Cavitation Erosion Performance of the INCONEL 625 Superalloy Heat-Treated via Stress-Relief Annealing

Abstract

Cavitation-induced degradation of metallic materials presents a significant challenge for engineers and users of equipment operating with high-velocity fluids. For any metallic material, the mechanical strength and ductility characteristics are controlled by the mobility of dislocations and their interaction with other defects in the crystal lattice (such as dissolved foreign atoms, grain boundaries, phase separation surfaces, etc.). The increase in mechanical properties, and consequently the resistance to cavitation erosion, is possible through the application of heat treatments and cold plastic deformation processes. These factors induce a series of hardening mechanisms that create structural barriers limiting the mobility of dislocations. Cavitation tests involve exposing a specimen to repeated short-duration erosion cycles, followed by mass loss measurements and surface morphology examinations using optical microscopy and scanning electron microscopy (SEM). The results obtained allow for a detailed study of the actual wear processes affecting the tested material and provide a solid foundation for understanding the degradation mechanism. The tested material is the Ni-based alloy INCONEL 625, subjected to stress-relief annealing heat treatment. Experiments were conducted using an ultrasonic vibratory device operating at a frequency of 20 kHz and an amplitude of 50 µm. Microstructural analyses showed that slip bands formed due to shock wave impacts serve as preferential sites for fatigue failure of the material. Material removal occurs along these slip bands, and microjets result in pits with sizes of several micrometers.

1. Introduction

Nickel-based superalloys are intended for the manufacture of components used in the nuclear energy production industry, the chemical and petrochemical industries, aircraft gas turbines, and other hydromechanical equipment. The service properties (mechanical, physical, and chemical) of these engineering materials can be improved through alloying, the kinetics of solidification processes, and thermal and thermochemical treatments [1,2].

The increase in operating speed of equipment involving pressurized fluids makes them prone to cavitation erosion, leading to reduced performance and lifespan. Cavitation erosion is characterized by the formation and collapse of voids or bubbles filled either with vapor or with a combination of vapor and gases within a liquid [1]. The collapse of these bubbles generates shock waves and/or localized high-amplitude microjets, which can reach velocities between 300 and 500 m/s [1,2,3]. The growth and collapse time of cavitation bubbles is on the order of milliseconds, and the microcurrents formed during their impulse exert surface pressures ranging from 1000 to 4000 MPa, which exceed the yield strength of most materials [2,4,5]. The accumulation of this load due to repeated collapses causes plastic deformation, mechanical work hardening, the formation of microcracks, and material loss, leading to accelerated wear and even damage to components [4,5,6].

The degradation process of metallic materials through cavitation progresses by wear and is mainly based on two responsible mechanisms:

• Degradation by fatigue;
• Degradation by exceeding the shear strength or the cleavage fracture stress.

Considering these mechanisms, the following guidelines can be indicated to increase resistance to cavitation erosion:

• Increasing the mechanical strength characteristics (Rm, Rp0.2, and HV);
• Achieving a homogeneous and as fine as possible distribution of structural obstacles: for example, local dislocation clustering leads to crack formation;
• Ensuring acceptable deformability of the material so that stress concentrations can be relieved through slip processes.

Several laboratory and small-scale testing methods have been developed for evaluating cavitation resistance, namely:

• The ASTM G32 Vibratory Cavitation Test (introduced in 1972);
• The TVA High Velocity Jet Cavitation Test.

Both tests have been widely used to classify the cavitation erosion resistance of over 200 industrial alloys. The results have shown that the high-velocity jet cavitation test produces very intense wear and can erode industrial materials in approximately 10 times less time compared to the vibratory cavitation test [1,2,3,4,5]. However, there are differences in hydrodynamics and fracture mechanics between the two methods [5]: the hydrodynamics of high-velocity water jet cavitation is achieved through the acceleration and deceleration of the liquid flow using ultra-high-pressure pumps (thousands of bars), whereas the hydrodynamics of ultrasonic cavitation relies on high-frequency sound waves to induce bubble formation in a cloud attached to the cavitated surface.

From the perspective of how the cavitated surface is stressed, the vibratory test produces a large number of repeated impacts (with high density and frequency), involving shock waves and microjets of various sizes and velocities at arbitrary points, mainly causing elastic–plastic deformation, fatigue, and fracture. In contrast, the high-velocity water jet test generates a much smaller number of repeated impacts, in well-defined surface areas, leading to mechanical work hardening (strain hardening) of the surface.

Improving the erosion resistance of materials involves optimizing both their bulk and surface properties [1,2,6]. Numerous research studies have investigated the effect of thermal, mechanical, and thermochemical treatments, as well as various high-performance coating techniques applied to substrates to reduce or prevent cavitation erosion. Some studies even suggest that the combined use of hardening techniques and coatings creates a synergistic effect, making the material even more resistant [6,7,8,9,10,11,12,13,14].

D.T. Yu et al. [15] demonstrated that solution annealing (1040 °C) followed by artificial aging (480 °C) applied to 17-4PH stainless steels has favorable effects on microstructure, reduces internal stresses, and improves resistance to both corrosion and cavitation erosion.

Lassi Raam et al. [16] studied the cavitation erosion resistance of AISI 420 stainless steel after applying quenching and partitioning heat treatment. The results suggest that resistance to cavitation erosion significantly increases due to the high hardness and the fraction of retained austenite. During cavitation, the retained austenite can absorb the energy from bubble collapse by transforming into strain-induced martensite.

L.M. Zhang et al. [17] developed a super-austenitic stainless steel coating reinforced with SiC and applied via HVAF spraying to provide high resistance to both corrosion and cavitation. The improvement of these properties is attributed to low porosity, increased alloying, and refined grain structure.

Mayur Pole et al. [18] showed that layers consisting of CrC particles (12.5% and 6% vol.) in combination with INCONEL 625 (Inc-625) powder, deposited under specific carrier gas conditions on a cast A27 steel substrate, exhibit pronounced plastic deformation, grain refinement, and increased hardness—factors that justify the improvement in cavitation erosion resistance.

T. Zhao et al. [19] determined the optimal values of linear energy used in the surface coating of the Hastelloy C-276 alloy via the CMT (cold metal transfer) technique, in order to prevent cavitation erosion damage in stainless steel components used in marine engineering.

Qi Liu et al. [20] demonstrated that the incubation period and the maximum depth of erosion penetration (MDE) in the GH 4738 superalloy have more favorable values compared to those of the Inconel 601 superalloy. Additionally, they found that in the severe cavitation region of the GH 4738 superalloy, selective damage occurred, and the Mo element was preferentially removed.

The material studied in the present work is an austenitic alloy that possesses an excellent combination of mechanical properties and high corrosion resistance. Its high mechanical strength is justified by the solid solution strengthening effect of its FCC crystal structure by Cr, Mo, and Nb [20]. Its chemical composition is responsible for resistance to pitting, crevice corrosion, intergranular corrosion, and stress corrosion cracking in chloride-ion environments [21]. Due to its favorable combination of properties, it is widely used in the manufacture of pressurized water nuclear reactor components, aerospace components, and other hydromechanical equipment [20,21,22,23]. However, its range of applications can still be extended, particularly in cavitating environments, and the specialized literature contains limited data on its behavior in such conditions.

Numerous components made from this material are manufactured through plastic deformation operations, machining, and/or welding. Following these processes—or even during the technological processing flow—stress-relief annealing heat treatment is applied in order to reduce residual stresses and increase dimensional stability. As a result of its application, equilibrium states are obtained that are stable from chemical, physical, mechanical, and structural points of view. In terms of origin, residual stresses can be: thermal, structural, and mechanical. Thermal stresses occur during non-uniform heating or cooling of the metallic product, which causes a thermal gradient across the section. Structural stresses are caused by phase transformations during heating or cooling, which involve changes in specific volume. Mechanical stresses are introduced by external actions inherent to technological processing operations.

The objectives of this study focus on investigating the mechanisms of initiation and analysing the progression of cavitation-induced surface damage in components made from this metallic superalloy, thermally treated via stress-relief annealing. The main sections of the paper are as follows:

• The effects of cavitation and prevention methods;
• The material used and the experimental technique;
• Evaluation and interpretation of experimental results;
• Final conclusions.

2. Experimental Details

Cavitation erosion experiments were conducted using a vibrating apparatus equipped with piezoceramic crystals, fully complying with the requirements of the ASTM G32-2016 standard [24]. The purpose of the test is to simulate this process under accelerated conditions conducted in the laboratory. The experimental setup is schematically shown in Figure 1. The cavitation process was rigorously controlled through software implemented on the computer operating the system, which ensured the maintenance of prescribed operating parameters: vibration amplitude (50 μm), resonance frequency (20 ± 0.3 kHz), water temperature (22 ± 1 °C), and power of the ultrasonic generator (500 W). The pH of the drinking water was measured using a PH TESTER—POMETER, before, at the beginning, and at the end of each intermediate testing period, without observing any differences that could affect the degree of material degradation. The pH values measured during the testing days ranged from 6.98 to 7.11. No measurements were taken regarding the dissolved gas content due to the short duration of the intermediate periods (maximum 15 min).

The sonotrode has an exponential shape, and the quartz transducer is fused to its terminal end. The test specimen is attached to the terminal part of the exponentially shaped sonotrode. The working parameters generate longitudinal vibrations, which are amplified and transmitted into the liquid as ultrasonic waves. The bubbles formed during these vibrations cause implosions at the surface of the specimens, leading to a cumulative effect that has a destructive impact on the surface.

To eliminate any doubts regarding the chemical and structural homogeneity of the material, the specimens were taken from the same round bar made of the Ni-based alloy, INCONEL 625.

The nominal chemical composition of the material is shown in

Table 1. Chemical composition of researched alloy, wt. %.

Chromium (Cr)	22.01
Iron (Fe)	4.48
Molybdenum	9.20
Niobium (Nb) + Tantalum (Ta)	3.39
Carbon (C)	0.018
Manganese (Mn)	0.12
Silicon (Si)	0.14
Sulfur (S)	0.002
Phosphorus (P)	0.008
Aluminum (Al)	0.28
Titanium (Ti)	0.31
Cobalt (Co)	0.07
Nickel (Ni)	Balance

, and its mechanical properties at room temperature for the as-delivered state (as rolled) are presented in

Table 2. Mechanical characteristics of INCONEL 625 at room temperature.

Ultimate tensile strength, Rm	948 N/mm2
Yield strength, Rp0.2	527 N/mm2
Elongation at break, A5	46%
Reduction in area, Z	51%
Hardness, HV	236 daN/mm2

.

Stress-relief annealing (Figure 2) is a heat treatment applied to reduce internal stresses that occur, develop, and partially remain in metal parts as a result of the technological processing they undergo.

Technological operations such as casting, welding, machining, and heat treatments are accompanied by volumetric changes in the parts, which manifest at submicroscopic, microscopic, and macroscopic scales. However, these changes do not occur freely, i.e., simultaneously and uniformly across all microvolumes or macroscopic regions of the parts. Instead, they are constrained, meaning that variations occurring in some microvolumes or regions hinder the free variation of neighbouring ones, and vice versa. As a result, a portion of the volumetric change is realized through plastic deformation, while the rest remains in the form of elastic deformation, which, according to Hooke’s Law, is associated with residual stress—either normal (σ_res = ε_res • E) or shear (τ_res = γ_res • G).

These elastic deformations and residual stresses have varying magnitudes and signs and balance each other out, such that the metallic part retains its integrity, thus forming a mechanically stable system. However, this equilibrium is unstable because any external action (further processing or operational loading) induces new stresses that add algebraically to the pre-existing ones and may lead to undesirable effects such as additional plastic deformation, cracking, or rupture, depending on the total stress level and the mechanical properties of the material. Hence, the need arises to reduce or eliminate residual stresses in parts that are to be further processed or used in service—an objective achieved through stress-relief annealing.

In principle, the elimination of residual stresses is based on transforming the accompanying elastic deformations into plastic ones. According to Hooke’s Law, the value of residual stresses is proportional to the elastic modulus. Therefore, if the elastic modulus is lowered, the residual stress level will also decrease. Lowering the elastic modulus is possible by heating the metallic material, since the modulus is a measure of the bonding forces between atoms, which weaken as atoms move further apart with increased temperature.

Moreover, Hooke’s Law states that residual stress cannot exceed the material’s yield strength, as this would lead to plastic deformation and dimensional changes in the part. Thus, an inequality is established between the residual stress and the yield strength:

σ_res ≤ R_p0.2

From this inequality, it follows that reducing the yield strength leads to a decrease in residual stress to a new value corresponding to the lowered R_p0.2. It is well known that as temperature increases, the yield strength decreases, meaning that heating also helps reduce residual stresses via this mechanism.

Additionally, it is known that if a metallic material subjected to a stress below its yield strength is heated for a period at a certain temperature, the stress decreases over time due to the phenomenon of stress relaxation through creep.

From the above, it is clear that stress relief is achieved through three complementary mechanisms: reduction in the elastic modulus, reduction in the yield strength, and relaxation via creep—all of which are favored by increasing temperature.

According to the heat treatment cycle diagram shown in Figure 2, heating to the process temperature (650–870 °C) [25] is performed at a constant, low rate of approximately 100 °C/h to avoid introducing high thermal stresses that could add to the existing residual stresses and potentially deform or crack the parts. The heating temperature was established based on data provided by the alloy’s T.T.T. (time–temperature–transformation) diagram in order to avoid the precipitation of carbides and other intermetallic phases [21].

The holding time, calculated using the relation t_hold = g_max/25 (where g_max is the maximum thickness in mm), is essential for achieving the desired degree of stress relief.

The cooling rate must also be kept low to avoid inducing thermal stresses in the already stress-relieved parts. Typically, a rate of 50–60 °C/h is used.

The values of the mechanical properties obtained following the application of this heat treatment are presented in

Table 3. Mechanical characteristics of INCONEL 625 at room temperature after stress-relief treatment.

Ultimate tensile strength, Rm	828 N/mm2
Yield strength, Rp0.2	464 N/mm2
Elongation at break, A5	51%
Reduction in area, Z	53%
Hardness, HV	161 daN/mm2

.

The cavitation test specimens (Figure 3) were machined from a cylindrical bar with a diameter of 25 mm, which had previously undergone stress-relief annealing. The specimens are equipped with a threaded section to ensure secure attachment to the head of the cavitation vibrator. After the turning operations, the cavitation-attack surface was ground and polished to a surface roughness of Ra = 0.2–0.5 μm.

For cavitation testing, the specimens were submerged in a container filled with industrial supply water. Since the temperature of the testing liquid significantly influences the degree of cavitation erosion [25,26,27], the water temperature in the container was kept constant at 22 ± 1 °C. Two sets of 3 specimens each were tested at successive intervals of 15 and 60 min. This relatively short duration was chosen to study the development of progressive damage while minimizing additional effects such as corrosion.

Another set of 3 specimens was exposed to cavitation attack for a total of 165 min, divided into one period of 5 min, one of 10 min, and twelve periods of 15 min each. After each testing interval, the specimens were washed with acetone, air-dried, and weighed using a Zatklady analytical balance with a precision of 10⁻⁵ g.

At the end of the cavitation tests, cross-sections of the damaged surfaces were taken, embedded in synthetic resin, and prepared for metallographic examination. Microstructural investigations were conducted using a TESCAN VEGA 3 LMU scanning electron microscope equipped with Bruker EDX Quantax.

Vickers hardness measurements were performed using a Zwick/Roell YHV-S apparatus (Ulm, Germany), with an applied load of 5 kgf (HV5).

3. Results and Discussion

3.1. Cavitation Curves

According to laboratory convention [28], cavitation erosion is evaluated based on diagrams (cumulative mass loss of expelled material, erosion rates, and the v(t) curve representing the averaged values of the erosion rates) and characteristic values (the maximum value of the cumulative mass Mₘₐₓ, and the final erosion rate vₛ, toward which the averaged velocity curve v(t) tends to stabilize); see Figure 4.

Since laboratory customs require justification of the experiment’s accuracy and data processing, the velocity diagram (Figure 4b) includes the dispersion band of the experimental values, corresponding to the standard deviation σ = 7.881 × 10⁻³, which represents an error of ±1.5%, bounded above and below by the curves S(t) and I(t), respectively.

The formulas used are as follows:

-

Cumulative mass loss [28]:

$${\mathrm{M}}_{\mathrm{i}}={\displaystyle \underset{\mathrm{i}=1}{\overset{n}{\Sigma }}}\mathsf{\Delta }{\mathrm{m}}_{\mathrm{i}}$$

(1)

-

Erosion rate:

$${\mathrm{v}}_{\mathrm{i}}={\displaystyle \frac{{\Sigma }_{\mathrm{i}=1}^n\mathsf{\Delta }{\mathrm{m}}_{\mathrm{i}}}{{\Sigma }_{\mathrm{i}=1}^{12}\mathsf{\Delta }{\mathrm{t}}_{\mathrm{i}}}}$$

(2)

-

Mean rate curve:

v(t) = A·(1 − e^−B·t) + A·B·t·e^−B·t

(3)

-

Standard deviation:

$$\mathsf{\sigma }={\left[{\displaystyle \frac{{\sum }_{\mathrm{i}=1}^n{\left({\mathrm{v}}_{\mathrm{i}}-\mathrm{v}(\mathrm{t}{)}_{\mathrm{i}}\right)}^2}{n-1}}\right]}^{{\displaystyle \frac{1}{2}}}$$

(4)

-

Curves defining the width of the dispersion band (approximation error of ± 1.5%):

S(t) = v(t) + 1.5·σ; I(t) = v(t) − 1.5·σ

(5)

Explanation of terms:

Δmi—mass of material lost due to erosion during the period “i”, in mg;

v_i—velocity value determined for the intermediate period “i”, in mg/min;

Δti—duration of the intermediate period “i” of cavitation (the first 5 min, the second 10 min, and the remaining 15 min);

n = 12—total number of intermediate periods;

A, B—scale and shape parameters, calculated iteratively to minimize the deviations of the experimental values (i.e., the standard deviation σ relative to the mean curve).

The dispersions of the experimental values, obtained from the three samples tested over 165 min, as well as their evolutions with cavitation exposure time—whether referring to mass losses (Figure 4a) or erosion rates (Figure 4b)—show dependence on the amount of fragile intermetallic phases, grain sizes, and uniformity of mechanical properties in the surface structure, especially hardness, which is the parameter with the greatest influence, as demonstrated in the studies of Bordeasu [28], Hobbs [29], Garcia [30], Franc [31], Steller [32], etc.

From the viewpoint of the behavior of the three samples over the cavitation attack duration, the two diagrams show interference and even overlaps of experimental values. These aspects are well highlighted by the erosion rate values at 60, 90, 120, and 165 min, Figure 4b. Of particular interest is the identical behavior of the three samples at 60, 75, 80, and 120 min, when erosion rates and mass losses decrease, and at 45, 105, and 150 min, when these increase. These similarities, even the identities in the behavior of the three samples, reflect the technological heat treatment regime with well-chosen and controlled technological parameters. Based on research data by other authors on austenitic structured metallic alloys [4,5,6,9,10], increases in mass loss are explained by the elimination of grains of fragile intermetallic phases and the γ solid solution phase. During the 0–15 min period, the losses are primarily due to the removal of surface roughness peaks and parts of the fragile intermetallic phases. The large losses after 45 min are also affected by the solid solution losses, as the coalescence of generated cracks enabled the expulsion of large amounts of material, forming craters. Decreases at various cavitation attack durations, with relative constancy after 75 min, are explained by a reduction in impact force intensity due to hardening by work hardening and the damping effect of air and subsequently of the air/water mixture in the formed cavities. This information is supported by macroscopic images in Figure 5 at 105 and 165 min.

The low value of the standard deviation σ = 7.881 × 10⁻³ of the experimental erosion rate values relative to the mean curve shows not only the accuracy of the technological heat treatment regime but also that of experimental cavitation testing program management.

A first assessment of the surface structure’s behavior under cyclic loading from shock waves and microjets generated by bubble implosion in the cavitation cloud can be made based on the photographic images in Figure 5 at significant durations, correlated with the erosion rate values from Figure 4b. These images were taken with a Canon PowerShot A480 camera. The dark color of the surface after 15 min of cavitation exposure is the effect of the destruction of roughness peaks, removal of abrasive dust, and plastic deformations generating shadows and a crack network created by the pressure force from the surface impact with shock waves and microjets. The microscopic cavities formed during this time are few and caused by the expulsion of grains of fragile intermetallic phases.

With increasing cavitation attack duration (at 45 and 60 min), the structure yields, with typical local fatigue-type stresses [29,30,31,32,33,34], more pronounced toward the periphery where the erosion ring forms. This ring, analyzed via scanning electron microscopy, besides deformations and deep interconnected cracks, shows multiple cavities of varying sizes, characteristic of both the expulsion of γ solid solution grains and intermetallic phases, which led to increased erosion rate values and a maximum in the curve v(t) (Figure 4b).

Continued cavitation attack leads to intensified destruction through increasing geometric dimensions (in plane and depth) of formed cavities, intensification of deformations, and cracks propagating into the surface, as a result of fatigue loading mechanisms, without major variations in erosion velocity and mass loss values (Figure 4) during the intermediate durations. This behavior also explains the slight asymptotic decrease in the curve v(t) toward the stabilization value v_s = 0.026 mg/min, corresponding to the linear evolution of cumulative losses over the 60–165 min interval (Figure 4a). The slight changes in the shape, size, and aspects of the exterior ring and central zone, shown in images at 105 and 165 min, correspond to erosion velocity values over the 90–165 min interval. This is due to the fact that the intensity of pressure forces, although contributing to cavity growth by smaller material expulsions, is dampened by water and air penetrating the cavities during the compression phase of the mechanical vibrator system on one hand and by hardening through mechanical work of the stressed layer on the other hand.

3.2. Microstructural Investigations

3.2.1. Metallographic Analysis

The microstructure obtained after the applied heat treatment consists of polyhedral γ solid solution grains, some of which exhibit annealing twins. Within the grains, along their boundaries, and at the twin boundaries, carbide particles and intermetallic phases are present (Figure 6). These observations are consistent with the findings of other studies [18,21].

Figure 7a–d shows the microstructure of cross-sections of the surfaces after cavitation exposure for varying durations. After 15 min of cavitation (Figure 7a), only localized plastic deformations appear on the surface. After approximately 60 min (Figure 7b), pits and microcracks form, which then coalesce and lead to material removal in certain areas. It is evident that, whereas at the start of the cavitation process the surface wear is minimal and uniform (Figure 7a), as the exposure time increases, the wear becomes non-uniform (Figure 7b,c). By the end of the testing period, the merging of adjacent craters promotes the detachment of significant amounts of material from the surface zone (Figure 7d).

3.2.2. EDX Energy-Dispersive X-Ray Analysis

These investigations provide quantitative data on the atomic and mass concentrations of the elements in the chemical composition. They were carried out in the central area of the surfaces of specimens tested for cavitation erosion for 165 min (Figure 8).

Figure 9 and

Table 4. Mass and atomic concentrations of the elements in the chemical composition.

Element	Series	Mass C. (wt.%)	Norm. C. (wt.%)	Atom C. (at.%)	Error (1 Sigma) (wt.%)
Al	K series	0.20	0.21	0.29	0.05
Si	K series	0.14	0.15	0.23	0.04
Ti	K series	0.29	0.31	0.40	0.04
Cr	K series	19.09	19.22	22.59	0.57
Mn	K series	0.12	0.12	0.22	0.04
Fe	K series	4.26	4.47	4.96	0.15
Co	K series	0.08	0.09	0.13	0.06
Ni	K series	60.26	62.14	64.26	1.47
Nb	K series	2.14	3.62	1.26	0.11
Mo	K series	6.93	7.28	4.68	0.47
Ta	L series	1.01	2.39	0.98	0.10
	Total	97.98	100	100

show, respectively, the dispersion spectrum and the concentrations of the alloying elements detected in the chemical composition of the surface after 165 min of cavitation erosion.

Key observations drawn from the analysis of these data are as follows:

• A decrease in the chromium concentration from approximately 22% to approximately 19%, attributed to the removal of carbide particles containing this alloying element.
• No change in the iron and titanium concentrations, elements that dissolve in the Ni-based solid solution matrix but also form carbides.
• A reduction in the molybdenum concentration from approximately 9% to approximately 7%, with molybdenum being a constituent of MC- and M₆C-type carbides.

3.2.3. Topography of the Cavitated Surfaces

In Figure 10, SEM images are shown of the surface topography of specimens that underwent stress-relief annealing and were then subjected to ultrasonic cavitation for 15 min, 60 min, and 165 min. From these images, it can be seen that at short cavitation exposure (15 min, Figure 10a), initiation of small pits occurs along the boundaries between the crystalline grains, where the fragile intermetallic phases are located. As erosion progresses, the attack continues along the twin boundaries and the grain boundaries of the γ solid solution matrix (Figure 10b). Due to their brittleness, these phases are less capable of absorbing the deformation energy caused by the stresses induced in the material by cavitation impact waves and, consequently, deteriorate more rapidly than the γ solid solution grains with an FCC crystal lattice.

This image demonstrates that the material-removal mechanism is governed by plastic deformation, accompanied by the formation of slip steps characteristic of the γ-matrix.

By the end of the cavitation test (165 min, Figure 10c,d), the fractured surface exhibits ductile features, typical of face-centered cubic (FCC) materials. Some zonal variations in cavity depth are observed (exceeding 18 µm), and evidence of local fatigue loading is highlighted by the formation of characteristic deformation lines (Figure 10d). These observations are in agreement with studies of the degradation mechanisms in austenitic metallic alloys [4,5,6,12,19].

4. Conclusions

Microstructural investigations, together with the analysis of cavitation erosion characteristics performed on stress-relief–annealed INCONEL 625, yield the following conclusions:

• Initial stage: Vibratory cavitation erosion induces localized plastic deformations on the surface due to shock waves and microjets.
• Progression: Microjets create small pits, and shock waves produce slip lines. These pits develop along slip lines, precipitate–matrix boundaries, and twin boundaries, eventually coalescing into cracks and craters.
• The high values of ductility characteristics (A, Z), associated with favorable values of mechanical strength (Rm, Rp0.2), justify the slow and uniform degradation of the material.
• Sites of initiation: Annealing twin boundaries, slip lines, and precipitate-particle interfaces are the primary initiation sites for cavitation erosion.
• Microstructural analyses showed that slip bands formed as a result of the impact from microjets and shock waves generated by the repetitive implosion of cavitation bubbles serve as preferential sites for fatigue failure of the material.
• Material removal occurs along the slip bands, with microjets causing the formation of craters measuring a few micrometers, from which radial microcracks propagate, leading to fatigue fracture.