Analysis of Influence of Ultrasonic Shot Peening on Surface Plastic Behavior of Superalloy

Abstract

This work focuses on the effects of ultrasonic shot peening (USP) on grain refinement and orientation behavior in the surface region of GH4151 superalloy. The microstructure evolution of the alloy under USP durations were studied. The effects of USP-induced grain refinement, orientation, and dislocation motion behavior were analyzed. The results indicated that during the USP process, the plastic deformation of the surface layer of superalloys is accompanied by changes in grain size and orientation. The random impact of the spheres on the surface area promotes grain refinement and grain rotation, enhancing the randomness of grain orientations and reducing the texture strength and the proportion of “soft” orientation distribution. Over a long period of treatment, a large number of spheres cause the slip planes and slip directions of each grain to rotate due to the additional shear stress from the impact, resulting in relatively consistent plastic deformation on the surface and the enhanced randomness of grain orientations, thus reducing the high texture strength introduced by previous machining processes. The understanding of dislocation pile-up behavior and the relationship between externally applied shear stress, pile-up characteristics, and grain refinement is essential for optimizing the USP process and achieving the desired material properties.

1. Introduction

USP treatment is commonly used to address the issue of fatigue crack initiation in alloys [1,2], typically occurring in the crack nucleation to microcrack stage (<1 mm). By suppressing crack initiation, it achieves an improvement in the fatigue life of alloys [3,4]. The initial propagation of fatigue microcracks is characterized by extension along slip bands, usually accompanied by plastic deformation processes. Microscopically, plastic deformation results from the relative sliding of one part of a crystal with respect to another along certain crystal planes and directions. Therefore, the crystallographic properties of materials have a certain influence on mechanical behavior during the initial stages of fatigue. Different materials exhibit different crystallographic characteristics [5,6], with the well-known three crystal structures being face-centered cubic (f.c.c) for Al, Cu, Ni, and γ-Fe, body-centered cubic (b.c.c) for α-Fe and β-Ti, and hexagonal close-packed (h.c.p) for α-Ti and Mg [7,8]. The elastic and plastic behavior of materials is dependent on their crystal structure, but even for the same lattice, significant variations can occur [9,10]. For instance, Cu exhibits significant anisotropy, while Al has a relatively smaller anisotropy, and α-Fe lies in between. High cycle fatigue typically occurs at low stress levels without macroscopic plastic deformation. Due to elastic anisotropy, the stress distribution between grains is non-uniform, with varying stress levels among the grains. Therefore, the behavior of microcrack propagation may differ for different types of materials [11,12]. In materials with a low stacking-fault energy, cross slip becomes difficult [13,14]. Al is a typical example of a material prone to cross slip, while nickel-based alloys exhibit resistance to cross slip. A limited activation of slip systems results in the continuous extension of microcracks along crystallographic directions until macroscopic cracking occurs [15,16]. Therefore, it is necessary to conduct research focusing on grain behavior. The present study employs the GH5151 nickel-based alloy, a high-temperature alloy resistant to 800 °C specifically designed for aero engine turbine disks, characterized by its high strength and resistance to deformation. The objective is to investigate the influence of USP on the plastic behavior of surface grains in alloys intended for turbine disks.

2. Materials and Methods

2.1. Materials

For the GH4151 (the chemical composition is shown in

Table 1. Chemical composition of the GH4151 alloy (mass fraction, %).

Co	Cr	Mo	W	Si	P	Al	Ti	Nb	Ni
15.00	11.05	4.60	3.10	<0.002	<0.006	3.80	2.80	3.35	Bal.

) alloy used in the experiment, the authors selected a solid solution temperature of 1080 °C and a solid solution time of 4 h. The standard double-stage aging heat treatment system was adopted, and the aging system was 840 °C × 6 h/air cooling + 760 °C × 32 h/air cooling. The sheets used in this work were cut into specific dimensions of 10 mm × 10 mm × 5 mm. Subsequently, the specimens were polished with #600, #800, #1500, and #2000 waterproof abrasive papers to achieve a relatively smooth surface. Then, they were placed in absolute ethanol for ultrasonic cleaning and to ensure a smooth and clean surface.

2.2. USP Treatment and Characterization Methods

USP has been proved to be an efficient surface treatment method for the surface grain refinement of metal and alloys [17,18]. Figure 1 shows a schematic illustration of the USP system. The spheres move at a high speed in the chamber and continuously impact the specimen, resulting in severe plastic deformation and grain refinement on the surface layer. During the USP treatment, ZrO₂ spheres (diameter 2.5 mm) are accelerated by a vibrating solid surface driven by a high-frequency ultrasonic signal (20 kHz) in an enclosed chamber made from stainless steel, employing 1000 spheres. The distance between the GH4151 specimen and the vibrator was 70 mm, and the processing durations were 20 min. The vibration intensity (%) was defined as the percentage of actual amplitude to the limit amplitude and set as 80%. The standard A-type Almen test pieces were applied to test the shot peening intensity, and the corresponding shot peening intensity was 0.25 mmA.

Electron back-scattered diffraction (EBSD) was used to characterize the cross-sectional microstructure changes, including grain size, dislocation density, and grain boundaries. The scan step of the EBSD was set to 90 nm to ensure the validity of the data, with an accelerating voltage of 20 kV. The dislocation distribution was characterized using a Titan ETEM G2 transmission electron microscope (TEM). The specimens were thinned down to small round disks with a diameter of 3 mm and a central hole through electrolytic double jet thinning. The composition of the double jet solution was 10 mL HClO₄ + 90 mL C₂H₅OH, and the voltage was set at 35 V.

3. Results

3.1. Effect of USP on Grain Size

EBSD was used to investigate the microstructure of the GH4151 matrix. Figure 2a,b show the position of the target grains (orientation) of the specimen after treatment by USP and non-peened (NP), in cross-sectional view; the average grain size was 18.9 μm and 12.6 μm, respectively (Figure 2c,d). Figure 2c,d show the TEM analysis of the microstructure of the top region of the specimen after different treatments. USP produced shear bands and localized dislocations, while NP showed only a few dislocation lines distributed throughout the entire TEM foil (Figure 2c,d) [19,20]. Compared with the specimen in Figure 2d, the grain size, standard deviation, and coefficient of variation of the specimen in Figure 2c are larger, and the grain size consistency is not as good as the grain size distribution after treatment. It is shown that USP treatment is helpful in ensuring the consistency of grain size distribution on the basis of reducing the grain size. Moreover, at the same depth of the cross-section, the grain size of the specimens in Figure 2d is significantly smaller than that of the specimens in Figure 2c, indicating that USP can promote the depth of the grain refinement layer.

3.2. Effect of USP on Grain Orientation

The material texture indicates the presence of the preferred orientations of grains along certain crystallographic directions, reflecting the material’s anisotropy, which significantly influences its mechanical properties and plastic deformation [21,22]. Figure 3a,b depict the pole figures obtained through EBSD for GH4151 alloy with different treatments. Under the conditions of USP, the material exhibits a reduced texture strength, with the maximum texture strength decreasing from 5.72 mud to 2.83 mud compared to the NP state. The texture becomes more randomly oriented, indicating that USP promotes the randomness of crystal orientations in GH4151 alloy [23]. The spheres, driven by the vibrating head, impart random impacts on the surface of the specimen, causing the rotation of surface grains and changes in their orientations. Figure 3c,d show the distribution of Schmid factors and the statistical relationship of Schmid factor quantities for the different specimens. The distribution of Schmid factors reflects the plastic deformation induced by USP, with a negative correlation between plastic strain and Schmid factors. The NP specimen exhibits the highest proportion of Schmid factors, close to 0.5, accounting for approximately 17.7%, indicating the characteristic of a soft orientation distribution. After USP, the proportion of Schmid factors close to 0.5 significantly decreases, indicating an increased proportion of γ-phase crystal {111} planes and <1–10> directions inclined at approximately 45° to the X-axis direction. This facilitates the easier glide of the γ-phase crystal {111}<1–10> slip systems and results in the formation of a large number of dislocations, thereby improving the fracture toughness, fatigue, and other mechanical properties of the surface layer [24,25].

4. Discussion

4.1. Analysis of Grain Refinement Mechanism

Under the influence of USP, superalloys undergo volume contraction, generating microscopic stresses within the material that induce plastic deformation [26,27]. Grain refinement is one of the main characteristics accompanying plastic deformation in the surface layer of alloys. Previous research [28,29] has acknowledged the influence of USP on grain structure, yet it has not elucidated the mechanisms responsible for these effects. The emphasis has predominantly been placed on macroscopic morphological characteristics and dislocation activities, with scant attention given to an analysis at the grain level. Figure 4 illustrates the mechanism of grain refinement during USP. The alloy of the pre-start period exhibits a scattered distribution of dislocations (Figure 4a) with a relatively low density. At the initial period, the motion of dislocations within grains leads to the formation of dislocation lines (Figure 4b) [30,31], accompanied by small magnitudes of microscopic stress. Under the influence of plastic deformation in the surface layer, dislocations inside the grains locally aggregate and form dislocation walls, dividing the original grains into individual dislocation cells (Figure 4d) [32]. At the later period, stress accumulates, causing the aggregation of dislocations into low-angle grain boundaries, resulting in grain refinement (Figure 4c). Solutes in the alloy tend to segregate near the dense regions of dislocations close to the grain boundaries. At the completion period, the distribution of precipitates near the grain boundaries increases. These fine and uniform precipitates can pin the grain boundaries, impeding dislocation motion and significantly enhancing the mechanical properties of high-temperature alloys [33]. According to the Hall–Petch equation [32,34],

$${\sigma }_{\mathrm{s}}={\sigma }_0+K{d}^{-{\displaystyle \frac{1}{2}}}$$

(1)

In the equation, ${\sigma }_{\mathrm{s}}$ is the yield strength of the polycrystal, $d$ is the average grain diameter, ${\sigma }_0$ is the resistance to deformation within the grains, which is equivalent to the yield strength of an individual grain, and $K$ is the influence coefficient of the grain boundaries on deformation, which is related to the grain boundary structure. The equation indicates that the smaller the grain size, the greater the yield strength, and the more difficult it is for dislocations to move.

USP has a significant impact on the distribution of dislocations in the surface layer of superalloys [30]. During the process of USP, multiple dislocation sources are activated in the surface layer of the alloy, releasing a series of dislocations. Under the action of the shear stress ${\tau }_0$ generated by the random impacts of the spheres when encountering obstacles, the dislocations accumulate and form dislocation pile-ups [32,35]. Within a dislocation pile-up, each dislocation is in equilibrium under the influence of two forces: the externally applied stress ${\tau }_0$ from the USP spheres and the stress field caused by other dislocations.

Since dislocations of the same sign repel each other on the same slip plane, the arrangement of dislocations within a dislocation pile-up exhibits a certain regularity, becoming sparser as the distance from the obstacle increases. The externally applied shear stress in USP is related to the distribution of dislocations within the pile-up. Considering the jth dislocation, it is influenced by two forces: the shear stress exerted by the sphere impact and the external stress, according to the Peierls model [27,36]. The effect of the externally applied shear stress on the jth dislocation is denoted as ${\tau }_{0}b$, where $b$ is the Burgers vector of the dislocation, and this force causes the jth dislocation to move in the positive x-direction. The other force acting on the jth dislocation is the stress field exerted by all other dislocations except for the jth dislocation. Thus, the force exerted by the stress field of the ith dislocation on the jth dislocation in the negative x-direction can be expressed as

$$F_{\mathrm{ij}}={\displaystyle \frac{\mu b^2}{2\pi (1-\nu )}}{\displaystyle \frac{1}{x_j-x_i}}$$

(2)

The sum of the forces exerted by all dislocations except for the jth dislocation on the jth dislocation can be expressed as follows:

$$F_s={\displaystyle \frac{\mu b^2}{2\pi (1-\nu )}}{\displaystyle \sum _{\begin{array}{l}i=1\\ i\ne j\end{array}}^n\frac{1}{x_j-x_i}}$$

(3)

At equilibrium, the sum of the forces acting on the jth dislocation is zero.

$$F_j={\tau }_{0}b-{\displaystyle \frac{\mu b^2}{2\pi (1-\nu )}}{\displaystyle \sum _{\begin{array}{l}i=1\\ i\ne j\end{array}}^n\frac{1}{x_j-x_i}}=0$$

(4)

When $D={\displaystyle \frac{\mu b}{2\pi (1-\nu )}}$ and ${\displaystyle \frac{{\tau }_0}{D}}={\displaystyle \sum _{\begin{array}{l}i=1\\ i\ne j\end{array}}^n\frac{1}{x_j-x_i}}$

This equation represents the equilibrium equation for the dislocation barriers. When n is large,

$$x_i={\displaystyle \frac{D{\pi }^2}{8n{\tau }_0}}{(i-1)}^2$$

(5)

Equation (5) indicates that the position of each dislocation $x_i$ is directly proportional to ${(i-1)}^2$, which means that the arrangement of dislocations within the barriers is non-uniform, with a denser arrangement closer to the obstacle.

The position $x_n$ of the nth dislocation can be approximately regarded as the total length of the dislocation barriers, denoted as L.

$$L\approx x_n=D{\displaystyle \frac{{\pi }^{2}n}{8{\tau }_0}}\approx 2D{\displaystyle \frac{n}{{\tau }_0}}$$

(6)

In the equation, when $D={\displaystyle \frac{\mu b}{2\pi (1-\nu )}}$, the dislocations in the barriers are edge-type dislocations, and when $D={\displaystyle \frac{\mu b}{2\pi }}$, the dislocations in the barriers are screw-type dislocations. The number of dislocations within a pile-up with a length of L can be expressed as

$$n={\displaystyle \frac{{\tau }_{0}L}{2D}}$$

(7)

Equation (7) demonstrates that the number of dislocations within a dislocation barrier is directly proportional to the externally applied shear stress ${\tau }_0$ introduced by the impacts of USP spheres, and the distance L, between the dislocation source and the barrier. When L is constant, the slip plane of the crystal experiences the effect of ${\tau }_0$, and the dislocation source continuously releases dislocations, gradually increasing the number of dislocations within the barriers. When the number of dislocations n reaches a certain threshold, the barriers can inhibit the release of dislocations from the dislocation source. To enable the continuous proliferation of dislocations from the source, it is necessary to increase the externally applied shear stress ${\tau }_0$, provided by USP. Therefore, the externally applied shear stress is a significant factor in the evolution of dislocation barriers into dislocation cells and the subsequent refinement of grains, while the USP process, influenced by the highly random impacts of spheres, greatly enhances the effect of this externally applied shear stress on the surface layer of the alloy, promoting grain refinement and plastic deformation in the surface layer.

4.2. Grain Plasticity and Orientation Analysis

Slip in GH4151 alloy crystals occurs under the applied shear stress of USP. Not all slip systems are simultaneously activated when the crystal is subjected to stress; rather, slip begins when the shear stress in a particular slip system reaches a certain critical value due to the external force from the spheres’ impact. When the external force, slip plane, and slip direction form angles of 45 degrees, the orientation factor reaches its maximum value of 0.5, indicating the highest shear stress and lowest yield strength. This configuration allows for an easier slip and exhibits maximum plasticity, known as the “soft” orientation. The applied external force $F$ and the shear stress $\tau$ in the slip direction can be expressed by Equation (8). When the critical shear stress reaches the threshold value ${\tau }_k$, the crystal starts to slip, and the relationship between the critical value ${\tau }_k$ and the angles $\lambda$ and $\phi$ is given by Equation (9). Compared to the NP specimens, the surface layer of the GH4151 alloy undergoes grain rotation influenced by the random angles of the spheres’ impact [37], resulting in a higher degree of grain rotation variability. The deviation from the 45-degree angle between the slip direction and the normal to the slip plane (Figure 5c) leads to a larger proportion of $\cos \lambda \cdot \cos \phi$, indicating a smaller orientation factor and a deviation from the “soft” orientation characteristic of the NP specimens. This deviation leans towards the “hard” orientation, thereby increasing the resistance to surface deformation [38,39].

$$\tau ={\displaystyle \frac{F\cos \lambda }{A/\cos \phi }}={\displaystyle \frac{F}{A}}\cos \lambda \cdot \cos \phi$$

(8)

$${\tau }_k=\cos \lambda \cdot \cos \phi$$

(9)

Due to the influence of machining and other processes, when the alloy surface experiences significant deformation, the orientations of each grain tend to become more uniform, thereby destroying the original disorder of grain orientations in the polycrystal (Figure 5b), resulting in a preferred orientation and increased texture strength. In most cases, the anisotropy caused by texture is harmful [40], as it causes an uneven deformation in the cold deformation process of the alloy, resulting in significant differences in microstress between grains in various directions (Figure 5a), which affects the dimensional accuracy of the formed components. After USP, the texture strength is weakened, and the randomness of the alloy surface orientation is improved. This is because during USP, the impact angle of the sphere is random. Over a long period of USP, a large number of spheres cause the slip planes and slip directions of each grain to rotate due to the additional shear stress from the impact, resulting in relatively consistent plastic deformation on the surface and the enhanced randomness of grain orientations, thus reducing the high texture strength introduced by previous machining processes [23].

5. Conclusions

This paper discussed the effects of USP on grain refinement and grain orientation behavior in the surface region of GH4151 superalloy. Based on the experimental results and the corresponding analysis, the conclusions arrived at are as follows:

• During the USP process, the plastic deformation of the surface layer of superalloys is accompanied by changes in grain size and orientation. The random impact of the spheres on the surface area promotes grain refinement and grain rotation, enhancing the randomness of grain orientations, reducing the texture strength and the proportion of “soft” orientation distribution;
• The USP treatment induces the formation of dislocation pile-up, where multiple dislocation sources are activated, releasing a series of dislocations. These dislocations accumulate and arrange themselves in a non-uniform manner within the barriers, with a denser arrangement closer to the barriers;
• The equilibrium of the dislocation pile-up is governed by the balance between the externally applied shear stress and the forces exerted by other dislocations. The position of each dislocation within the pile-up is proportional to the total length of the pile-up. The number of dislocations within a pile-up is directly proportional to the externally applied shear stress and the distance between the dislocation source and the barriers. To sustain dislocation proliferation, an increase in the externally applied shear stress provided by USP is necessary. This externally applied shear stress plays a crucial role in the evolution of dislocation pile-ups into dislocation cells and subsequently leads to grain refinement and plastic deformation in the surface layer of the alloy;
• The USP treatment has a significant impact on the distribution of dislocations in superalloys. The understanding of dislocation pile-up behavior and the relationship between externally applied shear stress, pile-up characteristics, and grain refinement is essential for optimizing the USP process and achieving the desired material properties.

However, it is regrettable that this study only conducted experiments and an impact analysis on GH4151 alloy materials and did not involve actual engineering application parts. In future work, the focus will be on aero engine components, including turbine disks. We will conduct multi-process parameter USP experimental research to establish a relationship between the theoretical model of USP process strengthening and experimental data.