2020 Volume 61 Issue 6 Pages 1070-1076
In order to clarify the effects of carbon concentration change on the growth of the compound layer in nitrided steel, the variation with time of the concentration of alloy elements on the surface layer and the phases of the compound layer were investigated in nitrided steels containing various amounts of carbon in the matrix. It was found that the variation with time of the carbon concentration in the compound layer was mainly responsible for the variation with time of the compound layer microstructure. Furthermore, it was discovered that the variation with time of the carbon concentration of the compound layer was caused by both the migration of carbon from the matrix to the compound layer and from the surface of the compound layer to the atmosphere. Thus, the gradient of the chemical potential of carbon in the through-thickness direction of the compound layer and the compound layer microstructure changed with nitriding time.
This Paper was Originally Published in Japanese in J. Japan Soc. Heat Treat. 59 (2019) 61–66. Captions of figures are slightly changed.
In recent years, the need for automobiles with higher fuel efficiency and reduced noise production has increased the demand for smaller and lighter transmission systems with additional stages. This, in turn, has increased the demand for automobile parts with higher accuracy and strength. Carburizing, which makes use of martensite transformation by quenching from the austenite region, is the most commonly used heat treatment for surface hardening. However, steel parts are inevitably deformed by the distortion produced by this heat treatment.
During nitriding and nitrocarburizing, steel is heated to a ferrite region of 450–600°C. Atomic nitrogen is formed from the decomposition reaction of NH3 gas and fed into the nitriding furnace. The formed N diffuses through the steel surface causing the precipitation of Cr, Al, and other element nitrides, and forming a solid solution in the matrix, thereby hardening the steel surface.1) As these heat treatments do not accompany phase transformation, the treated steel parts present small deformation and high assembly accuracy, and are effective in reducing noise production in automobiles.
When steel is nitrided, a layer referred to as “compound layer”, mainly composed of iron nitrides (i.e., the ε-Fe2-3N phase and the γ′-Fe4N phase), is formed on the steel surface. The compound layer has less deformability than that of the matrix and may easily peel and induce surface fatigue fracture.2,3)
The development of techniques for controlling nitriding furnace atmospheres and analyzing microstructures in recent years has confirmed that the fatigue strength of steel depends on the types of phases that comprise the compound layer.4–6) For example, Kobayashi et al.4) nitrided JIS G 4051 Grade S35C steel, a carbon steel for machine structural use, and prepared a nitrided steel with a compound layer consisting of the γ′ phase on the furnace atmosphere side and the ε phase on the matrix side. The nitrided steel presented higher surface fatigue strength than that of conventional nitrided steels with a compound layer mainly composed of the ε phase.4) To strengthen nitrided steel parts, therefore, it is vital to establish guidelines for controlling the phases of the compound layer.
Phase control of the Fe–N binary compound layer has not been widely addressed in the literature. Under general nitriding conditions, the phases in the compound layer are ε and γ′. From an Fe–N binary equilibrium phase diagram7) and an Fe–N binary Lehrer diagram,8) it is evident that the ε phase and the γ′ phase are stable in the high and low N concentration regions, respectively. The structure of the Fe–N compound layers may be therefore composed of either the γ′ phase or the ε phase on the furnace atmosphere side and the γ′ phase on the matrix side.
Some scholars have investigated on the composition of Fe–C–N alloys compound layers. For example, Slycke et al.9) and Kunze10) reported Fe–C–N ternary equilibrium phase diagrams for 575°C. Du11) and Hiraoka et al.12) reported calculated Fe–C–N ternary Lehrer diagrams. These diagrams show that C is an ε phase stabilizing element.
Commercial alloy steels might present phases different from those of Fe–N compound layers. Hiraoka et al.13) reported that when JIS G 4053 Grade SCM435 steel was nitrided, compound layers in equilibrium with the matrix presented either the ε phase or the γ′ phase. Asada et al.14) investigated the effect of Cr addition to the phase of the compound layer by using an Fe–Cr–0.2C alloy and reported that the compound layer in equilibrium with the matrix changed from being mainly composed of the γ′ phase to being mainly composed of the ε phase. However, these differences are difficult to interpret using the aforementioned equilibrium phase diagrams and pose major obstacles to the extended application of nitrided steels with controlled compound layer phases. Therefore, it is of significant importance for the industry to clarify the factors governing the formation of the compound layer during nitriding, and to establish a unified view on the mechanism of compound layer formation.
In commercial steels, a possible cause for the change in the compound layer microstructure when in equilibrium with the matrix is the effect of decarburization from the steel surface to the furnace atmosphere; decarburization is an unavoidable phenomenon that occurs when steel is heated. Mittemeijer and Somers15) reported that decarburization also occurred when steel was nitrided. This could indicate that the C concentration in the compound layer changes over time. In this study, therefore, we investigated the changes in the concentrations of alloying elements and phase distribution in the surface layer over time when steels with different matrix C concentrations were nitrided.
Table 1 shows the chemical composition of the steels used in the experiments. Based on JIS G4051 Grade S30C steel, Fe–C–N alloys with C contents of 0, 0.1, 0.3, and 0.8 mass%, hereinafter C00, C10, C30, and C80, respectively, were used as the experimental materials. A 10 kg ingot of each type of steel was melted in a vacuum induction furnace and hot forged into Ø25 mm bars. Then, based on JIS G0561, C00 and C10, C30, and C80 bars were normalized by heating to 925, 900, and 870°C, respectively, for 30 min. Specimens measuring 10(width) × 50(length) × 2(thickness) mm were machined from the bars and gas nitrided. A mixture of NH3 and N2 was used for gas nitriding. The specimens were held at 570°C for 0.5, 1.0, and 3.0 h and cooled in 80°C oil. The gas nitriding furnace atmosphere was maintained at atmospheric pressure. The partial pressure of H2 (PH2) produced by the decomposition of NH3, as expressed in eq. (1), was measured in real time with a thermal conductivity H2 sensor. A glass tube NH3 analyzer was used to measure the partial pressure of NH3 (PNH3). The nitriding potential16) (KN) was calculated via eq. (2) from PH2 and PNH3. The flow rates of NH3 and N2 were adjusted to maintain KN at 1.1 atm−1/2 (target value). Table 2 shows the average values of PH2, PNH3, and KN measured every 10 min. Figures 1(a) and 1(b) show the Lehrer diagrams of C00 and C80, respectively, as calculated by the integrated thermodynamic calculation system Thermo-Calc 2016b (thermodynamic databases: TCFE version 8.1 and Fe-DATA version 6). For comparison, an experimental Lehrer diagram8) is shown superimposed in Fig. 1(a). It was found that C00 and C80 were nitrided near the boundaries of the ε phase and the γ′ phase and in the ε phase region. The cross-sectional surface layers of the nitrided specimens were etched in a 3% nital solution and were then observed with an optical microscope and a scanning electron microscope (SEM). The phases of the compound layers were identified by electron backscatter diffraction (EBSD). The N, C, Si, and Mn contents of the compound layers were measured with a glow discharge optical emission spectrometer (GD-OES).
| \begin{equation} \text{NH$_{3}$} \leftrightarrow [\text{N}] + 3/2\text{H$_{2}$} \end{equation} | (1) |
| \begin{equation} K_{\text{N}} = P_{\text{NH}_{3}}/(P_{\text{H}_{2}})^{3/2}\quad [\textit{atm}^{-1/2}] \end{equation} | (2) |
Comparison between calculated Lehrer diagrams for (a) Fe–N binary alloy (C00) and (b) Fe–0.8C–N ternary alloy (C80) and experimental measurements.
Figure 2 shows the optical micrographs of the cross-sectional surface layers of the nitrided specimens. White unetched layers, apparently compound layers, were observed on the surface of all the steel specimens for all nitriding times. The thickness of the compound layers increased with increasing the matrix C content and nitriding time.
Light optical micrographs showing cross sections of test specimens after etching with 3% nital and nitriding for 0.5, 1.0, and 3.0 h at 570°C and KN = 1.1 atm−1/2.
Figure 3 shows the SEM images and the EBSD-determined phases of the surface layers of the nitrided C00 specimens. The C00 specimen nitrided for 0.5 h presented a single ε phase, a mixed ε+γ′ phase, and a single γ′ phase in the surface layer, the deeper layer, and the layer adjacent to the matrix, respectively. This phase distribution did not change as nitriding proceeded; however, both the single γ′ phase region and the single ε phase region grew.
Scanning electron microscope (SEM) images and phase map by electron back scatter diffraction (EBSD) showing cross sections of 0.0 mass%C steel specimens nitrided for 0.5, 1.0, and 3.0 h.
Figure 4 shows the SEM images and the EBSD-determined phases of the surface layers of the nitrided C10 specimens. The single ε phase and the mixed ε+γ′ phase regions were respectively observed in the surface layer and the deeper layer of the C10 specimen nitrided for 0.5 h. In the C10 specimens nitrided for 1.0 and 3.0 h, the single ε phase was observed in the surface layer. In the central region of the compound layer, the ε phase proportion decreased and the γ′ phase became the dominant phase. The γ′ phase was dominant in the region adjacent to the matrix, while the ε phase was heterogeneously distributed throughout the compound layer.
Scanning electron microscope (SEM) images and phase map by electron back scatter diffraction (EBSD) showing cross sections of 0.1 mass%C steel specimens nitrided for 0.5, 1.0, and 3.0 h.
Figure 5 shows the SEM images and the EBSD-determined phases of the surface layers of the nitrided C30 specimens. The entire compound layer of the C30 specimen nitrided for 0.5 h consisted of the ε+γ′ phase, where the ε phase was dominant. In the C30 specimens nitrided for 1.0 and 3.0 h, the γ′ phase was dominant near the center of the compound layer, as observed in the nitrided C10 specimens. The ε phase proportion in the region adjacent to the matrix was higher than that in the C10 specimens. A single ε phase, a single γ′ phase, and a mixed ε+γ′ phase were observed from the surface layer to the region adjacent to the matrix in the through-thickness direction.
Scanning electron microscope (SEM) images and phase map by electron back scatter diffraction (EBSD) showing cross sections of 0.3 mass%C steel specimens nitrided for 0.5, 1.0, and 3.0 h.
Figure 6 shows the SEM images and the EBSD-determined phases of the surface layers of the nitrided C80 specimens. The compound layers of the C30 specimens nitrided for 0.5 and 1.0 h entirely consisted of the ε+γ′ phase, where the ε phase was dominant. The phase of the C80 specimen nitrided for 3.0 h was similar to that of the C30 specimen nitrided for 3.0 h. However, the γ′ phase proportion near the center of the compound layer and the ε phase proportion in the region adjacent to the matrix were higher than those in the C30 specimen nitrided for 3.0 h.
Scanning electron microscope images (SEM) and phase map by electron back scatter diffraction (EBSD) showing cross sections of 0.8 mass%C steel specimens nitrided for 0.5, 1.0, and 3.0 h.
Figure 7 shows the variation with time of the N concentration profiles of the nitrided C00, C10, C30, and C80 specimens. The N concentration of the nitrided specimens was highest in the surface region, decreased with increasing the distance from the surface, and suddenly dropped at a distance equivalent to the compound layer thickness. The N concentration did not vary significantly with increasing nitriding time; it was approximately 7.5, 6.5, and 5.7 mass% in the surface, the center of the compound layer, and immediately before the sudden drop, respectively.
Profiles of nitrogen content measured by glow discharge optical emission spectrometry of steel specimens nitrided for 0.5, 1.0, and 3.0 h.
Figure 8 shows the variations in the C concentration profiles of the nitrided C00, C10, C30, and C80 specimens with time. The C10, C30, and C80 specimens presented higher C concentrations in the region adjacent to the matrix than in the matrix. It is presumed that this was caused by C enrichment from the matrix into the compound layer.17) The C concentration peaks tended to decrease with increasing nitriding time. The C concentration in the outermost surface was lower than in the matrix. This suggests that decarburization from the compound layer to the furnace atmosphere occurred during nitriding. The decrease in C concentration due to decarburization increased with increasing nitriding time.
Profiles of carbon content measured by glow discharge optical emission spectrometry of steel specimens nitrided for 0.5, 1.0, and 3.0 h.
Figure 9 shows the variation with time of the Si and Mn concentration profiles of the nitrided C80 specimens. Silicon and Mn concentrations were uniform in the through-thickness direction. The partition of Si and Mn to the compound layer was not observed, indicating that the compound layer grew in non-partition local equilibrium or paraequilibrium.
Profiles of silicon and manganese content measured by glow discharge optical emission spectrometry of 0.8 mass%C steel specimens nitrided for 0.5, 1.0, and 3.0 h.
As can be seen from the Fe–C–N ternary equilibrium phase diagrams and calculated Lehrer diagrams (Figs. 1 and 2), the stability of the ε phase increases as the C concentration in the compound layer increases. Since the partition of substitutional elements was not observed at the interface between the compound layer and the matrix, the variation with time of compound layer microstructure is expected to be strongly affected by interstitional elements. To explore the possibility of explaining the variation with time of the compound layer phases only by the variation in N and C concentrations, the compound layer microstructure was calculated from the measured N and C concentrations in the through-thickness direction and compared with the experimentally obtained results.
The calculations were performed using Thermo-Calc 2016b (thermodynamic database: TCFE version 8.1) and the measured N and C concentrations in the nitrided C80 specimens were regarded as representative of the nitrided specimens. Silicon and Mn were ignored, and it was assumed that the stable phases in the Fe–C–N ternary system were in complete equilibrium.
In Fig. 10, the equilibrium volume fraction of the ε phase calculated from the measured N and C concentrations in the nitrided C80 specimens is compared with that determined from the SEM-EBSD analysis. For each nitriding time, the volume fraction of the ε phase thermodynamically calculated from the N and C concentrations accurately agreed with the volume fraction of the ε phase determined by SEM-EBSD analysis. This fact supports the hypothesis that the variation with time of the microstructure of the compound layer is linked to the variation in the N and C concentrations in the compound layer.
Profiles of volume fraction of the ε phase in the compound layer of 0.8 mass%C steel specimens nitrided for 0.5, 1.0, and 3.0 h, measured by electron back scatter diffraction (EBSD) and calculated by Thermo-Calc 2016b software using nitrogen and carbon content profiles.
In this section, we discuss the effect of the C concentration in the matrix on the variation with time of the C concentration in the compound layer.
Figure 11 shows the equilibrium profiles of the N and C chemical potentials, μN and μC, calculated from the measured N and C concentrations of the nitrided C80 specimens. For each nitriding time, μN and μC in the steel surface adjacent to the nitriding furnace atmosphere were almost constant. While μN was higher in the steel surface than in the steel bulk, μC was lower in the steel surface than in the steel bulk. In addition, since the thickness of the compound layer increased with increasing nitriding time, the slopes of the μN and μC profiles in the through-thickness direction decreased with increasing nitriding time. Therefore, the driving force of the change in the N and C concentrations in the through-thickness direction of the compound layer decreased with increasing nitriding time.
Profiles of chemical potential in the compound layer of 0.8 mass%C steel specimens nitrided for 0.5, 1.0, and 3.0 h, calculated by Thermo-Calc 2016b software using nitrogen and carbon content profiles.
Based on the above findings, the relationship between the variation with time of the C concentration at each nitriding time and the microstructure of the compound layer is addressed in the following paragraphs.
The C concentration in the compound layer side of the interface between the matrix and the compound layer in the C10, C30, and C80 specimens nitrided for 0.5 h was higher than that in the matrix. This was probably because the driving force of the change in C concentration was large in the through-thickness direction of the compound layer and C enrichment from the matrix into the compound layer was significant appeared. The C concentration near the surface of the compound layer, where decarburization exerts its largest effect, was higher than that in the matrix and was approximately 0.5 mass% or greater. The C diffusion rate difference between the compound layer and the matrix can be cited as a cause for this phenomenon.18) Since the C diffusion rate became significantly smaller in the compound layer than in the matrix, the compound layer after 0.5 h of nitriding is considered to have consisted almost entirely of the ε phase under the effect of C enrichment from the matrix.
In the C10 and C30 specimens nitrided for 1.0 h, the C concentration dropped to 0.2–0.3 mass% near the center of the compound layer. An explanation regarding this finding is provided as follows. As the compound layer increases in thickness, the slope of the μC profile decreases in the through-thickness direction and the amount of C supplied by the matrix decreases. In addition, since the C extracted from the matrix during 0.5 h of nitriding is diluted in the compound layer, the effect of decarburization from the compound layer to the furnace atmosphere is relatively pronounced in the 0.5 h nitrided specimens. As a result, both the N and C concentrations decrease and the γ′ phase is stabilized near the center of the compound layer. In the C80 specimens, the C concentration is high in the matrix and a large amount of C is supplied into the compound layer. The C concentration near the center of the compound layer is thus relatively high (approximately 0.6 mass%). This is thought to have allowed the ε phase to stably exist throughout the compound layer, even after 1.0 h of nitriding.
As the compound layer increased in thickness after 3.0 h of nitriding, the amount of C supplied into the compound layer decreased even for the C80 specimens, and the effect of decarburization became relatively pronounced. This reduced the C concentration in the compound layer to the same level as that in the matrix and also reduced the C concentration near the center of the compound layer to approximately 0.3 mass%. The γ′ phase is thus considered to have been stabilized near the center of the compound layer.
In summary, it was clarified that the variation with time of the C concentration of the compound layer was mainly responsible for the variation with time of the compound layer microstructure. Moreover, it was shown that the C concentration in the compound layer varied with time because C enrichment from the matrix into the compound layer and decarburization from the surface near the compound layer into the furnace atmosphere were both affected by the change in the slope of the μC profile in the through-thickness direction.
To clarify the effect of C concentration change on the growth of the compound layer on the surface of nitrided steels, the variation with time of the concentration of alloying elements in the surface layer and the variation with time of the compound layer microstructure were investigated by nitriding steels of different matrix C concentrations. The conclusions reached are presented as follows:
