anodic dissolution of the passive film due to the dislocation activity (due to microplasticity) occurring in the material as a result of load application. The formation of pits on the specimen’s surface leads to stress concentration at these sites, which in turn might increase the dislocation activity (local high energy) and consequently enhance the film fracture rate and anodic dissolution following the slip-step dissolution model [7]. The experiment did not run until the specimen was fractured, and therefore, the pit transition to crack nucleation is not addressed in the present study. Investigating the pit transition to crack nucleation and propagation could be part of a subsequent study.
Additionally, it is acknowledged that the interruptions and cleaning of the sample with acetone might have influenced the pit growth rate or formation. To address this, an additional experiment with continuous 96 h exposure to the NaCl solution without interruptions will be carried out in future. This planned experiment will allow for a comparison of the number and size of pits with those observed in the interrupted tests, providing further insights into the impact of test interruptions.

5. Conclusions

The present study investigates the early stages of pit nucleation and growth in the TRIP Fe39Mn20Co20Cr15Si5Al1 (at.%) HEA. The specimen was subjected to a constant load equivalent to ~71% of the yield strength while immersed in a 3.5 wt.% NaCl environment at room temperature. The experiment was interrupted at different time intervals to investigates surface conditions. The SEM, EBSD, and three-dimensional digital stereomicroscopy techniques were utilized to investigate the susceptible sites for pit nucleation and growth. The pits were observed to nucleate at the γ–ε (interphase) and ε–ε (interlath/plates) boundaries. As the experiment progressed, the pit density on the surface was increased, and the maximum pit size (diameter) was found to be ~80 µm. Furthermore, the pit depth increased with each interruption.
Under the influence of the constant load, the microplastic deformation and incompatibility between different neighboring grains or phases possibly led to increased energy at the γ–ε (interphase) and ε–ε (interlath/plates) boundaries, evidenced by pit nucleation at such sites. These pits grew to a critical size and underwent anodic dissolution which possibly contributed to the increasing depth as the experiment progressed. Overall, this study provides insights into the initial stages of pit nucleation and growth in TRIP HEAs, emphasizing the role of microstructural features and mechanical loading in corrosion processes. Future work should investigate the transition from pit to crack nucleation and propagation to fully understand the long-term durability of these materials under similar conditions.

Author Contributions

Conceptualization, P.V. and N.K.; methodology, P.V. and N.K.; formal analysis, P.V.; investigation, P.V.; resources, N.K.; data curation, P.V. and N.K.; writing—original draft preparation, P.V.; writing—review and editing, P.V. and N.K.; visualization, P.V. and N.K.; supervision, N.K.; project administration, N.K.; funding acquisition, N.K. All authors have read and agreed to the published version of the manuscript.

Funding

This work was supported by the University of Alabama and the AIST foundation.

Data Availability Statement

The data presented in this study are available on request from the corresponding author due to privacy.

Acknowledgments

The authors express their gratitude for the support provided by the Alabama Materials Institute. They also thank Rajiv Mishra of the University of North Texas Denton, USA, for supplying the high-entropy alloy studied in this work.

Conflicts of Interest

The authors declare that they do not have any known competing financial interests or personal relationships that might have appeared to impact the research reported in this paper.

References

  1. Hines, J.G. On the propagation of stress-corrosion cracks in metals. Corros. Sci. 1961, 1, 21–48. [Google Scholar] [CrossRef]
  2. Meletis, E.I.; Hochman, R.F. A review of the crystallography of stress corrosion cracking. Corros. Sci. 1986, 26, 63–90. [Google Scholar] [CrossRef]
  3. McCafferty, E. Introduction to Corrosion Science; Springer: New York, NY, USA, 2010. [Google Scholar]
  4. Atrens, A.; Baroux, B.; Mantel, M. The secondary passive film for type 304 stainless steel in 0.5 M H2SO4. J. Electrochem. Soc. 1997, 144, 3697–3704. [Google Scholar] [CrossRef]
  5. Logan, H.L. Film-rupture mechanism of stress corrosion. J. Res. Natl. Bur. Stand. 1952, 48, 99–105. [Google Scholar] [CrossRef]
  6. Popov, B.N. Chapter 9—Stress Corrosion Cracking. In Corrosion Engineering; Elsevier: Amsterdam, The Netherland, 2015; pp. 365–450. [Google Scholar]
  7. Ford, F.P. The crack-tip system and its relevance to the prediction of cracking in aqueous environments. Environ. Induc. Crack. Met. 1990, 10, 139–162. [Google Scholar]
  8. Jones, R.H. Stress corrosion cracking. In ASM Handbook, Volume 13A—Corrosion: Fundamentals, Testing, and Protection; ASM International: Richland, WA, USA, 2003; pp. 346–366. [Google Scholar]
  9. Alyousif, O.M.; Nishimura, R. On the stress corrosion cracking and hydrogen embrittlement behavior of austenitic stainless steels in boiling saturated magnesium chloride solutions. Int. J. Corros. 2012, 1, 462945. [Google Scholar] [CrossRef]
  10. Forty, A.J.; Humble, P. The influence of surface tarnish on the stress-corrosion of α-brass. Philos. Mag. A J. Theor. Exp. Appl. Phys. 1963, 8, 247–264. [Google Scholar] [CrossRef]
  11. Engseth, P.; Scully, J.C. Repassivation studies on an austenitic stainless steel in chloride solutions. Corros. Sci. 1975, 15, 505–519. [Google Scholar] [CrossRef]
  12. Braun, R. Environmentally assisted cracking of alloy 7050-T7451 exposed to aqueous chloride solutions. Metall. Mater. Trans. A 2016, 47, 1367–1377. [Google Scholar] [CrossRef]
  13. Hua, T.S.; Song, R.G.; Zong, Y.; Cai, S.W.; Wang, C. Effect of solution pH on stress corrosion and electrochemical behaviour of aluminum alloy with micro-arc oxidation coating. Mater. Res. Express 2019, 6, 096441. [Google Scholar] [CrossRef]
  14. Cabrini, M.; Bocchi, S.; Urso, G.; Giardini, C.; Lorenzi, S.; Testa, C.; Pastore, T. Stress corrosion cracking of friction stir-welded AA-2024 T3 alloy. Materials 2020, 13, 2610. [Google Scholar] [CrossRef]
  15. Tang, Z.C.; Xu, W.; Zhao, D.Y.; Zhang, B. Improving the strength and SCC resistance of an Al-5Mg-3Zn alloy with low-angle grain boundary structure. J. Mater. Sci. Technol. 2023, 161, 63–73. [Google Scholar] [CrossRef]
  16. Ujjwal, K.; Al-Saadi, S.; Das, A.K.; Raman, R.K.S. Corrosion and stress corrosion cracking characteristics of 4043 aluminium alloy fabricated through directed energy deposition process. J. Alloys Compd. 2023, 976, 173154. [Google Scholar] [CrossRef]
  17. Freixes, M.L.; Zhou, X.; Aymerich-Armengol, R.; Vega-Paredes, M.; Peguet, L.; Warner, T.; Gault, B. Crack arrest markings in stress corrosion cracking of 7xxx aluminium alloys: Insights into active hydrogen embrittlement mechanisms. Scr. Mater. 2023, 237, 115690. [Google Scholar] [CrossRef]
  18. Rodrigues, P.S.; Zenóbio, I.R.; Da Silva, T.I.; Fernandes, C.Q.C.; De Sousa, T.G.; De Castro, J.A.; Da Fonseca, G.S.; Huguenin, J.A.O.; Ferreira, E.A. Effect of aging on corrosion resistance of AZ31 magnesium alloy. J. Mater. Eng. Perform. 2023, 33, 3413–3425. [Google Scholar] [CrossRef]
  19. Jia, H.; Li, J.; Li, Y.; Wang, M.; Luo, S.; Zhang, Z. Mechanical properties and stress corrosion cracking behavior of a novel Mg-6Zn-1Y-0.5Cu-0.5Zr alloy. J. Alloys Compd. 2022, 911, 164995. [Google Scholar] [CrossRef]
  20. Pan, Y.; Sun, B.; Chen, H.; Liu, Z.; Dai, W.; Yang, X.; Yang, W.; Deng, Y.; Li, X. Stress corrosion cracking behavior and mechanism of 2205 duplex stainless steel under applied polarization potentials. Corros. Sci. 2024, 231, 111978. [Google Scholar] [CrossRef]
  21. Ehrnstén, U.; Andresen, P.L.; Que, Z. A review of stress corrosion cracking of austenitic stainless steels in PWR primary water. J. Nucl. Mater. 2023, 588, 154815. [Google Scholar] [CrossRef]
  22. Pal, S. Role of residual stresses induced by U-bending on stress corrosion cracking susceptibility of F304 stainless steel. Adv. Mater. Process. Technol. 2023, 10, 315–328. [Google Scholar] [CrossRef]
  23. Martin, U.; Birbilis, N.; Macdonald, D.D.; Bastidas, D.M. Pit-to-crack mechanisms of 316LN stainless steel reinforcement in alkaline solution influenced by strain induced martensite. NPJ Mater. Degrad. 2023, 7, 88. [Google Scholar] [CrossRef]
  24. Joseph, S.; Kontis, P.; Chang, Y.; Shi, Y.; Raabe, D.; Gault, B.; Dye, D. A cracking oxygen story: A new view of stress corrosion cracking in titanium alloys. Acta Mater. 2022, 227, 117687. [Google Scholar] [CrossRef]
  25. Chen, T.; Zeng, Z.; Gao, H.; Liu, C.; Liu, Z.; Sun, Z.; Li, X. Influence of AC/DC on stress corrosion cracking of Ti-6Al-3Nb-2Zr-1Mo alloy in simulated marine solution. Int. J. Hydrog. Energy 2023, 55, 177–188. [Google Scholar] [CrossRef]
  26. Cao, S.; Zhu, S.; Samuel Lim, C.V.; Zhou, X.; Chen, X.; Hinton, B.R.W.; Boyer, R.R.; Williams, J.C.; Wu, X. The mechanism of aqueous stress-corrosion cracking of α + β titanium alloys. Corros. Sci. 2017, 125, 29–39. [Google Scholar] [CrossRef]
  27. Chiani, M.; Atapour, M. A study on the stress corrosion cracking susceptibility of friction stir welded Ti-6Al-4V alloy joints. Mater. Res. Express 2019, 6, 096598. [Google Scholar] [CrossRef]
  28. Moss, T.; Kuang, W.; Was, G.S. Stress corrosion crack initiation in Alloy 690 in high temperature water. Curr. Opin. Solid State Mater. Sci. 2018, 22, 16–25. [Google Scholar] [CrossRef]
  29. Telang, A.; Gill, A.S.; Teysseyre, S.; Mannava, S.R.; Qian, D.; Vasudevan, V.K. Effects of laser shock peening on SCC behavior of Alloy 600 in tetrathionate solution. Corros. Sci. 2015, 90, 434–444. [Google Scholar] [CrossRef]
  30. Li, D.; Li, K.; Gao, J.; Liu, Y.; Qin, C.; Li, J.; Li, Y.; Cao, W.; Zhai, Y.; Huang, G. Stress corrosion cracking of copper–nickel alloys: A review. Coatings 2023, 13, 1690. [Google Scholar] [CrossRef]
  31. Klapper, H.S.; Zajec, B.; Heyn, A.; Legat, A. Elucidating nucleation stages of transgranular stress corrosion cracking in austenitic stainless steel by in situ electrochemical and optical methods. J. Electrochem. Soc. 2019, 166, C3326–C3335. [Google Scholar] [CrossRef]
  32. Calabrese, L.; Bonaccorsi, L.; Galeano, M.; Proverbio, E.; Di Pietro, D.; Cappuccini, F. Identification of damage evolution during SCC on 17-4 PH stainless steel by combining electrochemical noise and acoustic emission techniques. Corros. Sci. 2015, 98, 573–584. [Google Scholar] [CrossRef]
  33. Huang, P.K.; Yeh, J.W.; Shun, T.T.; Chen, S.K. Multi-principal-element alloys with improved oxidation and wear resistance for thermal spray coating. Adv. Eng. Mater. 2004, 6, 74–78. [Google Scholar] [CrossRef]
  34. Cantor, B.; Chang, I.T.H.; Knight, P.; Vincent, A.J.B. Microstructural development in equiatomic multicomponent alloys. Mater. Sci. Eng. A 2004, 375–377, 213–218. [Google Scholar] [CrossRef]
  35. Tsai, M.; Yeh, J. High-entropy alloys: A critical review. Mater. Res. Lett. 2014, 2, 107–123. [Google Scholar] [CrossRef]
  36. Tong, C.; Chen, M.; Yeh, J.; Lin, S.; Chen, S.; Shun, T.; Chang, S. Mechanical performance of the AlxCoCrCuFeNi high-entropy alloy system with multiprincipal elements. Metall. Mater. Trans. A 2005, 36, 1263–1271. [Google Scholar] [CrossRef]
  37. Miracle, D.B.; Senkov, O.N. A critical review of high entropy alloys and related concepts. Acta Mater. 2017, 122, 448–511. [Google Scholar] [CrossRef]
  38. Chuang, M.; Tsai, M.; Wang, W.; Lin, S.; Yeh, J. Microstructure and wear behavior of AlxCo1.5CrFeNi1.5Tiy high-entropy alloys. Acta Mater. 2011, 59, 6308–6317. [Google Scholar] [CrossRef]
  39. Hsu, C.; Juan, C.; Wang, W.; Sheu, T.; Yeh, J.; Chen, S. On the superior hot hardness and softening resistance of AlCoCrxFeMo0.5Ni high-entropy alloys. Mater. Sci. Eng. A 2011, 528, 3581–3588. [Google Scholar] [CrossRef]
  40. Wang, Z.; Chen, S.; Yang, S.; Luo, Q.; Jin, Y.; Xie, W.; Zhang, L.; Li, Q. Light-weight refractory high-entropy alloys: A comprehensive review. J. Mater. Sci. Technol. 2023, 151, 41–65. [Google Scholar] [CrossRef]
  41. Zhou, Y.J.; Zhang, Y.; Wang, Y.L.; Chen, G.L. Solid solution alloys of AlCoCrFeNiTix with excellent room-temperature mechanical properties. Appl. Phys. Lett. 2007, 90, 181904. [Google Scholar] [CrossRef]
  42. Hsu, Y.; Chiang, W.; Wu, J. Corrosion behavior of FeCoNiCrCux high-entropy alloys in 3.5% sodium chloride solution. Mater. Chem. Phys. 2005, 92, 112–117. [Google Scholar] [CrossRef]
  43. Ayyagari, A.; Salloom, R.; Arora, H.S.; Mukherjee, S. Accelerated corrosion of high entropy alloys under tensile stress. arXiv 2021, arXiv:2106.03690. [Google Scholar]
  44. Varshney, P.; Mishra, R.S.; Kumar, N. Stress corrosion cracking of TRIP Fe39Mn20Co20Cr15Si5Al1 (at.%) high entropy alloy. In TMS 2021 150th Annual Meeting & Exhibition Supplemental Proceedings; Springer: Cham, Switzerland, 2021; pp. 742–750. [Google Scholar]
  45. Varshney, P.; Kumar, N. Stress corrosion cracking study of Fe39Mn20Co20Cr15Si5Al1 (at.%) compositionally complex alloy in 3.5 wt% NaCl salt solution. In TMS 2022 151th Annual Meeting & Exhibition Supplemental Proceedings; Springer International Publishing: Cham, Switzerland, 2022; pp. 1138–1145. [Google Scholar]
  46. Varshney, P.; Kumar, N. Investigating the susceptibility of Fe39Mn20Co20Cr15Si5Al1 high entropy alloy to stress corrosion cracking in a 3.5 wt.% NaCl environment. Corros. Sci. 2023, 222, 111422. [Google Scholar] [CrossRef]
  47. Varshney, P.; Mishra, R.S.; Kumar, N. Understanding the nature of passivation film formed during corrosion of Fe39Mn20Co20Cr15Si5Al1 high entropy alloy in 3.5 wt% NaCl solution. J. Alloys Compd. 2022, 904, 164100. [Google Scholar] [CrossRef]
  48. Nene, S.S.; Frank, M.; Liu, K.; Mishra, R.S.; McWilliams, B.A.; Cho, K.C. Extremely high strength and work hardening ability in a metastable high entropy alloy. Sci. Rep. 2018, 8, 9920–9928. [Google Scholar] [CrossRef] [PubMed]
  49. Mishra, R.S.; Nene, S.S. Some unique aspects of mechanical behavior of metastable transformative high entropy alloys. Metall. Mater. Trans. A 2021, 52, 889–896. [Google Scholar] [CrossRef]
  50. Polatidis, E.; Shukla, S.; Čapek, J.; Van Petegem, S.; Casati, N.; Mishra, R.S. Unveiling the interplay of deformation mechanisms in a metastable high entropy alloy with tuned composition using synchrotron X-ray diffraction. Mater. Today Commun. 2022, 30, 103155. [Google Scholar] [CrossRef]
  51. Sinha, S.; Nene, S.S.; Frank, M.; Liu, K.; Agrawal, P.; Mishra, R. On the evolving nature of c/a ratio in a hexagonal close-packed epsilon martensite phase in transformative high entropy alloys. Sci. Rep. 2019, 9, 13185. [Google Scholar] [CrossRef] [PubMed]
  52. Lu, B.T.; Zeng, Y.; Pang, X.; Luo, J.L. Effects of hydrogen and tensile stress on passivity of carbon steel. Corros. Eng. Sci. Technol. 2015, 50, 186–190. [Google Scholar] [CrossRef]
  53. Varshney, P.; Chhangani, S.; Prasad, M.J.N.V.; Pati, S.; Gollapudi, S. Effect of grain boundary relaxation on the corrosion behaviour of nanocrystalline Ni-P alloy. J. Alloys Compd. 2020, 830, 154616. [Google Scholar] [CrossRef]
  54. Tian, W.; Du, N.; Li, S.; Chen, S.; Wu, Q. Metastable pitting corrosion of 304 stainless steel in 3.5% NaCl solution. Corros. Sci. 2014, 85, 372–379. [Google Scholar] [CrossRef]
  55. Wu, K.; Byeon, J. Morphological estimation of pitting corrosion on vertically positioned 304 stainless steel using acoustic-emission duration parameter. Corros. Sci. 2019, 148, 331–337. [Google Scholar] [CrossRef]
  56. Scully, J.C. The interaction of strain-rate and repassivation rate in stress corrosion crack propagation. Corros. Sci. 1980, 20, 997. [Google Scholar] [CrossRef]
  57. Pradhan, S.K.; Bhuyan, P.; Mandal, S. Influence of the individual microstructural features on pitting corrosion in type 304 austenitic stainless steel. Corros. Sci. 2019, 158, 108091. [Google Scholar] [CrossRef]
  58. Kumar, N.; Fusco, M.; Komarasamy, M.; Mishra, R.S.; Bourham, M.; Murthy, K.L. Understanding effect of 3.5 wt.% NaCl on the corrosion of Al0.1CoCrFeNi high-entropy alloy. J. Nucl. Mater. 2017, 495, 154–163. [Google Scholar] [CrossRef]
  59. Filippova, V.P.; Kunavin, S.A.; Pugachev, M.S. Calculation of the parameters of the Lennard-Jones potential for pairs of identical atoms based on the properties of solid substances. Inorg. Mater. Appl. Res. 2015, 6, 1–4. [Google Scholar] [CrossRef]
  60. Feng, X.; Zhang, X.; Xu, Y.; Shi, R.; Lu, X.; Zhang, L.; Zhang, J.; Chen, D. Corrosion behavior of deformed low-nickel stainless steel in groundwater solution. Eng. Fail. Anal. 2019, 98, 49–57. [Google Scholar] [CrossRef]
  61. Pistorius, P.C.; Burstein, G.T. Metastable pitting corrosion of stainless steel and the transition to stability, Philosophical Transactions of the Royal Society of London. Ser. A Phys. Sci. Eng. 1992, 341, 531–559. [Google Scholar]
Crystals 14 00719 g001
Figure 1. (a) Computer-aided design (CAD) model of the flat dog-bone-shaped tensile specimen, (b) machined tensile specimen wrapped in chemical-resistant tape exposing only the area of interest in the gauge section during the test; (c) the constant load testing instrument (creep machine), which consists of an electrochemical setup mounted on the lower crosshead; and (d) a digitally magnified view of the electrochemical setup showing the different electrodes used during the test.
Figure 1. (a) Computer-aided design (CAD) model of the flat dog-bone-shaped tensile specimen, (b) machined tensile specimen wrapped in chemical-resistant tape exposing only the area of interest in the gauge section during the test; (c) the constant load testing instrument (creep machine), which consists of an electrochemical setup mounted on the lower crosshead; and (d) a digitally magnified view of the electrochemical setup showing the different electrodes used during the test.
Crystals 14 00719 g001
Crystals 14 00719 g002
Figure 2. EBSD results of the Fe39Mn20Co20Cr15Si5Al1 HEA: (a) an inverse pole figure (IPF) map with high-angle grain boundaries (HAGBs) (>15°) lines overlaid, (b) a phase map where green and red represent HCP and FCC phases, respectively, (c) a misorientation angle distribution (MAD) plot considering both phases together, (d) an MAD plot for the γ austenite phase, and (e) an MAD plot for the ε martensite phase.
Figure 2. EBSD results of the Fe39Mn20Co20Cr15Si5Al1 HEA: (a) an inverse pole figure (IPF) map with high-angle grain boundaries (HAGBs) (>15°) lines overlaid, (b) a phase map where green and red represent HCP and FCC phases, respectively, (c) a misorientation angle distribution (MAD) plot considering both phases together, (d) an MAD plot for the γ austenite phase, and (e) an MAD plot for the ε martensite phase.
Crystals 14 00719 g002
Crystals 14 00719 g003
Figure 3. An engineering stress vs. engineering strain plot for the Fe39Mn20Co20Cr15Si5Al1 (at.%) HEA when tested at 10−6 s−1 in an air environment, highlighting the yield strength and the stress used for applying a constant load to study pit nucleation and growth.
Figure 3. An engineering stress vs. engineering strain plot for the Fe39Mn20Co20Cr15Si5Al1 (at.%) HEA when tested at 10−6 s−1 in an air environment, highlighting the yield strength and the stress used for applying a constant load to study pit nucleation and growth.
Crystals 14 00719 g003
Crystals 14 00719 g004
Figure 4. SEM micrographs showing pits formation (black circled) on the specimen surface at different times (t), during the interrupted constant load experiment, performed on the flat dog-bone-shaped tensile specimen (Fe39Mn20Co20Cr15Si5Al1 (at.%) HEA). The time shown in the figure is the cumulative time.
Figure 4. SEM micrographs showing pits formation (black circled) on the specimen surface at different times (t), during the interrupted constant load experiment, performed on the flat dog-bone-shaped tensile specimen (Fe39Mn20Co20Cr15Si5Al1 (at.%) HEA). The time shown in the figure is the cumulative time.
Crystals 14 00719 g004
Crystals 14 00719 g005
Figure 5. High-magnification SEM micrographs of some of the pits which were nucleated at (a) t = 1.5 h, (b) t = 20 h, (c) t = 40 h, and (d) t = 98 h during the interrupted constant load tensile test on the Fe39Mn20Co20Cr15Si5Al1 (at.%) HEA specimen. The time mentioned here is the cumulative time.
Figure 5. High-magnification SEM micrographs of some of the pits which were nucleated at (a) t = 1.5 h, (b) t = 20 h, (c) t = 40 h, and (d) t = 98 h during the interrupted constant load tensile test on the Fe39Mn20Co20Cr15Si5Al1 (at.%) HEA specimen. The time mentioned here is the cumulative time.
Crystals 14 00719 g005
Crystals 14 00719 g006
Figure 6. Plot of pit area fraction (%) vs. Time (h) for the pits shown in Figure 5a,b. Positions 1 and 3 show the time the pits nucleated, while positions 2 and 4 show the end time of the constant load experiment. The SEM micrographs corresponding to the labeled numbers on the plot are shown on the top and right sides of the plot.
Figure 6. Plot of pit area fraction (%) vs. Time (h) for the pits shown in Figure 5a,b. Positions 1 and 3 show the time the pits nucleated, while positions 2 and 4 show the end time of the constant load experiment. The SEM micrographs corresponding to the labeled numbers on the plot are shown on the top and right sides of the plot.
Crystals 14 00719 g006
Crystals 14 00719 g007
Figure 7. Plot of pit depth (µm) vs. Time (h) for the pits shown in Figure 5a,b. Positions 1 and 3 show the time the pits nucleated, while positions 2 and 4 show the end time of the constant load experiment. The 3-dimensional profile micrographs corresponding to the labeled numbers on the plot are labeled and shown on the top and right sides of the plot.
Figure 7. Plot of pit depth (µm) vs. Time (h) for the pits shown in Figure 5a,b. Positions 1 and 3 show the time the pits nucleated, while positions 2 and 4 show the end time of the constant load experiment. The 3-dimensional profile micrographs corresponding to the labeled numbers on the plot are labeled and shown on the top and right sides of the plot.
Crystals 14 00719 g007
Crystals 14 00719 g008
Figure 8. (a,d) Phase maps across the pits shown in Figure 5a,b, overlaid with (>15°); (b,e) magnified view of the highlighted rectangles in (a,d) that consists of misorientation lines drawn on the maps; and (c,f) misorientation angle (°) vs. distance (µm) plots showing the point-to-point variation along the line drawn across the phase maps shown in (b,e), respectively. Numbers marked on the plot show the interface position through which the line passes.
Figure 8. (a,d) Phase maps across the pits shown in Figure 5a,b, overlaid with (>15°); (b,e) magnified view of the highlighted rectangles in (a,d) that consists of misorientation lines drawn on the maps; and (c,f) misorientation angle (°) vs. distance (µm) plots showing the point-to-point variation along the line drawn across the phase maps shown in (b,e), respectively. Numbers marked on the plot show the interface position through which the line passes.
Crystals 14 00719 g008
Disclaimer/Publisher’s Note: The statements, opinions and data contained in all publications are solely those of the individual author(s) and contributor(s) and not of MDPI and/or the editor(s). MDPI and/or the editor(s) disclaim responsibility for any injury to people or property resulting from any ideas, methods, instructions or products referred to in the content.

Share and Cite

MDPI and ACS Style

Varshney, P.; Kumar, N. Early Stages of Crack Nucleation Mechanism in Fe39Mn20Co20Cr15Si5Al1 High-Entropy Alloy during Stress Corrosion Cracking Phenomenon: Pit Initiation and Growth. Crystals 2024, 14, 719. https://doi.org/10.3390/cryst14080719

AMA Style

Varshney P, Kumar N. Early Stages of Crack Nucleation Mechanism in Fe39Mn20Co20Cr15Si5Al1 High-Entropy Alloy during Stress Corrosion Cracking Phenomenon: Pit Initiation and Growth. Crystals. 2024; 14(8):719. https://doi.org/10.3390/cryst14080719

Chicago/Turabian Style

Varshney, Pranshul, and Nilesh Kumar. 2024. "Early Stages of Crack Nucleation Mechanism in Fe39Mn20Co20Cr15Si5Al1 High-Entropy Alloy during Stress Corrosion Cracking Phenomenon: Pit Initiation and Growth" Crystals 14, no. 8: 719. https://doi.org/10.3390/cryst14080719

APA Style

Varshney, P., & Kumar, N. (2024). Early Stages of Crack Nucleation Mechanism in Fe39Mn20Co20Cr15Si5Al1 High-Entropy Alloy during Stress Corrosion Cracking Phenomenon: Pit Initiation and Growth. Crystals, 14(8), 719. https://doi.org/10.3390/cryst14080719

Note that from the first issue of 2016, this journal uses article numbers instead of page numbers. See further details here.

Article Metrics

Back to TopTop