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Article

Effects of Pre-Peening on Fatigue Performance of Gas-Nitrided SCM 440 Steel

by
Hao Chen
1,
Tai-Cheng Chen
2,3,
Wen-Han Chen
4,
Hsiao-Hung Hsu
5 and
Leu-Wen Tsay
1,4,*
1
Department of Optoelectronics and Materials Technology, National Taiwan Ocean University, Keelung 202301, Taiwan
2
Department of Material Research, National Atomic Research Institute, Taoyuan 325207, Taiwan
3
Neutron Scattering Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831, USA
4
Vincent Vacuum-Tech Co., Ltd., Taoyuan 326019, Taiwan
5
Iron and Steel Research and Development Department, China Steel Co., Ltd., Kaohsiung 812401, Taiwan
*
Author to whom correspondence should be addressed.
Metals 2025, 15(10), 1118; https://doi.org/10.3390/met15101118
Submission received: 9 September 2025 / Revised: 2 October 2025 / Accepted: 5 October 2025 / Published: 9 October 2025
(This article belongs to the Special Issue Fatigue and Fracture Behavior of Metallic Components and Structures Under Various Loading Conditions)
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Versions Notes

Abstract

Gas nitriding was implemented in the current work at a constant nitrogen potential (KN) of 2.0 for 8 h to enhance the fatigue properties of SCM 440 steel, and the results were compared with those of the substrate tempered at the nitriding temperature (475 °C). Fine particle peening (FPP) prior to nitriding imposed a refined structure and induced compressive residual stress (CRS) in the near-surface peened zone. The fine-grained structure provided numerous paths to enhance nitrogen diffusion inwards during nitriding. The compound layer formed on the nitrided SCM 440 steel primarily comprised a mixture of Fe3N and Fe4N; however, the pre-peened and nitrided (SPN) specimens exhibited a higher proportion of Fe3N and a thicker compound layer than the non-peened and nitrided (NPN) counterparts. In addition, FPP prior to nitriding increased both the case depth and the magnitude of the CRS field compared with nitriding alone. The fatigue limits of the substrate (SB), NPN, and SPN samples were approximately 750, 1050, and 1400 MPa, respectively. Gas-nitriding at 475 °C significantly improved the fatigue performance of SCM 440 steel. Moreover, pre-peening prior to nitriding further enhanced fatigue strength and life of the treated SCM 440 steel by introducing a deeper case depth and higher CRS field. Multiple cracks initiation at the outer surface of the SB sample accounted for its lowest fatigue limit among the tested samples. Surface microcracks and pits on the surface of the NPN specimen would be crack initiation sites and harmful to its fatigue resistance. These surface dents were considered to be responsible for fatigue crack initiation in the SPN specimens. Therefore, polishing after nitriding to reduce surface roughness and/or microcracks was expected to further increase the fatigue resistance and the reliability of nitrided SCM 440 steel.

1. Introduction

Metal fatigue, wear, corrosion, and their combinations are among the primary causes of failure in engineering components. The surface structure and properties of a component play a critical role in its service performance, as cracks are prone to initiate at the surface during prolonged operation. To enhance fatigue resistance under cyclic loading, both mechanical surface treatments (e.g., shot peening [1,2,3,4,5]) and thermo-chemical treatments (e.g., nitriding [6,7,8,9,10]) are widely employed. Gas nitriding, one of the most useful thermo-chemical processes, is highly effective in improving the abrasion and fatigue resistance of steel components such as extrusion dies, forging dies, valves, and gears [11,12,13]. This enhancement is mainly ascribed to the dissolution of nitrogen into the substrate and the precipitation of nitrides, which systematically increase surface hardness and introduce beneficial compressive residual stress (CRS) in the nitrided layer.
Nitriding has been widely demonstrated to enhance the fatigue strength and service life of steel components [6,7,8,9,10]. For example, gas nitriding significantly extends the high-cycle fatigue life of 4140 steel, with the improvement being more pronounced under lower stress levels [14]. For ion-nitrided 4340 steel, increasing the case depth and surface hardness leads to marked gains in fatigue strength and life compared with untreated steel [7]. Similarly, gas nitriding of 4135 steel at 580 °C produces thicker compound layers and stronger CRS fields with prolonged treatment time, both of which contribute to improved fatigue performance [9]. Plasma nitriding of 4140 steel also shows a similar trend, with CRS increasing alongside case depth as the nitriding duration is prolonged [10]. An important microstructural effect of nitriding is the shift in fatigue crack initiation sites from the surface to the subsurface, which is closely associated with the observed improvement in fatigue resistance [8,15,16,17,18]. In the high-cycle fatigue regime, subsurface crack initiation at inclusions is the dominant failure mode in nitrided 4340 steel [17,19]. However, in the low-cycle fatigue regime, the lower ductility of the compound layer relative to the steel core can be detrimental, reducing fatigue resistance in nitrided 4140 steel [20].
Fine particle peening (FPP), also referred to as micro-shot peening, involves bombarding a material surface with shot particles smaller than 200 μm in diameter [21,22,23,24]. Compared with conventional shot peening, FPP produces lower surface roughness [21]. Moreover, refined structure and high CRS are present in the surface layer of FPP-treated material [21]. It is noticed that FPP produces higher surface hardness, greater CRS but limited affected depth relative to conventional shot peening [25,26,27,28]. Notably, the magnitude of CRS induced by FPP increases with the hardness of the substrate, leading to a substantial enhancement in the fatigue limit, as demonstrated for 4135 steel [23]. It is reported that a fine-grained microstructure formed in the severely peened layer [5,27,28,29] plays a critical role in subsequent nitriding treatments [17,19,30]. The abundance of grain boundaries generated in this layer promotes nitrogen diffusion during nitriding, enabling deeper case depth. For example, severe peening prior to gas nitriding of low-alloy steel has been shown to facilitate the formation of a thicker compound layer [17]. Similarly, in 4140 steel, a nano-grained surface layer produced by surface mechanical attrition treatment (SMAT) significantly enhances nitrogen diffusion, resulting in a much thicker compound layer after ion nitriding compared with untreated steel [30]. FPP prior to gas-nitriding not only promotes nitrogen diffusion but also reduces the formation of pores and microcracks in the compound layer of gas-nitrided 4135 steel [22,24]. This pre-treatment has been shown to lower the friction coefficient, reduce wear loss, and further enhance the fatigue strength of nitrided 4135 steel relative to its un-peened counterpart [22,24].
The fatigue performance of a nitrided steel is governed by the core strength and the case depth, which are strongly dependent on the nitriding temperature and duration. Gas nitriding is typically conducted in an NH3-containing atmosphere at 500–600 °C. SCM 440 (AISI 4140) steel has been used extensively in the transmission parts. Surface modifications are applied to increase the wear and fatigue resistance of AISI 4140 steel. In prior work [19], peened+ gas-nitrided 4140 steel, which is nitrided at 520 °C for 8 h, has about the same fatigue strength of 1000 MPa as the quenched + 200 °C tempered 4140 substrate. Sub-surface crack initiation is responsible for the fatigue failure of the peened+ gas-nitrided 4140 steel [19]. With an increase in core strength by decreasing the nitriding temperature, the fatigue strength of nitrided steel is expected to be upgraded. In the current study, SCM 440 (AISI 4140) steel, with or without FPP as a pre-treatment, was gas-nitrided at a relatively low temperature of 475 °C for 8 h. As compared with conventional practice, lowered nitriding temperature in this work can increase the substrate fatigue strength while forming a nitrided surface layer, thereby enhancing fatigue strength/life relative to un-nitrided specimens and those nitrided at higher temperatures. Surface roughness was measured using a three-dimensional optical profiler. Rotating bending fatigue tests were conducted in laboratory air to assess both the combined and individual effects of FPP and gas nitriding on fatigue performance. Scanning electron microscopy (SEM) was utilized to examine the fatigue-fractured features, whereas electron backscatter diffraction (EBSD) mapping was applied to characterize their detailed microstructures.

2. Materials and Experiments

2.1. Preparation of Samples

SCM440 (AISI 4140) steel bars with a diameter of 10 mm were used as the experimental material in this study. The bars were supplied by China Steel Corporation, and their chemical composition is presented in Table 1. Figure 1 displays the gas-nitriding procedure with or without FPP. Figure 2 shows the flowchart showing the sample preparation used in current work. SCM440 (AISI 4140) steel bars were solution-treated in an industrial vacuum furnace (Vincent Vacuum-Tech Co., Ltd., Taoyuan, Taiwan) at 865 °C for 0.5 h, followed by oil quenching (Figure 1). The quenched specimens were subsequently tempered at 200 °C for 30 min, resulting in a hardness of approximately HV 500. A portion of these 200 °C tempered specimens underwent FPP at room temperature prior to nitriding and designated as shot-peened (SP) specimens, whereas those without FPP were designated as non-peened (NP) specimens. FPP was performed using amorphous powders as shots, with particle sizes between 80 and 120 μm, under an air pressure of 4 atm. The peening intensity, measured with an N-type Almen specimen, was approximately 0.126 mm in height. Gas nitriding was conducted in a controlled KN furnace at 475 °C for 8 h using NH3 as the nitriding gas and N2 as the carrier gas at a pressure of 600 Torr. KN was defined as the ratio of the partial pressures of NH3 and H2 in the nitriding atmosphere, expressed as KN = P(NH3)/P(H2)3/2. Depending on the pre-treatment, nitrided specimens were designated as SPN (shot-peened + nitrided) or NPN (non-peened + nitrided). For comparison, un-nitrided SCM440 steel specimens, designated as the substrate (SB), were austenitized and tempered at 475 °C/1 h to replicate the thermal history associated with gas nitriding.

2.2. Hardness Measurement and Fatigue Testing

Microhardness profiles of the investigated specimens were determined using an MVK-G1500 Vickers hardness tester (Mitutoyo, Kawasaki, Japan) with a 300 gf load and a dwell time of 15 s. To evaluate the nanohardness of the compound layer and the adjacent diffusion zone, a Hysitron TI 980 TriboIndenter (Bruker, Billerica, MA, USA) with a maximum load of 2000 μN was employed. To evaluate the combined effects of FPP pre-treatment and low temperature nitriding on fatigue resistance of SCM440 steel. Rotating bending fatigue tests were implemented in laboratory air with a frequency of 1500 cycles min−1. For comparison, identical tests were performed on the substrate tempered at 475 °C. Fatigue strength (S)–cycle number (N) (S–N) curves for each condition were constructed from repeated measurements. The geometry of the rotating bending fatigue test sample is presented in Figure 3.

2.3. Microstructural Observation

The surface topography of specimens, with and without FPP, was characterized using a VHX-7000 digital microscope (Keyence, Osaka, Japan). Phase analysis of the nitrided layer was performed using a D2 Phaser X-ray diffractometer (Bruker, Billerica, MA, USA) with Cu Kα radiation over a 2θ range of 20–100° at a scanning speed of 1°/min. A JSM-7100F field-emission SEM (JEOL, Tokyo, Japan) was employed to inspect both the cross-sectional microstructures and the fracture morphology of fatigued specimens. In addition, SEM equipped with NordlysMax2 EBSD detector (Oxford Instruments, Abingdon, UK) was employed to analyze the detailed phases within and near the compound layer of the gas-nitrided specimens. The chemical compositions in distinct zone of the sample were characterized by a JXA-8200 electron probe microanalyzer (EPMA, JEOL, Tokyo, Japan) operated under 1 micro spot size, 15 KV, 10−8 A.

2.4. Measurement of Residual Stress

Residual stress was evaluated with a µ-X360s residual stress analyzer (Pulstec, Hamamatsu, Japan) using Cr Kα radiation, which collected all diffracted beams from the surface of the irradiated specimen. Stress values were calculated by using the cos α method [31,32]. Lattice distortion was assessed, based on the full width at half maximum (FWHM) of the (103) Fe2N peak (2θ = 135.5°) and the (211) α peak (2θ = 156.4°) on the Debye ring. The difference in 2θ angle between the specimen’s Debye ring and that of the reference phase was used to compute the residual stress. The distribution of residual stress in the thickness direction was obtained by removing the surface layer of the sample using an EP-3 electrochemical polisher (Pulstec, Hamamatsu, Japan). Electro-polishing was interrupted periodically until the desired depth of material removal was obtained, which was measured by using a pin detector.

3. Results

3.1. Phase and Surface Morphology of the Nitrided Sample

The phases detected in the nitrided specimens (NPN and SPN) by XRD are demonstrated in Figure 4. In the NPN specimen (Figure 4a), the XRD pattern revealed primarily α-Fe along with Fe3N and Fe4N nitrides. In contrast, the α-Fe peaks disappeared in the SPN specimen (Figure 4b), where predominant Fe3N and a small amount of Fe4N were detected on the outer surface. These results suggest that FPP prior to nitriding promoted the formation of Fe-nitrides at the surface compared with the NPN condition. Figure 5 displays the surface appearance and three-dimensional topography of the inspected samples. SEM morphology of the NPN and SPN samples is shown in Figure 5a,c. According to the surface feature, the NPN sample displayed a relatively flat surface interspersed with microcracks and fine pits (Figure 5a). In contrast, FPP produced numerous shallow dents on the SPN surface, resulting in a typical peened morphology without observable microcracks. The three-dimensional surface topography (Figure 5b,d) was assessed by using Sa (arithmetical mean height), Sp (maximum peak height) and Sv (maximum pit depth). For the NPN specimen, the respective values of Sa, Sp, and Sv were 0.52, 3.02, and 2.58 μm. With FPP, the Sa, Sp, and Sv of the SPN specimen were of slightly higher values of 0.64, 3.05, and 2.75 μm in sequence. It was obvious that both nitrided specimens had similar surface textures, and FPP prior to nitriding did not significantly increase surface roughness.

3.2. EPMA Analysis of the SPN Sample

As listed in Table 2, the chemical compositions at various depths beneath the top surface of the SPN sample are measured by using EPMA. The N content reached as high as 4.27 wt% in the outermost region of the compound layer, decreasing to 3.76 wt% at a depth of approximately 6 μm. Within the diffusion zone, the N content declined from 0.34 wt% at 20 μm to 0.26 wt% at 50 μm. At greater depths into the substrate, the N concentration at 200 μm from the surface decreased to the same level as the substrate. The concentrations of other alloying elements showed minimal variation within the nitrided depth. The elevated N content near the outermost surface of the nitrided SCM440 steel was attributed to the formation of Fe-nitrides. In contrast, the N concentration in the diffusion zone immediately beneath the compound layer was substantially lower than that in the compound layer.

3.3. Hardness Profile of the Nitrided Specimens

Figure 6 illustrates the microhardness profile and nanohardness distribution measured from the nitrided surface toward the specimen center in both NPN and SPN conditions. As shown in Figure 6a, the surface region of nitrided SCM440 steel exhibited much higher hardness than the core, regardless of FPP treatment. The hardness gradually decreased from approximately HV 700 at the surface to HV 350 in the substrate. The case depth was defined as the distance from the outermost surface to the site where the hardness exceeded the substrate hardness (HV 350) by more than 50 HV. According to the specified hardness of HV 400, the case depth for NPN and for SPN samples was about 115 and 135 μm, respectively. This indicates that the SPN specimen possessed a slightly greater case depth than the NPN specimen. The nanohardness values measured in the surface region are presented in Figure 6b,c. The compound layer (CL, Fe-nitrides) of both specimens exhibited nanohardness in the range of 9.5–10 GPa, with the SPN layer being marginally harder than that of the NPN specimen. Moreover, at the same distance from the interface between the compound layer and diffusion zone (DZ), the DZ of the SPN specimen consistently possessed higher nanohardness than that of the NPN one.

3.4. EBSD Analysis

Figure 7 presents EBSD results of the surface-zone microstructures in cross-section for the NPN and SPN specimens. The band contrast maps (Figure 7a,b) illustrate the overall microstructure, the inverse pole figure (IPF) maps (Figure 7c,d) display grain orientations indicated by distinct colors, and the phase maps (PM) (Figure 7e,f) identify the phases present in the examined regions. In the NPN specimen (Figure 7a), the BC map reveals a thin surface layer composed of an aligned, refined granular structure. This layer can be further divided into two subzones: a relatively coarse-grained zone overlying a fine-grained zone. The layer was uneven in thickness, ranging from 1.3 to 3.0 μm. Beneath it, lath martensite packets with distinct orientations were observed, showing the characteristic of tempered SCM440 steel. In contrast, the SPN specimen (Figure 7b) exhibited a refined structure of uniform thickness (~8 μm) on its outer surface, consisting mainly of elongated grains. This refinement is consistent with previous reports that FPP induces nanostructures in the severely peened zones of AISI 4140 steel [19].
The IPF maps (Figure 7c,d) confirm that the surface layers of both specimens consist of fine grains in various orientations, with a tendency for alignment in parallel directions. The phase map of the NPN specimen (Figure 7e) reveals a thin compound layer, comprising Fe3N on top and Fe4N beneath. The SPN specimen exhibited the same phase constituents (Figure 7f) but with a significantly thicker compound layer than the NPN specimen. These phase analysis results are consistent with the XRD findings. As shown in Figure 6a, the hardness profiles indicate that the microhardness in the diffusion zone beneath the compound layer exceeded HV 600. However, the BC and phase maps did not show visible Fe-nitride precipitates in the diffusion zone; instead, lath packets of tempered martensite were observed. Those results indicated that the Fe-nitrides present in the diffusion zone were beyond the resolution limit of the EBSD technique employed in this study.

3.5. Residual Stress Measurements

Residual stress distributions as a function of case depth for the NPN and SPN specimens are shown in Figure 8. High CRS around the surface was advantageous to retard fatigue crack initiation from the external surface. As shown in Figure 8, both specimens exhibited the similar profiles, having peak stress in the subsurface zone. The peak stress was −500 MPa for the NPN sample and reached −700 MPa for the SPN sample. At a depth of ~100 μm from the surface, CRS of the NPN and SPN samples was about −150 and −300 MPa, respectively. In addition, the CRS field of the SPN specimen was obviously greater than the NPN specimen, which was beneficial to increase its fatigue resistance. Previous studies [19,24,26,28] have pointed out that FPP or micro-shot peening can produce high CRS into the severely peened zone, but within a restricted depth. The result also indicated that FPP prior to gas-nitriding could generate a deeper and higher CRS into the treated material, as compared with the non-peened + nitrided one.

3.6. Fatigue Tests

Figure 9 displays the S–N curves of the investigated specimens, as determined by rotating bending fatigue tests. After tempering at 475 °C/1 h, the SCM440 steel maintained a uniform core hardness of approximately HV 350. For the 475 °C tempered substrate (SB), the fatigue strength decreased with increasing cyclic stress, and the fatigue limit was approximately 750 MPa. Gas nitriding alone significantly enhanced the fatigue performance: the fatigue limit of the NPN specimen increased to ~1050 MPa, representing a marked improvement over the SB condition. The combined treatment of FPP and gas nitriding yielded the best performance, with the SPN specimen exhibiting a fatigue limit of ~1400 MPa, substantially higher than both the NPN and SB specimens. These results demonstrate that FPP prior to nitriding at 475 °C markedly improves the fatigue strength and life of SCM440 steel. Notably, the presence of the brittle compound layer, expected to improve wear resistance, did not have a markedly detrimental effect on the fatigue properties of the nitrided steel.

3.7. Fractured Surface Examinations

The fatigue-fractured morphology of SCM440 steel (SB) tested under a peak cyclic stress of 750 MPa is shown in Figure 10. This stress level typically initiated fatigue cracking at the external surface, with cracks propagating inward. In the SB specimen fatigued at 750 MPa, crack initiation occurred at the external periphery (Figure 10a). In a macro-view, a gradual change in appearance from a smooth zone with the directional traces transited to a coarse feature inter-dispersed with microcracks at the depth about 2.0 mm from the outer surface, then final rupture occurred at the depth about 2.8 mm from the crack initiation site. Quasi-cleavage dominated the fatigue fracture in the crack initiation zone of the SB sample (Figure 10b). In fact, vague fatigue striations could be observed in this region, as shown in Figure 10b. As the crack propagated inward, the crack growth rate increased, accompanied by the formation of tear ridges and a few secondary cracks (Figure 10c). At depths of around 2.0 mm from the outer surface, secondary cracks became more prevalent (not shown). Within a very short distance, the fracture appearance transitioned from quasi-cleavage to a dimple fracture morphology (Figure 10d).
Figure 11 presents the fatigue-fractured morphology of the NPN specimen subjected to cyclic stress at 1050 MPa. Detailed examinations revealed that all fractured NPN specimens exhibited crack initiation at the external surface followed by inward propagation. The macroscopic fracture surface (Figure 11a) showed crack initiation and growth from the outer surface, with surface roughness increasing as the crack advanced into the interior. A vague, thin layer was observed along the outer profile of the fracture surface (Figure 11a). Closer examination of the outer surface zone (Figure 11b) revealed that the rubbed profile was enveloped by a very thin compound layer surrounding the profile, which may be attributed to the effect of CRS. Beneath the rubbed zone, quasi-cleavage features accompanied by tear ridges, similar to those in Figure 10c, were observed. Compared with the SB specimen (Figure 10a), the higher applied stress in the NPN specimen reduced the extent of the smooth crack growth zone and produced more extensive and deeper microcracks prior to final rupture (Figure 11a). At a crack length of ~0.7 mm from the surface, predominant quasi-cleavage fracture interspersed with microcracks, mostly oriented normal to the crack growth direction, was observed (Figure 11c). At a depth of ~1.7 mm from the external surface, numerous deep secondary cracks were evident. As the crack propagated further inward, the proportion of quasi-cleavage fracture decreased while dimple fracture increased (Figure 11d). In this transition zone, deep secondary cracks along prior austenite grain boundaries were frequently observed before final rupture.
Figure 12 shows the fatigue-fractured morphology of the SPN specimen at various locations from the external surface. The results indicate that fatigue fracture in the SPN specimen was more likely to initiate at the surface and less likely to originate from subsurface inclusions. Under cyclic stresses at or above 1450 MPa, fatigue cracks initiated from the outer surface (Figure 12a). The macroscopic fracture appearance of the SPN specimen (Figure 12a) was similar to that of the NPN specimen (Figure 11a). Higher-magnification observations of the outermost zone (Figure 12b) revealed that the compound layer exhibited no signs of cracking or spalling and remained firmly adhered to the surface. Beneath the compound layer, lath structures with distinct orientations were observed (Figure 12b), corresponding to tempered martensite. Apart from the outermost zone, the fatigue-fractured features of the NPN and SPN specimens were similar at corresponding locations, as shown in Figure 11c,d.
It is widely recognized that subsurface fatigue crack initiation is strongly influenced by the amount and distribution of inclusions within the material. In the SPN specimen, subsurface fatigue crack initiation characterized by a fish-eye zone was also observed (Figure 12c). In such cases, the crack propagated inward more rapidly than it extended outward toward the surface. As shown in Figure 12d, a cleavage-like fracture pattern with emanating traces was observed from the inclusion located in the fish-eye zone. Examination of the surface region of the fractured SPN specimen (Figure 12e) revealed that the compound layer exhibited a rubbed appearance but remained intact, while the underlying diffusion zone displayed lath feature interspersed with limited intergranular fracture. As the crack propagated out of the fish-eye zone toward the interior (Figure 12a), a change in fracture morphology occurred, with detailed features shown in Figure 12f. Within the fish-eye zone, the fracture surface appeared relatively smooth with distinct tear ridges, while outside this region it was noticeably rougher and contained numerous microcracks. The lath packet structures inside the fish-eye zone were replaced by irregular fine facets outside it, resulting in increased surface roughness. These microcracks became longer and coarser with increasing distance from the fish-eye zone and were often oriented normal to the crack growth direction (Figure 12f). This change in fracture morphology is not yet fully understood but may be related to the increased crack growth rate. As the crack extended further into the interior, the fracture features resembled those observed at corresponding locations in the NPN specimen.

4. Discussion

Surface appearance and roughness values for the NPN and SPN specimens are shown in Figure 5. Although the SPN specimen had slightly high surface roughness, the difference between the NPN and SPN specimens was limited. Pre-peening with fine particles prior to gas nitriding produced only a minor increase in surface roughness relative to the NPN condition. As illustrated in Figure 5d, the relatively high roughness of the SPN specimen resulted from fine surface dents and debris. By contrast, Figure 5c reveals microcracks at the outermost surface of the NPN specimen, which are anticipated to adversely affect its fatigue resistance. Such surface microcracks have also been reported in conventional gas-nitrided steels.
EPMA analysis revealed the nitrogen concentrations in distinct zones of the SPN specimen, including the compound layer, diffusion zone, and substrate, as listed in Table 2. The nitrogen content reached as high as 4.27 wt% near the outermost surface, decreasing to 3.76 wt% at a depth of 6 μm from the top surface. Within the diffusion zone, the nitrogen content dropped sharply to 0.34 wt% at a depth of 20 μm and further to 0.26 wt% at 50 μm from the outer surface. With increasing depth into the core of the SCM440 steel, the nitrogen concentration decreased to a negligible level. The nitrogen content in the diffusion zone was significantly lower than that in the compound layer, even at locations immediately beneath it. Similar trends have been reported for QPQ-treated 4140 steel [33], where the nitrogen concentration in the compound layer was up to ten times higher than that in the adjacent diffusion zone. This large difference can be attributed to the formation of Fe-nitrides in the compound layer, as shown in Figure 4 and Figure 7. Apart from nitrogen, the concentrations of other alloying elements exhibited minimal variation throughout the analyzed depth. In addition, the reduced carbon content near the top surface may be partly associated with the high nitrogen content in the compound layer.
Figure 7 reveals that the compound layer of the SPN specimen was thicker and finer in structure than that of the NPN specimen, indicating that FPP prior to nitriding promoted the formation of fine Fe-nitrides on the surface of SCM440 steel. Under the same nitriding duration, the SPN specimen exhibited a greater case depth than the NPN specimen (Figure 6a). This suggests that the fine-grained structure induced by FPP facilitated nitrogen diffusion inward, resulting in an increased case depth. Nanohardness measurements of the NPN and SPN specimens revealed that the compound layer of SCM440 steel was slightly harder than the adjacent diffusion zone, consistent with Fe-nitrides being harder than nitrogen-enriched tempered martensite. At a distance of ~6 μm from the interface between the compound layer and diffusion zone, the NPN specimen exhibited slightly lower nanohardness than the SPN specimen. These findings are consistent with the microhardness profiles shown in Figure 6a and further support the conclusion that the refined structure produced by FPP provides additional paths for nitrogen diffusion. Consequently, the lower nitrogen supplement in the NPN specimen limited the formation of a thick compound layer and resulted in a comparatively lower hardness in the diffusion zone compared with the SPN specimen.
The fatigue crack initiation site of a bearing steel is reported to be associated with its inclusion size and depth in the steel [34]. Coarse inclusion and low depth to the surface will shorten the fatigue life of the steel [34]. The crack initiation site of nitro-carburized + post-oxidized 35CrMo steel is found to be on the surface in short fatigue life and from the internal inclusion in very high fatigue life regime [35]. The fatigue crack initiation of the nitrided and nitride + shot-peened 4140 steel is on the surface in low cycle fatigue regime, due to the low ductility of the case layer [20]. Removing the compound layer by shot-peening results in a marked increase in 35% fatigue limit, as compared with the 580 °C/1 h tempered 4140 substrate [20]. Moreover, the fatigue limit of a gas-nitrocarburized steel is known to be less affected by its surface conditions, which include the as-nitrided, oxide-removed and oxide + compound-removed conditions [36].
The fatigue limits of the SB, NPN, and SPN specimens were approximately 750, 1050, and 1400 MPa, respectively (Figure 9). The SPN specimen exhibited the highest fatigue resistance, whereas the SB specimen (SCM440 steel tempered at 475 °C) showed the lowest. The enhanced fatigue strength and life of both the NPN and SPN specimens confirm that nitriding can effectively improve the fatigue performance of SCM440 steel. Furthermore, pre-peening prior to nitriding significantly increased the fatigue limit compared with the non-peened condition. Notably, the SPN specimen exhibited a fatigue limit nearly double that of the tempered substrate. This enhancement is primarily associated with the high CRS field and the refined surface structure, as revealed in Figure 8.
As reported in previous works [19,33], the fatigue limit of SCM440 (AISI 4140) steel increases with decreasing tempering temperature. For the investigated steel with a hardness of HV 500, obtained by solution treatment and tempering at 200 °C, the fatigue limit can exceed 900 MPa [19]. With shot peening, the fatigue limit of the 200 °C tempered 4140 steel increases to ~1150 MPa [19], as shown in Figure 9, indicating that FPP is beneficial for enhancing fatigue resistance. In the case of peening followed by nitriding at 520 °C (SPN520 in Figure 9), the fatigue performance was similar to that of the 200 °C tempered 4140 steel [19]. It is known that shot-peened and nitrided 4140 steel consists of a hardened case and a soft core, which differs from fully hardened steel with HV 500 hardness throughout the cross-section. The fatigue limit of nitrided SCM440 steel also increased as the nitriding temperature decreased (Figure 9). When the nitriding temperature was reduced from 520 to 475 °C, the fatigue limit of shot-peened + nitrided SCM440 steel increased from ~1000 MPa to ~1400 MPa. This remarkable improvement in fatigue resistance is attributed in part to the increased core strength and the presence of a high CRS field.
Fatigue cracks in the SB specimen were found to initiate at the external periphery, with quasi-cleavage in the crack initiation zone accounting for the fatigue failure. The fracture appearance of the SB specimen gradually transitioned from a smooth zone with directional traces to a coarse region interspersed with microcracks. Similarly, all fatigue-fractured NPN specimens exhibited crack initiation at the external surface followed by inward propagation. However, instead of quasi-cleavage at the crack initiation site, the outer surface zone of the NPN specimens displayed a rubbed surface feature coated with a very thin compound layer along the outer profile. No shattering or spalling of the brittle compound layer was observed, which may be attributed to the activity of CRS in the NPN specimens.
In the SPN specimen, fatigue cracks were more likely to initiate at the surface and less likely to originate from subsurface inclusions. The compound layer remained firmly adhered to the surface, and beneath it, lath structures with distinct orientations were observed, corresponding to tempered martensite. These differently oriented lath structures in the fatigue-fractured SPN specimen may have altered the crack growth path within the diffusion zone, thereby contributing to the increased fatigue resistance. Subsurface crack initiation, characterized by a fish-eye zone, was rarely observed in the SPN specimen, and when present, the crack tended to propagate inward more rapidly than outward toward the surface. For both nitrided conditions, fatigue cracks generally initiated at the outermost surface. As shown in Figure 5a, the top surface of the NPN specimen contained microcracks and fine pits, while Figure 5c shows that the SPN specimen exhibited numerous fine dents with minor debris. As illustrated in Figure 11b and Figure 12b,d, such superficial defects on the compound layer surface acted as preferred crack initiation sites in nitrided SCM440 steel. Therefore, surface fatigue crack initiation in nitrided SCM440 steel was associated with those surface defects present in the compound layer. Therefore, polishing after nitriding to reduce surface roughness and remove microcracks was expected to further increase the fatigue resistance and the reliability of nitrided SCM 440 steel. Removing these surface defects by SMAP (Shot Machine A One Polish) treatment [37] may be an effective way to further improve the fatigue properties of FPP + nitrided SCM440 steel.

5. Conclusions

  • Pre-peening before nitriding could impose a refined structure and induce compressive residual stress (CRS) in the near-surface peened zone. The fine-grained structure provided numerous paths to enhance nitrogen diffusion inwards during nitriding. Therefore, the SPN specimen exhibited a higher fraction of iron nitrides in the compound layer, thicker compound layer and deeper case depth than the NPN one.
  • The fatigue limits of the SB, NPN, and SPN specimens were approximately 750, 1050, and 1400 MPa, respectively. Comparing the fatigue limits of the SB and NPN samples, gas-nitriding significantly improved the fatigue properties of SCM 440 steel. Nitriding at 475 °C increased the fatigue limit of the peened + nitrided SCM440 steel to 1400 MPa. This significant improvement in fatigue performance of the peened + nitrided SCM440 steel can be attributed to the combined effects of a high CRS field and a hardened case, both of which retard surface crack initiation and growth.
  • Fatigue cracks were prone to initiate at the external periphery of all tested specimens. In the SB specimen, quasi-cleavage at the crack initiation zone led to fatigue failure. In contrast, the NPN and SPN specimens exhibited a rubbed surface feature at the crack initiation site instead of quasi-cleavage. In both nitrided specimens, the compound layer remained firmly adhered to the surface, with no evidence of shattering or spalling, likely due to the presence of CRS.
  • Surface morphology displayed a relatively flat surface interspersed with microcracks and fine pits in the NPN specimen. By contrast, numerous shallow dents without microcracks were obtained on the top surface of the SPN specimen. Surface microcracks and pits in the NPN specimen could be fatigue initiation sites and harmful to its fatigue resistance. Moreover, these surface dents were considered to be responsible for fatigue crack initiation in the SPN specimens. Therefore, polishing after nitriding to reduce surface roughness and or microcracks was expected to further increase the fatigue resistance and the reliability of nitrided SCM 440 steel.

Author Contributions

Conceptualization, L.-W.T.; Methodology, L.-W.T.; Validation, W.-H.C. and H.-H.H.; Formal analysis, H.C.; Investigation, H.C., T.-C.C., W.-H.C. and H.-H.H.; Resources, T.-C.C. and W.-H.C.; Writing—original draft, L.-W.T.; Writing—review & editing, T.-C.C.; Visualization, T.-C.C.; Supervision, L.-W.T.; Project administration, L.-W.T.; Funding acquisition, L.-W.T. All authors have read and agreed to the published version of the manuscript.

Funding

This research was supported by China Steel Corporation (Contract No. 113I26031).

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

The original contributions presented in this study are included in the article. Further inquiries can be directed to the corresponding author.

Acknowledgments

The authors sincerely thank National Taiwan University for support with EPMA analysis (EPMA000300), National Cheng Kung University for providing access to the optical profiler (OTHER002200), and National Taiwan University of Science and Technology for assistance with nano-hardness measurements (OTHER001900). The fine particle peening was carried out by Vincent Vacuum Tech, and the residual stress evaluations were undertaken by Likuan Technology Corp.

Conflicts of Interest

Authors Wen-Han Chen and Leu-Wen Tsay were employed by the company Vincent Vacuum-Tech Co., Ltd. Author Hsiao-Hung Hsu was employed by the company China Steel Corporation. The remaining authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

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Figure 1. Schematic diagram showing the gas-nitriding process with or without FPP in this work.
Figure 1. Schematic diagram showing the gas-nitriding process with or without FPP in this work.
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Figure 2. Flowchart showing the sample preparation used in this work.
Figure 2. Flowchart showing the sample preparation used in this work.
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Figure 3. Geometry of the rotating bending fatigue test sample used in the current study (unit: mm).
Figure 3. Geometry of the rotating bending fatigue test sample used in the current study (unit: mm).
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Figure 4. XRD spectrum detected from the external surface of the (a) NPN and (b) SPN samples.
Figure 4. XRD spectrum detected from the external surface of the (a) NPN and (b) SPN samples.
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Figure 5. Surface morphologies of (a) NPN and (c) SPN samples observed by SEM, and 3D topography of the (b) NPN and (d) SPN sample obtained using a digital microscope.
Figure 5. Surface morphologies of (a) NPN and (c) SPN samples observed by SEM, and 3D topography of the (b) NPN and (d) SPN sample obtained using a digital microscope.
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Figure 6. (a) The micro-Vickers hardness profile of the NPN and SPN samples from the outmost to the interior; the nano-hardness of the compound layer and diffusion zone of the (b) NPN and (c) SPN samples.
Figure 6. (a) The micro-Vickers hardness profile of the NPN and SPN samples from the outmost to the interior; the nano-hardness of the compound layer and diffusion zone of the (b) NPN and (c) SPN samples.
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Figure 7. EBSD analysis of the investigated samples in cross-sectional view: (a) BC map of the NPN sample; (b) BC map of the SPN sample; (c) IPF of the NPN sample; (d) IPF of the SPN sample; (e) phase map of the NPN sample; (f) phase map of the SPN sample.
Figure 7. EBSD analysis of the investigated samples in cross-sectional view: (a) BC map of the NPN sample; (b) BC map of the SPN sample; (c) IPF of the NPN sample; (d) IPF of the SPN sample; (e) phase map of the NPN sample; (f) phase map of the SPN sample.
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Figure 8. The residual stress distribution of the NPN and the SPN samples in the thickness direction from the surface to the interior.
Figure 8. The residual stress distribution of the NPN and the SPN samples in the thickness direction from the surface to the interior.
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Figure 9. Fatigue stress (S) versus cycle (N) curves of the NPN and SPN samples compared with the SCM 440 steel tempered at 475 °C/1 h (SB). Prepeening + nitriding at 520 °C (SPN 520) and peening the SCM 440 steel tempered at 200 °C/1 h (SP) adapted from Ref. [19] are included.
Figure 9. Fatigue stress (S) versus cycle (N) curves of the NPN and SPN samples compared with the SCM 440 steel tempered at 475 °C/1 h (SB). Prepeening + nitriding at 520 °C (SPN 520) and peening the SCM 440 steel tempered at 200 °C/1 h (SP) adapted from Ref. [19] are included.
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Figure 10. SEM fracture appearance of the SB samples fatigue-fractured at 800 MPa, 4.35 × 105 cycles: (a) the macro-fractured appearance, (b) quasi-cleavage at the crack initiation site, (c) quasi-cleavage fracture inter-dispersed micro-cracks and (d) transition in fracture mode from quasi-cleavage to dimple fracture before final fracture.
Figure 10. SEM fracture appearance of the SB samples fatigue-fractured at 800 MPa, 4.35 × 105 cycles: (a) the macro-fractured appearance, (b) quasi-cleavage at the crack initiation site, (c) quasi-cleavage fracture inter-dispersed micro-cracks and (d) transition in fracture mode from quasi-cleavage to dimple fracture before final fracture.
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Figure 11. SEM photographs showing the fatigue-fractured features of the NPN sample fatigue-fractured at 1050 MPa, 1.08 × 106 cycles: (a) macro-fractured appearance, (b) fracture feature of the surface zone, (c) enlarged view at the site indicated in Figure 10a, showing parallel secondary cracks and (d) enlarged view at the site indicated in Figure 10a, showing mixed mode fracture with secondary cracks along prior austenite grain boundaries.
Figure 11. SEM photographs showing the fatigue-fractured features of the NPN sample fatigue-fractured at 1050 MPa, 1.08 × 106 cycles: (a) macro-fractured appearance, (b) fracture feature of the surface zone, (c) enlarged view at the site indicated in Figure 10a, showing parallel secondary cracks and (d) enlarged view at the site indicated in Figure 10a, showing mixed mode fracture with secondary cracks along prior austenite grain boundaries.
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Figure 12. SEM morphologies of the SPN sample: (a) macro-fractured appearance, (b) fracture feature of the surface zone, (c) macro-view showing subsurface crack initiation, (d) enlarged view at the site indicated in Figure 11c, showing inclusion induced quasi-cleavage fracture, (e) enlarged view at the site indicated in Figure 11c, showing the fracture appearance feature around the surface and (f) enlarged view at the site indicated in Figure 11c, showing the transition in fracture morphology. (a,b) fatigue-fractured at 1500 MPa, 1.88 × 105 cycles; (c,f) fatigue-fractured at 1450 MPa, 8.97 × 105 cycles.
Figure 12. SEM morphologies of the SPN sample: (a) macro-fractured appearance, (b) fracture feature of the surface zone, (c) macro-view showing subsurface crack initiation, (d) enlarged view at the site indicated in Figure 11c, showing inclusion induced quasi-cleavage fracture, (e) enlarged view at the site indicated in Figure 11c, showing the fracture appearance feature around the surface and (f) enlarged view at the site indicated in Figure 11c, showing the transition in fracture morphology. (a,b) fatigue-fractured at 1500 MPa, 1.88 × 105 cycles; (c,f) fatigue-fractured at 1450 MPa, 8.97 × 105 cycles.
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Table 1. Chemical compositions of the SCM440 steel bar.
Table 1. Chemical compositions of the SCM440 steel bar.
CMnPSCrMoSiNiFe
0.420.810.0110.0061.010.170.230.01Bal.
Table 2. The chemical compositions in distinct zone of the SPN sample.
Table 2. The chemical compositions in distinct zone of the SPN sample.
LocationDistance from Surface (μm)Chemical Composition in wt%
CNOMnCrMoSiFe
Compound layer20.364.270.150.500.770.120.23Bal.
60.403.760.090.480.750.130.23Bal.
Diffusion zone200.400.340.040.821.010.120.23Bal.
500.420.260.030.910.850.080.19Bal.
1500.440.110.030.760.930.120.20Bal.
Base2500.420.000.000.721.000.100.21Bal.
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Chen, H.; Chen, T.-C.; Chen, W.-H.; Hsu, H.-H.; Tsay, L.-W. Effects of Pre-Peening on Fatigue Performance of Gas-Nitrided SCM 440 Steel. Metals 2025, 15, 1118. https://doi.org/10.3390/met15101118

AMA Style

Chen H, Chen T-C, Chen W-H, Hsu H-H, Tsay L-W. Effects of Pre-Peening on Fatigue Performance of Gas-Nitrided SCM 440 Steel. Metals. 2025; 15(10):1118. https://doi.org/10.3390/met15101118

Chicago/Turabian Style

Chen, Hao, Tai-Cheng Chen, Wen-Han Chen, Hsiao-Hung Hsu, and Leu-Wen Tsay. 2025. "Effects of Pre-Peening on Fatigue Performance of Gas-Nitrided SCM 440 Steel" Metals 15, no. 10: 1118. https://doi.org/10.3390/met15101118

APA Style

Chen, H., Chen, T.-C., Chen, W.-H., Hsu, H.-H., & Tsay, L.-W. (2025). Effects of Pre-Peening on Fatigue Performance of Gas-Nitrided SCM 440 Steel. Metals, 15(10), 1118. https://doi.org/10.3390/met15101118

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