Fatigue behavior of an ultrafine-grained metastable CrMnNi steel tested under total strain control

Accepted Manuscript Fatigue behavior of an ultrafine-grained metastable CrMnNi steel tested under total strain control M. Droste, C. Ullrich, M. Motylenko, M. Fleischer, A. Weidner, J. Freudenberger, D. Rafaja, H. Biermann PII: DOI: Reference:

S0142-1123(17)30393-6 https://doi.org/10.1016/j.ijfatigue.2017.10.001 JIJF 4475

To appear in:

International Journal of Fatigue

Received Date: Revised Date: Accepted Date:

30 May 2017 29 September 2017 1 October 2017

Please cite this article as: Droste, M., Ullrich, C., Motylenko, M., Fleischer, M., Weidner, A., Freudenberger, J., Rafaja, D., Biermann, H., Fatigue behavior of an ultrafine-grained metastable CrMnNi steel tested under total strain control, International Journal of Fatigue (2017), doi: https://doi.org/10.1016/j.ijfatigue.2017.10.001

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Fatigue behavior of an ultrafine-grained metastable CrMnNi steel tested under total strain control M. Drostea,*, C. Ullrichb, M. Motylenkob, M. Fleischera,d, A. Weidnera, J. Freudenbergerb,c, D. Rafajab, H. Biermanna a

b

c d

Institute of Materials Engineering, Technische Universität Bergakademie Freiberg, GustavZeuner-Straße 5, 09599 Freiberg, Germany Institute of Materials Science, Technische Universität Bergakademie Freiberg, GustavZeuner-Straße 5, 09599 Freiberg, Germany IFW Dresden, Helmholtzstraße 20, 01069 Dresden, Germany Schmiedeberger Gießerei GmbH, Altenberger Straße 59a, 01744 Dippoldiswalde, Germany

* Corresponding author:

E-mail adress: [email protected] telephone number: 0049 3731 392336

Keywords: ultrafine-grained, martensitic transformation, metastable steel Abstract An ultrafine-grained (UFG) microstructure in a metastable austenitic CrMnNi steel was achieved using a thermo-mechanically controlled process by rotary swaging and subsequent reversion annealing. The material with an average grain size of 0.7 m was cyclically deformed

in

total

strain

controlled

tests

at

strain

amplitudes

in

the

range

of

0.3 % ≤ t/2 ≤ 1.2 %. This treatment increased the cyclic stress amplitudes as well as the fatigue life in comparison with the conventionally grained counterpart. For strain amplitudes t/2 ≥ 0.4 % a martensitic phase transformation occurred, which was observed in situ by a ferrite sensor as an increase of the ’-martensite fraction. The microstructure changes, and the deformation mechanisms in particular, were investigated by means of electron backscatter diffraction, scanning electron microscopy in transmission mode and transmission electron microscopy that revealed the formation of small ’-nuclei which rapidly grew until the entire austenitic grain was transformed.

1. Introduction Due to the necessity for lightweight construction there is a persistent need for high-strength steels which enable a reduced material thickness of construction elements and components. In particular, austenitic stainless steels exhibit outstanding mechanical properties combining excellent ultimate tensile strength (UTS) and a good ductility. Even more, these properties 1

can be adjusted by the chemical composition of the steels that controls the dominating deformation mechanisms, i.e., the formation of deformation-induced ’-martensite or twinning [1–7]. However, a common disadvantage of austenitic stainless steels is their quite low yield strength. Beside the solid solution, precipitation or work hardening, the grain refinement is an excellent method for eliminating this drawback [e.g. 8]. The most common way for metastable austenitic stainless steels to produce an UFG microstructure is a well-defined thermomechanically controlled processing (TMCP) [9–16]. After a cold deformation step to induce a high amount of ’-martensite, an annealing treatment yields the reversion of the martensite back into austenite [9–16]. The resulting microstructure strongly depends on the parameters of cold deformation and annealing, e.g. annealing temperature and time. The reversion process is well investigated. For interstitial free steels it is known that the reversion is based on a shear process without any diffusion [17–20]. Consistently, this reversion path was reported by Weidner et al. [21] for a steel of similar chemical composition as the present material by investigations on the texture evolution during TMCP. However, the dominant deformation mechanisms in austenitic steels strongly depend on the stacking fault energy (SFE) that decides whether the plastic deformation occurs predominantly via deformation-induced phase transformation, twinning or dislocation glide. Hence, the SFE is an important parameter affecting the deformation behavior and the development of the microstructure during deformation, which is already well investigated for the conventionally grained (CG) counterpart of the present steel for static [3,5,22–27] as well as for cyclic [1–3,28–30] loading. But since the literature discusses the influence of grain size on the martensitic transformation as well as on the apparent SFE [31–33], i.e. an increasing apparent SFE with decreasing grain size, the question arises whether the deformation behavior and the characteristics of ’-martensite formation differ for the UFG and the CG states of the present austenitic steel. For the latter, the ’-martensite nucleation usually occurs at intersecting deformation bands via the intermediate -martensite [34–36]. For submicron austenitic grains, on the other hand, authors reported either no phase transformation at all [37,38] or a change in the nucleation mechanism, i.e. the formation of ’-martensite at grain boundaries 2

or twins [39,40]. However, these findings are based on quasi-static tensile tests. Previous studies dealing with the fatigue behavior of UFG austenitic steels are almost exclusively based on stress controlled fatigue tests [41–46] which either show almost no phase transformation in the reversion-annealed ultrafine grains at all [41–45] or the ’-martensite formed in the first cycles leads to a strong decrease in the strain amplitude [46] suppressing a further phase transformation. Hence, these studies are focusing on other aspects of the microstructure evolution as well as on the fracture mode and the stress-based fatigue life which is enhanced for UFG steels compared to the CG counterparts due to their higher strength. Anyway, because of the very limited data for UFG austenitic steels tested in strain control [42,47] one has to take into account several investigations done on other UFG metals for strainbased fatigue life considerations. Usually, the higher strength of UFG metals causes an enhanced lifetime in the high-cycle fatigue (HCF) regime whereas the lower ductility causes an inferior lifetime in the low-cycle fatigue (LCF) regime [48–50]. This leads to an intersection of the total strain amplitude-fatigue lifetime curves of the UFG and CG states, respectively. The main goal of the present study is the characterization of the fatigue behavior of an UFG metastable CrMnNi steel exhibiting a low content of interstitials under total strain control. This includes (i) the cyclic stress response, (ii) the ’-martensite nucleation and evolution of the phase fraction, (iii) the fatigue life, and (iv) microstructure investigations of the initial and failed states. Furthermore, these results obtained for the UFG steel are compared with the characteristics of a CG counterpart of similar chemical composition which exhibits an average grain size of 14 m.

2. Experimental details The chemical compositions of the investigated high-alloy metastable austenitic CrMnNi steels are given in Table 1. The material for the ultrafine-grained state was cast (Institute of Iron and Steel Technology, TU Freiberg, Germany) using a cylindrical slightly tapered die exhibiting a minimum diameter of 50 mm and a maximum diameter of 70 mm at a length of 400 mm. For the present investigations a 140 mm long segment (min. diameter 51 mm) of 3

this cast cylinder was turned to a diameter of 50 mm to remove the cast skin. The next step was hot forging at 1200 °C to a diameter of 27 mm. Afterwards, the material has undergone a TMCP whose parameters were varied in terms of the degree of deformation (final diameters of 11.6, 10.4 and 8.3 mm) and the annealing time (2.5 – 20 min) in order to achieve an UFG microstructure. The different sample states were examined by means of back-scattered electron (BSE) imaging to determine the most suitable parameters for a homogeneous UFG microstructure. The finally selected processing steps are shown in Fig. 1. The metastable austenitic material was cold deformed by rotary swaging in 11 passes, reducing the diameter from 27 mm to 8.3 mm to receive a high amount of deformation-induced ’-martensite. Between the single steps, the rods were cooled to room temperature in order to avoid too high temperature. The last process step was an annealing at 700 °C for 5 minutes which led to a reversion of the martensite back to austenite and thereby to the formation of desired UFG microstructure. Table 1: Chemical compositions of investigated steels in wt%. The carbon concentration was determined by the combustion infrared detection technique, the concentration of nitrogen by the inert gas fusion infrared and thermal conductivity detection and the other elements by glow discharge optical emission spectroscopy (GDOES), respectively. Fe

Cr

Mn

Ni

C

N

Si

UFG

bal.

16.6

7.1

6.4

0.04

0.02

1.0

CG

bal.

16.5

6.4

6.8

0.04

0.04

1.0

Fig. 1: Processing steps to achieve an UFG microstructure by reversion annealing.

4

The CG counterpart was cast (ACTech, Freiberg, Germany) and gas atomized (TLS, Bitterfeld, Germany). The steel powder exhibited a mean diameter d 50 of 28.1 m and was sintered by hot pressing under vacuum at 1250 °C for 30 minutes at a pressure of approximately 30 MPa (Fraunhofer IKTS, Dresden, Germany). The heating and cooling rates were set to 10 and 5 K/min, respectively. The resulting material exhibited a porosity < 1 % according to the Archimedean principle. The cylindrical fatigue specimens with a gauge length of 14 mm and a gauge diameter of 5 and 6 mm for the UFG and CG material, respectively, were ground and polished mechanically. The specimens were cyclically deformed under total strain control following a triangular strain function with a constant strain rate of 4×10-3 s-1. The strain amplitudes were varied in the range 0.3 % ≤ t/2 ≤ 1.2 %. All tests were performed under symmetrical push-pull conditions (load ratio R = -1). The tests were carried out on servohydraulic testing systems (MTS Landmark 100, MTS Landmark 250) and accompanied by measurements using a Feritscope® (Fischerscope® MMS® PC) to determine the evolution of the ’-martensite formation in situ during cyclic loading. Talonen et al. [51] compared the Feritscope® readings in Fe-% (ferromagnetic fraction) with other techniques for measuring the ’-martensite content and concluded that the measured values have to be multiplied by a factor of 1.7 to obtain the ’-martensite fraction in mass-%. However, this linear correlation is only valid up to 55 Fe-%. Because some ’-martensite fractions presented in this work exceeded this boundary value, every designation of the ’-martensite content  is given in Fe-% and represents the Feritscope® readings without any correction factor. For reference, quasi-static tensile tests were carried out at a constant cross-head velocity of 2 mm/min which leads to an average strain rate of about 2×10-3 s-1. The microstructures of the initial states as well as of failed specimens were characterized by means of scanning electron microscopy (SEM) using a field-emission Tescan Mira 3 FE-SEM. The microscope was operated at an acceleration voltage of 20 kV for the EBSD (EDAX, Ametec) measurements and at 30 kV for the investigations performed in the transmission mode (t-SEM), respectively. The electron backscatter diffraction (EBSD) measure5

ments with a step size of 70 nm were performed for the determination of the average grain size

that was area-weighted according to the following relationship

where Ai and di are the area and diameter of the ith grain, respectively. Misorientations ≥ 12° were designated a grain boundary. The samples for the EBSD investigations were cut parallel to the loading axes and prepared by conventional grinding and polishing steps followed by final vibration polishing (SiO2 suspension of 0.02 m grade for 24 h). The transmission electron microscopy (TEM) foils were prepared by cutting thin slices from the bulk material using a diamond wire cutting machine. Afterwards, they were mechanically grounded and finally thinned by electrolytic etching. The TEM investigations were conducted on a transmission electron microscope JEOL JEM-2200FS that is equipped with a fieldemission gun and an EDX detector. The TEM was operated at 200 kV.

3. Results and discussion 3.1

Microstructure of the reversion annealed state

The microstructure of the UFG steel after reversion annealing is shown in Fig. 2. This initial state was almost fully austenitic (cf. Fig. 2a) and exhibited an average grain size of = 0.7 ± 0.4 m as determined from the grain map shown in Fig. 2b. On grain boundaries as well as in the interior of austenitic grains precipitates were observed which have been formed during the reversion annealing step at 700 °C [52,53]. Their size differed from a few nanometer to approximately 100 nm. Energy dispersive X-ray spectroscopy (EDX) and selected area electron diffraction (SAED) in TEM yielded that the precipitates shown in Fig. 2c are Cr23C6 carbides, which have an orientation relationship with the austenitic matrix. In addition, it is known from literature that the interfaces are partially coherent [53]. Furthermore, many grains contain a few individual stacking faults already in the reversion annealed state (Fig. 2d).

6

Fig. 2: Microstructure investigations of the reversion annealed state. a) EBSD phase map, gray (band contrast)—austenite, yellow—-martensite and blue—’-martensite, b) EBSD orientation map, c) TEM bright field image of Cr23C6 carbides and SAED pattern showing orientation relationship of Cr23C6 carbide (faint spots) and the austenitic matrix (intense spots), d) t-SEM image of stacking faults inside an austenitic grain.

3.2

Tensile properties

The stress-strain curves of the UFG and the CG steel obtained in quasi-static tensile tests are shown in Fig. 3. According to the plot of engineering stress versus engineering strain (Fig. 3a) it is evident, that the UFG material exhibits a markedly higher yield strength y than the CG steel, according to the Hall-Petch relationship:

7

where 0 is the friction stress, k a constant and D the grain size. The yield strength of 923 MPa, which in this study is specified as the stress at 0.5 % plastic strain due to microplasticity effects, is 2.5 times higher compared to the respective CG value of 365 MPa. The data in the literature concerning the dependence of the yield strength on the grain size in austenitic stainless steels is dissenting in terms of a quantitative correlation. Some authors found a smaller increase of yield strength with decreasing grain size [54–56] while others reported comparable relationships [9,57] and a similar yield strength at an average grain size of about 0.7 m. However, the discrepancies could be caused by several reasons like i) different approaches for grain size estimation, ii) different chemical compositions, especially regarding the carbon and nitrogen contents [56], iii) different precipitation behavior during annealing or iv) different TMCP procedures which for instance lead to different fractions of retained austenite. However, no study reported a higher increase of the yield strength, i.e. the present strengthening due to grain refinement seems to be very effective for the UFG steel. The ultimate tensile strength (UTS) is higher for the UFG steel as well, although the strain hardening is less pronounced. This lower strain hardening of the UFG state leads to an earlier onset of necking and accordingly to a considerably smaller uniform elongation and elongation to fracture (cf. Fig. 3). These findings are in agreement with results on UFG austenitic stainless steels reported in the literature [39,54,56,58].

8

Fig. 3: Mechanical properties obtained from tensile testing. Engineering stress plotted versus engineering strain (yield strength Rp0.5, ultimate tensile strength UTS, uniform elongation UE, and strain to rupture A r).

3.3

Cyclic behavior

Fig. 4 shows the cyclic deformation curves and the ’-martensite evolution during fatigue for the UFG as well as for the CG state. As expected, the latter exhibits considerably lower stress amplitudes at the beginning of the cyclic loading. For strain amplitudes t/2 ≤ 0.5 %, an initial hardening regime followed by cyclic softening is observed in CG specimens. This behavior is well known for stainless steels, and is related to an increasing dislocation density and their interactions in the beginning and, subsequently, a rearrangement and annihilation of dislocations, respectively, in the stage of cyclic softening [59,60]. After a certain number of cycles a pronounced secondary hardening regime occurred caused by the onset of the deformation-induced ’-martensite formation (cf. Fig. 4c) [28]. With increasing strain amplitude, the onset is shifted to a lower number of cycles and the hardening gets more pronounced due to the higher amount of ’-martensite formed.

9

Fig. 4: Cyclic deformation curves (a, b) and corresponding ’-martensite evolution (c, d) at different strain amplitudes for the CG reference material (a, c, according to Glage [61]) and the UFG state (b,d). The UFG material exhibited much higher stress amplitudes right from the beginning due to the much higher yield stress. For strain amplitudes t/2 ≤ 0.5 % a mild initial cyclic hardening followed by a stage of cyclic softening occurred. The onset of the martensitic phase transformation led to a stabilization of the stress amplitude in case of t/2 = 0.4 % and to a slight secondary hardening in case of t/2 = 0.5 %, respectively. The secondary hardening, however, was much less pronounced than in case of the CG material (compare Fig. 4a and b). Higher strain amplitudes lead to cyclic hardening right from the beginning up to a certain maximum and, thereafter, softening until failure occurred. The cyclic hardening is most probably due to an increase in the dislocation density and an early onset of ’-martensite formation (cf. Fig. 4d). In contrast, the cyclic softening occurred despite an ongoing ’martensite formation. As will be shown later on, this is assumed to stem from a grain coarsening of the martensitic grains due to a rearrangement of low angle grain boundaries

10

(LAGB). However, even prior to the maximum of the stress amplitude the martensitic phase transformation obviously had a smaller influence on the stress amplitude compared to the CG material, i.e. the hardening effect due to the formation of ’-martensite is considerably smaller. In addition, this statement also holds true comparing the medium and small strain amplitudes. The reason for this dissimilar hardening behavior is the origin of the hardening effect due to the phase transformation. In this kind of TRIP steels with a low content of interstitials the hardening effect is only to a small extend a result of the higher hardness of the formed ’-martensite as shown by nanoindentation [62]. According to Martin et al. [25] and Weidner et al. [62], a significant contribution to the hardening effect is supplied by the reduced mean free path for the dislocations due to the formation of ’-martensite inside the deformation bands of the austenite in the CG state. This mechanism is regarded to be minor in the UFG steel, since the ’-martensite nuclei are expected to grow rapidly into the entire grain, as it will be discussed later. Thus, their action as a barrier for dislocation movement in the parent austenite grain is limited to a few cycles. The second substantial contribution to the hardening effect in the CG state is the smaller size of the ’-martensite grains compared to the grain size of their austenitic matrix, as was shown by Weidner et al. [62]. But since the grain sizes of UFG austenite and martensite are in the same order of magnitude, this grainrefinement hardening effect should also be negligible. Accordingly, the influence of both hardening mechanisms acting in the CG material, i.e. the reduced mean free path in deformation bands and the grain-refinement, is considerably decreased for the UFG steel leading to a reduced cyclic hardening capacity due to the phase transformation.

11

Fig. 5: The ’-martensite evolution plotted versus the cumulated plastic strain p of the CG (a, data after Glage [61]) and UFG (b) state, respectively. In metastable austenitic steel with conventional grain size, the deformation-induced formation of ’-martensite usually occurs at intersecting deformation bands containing a high amount of -martensite [63]. From the point of view of the crystal structure, the -martensite is formed through a certain arrangement of stacking faults on {111} planes in the fcc austenite. If these stacking faults are located on every second {111} lattice plane of austenite, the stacking sequence of lattice planes changes from ABCABC typical for the fcc crystal lattice to ABABAB, which is characteristic for the hexagonal lattice of -martensite [64]. However, since the ’martensite formation is triggered by the double shear mechanism of two intersecting stacking faults on different slip planes [35,65] a certain threshold value of the cumulated plastic strain p,th has to be exceeded to trigger the fatigue-induced transformation, which is the reason for the incubation period. Fig. 5 illustrates the relationship between the ’-martensite content  and the cumulated plastic strain p. As already observed by other authors [28,60,66–69], the threshold value p,th increases with decreasing strain amplitudes for both materials. This is due to a decrease in lattice defects and density of deformation bands with decreasing strain amplitudes, which leads to a reduction of nucleation sites for the formation of ’-martensite [60,66,67]. The ’-martensite in the UFG steel seems to form in a similar manner. Just instead of deformation bands, the grains contain numerous interacting stacking faults. Nevertheless, there are differences between the UFG and CG material when comparing the threshold values p,th 12

of the same strain amplitudes. For the medium and high strain amplitudes t/2 ≥ 0.5 % the threshold values are smaller for the UFG state, which is assumed to be caused by the higher applied stress amplitudes. These high stress amplitudes lead to very small ’-martensite nuclei arising from at least two individual stacking faults on different slip planes, triggering the phase transformation at an earlier stage of deformation compared to the CG material. For the strain amplitudes t/2 < 0.5 % a contrary effect is observed, i.e. the threshold value p,th is higher for the UFG material. For t/2 = 0.3 % actually only in the CG material a fatigueinduced martensitic phase transformation is measured. A reason for this behavior could be the distinct decrease of cyclic stresses with decreasing strain amplitude in case of the UFG state. In contrast, in the CG material minor differences were obtained between the highest and smallest strain amplitudes regarding the cyclic stress until the threshold value is exceeded.

3.4

Fatigue lives

The curves showing the total strain amplitude t/2 vs. reversals to failure 2Nf (Fig. 6a) were calculated according to the Manson-Coffin and Basquin relationship:

where

represents the fatigue strength coefficient, E the Young’s mo ulus, b the fatigue

strength exponent,

the fatigue ductility coefficient and c the fatigue ductility exponent. The-

se fatigue lifetime parameters were determined with respect to compatibility following an approach of Niesłony et al. [70]. The respective stress amplitudes /2, elastic strain amplitudes e/2 and the plastic strain amplitudes pl/2 were determined at 0.5 Nf.

13

Fig. 6: a) Total strain based fatigue life curves and b) total strain amplitudes plotted versus the cumulative plastic strain at the end of lifetime. As discussed earlier, regarding total strain controlled testing UFG materials usually exhibit higher fatigue lives in the HCF regime and lower fatigue lives in the LCF regime compared to their CG counterparts due to their higher strength and lower ductility, respectively. Furthermore, several studies on CG metastable steels already investigated the influence of the austenite stability and thereby of the martensitic phase transformation on the fatigue life, either by a variation of the chemical composition [1,2] or by variation of the test temperature [28,71]. All authors found that an increasing ’-martensite fraction leads to a fatigue life enhancement in the HCF regime because of the higher strength of the material, but to a fatigue life reduction in the LCF regime due to a decrease in fatigue ductility. Thus, taking into account these results, i.e. the lifetime dependence regarding the grain size and ’-martensite fraction, and set in context with the present ’-martensite evolutions, it is expected that the UFG TRIP steel exhibits a lower fatigue life in the LCF regime and a higher fatigue life in the HCF regime than the CG material. However, according to the total strain amplitude vs. reversals to failure curves in Fig. 6a, this expectation indeed holds true for the latter but for high strain amplitudes the UFG steel exhibits a higher lifetime as well. Chlupová et al. [47] report similar results for a 301LN stainless steel which, however, could just show a tendency due to the limited data on the CG counterpart. However, the observed behavior could be based on a more homogenous strain distribution at high strain amplitudes with medium or high fractions of ’-martensite in case of the UFG

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state. Since the hardness of the ’-martensite is only slightly higher compared to the hardness of the austenite [62], the resistance against deformation strongly depends on the grain size. Hence, the resistance against deformation is more or less comparable for the two phases in the UFG state. As a consequence, under cyclic loading and at a given total strain amplitude both phases should have to bear their fair share of strain according to the phase distribution. In contrast, in the CG state the martensitic phase exhibits a substantially smaller grain size and thus a higher resistance against deformation than the austenitic phase, which is why the latter has to carry a disproportionately high fraction of strain compared to its phase fraction. Hence, the strain distribution is more inhomogeneous and the strain is localized in the austenitic phase. With increasing ’-martensite content this mismatch increases, leading to shorter fatigue lives at high strain amplitudes. Accordingly, the lower ductility of the UFG material is overcompensated by the more homogeneous strain distribution. This assumption is supported by Fig. 6b which illustrates the cumulative plastic strain at the end of fatigue life p,Nf plotted versus the total strain amplitude. For the UFG as well as for the CG state the cumulative plastic strain at failure p,Nf decreases with increasing strain amplitude. Moreover, for small strain amplitudes a lower cumulative plastic strain at failure p,Nf is observed for the UFG material. With no or a low amount of ’-martensite, respectively, the ability of bearing plastic strain is primary a matter of the ductility of the austenite. As the ductility is higher for the CG state, the UFG steel fails at a lower cumulative plastic strain. However, the plastic strain amplitude pl/2 of the latter is considerably smaller due to the higher strength. Hence, the plastic strain per cycle is lower, leading to a superior fatigue life of the UFG material. On the other hand, for high strain amplitudes with high amounts of ’-martensite the cumulative plastic strain at failure is the same for the UFG and for the CG material, respectively. This is most probably due to the more homogeneous strain distribution between the two phases in case of the UFG state, in contrast to a strain localization in the austenitic phase in case of the CG state.

15

Fig. 7: Fatigue life curves in terms of stress amplitude at 0.5 Nf vs. reversals to failure. The stress amplitude /2 at 0.5 Nf vs. number of reversals to failure 2Nf shown in Fig. 7 is well described using the Basquin relationship:

The UFG state exhibited a markedly higher lifetime for all studied strain amplitudes. Hence, it can be assumed that stress controlled fatigue tests will result in significantly higher cycles to failure for the UFG state. This statement is confirmed by many investigations carried out by other authors on UFG metals, cf. [49,50]. Thus, an investigation on dual phase ferrite/martensite low carbon steel [72] and several studies on ultrafine-grained austenitic Cr-Ni steel [43,45] are based on stress-controlled fatigue tests and confirm the higher fatigue life compared to the respective CG counterparts for stress-controlled tests.

3.5

Microstructure evolution

Fig. 8 illustrates the development of the microstructure after cyclic loading with increasing strain amplitudes in terms of EBSD phase maps (confidence index CI ≥ 0.1) and grain size distributions. As expected, the fraction of the austenitic phase decreases with increasing strain amplitude. At t/2 = 0.8 % merely a fraction of 8 % remained austenitic. Furthermore, all cyclically deformed microstructures exhibit a certain amount of -martensite which dissents from results obtained in quasi-static tensile tests by Kisko et al. [39] on a 15Cr-9Mn-Ni-Cu stainless steel who observed -martensite solely for grains with a size equal 16

or above 3 m. The present -martensite fraction reaches a maximum of about one-third of the volume at a strain amplitude of t/2 = 0.5 %. At higher strain amplitudes, the ’martensite fraction increases at the expense of -martensite content. However, these results are a strong indicator that in the present material the nucleation of ’-martensite during fatigue takes place via the intermediate -martensite as already known for the CG state [3]. The ’-martensite fraction increases with increasing strain amplitude up to 89 % of the indicated area for t/2 = 1.2 %.

Fig. 8: Microstructure development with increasing strain amplitude. (a-e) Phase distribution of the reversion annealed (a) and cyclically deformed states (b-e), red—austenite, yellow—-martensite, blue—’-martensite, black—CI < 0.1. (f,g) Normalized grain size distribution of the austenite (f) and ’-martensite (g) (not filtered regarding CI). Furthermore, Fig. 8 illustrates the normalized grain size distributions (class width of 75 nm), i.e. the percentage of grains within a particular diameter range in relation to the total number of grains, of the austenitic (Fig. 8f) and the ’-martensite (Fig. 8g) phase for the reversion annealed and cyclically deformed states. The strain amplitude of t/2 = 0.8 % was excluded from these considerations due to inhomogeneity of the scanned area causing a distorted impression of the grain size distribution of the martensitic grains. However, it is obvious that the grain sizes of both phases follow a logarithmic normal distribution function. The mean austenitic grain size decreases with increasing strain amplitude, i.e. with a decreasing frac-

17

tion of the austenitic phase. The strain amplitude of t/2 = 0.3 % with an austenite fraction of still 84 % exhibits almost no deviation compared to the reversion annealed state whereas the differences considerably increase when higher amounts of - and ’-martensite are formed. Therefore, the ’-martensite formation seems to start in the biggest austenitic grains whereas the smallest ones exhibit the biggest resistance against phase transformation. In principle, this trend is in accordance with results obtained by other authors [37–39,58] attributing this behavior to the higher strength of the austenite, which counteracts the volume expansion accompanying the →’ phase transformation. Nevertheless, some authors [39,40] found a reversal of this trend if a certain grain size was undercut. It was suggested that this behavior stems from a change in the mechanism of ’-martensite formation, more precisely the nucleation at grain boundaries and twins instead of intersecting deformation bands as in coarser grains. Though, a contribution of the retained deformed austenite to this more intense phase transformation cannot be ruled out [39]. However, the present results do not indicate such a trend reversal, i.e. the smaller the grain size the less probable is the formation of ’martensite. On the other hand, the grain size distributions of the ’-martensite grains do not reflect the latter statement as the distribution curve for the higher strain amplitude t/2 = 1.2 % and the medium strain amplitude t/2 = 0.5 % are more or less similar. If the ’-martensite formation really starts in the biggest austenitic grains, it should lead to a pronounced difference of the distribution curves, i.e. bigger grain sizes at lower ’-martensite fractions – unless one austenitic grain would be the origin of two or more martensitic grains. Hence, to enable a well-defined statement, these grain size distributions should be further investigated which will be done by interrupted monitoring in future work.

18

Fig. 9:

t-SEM micrographs of the microstructure cyclically deformed at t /2 = 0.8 % showing individual stacking faults in austenitic grains (a,b), ’-martensite nuclei marked as I and II (c), and dislocations interacting with precipitates in martensitic grains (d,e). (a-d) Bright field images. (e) Dark field image.

Selected images of further microstructure investigations by means of t-SEM are shown in Fig. 9 for a specimen cyclically deformed at a total strain amplitude of 0.8 %. The grains in Fig. 9a and b contain individual stacking faults within an austenitic structure. Likewise to these, all observed stacking faults of this deformed state start and end at grain boundaries or other interfaces (as stacking faults), respectively. Another grain exhibiting a high density of stacking faults is shown in Fig. 9c. Moreover, an ’-martensite nucleus of very small size of about 100 nm (marked as I) seems to form inside the grain close to another nucleus (II). This condition was rarely to find in the microstructure as most grains were either still austenitic or already transformed to ’-martensite entirely. Hence, the high stresses, which occurred in the UFG state seem to enable small nuclei which rapidly propagate across the entire grain. In agreement with this finding, Kisko et al. [39] assumed a rapid growth subsequent to the nucleation at grain boundaries. Moreover, for grain sizes in the range of 0.5 to 1 m, Misra et 19

al. [73] found a preferred nucleation at slip bands which was related to a high density of dislocation tangles. The assumption that the ultrafine grains do not transform at once but via small nuclei which grow to the former austenite grain boundary is supported by TEM investigations shown in Fig. 10a-c. The bright field image of Fig. 10a shows an austenitic grain containing a high amount of stacking faults (spot II, Fig. 10c) and an area transformed to ’martensite exhibiting a bcc lattice structure (spot I, Fig. 10b). The related SAED patterns prove the Nishiyama-Wassermann orientation relationship between these two regions:

Hence, the ’-martensite originates from the austenitic grain, which was not transformed entirely. Instead, the fcc phase still remains to some extent (spot , Fig. 10c). However, there is a distinct difference compared to the ’-martensite formation in CG steels which occurs in deformation bands containing a high amount of stacking faults. Such deformation bands have not been observed in the UFG state. Rather the density of stacking faults in the whole grain seems to increase until so-called -martensite and finally the phase transformation to ’-martensite occurs. However, the presence of stacking faults in ultrafine austenitic grains as well as the formation of ’-martensite indicate a negligible influence of grain size on the apparent SFE and on the occurrence of the particular deformation mechanisms, cf. [39]. The influence of the precipitates on the dislocation movement in martensitic grains is shown in Fig. 9d and e. Obviously, the dislocations strongly interact with the carbides. Orowan bowing seems to be the dominating mechanism for passing the precipitates even if they are of small size.

20

Fig. 10: TEM investigations of the microstructure of the sample cyclically deformed at t /2 = 0.8 %. a, d) bright field images, b, c) SAED patterns of the spots marked as I and II in (a), e, f) SAED patterns of the spots marked as III and IV in (d). The bright field image in Fig. 10d illustrates a characteristically martensitic grain which is divided in two parts by a LAGB (compare the SAED patterns in Fig. 10e and f). Probably it is a result of growth of two different martensite nuclei. However, LAGBs like this are assumed to be the reason for the cyclic softening observed for high strain amplitudes after reaching a maximum of stress amplitude (Fig. 4b). According to Niendorf et al. [74,75], LAGBs in a bcc UFG interstitial-free steel are not stable under cyclic loading. The authors as well tested in total strain control and found cyclic softening for states with a high amount of LAGB [75]. This cyclic softening is regarded to be a consequence of the depletion of LAGB during cyclic straining [74]. Due to the plastic deformation, the LAGBs interact with the dislocations and disappear which leads to a lower defect density inside the respective grains. Hence, if the ’martensite fraction and the amount of LAGBs is high enough the depletion of the latter could cause the cyclic softening effect.

4. Conclusions An ultrafine-grained metastable CrMnNi steel with an initial average grain size of 0.7 m has been obtained by thermo-mechanically controlled processing via rotary swaging and annealing, and was tested in terms of tensile loading and total strain controlled fatigue. Further-

21

more, the behavior was compared to a conventionally grained counterpart. The results can be summarized as follows: 

The UFG steel exhibits a considerably higher yield strength, higher tensile strength, and a good ductility. Nevertheless, the uniform elongation is reduced as compared to the CG reference material due to a reduced strain hardening.



The cyclic stress amplitude is higher but the cyclic hardening effect caused by the ’martensite formation is less pronounced for the UFG state.



The fatigue life is higher for the UFG state than for the CG material, which at low strain amplitudes is due to the higher strength and at high strain amplitudes due to the more homogenous strain distribution between the two phases  and ’.



For all strain amplitudes t/2 ≥ 0.4 % fatigue induced ’-martensite is formed via the intermediate -martensite. The transformation does not occur simultaneously for the whole grain but small nuclei are formed which rapidly grow until the entire austenitic grain is transformed.

Acknowledgements The authors thank the German Research Foundation (DFG) for its financial support of the Collaborative Research Centre “TRIP-Matrix-Composite” (CRC 799, projects B1, B2 and B3). They also wish to thank D. Seifert (IFW Dresden) for conducting the hot forming and rotary swaging, K. Becker for preparation of the specimens for SEM investigations, and A. Leuteritz (IWW, TU Freiberg) for preparation of the TEM foils.

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Highlights  Remarkable high yield strength, tensile strength, and a good ductility. 

Higher stress amplitude but lower cyclic hardening due to ’-martensite formation.



Higher fatigue life than CG material at all tested strain amplitudes.



’-martensite nuclei form inside austenitic grains via -martensite.

27