Effect of Pre-Modification by Water Jet Blasting Prior to Nitriding Vein-Like Precipitates in the Composite Diffusion Layer of M50NiL Steel

Abstract

In this study, M50NiL steel was carburized (C), nitrided (N), and compound-carburized then nitrided (C + N). Vein-like grain boundaries (VLGBs) were observed in the diffusion layers of both the N and C + N states due to the limited opportunity for diffusion. Transmission electron microscopy (TEM) observation revealed that the VLGB organization differed in the N and C + N states. The VLGB organization consisted mainly of Fe₄N in the N state and Fe₃C and Fe₄N in the C + N state. When the C state was pre-modified by a 200 MPa water jet and then nitrided (C + 200P + N), the increase in dislocation density resulted in a dislocation entanglement phenomenon that split the grains to form subcrystals. The increases in grain boundaries and dislocation density promoted the diffusion of atoms, and thus the VLGB structure was not observed in the diffusion layer of the C + 200P + N state.

1. Introduction

M50NiL martensitic steel is a low carbon secondary hardening steel with high contents of Cr, Mo, Ni and V alloying elements [1]. Its fracture toughness is more than twice that of M50 steel [2]. A new generation of high-strength aero-engine gears requires bearing steel with excellent surface properties in order to operate under heavy loads at high speeds, high temperatures and other harsh service conditions. Carburizing + plasma nitriding (compound carburizing) is a thermochemical surface treatment process [3,4,5]. Compared to other surface thermochemical treatments, such as carburizing [6,7], nitriding [8,9,10], and carbonitriding [11,12], compound-carburizing treatments produce a large hardness gradient to maintain operation in service and prevent the reduction of the residual compressive stress gradient field [3]. Compared with carburized specimens, the microhardness and wear resistance of compound-carburized surfaces are better [5]. In plasma nitriding, carbonitriding and compound-carburizing processes, a diffusion layer usually forms adjacent to the substrate in addition to a compound layer. The compound layer plays an important role in improving wear resistance and corrosion resistance [13,14,15,16]. However, the fatigue properties are closely related to the diffusion layer [5]. The plasma-nitriding process results in the formation of vein-like precipitates along the grain boundaries in the diffusion layer, which are detrimental to the fatigue performance and can become weak spots for cracking during fatigue loading [17].

For the formation of VLGB tissue, the initial carbide-to-nitride transition occurs during the subsequent plasma-nitriding process in the compound infiltration process. This produces many reactive C atoms, which mainly diffuse towards the matrix region using the grain boundaries as channels [18]. Chen et al. [19,20] studied the structure of vein-like grain boundaries in the diffusion layer during the carbon-nitriding process and found that during the plasma carbon-nitriding process, the initial transformation of carbide to nitride released many carbon atoms that diffused along the grain boundaries into the matrix. However, the number of grain boundaries is limited, and carbide or nitride is formed at the grain boundaries. Since grain boundaries are the main diffusion channels for atoms, nitriding pre-treatment can be considered before nitriding, and the increase in the number of grain boundaries will promote the diffusion of C and N atoms into the matrix region [21]. Therefore, surface deformation strengthening of the metal surface before nitriding, ultrasonic surface rolling process (USRP) [22,23,24], surface mechanical abrasion treatment (SMAT) [25] and ultrasonic nanocrystal surface modification (UNSM) [26,27], are imperative. However, due to the complex shape of aero-engine parts, the use of the above deformation strengthening methods has limitations. Water jet blasting (SP) is a traditional surface-modification method, and specimen SP treatment can increase the number of grain boundaries by changing the microstructure of the deformation layer and forming subcrystals through an increase in dislocation density to achieve grain refinement [28,29,30].

A detailed study on the early stage of composite carburizing of M50NiL steel was conducted. Vein-like precipitates were observed in the diffusion layer, but the mechanism of their formation was not discussed in detail [2,3,4,5]. In this study, carburized M50NiL steel was subjected to plasma-nitriding treatment and water jet shot pre-treatment followed by nitriding treatment. Since C + N has a better modification effect than a single surface-modification treatment, the results of this study are expected to provide an experimental and theoretical basis for improving the heat-treatment process of M50NiL steel. Therefore, the following questions must be answered in this paper.

(1)

How does the VLGB organization form in the diffusion layer after nitriding treatment and compound-carburizing treatment?

(2)

How does the water jet shot peening pre-treatment of carburized M50NiL steel before plasma nitriding inhibit the formation of VLGB tissue?

(3)

What is the difference between the VLGB tissue in the N state and the VLGB tissue in the C + N state?

2. Experimental Materials and Procedures

2.1. Materials

The material used in this paper was M50NiL steel, which was prepared by the vacuum induction melting (VIM) and vacuum arc remelting (VAR) processes. The microstructure of the raw material was characterized as martensitic by electron back scatter diffraction (EBSD), as shown in Figure 1a,b. The chemical composition of M50NiL steel is shown in

Table 1. Chemical composition of M50NiL steel (wt%).

Element	C	Cr	Ni	Mo	V	Mn	Si	P	S	Fe
Content	0.13	4.14	3.37	4.07	1.24	0.29	0.19	0.005	0.002	balance

.

2.2. Experimental Methods

Before carburizing, the rod specimens were ground and polished, followed by ultrasonic cleaning in ethanol for 30 min. A low-pressure vacuum carburizing treatment was performed on the rod specimens in a double-chamber vacuum carburizing furnace (WZST-20 type) produced by the Beijing Institute of Mechanical and Electrical Research. Vacuum carburizing was divided into two stages: (1) pulsed strong carburizing stage at 950 °C for 1 h in 50% acetylene and 50% high purity nitrogen at flow rates of 4 L/min. The carburizing treatment is to maintain pressure at 1500 Pa for 20 s after injecting C₂H₂ and N₂ gas mixture, and then vacuum in the hot chamber, this process is repeated before the diffusion stage; (2) diffusion stage at 950 °C for 2 h. The carburized specimens were air-cooled to room temperature in a cold chamber filled with N₂, followed by a series of heat treatments as follows: (1) vacuum annealing of all specimens at 750 °C for 3 h; (2) solution heat treatment of all specimens at 1100 °C for 30 min, followed by gas quenching by passing N₂ to cool the specimens to room temperature; (3) deep cooling treatment at −73 °C for 2 h; and (4) tempering at 540 °C. The specimens in this state were called C, as shown in

Table 2. Heat treatments of C specimens.

Parameter	Vacuum Annealing	Gas Quenching (N2)	Deep Cooling Treatment	Tempering
Temperature/°C	750	1100	−73	540
Time/h	3	0.5 (0.4 MPa)	2	2 (3 times)

. The C and the original specimens were nitrided in a plasma-nitriding furnace (LDMC-30AZ, Ande Heat Treatment Equipment Limited Company, Wuhan, China) at a nitriding temperature of 520 °C and air pressure of 290 Pa, as shown in

Table 3. Nitrided treatments of the C and the original specimens.

Sample	Temperature/°C	Time/h	Preesure/Pa
N	520	7	290
C + N

. The specimens were cooled to room temperature in the furnace after holding at 520 °C for 7 h. The specimens in this state were called C + N and N, respectively. The last three groups of samples were treated with water jet shot peening at different pressures on C, denoted as C + 100P, C + 200P, and C + 300P respectively. Then, nitriding was treated at 520 °C, denoted as C + 100P + N, C + 200P + N, and C + 300P + N respectively. Specific parameters are shown in

Table 4. Water jet parameters.

Sample	Pressure (MPa)	Shaft Speed (r/min)	Target Distance (mm)	Feed Rate (mm/min)
C-100P	100	30	10	8
C-200P	200
C-300P	300

. Abrasive water jet blasting was performed using SQ1313 CNC (Jiangsu Xusheng Water Jet Technology Company, Yixing, China) water jet equipment, which consisted of CNC machine tools, abrasive conveying systems, water supply systems, and so forth. The highest working water pressure reached 400 MPa. The basic principle relies on the self-weight of the abrasive and the negative pressure generated by the high-speed water flow. Abrasive particles are “pumped” into the mixing chamber and mixed with the high-pressure water carrying a very great amount of energy. The mixture is then sprayed onto the metal surface at high speed to produce plastic deformation and achieve the modification effect, as shown in Figure 2.

2.3. Characterization

The cross sections of the samples were etched with 4% nitric acid alcohol solution, and the microstructure of the specimens was observed using field emission scanning electron microscopy (SEM, SUPRA40-41-90, SEM, SUPRA40-41-90, Zeiss, Jena, Germany), followed by EBSD (SUPRA40-41-90, Zeiss, Jena, Germany) to identify the tissue structure of the raw material using a step size of 0.4 μm. The hardness distribution was determined by a microhardness tester (HV1000, HV1000, Lunjie Motor Instrument Company, Shanghai, China) by applying a load of 0.98 N and a dwell time of 20 s. The phase structure of the surface-treated layers was determined using X-ray diffraction (XRD, XRD, Rigaku Ultima IV, Tokyo, Japan) and Cu Kα radiation. The specimens were mechanically ground to less than 60 μm and then ion-milled to approximately 200 nm for transmission electron microscopy (TEM) and high-resolution transmission electron microscopy (HRTEM) (Tecnai G2F20STWIN, Tecnai G2F20STWIN, FEI, Hillsboro, OR USA, 200 kV) to observe the vein-like precipitates in the diffusion layer.

3. Results and Discussion

3.1. Microstructure after C, N and C + N Treatments

The solid-state diffusion process is often accompanied by the occurrence of phase changes and formation of some new phases. After C, N and C + N treatments, precipitates rich in alloying elements formed in the modified layers because M50NiL steel has many strong carbon- and nitride-forming elements (Cr, Mo, V). For the C state, reaction diffusion and downhill diffusion occur at 950 °C. This generates carbides rich in alloying elements and carbon atoms in the octahedral interstices of α-Fe along the gradient direction to form interstitial solid solutions. As shown in Figure 3a for the distribution of carbides after C treatment, carbides are uniformly distributed in the carburized layer in a granular form. The carbide of the carburized layer was observed by TEM, as shown in Figure 4, and the alloying elements combined with C atoms to generate carbide. According to the results of EDS, the enrichment phenomenon of V is more obvious, because the binding ability of V to C is greater than that of Mo and Cr. For the N state, reaction diffusion occurred first in the surface layer and formed a compound layer of approximately 6 μm, as shown in Figure 3b. Then, in the diffusion layer, N atoms combined with Cr, Mo and V to form nitrides rich in alloying elements or solid solutions in the octahedral interstices of α-Fe, as shown in Figure 5. In addition to the formation of compound layers and precipitation of nitrides rich in alloying elements, coarse vein-like precipitates formed in the diffusion layer, which is shown in Figure 3c. For specimens in the C + N state, the formation of the VLGB organization was complicated by the precipitation of carbides after C treatment. The carbides in the carburized layer and the C atoms in the α-Fe octahedral gap became unstable due to the introduction of N atoms, and the decomposition of alloy carbides and the devolatilization of C atoms occurred [2]. As shown in Figure 6, the grainy precipitates are carbon-nitrides rich in alloying elements formed by the replacement of C atoms by N atoms. Atoms were desolvated to form an N-containing α-Fe interstitial solid solution, which was subsequently verified by XRD.

3.2. Formation Mechanism of the VLGB Structures

The decomposed alloy carbide and the desolvation phenomenon release many free C atoms, and the grain boundaries are the main channels for atomic diffusion. These C atoms and some N atoms in the C + N process diffuse along the grain boundaries towards the matrix region [18]. However, the number of grain boundaries is also limited, and not all C and N atoms can diffuse smoothly through the grain boundaries towards the matrix region. They are enriched at the grain boundaries to form VLGBs, as shown in Figure 7a,b, which shows the VLGB organization. During the C + N treatment, the carburizing process precipitated many uniformly distributed carbides of alloying elements in the carburizing layer. The alloying elements gathered at the granular carbides, and the carbon-nitrides of alloying elements formed in the following nitriding process [31], as shown in Figure 6 and Figure 7c,d. These granular precipitates were enriched with alloying elements. The surrounding VLGB tissue formed mainly when C and N atoms diffused along the grain boundaries, the matrix contained a large amount of Fe, the alloying elements were already enriched at the grain precipitates, and the C and N atoms combined with the Fe atoms to form Fe-rich compounds [32]. EDS analysis of the VLGB tissue is shown in Figure 7c. It was initially determined to be an Fe-rich compound; the elemental composition of the grain precipitates around the VLGB tissue differed from that of the VLGB tissue. The grain precipitates were carbon-nitrides enriched with alloying elements. Specific VLGB tissue was further analyzed by TEM.

3.3. XRD Analysis

As shown in Figure 8, the X-ray diffraction diagrams of the three states of C, N and C + N showed that the surface layer of the N specimen was a compound layer consisting of γ′ (Fe₄N) and ε(Fe_2–3N) [2]. The compound layer was observed in Figure 3b, where it hindered the diffusion of N atoms to the surface [18]. Without the presence of the compound layer, nitrides rich in alloying elements formed at the compound layer as well as more diffusion channels. This was because the original location with grain boundaries had fewer diffusion channels due to the formation of the compound, and the N atoms mainly diffused inwards through the grain boundaries. Thus, the N atoms were more likely to be enriched at the grain boundaries to precipitate VLGB tissue. Comparing Figure 3b and Figure 7b shows the N state. The surface of the C specimen was mainly composed of α-Fe, and after the C + N treatment, the N atoms had a smaller atomic radius than the C atoms and had a higher maximum solid solution in α-Fe. Thus, the diffraction peaks of α-Fe became wider, which proved that many N atoms were in solid solution in the octahedral interstices of α-Fe where they replaced the N atoms that were originally in solid solution in the octahedral interstices of α-Fe [19]. This indirectly proved that the precipitation of the VLGB structure was closely related to the diffusion of C atoms.

3.4. TEM Observations of VLGB Tissue

To distinguish between the VLGB tissue in the C + N and N states and to verify that the C and N atoms mainly diffused along the grain boundaries and were enriched at the grain boundaries, TEM observations were performed for the VLGB tissue in the N and C + N states. As shown in Figure 9a, the VLGB tissue formed by the plugging of N atoms at the grain boundaries was observed by TEM, and the diffraction spots shown in Figure 9b were obtained by electron diffraction of selected areas at the VLGB tissue marks in Figure 9a, which were mainly composed of Fe₄N. For the C + N state, since the enrichment of C atoms was accompanied by the enrichment of N atoms at the grain boundaries, the TEM images of the VLGB tissue in the diffusion layer in the C + N state is shown in Figure 10a at the red dashed marker. Additionally, a high-resolution image of the yellow box is shown in Figure 10b where there is a clear demarcation line between the two phases shown by the red dashed line. Selected area electron diffraction was performed on the c and d regions of Figure 10b to obtain Figure 10c,d. The VLGB tissue in the C + N state was different from that in the N state. The VLGB tissue in the C + N state was composed of Fe₃C and Fe₄N, and other scholars found that this VLGB tissue was composed of different physical phases [20].

3.5. Mechanism of Inhibiting VLGB Tissue Precipitation

Considering that grain boundaries are the diffusion channels for C and N atoms, this experiment considered the use of water jet blasting to pre-modify the C-treated specimens prior to nitriding so that the surface underwent plastic deformation to produce grain refinement [30], the concentration of C and N atoms per unit area is reduced by increasing the number of grain boundaries after refining, and it is not easy to enrich and form VLGB structures at grain boundaries.

3.5.1. Microstructure of Water Jet Blasting and Water Jet Blasting + Nitriding on C

In this study, 100 MPa, 200 MPa, and 300 MPa were selected as the applied pressures, and the carburized specimens were subjected to different pressures of water jet shot peening followed by plasma nitriding. The VLGB organization precipitated along the grain boundary and appeared in the diffusion layer of the 100 MPa treated specimens, which was 225 μm from the surface, as shown in Figure 11a. When the pressure changed to 200 MPa and 300 MPa, the VLGB tissue disappeared at the same location in the diffusion layer, and only uniformly distributed granular carbonitrides were present, as shown in Figure 11b,c. Figure 12a–e shows the microstructure of the surface layer to 50, 100, 150, and 200 μm after C + 200P treatment, and Figure 12f–j shows the microstructure of the surface layer to 50, 100, 150, and 200 μm after C + 200P + N treatment. Comparing the two states shows that the modified layer was composed of uniformly distributed granular precipitates, and there was no obvious difference in microstructure. The microstructure of the modified grains was not significantly different from that of the original austenite grains, and these carbides were very stable when heated and not easily dissolved into the original austenite grains. Thus, this prevented the regrowth of the modified grains during the nitriding process and reduced the channels for the diffusion of C and N atoms, which led to their enrichment at the grain boundaries. The cause of the disappearance of VLGB organization is analyzed in detail below.

3.5.2. Grain Refinement Promotes the Diffusion of C and N Atoms

After surface water jet blasting to pre-modify the C state, a certain depth of the modified layer formed. In the modified layer, grain refinement occurred, and the degree of grain refinement differed for different degrees of plastic deformation under different pressures. Grain refinement formed high-speed diffusion channels (grain boundaries) in the material, and the presence of more grain boundaries promoted the diffusion of C and N atoms in the crystal [33]. The increase in the number of grain boundaries after surface modification put the material in a high-energy, sub-stable state [34], and the material tended to change to a thermodynamically stable state during the subsequent nitriding process. The higher energy after surface modification acted as activation energy for diffusion and thus reduced the energy required for diffusion. Thus, the activation energy for diffusion along the grain boundaries was much smaller than in the lattice, which enhanced the diffusion rates of C and N atoms.

Figure 13a shows the microstructure of the C + 100P specimen; the depth of the modified layer was only approximately 100 μm at a pressure of 100 MPa. When the pressure was increased to 200 and 300 MPa, the depth of the modified layer increased to approximately 300 μm, and the depths of the modified layers of C + 200P and C + 300P were approximately the same, as shown in Figure 13b,c. The above description proved that the degree of plastic deformation under a pressure of 100 MPa was less than in the two states of C + 200P and C + 300P, which led to insufficient grain refinement to provide sufficient channels for the diffusion of C and N atoms. Thus, there was still the formation of VLGB tissue in the diffusion layer of C + 100P + N, as shown in Figure 11a, and it was still possible to observe C and N atoms in the enrichment phenomenon at the grain boundaries. The next TEM analysis of the three states of C, C + 100P and C + 200P revealed the superficial region of the carburized layer in the C state, as shown in Figure 14a. This showed a set of single-crystal diffraction spots without grain refinement after the selected area electron diffraction of its yellow marked position. Since the modification effects of C + 200P and C + 300P were approximately the same, the selected C+ 100P and C + 200P were selected for TEM characterization to further analyze the different degrees of plastic deformation caused by different pressures and thus the different degrees of grain refinement, as shown in Figure 14b,c for the TEM of the modified layer local area of C + 100P and C + 200P, respectively. Then, the diffraction rings appeared after the selected electron diffraction of the yellow marked positions, and the appearance of diffraction rings represented the appearance of diffraction rings at C + 100P and C + 200P specimens. They showed grain refinement at 100 MPa and 200 MPa, but the diffraction rings of C + 100P were not continuous compared to that at C + 200P. This proved that the degree of plastic deformation was not sufficient, and thus the degree of grain refinement was not as good as that of C + 200P. The degree of deformation produced by C + 100P was less than that of C + 200P [35], and the pre-modification before nitriding at a pressure of 200 MPa inhibited the formation of VLGB tissue. As shown in Figure 15a, which is a TEM image of the carburized layer in the C state, the boundary of the original slate-martensite within the carburized layer could be clearly observed. After the pre-modification of C by 200 MPa water jet blasting, the shape of the original slate-martensite gradually disappeared due to the significant plastic deformation. Additionally, the martensitic lath martensite were divided into small subcrystals by the surrounding dislocation entanglement, which was also found by Hosseini [36]. The formation of subcrystals led to an increase in the number of grain boundaries and thus provided more channels for the diffusion of C and N atoms. V. Lacaille et al. [37] enhanced the nitriding efficiency of pure iron by pre-modifying it with shot peening. The nanocrystals formed by this surface deformation strengthening technique provided a very significant advantage by increasing the number of grain boundaries and enhancing the diffusion phenomenon.

According to the Arrhenius formula [38]:

$$D=D_0{\mathrm{e}}^{\frac{RT}{Q}}$$

(1)

where D₀ denotes the diffusion constant, Q is the diffusion activation energy, R is the gas constant, T is the thermodynamic temperature, and D is the diffusion coefficient.

According to Equation (1), for C + N, C + 100P + N and C + 200P + N, when no pre-modification by water jet blasting was performed, the diffusion of C and N atoms in the C + N state required the largest diffusion activation energy. At the same thermodynamic temperature T, it had the smallest diffusion coefficient, and thus C and N atoms were easily plugged at the grain boundaries, which led to the formation of VLGB tissue. For the C + 100P and C + 200P states, since their carburizing and nitriding temperatures were the same, the difference was the pressure parameter of water jet blasting. When the pressure was 100 MPa, the modified layer was shallow, and the degree of plastic deformation was not as good as that at 200 MPa. After pre-modification at 200 MPa, the number of grain boundaries increased due to the formation of subcrystals, as shown in Figure 15b, and the grain boundaries. The increase in the number of grain boundaries and the increase in the degree of plastic deformation put it in a high-energy, sub-stable state, and this higher energy acted as the energy required for diffusion. Thus, it was easier for C + 200P to overcome the energy barrier required for diffusion, and the C + 200P state with a large degree of deformation required less diffusion activation energy than the C + 100P state. As a result, the diffusion coefficient of the C + 200P state was large, and the diffusion rates of C and N atoms were larger than those of the C + 100P and C + N states. The atoms diffused mainly along the grain boundaries towards the core, so no VLGB organization formed.

3.5.3. Variation in the Hardness Gradient of the Modified Layer

The movement of dislocations was closely related to the occurrence of plastic deformation, and the increase in the degree of plastic deformation was accompanied by an increase in hardness. Then, combined with Equation (2) [39], the results showed that there was a positive correlation between the dislocation density and the increase in hardness. The surface water jet shot peening of carburized M50NiL steel occurred after the surface-hardening phenomenon, and thus the surface microhardness increased. As shown in Figure 16a, the surface microhardness increased to 780 HV0.1 after the C + 100P treatment, which was an approximately 4% increase, and to 810 HV0.1 after the C + 200P treatment, which was an approximately 8% increase. It is proved that after C was pre-modified by water jet blasting at 200 MPa, Figure 14c showed that the surface layer underwent severe plastic deformation that led to grain refinement (work hardening). Thus, the surface microhardness of C + 200P was the largest.

$$H=3\sqrt{3}\alpha Gb\sqrt{\rho }$$

(2)

where H is the Knoop hardness (GPa), α is a constant with a value of 0.3, G is the shear modulus (80 GPa for steel), b is the Burgers vector (0.2476 nm for steel), and ρ is the dislocation density (m⁻²).

After pre-modification of C by water jet blasting, in addition to grain refinement, the dislocation density within the modified layer changed. From Equation (2), the dislocation density of the surface layer of C was calculated as 6.1 × 10²⁴/m². After C + 100P treatment, the dislocation density of the surface layer increased to 6.6 × 10²⁴/m², and after C + 200P treatment, the dislocation density of the surface layer increased to 7.1 × 10²⁴/m². The 200 MPa water jet modification increased the dislocation density relative to C by 16.4%. The dislocation density of C + 100P increased by 8.2% relative to that of C. According to Figure 16(a)(c), the surface hardness and dislocation density increment of C + 200P were approximately twice those of C + 100P. As shown in Figure 15b, the red marks are subcrystals formed by the surrounding dislocations. Thus, the dislocation density near the subcrystal grain boundaries was greater, and the increase in dislocation density proved that the material underwent severe plastic deformation under a pressure of 200 MPa. The results are shown in Figure 16b. After C + N treatment, the microhardness at 225 μm from the surface was 750 HV0.1, which was approximately the same as the surface microhardness of C. After C + 100P + N treatment, the microhardness at 450 μm from the surface was 775 HV0.1, which was approximately the same as the surface hardness of C + 100P. After C + 200P + N treatment, the microhardness at a distance of 550 μm from the surface was 812 HV0.1, which was approximately the same as that of C + 200P. Combined with the microhardness from Figure 16d, the thickness of the nitride layer of C + N was 225 μm, and the thickness of the nitride layer of C + 100P + N was 450 μm, which was 225 μm more than that of C + N. Additionally, the thickness of the nitride layer of C + 200P + N was 550 μm, which was 325 μm more than that of C + N. After pre-modification by water jet blasting, the dislocation density and the increase in grain boundaries led to an increase in the number of N atoms. The increases in the numbers of grain boundaries and dislocations both reduced the diffusion activation energy required for the diffusion of C and N atoms and provided channels for the diffusion of C and N atoms. This significantly promoted the diffusion of C and N atoms and thus increased the depth of the nitride layer after pre-modification of C and then nitriding [40].

4. Conclusions

In this study, M50NiL steel was treated with N, C, C + 100P, C + 200P, C + N, C + 100P + N, and C + 200P + N. The main findings are as follows:

(1)

Due to the limited number of diffusion channels, VLGB tissue precipitated along the grain boundaries in the diffusion layers of N, C + N, and C + 100P + N, and no VLGB tissue precipitated in the diffusion layers of C + 200P + N and C + 300P + N.

(2)

The plastic deformation of C + 200P was greater than that of C + 100P, which led to an increase in the surface microhardness and dislocation density of C + 200P that was approximately twice as much as that in the case of C + 100P. The increase in dislocation density formed dislocation entanglements and thus refined the original coarse grains to produce subcrystals, which provided more diffusion channels and thus inhibited the formation of VLGB tissue.

(3)

The VLGB organization differed in the N and the C + N states. The VLGN organization was mainly composed of Fe₃C in the N state and Fe₃C and Fe₄N in the C + N state.