Crack propagation behavior of solution annealed austenitic high interstitial steels

International Journal of Fatigue 79 (2015) 25–35

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Crack propagation behavior of solution annealed austenitic high interstitial steels Michael Schymura ⇑, Robert Stegemann 1, Alfons Fischer University of Duisburg-Essen, ITM, Materials Science and Engineering, Duisburg, Germany

a r t i c l e

i n f o

Article history: Received 8 January 2015 Received in revised form 17 April 2015 Accepted 26 April 2015 Available online 6 May 2015 Keywords: Austenitic high interstitial steels Stable crack propagation Planar slip

a b s t r a c t Austenitic stainless steels provide a beneficial combination of chemical and mechanical properties and have been used in a wide field of applications for over 100 years. Further improvement of the chemical and mechanical properties was achieved by alloying nitrogen. But the solubility of N within the melt is limited and can be increased in substituting Ni by Mn and melting under increased pressure. In order to avoid melting under pressure and decrease production costs, a part of N can also be substituted by C. This leads to austenitic high interstitial steels (AHIS). Within the solution annealed state strength and ductility of AHIS is comparable or even higher of those of AHNS and can be further improved by cold working. Unfortunately the endurance limit does not follow this trend as it is known from cold-worked austenitic CrNi steels. This is due to the differences of the slip behavior which is governed by the stacking fault energy as well as other near field effects. Construction components operating under cyclic loads over long periods of time cannot be considered being free of voids or even cracks. Thus the crack propagation behavior is of strong interest as well. This contribution presents the tensile, fatigue, crack propagation and fracture toughness properties of AHNS and AHIS in comparison to those of CrNi-steels. The differences are discussed in relation to microstructural characteristic as well as their alterations under cyclic loading. Ó 2015 Elsevier Ltd. All rights reserved.

1. Introduction Austenitic steels are the material of choice in a wide field of applications. The beneficial combination of strength, ductility, and corrosion resistance finds applications in biomedical, automotive, mechanical and process engineering [1–3]. In accordance with the chemical composition e.g. CrNi, CrNiMo, CrNiMoN, CrNiMnMoN as well as MnC (Hadfield type of steels), CrMnN, CrMnCN and CrMnMoN are common for austenitic steels. CrNiMoN, CrNiMoMnN, CrMnN and CrMnMoN are considered being high nitrogen steels [4] while CrMnCN represent high interstitial steels [5]. For CrNi, CrNiMo, and CrNiMoN an overview of both fatigue and crack propagation as well as the related microstructural evolution is given by [6–15]. CrNiMnMoN were investigated by [15,16] accordingly, while the fatigue behavior is mainly characterized by cyclic softening. In CrNi steels wavy slip and strain induced phase transformation from fcc to a0 -Martensite was reported ⇑ Corresponding author at: Schoellerwerk GmbH, Im Kirschseiffen 1, 53940 Hellenthal, Germany. Tel.: +49 2482 81 182; fax: +49 2482 81 109. E-mail address: [email protected] (M. Schymura). 1 Present address: BAM, Berlin, Germany. http://dx.doi.org/10.1016/j.ijfatigue.2015.04.014 0142-1123/Ó 2015 Elsevier Ltd. All rights reserved.

[17]. CrNiMo steels also showed wavy slip but only a weak tendency to strain induced phase transformation [7,15,18]. Nitrogen leads to an increase of strength resulting from solid solution hardening and an increase of ductility by the higher density of free electrons, which enhances the metallic character of the interatomic bonds [4]. Planar slip is promoted by dissolved N as a further consequence of this near-field effect [19,20]. Microstructural investigations of MnC are mainly conducted after rolling, tensile, or tribological loading. Srain induced c ? e transformation, twinning and dislocation cells were reported [21–23]. The phase transformation c ? e is prone for MnC steels with a stacking fault energy (SFE) below 20 mJ/m2 [24,25]. Similar to N-free Ni-alloyed austenitic steels the mechanical behavior of alloys with a SFE > 20 mJ/m2 is characterized by cyclic hardening within the first cycles followed by cyclic softening until fracture [25–27], while the cyclic behavior is governed by wavy slip [28]. In contrast the high nitrogen steels of CrMnN-type only show dislocation arrangements typical for planar slip within the plastic zone in front of the crack tip as reported by Kamenova et al. [29]. A more detailed investigation on the fatigue behavior of CrMnMoN was reported by [30–32]. Cyclic softening prevailed while solely planar slip and the tendency to strain induced

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Table 1 Chemical composition in wt.% and solution annealing temperature Tsa. Ni-austenite

a

Mn-austenite

Alloyed with

CrNi

CrNiMo

MnC

CrMnN

Brand name

AISI 304

AISI 316L

Hadfield steel

P 900a

P 900 Na

CrMnMoN

CrMnCN

P 2000a

CarNit

Designation

Ni0.07

NiMo0.09

GCMn1.20

NMn0.71

NMn0.90

NMnMo0.85

CNMn1.07

CNMn0.96

CNMn0.85

Tsa [°C] Al C Co Cr Cu Fe Mn Mo N Nb Ni P S Si Ti V C+N N/C C/N

1050 0.003 0.026 0.12 17.86 0.352 Bal. 1.86 0.30 0.040 0.02 8.32 0.029 0.024 0.56 0.01 0.09 0.066 1.536 0.651

1070 – 0.018 0.11 16.56 0.27 Bal. 1.75 2.04 0.076 – 10.15 0.030 0.024 0.38 – – 0.094 4.222 0.237

1050 0.034 1.200 0.01 0.10 0.02 Bal. 12.17 – 0.000 – 0.05 0.011 0.004 0.42 – 0.01 1.200 0.000 1

1085 0.013 0.086 0.02 18.16 0.03 Bal. 19.32 0.06 0.627 – 0.35 0.021 0.001 0.40 – 0.06 0.713 7.291 0.137

1085 0.015 0.065 – 18.18 – Bal. 18.93 0.04 0.830 – 0.38 0.018 0.001 0.30 – 0.06 0.895 12.769 0.078

1150 0.001 0.100 0.01 17.26 0.04 Bal. 12.30 3.03 0.750 0.01 0.14 0.017 0.008 0.81 – 0.03 0.850 7.500 0.133

1150 0.0086 0.489 – 18.82 – Bal. 18.89 0.07 0.578 – 0.41 0.020 0.002 0.43 – 0.05 1.067 1.183 0.845

1150 0.008 0.344 – 18.20 – Bal. 18.89 0.06 0.614 – 0.34 0.018 0.002 0.30 – 0.04 0.958 1.782 0.561

1150 0.007 0.260 – 18.26 0.03 Bal. 18.52 0.04 0.590 0.01 0.26 0.018 0.001 0.26 – 0.05 0.850 2.269 0.441

Brand names are registered trademarks of Energietechnik Essen GmbH, Essen, Germany.

2. Materials and methods Table 2 Compact C(T) specimen proportions. Material designation

Ni0.07 NiMo0.09 NMn0.71 NMnMo0.85 CNMn0.85 CNMn1.07 CNMn0.96 NMn0.90 GCMn1.20

2.1. Materials Proportion (mm) W

B

a0

51.0

25.4

10.2

51.0 33.2

16.8 15.3

10.2 6.6

W – specimen width, B – specimen thickness, a0 – original crack size.

twinning and c ? e transformation are reported being typical. This also holds true for CrMnN and CrMnCN steels as shown by Ref. [19,33]. In conclusion it became clear that the fatigue limit of solution annealed austenitic steels is mainly ruled by the amount of interstitially dissolved alloying elements like C and N. However, there is a lack of information on stable crack propagation behavior and the resulting microstructural changes in such steels. Thus this contribution should present the stable crack growth behavior of Ni-free N-alloyed CrMnC-steels and discuss the differences to CrNiC-as well as to N-free ones.

Nine austenitic steels from the above mentioned groups CrNi, CrNiMo, as well as MnC, CrMnN, CrMnCN and CrMnMoN with different sums and ratios of C and N were investigated. The designation was chosen with respect to the main alloying elements as well as to the sum of the interstitials C and N. The content of C, N or C + N alloying is brought about by the designation when the interstitial content of the referred species is above 0.15 wt.%. The number given in the designation represents the sum of C + N in wt.%. E.g. CNMn0.85 refers to the group of Mn-austenites with C > 0.15 wt.%, N > 0.15 wt.% and C + N = 0.85 wt.%. The chemical composition and solution annealing temperature Tsa of the investigated alloys is given in Table 1. The blanks before solution annealing were wrought bars in case of Ni0.07, NiMo0.09, NMn0.90, NMnMo0.85, CNMn0.85, CNMn0.96 and CNMn1.07, a wrought retaining ring in case of NMn0.71, and centrifugally cast tubes in case of GCMn1.20 (G  Guss, german for cast). After heat treatment the grain size, hardness, and tensile properties were measured according to [34–37] respectively. 2.2. Fatigue tests To characterize the fatigue behavior uniaxial total strain controlled (ea,t) tension–compression tests were carried out using

Fig. 1. Sequence of metallographic specimen generation after crack propagation measurements.

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the ‘‘cone head’’ sample geometry and experimental procedure described in Ref. [19]. Therefore, fatigue specimens were machined from blanks and afterward ground and polished within the measurement length solely in axial direction. All tests were run in laboratory air at room temperature either until fracture (Nf) or up to 2  106 load cycles (N). The cyclic stress–strain hysteresis were analyzed as to ea,el and ea,pl, using an in-house developed MatLabÒ-routine, by means of the classical equations of Basquin [38] and Manson–Coffin [39,40], which can be summed up to

ea;t ¼ ea;el þ ea;pl ¼ r0f =Eð2Nf Þb þ e0f ð2Nf Þc

(182 Keithley Instruments Inc., Cleveland, OH, USA) was used to measure the resulting potential. The loading clevis as specified in [44] was modified in order to validate the potential measurement. Electrically isolated cooling adapters were attached between the load frame and the loading clevis in order to ensure that the CT-specimen are free from external currents and to maintain a DT 6 2 K. The stable crack propagation experiments were halted at a crack length of aQ = a0 + (W  a0)/2 followed by fracture toughness tests acc. to [44]. 2.4. Metallography and microscopy

with r0 f to be the true fracture strength and e0 f to equal to ln(1/(1  Z)) as reported by Morrow [41] and Mitchell [42]. 2.3. Crack propagation measurements Stable crack propagation measurements were carried out by means of CT-specimens as described in [43,44] and Table 2. The specimen were machined from cuboid-shaped, solution annealed, and quenched blanks of the dimensions (1.2W + 2 mm)  (1.25W + 2 mm)  (B + 2 mm), while the notch was generated by electro discharge machining (EDM). Force controlled sine-wave loading with the frequency of f = 10 Hz at Rr = ru/ro = 0.1 was applied by means of a commercial servo hydraulic test rig (load frame, hydro pulse actor PLm 100 K and load cell PM 100 Rn: Carl Schenck Maschinenfabrik GmbH, Darmstadt, Germany; servo controller FlexTest40 and controller software MTS 793.10: MTS Systems Corporation, Eden Prairie, MN, USA). The direct current potential method, as described in [43,45,46], was used to continuously measure the crack length. A current of 40 A was applied by a power source (SM 1540 D Delta Elektronika B.V., Zierikzee, The Netherlands) while a nV-amplifier

Longitudinal cross sections were prepared according to Fig. 1. After wet cutting (Accutom-50, Struers, Willich, Germany) the samples were embedded (Epovit, Buehler, Düsseldorf, Germany), ground by means of SiC-abrasive down to 1200 mesh size (Struers, Willich, Germany) followed by polishing down to 1 lm grain size in polycrystalline diamond suspension (ATM GmbH, Memmelzen, Germany). Etching in 50 ml HCl, 50 ml aqua dest., 4 g KCl for 20 s finalized the preparation for light optical microscopy (LOM) (Microscope: BX-10, Olympus, Düsseldorf, Germany; Camera: DFC 420, Leica, Heerbrugg, Schweiz; Documentation Software: ImageAccess Standard 7, Imagic Bildverarbeitung AG, Glattbrugg, Schweiz) as well as hardness measurements. In order to gain sufficient quality for electron backscatter diffraction analyses (EBSD) 0.05 lm colloidal silica suspension (Masterprep Polishing Suspension, Buehler, Düsseldorf, Germany), vibratory polishing (Vibromet, Buehler, Düsseldorf, Germany), and finally electro-polishing in A2 (700 ml ethanol, 110 ml aqua dest., 100 ml Butyldiglycol, 78 ml perchloric acid) at 25 °C and 25 V for 35 s (LectroPol-5, Struers, Willich, Germany) were applied. LOM and scanning electron microscopy (SEM) were used for

Fig. 2. Selected LOM images prior to mechanical testing.

Table 3 Grain size, hardness and tensile properties of the investigated steels in solution annealed state. Ni-austenite

Mn-austenite

Designation

Ni0.07

NiMo0.09

GCMn1.20

NMn0.71

NMn0.90

NMnMo0.85

CNMn1.07

CNMn0.96

CNMn0.85

GS (lm) Hardness [HV10]

62 ± 43 153 ± 5

45 ± 36 162 ± 11

460 ± 363 175 ± 7

68 ± 46 278 ± 14

37 ± 32 267 ± 9

137 ± 91 271 ± 7

133 ± 84 278 ± 13

60 ± 35 271 ± 9

68 ± 47 270 ± 4

Rp0.2 (MPa) Rm (MPa) Z (%) A (%) Ag (%) E (GPa) Ws (J/mm3)

258 683 77 89 70 194 504

283 609 83 76 53 193 406

334 804 33 43 42 202 266

501 875 76 80 54 198 624

611 995 74 70 44 194 630

633 1022 75 77 51 194 701

582 1044 64 76 61 194 711

599 1028 71 76 54 197 693

586 1003 69 74 52 195 669

SFE (mJ/m2) DF (1022/cm3)

20 N.a.

45 N.a.

41–68 Not measurable

20 1.1

35–45 0.6

30–50 1.0

43–59 2.0

27–38 2.9

36–43 2.8

GS – grain size, Rp0.2 – yield strength, Rm – tensile strength, Z – reduction of area to fracture, A – elongation to fracture, Ag – uniform elongation, E – Young0 s modulus, Ws – specific tensile fracture energy, SFE – stacking fault energy, DF – density of free electrons. SFE values are taken or computed from [56,60,61] computed values are marked by ; DF values of CNMn0.85, CNMn0.96 and CNMn1.07 are taken from [56]; DF values of NMn0.71, NMn0.90 and NMnMo0.85 are estimated from [5,56] and marked by .

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fractography as well as microstructural characterization before and after mechanical testing. SEM was performed using a LEO 1530 GEMINI (Carl Zeiss Microscopy GmbH, Munich, Germany) field emission gun supplemented with a combined microanalysis system (AMETEK – EDAX, Ametek GmbH, Wiesbaden, Germany) consisting of the Apollo X Silicon drift detector with TEAM 3.11 software for energy dispersive spectroscopy (EDS) and a Digiview

IV CCD camera with OIM 6.2 data collection and analysis software for EBSD. Hardness measurements and EBSD measurements were carried out before as well as after the stable crack propagation experiments. The positions for such analyses were chosen as to the calculated radii of the plastic zone rpl being 5 mm and 10 mm, respectively. Hardness was measured using a load of 98.07 N,

(a) Pole figure maps; measurement resolution is 2 µm

(b) Kernel average missorientation map; measurement resolution is 2 µm, missorientation is angle 5°, number of compared neighbours is 37 (3 rings) Fig. 3. EBSD analyses of selected steels prior to mechanical testing.

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3.2. Tensile properties The tensile properties of the investigated steels are given in Table 3. Yield strength Rp0.2 and tensile strength Rm, increase with C + N up to about 1.1 C + N wt.% (Fig. 4). However, GCMn1.20 exclusively contains carbon and no N and shows the biggest grain size. Reduction of area at fracture Z decreases linearly with increasing C-content. In spite of this specific tensile fracture energies higher than 650 J/mm3 were measured for NMnMo0.85 and the CrMnCN group of steels, which is in good accordance to Ref. [47]. 3.3. Strain controlled fatigue tests Fig. 4. Influence of interstitial content on the tensile properties of the investigated steels.

which equals HV10. EBSD analyses contain kernel average misorientation (KAM) as well inverse pole figure (IPF) maps. IPF maps were analyzed with the loading direction being parallel to ‘‘RD – rolling direction’’. Different KAM maps were generated as follows: The EDAX system works with a hexagonal data collection grid, hence for KAM the orientation of a center point to the next neighbor points located on a surrounding ring with the resolution as radius is compared. The shown KAM angle was chosen to be 5° (designation is KAM 5°). The number of surrounding rings (1, 2, 3, . . ., 10) was chosen on purpose (designation is KAM 5° XR, where X equals to the number of considered surrounding rings). It has to be mentioned that a reasonable comparison of two KAM maps measured with different resolutions is possible only, if the considered volume remains constant. Thus KAM 5° 1R measurements with a resolution of 2 lm can only be compared quantitatively with KAM 5° 10R at a resolution of 0.2 lm. 3. Results 3.1. Microstructure before mechanical testing The initial microstructure shown in Fig. 2 is typical for solution annealed austenitic steels. The grain size differed markedly depending on the various production routes. The hardness values are higher than 250 HV in case of all CrMn- and between 150 and 180 HV10 for the CrNi-austenites and GCMn1.20 (Table 3). EBSD analyses of the solution annealed state revealed no crystallographic anisotropy. The kernel average misorientation values (KAM) do not exceed 1° (Fig. 3) proving the microstructure being technically free of any residual strains after solution annealing and before mechanical testing.

In addition to Ref. [19] here the investigated number of steels is extended by Ni0.07, NiMo0.09, NMn0.90 and NMnMo0.85. The N-free CrNi(Mo) steels as well as GCMn1.20 show cyclic hardening within the first load cycles followed by cyclic softening until fracture at total strain amplitudes ea,t > 0.3%. NiMo0.09 and GCMn1.20 depict cyclic softening at ea,t < 0.3% and ea,t < 0.15%, respectively, while Ni0.07 reveals cyclic hardening at all ea,t. The N-containing CrMn(MoC)N austenitic steels demonstrate cyclic softening after the first load cycle until fracture. All fatigue results are in good agreement with previous studies [7,19,26,30,48]. The endurance limit rD, the corresponding total strain amplitude ea,t(rD), and the best fit coefficients of the Manson–Coffin type of analyses are given in Table 4 for completeness. Fig. 5 shows the influence of C + N on the endurance limit of stainless steels. In [19,33] the maximum of the fatigue limit was found at about 0.96 C + N wt.%, which is now further supported by NMn0.90 and NMnMo0.85. Thus rD does not depend on the N/C-ratio. 3.4. Crack propagation measurements The stable crack propagation data were evaluated by means of the so-called seven point increment polynomial technique described in [43] while the fracture toughness values were derived according to [44]. The fundamental principles for such analyses are the definitions of the stress intensity factor K given by Irwin [49] and the stress intensity factor range DK by Rice [50] as follows:

K ¼ rN

pffiffiffiffiffiffi pffiffiffiffiffiffi paYða=WÞ and DK ¼ DrN paYða=WÞ

with rN – the nominal (net-section) stress, DrN – the nominal (net-section) stress range during cyclic loading, a – the crack length, W – the sample width and Y – the stress intensity factor coefficient for CT-specimen as given by Newman and Srawley [51,52]. With respect to the sample geometry and the mechanical properties of the investigated steels a valid KIc fracture toughness could not be measured. Fig. 6 shows the resulting crack growth curves and the preliminary fracture toughness (points marked by triangles),

Table 4 Fatigue data and best fit coefficients of Manson–Coffin analyses. Ni-austenite

Mn-austenite

Designation

Ni0.07

NiMo0.09

GCMn1.20

NMn0.71

NMn0.90

NMnMo0.85

CNMn1.07

CNMn0.96

CNMn0.85

rD (MPa) ea,t (rD) (%)

190 0.1

173 0.09

220 0.12

290 0.15

389 0.20

347 0.18

313 0.16

390 0.20

320 0.17

917 0.085 0.4 0.67

2299 0.155 1.418 0.6

2416 0.135 1.359 0.6

2453 0.135 1.37 0.62

2423 0.15 1.03 0.58

2520 0.15 1.245 0.6

2318 0.14 1.363 0.61

Best fit coefficients of Manson-Coffin analyses 1932 2242 b 0.145 0.165 0 ef 1.474 1.784 c 0.74 0.65

r0 f (MPa)

rD – endurance limit, ea,t(rD) – total strain amplitude at endurance limit, r0 f – fatigue strength coefficient, b – fatigue strength exponent, e0 f – fatigue ductility coefficient, c – fatigue ductility exponent.

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Fig. 5. Axial endurance limit vs. the sum of interstitials. data taken from ⁄[7], ⁄⁄[30], ⁄⁄⁄ [15,16].

Fig. 6. Fatigue crack growth curves da/dN–DK of the investigated materials.

Fig. 8. m – Paris exponent vs. the sum of interstitials.

Fig. 9. Reciprocal Paris exponent vs. endurance limit of the investigated steels.

Fig. 7. Selected SEM micrographs of the forced fracture surfaces of the investigated steels.

Table 5 Crack propagation, fracture properties and best fit coefficients of Paris analyses. Ni-austenite Designation p KQ (MPa m) p DKQ (MPa m)

Mn-austenite

Ni0.07

NiMo0.09

GCMn1.20

NMn0.71

NMn0.90

NMnMo0.85

CNMn1.07

CNMn0.96

CNMn0.85

64 57.6

61 54.9

58 52.2

84 75.6

45 40.5

142 127.8

96 86.4

78 70.2

96 86.4

4.27 2.73  1014

7.58 5.69  1019

3.49 3.83  1014

3.64 3.18  1014

3.43 5.63  1014

3.52 2.77  1014

3.83 1.65  1014

3.34 7.48  1014

Best fit coefficients of Paris analyses m 5.96 C 1.46  1016

p KQ – preliminary fracture toughness, DKQ – preliminary cyclic fracture toughness at R = ru/r o = 0.1, C – Paris coefficient (corresponds to da/dN at DK = 1 [MPa m]), m – Paris exponent (corresponds to the slope of the da/dN–DK – curve within regime II).

M. Schymura et al. / International Journal of Fatigue 79 (2015) 25–35

DKQ = Rr  KQ. For the main focus of this paper DKQ is given but will not be discussed. Still it should be mentioned that all steels investigated showed ductile behavior as expected from the tensile data and the measured DKQ-values (Fig. 7). In general the N-alloyed, Ni-free CrMn(MoC)N steels display crack propagation rates being about one order of magnitude smaller compared to the CrNi(Mo) steels as well as to GCMn1.20. Hence, the stable crack propagation (regime II) behavior was further analyzed with respect to the Paris–Erdogan equation [53]

da=dN ¼ CðDKÞm with da/dN – crack propagation per load cycle, C – Paris coefficient and m – Paris exponent, respectively. The results are given in Table 5. The Paris exponent m, representing the slope of the crack growth curves, decreases with the sum of interstitials, reaching a minimum at about 0.85 C + N wt.% (Fig. 8). In case of the investigated CrMn(MoC)N steels m ranges only from 3.3 to 3.8 while it varies between 4.2 and 7.6 for CrNi(Mo) and GCMn1.20. Any correlation between the Paris coefficient is evident neither for C + N nor for N/C.

strain. Since under stable crack growth a plastic zone is generated at the propagating crack tip such differences of the slip characteristics at high plastic strain amplitudes should also influence the stable crack growth. Thus in order to be able to compare the stable crack propagation behavior certain positions along the stable fatigue crack must be chosen, which are at least theoretically characterized by similar plastic strains. The theoretical size of the plastic zone was therefore calculated acc. to rpl = (DK/(Rr  Rp0.2))2. Then SEM micrographs were taken at certain rpl-values. Here Fig. 10 shows characteristic fatigue crack surfaces at about rpl  5 mm, which reveals distinct differences as to the stable crack propagation appearances. The transcrystalline fracture surfaces have a smooth appearance in case of the CrNi(Mo) steels as well GCMn1.20, while there is a much rougher surface visible in case of the CrMn(MoC)N steels. Not surprisingly the first two represent those steels with wavy-slip while the latter belong to the planar-slip group. Wavy slip is characterized by the generation of dislocation cells under fatigue loading [7,54]. Dislocations are

4. Discussion In order to evaluate the properties of metals with and without cracks one can relate the reciprocal 1/m to the fatigue limit as given in Fig. 9. According to this the N-alloyed, Ni-free steels perform better as to both criteria compared to the Ni-alloyed or the N-free ones. The most favorable combination of both the reciprocal of m and rD lies between 0.85 and 0.96 C + N wt.%. Obviously the earlier reported [19,20] slip characteristics as well as the amount of interstitially dissolved alloying elements do not only affect the fatigue limit but also have a marked influence on the crack propagation. From the authors’ earlier work it has been shown that the Ni-alloyed and N-free steels show wavy slip at high cyclic plastic strain amplitudes while the Ni-free and N-alloyed ones always provided planar slip regardless of the plastic fraction of the cyclic

(a) wavy slip (NiMo0.09 and GC1.20)

31

Fig. 11. Selected hardness profiles of the investigate steels.

(b) planar slip (CNMn1.07 and NMn0.71

Fig. 10. Selected SEM micrographs of the fatigue fracture surfaces of the investigated steels.

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(a) NiMo0.09, rpl = 5 mm

(b) NiMo0.09, rpl = 10 mm Fig. 12. Selected EBSD mappings of NiMo0.09 below the fracture surfaces of the investigated steels (mapping orientation: bottom – fracture surface, top – bulk; IPF maps are given for easier localization of the KAM 5° 10R maps within the KAM 5° 3R and 1R overview).

M. Schymura et al. / International Journal of Fatigue 79 (2015) 25–35

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(a) NMnMo0.85, rpl = 5 mm

(b) NMnMo0.85, rpl = 10 mm Fig. 13. Selected EBSD mappings of NMnMo0.85 below the fracture surfaces of the investigated steels (mapping orientation: bottom – fracture surface, top – bulk; IPF maps are given for easier localization of the KAM 5° 10R maps within the KAM 5° 3R and 1R overview).

blocked with such cell structures, which allows for transgranular crack propagation along such cell walls. This leads to the smooth appearance of the fatigue fracture surface. In contrast to this the

propagating fatigue cracks are limited to distinct slip planes – as recently observed in [55] as well – for solely planar sliding materials. Due to the fact that in addition these dislocations remain

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mobile, more energy is needed for crack propagation. Thus several synergistic effects lead to smaller crack propagation rates in CrMn(MoC)N steels. With increasing C + N-content an increase of interatomic short range order effects is observed as well, hindering dislocation sliding [4,5]. In [56] the maximum effects within the CrMnCN group was found at 0.85 C + N wt.%. With such increasing short range order the necessary energy and, therefore, the shear stress to initiate and propagate dislocation sliding increases [28]. This shear stress also depends on the orientation relation between the slip plane and the loading direction and hence on the grain orientation of each affected grain. Finally the limitation to certain sliding planes forces the crack tip to change orientation (deflection) even within one grain, which further increases the energy needed to propagate over a certain length. Thus the Ni-free, N-alloyed steels gain the lower crack propagation rate by the higher content of interstitials C + N which, (a) distinctly increases the effect of dislocation–interstitial interaction, (b) in parallel provides a more metallic character of the interatomic bonds at intermediate N/C-ratios and, therefore, (c) does only allow for planar slip. It should be mentioned here that all steels investigated are in about the same range of stacking fault energy (Table 3), which therefore does not work as predominant – still not negligible – criterion as has been discussed earlier [19,20,57]. Crack propagation leads to cyclic plastic deformation of the microstructure within the plastic zone, which size gives a theoretical measure of the affected volume at a certain crack length and load a.k.a. stress intensity range value. Certainly cyclic plastic deformation is characterized by an increase of the lattice-defect density close to the crack surface. In addition this should be more distinct at rpl  10 mm if compared to rpl  5 mm. Since any former deformation increases the defect density it should lead to distinct work-hardening effect of such steels if e.g. the gradient of hardness values is measured at such distance from the crack surface. Fig. 11 shows such hardness profiles of the investigated materials, measured as shown in Fig. 1. There is no hardness gradient visible at all within the standard deviation of the reference measurements of the undeformed bulk. Thus either there was no deformation, or it was so weak that no distinct strain hardening was measurable, or the size of the plastic zone is smaller than 0.5 mm. Thus a much more sensitive method for measuring the residual strains within a crystalline solid solution like e.g. EBSD and the local misorientation within lattices [58] should clarify that. The measures of the kernel average misorientation in Figs. 12 and 13 indicate that for no steel investigated the plastic strains are homogenously distributed. A distinct localization can be seen at internal interfaces like grain and phase boundaries within the theoretically affected zone. Still the largest strains appear directly at the fatigue crack surface – always at the bottom of each KAM figure – but within only some lm and not exceeding the grain size. Obviously this localization is more distinct for the CrMn(MoC)N steels while a more homogenous distribution is observed for the CrNi(Mo) steels and GCMn1.20 (compare KAM 5° 3R and KAM 5° 1R). This again appears to be based on the differences in slip behavior. Accumulation of dislocations within cell walls hindering and inhibiting dislocation movement result in higher residual strains within the lattice. Planar slip allows for such highly localized plastic strain but with still mobile dislocations. Furthermore the possibility for the extinction of inverse dislocations on the same sliding plane is given [59]. Thus despite its more distinct localization of strains in the vicinity of the fatigue fracture face planar slip results in a smaller crack propagation rate. Furthermore no explicit difference is observed between rpl  5 mm and rpl  10 mm. The volume

adjacent to the crack surface showing a higher KAM angle is about 30 lm for the CrMn(MoC)N steels (KAM 5° 10R in Fig. 13), while being about 60 lm for CrNi(Mo) steels and GCMn1.20 (KAM 5° 10R in Fig. 12). Twinning as could be expected from fatigue investigations at high cyclic plastics strain amplitudes >0.5% [19] is not observed. The deformation state within the high KAM areas close to the fatigue crack surface is rather comparable to the deformed states of samples after loading at strain amplitudes between 0.2% < ea,t < 0.5%. It should be mentioned here that the fatigue limit is reached at ea,t 6 0.2%. Thus stable crack propagation obviously takes place slightly above the fatigue limit of strain-controlled fatigue tests. High resolution microstructural analyses by means of TEM should give a deeper insight into this and will be carried out in the future.

5. Conclusions – CrMn(MoC)N-alloyed austenitic steels provide lower crack propagation rates compared to conventional CrNi(Mo)-alloyed austenitic stainless steels. – The increase of interatomic short range order brought by the high content of interstitials, N or C + N, and forcing planar slip contributes to the improved crack propagation rates. – The lowest crack propagation rates were observed at C + N = 0.85 wt.% being independent of N/C-ratio.

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