Unusual Iron Nitride Formation Upon Nitriding Fe-Si Alloy

Cast rods of Fe-2 at. pct Si and Fe-4.5 at. pct Si alloys with 10 mm diameter and 100 mm length were produced by induction melting of elemental granules of iron (purity: 99.98 wt pct) and pieces of silicon (purity: 99.999 wt pct) under a protective argon atmosphere. The amounts of silicon and light element impurities in the produced casts were analyzed. The silicon content was determined by inductively coupled plasma-optical emission spectroscopy (ICP-OES), the nitrogen and oxygen contents were measured by applying carrier gas hot extraction. Sulfur and carbon contents were determined by the combustion method. The thus obtained results of the chemical analysis are shown in Table I.

A cast rod of Fe-2 at. pct Si alloy was cut into pieces, which were cold rolled to produce sheets having a thickness of approximately 1 mm. From these sheets, rectangular specimens with lateral dimensions of 10 mm × 15 mm were cut. The surfaces of the specimens were ground and polished down to 1-μm diamond finish. These prepared specimens were placed into a quartz tube, which was evacuated, filled with argon gas, encapsulated and subsequently annealed at 750 °C (1023 K) for 2 hours to recrystallize the specimens.[44] Afterwards, the specimens were ultrasonically cleaned in ethanol and subjected to the nitriding treatment.

Attempts to cold roll pieces of the Fe-4.5 at. pct Si alloy were not successful because of cracking of the material while rolling, which is attributed to its extreme brittleness. Therefore, approximately 1-mm-thick specimens were cut directly from the cast rod and prepared (see above) for nitriding. As for the Fe-2 at. pct Si specimens, the Fe-4.5 at. pct Si specimens were subjected to annealing at 750 °C (1023 K) for 2 hours in argon to relieve any deformation induced in the specimens during cutting and also to remove any casting-induced elemental segregations.

The specimens were nitrided in a laboratory-scale, vertical quartz tube furnace having an inner diameter of 28 mm, equipped with a facility to hang the specimen in the middle of the furnace and a water-filled flask to quench the specimens after completing the nitriding treatment. The furnace temperature was controlled within ± 1 K variation, and the flow rates of ammonia and hydrogen gases were controlled by MKS mass flow controllers. For the nitriding treatment, high-purity ammonia (99.99 pct) and hydrogen (99.999 pct) gases were used. Before starting, the furnace was purged with nitrogen gas to remove all the oxygen and other adsorbed gases from the furnace walls. After setting the required flow rates of ammonia (195 ml/min) and hydrogen (305 ml/min) and allowing the gas mixture to reach a steady-state condition in the furnace, the specimen was hung up in the uniform temperature zone of the furnace. The total gas flow rate of 500 ml/min resulted in a linear gas velocity of 13.5 mm/s, at room temperature, for the quartz tube furnace used. At this linear gas velocity, a change of the chemical composition of the gas mixture due to dissociation of ammonia can be neglected.[45] These nitriding conditions result in a nitriding potential, r_N, of 0.82 atm^−1/2.[46] For the applied temperature of 550 °C (823 K) and r_N, only γ′ iron nitride formation is expected upon nitriding pure iron according to the Lehrer diagram.[47] After completing the nitriding treatment, the specimens were quenched in a water-filled flask. To minimize the risk of specimen surface oxidation during water quenching, before quenching the water in the quenching flask was purged with nitrogen gas for about 20 min to remove/minimize the amount of oxygen dissolved in the water. To trace the evolution of the compound layer with time, specimens of both alloys were nitrided for three different times (15 min, 2 hours and 20 hours) at 550 °C (823 K) and r_N = 0.82 atm^−1/2.

X-ray diffractograms recorded from the surfaces of the un-nitrided and nitrided [for different times at 550 °C (823 K) employing a nitriding potential of 0.82 atm^−1/2] Fe-2 at. pct Si and Fe-4.5 at. pct Si alloy specimens are shown in Figures 1 and 2, respectively. Evidently, distinct formation of γ′ takes place upon nitriding; the amount of γ′ increases with nitriding time. Although ε iron nitride formation was not expected for pure iron under the employed nitriding conditions (see the so-called Lehrer diagram[47]), some ε nitride developed as well. EBSD investigations confirmed the occurrence of ε (see Figure 7 in Section III–C).

Light optical micrographs recorded from the cross sections of Fe-2 at. pct Si alloy specimens nitrided for different times are shown in Figure 3. After 15 min of nitriding, a thin γ′ layer had formed on the surface with extensions along the ferrite grain boundaries intersecting the specimen surface (Figure 3(a)). With increasing nitriding time (2 hours), the thickness of both the grain boundary and the surface γ′ layers increased, and protrusions of γ′ developed from both the grain boundary and surface γ′ layers (Figure 3(b)). After 20 hours of nitriding, nearly all the ferrite grain boundaries near the specimen surface were covered with the γ′ phase, and the γ′ protrusions had become more pronounced (Figure 3(c)). Similar observations were made from cross sections of the nitrided Fe-4.5 at. pct Si alloy specimens (Figure 4). In some ferrite grains a “teeth-like,”/“comb-like” arrangement of the protrusions from the γ′ layer is observed (Figure 5(a)), whereas at other locations a compact layer morphology is maintained (Figure 5(b)).

EBSD was performed on the cross sections of the nitrided specimens. Phase and orientation maps recorded from the 20-hours nitrided Fe-2 at. pct Si alloy specimen are shown in Figures 6(a) and (b), respectively. It follows that the grain boundary γ′ phase, growing from any location at the original grain boundary, has different orientations at both sides of the original ferrite grain boundary. Similar observations were made for the grain boundary γ′ phase in the 20-hours nitrided Fe-4.5 at. pct Si specimen (cf. Figure 7b and d). Furthermore, for this specimen, each “tooth” of the “teeth-like” γ′ growing into a single ferrite grain exhibits the same two (see what follows) orientations with respect to that of the ferrite grain (Figure 7(a) and (c)). Each “tooth” of the “teeth-like γ′” is composed of two γ′ grains separated by a Σ3 twin boundary, as determined by the EBSD analysis (Figure 8).

According to the procedure described in Section II–C–5, the angles between 〈100〉 directions of ferrite and 〈100〉 directions of γ′ were determined. The thus obtained results for the Fe-2 at. pct Si and Fe-4.5 at. pct Si alloy specimens nitrided for 2 and 20 hours are shown in Figures 9 and 10, respectively. From comparison of these experimentally determined sets of angular values with those calculated theoretically for the different commonly observed ORs for FCC/BCC systems (see Table II), it follows that a (near) N–W OR prevails for the ferrite/γ′ system.

At short nitriding times (2 hours), in both Fe-2 at. pct Si and Fe-4.5 at. pct Si alloys for some γ′ precipitates, occurrence of a (near) Bain OR was observed (Figures 9(a) and 10(a): see the angular data sets in the dashed rectangles and compare with Table II).

The elemental (Fe, Si, N) concentration-depth profiles, as determined by EPMA, are shown in Figures 11 and 12 for the 2- and 20-hours nitrided specimens of Fe-2 at. pct Si and Fe-4.5 at. pct Si alloys, respectively, superimposed on the corresponding light optical micrographs. It follows that no partitioning of Si had occurred: the N content of the γ′ phase is higher than the N content expected for pure γ′-Fe₄N (cf. dashed, horizontal black lines in Figures 11 and 12), which hints at the presence of silicon nitride precipitates within the γ′. Assuming that all Si is present as stochiometric silicon nitride, Si₃N₄, within the γ′-Fe₄N phase (as fine precipitates), the N content of such a γ′-Fe₄N + Si₃N₄ mixture can be calculated using the following Eq. [1]:where “at. pct Si” and “at. pct Fe” denote the atomic percentages of Si and Fe in the un-nitrided specimen. The thus calculated theoretical N content together with the N contents determined experimentally for the γ′ phase regions, measured by EPMA, is shown in Table III.¹ The theoretically predicted and the measured N contents agree very well (Table III) for the 2- and 20-hours nitrided specimens, indicating that all Si originally present in dissolved state in the ferrite matrix has been incorporated as Si₃N₄ within a γ′-phase matrix. The N content of the developed γ′-phase region in the 15-min nitrided Fe-2 at. pct Si alloy (Table III) is only about 20 at. pct, which is as expected for a single-phase γ′ region. Apparently γ′ forms initially with all Si from the original ferrite matrix dissolved in it, and later this dissolved Si precipitates out as silicon nitride, i.e., γ′ has a negligible solubility for Si.

The TEM BF image and corresponding SADP recorded from the same location are shown in Figures 13(a) and (b), respectively. The BF image shows the presence of two kinds of precipitates in the γ′ matrix: a thick irregular or nearly rectangular blocky type (shown with a white arrow in Figure 13(a)) and a platelet type (shown with a black arrow in Figure 13(a)). A DF image formed by the diffuse scattered intensity around the transmitted 000 spot (location of the objective aperture is shown with the white dotted circle in Figure 13(b)) reveals the precipitates, observed in the BF image, as bright objects against a darker background (Figure 13(c)). This observation is compatible with an amorphous nature of the observed precipitates. EELS elemental analysis demonstrated the enrichment of both Si and N at the precipitate locations (Figure 14). On this basis, it is concluded that the precipitates are amorphous silicon nitride, as also observed for silicon nitrides developing in a ferrite matrix.[16,36] It can be speculated that the platelet-type silicon nitride has possibly developed at defects as planar faults in the γ’, whereas the blocky type of silicon nitride could have developed by homogeneous nucleation.

The elemental concentration-depth profiles measured on 2 and 20-hours nitrided Fe-2 at. pct Si alloy specimens show only a very small N content in the ferrite matrix underneath the γ′ phase region. It is therefore concluded that no silicon nitride has precipitated in the diffusion zone (Figure 11: N content of ferrite is nearly zero; if all Si had precipitated in the ferrite matrix as Si₃N₄, a minimum N content of 2.6 at. pct would be expected for the ferrite matrix). Silicon nitride precipitates only very slowly in the nitrided ferrite matrix of Fe-Si alloys.[26] This even induces the occurrence of ideally weak nitriding kinetics.[16] If more Si is dissolved initially, the driving force for silicon-nitride precipitation is larger. Indeed, whereas in the case of Fe-4.5 at. pct Si alloy, after 2 hours of nitriding, the ferrite matrix contains a negligible amount of N (Figure 12(a): N content of ferrite is nearly zero), after 20 hours of nitriding a significant amount of N is detected in the ferrite matrix indicating that now the precipitation of silicon nitride has set in (Figure 12(b)).