11.2 Fundamental Microstructure Defects, Their Activity and Configurations in Austenitic Steels
11.2.3 Arrangement of the Stacking Faults in Austenite
11.2.3.3 Coexistence of Different Stacking Fault Arrangements
(11.14), see Sect.11.4. The effect of perfect dislocations on the line broadening is described for all phases using (11.15).
The Rietveld refinement of the diffraction pattern of the fine-grained sample X2CrMnNi16-6-6 compressed to 15%, cf. previous simulation using the DIFFaX model, revealed an intrinsic stacking fault probability inε-martensite of about 10%, which agrees well with the value of 1−Phcp obtained from the DIFFaX model.
The volume faction of ε-martensite correlates directly withPhcp. In the austenite, the intrinsic and extrinsic SF probabilities calculated using the DIFFaX model and the Warren model can be compared directly as well. The parameterPtwincan hardly be compared with the microstructure parameters obtained from the Warren model, because the twinning does not cause any shift of the diffraction lines but only a line broadening, see (11.11). However, as it can be seen from (11.11), the densities of intrinsic (isolated) SFs (α) and extrinsic SFs (β) terminating the twins cannot be determined simultaneously solely from the line broadening, but only in conjunction with the analysis of the line shift, or from the line asymmetry [30]. However, as the density of the twin boundaries in the steels under study is much lower than the density of intrinsic SFs, it cannot be determined reliably using the combined analysis of the line shift and line broadening. The analysis of the line asymmetry fails as well, as the asymmetry of the diffraction lines is not very pronounced [33].
Another result of the Rietveld refinement was that the distance between the adjacent{111}f cclattice planes within the faulted stacking sequences was smaller than the interplanar spacing within the regular stacking sequences. Consequently, the ε-martensite had the c/a ratio of 1.62, which is slightly below the value of c/a = 2√
2/3 = 1.633 that corresponds to a pseudo-cubic hcpcrystal structure witha=af cc/√
2 andc=2d111f cc=2af cc/√ 3.
Fig. 11.10 EBSD mapping of deformation bands in compressed X5CrMnNi16-6-9 steel, austenite in grey,ε-martensite in yellow, twin boundaries in red, unindexed pixels in black. Adopted from [34]
Fig. 11.11 HRTEM image of a deformation band, where atomic arrangements of both twins and ε-martensite are found coexisting in close vicinity
the HRTEM micrographs. Figure11.11illustrates the different stacking sequences of the former{111}f cclattice planes, which can be described as an irregular ordering of stacking faults. Extended sequences AB⊥AB⊥AB⊥AB of the lattice planes{111}f cc
are interpreted ashcpε-martensite, extended sequences ABCAB⊥A⊥C⊥B⊥A⊥C⊥ as twins. The irregular sequences of SFs produce streaks in the reciprocal lattice of austenite along the faulted 111f cc direction, which include the reciprocal lattice points ofε-martensite (see FFT2 in Figs.11.11and11.12b). The streaks stem from superimposed truncation rods produced by thin slabs of faulted austenite [28,35, 36]. The presence of extended slabs ofε-martensite or extended twins leads to the
Fig. 11.12 TEM characterization of deformation bands in PM X5CrMnNi16-6-9 steel.aBright- field image of parallel deformation bands,bcorresponding SAED, recorded with a 200 nm aperture, cindexing of spots assigned to austenite matrix, twin andε-martensite. Adopted from [34]
fragmentation of the streaks and to the formation of separated diffraction spots that correspond toε-martensite or to twins in austenite (see FFT2 and FFT3 in Fig.11.11).
These microstructure features are interpreted in a very similar way by all diffraction methods, i.e., EBSD (Fig.11.10), selected area electron diffraction (SAED) in TEM (Fig.11.12) and XRD (Fig.11.9).
The crystallographic explanation of these diffraction phenomena is based on the stacking fault models from Figs. 11.6 and11.7, and on the orientation relation- ships between austenite,ε-martensite and twinned austenite that can be described as (0001)hcp||
111¯
f cc||
111¯
twinand 0110¯
hcp|| [110]f cc||
101¯
twin, cf. Fig.11.12b, c. The deformation bands visible in the TEM micrograph (Fig. 11.12a) contain remainders of the austenite matrix, uncorrelated SFs and areas, which are inter- preted asε-martensite or as twins in austenite. In the SAED pattern (Fig.11.12b), the remainders of the austenite matrix, theε-martensite areas and the twinned austen- ite produce intense diffraction spots, the uncorrelated SFs streaks along the
111¯
f cc
reciprocal space direction.
From the thermodynamic point of view, the simultaneous occurrence of twinned austenite andε-martensite can be discussed in terms of SFE and interface energy [22].
In the case of a low thermodynamic driving force (Gγ→ε∼=0, cf. Sect.11.2.3.1), all arrangements of stacking faults corresponding to isolated intrinsic and extrinsic SFs,ε-martensite and twins result in the same overall energy. Therefore, the transi- tion between the transformation-induced plasticity (TRIP) and the twinning-induced plasticity (TWIP) is continuous. Assuming that there are no other obstacles for the phase transition, the TRIP effect and the generation of a local hexagonal stack- ing sequence are preferred forGγ→ε <0, while forGγ→ε >0, the twinning dominates.
An important consequence of this microstructure model is an improved atomistic description of theε-martensite formation. Although SFs were always regarded as nucleation precursor [21] of theε-martensite formation, theε-martensite has been treated as a metastable martensitic phase occurring along theγ →ε→αtransfor- mation path [37,38] and as a product of bulk martensitic transformation of austenite
[39]. The current understanding ofε-martensite, which evolves during plastic defor- mation of austenite, is that it is rather a heavily faulted austenite than a distinct phase.
Using electron microscopy, theε-martensite was found to appear within deforma- tion bands [31,40], where a high density of stacking faults is present. Consequently, the deformations bands and hence theε-martensite form during the plastic defor- mation of austenite, when a low SFE facilitates the SF formation as a competing mechanism to the dislocation mobility. On the other hand, the deformation bands and theε-martensite act as repositories of dissociated dislocations, which are bound to their slip planes. Hence, the dynamic recovery is inhibited, as the annihilation would be possible only for perfect dislocations, which are not stable, as their dissociation is facilitated by the low SFE.