Deformation mechanisms induced by nanoindentation tests on a metastable austenitic stainless steel: a fib/sim investigation

Deformation mechanisms induced by nanoindentation tests on a metastable austenitic stainless steel: a FIB/SIM investigation

I. Sapezanskaia^1,2, J.J. Roa^1,3*, G. Fargas^1,3, M. Turon-Viñas^1,3, T. Trifonov³, R. Kouitat Njiwa²,

A. Redjaïmia², A. Mateo^1
1 CIEFMA - Dept. of Materials Science and Engineering, Universitat Politècnica de Catalunya, EEBE - Campus Diagonal Besòs, 08019 Barcelona, Spain

² Institut Jean Lamour, UMR 7198 CNRS, Université de Lorraine, Parc de Saurupt, 54011, Nancy Cedex, France

³ Centre for Research in NanoEngineering, Universitat Politècnica de Catalunya, EEBE-Campus Diagonal Besòs, 08019 Barcelona, Spain

* Corresponding author, e-mail: [e-mail]

Abstract:

Metastable austenitic stainless steels are materials that can undergo austenite to martensite phase transformation when subjected to deformation and thus they represent a multiphase material with interesting mechanical properties. Different electron microscopy techniques are widely applied for the characterization of their deformation mechanisms at micrometric length scale. In doing so, Scanning Ion Microscopy (SIM) imaging, performed with a Focused Ion Beam (FIB), can be useful to evaluate microstructural features induced by different stress fields and, in certain cases, may substitute the conventional Transmission Electron Microscopy (TEM) technique.

In this work, nanoindentation experiments (both monotonic and cyclic) were carried out on AISI 301LN metastable steel in order to induce localized deformation of individual austenitic grains. The activated plastic deformation mechanisms were evaluated by using different advanced characterization techniques (Electron BackScattered Diffraction (EBSD) and TEM), but mainly by FIB/SIM. FIB/SIM 3D-tomography was also conducted to reconstruct the deformation structure under the residual imprint. These observations, surprisingly, showed the existence of a good correlation between SIM and TEM images, concerning phase transformation and plastic zone development.

Keywords: Metastable stainless steels, phase transformation, monotonic and cyclic indentation tests, transmission electron microscopy, electron backscattered diffraction, focused ion beam.

1. 
   Introduction
   

Austenitic stainless steels, due to their excellent properties, are widely employed in many industry sectors, such as transport, food or chemical, among others [¹]. They offer the advantage of being highly resistant to corrosion, as well as of exhibiting a good combination of strength and ductility. Metastable austenitic stainless steels are distinguished by the capability to undergo a partial phase transformation, from the relatively soft austenite to the harder martensite, when mechanically stressed. This phenomenon, known as TRIP (Transformation Induced Plasticity) effect [²], contributes to an even further increase in both strength and ductility [³]. Therefore, metastable steels combine a relatively easy plastic conformation in the initial annealed state with a high strength in the final service condition. Furthermore, the resulting high work hardening rate gives the metastable stainless steels the capacity of great energy absorption under impact, interesting for crash worthiness [⁴]. However, the major drawback is the difficulty to accurately predict their mechanical behavior, since it depends predominantly on their microstructure. In this regard, the understanding of the exact interrelation between microstructure and mechanical properties is crucial for a targeted tailoring of those steels, according to their respective application and requirements.

In this context, the study of the local mechanical properties, as well as of that of the main plastic deformation mechanisms, activated under different stress fields, becomes relevant. So, the complex deformation mechanisms in metastable austenitic stainless steels typically involve linear defects (dislocations), planar defects, such as stacking faults (SFs) and/or shear bands [⁵,⁶], mechanical twins, as well as the phase transformation phenomena (austenite to martensite) [⁷,⁸,⁹]. It is required to mention that during the last years plenty of works have been conducted in order to better understand the macroscopic phase transformation on TRIP steels by using advanced characterization techniques like in-situ X-ray diffraction [¹⁰,¹¹,¹²,¹³,¹⁴], TEM [Error: Reference source not found,¹⁵,¹⁶,¹⁷,¹⁸], among others. More recently, Electron Channeling Contrast Imaging (ECCI) has been used for metallic systems, but it was not suitable to clearly distinguish the main microstructural features [¹⁹,²⁰,²¹], while Scanning Ion Microscopy (SIM) in a Focused Ion Beam (FIB) provides a better contrast, as reported in Ref. [²²].

Several recent studies in ductile materials [²³,²⁴,²⁵,²⁶,²⁷] have evaluated the plastic zone affected during the indentation process by using different advanced characterization techniques. In the majority of these studies, testing was performed within single grains and unique phases. Little information is available within the literature on the interaction of the plastic zone with the grain boundary.

Within this framework, the purpose of this study was (is) to analyze the deformation mechanisms, activated under certain stress fields (monotonically and cyclic), at the nanometric length scale, by using the nanoindentation technique for TRIP steel specimens. In doing so, advanced characterization techniques (e.g., TEM, SIM/FIB and 3D-FIB tomography) were used to document the main deformation mechanisms.

2. Experimental procedure

2.1. Material

The metastable stainless steel studied is an AISI 301LN (equivalent to EN 1.4318). It was supplied as 1.5 mm thick sheets by Outokumpu (Finland), in two different conditions: annealed (fully austenitic) and skin-passed up to 15% of thickness reduction. The chemical composition is given in Table 1. The specimen was annealed at 1100 ºC for 1 hour. The heating and cooling rate was kept constant to 5 ºC·min^-1.

Prior to the micromechanical and microstructural characterization, specimens were polished using diamond suspensions with gradually decreasing particle sizes from 30 to 1 µm. After mechanical polishing, in order to remove the phase transformation induced during the mechanical polishing process, all samples were electrochemically polished at 15ºC and a constant voltage of 23 V.

2.2. Micromechanical properties: Monotonic and cyclic tests

Local deformation was induced by means of the nanoindentation technique, with an ultra-nanohardness tester (UNHT) from CSM Instruments by using a Berkovich tip indenter.

Two sets of experiments were conducted under displacement control mode, working at 200 nm of maximum penetration depth in order to induce different stresses within the austenitic grains: i) monotonic tests with a loading/unloading indentation rate of around 10 mN/min, and ii) cyclic loading experiments (~ 200 cycles). In both experiments, after the first cycle, the maximum applied load was held constant during 10 s in order to stabilize the deformation induced during the indentation process.

2.3. Microstructural characterization

2.3.1. Electron backscatter diffraction, EBSD

Prior to nanoindentation tests, electron backscatter diffraction (EBSD) mapping was performed to select the grains of interest (grains oriented with their <111> axis parallel to the loading axis). EBSD measurements were conducted in a field emission scanning electron microscopy (FESEM) JEOL 7001F equipped with an orientation imaging microscopy system. EBSD measurements were performed with a constant scanning step (100 nm) at an acceleration voltage of 20 kV.

2.3.2. Focused ion beam

A dual beam workstation Zeiss Neon 40, working at 30 keV, was used to prepare cross sections of the residual nanoindentation imprints, to perform SIM imaging, as well as to extract TEM lamellae. The coarse and fine milling were performed with a maximum current of 900 and 200 pA, in order to prevent damage or phase transformation due to the Ga⁺ implantation [²⁸]. SIM images were acquired at 50 pA and a pixel dwell time of 18 µs in line integration mode.

The tomography was performed by sectioning the sample in 25 nm steps by the FIB, at a current of 100 pA. The dimensions of the total examined volume were 4.3 x 3.7 x 1.8 µm. The 3D-reconstruction was performed with the software FEI Avizo 8.0^.

2.3.3. Transmission electron microscopy

In order to correlate the SIM contrast with microstructural features, TEM investigations were performed by using Philips CM200 microscope, working at 200 kV, to obtain Bright (BF) and Dark (DF) Field images, as well as Selected Area Electron Diffraction (SAED) patterns from the deformed regions of interest.

3. Results and discussion

To understand the main plastic deformation mechanisms in metastable stainless steels, it is important to consider that the contrast in SIM images depends on both the local crystallographic orientation of different grains of the same phase, as well as the mass density of the different phases [²⁹]. Since α’-martensite has a lower density than austenite, ions can channel deeper into the material before interacting with atoms and inducing the emission of secondary electrons (SEs). Within this context, martensite is expected to appear, in SIM images, darker than austenite, because of its lower density.

3.1. Plastic deformation mechanisms for the as-received specimens

Figure 1 shows the microstructure of macroscopically deformed samples, i.e., by industrial skin-passing. A large area of the sample surface was imaged by the SIM detector, as it is represented in Figure 1a. In this image, a heterogeneous microstructure with different prominent features (i.e., shear bands, forest dislocations, etc.) is clearly evident, which is similar to that observed in preliminary works on deformed metastable stainless steel, as reported in references [³⁰,³¹].

Figure 1b exhibits a detailed magnified image of the plastic deformation features from the white square in Figure 1a, where several parallel shear bands appear. Attempting to get deeper knowledge of the deformation scenario, a cross section from the center of the shear bands (green dashed line in Figure 1b) was prepared and subsequently examined using FIB. In order to reveal the microstructure in the trench of interest, the zone of study was exposed to the ion beam during several seconds to produce ion etching and afterwards it was observed by SIM.

Figure 1c shows that the aforementioned shear band microstructure is not only a surface feature and that it is composed by platelets of some tenths of nanometers. Moreover, this image presents a grey shade gamut, which may be assigned to different phases (inset in Figure 1c). As it is evident from the gray scale SIM image, the shear bands exhibit a darker contrast than the austenitic phase, which may be associated with a volume change generated during the _ phase transformation, as reported in references [³²,³³,³⁴]. From the abovementioned image, it is manifest that metastable stainless steels present a complex deformation microstructure, being one of the predominant deformation mechanisms the cross-shear bands, with an angle between the primary and secondary band systems around 71º (see Figure 1d). This value corresponds to the angle between two {111} slip planes viewed along the _direction.

To acquire more information on the plastic deformation mechanisms, a lamella of the studied area was extracted by FIB and examined by TEM. Using SAED pattern recorded along _coupled with dark field (DF) observation mode, from - and ’-martensite, Figures 1e and 1f where obtained, respectively. Both images allow distinguishing unambiguously between both martensitic phases. In doing so, the DF-TEM images reveal two coexisting plastic deformation mechanisms: microscopic shear bands and forest dislocations inside the bands, and in particular at the band intersection sites. The analysis of the electron diffraction patterns (inset Figure 1e) exposes that the shear band interaction region corresponds to: i) ’-martensite, mainly developed at the band intersections, and ii) -martensite with platelet shape inside the shear bands. This presence of -martensite as platelets has been previously reported by Talonen et al. [Error: Reference source not found]. Also these TEM observations are in agreement with a known mechanism concerning the formation of martensite embryos in metastable austenitic stainless steels [³⁵,³⁶].

The calculated lattice parameter for ’-martensite reported in Table 2 is in fair agreement with the work published in Refs. [³⁷,³⁸]. The lattice parameters (a,c) of the hexagonal unit cell have been determined through the diffraction pattern matches with the values that were obtained using the following expression [³⁹].

a = (2a_^ and c = 2·(3a_)^1/2

where a_ is the lattice parameter of austenite (a_ is around 3.59 nm). The lattice parameters values are coincident with those reported by Hedström et al. [Error: Reference source not found], Petein [⁴⁰] and Tavares et al. [⁴¹].

On the other hand, different artifacts produced by FIB technique have been reported due to the high energy of the ion beam [⁴²]. Therefore, a reference TEM sample was achieved by twinjet polishing in order to be compared with the FIB lamella. A huge density of microscopic shear bands appears in Figure 2, which is a TEM image from this reference sample. This microstructure resembles that observed in Figure 1, obtained from the sample lifted out by FIB and imaged by SIM. This correspondence between both images, in Figures 1 and 2, allows concluding that the low-energy FIB beam applied in the present study did not change the original microstructure, and thus the observed microstructure cannot be considered as an artifact.

3.2. Plastic deformation under different stress fields: monotonic and cyclic nanoindentation tests

After confirming the agreement consistency between FIB-SIM and TEM images for macroscopically deformed samples, the next step was the characterization of the plastic deformation induced by monotonic nanoindentation.

Figure 3 corresponds to a cross section of the residual nanoindentation imprint produced by a monotonic test in an austenitic grain observed along <111> axis. SIM image is depicted in Figure 3a, where the region immediately under the imprint, delimited by a white dashed line, corresponds to high deformed austenite. In this region, high dislocation density yields a brighter contrast within the SIM images than the less deformed areas. This phenomenon may be related to local lattice rotation caused by defects that hinder the channeling process, resulting in a higher fraction of SEs emitted near the surface. Thus, the plastic deformation zone, which contains a high dislocation density, appears brighter.

In order to get more information related to the main deformation mechanisms induced under monotonic indentation tests, TEM lamella was extracted and subsequently observed (Figures 3b to 3d). Figure 3b is a BF-TEM image where several deformation bands are clearly seen. On the other hand, the corresponding DF-TEM image (Figure 3c) reveals that the dark patch corresponds to a small α’-martensitic grain, with a size around 0.02 µm² with a shape similar to the plastic zone induced by the indentation. The martensitic character of the entire deformed region examined was corroborated through analysis of SAED patterns, pointing then out the phase transformation induced under the stress field (inset between figures 3c and 3d). This finding is in close agreement with previous work conducted by Kim et al. [⁴³], where the different discontinuities in the loading-unloading curve were attributed to a gradual phase transformation (from a metastable austenite grain to martensitic grain) during the nanoindentation process.

Figure 4 illustrates another residual imprint, but this one produced as a consequence of a cyclic test performed at the nanometric length scale inside a <111> austenitic grain. Figure 4a shows the corresponding local crystallographic mapping determined by means of EBSD. Several slip traces near the residual imprint can be observed in Figure 4b. In this figure, dashed lines indicate the exact positions where selected series of cross-sections were milled through the entire imprint. These cross-sections were subsequently observed by SIM (Figures 4c to 4f), in order to discern the evolution of the main deformation mechanisms.

Figure 4c shows that, even at the corner of the imprint (section corresponding to line 1 in Figure 4b), inhomogeneous contrast is observed and two different areas can be distinguished within the austenite grain. In contrast to the case of monotonic indentation test presented in Figure 3, after cyclic nanoindentation the austenitic grain exhibits a highly-deformed substructure, which expands around the residual imprint and along the grain boundary. As it is evident when cross-sections from the corner (Fig. 4c) to the center of the imprint (Fig. 4f) are compared, the closer to the indentation center the more pronounced becomes the presence of a dark α’-martensitic zone. This zone grows from the center of the residual imprint, where the plastic stress becomes maximum, with an angle close to 70.5˚ (thus, parallel to a {111} slip plane), towards the grain boundary. Moreover, several dark contrast bands were observed parallel to the grain boundary, as can be seen in Figure 4e.

Attempting to get more detailed knowledge on the deformation scenario induced by cyclic indentation, TEM lamellae were directly extracted by FIB from the center of residual imprints presented in Figure 4f. Figure 5a is a BF-TEM image, whereas 5b and 5c are DF-TEM views of the main deformation features present in the region of interest. Under these observations, the region where austenite has transformed to ’-martensite, after the 200 indentation cycles, has around 3.1 µm², which exceeds its counterpart after single indentation by more than two orders of magnitude. This fact may indicate that during each individual indentation test, the martensite induced just below the residual imprint grows continuously, due to the accumulation of the plastic deformation.

The diffraction pattern and the morphology of the martensitic phase show that the latter is of cell-type. Dislocation cell-type martensite is known to form after severe deformation by break up of lath martensite and its diffraction pattern is a superposition of many differently oriented martensitic phases, resulting from crystallographic distortion due to the plastic deformation mechanisms activated during the deformation process, as reported in Refs. . It is an interesting finding that this type of martensite can also be induced by cyclic nanoindentation, since it shows that cyclic nanoindentation is a suitable testing mode to study the evolution of plasticity on individual austenitic grains as a function of their local crystallographic orientation. Dislocation cell-type martensite typically contains a high dislocation density, as confirmed by the diffuse DF image presented in Figure 5b.

Placing the focus on the grain boundary region, BF TEM-image (Figure 5d) shows that plasticity transfer took place through this boundary, whereas the corresponding DF image reveals well defined martensitic laths (Figure 5e). From the diffraction pattern in Figure 5f, it can be concluded that these martensitic laths adopt Kurdjumov-Sachs [⁴⁴] orinentation relationship (_; _) with the surrounding austenitic phase, and apparently have evolved from it. As it has been mentioned and presented along this section, SIM images (Figure 4) yield a better contrast and definition of the deformed microstructure, especially for the dislocation cell-type martensite.

In order to achieve a better understanding of the dimensions of the plastic zone, a 3D reconstruction was carried out with a total of 72 SIM images of the residual imprint produced by cyclic indentation (dashed rectangle region presented in Figure 4b, number 4). Figure 6 shows several images taken from the reconstructed volume, indicating that the growth of the martensitic area occurs at the expense of the austenitic plastic zone, confirming the phase transformation reaction_. This is in agreement with the common assumption that the presence and interaction of dislocations reduces the nucleation barrier for martensitic formation, as reported in Refs. [⁴⁵,⁴⁶]. 3D-FIB/SIM tomography visualizes particularly well that both the martensitic phase and austenite deformation disseminate along the grain boundary.

4. Conclusions

The present study highlights the usefulness of SIM imaging. Particularly, regarding the evolution of deformation in metastable austenitic steels, the contrast provided by SIM observations can be directly related to the accumulation of plastic deformation and phase transformation, allowing thus a detailed characterization. Compared to other advanced characterization techniques, such as EBSD to evaluate the crystallographic orientation at the surface and TEM to determine the main deformation mechanisms at the microstructural level, SIM results are faster and easier. Moreover, 3D-FIB/SIM tomography provides complete information about the microstructural evolution and the main plastic deformation features in a volume. This makes FIB not only a powerful technique for sample preparation at small scale, but also for contrast-enhanced imaging.

Concerning the main deformation mechanisms observed in the AISI 301LN steel, for the skin-pass samples shear bands as well as forest dislocation are predominant, while for the deformed material under different stress fields produced by nanoindentation, the phase transformation from austenite to ’- and -martensite is the most relevant mechanism.

Acknowledgements

The authors acknowledge the support during the experimental procedure of S. Migot, C. Gendarme and I. López-Insa. I. Sapezanskaia would like to thank DOCMASE program for its financial support and J. Säynäjäkangas and A. Kalapudas, from Tornio Research Centre, Outokumpu (Finland), for providing the steel samples.

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