Introduction

Duplex stainless steel (DSS) alloys, which consist of austenite (face-centred cubic) and ferrite (body-centred cubic) phases, exhibit excellent mechanical and corrosion resistance properties. Austenite-matrix DSS are used for pipework systems in nuclear power plants, whereas ferrite-matrix DSS have been more often employed in the oil and gas industries.[1] The toughness of DSS alloys can be reduced by age-hardening at temperatures between 280 and 500°C This phenomenon, known as “475°C embrittlement”, arises from spinodal decomposition and precipitation in the ferrite phase that impedes dislocation movement and encourages {112}<111> deformation twinning.[2] The effect is a ductile to brittle transition, like that in ferrite at low temperatures without ageing due to thermally activated dislocation behaviour.[3,4]

Deformation twinning can affect fracture by initiating cracks, providing preferential propagation paths along the twin/parent interface, or blocking or deflecting cleavage crack propagation. Ferrite cleavage cracks propagate on {112}, {110}, and {001},[2,4] but may deflect along deformation twin/parent grain interfaces and austenite-ferrite phase boundaries.[5,6,7] Ferrite’s deformation twinning may also be significant in environment-assisted fatigue of embrittled DSS.[8] Numerous fractographic analyses have investigated the role of twins in ferrite crack propagation,[3,9] but no high-resolution studies have investigated the cracking of the twin/parent grain interface.

On-axis transmission Kikuchi diffraction (TKD) within a scanning electron microscope (SEM) can provide detailed information, by analysis of the transmitted and diffracted electrons, about crystallographic orientations, grain boundaries, strain fields and lattice defects, including dislocation arrangements, at a spatial resolution of a few nanometres.[10] On-axis TKD requires a shorter exposure time for Kikuchi pattern collection, with less distortion than off-axis TKD, making pattern indexing more accurate.[11] However, like other TKD and electron backscatter methods, Kikuchi pattern degrades with plastic deformation, which affects indexing accuracy.[11,12]

In this study, on-axis TKD characterisation has been employed, with both bright- and dark-field imaging, to reveal local lattice distortions and dislocation arrangements at the interface between the parent grain and a deformation twin. This boundary is important in the nucleation and propagation of microcracks in the ferrite of age-hardened duplex stainless steels. Our work provides a novel perspective on the role of deformation twinning in the embrittlement of DSS.

Methodology

A duplex stainless steel sample (Zeron 100, a super-duplex stainless steel with a ferrite matrix) was aged at 475°C for 100 h to induce age-hardening. The 4 × 1 × 1.5 mm3 dog-bone sample was mirror polished and tested in tension to a strain of 22% at the material’s ultimate tensile stress of about 850 MPa (see Ref. [12] for more details). The sample was cut and sectioned using a slow-speed saw before the thickness was reduced to ~ 0.25 mm by slowly grinding the unpolished surface using 240-grit SiC paper. A ~ 3 mm diameter disc [Fig. 1(a)] specimen was punched out and thinned by twin-jet electropolishing.

Figure 1
figure 1

(a) SEM image of the age-hardened duplex stainless steel 3 mm disc after thinning using twin-jet electropolishing. The area around the electropolished hole in (b) is electron-transparent. (c) The Bruker eFlash detector is composed of a 34 mm × 25.5 mm phosphor screen and three 5.2 mm × 4.4 mm transmitted electron diodes. The central low-angle transmission electron diodes were positioned perpendicular to the SEM pole piece to collect bright-field (BF) images in the transmitted electron imaging mode. Two high-angle diodes were used for dark-field (DF) images. (d) In the transmission Kikuchi diffraction (TKD) mode, the phosphor screen was perpendicular to the pole piece axis to capture the on-axis transmission Kikuchi diffraction patterns originating from the sample. An on-axis TKD pattern was projected onto the phosphor screen.

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Electropolishing, a process that uses an electrochemical reaction to remove a thin layer of material, was employed to create a smooth, mirror-like finish on the sample surface. This process minimises the introduction of mechanical damage and deformation, which could potentially affect the subsequent microscopic observations. For electropolishing, we used perchloric acid (HClO4) in ethanol (C2H6O) for 10 min at − 30°C (liquid nitrogen coolant) and an electropolishing voltage of 14 V. The perforation, approximately 70 µm in diameter, was surrounded by an electron-transparent region that extended about 70 µm radially [Fig. 1(b)].

A Zeiss Merlin FEG-SEM with a Bruker eFlash detector was used to study the specimen. The detector’s three ARGUS diodes were used to acquire high-resolution dark-field (DF) and bright-field (BF) images at a spatial resolution of up to 1 nm in the transmission electron imaging mode [Fig. 1(c)]. BF imaging has the benefit of atomic number contrast and the capability to observe dislocations due to its sensitivity to lattice distortion.[13] ImageJ software was used to measure the distances in the images.

The Bruker eFlash detector CCD camera was placed under the specimen to collect transmitted Kikuchi patterns [Fig. 1(d)]. The 800 × 600-pixel patterns were collected at 3 nA/30 kV beam conditions, 42 ms exposure time per pattern, and 5 nm step size. The CCD was tilted by 1° and positioned 18 mm (DD) from the specimen. The working distance (WD) was 5 mm. The indexed TKD patterns were analysed using MTEX[14] to (i) identify grain boundaries, (ii) plot orientation maps and pole figures, (iii) calculate the local misorientation relative to the grain mean orientation, and (iv) calculate the parent-twin boundary disorientation. MTEX was also used for trace analysis of deformation twins and cracks.

Results and discussion

A deformation twin that terminated within a ferrite grain was observed. The DF image [Fig. 2(a)] clearly shows the twin, which has irregular interfaces with its parent grain, and there is a crack along one side of the twin. The DF image does not reveal the austenite/ferrite phase boundary. This is indicated by the dashed white line, derived from the BF observation [Fig. 2(b)]. Conversely, the BF image does not show the crack, indicated by the red dashed line traced from the DF image. The BF image also reveals the deformation twin due to its diffraction contrast (i.e. crystal orientation change) and shows local lattice distortions within the twin [white arrows in Fig. 2(a) and (b)] and dislocations, including traces of \(\left(0\overline{1 }1\right)\left[\overline{1 }11\right]\) and \(\left(112\right)\left[11\overline{1 }\right]\) slip planes, which differ by 1.5°. A similar example can be found in the Supplementary Information, Fig. S4.

Figure 2
figure 2

\((\overline{2 }11)[111]\) Deformation twin, terminated mid-grain, imaged using the DF diodes (a) and the BF diode (b). The crack trace from the DF image is drawn on the BF image using a red dashed line. A TKD orientation map was overlaid over the twin tip, as shown in detail in (c). (c) TKD orientation map for the \((\overline{2 }11)[111]\) deformation twin and the parent grain with a grain boundary, which has a 60° disorientation, highlighted in white. The crystal orientation of the parent grain is visualised with a black cube and the twin with a red cube. The IPF orientation colour scale is shown in (f). (d) Misorientations at the deformation twin and parent grain were calculated relative to their mean orientations. (e) High-resolution BF image of the twin tip with a white arrow pointing to the dislocation-free zone (DFZ). The field of view within the dashed box was used to calculate the stereographic projection pole figure for {112} in (f), where the twin trace and the trace normal are shown using dashed and continuous black lines, respectively. The twin habit plane (the plane intersecting the trace normal) was identified and marked with a black arrow. See the Supplementary Information, Table S1, for more details about the parent grain and the twin orientation.

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The irregularity of the twin/parent interface is due to the two-step growth mechanism of deformation twins in ferrite. The first step is thickening, facilitated by the interaction between the sessile and mobile twin partial dislocations at the coherent twin/parent boundary.[15,16] The second step involves dislocation movement along the direction of the twin shear,[17,18] which causes the twin/parent boundary to extend[19] with mixed coherent and non-coherent boundaries.[20] This intermittent growth mechanism accounts for the observed irregularity of the twin/parent boundary [Fig. 2(a)] and leads to a significant residual strain.[7,21] The BF image also reveals dislocation arrangements around the twin, resulting from plastic deformation to accommodate the twin’s growth.

The TKD map [Fig. 2(c)] shows the orientations of the parent grain and twin, and the (112) pole figure [Fig. 2(f)] confirms the characteristic orientation relationship between the parent grain and the twin tip. There is a significant misorientation between the twin tip and its more distant parts, which occurred after twin growth [Fig. 2(d)]. The cracked interface did not extend to the twin tip, and the change in orientation along the twin was attributed to the accommodation of the residual strains of the twin by the compliant and cracked thin foil. There are significant local misorientations near the parent–twin interfaces [Fig. 2(d)], particularly near the cracked boundary, which is consistent with the strains from the interaction of discontinuous cracks along the irregular interface. The BF image of the twin tip [Fig. 2(e)] shows a dislocation-free zone (DFZ) ahead of the twin tip, indicated by the white arrow, with a size of approximately 240 ± 18 nm. The dislocation density is high near each side of the DFZ and decreases away from the tip.

A second example is shown in Fig. 3(a), where a crack propagating along the parent–twin interface has deviated to cross the deformation twin, following the \((2\overline{1 }1)\) plane (see the Supplementary Information, Fig. S1 for more parent/twin analysis). This propagation behaviour is likely due to favourable local stress conditions and crystallographic orientations, similar to the first example, where there is a higher degree of local misorientation at the cracked twin/parent interface. The BF image [Fig. 3(b.i)] shows contrast from significant lattice distortions on one side of the crack. This is also apparent in the calculated local lattice misorientations [Fig. 3(b.ii)]. The misorientation and dislocation density decrease with increasing distance from the crack tip.

Figure 3
figure 3

\((12\overline{1 })[\overline{1 }11]\) Deformation twin, with a continuous crack along the parent-twin boundary before crossing into the twin, imaged using (a.i) DF and (a.ii) BF modes. (b.i) BF image of the crack tip and (b.ii) local misorientation around the crack. The curved lines, seen at the top left of the image in (b.i), are bend contours. These are formed due to a slight bending in the thin foil, affecting the BF image’s sharpness but not the TKD patterns. (c.i) BF image of the deformation twin and (c.ii) local misorientation at the deformation twin with the twin boundary highlighted with white dashed lines. (c.i) and (c.ii) have the same length scale as shown in (c.i). See the Supplementary Information, Table S1, for more details about the parent grain and the twin orientation.

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On exiting the twin (which had a thickness of 255 ± 21 nm), the crack continued for 1.1 µm into the parent grain along the \(\left(\overline{1 }2\overline{1 }\right)\) plane, before deviating as it alternated between the \((001)\) and \((010)\) planes (see the Supplementary Information, Figs. S2 and S3, for more details about trace analysis). This is because as the crack propagates further and influences the stress field itself, it transitions to the {100} planes that are more energetically favourable.[4] In addition, the (001) and (010) planes are energetically equivalent, and such local alternating behaviour is commonly observed when the crystallographically preferred planes are not aligned with the overall crack path driven by the overall stress state.[22,23]

While no discernible dislocation-free zone (DFZ) is observed at the crack tip in this example, a DFZ (approximately 133 ± 10 nm) is present ahead of the \((12\overline{1 })[\overline{1 }11]\) twin, as shown in Fig. 3(c). Although the DFZ, in this case, is not directly involved in the observed crack propagation, it provides insights into the behaviour of dislocations emitted from sources ahead of the twin. The tip stress field repels dislocations, reaching equilibrium when the repulsive force is balanced by lattice friction.[24] This balance forms a DFZ at the twin tip. The size of the DFZ observed at crack tips has been observed to vary with the applied stress intensity factor and lattice friction stress.[25] This suggests that future studies of the DFZ can provide information about the local stress conditions at the twin tip.

In age-hardened duplex stainless steel, the interface between the parent grain and a deformation twin is a critical region where lattice distortions and dislocation arrangements significantly impact the material’s fracture behaviour. Deformation twins grow in response to a macroscopically imposed strain, with twin growth leading to highly localised strain accumulation that increases with the deformation twin thickness.[7] Previous studies have postulated that this provides a sufficient driving force for concurrent twin nucleation[26] and cleavage crack nucleation[27] as local strain-relaxing mechanisms. Observations of “tongue-like” features on fracture surfaces have suggested that deformation twins nucleate ahead of a propagating crack to provide pre-cleavage microcracks along the twin/matrix interface that cause the crack to deviate locally.[4,9] In addition, post-test analysis[3] and modelling (e.g. crystal plasticity finite element analysis)[4,28] have led to the view that microcracks nucleate at intersecting deformation twins and slip bands before propagating along the deformation twinning boundary.

In essence, deformation twins in age-hardened duplex stainless steel create conditions favouring brittle fracture by inducing local stress concentrations. The ∑3 {60°} parent/twin boundary, with a twin {112} habit plane, has low energy;[29] thus, as in the current observations, fracture occurs to reduce the elastic strain energy caused by the residual stress field associated with the deformation twin growth. The twin growth mechanism introduces an irregular interface between the twin and parent grain; therefore, cleavage cracking on the twin habit plane is discontinuous. Interactions of the discontinuous and non-coplanar cracks lead to the local misorientations observed along the twin interface [Figs. 2(d), 3(c.ii)]. This suggests that the cracking of the twin/parent grain interface is intrinsically linked to twin growth and is driven by the relaxation of local stresses along the irregular twin/parent boundary.

Conclusions

In summary, this study has investigated deformation twins and cleavage cracking in the ferrite phase of age-hardened duplex stainless steel (DSS) alloy. Using on-axis transmission Kikuchi diffraction (TKD) in a scanning electron microscope (SEM), we have observed dislocation arrangements around the twin, the dislocation-free zone (DFZ) at the twin tip, and most importantly, lattice distortions at cracked and non-cracked twin/parent interfaces.

In the first example (Fig. 2), a crack propagates along the parent–twin interface, likely due to favourable local stress conditions and crystallographic orientations. However, in the second example (Fig. 3), the crack deviates to cross the deformation twin. This deviation is due to local variations in the stress state due to the deformation twinning and age-hardening process. We deduce that the elastic strains from intermittent twin growth encourage cleavage of the low-energy ∑3 twin/parent interface. This explains the inherent tendency for fracture along the twin/parent boundaries in ferrite.