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Micro-mechanisms of microstructural damage due to low cycle fatigue in CoCuFeMnNi high entropy alloy
International Journal of Fatigue 130 (2020) 105258
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Micro-mechanisms of microstructural damage due to low cycle fatigue in CoCuFeMnNi high entropy alloy Fateh Bahadur, Krishanu Biswas, N.P. Gurao
T
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Department of Materials Science and Engineering, IIT Kanpur, Kanpur 208016, India
ARTICLE INFO
ABSTRACT
Keywords: Advanced materials Low cycle fatigue Microstructures Dislocations Damage
Low cycle fatigue behavior of equiatomic CoCuFeMnNi high entropy alloy was investigated under fully reversible strain control mode followed with detailed microstructural characterization using electron back scatter diffraction. There is an increase in the intragranular misorientation and geometrically necessary dislocation density with increase in the strain amplitude that leads to accumulation of damage and intergranular cracking. The interaction of dislocations with copper rich nano-clusters, solute environment and grain boundaries affect slip reversibility, thereby deciding cyclic deformation behaviour. Annealing twin boundaries are resistant to damage and increasing their population density by grain boundary engineering can improve performance of CoCuFeMnNi alloy.
1. Introduction Design and development of new alloys for achieving excellent mechanical properties like high specific strength, optimum combination of strength and ductility, excellent fracture toughness at low temperature, superior fatigue and creep resistance has been pursued to cater the stringent demands of the engineering industries in the recent past. Recently, a new class of materials known as high entropy alloys (HEAs) or complex concentrated alloys (CCAs) that are multi-principal multicomponent alloys containing five or more elements in equal or near equal proportions have been developed [1,2]. Their high configurational entropy of mixing favors the formation of solid solution over brittle intermetallic phases at high temperature [1–4]. This innovative alloy design strategy overcomes the conventional approach in which properties are primarily governed by principal element, thus, providing enormous opportunities in terms of physical and functional properties by permitting extensive and flexible tuning of alloy composition. The past decade of dedicated research revealed that many HEAs possess unique properties in comparison with the conventional alloys for instance, high thermal and microstructural stability [5,6], high hardness [6], high strength at elevated temperature [7,8], superior fracture toughness [9], and excellent wear resistance [10]. Among HEAs, equiatomic CoCrFeMnNi Face Centered Cubic (FCC) HEA (Cantor alloy) exhibits high tensile strength (1130 MPa), sufficient ductility (65%), and higher fracture toughness (200 MPa.m1/2 ); especially at cryogenic temperatures due to the presence of deformation twins and has been studied in great details [4,9,11–13].
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Most of these studies regarding mechanical behavior are usually investigated in uniaxial loading, [14–19] which are necessary but still insufficient to broaden the application range of HEAs for structural applications. Most engineering applications involve load reversal and cyclic loading in service, and therefore, it is essential to study them under cyclic loading. In particular, low cycle fatigue behavior of metals and alloys is important from technological perspective as well as from a scientific point of view to investigate the accumulation of damage wherein dislocations and their interaction with each other and other defects play an important role. Low cycle fatigue involves plastic strain and is generally investigated in strain control mode. Deformation behaviour and failure of different engineering materials like IF steel [20], dual phase steel, low alloy steel, TWIP/TRIP steel [21–24], stainless steel [25–27], commercially pure titanium, titanium alloy [28,29], and Ni-based superalloys [30,31] has been well-established under low cyclic fatigue regime. However, the research on fatigue behavior, particularly pertaining to low cycle fatigue of high entropy alloys, has been limited. Hemphill et al. [32] have examined the high cycle fatigue behavior of as-cast and wrought Al0.5CoCrCuFeNi alloy using four-point bending test at room temperature and found that the fatigue endurance limit lies in the range of 540–945 MPa. A similar composition of Al0.5CoCrCuFeNi alloy with high purity elements has also been studied by Tang et al. [33] who showed that the use of high purity elements results in better fatigue performance. This was attributed to the formation of nano-twins during cyclic loading in the alloy containing high purity elements leading to higher endurance limit and prolonged fatigue life compared to many
Corresponding author. E-mail address: [email protected] (N.P. Gurao).
https://doi.org/10.1016/j.ijfatigue.2019.105258 Received 26 June 2019; Received in revised form 12 August 2019; Accepted 28 August 2019 Available online 29 August 2019 0142-1123/ © 2019 Elsevier Ltd. All rights reserved.
International Journal of Fatigue 130 (2020) 105258
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Nomenclature
R APT IPF GROD KAM GND
List of symbols and acronyms Δε/2 c n
strain amplitude fatigue ductility exponent strain hardening co-efficient
conventional alloys. Niendorf et al. [34] reported the low cycle fatigue behavior of Fe50Mn30Co10Cr10 alloy showing remarkable changes in the deformation mechanisms under monotonic and cyclic loading conditions. The alloy showed poor strain hardening in tension but occurrence of strain-induced martensitic transformation in low cycle fatigue regime contributed to better fatigue performance. This peculiar behavior was attributed to partial deformation reversibility and the presence of the planar nature of slip. In other investigations, microstructural engineering approaches like grain refinement were employed to improve fatigue performance [35,36]. Multi-phase high entropy alloys investigated till date show remarkable fatigue strength compared to the conventional alloy systems. But the fundamental understanding of the fatigue deformation mechanism of single phase high entropy alloys, especially under low cycle fatigue regime has not been studied yet. Accordingly, the present study focuses on the low cycle fatigue behavior of single phase FCC CoCuFeMnNi HEA system developed by Tazuddin et al. [37], under strain control conditions. It was shown that operation of planar partial slip along with conventional octahedral slip led to the unique Goss-Brass type texture evolution during cold rolling of the alloy [38]. An atom probe tomography study by Sonkusare et al. [39] revealed the presence of copper rich nano-clusters (2.5 nm average
strain ratio atom probe tomography inverse pole figure grain reference orientation deviation kernel average misorentation geometrically necessary dislocation
size) in the FCC matrix of the alloy. Other investigations on the alloy have established the importance of the copper rich nano-clusters and their interactions with dislocations in tension and high pressure torsion in monotonic and cyclic mode [40,41]. Thus, the deformation behavior of the CoCuFeMnNi high entropy alloy is studied for different loading conditions. However, the cyclic deformation of the alloy is still unexplored. The present investigation deals with understanding the deformation mechanism and microstructural evolution of CoCuFeMnNi alloy subjected to strain controlled low cycle fatigue using various characterization tools like scanning electron microscopy (SEM), electron back scatter diffraction (EBSD) and back scattered electron imaging showing electron channeling contrast. 2. Experimental procedure 2.1. Material and methods Equiatomic CoCuFeMnNi alloy was prepared by arc melting buttons of ~12 g of equimolar mixture of elements with greater than 99.9% purity obtained from Sigma Aldrich. Alloy buttons were remelted at least four times to obtain compositional homogeneity. Four such alloy
Fig. 1. (a) X-ray diffraction pattern, (b) normal direction inverse pole figure map and (c) elemental distribution map from atom probe tomography of the homogenized CoCuFeMnNi HEA sample. 2
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buttons were remelted again and suction cast into a billet of dimension 10 × 10 × 60 mm3. The billet was sealed in a quartz tube with argon and homogenized at 1273 K for 24 h followed by water quenching. Subsequently, the homogenized billet was cold rolled to 50% reduction in thickness at room temperature. The rolled specimen was subjected to annealing for 2 h at 1273 K to obtain a fully crystallized microstructure. Tensile and fatigue samples were machined from the annealed billet by electric discharge machining according to the ASTM E8 (10 mm × 2.5 mm × 2.5 mm) and ASTM E-606 (8 mm × 5 mm × 2.5 mm) standards respectively.
diffraction (EBSD, Oxford Instruments, U.K.) detector. EBSD scans were acquired with Aztec HKL software. The post-processing of the data was performed using TSL-OIM version 7.2 software. The EBSD samples were prepared by polishing to 2000 grit size using emery paper and then cloth polishing using the 0.5 μm alumina suspension followed with vibromet polishing (Buehler, USA) using colloidal silica suspension with a particle size of 50 nm. Backscattered electron imaging was carried out to obtain electron channeling contrast [44] by tilting the sample 10 degree at a working distance of 10 mm and rotating the sample to reveal the dislocation structure near the fracture tip of the fatigue tested samples.
2.2. Tensile and fatigue test
3. Results
The uniaxial tensile test was carried out at room temperature with a strain rate of 1 × 10 3s 1 using Instron-1195 Universal Testing Machine. Low cycle fatigue tests were also performed under straincontrol conditions at different strain amplitude (Δε/2) of 0.5%, 0.75%, 1% and 1.5% with fully reversible cyclic strain condition (R = −1) up to failure. All fatigue tests were conducted with the strain rate of 5 × 10−3 s−1 using the 100 kN Universal Testing Machine (BiSS, Bangalore, India).
3.1. Initial material X-ray diffraction (XRD) pattern of the homogenized CoCuFeMnNi alloy shown in Fig. 1(a) indicates the presence of FCC single phase with a lattice parameter of 0.361 nm. The normal direction (ND) inverse pole figure (IPF) map of the investigated HEA in Fig. 1(b) shows a fully recrystallized microstructure with annealing twins and an area average grain size of ~47 µm. Atom probe tomography carried out on the homogenized sample provided elemental distribution at the atomic length scale. Fig. 1(c) shows the elemental distribution from atom probe tomography indicating the homogeneous distribution of cobalt, iron, manganese, and nickel while there is heterogeneous distribution of copper, characterized by the presence of near spherical copper rich nano-clusters with an average diameter of 2.5 nm.
2.3. X-ray diffraction and microstructure analyses The crystal structure of the homogenized HEA was examined by Xray diffraction (PANalytical, Netherland), with a step size of 0.02 degree/ s, operated at 45 kV and 40 mA using CuKα (λ = 0.15402 nm) radiation. The grain morphology, fractography and compositional analyses of the synthesized sample were carried out utilizing field emission scanning electron microscope (FE-SEM) (JSM-7100F, JEOL, Japan) equipped with energy dispersive spectroscopy facility (EDS, Oxford Instruments, UK) operated at 20 kV. Atom probe tomography (APT) was carried out using the local electrode atom probe (LEAP 4000X HR™, Cameca Instruments) to obtain near-atomic scale elemental distribution. APT samples were prepared by using standard procedure [42,43] using a dual-beam focused ion beam (FEI Helios Nanolab 660, USA). A volume of 60 × 60 × 100 nm3 was acquired, through a pulse repetition rate of 250 kHz and pulse energy of 30 pJ. The tip temperature was maintained at around 60 K. The APT data was evaluated with the help of IVAS 3.6.10a software provided by Cameca Instruments.
3.2. Tensile and fatigue tests Engineering stress - engineering strain behaviour of the CoCuFeMnNi alloy as shown in Fig. 2(a) indicates that the investigated alloy exhibits yield strength of 324 ± 20 MPa and ultimate tensile strength of 710 ± 35 MPa. Cyclic tests performed with fully reversible strain amplitude (Δε/2) of 0.5%, 0.75%, 1%, and 1.5% showed a unique evolution of stress amplitude with number of cycles that is depicted in Fig. 2(b). The stress amplitude at the onset of cyclic loading increases with an increase in the strain amplitude. However, the evolution of the stress amplitude with number of cycles is different for 0.5% and 0.75% (low) strain amplitude compared to the 1% and 1.5% (high) strain amplitude with the former two showing a significantly higher number of cycles to failure than the latter two. There is initial hardening; that is an increase in stress amplitude with the number of cycles for all the four samples. However, 0.5% and 0.75% strain amplitude samples show softening followed with fluctuations in the stress amplitude. In
2.4. Electron back scatter diffraction (EBSD) and electron channeling contrast imaging Electron back scatter diffraction (EBSD) data acquisition was carried out in FE-SEM from JEOL equipped with electron back scatter
Fig. 2. (a) Engineering stress – engineering strain diagram and (b) variation of stress amplitude with number of cycles in low cycle fatigue test at different strain amplitude for CoCuFeMnNi HEA. 3
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necessary dislocation density (GND) were determined from the EBSD scan and are listed in Table 1. It is observed that there is an increase in the average KAM and GAM value with increase in the strain amplitude. Similarly, the geometrically necessary dislocation density is also found to increase with the increase in strain amplitude. Fig. 3(e) shows that the kernel average misorientation (KAM) distribution shifts to the right and becomes broader as strain amplitude increases, indicating an increase in dislocation density and more heterogeneous distribution of the dislocations with increasing strain amplitude [45]. 3.4. Fractography Detailed analysis of the fractured surface of the fatigue tested samples was carried out using a scanning electron microscope. The illustrative SEM fractographs at lowest (0.5%) and highest (1.5%) strain amplitude are shown in Fig. 4(a–c) and Fig. 4(d–f) respectively. Fig. 4(a) shows formation of striation marks with different orientation observed within different grains. Fig. 4(b) and (c) display crack propagation path with higher number of cycles, which are found to be mostly intergranular in nature. Fig. 4(d) shows the presence of void or inhomogeneity in the material that can act as local stress raisers at grain boundaries, where crack initiation will take place. Once the cracks are initiated the stable propagation of the crack with each cycle is a clear representation of striation mark as shown in Fig. 4(e). However, as shown in Fig. 4(f), the spacing between the striation marks was found to increase with strain amplitude. The increase of strain amplitude causes faster crack growth rate. . Therefore, the mechanism of fatigue failure in the HEA is similar to that in conventional alloys based on crack initiation, stable crack propagation, and final fracture [46]. 4. Discussion
Fig. 3. Loading axis inverse pole figure map near the fracture tip of samples tested in fatigue at strain amplitude of (a) 0.5% (b) 0.75% (c) 1% (d) 1.5% and (e) Kernel average misorientation distribution from all the fatigue tested samples.
4.1. Cyclic stress response The variation of stress amplitude with strain amplitude shows two distinct behavior for the equiatomic CoCuFeMnNi alloy in the present investigation. The stress amplitude reaches a maximum and fluctuates over a prolonged cyclic stage and then decreases sharply for strain amplitude of 0.5% and 0.75%. However, at higher strain amplitudes of 1.0% and 1.5%, value of stress amplitude increases but the regime of high-stress amplitude is very small, and the sample fails after few cycles compared to low strain amplitude samples. Chen et al. [47] have shown that long range motion of dislocations at higher strain amplitude shows a distinct cyclic hardening/softening behaviour compared to short range motion of dislocations at low strain amplitude in Hadfield steel. Agrawal et al. [40] have reported four factors responsible for the strain hardening behavior of the CoCuFeMnNi HEA under monotonic condition, namely; (i) dislocation multiplication, (ii) grain boundary strengthening, (iii) solid solution strengthening and (iv) interaction between dislocations and copper rich nano-clusters. Cyclic loading leads to forward and backward movement of dislocations characterized by long range motion at higher strain amplitude and short-range motion at low strain amplitude. All the aforementioned four interactions play an important role in the reversible nature of dislocation motion in the CoCuFeMnNi high entropy alloy. Grain boundaries act as barrier for the motion of dislocations leading to formation of dislocation pile-up at grain boundaries. During reverse loading, the dislocations can move in the reverse direction at lower stress and annihilate each other, contributing to cyclic softening. This happens for materials with low strain hardening exponent (n ≤ 0.15), while for materials with higher strain hardening co-efficient (n > 0.15) dislocation accumulation occur, leading to cyclic hardening [48]. The alloy under investigation shows higher strain hardening in tension [38] and is a cyclic hardening material at all the strain amplitudes. Both solute atmosphere and copper rich nanoclusters are likely to pin dislocations and reduce the mobility of dislocation. This, in turn, can cause an increase in dislocation density leading to accumulation of intragranular misorientation that can cause accumulation
Table 1 Variation of average kernel average misorientation (KAM), grain average misorientation (GAM) and geometrically necessary dislocation (GND) density for the low cycle fatigue tested samples at different strain amplitude. Strain amplitude
Average KAM°
Average GAM°
Average GND density × 1013 (m−2)
0.5% 0.75% 1% 1.5%
0.48 0.56 0.69 0.91
1.07 1.51 1.53 1.97
2.1 3.2 5.3 5.7
addition, the 1% and 1.5% samples show a decrease in stress amplitude till failure. Despite the fluctuation in stress the amplitude, the 0.5% and 0.75% samples show overall increase in the stress amplitude at the onset of failure. 3.3. Microstructural evolution Fig. 3(a d) shows the loading axis inverse pole figure (IPF) maps of the deformed samples near the fracture tip. The IPF map is characterized by color gradient within the grains indicating slip activity and dislocation dominated deformation in all the samples. Low strain amplitude samples are characterized by voids near grain boundaries and triple junctions while 1% and 1.5% strain amplitude samples show profuse intergranular cracking accompanied with few transgranular cracks near the fracture tip. In order to understand the origin of damage in the fatigue tested samples and establish a link between the operation of slip and damage, intragranular misorientation parameters like kernel average misorientation (KAM), grain average misorientation (GAM) and geometrically 4
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Fig. 4. Secondary electron images of the fractured surface of the low cycle fatigue tested CoCuFeMnNi HEA samples for (a–c) 0.5% and (d–f) 1.5% strain amplitude.
Fig. 5. (a) Coffin-Manson plot and (b) Hysteresis loop at half-life for different strain amplitude for low cycle fatigue tested CoCuFeMnNi HEA samples.
of damage and final failure. At low strain amplitudes, the path length of dislocations is smaller and there is a chance of an edge dislocation to get trapped in the copper rich nano-clusters. At higher strain amplitude, an edge dislocation is likely to cut through the cluster due to the higher slip length. In addition, the overall probability of interaction of an edge dislocation with a copper rich nano-cluster is higher at higher strain amplitude due to higher dislocation density at higher strain amplitude. A balance between dislocation accumulation by the interaction of dislocations with other dislocations, solute environment and nano-cluster and the annihilation of dislocations contribute to a steady state at low strain amplitude. However, for higher strain amplitude, dislocation accumulation is dominant, and the sample fails after few cycles. Fig. 5(a) displays the Coffin-Manson plot, which reveals a logarithmic behavior between strain amplitude and cycles to failure. The determined constant value of c (slope of the curve) is 0.34, which is significantly lower than the conventional range for steels (0.5–0.7) [49]. Recently, a new design strategy has been proposed to improve low cycle fatigue life of FeMnSi alloys, relying on increasing slip planarity by solute pinning that aids in slip reversibility and hence a higher fatigue life [50]. It has already been established that planar partial slip plays an important role in determining deformation behavior of CoCuFeMnNi high entropy alloy and it
is expected that the planar partial slip with smaller slip length will lead to better fatigue performance. However, at higher strain amplitude, it is expected that conventional octahedral slip dominates reducing slip reversibility and contributing to early failure. A similar observation of the transition from dominant partial slip to octahedral slip has been reported in the monotonic deformation by rolling using viscoelastic self-consistent simulations by Tazuddin et al. [38]. Fig. 5(b) illustrates the half-life hysteresis loops plotted in relative co-ordinates where the upper branches of hysteresis loops were almost congruent, revealing a perfect Masingbehavior [23]. This is a clear indication that the dislocation structure might remain unchanged during fatigue at various strain amplitudes. With respect to dislocation motion, the distribution of effective obstacles in Masing-type material may remain constant during fatigue deformation. 4.2. Intragranular misorientation and damage Figs. 6 and 7 show grain reference orientation deviation (GROD) and kernel average misorientation (KAM) maps near the fracture tip of the fatigue tested samples. The GROD map shows homogeneous evolution of GROD for 0.5% strain amplitude sample while all other samples show the heterogeneous distribution of GROD. These samples 5
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Fig. 6. Grain reference orientation deviation (GROD) map of low cycle fatigue tested samples at strain amplitude of (a) 0.5% (b) 0.75% (c) 1.0% (d) 1.5%.
Fig. 7. Kernel average misorientation (KAM) map of low cycle fatigue tested samples at strain amplitude of (a) 0.5% (b) 0.75% (c) 1.0% and (d) 1.5%.
are characterized by grains with lower GROD values as well as grains with very high GROD values, clearly indicating partitioning of strain between the grains. The KAM map shows higher but homogeneous distribution of KAM values for 0.5% sample while other samples show
regions of low KAM and regions of very high KAM. The KAM value is highest near voids and cracks indicating initiation of cracks and voids due to strain incompatibility between neighboring grains and related to geometry necessary dislocations (GND) during the course of cyclic 6
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Fig. 8. Geometrically necessary dislocation (GND) density map of low cycle fatigue tested samples at strain amplitude of (a) 0.5% (b) 0.75% (c) 1.0% and (d) 1.5%.
Fig. 9. Back scattered images with channeling contrast for low cycle fatigue tested samples at (a) 0.5% and (b) 1.5% stain amplitude.
deformation. The strain incompatibility between neighboring grains induces the local misorientation within the grains and can be correlated with the magnitude of the macroscopic plastic strain. Kamaya et al. [51] have proposed that the effect of plastic strain on cracking can be correlated with the evolution of geometrically necessary dislocation (GND) density in the microstructure. The GND density maps near the fracture tip of the fatigue tested samples obtained from EBSD are shown in Fig. 8(a–d), There is a clear indication of increase in GND density in the vicinity of grain boundaries due to significantly strong incompatibility between neighboring grains. The BSE image with electron channeling contrast in Fig. 9(a) near the fracture tip at 0.5% strain amplitude sample hardly shows any contrast within the grains and only orientation contrast is observed. Insignificant change was observed after changing the diffraction condition by rotating the sample, indicating very little dislocation density. In fact, this was uniform throughout the sample with only the orientation contrast visible for different grains. However, for the high strain amplitude deformed sample shown in Fig. 9(b), significant contrast was observed, and there was drastic change in the contrast for different diffraction conditions. This clearly indicates higher dislocation density and more heterogeneous distribution of dislocations at higher strain amplitude. The distribution of dislocation is measured in terms of strain heterogeneity due to locally different stress state during plastic deformation of the material. Due to the different crystallographic
orientation of grains, some grains are more preferentially oriented where the dislocation glide can take place easily than other grains during cyclic deformation. These preferential dislocation motion induces local plastic deformation and creates local stress concentrations and strain gradient. This results in locally different regions of cyclic softening and hardening behavior during cyclic deformation [52]. 4.3. Role of slip and grain boundary character in damage Fig. 10(a–d) shows Schmid factor maps determined using octahedral slip for tension/compression at various strain amplitude (0.5–1.5%) for a range of (0.3–0.5) Schmid factor value. Fig. 10(e and f) displays the variation of the relative fraction with Schmid factor and strain amplitudes to understand the activity of slip systems. The plot indicates distinct value of Schmid factor associated with cyclic plastic stain accommodation in terms of slip activity. It is expected that only single slip system is active for Schmid factor greater than 0.41 and multi-slip is operative for Schmid factor in the range of 0.32–0.41 while no slip is expected for a value less than 0.32 based on slip line trace analysis [53]. The distribution of Schmid factor in all the fatigue tested samples clearly show that with the increase in strain amplitude, there is a reduction in fraction of grains undergoing single slip and most grains undergo multi-slip thereby reducing slip reversibility [53]. Interaction between dislocations and grain boundaries play an 7
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Fig. 10. (a–d) Schmid factor map, (e) Schmid factor distribution for cyclic deformed samples and (f) Schmid factor domain varying as function of strain amplitude.
fraction of special boundaries in general random grain boundary networks. Fig. 11(a and b) shows that the annealing twins generated during processing can act as barrier for crack propagation and void nucleation unlike random high angle grain boundaries during cyclic deformation. Thus, increasing the fraction of special twin boundaries at the expense of random high angle grain boundaries by grain boundary engineering can improve the low cycle fatigue performance of CoCuFeMnNi HEA. Previous researchers have also reported that the grain size, annealing twins and grain boundary network can improve low cycle fatigue performance of conventional alloys [54,55]. Fig. 12(a) and (b) show the interaction of dislocations with nanoclusters. At low stain amplitude the dislocations are pinned by the nanoclusters. However, at higher strain amplitude or high stress applied during cyclic loading, the copper rich nano-clusters get sheared. Fig. 12(c) shows the crack initiation taking place at grain boundary triple junction due to the presence of high geometrically necessary dislocation density around the triple point... There is a greater chance for the crack to initiate at the triple junction and follow an intergranular path. Fig. 11. Loading axis inverse pole figure (IPF) map near crack growth region of the 1.5% stain amplitude fatigue tested CoCuFeMnNi HEA.
5. Conclusions
important role during cyclic deformation and influenced the crack propagation. Therefore, engineering the grain boundary network can improve the mechanical properties of the materials by increasing the
Low cycle fatigue behavior of CoCuFeMnNi high entropy alloy was investigated at different strain amplitudes, and detailed microstructural characterization of the failed samples was carried out to decipher the 8
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Fig. 12. Schematic showing interaction of dislocation with copper rich nano-clusters at (a) low strain amplitude (b) high strain amplitude and (c) crack initiation and propagation in CoCuFeMnNi high entropy alloy subjected to low cycle fatigue.
micro-mechanisms of damage. The important findings from the investigation are provided here.
[2]
1. There is a decrease in low cycle fatigue life with an increase in strain amplitude for CoCuFeMnNi high entropy alloy indicating poor low cycle fatigue behavior manifested by lower Coffin-Manson exponent compared to steels. 2. There is cyclic hardening followed with saturation for low strain amplitude while only softening regime is observed post initial hardening for high strain amplitude regime samples. 3. Higher strain amplitude leads to the evolution of higher intragranular misorientation contributing to the accumulation of damage near grain boundaries. 4. The planar nature of dislocation aids in the reversible motion of dislocations while the presence of solute atmosphere and copper rich nano-clusters retard the reversibility of dislocations. Higher dislocation density at higher strain amplitude leads to higher damage accumulation near grain boundaries which are weakened by the presence of clusters themselves and lead to intergranular failure. 5. Coherent annealing twin boundaries show little damage compared to random high angle grain boundaries indicating that increasing population of special boundaries may improve the low cycle fatigue performance of CoCuFeMnNi high entropy alloy. Further work is required to unequivocally establish the micro-mechanisms of low cycle fatigue failure in CoCuFeMnNi high entropy alloy.
[3]
[4] [5]
[6]
[7] [8] [9] [10] [11]
Acknowledgements
[12]
The authors would like to thank Science and Engineering Research Board under Department of Science and Technology (DST) for funding, IIT Kanpur for experimental facilities and Dr. K.G. Pradeep for Atom Probe Tomography measurements.
[13]
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Micro-mechanisms of microstructural damage due to low cycle fatigue in CoCuFeMnNi high entropy alloy Fateh Bahadur, Krishanu Biswas, N.P. Gurao
T
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Department of Materials Science and Engineering, IIT Kanpur, Kanpur 208016, India
ARTICLE INFO
ABSTRACT
Keywords: Advanced materials Low cycle fatigue Microstructures Dislocations Damage
Low cycle fatigue behavior of equiatomic CoCuFeMnNi high entropy alloy was investigated under fully reversible strain control mode followed with detailed microstructural characterization using electron back scatter diffraction. There is an increase in the intragranular misorientation and geometrically necessary dislocation density with increase in the strain amplitude that leads to accumulation of damage and intergranular cracking. The interaction of dislocations with copper rich nano-clusters, solute environment and grain boundaries affect slip reversibility, thereby deciding cyclic deformation behaviour. Annealing twin boundaries are resistant to damage and increasing their population density by grain boundary engineering can improve performance of CoCuFeMnNi alloy.
1. Introduction Design and development of new alloys for achieving excellent mechanical properties like high specific strength, optimum combination of strength and ductility, excellent fracture toughness at low temperature, superior fatigue and creep resistance has been pursued to cater the stringent demands of the engineering industries in the recent past. Recently, a new class of materials known as high entropy alloys (HEAs) or complex concentrated alloys (CCAs) that are multi-principal multicomponent alloys containing five or more elements in equal or near equal proportions have been developed [1,2]. Their high configurational entropy of mixing favors the formation of solid solution over brittle intermetallic phases at high temperature [1–4]. This innovative alloy design strategy overcomes the conventional approach in which properties are primarily governed by principal element, thus, providing enormous opportunities in terms of physical and functional properties by permitting extensive and flexible tuning of alloy composition. The past decade of dedicated research revealed that many HEAs possess unique properties in comparison with the conventional alloys for instance, high thermal and microstructural stability [5,6], high hardness [6], high strength at elevated temperature [7,8], superior fracture toughness [9], and excellent wear resistance [10]. Among HEAs, equiatomic CoCrFeMnNi Face Centered Cubic (FCC) HEA (Cantor alloy) exhibits high tensile strength (1130 MPa), sufficient ductility (65%), and higher fracture toughness (200 MPa.m1/2 ); especially at cryogenic temperatures due to the presence of deformation twins and has been studied in great details [4,9,11–13].
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Most of these studies regarding mechanical behavior are usually investigated in uniaxial loading, [14–19] which are necessary but still insufficient to broaden the application range of HEAs for structural applications. Most engineering applications involve load reversal and cyclic loading in service, and therefore, it is essential to study them under cyclic loading. In particular, low cycle fatigue behavior of metals and alloys is important from technological perspective as well as from a scientific point of view to investigate the accumulation of damage wherein dislocations and their interaction with each other and other defects play an important role. Low cycle fatigue involves plastic strain and is generally investigated in strain control mode. Deformation behaviour and failure of different engineering materials like IF steel [20], dual phase steel, low alloy steel, TWIP/TRIP steel [21–24], stainless steel [25–27], commercially pure titanium, titanium alloy [28,29], and Ni-based superalloys [30,31] has been well-established under low cyclic fatigue regime. However, the research on fatigue behavior, particularly pertaining to low cycle fatigue of high entropy alloys, has been limited. Hemphill et al. [32] have examined the high cycle fatigue behavior of as-cast and wrought Al0.5CoCrCuFeNi alloy using four-point bending test at room temperature and found that the fatigue endurance limit lies in the range of 540–945 MPa. A similar composition of Al0.5CoCrCuFeNi alloy with high purity elements has also been studied by Tang et al. [33] who showed that the use of high purity elements results in better fatigue performance. This was attributed to the formation of nano-twins during cyclic loading in the alloy containing high purity elements leading to higher endurance limit and prolonged fatigue life compared to many
Corresponding author. E-mail address: [email protected] (N.P. Gurao).
https://doi.org/10.1016/j.ijfatigue.2019.105258 Received 26 June 2019; Received in revised form 12 August 2019; Accepted 28 August 2019 Available online 29 August 2019 0142-1123/ © 2019 Elsevier Ltd. All rights reserved.
International Journal of Fatigue 130 (2020) 105258
F. Bahadur, et al.
Nomenclature
R APT IPF GROD KAM GND
List of symbols and acronyms Δε/2 c n
strain amplitude fatigue ductility exponent strain hardening co-efficient
conventional alloys. Niendorf et al. [34] reported the low cycle fatigue behavior of Fe50Mn30Co10Cr10 alloy showing remarkable changes in the deformation mechanisms under monotonic and cyclic loading conditions. The alloy showed poor strain hardening in tension but occurrence of strain-induced martensitic transformation in low cycle fatigue regime contributed to better fatigue performance. This peculiar behavior was attributed to partial deformation reversibility and the presence of the planar nature of slip. In other investigations, microstructural engineering approaches like grain refinement were employed to improve fatigue performance [35,36]. Multi-phase high entropy alloys investigated till date show remarkable fatigue strength compared to the conventional alloy systems. But the fundamental understanding of the fatigue deformation mechanism of single phase high entropy alloys, especially under low cycle fatigue regime has not been studied yet. Accordingly, the present study focuses on the low cycle fatigue behavior of single phase FCC CoCuFeMnNi HEA system developed by Tazuddin et al. [37], under strain control conditions. It was shown that operation of planar partial slip along with conventional octahedral slip led to the unique Goss-Brass type texture evolution during cold rolling of the alloy [38]. An atom probe tomography study by Sonkusare et al. [39] revealed the presence of copper rich nano-clusters (2.5 nm average
strain ratio atom probe tomography inverse pole figure grain reference orientation deviation kernel average misorentation geometrically necessary dislocation
size) in the FCC matrix of the alloy. Other investigations on the alloy have established the importance of the copper rich nano-clusters and their interactions with dislocations in tension and high pressure torsion in monotonic and cyclic mode [40,41]. Thus, the deformation behavior of the CoCuFeMnNi high entropy alloy is studied for different loading conditions. However, the cyclic deformation of the alloy is still unexplored. The present investigation deals with understanding the deformation mechanism and microstructural evolution of CoCuFeMnNi alloy subjected to strain controlled low cycle fatigue using various characterization tools like scanning electron microscopy (SEM), electron back scatter diffraction (EBSD) and back scattered electron imaging showing electron channeling contrast. 2. Experimental procedure 2.1. Material and methods Equiatomic CoCuFeMnNi alloy was prepared by arc melting buttons of ~12 g of equimolar mixture of elements with greater than 99.9% purity obtained from Sigma Aldrich. Alloy buttons were remelted at least four times to obtain compositional homogeneity. Four such alloy
Fig. 1. (a) X-ray diffraction pattern, (b) normal direction inverse pole figure map and (c) elemental distribution map from atom probe tomography of the homogenized CoCuFeMnNi HEA sample. 2
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buttons were remelted again and suction cast into a billet of dimension 10 × 10 × 60 mm3. The billet was sealed in a quartz tube with argon and homogenized at 1273 K for 24 h followed by water quenching. Subsequently, the homogenized billet was cold rolled to 50% reduction in thickness at room temperature. The rolled specimen was subjected to annealing for 2 h at 1273 K to obtain a fully crystallized microstructure. Tensile and fatigue samples were machined from the annealed billet by electric discharge machining according to the ASTM E8 (10 mm × 2.5 mm × 2.5 mm) and ASTM E-606 (8 mm × 5 mm × 2.5 mm) standards respectively.
diffraction (EBSD, Oxford Instruments, U.K.) detector. EBSD scans were acquired with Aztec HKL software. The post-processing of the data was performed using TSL-OIM version 7.2 software. The EBSD samples were prepared by polishing to 2000 grit size using emery paper and then cloth polishing using the 0.5 μm alumina suspension followed with vibromet polishing (Buehler, USA) using colloidal silica suspension with a particle size of 50 nm. Backscattered electron imaging was carried out to obtain electron channeling contrast [44] by tilting the sample 10 degree at a working distance of 10 mm and rotating the sample to reveal the dislocation structure near the fracture tip of the fatigue tested samples.
2.2. Tensile and fatigue test
3. Results
The uniaxial tensile test was carried out at room temperature with a strain rate of 1 × 10 3s 1 using Instron-1195 Universal Testing Machine. Low cycle fatigue tests were also performed under straincontrol conditions at different strain amplitude (Δε/2) of 0.5%, 0.75%, 1% and 1.5% with fully reversible cyclic strain condition (R = −1) up to failure. All fatigue tests were conducted with the strain rate of 5 × 10−3 s−1 using the 100 kN Universal Testing Machine (BiSS, Bangalore, India).
3.1. Initial material X-ray diffraction (XRD) pattern of the homogenized CoCuFeMnNi alloy shown in Fig. 1(a) indicates the presence of FCC single phase with a lattice parameter of 0.361 nm. The normal direction (ND) inverse pole figure (IPF) map of the investigated HEA in Fig. 1(b) shows a fully recrystallized microstructure with annealing twins and an area average grain size of ~47 µm. Atom probe tomography carried out on the homogenized sample provided elemental distribution at the atomic length scale. Fig. 1(c) shows the elemental distribution from atom probe tomography indicating the homogeneous distribution of cobalt, iron, manganese, and nickel while there is heterogeneous distribution of copper, characterized by the presence of near spherical copper rich nano-clusters with an average diameter of 2.5 nm.
2.3. X-ray diffraction and microstructure analyses The crystal structure of the homogenized HEA was examined by Xray diffraction (PANalytical, Netherland), with a step size of 0.02 degree/ s, operated at 45 kV and 40 mA using CuKα (λ = 0.15402 nm) radiation. The grain morphology, fractography and compositional analyses of the synthesized sample were carried out utilizing field emission scanning electron microscope (FE-SEM) (JSM-7100F, JEOL, Japan) equipped with energy dispersive spectroscopy facility (EDS, Oxford Instruments, UK) operated at 20 kV. Atom probe tomography (APT) was carried out using the local electrode atom probe (LEAP 4000X HR™, Cameca Instruments) to obtain near-atomic scale elemental distribution. APT samples were prepared by using standard procedure [42,43] using a dual-beam focused ion beam (FEI Helios Nanolab 660, USA). A volume of 60 × 60 × 100 nm3 was acquired, through a pulse repetition rate of 250 kHz and pulse energy of 30 pJ. The tip temperature was maintained at around 60 K. The APT data was evaluated with the help of IVAS 3.6.10a software provided by Cameca Instruments.
3.2. Tensile and fatigue tests Engineering stress - engineering strain behaviour of the CoCuFeMnNi alloy as shown in Fig. 2(a) indicates that the investigated alloy exhibits yield strength of 324 ± 20 MPa and ultimate tensile strength of 710 ± 35 MPa. Cyclic tests performed with fully reversible strain amplitude (Δε/2) of 0.5%, 0.75%, 1%, and 1.5% showed a unique evolution of stress amplitude with number of cycles that is depicted in Fig. 2(b). The stress amplitude at the onset of cyclic loading increases with an increase in the strain amplitude. However, the evolution of the stress amplitude with number of cycles is different for 0.5% and 0.75% (low) strain amplitude compared to the 1% and 1.5% (high) strain amplitude with the former two showing a significantly higher number of cycles to failure than the latter two. There is initial hardening; that is an increase in stress amplitude with the number of cycles for all the four samples. However, 0.5% and 0.75% strain amplitude samples show softening followed with fluctuations in the stress amplitude. In
2.4. Electron back scatter diffraction (EBSD) and electron channeling contrast imaging Electron back scatter diffraction (EBSD) data acquisition was carried out in FE-SEM from JEOL equipped with electron back scatter
Fig. 2. (a) Engineering stress – engineering strain diagram and (b) variation of stress amplitude with number of cycles in low cycle fatigue test at different strain amplitude for CoCuFeMnNi HEA. 3
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necessary dislocation density (GND) were determined from the EBSD scan and are listed in Table 1. It is observed that there is an increase in the average KAM and GAM value with increase in the strain amplitude. Similarly, the geometrically necessary dislocation density is also found to increase with the increase in strain amplitude. Fig. 3(e) shows that the kernel average misorientation (KAM) distribution shifts to the right and becomes broader as strain amplitude increases, indicating an increase in dislocation density and more heterogeneous distribution of the dislocations with increasing strain amplitude [45]. 3.4. Fractography Detailed analysis of the fractured surface of the fatigue tested samples was carried out using a scanning electron microscope. The illustrative SEM fractographs at lowest (0.5%) and highest (1.5%) strain amplitude are shown in Fig. 4(a–c) and Fig. 4(d–f) respectively. Fig. 4(a) shows formation of striation marks with different orientation observed within different grains. Fig. 4(b) and (c) display crack propagation path with higher number of cycles, which are found to be mostly intergranular in nature. Fig. 4(d) shows the presence of void or inhomogeneity in the material that can act as local stress raisers at grain boundaries, where crack initiation will take place. Once the cracks are initiated the stable propagation of the crack with each cycle is a clear representation of striation mark as shown in Fig. 4(e). However, as shown in Fig. 4(f), the spacing between the striation marks was found to increase with strain amplitude. The increase of strain amplitude causes faster crack growth rate. . Therefore, the mechanism of fatigue failure in the HEA is similar to that in conventional alloys based on crack initiation, stable crack propagation, and final fracture [46]. 4. Discussion
Fig. 3. Loading axis inverse pole figure map near the fracture tip of samples tested in fatigue at strain amplitude of (a) 0.5% (b) 0.75% (c) 1% (d) 1.5% and (e) Kernel average misorientation distribution from all the fatigue tested samples.
4.1. Cyclic stress response The variation of stress amplitude with strain amplitude shows two distinct behavior for the equiatomic CoCuFeMnNi alloy in the present investigation. The stress amplitude reaches a maximum and fluctuates over a prolonged cyclic stage and then decreases sharply for strain amplitude of 0.5% and 0.75%. However, at higher strain amplitudes of 1.0% and 1.5%, value of stress amplitude increases but the regime of high-stress amplitude is very small, and the sample fails after few cycles compared to low strain amplitude samples. Chen et al. [47] have shown that long range motion of dislocations at higher strain amplitude shows a distinct cyclic hardening/softening behaviour compared to short range motion of dislocations at low strain amplitude in Hadfield steel. Agrawal et al. [40] have reported four factors responsible for the strain hardening behavior of the CoCuFeMnNi HEA under monotonic condition, namely; (i) dislocation multiplication, (ii) grain boundary strengthening, (iii) solid solution strengthening and (iv) interaction between dislocations and copper rich nano-clusters. Cyclic loading leads to forward and backward movement of dislocations characterized by long range motion at higher strain amplitude and short-range motion at low strain amplitude. All the aforementioned four interactions play an important role in the reversible nature of dislocation motion in the CoCuFeMnNi high entropy alloy. Grain boundaries act as barrier for the motion of dislocations leading to formation of dislocation pile-up at grain boundaries. During reverse loading, the dislocations can move in the reverse direction at lower stress and annihilate each other, contributing to cyclic softening. This happens for materials with low strain hardening exponent (n ≤ 0.15), while for materials with higher strain hardening co-efficient (n > 0.15) dislocation accumulation occur, leading to cyclic hardening [48]. The alloy under investigation shows higher strain hardening in tension [38] and is a cyclic hardening material at all the strain amplitudes. Both solute atmosphere and copper rich nanoclusters are likely to pin dislocations and reduce the mobility of dislocation. This, in turn, can cause an increase in dislocation density leading to accumulation of intragranular misorientation that can cause accumulation
Table 1 Variation of average kernel average misorientation (KAM), grain average misorientation (GAM) and geometrically necessary dislocation (GND) density for the low cycle fatigue tested samples at different strain amplitude. Strain amplitude
Average KAM°
Average GAM°
Average GND density × 1013 (m−2)
0.5% 0.75% 1% 1.5%
0.48 0.56 0.69 0.91
1.07 1.51 1.53 1.97
2.1 3.2 5.3 5.7
addition, the 1% and 1.5% samples show a decrease in stress amplitude till failure. Despite the fluctuation in stress the amplitude, the 0.5% and 0.75% samples show overall increase in the stress amplitude at the onset of failure. 3.3. Microstructural evolution Fig. 3(a d) shows the loading axis inverse pole figure (IPF) maps of the deformed samples near the fracture tip. The IPF map is characterized by color gradient within the grains indicating slip activity and dislocation dominated deformation in all the samples. Low strain amplitude samples are characterized by voids near grain boundaries and triple junctions while 1% and 1.5% strain amplitude samples show profuse intergranular cracking accompanied with few transgranular cracks near the fracture tip. In order to understand the origin of damage in the fatigue tested samples and establish a link between the operation of slip and damage, intragranular misorientation parameters like kernel average misorientation (KAM), grain average misorientation (GAM) and geometrically 4
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Fig. 4. Secondary electron images of the fractured surface of the low cycle fatigue tested CoCuFeMnNi HEA samples for (a–c) 0.5% and (d–f) 1.5% strain amplitude.
Fig. 5. (a) Coffin-Manson plot and (b) Hysteresis loop at half-life for different strain amplitude for low cycle fatigue tested CoCuFeMnNi HEA samples.
of damage and final failure. At low strain amplitudes, the path length of dislocations is smaller and there is a chance of an edge dislocation to get trapped in the copper rich nano-clusters. At higher strain amplitude, an edge dislocation is likely to cut through the cluster due to the higher slip length. In addition, the overall probability of interaction of an edge dislocation with a copper rich nano-cluster is higher at higher strain amplitude due to higher dislocation density at higher strain amplitude. A balance between dislocation accumulation by the interaction of dislocations with other dislocations, solute environment and nano-cluster and the annihilation of dislocations contribute to a steady state at low strain amplitude. However, for higher strain amplitude, dislocation accumulation is dominant, and the sample fails after few cycles. Fig. 5(a) displays the Coffin-Manson plot, which reveals a logarithmic behavior between strain amplitude and cycles to failure. The determined constant value of c (slope of the curve) is 0.34, which is significantly lower than the conventional range for steels (0.5–0.7) [49]. Recently, a new design strategy has been proposed to improve low cycle fatigue life of FeMnSi alloys, relying on increasing slip planarity by solute pinning that aids in slip reversibility and hence a higher fatigue life [50]. It has already been established that planar partial slip plays an important role in determining deformation behavior of CoCuFeMnNi high entropy alloy and it
is expected that the planar partial slip with smaller slip length will lead to better fatigue performance. However, at higher strain amplitude, it is expected that conventional octahedral slip dominates reducing slip reversibility and contributing to early failure. A similar observation of the transition from dominant partial slip to octahedral slip has been reported in the monotonic deformation by rolling using viscoelastic self-consistent simulations by Tazuddin et al. [38]. Fig. 5(b) illustrates the half-life hysteresis loops plotted in relative co-ordinates where the upper branches of hysteresis loops were almost congruent, revealing a perfect Masingbehavior [23]. This is a clear indication that the dislocation structure might remain unchanged during fatigue at various strain amplitudes. With respect to dislocation motion, the distribution of effective obstacles in Masing-type material may remain constant during fatigue deformation. 4.2. Intragranular misorientation and damage Figs. 6 and 7 show grain reference orientation deviation (GROD) and kernel average misorientation (KAM) maps near the fracture tip of the fatigue tested samples. The GROD map shows homogeneous evolution of GROD for 0.5% strain amplitude sample while all other samples show the heterogeneous distribution of GROD. These samples 5
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Fig. 6. Grain reference orientation deviation (GROD) map of low cycle fatigue tested samples at strain amplitude of (a) 0.5% (b) 0.75% (c) 1.0% (d) 1.5%.
Fig. 7. Kernel average misorientation (KAM) map of low cycle fatigue tested samples at strain amplitude of (a) 0.5% (b) 0.75% (c) 1.0% and (d) 1.5%.
are characterized by grains with lower GROD values as well as grains with very high GROD values, clearly indicating partitioning of strain between the grains. The KAM map shows higher but homogeneous distribution of KAM values for 0.5% sample while other samples show
regions of low KAM and regions of very high KAM. The KAM value is highest near voids and cracks indicating initiation of cracks and voids due to strain incompatibility between neighboring grains and related to geometry necessary dislocations (GND) during the course of cyclic 6
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Fig. 8. Geometrically necessary dislocation (GND) density map of low cycle fatigue tested samples at strain amplitude of (a) 0.5% (b) 0.75% (c) 1.0% and (d) 1.5%.
Fig. 9. Back scattered images with channeling contrast for low cycle fatigue tested samples at (a) 0.5% and (b) 1.5% stain amplitude.
deformation. The strain incompatibility between neighboring grains induces the local misorientation within the grains and can be correlated with the magnitude of the macroscopic plastic strain. Kamaya et al. [51] have proposed that the effect of plastic strain on cracking can be correlated with the evolution of geometrically necessary dislocation (GND) density in the microstructure. The GND density maps near the fracture tip of the fatigue tested samples obtained from EBSD are shown in Fig. 8(a–d), There is a clear indication of increase in GND density in the vicinity of grain boundaries due to significantly strong incompatibility between neighboring grains. The BSE image with electron channeling contrast in Fig. 9(a) near the fracture tip at 0.5% strain amplitude sample hardly shows any contrast within the grains and only orientation contrast is observed. Insignificant change was observed after changing the diffraction condition by rotating the sample, indicating very little dislocation density. In fact, this was uniform throughout the sample with only the orientation contrast visible for different grains. However, for the high strain amplitude deformed sample shown in Fig. 9(b), significant contrast was observed, and there was drastic change in the contrast for different diffraction conditions. This clearly indicates higher dislocation density and more heterogeneous distribution of dislocations at higher strain amplitude. The distribution of dislocation is measured in terms of strain heterogeneity due to locally different stress state during plastic deformation of the material. Due to the different crystallographic
orientation of grains, some grains are more preferentially oriented where the dislocation glide can take place easily than other grains during cyclic deformation. These preferential dislocation motion induces local plastic deformation and creates local stress concentrations and strain gradient. This results in locally different regions of cyclic softening and hardening behavior during cyclic deformation [52]. 4.3. Role of slip and grain boundary character in damage Fig. 10(a–d) shows Schmid factor maps determined using octahedral slip for tension/compression at various strain amplitude (0.5–1.5%) for a range of (0.3–0.5) Schmid factor value. Fig. 10(e and f) displays the variation of the relative fraction with Schmid factor and strain amplitudes to understand the activity of slip systems. The plot indicates distinct value of Schmid factor associated with cyclic plastic stain accommodation in terms of slip activity. It is expected that only single slip system is active for Schmid factor greater than 0.41 and multi-slip is operative for Schmid factor in the range of 0.32–0.41 while no slip is expected for a value less than 0.32 based on slip line trace analysis [53]. The distribution of Schmid factor in all the fatigue tested samples clearly show that with the increase in strain amplitude, there is a reduction in fraction of grains undergoing single slip and most grains undergo multi-slip thereby reducing slip reversibility [53]. Interaction between dislocations and grain boundaries play an 7
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Fig. 10. (a–d) Schmid factor map, (e) Schmid factor distribution for cyclic deformed samples and (f) Schmid factor domain varying as function of strain amplitude.
fraction of special boundaries in general random grain boundary networks. Fig. 11(a and b) shows that the annealing twins generated during processing can act as barrier for crack propagation and void nucleation unlike random high angle grain boundaries during cyclic deformation. Thus, increasing the fraction of special twin boundaries at the expense of random high angle grain boundaries by grain boundary engineering can improve the low cycle fatigue performance of CoCuFeMnNi HEA. Previous researchers have also reported that the grain size, annealing twins and grain boundary network can improve low cycle fatigue performance of conventional alloys [54,55]. Fig. 12(a) and (b) show the interaction of dislocations with nanoclusters. At low stain amplitude the dislocations are pinned by the nanoclusters. However, at higher strain amplitude or high stress applied during cyclic loading, the copper rich nano-clusters get sheared. Fig. 12(c) shows the crack initiation taking place at grain boundary triple junction due to the presence of high geometrically necessary dislocation density around the triple point... There is a greater chance for the crack to initiate at the triple junction and follow an intergranular path. Fig. 11. Loading axis inverse pole figure (IPF) map near crack growth region of the 1.5% stain amplitude fatigue tested CoCuFeMnNi HEA.
5. Conclusions
important role during cyclic deformation and influenced the crack propagation. Therefore, engineering the grain boundary network can improve the mechanical properties of the materials by increasing the
Low cycle fatigue behavior of CoCuFeMnNi high entropy alloy was investigated at different strain amplitudes, and detailed microstructural characterization of the failed samples was carried out to decipher the 8
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Fig. 12. Schematic showing interaction of dislocation with copper rich nano-clusters at (a) low strain amplitude (b) high strain amplitude and (c) crack initiation and propagation in CoCuFeMnNi high entropy alloy subjected to low cycle fatigue.
micro-mechanisms of damage. The important findings from the investigation are provided here.
[2]
1. There is a decrease in low cycle fatigue life with an increase in strain amplitude for CoCuFeMnNi high entropy alloy indicating poor low cycle fatigue behavior manifested by lower Coffin-Manson exponent compared to steels. 2. There is cyclic hardening followed with saturation for low strain amplitude while only softening regime is observed post initial hardening for high strain amplitude regime samples. 3. Higher strain amplitude leads to the evolution of higher intragranular misorientation contributing to the accumulation of damage near grain boundaries. 4. The planar nature of dislocation aids in the reversible motion of dislocations while the presence of solute atmosphere and copper rich nano-clusters retard the reversibility of dislocations. Higher dislocation density at higher strain amplitude leads to higher damage accumulation near grain boundaries which are weakened by the presence of clusters themselves and lead to intergranular failure. 5. Coherent annealing twin boundaries show little damage compared to random high angle grain boundaries indicating that increasing population of special boundaries may improve the low cycle fatigue performance of CoCuFeMnNi high entropy alloy. Further work is required to unequivocally establish the micro-mechanisms of low cycle fatigue failure in CoCuFeMnNi high entropy alloy.
[3]
[4] [5]
[6]
[7] [8] [9] [10] [11]
Acknowledgements
[12]
The authors would like to thank Science and Engineering Research Board under Department of Science and Technology (DST) for funding, IIT Kanpur for experimental facilities and Dr. K.G. Pradeep for Atom Probe Tomography measurements.
[13]
[14]
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