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Vol. 9. Issue 2.
Pages 2331-2337 (March - April 2020)
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Vol. 9. Issue 2.
Pages 2331-2337 (March - April 2020)
Original Article
DOI: 10.1016/j.jmrt.2019.12.064
Open Access
On nitrogen diffusion during solution treatment in a high nitrogen austenitic stainless steel
Rui Zhoua, Derek O. Northwoodb, Cheng Liua,
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Corresponding author.
a College of Mechanical Engineering, Yangzhou University, Yangzhou, Jiangsu, PR China
b Mechanical, Auto and Materials Engineering, University of Windsor, Windsor, Ontario, Canada
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Tables (1)
Table 1. Chemical composition of HNASS (wt.%).

The generation and microstructural characteristics of the surface layer formed during solution treatment at 1200°C of a high nitrogen austenitic stainless steel, have been investigated using optical microscopy, X-ray diffraction (XRD) and scanning electron microscopy (SEM) with energy dispersive X-ray spectroscopy (EDS). Contrasting with a bulk microstructure of austenite and ferrite, a multi-layer surface structure is obtained, which consists of a white outer-surface layer which is ferrite and a grey subsurface layer with austenite. Formation of the multi-layer structure occurs through phase transformations between austenite and ferrite due to the various levels of nitrogen diffusion. With increasing solution treatment time, the thicknesses of both the surface and the subsurface layer increase. Hardness tests show that the microhardness increases from the surface layer, to the subsurface layer and then to the bulk. This is attributed to the combined effects of solid solution strengthening from nitrogen and the grain size of austenite or ferrite.

High nitrogen austenitic stainless steel
Solution treatment
Multi-layer structure
Nitrogen diffusion
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High nitrogen austenitic stainless steels (HNASS), developed by partial or total replacement of nickel (Ni) through nitrogen (N), have been widely applied in manufacturing and various other fields including biomedical, ocean engineering and armor because of a good combination of mechanical properties, high resistance to corrosion and low cost [1–5]. N is known as a strong austenitic stabilizer element. The interstitial N atom has a large solid solution strengthening effect, leading to an increase in the strength of stainless steels without deteriorating their ductility and toughness [6–9]. However, the microstructure of these steels may contain some secondary phases in addition to austenite, such as δ ferrite, nitrides (Cr2N), the χ phase (an Fe, Cr, Mo intermetallic compound) and the σ phase (σ–Fe–Cr) due to the segregation of Cr, Mo, N and Mn during casting and forging [10–12], which worsens the mechanical performance and causes premature failure [13–15].

Solution treatment is a well-established process to reduce, or prevent, the formation of the second phases and improve the strength and ductility [16–18]. It was reported by Hao et al. that the banded microstructure of an austenitic stainless steel formed after hot rolling with a slow cooling rate, could be eliminated through solution treating at 1100°C for 30min [19]. Cojocaru et al. found that the σ phase decreased with increasing solution treatment temperature and could be completely removed from a duplex stainless steel if the temperature was above 1000°C [14]. Ren et al., however, observed the formation of σ phase and a reduction in austenite when the solution treatment temperature was higher than 1050°C in both Cr22Ni5Mo3MnSi and Cr30Ni7Mo3MnSi stainless steels [20]. It was shown that the austenitic grain size increased and the nitrides dissolved after a solution treatment at 1150°C for 1h and water quenching, leading to a decrease in the hardness in an Fe-18Cr-16Mn-2Mo-0.66N austenitic stainless steel [21]. Although the effects of solution treatment on microstructure, grain size and mechanical properties of stainless steels have been investigated, relatively little work has been carried out to gain a better understanding of the surface microstructure of HNASSs during solution treatment at high temperature so as to improve their mechanical behavior and assist in the development of industrial processing. Ran et al. demonstrated that the outward diffusion of both Mn and N occurred during solution treating at 1100°C, which caused a formation of ferritic diffusion layer on the surface of a Cr–Mn–Ni–N duplex stainless steel [22]. It was believed that the microstructure and the N content of HNASS were closely related to the external nitrogen atmosphere during solution treatment [23,24].

In the present work, the generation and microstructural characteristics of the surface layers in a HNASS have been examined after solution treatment. Possible mechanisms for the formation of the multilayer structure are discussed.

2Experimental details

The chemical composition of the hot-rolled HNASS used in this study is given in Table 1. Samples with dimensions of 20×20×10mm, were cut longitudinally to the rolling direction and treated in a SG-QF1700 atmosphere muffle furnace operating at a base pressure of −0.1MPa at 1200°C for 30, 60, 90, 120 and 240min respectively, and then quenched into water at room temperature. The microstructure was observed by optical microscopy (OM, LEICA GQ-300) and a Gemini-SEM-300 scanning electron microscopy (SEM). The distribution of chemical elements was analyzed by energy dispersive X-ray spectroscopy (EDS) system. X-ray diffraction (XRD) analysis was carried out to determine the layer constituent phases through peeling off the sample layer by layer according to the microstructure observation. An image-pro plus 6.0 software was used to determine the depth of each layer, phase content and grain sizes of the austenite and the ferrite obtained by solution treatment.

Table 1.

Chemical composition of HNASS (wt.%).

Si  Mn  Cr  Mo  Ni  Fe 
0.02  0.21  19.44  0.018  0.01  0.65  18.93  1.63  0.31  Bal. 
3Results and discussion

Fig. 1 shows typical OM micrographs of the as-received sample. As seen in Fig. 1(a), a mixed microstructure comprising blocky austenite (γ, grey), banded ferrite (δ, white) and chromium nitrides (Cr2N, black) is formed. This observation is consistent with our previous results [25,26]. It is clear from SEM morphology in Fig. 1(b) that there are two kinds of nitrides in HNASS. One is dot, which is mainly distributed along the austenitic boundaries or at the interfaces between γ and δ phase (marked by dash arrow). Another is dendritic, which is generally found within the ferritic band (marked by solid arrow).

Fig. 1.

(a) OM and (b) SEM micrograph of as-received sample.


The OM microstructures after solution treatment at 1200°C for all processing times are shown in Fig. 2. They show a gradually changing microstructure from the surface to the bulk. It can be seen in Fig. 2(a) that the microstructure in the bulk consists of the white δ phase and the grey γ phase, but no nitrides. This suggests that all nitrides have been removed from the bulk through solution treatment for 30min. There is a white, bright surface layer, but the detailed microstructure cannot be revealed by OM. A subsurface layer between the surface and the bulk is dark grey, with a microstructure similar to the γ phase in the bulk. With increasing solution treatment time, both the white surface layer and the subsurface layer become thicker (Fig. 2(b)–(e)).

Fig. 2.

OM micrographs of samples treated by solution treatment at 1200°C for (a) 30min, (b) 60, (c) 90min, (d) 120min and (e) 240min.


The XRD analysis in Fig. 3(a) reveals that after solution treatment, a precipitates-free structure on the surface layer is obtained, consisting of only δ (α) phase with equiaxed grains (Fig. 3(b)) from Z direction (vertical to the rolling direction in Fig. 3(a)). A fully γ phase is observed in the subsurface layer, which also can be proved by microstructural features such as austenitic grains with annealing twins seen in Fig. 3(c). Ferrite was not detected by means of the XRD.

Fig. 3.

XRD patterns and microstructures of different layers in sample by solution treated at 1200°C for 120min.


Based on the above observations, a multi-layer structure is formed in HNASS during solution treatment. This may be caused by a complex process of localized phase transformation and N diffusion. Clearly, a homogeneous ferritic layer obtained on the surface has two main consequences. First, the high solution treatment temperature with a very low N potential environment leads to the N loss from the sample surface. This is due to the low activity of N in the vacuum atmosphere compared to the bulk, which drives a N concentration gradient from the bulk to the surface. It has been postulated that the γ phase close to the δ phase becomes the source of N since the solubility of N is higher in the γ phase than in the δ phase, while the δ phase only acts as a pathway for N diffusion [27,28]. Therefore, N atoms start to diffuse from the γ phase, pass through the δ phase to develop molecular N by the reaction of [N]+[N]N2 on the surface, and enter into the outside atmosphere to balance the N potential. Secondly, for solution treatment under a negative pressure, the outward diffusion of N can be accelerated. With decreasing N content in the γ phase, the transformation of γδ occurs in order to maintain the necessary surface nitrogen content for austenite stability. Therefore, a fully ferritic layer is finally obtained due to the N loss from the surface of HNASS during solution treatment. A similar phenomenon has been found in the aging process, in that the nitrogen-depleted γ phase around nitrides, can transform to the δ phase, due to the precipitated nitrides consuming the N content of the γ phase [27,29].

The influence of solution treatment time on layer thickness is shown in Fig. 4(a). It is noted that the thickness of the surface layer increases with increasing solution treatment time. This because, at this high temperature, more nitrogen atoms can successively diffuse from the γ phase at or near the surface and rapidly go through the ferritic region towards the surface as the time increases, resulting in more γ phase transforming into δ phase. Thus, the fully ferritic surface layer grows with time.

Fig. 4.

(a) Relationship between layer thickness and time and (b) plot of experimental X2 in (a) vs time for 1200°C solution treatment.


According to Fick’s second law, the relationship between layer depth from the surface and N content is as shown in Eq. (1) as following:

where CB is the N content of the bulk, CS is the N content on the surface, X is the distance to the surface, D is the N diffusion coefficient, and t is the solution treatment time. In the present experiments on HNASS, the following assumptions are made based on the schematic phase diagram shown in Fig. 5(a):
  • (a)


  • (b)


  • (c)

    When a fully ferritic surface layer with a thickness of X is formed, CX =  Cδ =0.05wt.%, which is the maximum N content of the δ phase at 1200°C.

Fig. 5.

(a) Schematic phase diagram in Fe–Cr–Mn–Mo–N system [30] and (b) schematic nitrogen content profile in the multi-layer at 1200°C.


Thus, the relationship between t and X can be achieved in Eq. (2):

where K is the constant, D can be calculated using the following Eq. (3):
where D0 is the maximal diffusion coefficient, R is the universal gas constant, and Q is the activation energy for diffusion. They are the diffusion constants for N during solution treatment at 1200°C. This suggests that a plot of the square of the depth of surface layer (X2) vs solution treatment time (t) should be linear. Taking the data in Fig. 4(a), and plotting X2vs t is for the fully ferritic surface layer. Fig. 4(b), a linear plot is obtained, which is in a good agreement with the calculations using Eq. (2).

The 100% γ phase in the subsurface layer corresponding for solution treatment at 1200°C is shown in Figs. 2 and 3, which is different from the (γ+δ) phases in the bulk (Figs. 1 and 2). This can be explained by the inward diffusion of N atoms. It is known that the ferritic surface layer is formed by transforming the γ into the δ phase through N diffusing outwards. With the δ phase the surface layer formation, N is partitioned into the γ phase, due to the much lower solubility for N in the δ phase than in the γ phase [30]. As a result, N will diffuse inward to the subsurface when the surface layer starts to grow. The N diffusion in the inward direction has been found at temperatures above 1000°C during welding in a duplex stainless steels by Hosseini and Meyer et al. [27,31]. They consider that this hinders nitride formation because N can diffuse into the γ phase easily due to the high temperature. This may be the main reason why nitrides could not be found within the ferritic surface layer after high temperature solution treatment in this investigation (Figs. 2 and 3).

Fig. 5(b) shows that there are two kinds of N content gradients during solution treatment. One is formed in the ferritic surface layer with the concentration difference of Cδ and Cs (Cs=0). Another is obtained in the subsurface layer. With increasing N content caused by diffusion from the interface to the bulk, the γ phase in the subsurface close to the interface becomes stable and the δ phase transforms into the γ phase simultaneously. It is postulated that the N gradient from Cγ to CB has a significant effect on the ferritic transformation [32]. Consequently, the δ phase fraction decreases and a fully austenitic subsurface layer is formed. Such a phenomenon in the high nitrogen stainless steels has been reported in the literature. Wang et al. have found that a single γ phase is formed at the high N content region, conversely, a δ phase precipitates at the low nitrogen content region of a Fe-18Cr-15Mn-2.5Mo-N steel [30]. For a 15-5PH martensitic stainless steel, the transformation of δ to γ phase is observed on the surface due to the increase of nitrogen content by plasma nitridation, and the thickness of this layer increases with solution treatment temperature [33].

A typical EDS result for the N concentration within the surface layer, and its adjacent subsurface layer, is illustrated in Fig. 6. The red number in the graph shows the weight percentage value of N measured in the black square spot. This data confirms the existence of two N content gradients within the two layers, as illustrated schematically in Fig. 5(b). The closer to the interface, the higher is the N content in either the surface, or the subsurface, layer. The N concentration in the fully austenitic subsurface layer is higher than that in single ferritic surface layer. Additionally, in a particular γ phase area, which is close to the interface between the γ and the δ phases, is significantly higher than that of the γ phase area in the interior subsurface layer. This confirms the occurrence of nitrogen partitioning from the δ phase to the surrounding γ phase areas during the δ-to-γ phase transformation. Although the EDS technique for the N or C distribution analysis can not be used to measure the N content quantitatively, it has been applied by some researchers to qualitatively compare the level of N or C contents [27,34]. Fig. 2 shows that the surface layer has a flat interface line with the subsurface layer after solution treatment. However, a closer examination of the layer boundary in Fig. 6 shows a concavity of the arc formed at the layer boundary line, demonstrating that N diffusion occurs across the layer boundary. This is related to the higher diffusivity of N atoms in the δ phase than in the γ phase [35,36].

Fig. 6.

SEM micrograph with EDS point analysis of different layers for sample solution treated at 1200℃ for 120min.


Owing to the small thickness of the surface layer that is produced, it is impossible to obtain microhardness values for short time solution treatments. Therefore, the microhardness profiles were only determined for samples for 120min (see Fig. 2(d) for location of hardness profiles). It shows a gradual hardening from the surface to the subsurface layer, then to the bulk (Fig. 7). A surface layer with a microhardness 100HV lower than that of the bulk was obtained. N significantly increases the strength and hardness in stainless steel by solid solution strengthening [23,24]. It is not surprising that the lowest hardness value is obtained in the fully ferritic surface layer: see Eq. (4) which describes the correlation between N content and microhardness (HV is the microhardness value, N% is the nitrogen content in weight percentage) [37]. This low hardness is caused by the low solubility of N in the δ phase (Fig. 6), which decreases the effect of solid solution strengthening.

Fig. 7.

Microhardness distribution along dash lines in Fig. 2(d) after solution treatment at 1200°C for 120min.


The hardness values in the subsurface layer with 100% γ phase show an increase relative to the surface layer, but are lower than those measured in the bulk. This is uncommon for high N content stainless steel, as shown in Fig. 5(b). This can be explained at least in part by the results for grain size.

Fig. 8 shows a grain growth in both the surface layer and the subsurface layer, compared with the bulk, after solution treatment for 120min. The largest average grain size (an average value of all grains shown in the locations in Fig. 2(d) using image-pro plus 6.0 software) of 120μm is obtained in the austenitic subsurface layer, where exhibits a considerable variation ranging from 35 to 260μm. It should be pointed out that grain size is anticipated to play an important hardening role since there are no nitrides in the subsurface layer. As a result, the hardness in this area is governed by both the content of N and the grain size of γ phase. A high N content in the subsurface layer can increase the hardness. However, some abnormally coarse austenitic grains give rise to a low microhardness value.

Fig. 8.

Grain size in areas marked by rectangle patterns in Fig. 2(d) after solution treatment at 1200°C for 120min.


Fig. 4(a) shows that the thickness of the surface layer increases from 78 to 179μm, while the subsurface layer depth increases from 206 to 736μm, when the solution treatment time increased from 30 to 240min. This suggests that the subsurface layer thickness is larger than that of the surface layer for all processing times. This thickness difference seems to increase with increasing solution time. There is about a 560μm difference in the thickness between the surface and the subsurface layer after 240min. As previously reported [38], crystal defects such as dislocations and vacancies can act as trap sites for N, decreasing the N diffusion depth. The trap sites density for N atom in the subsurface layer is lower than the surface layer because of the lower level of crystal defects within coarser austenitic grains. Thus, the larger the austenitic grains size, the thicker the subsurface layer. Also, it should be noted that, a larger N content difference between Cγ and CB obtained within the subsurface layer compared to it between Cδ and CS (about 0.05%) in the surface layer, as shown in Fig. 5(b), may promote the phase transformation from γ to δ phases, resulting in a greatly increased depth of the subsurface layer with solution time.


A multi-layer structure consisting of a ferritic surface layer and an austenitic subsurface layer, was obtained in a 19.44% Mn, 18.93% Cr, 1.63% Mo and 0.65% N stainless steel by solution treatment at 1200°C for holding times ranging from 30 to 240min under a negative pressure vacuum. There is no evidence of nitride precipitates in the microstructure of the multi-layer. The formation of a multi-layer structure is attributed to the transformation between austenite and ferrite resulting from nitrogen diffusion at high temperatures. Both the subsurface layer and the surface layer increase in thickness with increasing solution treatment time. However, the subsurface layer (206–736μm) is always thicker than the surface layer (78–179μm). Coarse austenite with a high nitrogen content formed in the subsurface layer exhibits a hardness higher than ferritic surface layer, but lower than that of the bulk.

Conflicts of interest

The author declares no conflicts of interest.

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