3.1. Microstructures of Nitride Layers
The cross-sectional microstructure of the 1Cr17Ni2 steel sample after nitriding is shown in
Figure 1.
Figure 2 shows the distribution of N element in the cross-section of the sample under different nitriding processes using EDS. As depicted in the figure, it is clear that with the increase in nitriding time, there is a significant increase in the thickness of the nitride layer. The steel surface forms a thin and uniform compound layer [
1]. It can be seen from the figure that the nitride layer surface is composed of three parts, namely the diffusion zone, the heat-affected zone, and the substrate. The interface between the diffusion zone and the heat-affected zone is smooth, while the interface between the heat-affected zone and substrate is blurred. The depth of the nitrided layer for different nitriding processes is shown in
Table 2. It also confirmed that as the temperature rises, the depth of the nitrided layer in the sample nitrided for 8 h increases significantly with the increase in nitriding temperature, while the sample nitrided for 16 h shows an increasing trend, although this is not very obvious, with the increase in nitriding temperature, and the same situation also occurs in Landgraf’s study [
20]. It can also be observed from the figure that as the temperature increases, the particles in the matrix gradually grow, which is due to the increase in the driving force for grain growth caused by the temperature increase. At the same time, with the extension in nitriding time, the modification depth also continues to deepen due to the introduction of more nitrogen into the material. Due to the influence of diffusion kinetics, nitrogen enters from the high-concentration part of the material into the lower-concentration part, and the higher temperature provides energy for nitrogen to diffuse inward, resulting in the further inward diffusion of nitrogen, which explains why the modification depth continuously increases with the extension in nitriding time.
As shown in
Figure 3, the XRD spectra of samples with different nitriding processes exhibit significant differences in peak intensity. During the nitriding process, the diffusion and chemical reactions of Fe, Cr, Ni, and N elements lead to changes in the surface phase composition of the samples. The nitride layer mainly consists of a compound layer and a diffusion zone. The compound layer is primarily composed of γ′-Fe
4N, a small amount of ε-Fe
2-3N, and CrN, while the diffusion zone is primarily composed of α′
N-Fe [
21,
22], where N and C atoms dissolve interstitially in the α-Fe lattice. The composition of the nitride layer is the same as the phase composition obtained by Li junyang et al. for 2Cr13 martensitic stainless steel after ion nitriding [
23]. γ′-Fe
4N is formed due to the diffusion of N atoms into Fe during the nitriding process. It is worth noting that the diffraction peaks of the γ′-Fe
4N phase show a slight shift towards lower angles, mainly due to chemical reactions and element diffusion during nitriding. On one hand, the formation of nitrides and metal compounds reduces the content of pure metal phases. On the other hand, N atoms as external interstitial atoms enter the metal, while alloying elements usually act as substituting atoms, resulting in the expansion of lattice parameters. This is consistent with the increase in equilibrium lattice constant observed in the experimental calculations in
Section 3.3 of this paper. Alloying elements diffuse and interact atomically during the nitriding process, leading to the formation of metal compounds, and they play a special role in phase transformations. It is worth noting that at lower nitriding temperatures and shorter nitriding times, the peak intensity of γ′-Fe
4N is higher. With an increase in nitriding temperature and extension of nitriding time, the peak intensity of γ′-Fe
4N gradually weakens while the peak intensity of ε-Fe
2-3N increases. This is because an increase in nitriding temperature and time leads to an increase in the concentration of N solid solution into the matrix. At the same time, the higher nitriding temperature makes the nitriding kinetics more active, which allows for a sufficient amount of N to react and form nitrides. Also, with an increase in N concentration, the reaction between γ′-Fe
4N and N favors the formation of ε-Fe
2-3N. Furthermore, a narrowing of the peaks is observed, indicating that the grain has grown to a certain extent, which is consistent with the observations in
Figure 1. Additionally, due to the high Cr content of the 1Cr17Ni2 stainless steel, the surface phase of the sample contains CrN. The theoretical reaction products of Cr and N include CrN and Cr
2N. However, only CrN is observed in the XRD pattern, which agrees with the experimental and theoretical results. According to the first-principle calculations, the energy change for the precipitation of CrN is −0.19 eV, while that of Cr
2N is 9.86 eV [
24]. This suggests that the precipitation of CrN can occur spontaneously, while the precipitation of Cr
2N is more difficult. This also explains why CrN is preferentially precipitated in the nitriding layer on the surface of 1Cr17Ni2 stainless steel.
3.2. Microhardness after Nitriding
As shown in
Figure 4, measuring the hardness along the surface towards the center of the sample and taking the average of the hardness of three test points at the same depth every 10 μm as the hardness at that depth, the microhardness profiles of the nitride layer profile after nitriding treatment at different temperatures for 1Cr17Ni2 steel show a parabolic distribution overall, reaching a peak hardness of approximately 1000 Hv
0.1. The surface hardness of the nitride layer was greatly improved after nitriding treatment. It can be seen from the figure that as the depth of the nitride layer increases, the hardness first increases, reaches a peak, and then slowly decreases before suddenly dropping. A hardness gradient distribution appears at the interface between the nitride layer and the substrate, and the gradient becomes more gradual as the temperature increases. Upon the careful observation of
Figure 4a,b, it can be seen that a peak in microhardness occurs within a depth range of 30 μm to 90 μm, which coincides with the depth range of the white bright layer when compared to
Figure 1. The sample nitride at 460 °C for 16 h exhibits the highest surface hardness, reaching 1338 Hv
0.1. Comparing the samples that were nitrided for 8 h and 16 h, it can be seen that samples nitride at 460 °C generally have the highest surface hardness. According to
Figure 3, it can be seen that there are a large amount of high-hardness α′
N, ε-Fe
2-3N, γ′-Fe
4N, and a small amount of CrN phases generated in the surface phase at 460 °C, while the CrN content is higher in the surface phase at 500 °C and 550 °C, reducing the surface hardness at these two temperatures [
25]. Meanwhile 460 °C is more suitable for nitrogen solid solution, while at 500 and 550 °C, a small amount of compound precipitation occurs, resulting in a slight decrease in hardness [
26].
3.3. Friction and Wear Mechanism of 1Cr17Ni2 Steel after Plasma Nitriding
This study used a friction tester to determine the friction coefficients of specimen surfaces after different processing techniques. As shown in
Figure 5, the friction coefficients of different samples generally increased over time before stabilizing. The significant changes in friction coefficients among different techniques mainly occurred within the first 800 s of wear, after which they gradually stabilized. The original specimen exhibited the lowest friction coefficient, with an overall trend of initially increasing and then decreasing.
The variation in friction coefficient during wear depends on several factors, including the phase composition, the hardness of the worn surface, and the wear mechanisms. The local fluctuations in friction coefficient during wear can be explained by the mechanism of “oxidation–separation–reoxidation” in oxidative wear [
27]. Heat generated by friction during wear causes debris to oxidize, forming a dense oxide film that effectively prevents direct contact between the mating surfaces, thus lubricating and reducing the friction coefficient.
The wear rate of the nitrided surface layer initially increases and then decreases as the nitriding temperature rises, and this phenomenon holds true when nitriding for 16 h as well. As evident from the graph, the wear rate decreases with the increase in nitriding time. Observing the changes in the stable friction coefficient, it is evident that its variation is inversely proportional to the surface wear rate: as time progresses, the friction coefficient of the 16 h nitrided surface is consistently higher than that of the 8 h nitrided samples. The wear resistance after nitriding experiences a significant improvement compared to unnitrided samples. Specifically, the wear rate of the original sample is 2.31 × 10−5 mm3/(N·m), whereas the wear rates of the nitrided samples decrease by two orders of magnitude compared to the original sample.
Moreover, as depicted in
Figure 6, under the same nitriding temperature, the wear rate of samples nitrided for 16 h is generally lower than that of samples nitrided for 8 h. Concurrently, the friction coefficient of the 16 h nitrided samples is generally higher than that of the 8 h nitrided samples. This indicates that the friction performance of the samples improves significantly with the increase in nitriding time.
The three-dimensional morphology of wear after 2400 s at 300 rpm for different nitriding processes is depicted in
Figure 7, showing similar wear profiles for all specimens without significant macroscopic plow grooves.
Table 3 provides the depths of wear marks under different processes, indicating a significant reduction in wear mark depth after nitriding compared to the original specimen. At the same sliding speed, the depth of wear marks increased with time and temperature. Among them, nitride at 460 °C exhibited the lowest wear mark depth, which increased with nitriding temperature, leading to an increase in both wear rate and wear mark depth.
A comparison between the wear rate and hardness revealed an inverse relationship, indicating that increasing hardness can significantly reduce wear rate and improve wear resistance [
28,
29]. However, hardness is not the sole determinant of wear resistance, as co-nitrided layers with optimal toughness exhibit the best wear resistance.
As shown in
Table 3, the wear scar of the nitrided sample is significantly reduced compared to the original sample. At the same sliding speed, the depth of the wear scar increases with the increase in time and temperature. Among them, the 460 °C has the shallowest wear scar, and the wear scar depth increases with the increase in nitriding temperature, resulting in an increase in both wear rate and wear scar depth. At the same time, through the comparison of wear rate and hardness, it is found that the wear rate is roughly inversely proportional to the hardness, and increasing the hardness can significantly reduce the wear rate and improve the wear resistance.
In order to observe the micro-morphology of the wear scar in more depth and analyze the micro-mechanism of wear in detail. The surface friction and wear morphology of the nitrided
Figure 8 1Cr17Ni2 steel was analyzed.
Table 4 shows the surface morphology of the original sample after friction and wear, and
Figure 8b shows the EDS test results for points A, B, and C. As can be seen from the figure, as the friction and wear progress, the surface of the wear scar is covered with an oxide film, which is formed by the compaction of wear debris generated during the wear process during movement. This oxide film can effectively prevent direct contact between the sample and the friction pair, thereby reducing the friction coefficient. The results of energy spectrum analysis also indicate that the wear debris contains oxides, suggesting that the type of oxide is Fe
2O
3. With the cycle of the friction process, the oxide film gradually thickens and fatigue cracks are generated. Over time, the fatigue cracks gradually increase and fatigue peeling occurs, exposing a new surface to contact with the friction pair, eventually reaching an equilibrium. There are shallow ploughing grooves on the surface of the wear scar (point A in
Figure 8a), indicating that there is abrasive wear on the surface. Meanwhile, observing point B in
Figure 8a, it can be seen that there is slight tearing on the surface of the wear scar, indicating that there may be adhesive wear on the surface. The adhesive ensures that the surface wear rate does not increase rapidly. In summary, the main wear mechanisms of 1Cr17Ni2 steel in the quenched state are slight abrasive wear and oxidative wear, in addition to slight adhesive wear.
In order to reveal the wear mechanism, it is necessary to observe the wear behavior first.
Figure 9 shows the surface friction and wear morphology after different nitriding processes, with typical areas of wear scars enlarged. As shown in
Figure 9a, there is a slight ploughing groove inside the wear scar, indicating that there is slight abrasive wear during the wear process. By magnifying the edge morphology of the wear scar, it can be seen that there are obvious plastic deformation zones and crack generation zones at the edge of the wear scar, indicating that fatigue wear has occurred at the edge of the wear scar. From
Figure 9b, it can be seen that there is a clear ploughing pattern inside the wear scar, indicating that there is abrasive wear during the wear process. At the same time, by zooming in on the image of the edge of the wear scar, it can be seen that there are significant plastic deformation zones and crack formation zones at the edge of the wear scar, indicating that fatigue wear has occurred at the edge of the wear scar. Compared to the wear effect of the original sample, there are no cracks at the center of the wear scar after nitriding, indicating that the nitrided layer has good fatigue performance. The reason for the formation of cracks at the edge is that the wear debris in the center and the oxides generated during the wear process are pushed to the edge of the wear scar by the friction pair during the wear process and are compacted. As the friction and wear test time increases, the compacted wear debris at the edge of the wear scar has poor fatigue performance and gradually produces fatigue cracks. As time further increases, the cracks expand and eventually fall off. In summary, the main wear mechanisms of the sample nitrided at 500 °C for 8 h are particle wear and oxidative wear. Upon observation of
Figure 9c, it is evident that the wear morphology of 550 °C nitriding for 8 h bears a resemblance to that of 460 °C nitriding for 16 h. However, in terms of wear depth and wear rate, the 460 °C nitriding for 8 h exhibits a lower wear rate and a shallower wear depth, potentially related to the surface phase structure. By comparing the SEM images of the three processes, it can be concluded that as the nitriding temperature increases, the wear mechanism changes from abrasive wear and oxidative wear to oxidative wear.
Figure 10 shows the surface and edge micro-morphology of the friction and wear experiments after nitriding at 500 °C for 8 h and 500 °C for 16 h. From
Figure 10b, it can be seen that the ploughing shape inside the wear scar has decreased, and almost no obvious ploughing marks are visible, indicating that no abrasive wear has occurred at this time. Cracks also appear on the edge of the wear scar, and the mechanism of occurrence is consistent with that of 500 °C nitriding for 8 h, indicating that the main wear mechanism of the 500 °C nitriding for 16 h sample is oxidative wear. It can be seen that with the extension of nitriding time, the wear mechanism changes from abrasive wear and oxidative wear to oxidative wear.
3.4. Crystal Cell Optimization
α′
N (over-saturated nitrogen martensite) is the martensitic structure formed when nitrogen is in a supersaturated state solidly dissolved in the octahedral interstitial sites of α-Fe. Therefore, this study selects α-Fe as the basis for constructing the model. Many studies on surfaces have focused on the atomic level, including surface segregation and surface magnetic structures of iron-chromium alloys. Therefore, we adopt the surface structure of the bcc Fe-Cr system as a model reference for the surface of 1Cr17Ni2 stainless steel [
16,
30]. The body-centered cubic (bcc) structure of α-Fe allows the interstitial atoms to more easily occupy octahedral sites and maintain a relatively stable structure. Alloying elements in martensitic stainless steel mainly exist in the form of substitutional solid solution. As the temperature increases, the energy of the system also increases, making it easier for alloying elements to move, thus increasing the likelihood of reacting with nitrogen. However, most alloying elements have significantly different crystal structures from α-Fe. Therefore, it is almost impossible for the reaction to occur when the temperature cannot reach the critical value. First-principles calculations are based on the system energy at absolute zero, so it is assumed that the Cr atoms are randomly and disorderly distributed, while nitrogen is always located in octahedral interstitial sites. Therefore, based on the Cr content in the 1Cr17Ni2 martensitic stainless steel, a 2 × 2 × 2 supercell is established as the calculation model.
To investigate the influence of different concentrations of nitrogen solid solution on the nitriding of 1Cr17Ni2 martensitic stainless steel, a nitriding model was established for the supercell due to the mass percentage range of 16% to 18% for chromium (Cr) in 1Cr17Ni2 stainless steel. The model was Fe
13Cr
3, with a mass percentage of 17.65% for Cr, which is close to the real composition. Three original Fe atoms in the supercell were replaced by Cr atoms, taking into account the symmetry of the crystal structure. Additionally, varying numbers of N atoms were added to the octahedral interstitial sites. The nominal chemical formula for this series of structures is Fe
13Cr
3N
n. The optimized structure model of Fe
13Cr
3N
n after computational structural optimization is shown in
Figure 11. Meanwhile, the composition table of Cr and N contents in the Fe
13Cr
3N
n structural series with varying nitrogen contents is shown in
Table 5.
Since there are two main forms of alloy elements in steel, most of the alloy elements will replace the lattice points of iron to form substitutional solid solutions; in addition, there will be alloy elements undergoing precipitation reactions to form alloy nitrides. Due to the significant difference in lattice constants between most alloy nitrides and the -Fe matrix, the formation of such alloy nitrides requires a significant driving force, making it generally difficult for them to form. The results obtained from first-principles calculations are based on the assumption that alloy elements in the alloy are randomly distributed and nitrogen atoms are located in the interstitial positions at 0 K. When constructing the Fe-Cr system, the influence of chromium element content on the stability of nitrogen-containing martensitic phase needs to be considered first, so the content of other elements is transformed into chromium equivalents according to the electronegativity rules.
Assuming that the chromium content remains unchanged in each model, a body-centered cubic (BCC) structure of α-Fe single crystal cell is established, and an ideal Fe-Cr series structure is then constructed based on this. Under the premise of maintaining symmetry, the different amounts of nitrogen atoms are sequentially added to the octahedral interstitial positions around the BCC atoms. The total energy of the supercell for different Fe-Cr-N systems is calculated through first-principle calculations, and the effect of nitrogen on the stability of the system structure is analyzed. The equilibrium lattice constants obtained after structure optimization are shown in
Table 6.
By analyzing the data in
Table 6, it is easy to determine that the original α-Fe single crystal cell has a standard body-centered cubic structure. Under the condition of maintaining symmetry, by sequentially adding N atoms to the octahedral interstitial sites of the α-Fe supercell, the significant distortion of the original crystal structure can be observed, resulting in a transformation from cubic system to tetragonal system. Moreover, with an increase in nitrogen content, there is a continuous expansion in the cell volume. This also corresponds to the leftward shift of the γ′-Fe
4N diffraction peak in the XRD.
3.5. Structural Stability Analysis of the Fe-Cr-N Series
The total energy of the supercell system obtained from first-principles calculations provides an intuitive representation of the variation in structural energy of the Fe-Cr-N system. In this study, the cohesive energy (Ec) and the formation energy (Ef) were used to measure the thermodynamic stability and the ease of formation of the Fe-Cr-N system structures.
Formation energy refers to the energy released or absorbed when a reaction occurs between stable existing solid-state reactants and products. For exothermic reactions, the formation energy is negative. Cohesive energy refers to the energy released or absorbed when individual atoms combine from their free state to form a crystalline structure. The smaller the absolute value of cohesive energy, the less stable the system.
For a compound with the molecular formula
AxByCn, the calculation formulas for cohesive energy and formation energy are as follows [
31,
32]:
where
Etot is the total energy of the supercell (eV) calculated by first principles [
33,
34];
Esolid is the average energy (eV) of a single atom under the condition of stable solid crystal formation from a single substance;
Eatom is the energy of a single atom in a free state (eV); and
x,
y, and
n are the number of
A,
B, and
C atoms (where
A is Fe, and the range of values for
x is 13;
B is Cr, and the range of values for
y is 3; and
C is N, and the range of values for
n is 0, 2, 4, 6). The specific energy values required for the calculations are shown in
Table 7.
Regarding the calculations of binding energy and formation energy in the iron–chromium–nitrogen system, we found that in the series of iron–chromium–nitrogen structures, the total energy of the system decreases as the nitrogen atoms are incorporated. From the energy
Table 8 system, we can see that the binding energies are negative, indicating that this series of structures can be formed by the combination of free atoms.
As the number of nitrogen atoms increases, the difficulty of binding increases. Looking at the formation energy, in the original α-Fe supercell and the supercell where only one Fe atom is replaced by a Cr atom, the formation energies for systems without nitrogen atoms and with two incorporated nitrogen atoms are negative, indicating good stability of these two structures. The formation energies for four and six incorporated nitrogen atoms are positive, but their values are small, suggesting the possibility of forming metastable states of these structures. In the supercell systems where two to four Fe atoms are replaced by Cr atoms, all the formation energies are positive. However, when the nitrogen content is two atoms, the formation energy decreases, and the absolute value of the binding energy increases significantly, indicating that the supercell with a nitrogen content of two atoms is more stable than the original supercell structure. But looking at the trend, for the same supercell system, the formation energy increases as the number of incorporated nitrogen atoms increases, and nitrideing systems with a lower nitrogen content have better stability. The absolute value of the binding energy gradually decreases as the number of nitrogen atoms increases, indicating that systems with lower nitrogen solubility have better stability. The specific energy calculation values for the relevant unit cells are shown in
Table 8.
3.6. Analysis of Fe-Cr-N State Density and Differential Charge Density of Fe-Cr-N
The Fermi level (EF) is the highest energy level occupied by electrons at absolute zero temperature, and many properties of crystals are determined by the distribution of valence electrons near the Fermi level. Two peaks can be observed on both sides of the Fermi level, and the DOS between the two peaks is non-zero. This low-energy DOS region is called a pseudo-gap, which is typically represented by a V-shaped structure [
33]. The stability of a system can be judged to some extent by the presence of a pseudo-gap near the Fermi level, or by the magnitude of the electron density of states M(EF) near the Fermi level. The higher the coincidence between the pseudo-gap and the Fermi level, or the lower the electron density of states near the Fermi level, the lower the system energy, indicating a more stable state.
The comparison chart of total density of states for the Fe
13Cr
3N
n (n = 0, 2, 4, 6) series of structures is shown in
Figure 12. Pseudo-gaps exist near the Fermi level for Fe
13Cr
3 systems with different nitrogen contents. It can be seen that the introduction of nitrogen atoms in the system causes a density of states peak at around −16 eV. The Fe
13Cr
3 and Fe
13Cr
3N
2 systems in the figure show a pseudo-gap that coincides with the Fermi level. By comparison, increasing the nitrogen content causes fluctuations in the electron density near the Fermi level. M(EF) increases when two nitrogen atoms are introduced but starts to decrease when four nitrogen atoms are introduced. Moreover, the pseudo-gap peak shifts away from the Fermi level towards higher energy. This indicates that when only two nitrogen atoms are added to the system, the large lattice distortion reduces the structural symmetry, leading to a decrease in system stability. As the nitrogen content continues to increase, the symmetry of the unit cell partially recovers, but the stability of the system decreases, consistent with the previously calculated formation energy and binding energy results.
The partial density of states for Fe, Cr, and N in the Fe
13Cr
3N
n system is shown in
Figure 13. The PDOS peaks near the Fermi level for the Fe
13Cr
3 alloy appear at around ±1.4 eV, with a peak value of approximately 33 e/eV. For the Fe
13Cr
3N
2 alloy, the PDOS peaks near the Fermi level appear at around ±1.3 eV, reaching a peak value of around 30 e/eV. The PDOS peaks near the Fermi level for Fe
13Cr
3N
4 appear at around ±2 eV, with a peak value close to 36 e/eV. For the Fe
13Cr
3N
6 alloy, the PDOS peaks near the Fermi level appear at around ±2.1 eV, with a peak value of approximately 40 e/eV. The majority of these alloy’s PDOS peaks originate from the 3d orbitals of Fe and Cr atoms. In the Fe
13Cr
3N
n alloy, the peak at −72 eV mainly originates from the s orbital of Cr atoms, the peak at −43 eV mainly comes from the p orbital of Cr atoms, and the peak at −16 eV mainly originates from the s orbital of N atoms. The DOS curve at 11eV primarily originates from the p orbitals of Fe and Cr atoms. At −7 eV, there is a hybridization phenomenon between the p orbitals of N atoms and the s orbitals of Fe and Cr atoms, indicating a strong covalent interaction among Fe, Cr, and N atoms. The overlap of DOS near the −7 eV level with the Fermi level for Fe
13Cr
3N
4 indicates a relatively stable structure. The narrowing of the DOS width for Fe
13Cr
3N
6 suggests a decrease in covalent interaction.
As shown in
Figure 14, it can be observed that in the Fe
xCr
yN
n series structures, in series structures such as Fe
13Cr
3N
n, there is a strong bonding interaction between Fe and Cr atoms on different crystal planes. In the Fe
xCr
yN
n system with n = 0, Cr atoms gain electrons and carry negative charges, while Fe atoms lose electrons and carry positive charges. After the N atoms infiltrate into the crystal cell, the bonding interaction between Fe and Cr atoms weakens, and the bonding interaction between Cr-N and Fe-N atoms strengthens. When only two N atoms enter the system, Cr atoms are still in a relatively electron-attracting state, while Fe atoms gradually transition from being electron-losing to electron-attracting. As the number of N atoms in the system increases, N becomes the side that is more prone to losing electrons. When six N atoms are infiltrated, Fe atoms in the face-centered position are more likely to gain electrons compared to Cr atoms, carrying a higher negative charge. The analysis of Fe-Cr-N state density and differential charge density of Fe-Cr-N explains the mechanism of surface modification hardening.
This experiment also calculated the elastic properties of the nitride phases of 1Cr17Ni2 steel using first-principles computational methods to further verify the experimental results. The mechanical properties of Fe
13Cr
3N
2, γ′-Fe
4N, ε-Fe
2-3N, and CrN were calculated, including bulk modulus (B), shear modulus (G), Young’s modulus (E), Pugh ratio (G/B), Poisson’s ratio (V), and hardness (Hv). According to the generalized Hooke’s law, the stiffness coefficient calculation results and mechanical property calculation results are shown in
Table 9 and
Table 10.
As shown in
Table 10, the Fe
13Cr
3N
2 phase exhibits the highest hardness, reaching up to 16.53 GPa, which is in good agreement with the actual measurement values obtained by Yixue Wang. The Pugh ratio G/B of different nitrided phases in 1Cr17Ni2 steel is less than 0.57, so they all have good toughness and are good ductile phases. At the same time, Poisson’s ratio v indicates that Fe
13Cr
3N
2 and ε-Fe
2-3N have strong rigidity and high microhardness, which also matches the phase composition of the high-hardness nitrided surface at 460 °C. The ratio of hardness to Young’s modulus can better display the wear resistance of materials than a single hardness. The higher the value of H/E, the better the wear resistance of the material [
35]. The higher the value of H
3/E
2, the lower the probability of plastic deformation. From
Table 10, it can be seen that Fe
13Cr
3N
2 and ε-Fe
2-3N phases have lower wear rates, which can also be verified by the wear rate calculated in
Figure 6. It can be inferred that by adjusting the nitriding temperature and time, a nitrided layer with excellent toughness and hardness can be obtained. This result is also consistent with the conclusion reached by Gonzalez-Moran in this article that nitrogen-expanded martensite has excellent wear resistance and is a rigid elastic solid with high elastic properties and resistance to plastic deformation [
36,
37,
38,
39,
40].