Effect of Laser Shock Peening on the Fatigue Life of 1Cr12Ni3Mo2VN Steel for Steam Turbine Blades

Abstract

In the present study, laser shock peening (LSP) was employed to enhance the rotating bending fatigue life of 1Cr12Ni3Mo2VN martensitic stainless steel used in steam turbine blades, addressing the issue of insufficient fatigue performance in these components. The aim of this study was to investigate the effect of LSP on the microhardness, residual stress, and rotating bending fatigue life of 1Cr12Ni3Mo2VN steel samples. The microhardness of LSP-treated samples was increased by 10.5% (LSP-3J sample) and 15.3% (LSP-4J sample), respectively, compared to high-frequency hardening samples. The residual compressive stress of the LSP-4J sample was the largest, reaching −689 MPa, and the affected layer depth was about 800 μm. Fatigue tests showed that the number of cycles at the fracture point for the LSP-3J and LSP-4J samples increased by 163% and 233%, respectively. The fatigue fracture morphology of the four samples showed that the microhardness and residual compressive stress distribution introduced by LSP could effectively inhibit the initiation of surface cracks, slow down the crack growth rate, and improve the rotating bending fatigue life of 1Cr12Ni3Mo2VN.

1. Introduction

As an important part of the energy industry, electricity has become an important support for people’s daily lives and social production. Steam turbines are one of the key equipment in a power plant, and the steam turbine blade is the key component in the steam turbine that converts the thermal energy of high temperature steam into mechanical energy [1]. However, sporadic incidents of turbine blade failures have led to power plant shutdowns and significant financial losses [2,3,4,5]. Failure analysis of many low-pressure last-stage blades of steam turbines revealed that these blades experienced the combined effects of huge centrifugal forces, alternating steam loads, and vibrations during service [6,7,8,9,10,11]. Therefore, the blade material must possess good toughness, strength, and corrosion resistance to withstand such complex service environments. Martensitic heat-resistant stainless steel [12] is commonly used for steam turbine blades due to its high strength, good ductility, toughness, and corrosion resistance. However, the actual service life of blades fabricated from this material still does not meet expectations. For example, Yuwei W. et al. [13] found fatigue cracks in 80 out of 98 low-pressure, final-stage blades from a steam turbine operating for less than two years. Adnyana D.N. et al. [14] documented that the failure of steam turbines within a few years was due to the cracks in the blades under the combined action of cavitation erosion and fatigue. Krechkovska H. et al. [8] attributed premature blade failure to fatigue cracks at the blade edges, leading to a change in the intrinsic oscillation frequency of the blade and resulting in resonant overload. Hence, specific processes and methods are needed to improve the fatigue performance of blade materials and ensure their safety during use.

Many theoretical studies have shown that various mechanical surface treatment techniques can significantly improve fatigue life by altering surface hardness and residual stress to inhibit the generation of cracks, including shot peening, surface carburizing, ultrasonic surface rolling processing, laser shock peening, etc. [15,16,17]. While shot peening enhances the fatigue properties of metallic materials through work-hardening effects and the introduction of residual compressive stresses, it may also induce the formation of microcracks on the material’s surface due to the introduction of large depths of residual stress layers [18,19,20]. Similarly, low-temperature carburizing enhances fatigue performance through residual compressive stresses, but may result in stress relaxation due to the formation of microcracks in the outermost layer [21]. Ultrasonic surface rolling processing [22,23], an advanced surface modification technology, can form a gradient nanostructure on the metal surface to provide superior mechanical properties and enhance material fatigue life. However, due to its structure and principle, it may not be adaptable for complex structural workpieces, such as curved components. As a frontal surface strengthening technology, laser shock peening reduces the growth rate of fatigue cracks by introducing high residual compressive stress [24,25]. Moreover, it improves the fatigue life of metals by enhancing their surface integrity [26]. For example, Li X. et al. [27] found that the residual compressive stress induced by laser shock peening suppressed the initiation of cracks by reducing the effective stress intensity factor and improved the fatigue performance of U75VG steel track welded joints. Compared to shot peening, laser shock peening can introduce greater residual compressive stress and a deeper residual stress-affected layer of more than 1 mm [28]. Compared with surface carburizing, laser shock peening does not compromise the material’s corrosion resistance [29]. Furthermore, laser shock peening is more suitable for strengthening curved surface components such as aviation blades than ultrasonic surface rolling processing [30]. Considering the advantages of LSP over other strengthening methods, researchers have applied LSP to steam turbine blades. For example, simulation and experimental investigations by Fameso F [31] revealed that the energy intensity of LSP was the main influencing factor for increasing the residual stress in blade materials. In addition, Sundar R [32] found that the depth of the residual stress layer generated in the blade root after LSP treatment was more than 900 μm. However, though the blade material introduced a large residual stress after LSP, researchers [31,32,33,34] did not perform fatigue experiments on the strengthened blade material to verify its fatigue life, indicating the need for further research on the influence of LSP.

1Cr12Ni3Mo2VN is a martensitic heat-resistant stainless steel with better high-temperature oxidation resistance and performance stability than ordinary martensitic stainless steel, which is mainly used for the manufacturing of the final-stage blades of the 1000 MW supercritical unit steam turbine [12]. However, the blades made of it had experienced failure prematurely. Therefore, the failure mode of the blade made of 1Cr12Ni3Mo2VN steel was studied by failure detection and simulation analysis in this study. Secondly, aiming at the failure mode, the LSP experiments were designed. The mechanism of LSP to the fatigue life of 1Cr12Ni3Mo2VN steel was investigated by performing a rotating bending fatigue test and fracture morphology detection. The purpose of this study is to apply LSP to steam turbine blade materials, and its effect on improving the fatigue life of 1Cr12Ni3Mo2VN steel is verified, which provides a solid theoretical basis for the development of high-performance turbine blades.

2. Failure Analysis of Steam Turbine Blades

2.1. Analysis of Blade Fracture Morphology

In this study, the blade that experienced failure was composed of 1Cr12Ni3Mo2VN stainless steel and had been in service for approximately 35,000 h. Starting from the sixth year of operation of the steam turbine, cracks were discovered on the inlet edge of 16 out of 94 low-pressure final stage blades, situated 100–140 mm away from the blade shroud band. Figure 1 displays one of the aforementioned failed blades and indicates the position of the crack, and the height of the blade is about 909 mm.

The fracture, which was 130 mm from the blade tip, was observed under the Crossbeam 550 Zeiss scanning electron microscope (Oberkochen, German), and the fracture morphology is shown in Figure 2. Fatigue signs were observed based on these fracture surfaces. Macroscopic observation of the crack fracture shown in Figure 2a revealed a crack initiation region, crack propagation region, and transient crack region. Figure 2b was a magnified observation of the yellow elliptical area b in Figure 2a that revealed the fracture as a deconstruction step with secondary cracks and fatigue strips. The fracture morphology of the blade indicated that the cracks gradually expanded due to the action of obvious alternating loads, and the failure mode of the blade fracture was fatigue.

2.2. Simulation Analysis of Alternating Load

In Section 2.1, the fracture morphology of the blade indicated that the failure mode of the low-pressure, last-stage blade was due to fatigue. Considering that fatigue characterizations such as fatigue striation in Figure 2 and the steam turbine blades are mainly subjected to centrifugal force and alternating steam loads during operation, it can be inferred that the alternating steam load played a significant role in causing the fatigue failure of the blade. Therefore, simulations of stresses and strains of alternating steam loads were performed on the blade model using Abaqus. The research object was 1Cr12Ni3Mo2VN, the low-pressure, last-stage moving blade steel of the steam turbine. According to the requirements of GB/T 8732-2014 for blades made of martensitic steel [35], the mechanical properties of 1Cr12Ni3Mo2VN steel are shown in

Table 1. Mechanical properties of 1Cr12Ni3Mo2VN steel.

Tensile Strength, Rm (MPa)	Yield Strength, Rp0.2(MPa)	Elongation, Δ (%)	Reduction of Area, A (%)	Impact, Kv2 (J)	Hardness, HB
≥1100	≥860	≥13	40	≥54	≤331–363

. The elastic modulus was 198 GPa, the Poisson’s ratio was 0.3, and the density was 7800 kg/m³. In the simulation calculation, it is assumed that the materials are isotropic [36]. The simplified model of the blade is discretized into 203,002 nodes and 185,678 grid cells using Hypermesh 2021 software, of which 175,181 are hexagonal grid cells and 10,497 are pentagonal grid cells. The edge length of the smallest grid cell is 1.32 mm, which is less than the minimum thickness of the inlet side of 2 mm. The mesh was imported into Abaqus, the completely fixed boundary conditions were set at the bottom of the wheel in the simplified model, and the contact pairs were established at the contact surface of the blade and the wheel. An alternating steam load with a pressure of 0.3 MPa on the positive arc surface was applied.

The simulation results of alternating stress and strain of the simplified model are shown in Figure 3. As shown in Figure 3, regarding the increase of steam pressure in a single cycle, the stress and logarithmic strain values at the inlet edge near the tip of the blade (120–260 mm) were the largest. If the supporting effect of the lacing wire were considered, the maximum stress position would move further toward the blade tip, coinciding with the actual crack initiation position. Combining the SEM morphology of the crack fracture of the blade (Figure 2) and the simulation results of the blade’s alternating load, it was concluded that due to the alternating stress of the blade during the rotation process, fatigue cracks could easily form in the inlet edge of the blade 100–140 mm away from the shroud band, leading to failure. Therefore, LSP was performed on the blade material, and the alternating load was equivalent to the rotational bending load to explore the effect of LSP on the rotational bending fatigue life of 1Cr12Ni3Mo2VN.

3. Experimental Methods

The experimental work plan is shown in Figure 4. The 1Cr12Ni3Mo2VN sample was pretreated and subjected to high-frequency hardening (HFH); then, two high-frequency hardening samples were selected for laser shock peening. The untreated, HFH, and LSP samples were prepared by wire cutting, and the residual stress and hardness were measured, respectively. At the same time, the four samples of samples were subjected to the rotating bending fatigue test, the experimental results were recorded, and the fracture morphology was analyzed.

3.1. Material and Sample Preparation

In this study, the steam turbine blade material 1Cr12Ni3Mo2VN steel was selected as the research object. After electric furnace smelting and electroslag remelting, it was forged into a Φ16 mm round bar. According to the reference standard GB/T 8732-2014 [35], the main chemical compositions of 1Cr12Ni3Mo2VN steel is shown in

Table 2. The main chemical compositions of 1Cr12Ni3Mo2VN steel (wt.%).

C	Si	Mn	P	S	Cr	Ni	Mo	N	V
0.1~0.17	<0.3	0.5~0.9	<0.02	<0.015	11~12.75	2~3	1.5~2	0.01~0.05	0.25~0.4

. All samples were quenched and tempered, and their mechanical properties are shown in Table 1. Part of the quenched and tempered sample was treated using high-frequency hardening to make its hardness reach 45 HRC.

In accordance with the ISO 1143 standard [37] for rotating bending fatigue testing of metal materials, the dimensions of the fatigue sample size (Figure 5) were measured in millimeters. All fatigue samples were soaked in industrial ethanol, and the surface was cleaned by ultrasonic waves to reduce experimental errors. Five samples were prepared for each sample in different processes, three of which were used for fatigue experiments. Another was used to test the residual stress along the radial depth at the small diameter of the sample, and the remaining one was used to measure the microhardness along the radial depth at the small diameter of the sample.

3.2. LSP Principle and Process

The LSP experiment was carried out at room temperature (25 °C). The LAMBER-12 pulsed laser system (Beijing, China) with an output wavelength of 1064 nm was used to perform LSP treatment on two HFH samples. When the laser beam irradiated the absorbing layer on the surface of the sample, the generated plasma rapidly expanded to form a high-pressure shock wave, which was propagated into the material through the action of the constraining layer, causing the surface of the material to undergo plastic deformation to obtain residual compressive stress. When the laser-induced plasma shock wave acts on the surface of the specimen and the peak pressure of the shock wave satisfies the Hugoniot elastic limit of 2.0 to 2.5 times, the best impact effect can be obtained [38]. The dynamic elastic limit σ_HEL of the material is expressed as:

$${\sigma }_{HEL}=\frac{1-\upsilon }{1-2\upsilon }{\sigma }_0$$

(1)

In Equation (1): υ is Poisson’s ratio, and υ is taken as 0.3; σ₀ is the yield strength of material, and it is 800 MPa. The dynamic elastic limit of 1Cr12Ni3Mo2VN steel is 1400 MPa, calculated in Equation (1). The laser energies were based on the laser energy range derived from the peak pressure equation for laser-induced plasma shock waves studied by Fabbro et al. [39]. The peak pressure P_max of the laser-induced plasma shock wave is:

$$P_{max}=0.01\sqrt{\frac{\alpha }{2\alpha +3}}\sqrt{Z}\sqrt{I_0}$$

(2)

In Equation (2), α is the efficiency of laser-induced plasma interaction with the sample, and α is taken as 0.2; Z is the composite acoustic impedance of the confining layer medium and the specimen, and it is 3.2 × 10⁵ g/cm²·s; I₀ is the laser power density, its unit is GW/cm². According to 2σ_HEL < P_max < 2.5σ_HEL and Equation (1), the peak pressure range of the 1Cr12Ni3Mo2VN steel material can be obtained as 2800 MPa ≤ P_max ≤ 3500 MPa. The pulse energy W is calculated as follows:

$$W=\frac{D^2·I_0·\tau ·\pi }{4}$$

(3)

In Equation (3), D is the laser spot diameter; τ is the pulse width. Therefore, during LSP treatment, parameters were set as follows: a pulse width of 10 ns, a laser beam with a diameter of 3 mm, a pulse energy of 3 J and 4 J, and a repetition rate of 1 Hz. The LSP schematic diagram is shown in Figure 6. Black glue with a thickness of 0.13 mm was used as the absorbing layer, and a water layer with a thickness of 2 mm was used as the constraining layer to prevent the surface of the sample from being ablated.

3.3. Surface Integrity Measurements

Microhardness was measured using a FALCON 500 Vickers microhardness tester (Maastricht, The Netherlands) with a test load of 100 g and a dwell period of 10 s. Cross-sections of the LSP samples (the yellow area in Figure 6) in the minimum diameter (Φ6 mm) were prepared for coarse grinding, precise grinding, and polishing.

A high-power HDS-I X-ray diffraction residual stress tester (Handan, China) equipped with a high-precision Cr target was used to measure the residual stress on the surface of the sample. The residual stress was measured using the roll fixed Ψ method, and the Ψ angles were 0° and 45°. The scanning start angle and end angle of 2θ are 159° and 149°, respectively. The residual stress in the axial direction of different depths was measured using the electrolytic polishing method, and the area at the minimum diameter (Φ6 mm) of the sample was taken as the electrolytic area. The black tape was used to protect the non-measurement area of the arc section, which was conducive to the depth positioning of the electrolytic area by the micrometer. Layer by layer Electro-polishing was carried out at a voltage of 15 V using a saturated NaCl solution. The layer removal size was 25 μm in the first 200 μm depths to get more detailed stress information of the subsurface. Then, the spacing increased from 50 μm to 100 μm in the depth range of 200–1000 μm.

3.4. Rotating Bending Fatigue Tests

According to the ISO 1143 rotating bar bending fatigue test method, four samples were tested using QBWP-6000J supported beam rotating bending fatigue testing machine (Changchun, China). At room temperature (25 °C), the loading force was set to 70 N, the cycle frequency was 50 Hz, and the stress ratio was −1. The number of cycles of all fatigue samples under the specified stress conditions was recorded. At least 3 samples were prepared for each set of experiments to reduce errors. Four sets of fracture samples were prepared using the wire electric discharge method and subsequently ultrasonically cleaned with alcohol to prepare scanning electron microscope samples. The fracture surfaces of all fatigue fractures were observed by a Crossbeam 550 Zeiss SEM at 10–15 kV.

4. Results and Discussion

4.1. Microhardness Distribution

It is now understood that the microhardness of a material is one of the indicators that affect its fatigue life [40]. The microhardness distributions of samples in different processes are shown in Figure 7. It can be seen that that the microhardness of the untreated sample fluctuated between 540 HV_0.1 and 550 HV_0.1 after high-frequency hardening. After LSP, the surface hardness of the impact-strengthened samples under different laser energies increased by 10.5% (LSP-3J sample) and 15.3% (LSP-4J sample) compared with the HFH sample. As depicted by the blue curve (LSP-3J sample) and the green curve (LSP-4J sample), the microhardness gradually decreased with increased depth and finally approached the matrix hardness of the HFH sample. The surface hardness of the LSP-4J sample was the highest (628.5 HV_0.1), and the affected layer depth was about 800 μm. As shown in Figure 7, LSP could significantly increase the microhardness of the material surface and form a hardness field of hundreds of microns on the surface.

4.2. Residual Stress Distribution

It has been established that the laser-induced shock wave creates a residual compressive stress on the material’s surface during the LSP process. This residual compressive stress effectively counteracts tensile stresses and works in conjunction with the surface hardening layer to prevent the initiation and propagation of fatigue cracks [41]. The residual stress distribution of samples in different processes is shown in Figure 8. The residual compressive stresses on the surface of the untreated sample and the HFH sample were −240 MPa and −360 MPa, respectively, and gradually fluctuated in the range of −80 MPa to 100 MPa after 50 μm; this phenomenon was due to the finishing and polishing processes during the preparation phase of the sample [26,42]. In addition, the surface residual stresses of LSP-3J and LSP-4J samples reached −646.7 MPa and −689 MPa, respectively. With increasing depth, the residual stress in the LSP-4J sample gradually approaches the stress level of the material’s matrix at approximately 800 μm, corresponding to its microhardness distribution. Compared to the untreated and HFH samples, the LSP-treated samples exhibited significantly larger residual compressive stress on the surface layer. Moreover, this residual compressive stress extended to a depth of approximately 800 μm. However, as the depth increased, the gradient of the residual compressive stress became smaller, gradually decreasing towards the substrate level.

4.3. Fatigue Life

The rotating bending fatigue test was conducted on the untreated, HFH, and LSP samples. These samples were subjected to cycling at room temperature (25 °C) with a load of 70 N until they experienced failure. As shown in Figure 9, the average fatigue lives of the untreated samples, the HFH samples, LSP-3J samples, and LSP-4J samples were 59,679, 65,962, 14,932, and 193,264 cycles, respectively. Compared with the untreated samples, HFH, LSP-3J, and LSP-4J samples exhibited significant improvements in their cycle numbers, increasing by 10.5%, 163%, and 233%, respectively. The experimental results demonstrated that though HFH exhibited enhancement in the fatigue life of 1Cr12Ni3Mo2VN, the improvements achieved through LSP were more significant. As shown in Figure 7 and Figure 8, the LSP-treated samples, specifically LSP-3J and LSP-4J, possessed a deep surface-hardened layer and a large residual compressive stress field, which were absent in the untreated and HFH samples. These changes were considered key factors contributing to the observed improvement in fatigue life. Yong W. et al. [26] emphasized the beneficial impact of surface hardening and the residual compressive stress introduced by LSP, resulting in a great increase in the fatigue life of their specimens. Combining the data from Figure 7, Figure 8 and Figure 9, it can be concluded that both the microhardness and the residual stress distribution significantly affect the fatigue life of the material. However, the enhancement achieved through the residual compressive stress field is particularly noteworthy. In conclusion, the reasonable application of LSP can effectively improve the fatigue life of 1Cr12Ni3Mo2VN martensitic stainless steel.

4.4. Fracture Analysis

The surface topographies of fatigue specimens in different processes are shown in Figure 10, Figure 11, Figure 12 and Figure 13. From the macro view of the fracture (Figure 10a, Figure 12a and Figure 13a), the material fatigue damage could be divided into three regions, namely the crack initiation region (Region I), the crack propagation region (Region II), and the transient crack region (Region III). These regions corresponded to the three stages of crack initiation, crack stable growth, and crack growth out of the stable state experienced during the fatigue failure of the material.

Figure 10 shows the fracture morphology of the untreated sample, with Figure 10b,c providing enlarged views of the marked areas b and c in Figure 10a, respectively. Additionally, Figure 10d offers an enlarged view of the marked area d in Figure 10c. As depicted in Figure 10a, the macro-fracture of the untreated sample exhibits the characteristics of multiple fatigue sources. Figure 10b provides a closer view of the fatigue striations in the crack propagation region. Furthermore, Figure 10c reveals the radiation extending from the crack initiation point into the material, while micro-cracks can be observed in Figure 10d. This observation suggests that the cracks originated on the surface and then propagated radially, ultimately leading to instant fracture. In summary, the cyclic bending load induced the slipping of crystal grains on the surface, causing stress concentration, the formation of microcracks, and the development of fatigue sources.

Figure 11a displays the macroscopic fracture morphology of the HFH sample. Figure 11b,c are the corresponding magnified views of marked areas b and c, respectively, in Figure 11a. The fatigue source area was visible on the surface in Figure 11a, showing multiple initiation points for fatigue cracks. The fatigue source area and the crack propagation region were relatively flat, while the transient crack region appeared non-uniform. The fatigue crack propagated upward, exhibiting a radiation pattern, until it fractured completely. Figure 11b shows the typical fatigue striation after the crack propagation region was enlarged. Figure 11d offers an enlarged view of the marked area d in Figure 11c. The radial lines in Figure 11c and the notch feature in Figure 11d indicated that the cracks originated from the discontinuous positions of the machining tool marks. From Figure 11, it was inferred that the rotating bending load subjected the surface to higher tensile stress than the interior, resulting in increased stress concentration due to the unevenness of the surface tool marks. This condition increased susceptibility to the generation of fatigue sources and the initiation of fatigue cracks.

Figure 12 depicts the fatigue fracture morphology of the LSP-3J sample. It was observed that the fatigue source area covered a significant portion in the lower left corner in Figure 12a, displaying multiple cracks growing together. These cracks propagated inward in a radiation pattern until a complete fracture occurred. Compared to the first two groups of samples, one fatigue source emerged in the subsurface, and the crack propagation region was relatively larger and flatter. Other fatigue sources exhibited clear fracture steps in Figure 12a, suggesting that crack initiation resulted from stress concentration caused by the surface tool mark’s discontinuity. Figure 12b is the magnified view of the marked area b in Figure 12a and shows the fatigue striation spacing measured about 0.63 μm smaller than the HFH sample (about 1.12 μm). Figure 12c offers an enlarged view of the marked area c in Figure 12a,d, which is the magnified view of Figure 12c. It can be seen that the fatigue source appeared in the subsurface below the surface, indicating that the residual compressive stress and microhardness introduced by LSP overcame the surface stress concentration. Consequently, the fatigue source transferred from the sample’s surface to the subsurface, extending the sample’s life.

Figure 13a displays the macroscopic fracture morphology of the LSP-4J sample. Compared to the LSP-3J sample, the number of fatigue sources in the LSP-4J sample was lower. The fatigue sources were observed on both the surface and subsurface in Figure 13a, indicating a multi-source fatigue pattern. The surface fatigue source propagated internally in a splitting manner. Figure 13b shows an enlarged view of the fatigue strips in the crack propagation region in Figure 13a with a spacing of about 0.53 μm. As shown in Figure 13c, which is the corresponding magnified view of the marked area c in Figure 13a, the subsurface fatigue source exhibited a clear radiation trend in the crack propagation region, eventually leading to an instantaneous fracture in the uppermost area. The spacing of the fatigue striations on this surface decreased with increasing laser energy. Figure 13c shows that the fatigue source due to surface stress concentration was also transferred to the subsurface due to the residual compressive stress and microhardness. Figure 13d offers an enlarged view of marked area d in Figure 13a, where the discontinuous surface machining tool mark can be seen. Combining the fatigue fracture diagrams of all four sample groups, it can be inferred that the LSP-induced residual compressive stress and microhardness not only improved the local stress concentration caused by discontinuous tool marks on the surface, but also reduced the spacing of fatigue striations, thereby slowing down the movement of continuous slip bands. As a result, reasonable LSP treatment could effectively enhance the rotating bending fatigue life of 1Cr12Ni3Mo2VN.

5. Conclusions

In this study, laser shock peening was first carried out on the 1Cr12Ni3Mo2VN steel after high-frequency hardening. The microhardness and residual stress of samples subjected to different processes had been compared and analyzed. Subsequently, rotating bending fatigue tests were performed on these samples. The results, combined with the fracture morphology obtained from the fatigue test, demonstrate that laser shock peening can effectively enhance the fatigue life of 1Cr12Ni3Mo2VN martensitic stainless steel. The key conclusions drawn from our investigation are as follows:

(1)

The LSP process led to a notable rise in surface hardness. Remarkably, the LSP-4J samples displayed the highest surface hardness of 628.5 HV_0.1, extending to an approximate depth of 800 μm.

(2)

The LSP process introduced high residual compressive stress and large depths of residual stress layers in the surface. The LSP-4J samples had the highest surface residual compressive stress of -689 MPa, and the depth of the residual compressive stress layer in the LSP-4J sample nearly corresponded with the microhardness distribution, approximately 800 μm in depth.

(3)

The fatigue life of 1Cr12Ni3Mo2VN steel had significant enhancement after LSP treatment. Compared to untreated samples, the increase in cycles to fracture for HFH samples was not substantial. In contrast, the LSP-3J and LSP-4J samples had impressive increases of about 163% and 233% in cycles to fracture, respectively.

(4)

The microhardness and residual compressive stress induced by LSP were crucial in mitigating local stress concentration caused by surface discontinuities. Furthermore, these enhancements reduced the spacing of fatigue striations, resulting in a slower crack expansion rate, and ultimately improved the fatigue life of 1Cr12Ni3Mo2VN steel.