Chapter 4: Diffraction‐line shifts and broadenings in continuously and discontinuously
4.4 Discussion
Fe‐4.5 at.%Cr thin foil after 135 min of nitriding at 500 °C (Fig. 4.7) can be compared with the 211 lattice‐plane spacing changes determined from the measured peak positions in Table 4.1.
A full comparison of measured and predicted lattice‐parameter changes (diffraction‐peak shifts) requires recognizing the following contributions to the lattice‐parameter change:
i): The ferrite lattice contracts by the depletion of Cr from the solid solution during CrN precipi‐
tation: the {211} lattice‐plane spacing changes ∆ in Table 4.1 (due to elastic misfit accom‐
modation; see above) have been given relative to pure α‐Fe.
The following (minor) contributions occur, but have not been quantified:
ii): As the current specimens were not denitrided, an additional dilatation due to (equilibrium and excess) N dissolved in the ferrite matrix (for details about excess N see e.g. Ref. [10]) can occur for the specimens (as observed after cooling to room temperature). In Table 4.1, the ef‐
fect on the {211} lattice‐plane spacing by dissolution of the equilibrium amount of N in ferrite (Δd dissolved N) has been given. For the Fe‐2.0 at.%Cr specimen, the measured peak shift (Δdmeasured) is a little smaller than would be expected taking into account the dilatation both by CrN precipitation and by dissolution of the equilibrium amount of N in ferrite (Δd total). This sug‐
gests that due to the relatively high nitriding temperature of 580 °C, already after 1 h of nitrid‐
ing some relaxation has occurred (i.e. the precipitation‐induced misfit is only partially elastically accommodated). On the other hand, the Fe‐4.5 at.%Cr specimen was nitrided at the relatively low temperature of 500 °C, and thus, less relaxation is expected. Indeed, the measured peak shift (Δdmeasured) for this specimen is (even) larger than would be expected taking into account the dilatation both by CrN precipitation and by dissolution of the equilibrium amount of N in ferrite (Δd total). This additional peak shift is likely due to the presence of excess N as indicated in [10].
iii): Precipitation of α”‐Fe16N2 from dissolved N will occur from N‐supersaturated ferrite at room temperature (see e.g. Ref. [98], [155], [156]). This additional precipitation process at room tem‐
perature will also lead to additional line shifts and broadening changes [118], [131], [132]. How‐
ever, the misfit between ferrite and α”‐Fe16N2 is considerably smaller ( 0.0300 (lattice pa‐
rameters of α”‐Fe16N2 from Ref. [156])) than the misfit between ferrite and CrN ( 0.1464, see Chapter 2) and thus its influence on the measured lattice‐parameter change is minimal.
iv): An additional misfit between CrN and ferrite results from the differences in thermal expan‐
sion coefficient of CrN precipitates and ferrite matrix ( 0.0065 for cooling from 580 °C to
25 °C). Also this contribution to the lattice‐parameter change can be neglected compared to the magnitude of the elastically accommodated volume misfit present after full initial precipitation of CrN in the alloys (see above).
In addition to the observed peak shifts due to elastic precipitate/matrix misfit accommodation, the asymmetric tail of the main reflection (towards higher diffraction angles) represents the severely distorted nature of the regions in the vicinity of the precipitates (Chapter 2 and Refs. [13], [78]).
4.4.2 Continuous Coarsening
The nitride precipitates in the Fe‐2.0 at.%Cr alloy upon annealing show considerable growth in length from the initially, fully coherent state to the incoherent state, as evidenced by TEM (cf. Fig. 4.5a) and b)). The coarsening process is accompanied by a loss of elastic accommodation of the precipitate/matrix (volume) misfit and thereby the overall, long‐range, elastic lattice ex‐
pansion decreases (Fig. 4.3). This is accompanied by a “back”‐shifting of the CP ferrite main peak towards the unstrained position as revealed in Fig. 4.2. Cr stays depleted from the ferrite matrix and thus does not contribute further to the peak shift as a function of aging time. Other (subor‐
dinate) causes of peak shift as a function of aging time are the following (cf. discussion in sec‐
tion 4.4.1):
i): During relaxation of the precipitate/matrix volume misfit and the coarsening of the nitride platelets (reduction in interfacial area) the amount of excess N in the alloy decreases (desorption) and thus its (already minor) contribution to the peak shift decreases with increas‐
ing annealing time.
ii): The thus decreased amount of dissolved N with increased annealing time can lead to a de‐
creased development of α”‐Fe16N2 precipitates at room temperature. Also this (already minor) contribution to the peak shift decreases with increasing annealing time.
iii): The minor contribution of the thermal misfit does not change. Its corresponding peak shift thus remains minimal (see section 4.4.1).
In an end stage of coarsening and apart from the immediate vicinity of the precipitates, the mis‐
fit will be (nearly) fully relaxed and thus the ferrite matrix shows a lattice spacing close to that of pure ferrite. Therefore, the emerging CC‐peak develops at the approximate position of pure
α‐Fe. However, the heterogeneity of the relaxation process and its locally varying degree of
completion involve the presence of a pronounced amount of microstrain and thus the occur‐
rence of an appreciable amount of diffraction‐line broadening as exhibited by the CC‐peak. Ob‐
viously, upon increasing (aging) time, the strain variation decreases leading to a sharpening of the corresponding reflection (Fig. 4.2).
The developing incoherency of the precipitate/matrix interface randomizes the phase relation of X‐rays scattered by the matrix and those by the precipitates and thus incoherent diffraction of matrix and precipitates now occurs [58], [59]. Indeed, the TEM SADP in Fig. 4.5b) now reveals separate diffraction spots from matrix and precipitates. In the XRD patterns, individual precipi‐
tate reflections cannot be observed due to the low volume fraction.
4.4.3 Discontinuous Coarsening
If the alloy contains sufficient alloying element and if the misfit and the amount of precipi‐
tate/matrix interfacial area is high enough, then discontinuous coarsening is favored over the (usual) continuous coarsening. Indeed DC governs the coarsening process in the nitrided Fe‐4.5 at.%Cr alloy (see Fig. 4.10). The overall DC reaction in the nitrided Fe‐4.5 at.%Cr specimen can be expressed as (see Refs. [21], [150], [152]):
′ ′ fine CP → coarse DC lamellae
where, ′ denotes the strained, N‐supersaturated ferrite matrix, ′ the strained, coherent CrN precipitates, the relaxed ferrite lamellae without supersaturation and the relaxed, coarse incoherent CrN precipitate lamellae. The energy difference of the continuously precipi‐
tated and discontinuously coarsened states, i.e. the driving force for DC, for the present speci‐
mens is comprised of three contributions [21]; reduction of the interfacial area (reduction of interfacial energy), the loss of ferrite matrix (nitrogen) supersaturation (i.e. loss of excess N;
reduction in chemical energy), and the relaxation of the long‐range strain‐fields due to the co‐
herent to incoherent transition (release of strain energy).
Actually, in specimens where DC can occur, both continuous and discontinuous coarsening take place simultaneously but with different reaction rates [19], [22], [23], [25]: The DC reaction can be rate controlled by diffusion of the atoms along the mobile grain boundary that sweeps through the specimen and in the wake of which the cellular microstructure develops, whereas the CC reaction is governed by volume diffusion processes, necessary for the growth and disso‐
lution of precipitates. Consequently, the rate of the DC reaction is (much) higher than that of the CC reaction. Indeed, whereas an end stage of CC at 580 °C is reached after hundreds of hours, a final state of DC at 500 °C is reached after hundreds of minutes (cf. Fig. 4.2 and Fig. 4.7).
Therefore, in addition to the development of relatively slowly relaxing and thus only partly re‐
laxed regions surrounding the continuously coarsened precipitates, in advance of the reaction front of the DC reaction, fully relaxed matrix regions (ferrite lamellae) occur behind the DC reac‐
tion front. Thus, in a specimen that undergoes DC, in principle two types of ferrite regions con‐
tribute to the second diffraction peak at higher diffraction angle (cf. section 4.3.2): i) the contin‐
uously coarsened regions (decreasing in volume upon continued aging) corresponding with a broad peak (cf. section 4.4.2), and ii) the fully stress‐relieved ferrite in the discontinuously coarsened part (increasing in volume upon aging) corresponding with a sharp peak. At an end stage of aging (i.e. after 300 min of aging at 500 °C) the contribution of continuously coarsened regions has become minimal (cf. Fig. 4.10). Then, the emerging second peak is dominated by the discontinuously coarsened ferrite and is therefore sharp. Also, due to the presence of a clear boundary (see Refs. [23], [25]) between the ferrite in the continuously coarsened region and the ferrite in the discontinuously coarsened regions, both will always diffract incoherently. For the DC microstructure, i.e. with incoherently diffracting precipitates of relatively large dimensions, separate precipitate reflections of CrN and (DC‐) ferrite matrix occur (see Fig. 4.11).