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Pile-up and sink-in nanoindentation behaviors in AlCoCrFeNi multi-phase high entropy alloy
Author’s Accepted Manuscript Pile-up and sink-in nanoindentation behaviors in AlCoCrFeNi multi-phase high entropy alloy Gokul Muthupandi, Ka Ram Lim, Young-Sang Na, Jieun Park, Dongyun Lee, Hanjong Kim, Seonghun Park, Yoon Suk Choi www.elsevier.com/locate/msea
PII: DOI: Reference:
S0921-5093(17)30497-5 http://dx.doi.org/10.1016/j.msea.2017.04.045 MSA34946
To appear in: Materials Science & Engineering A Received date: 8 March 2017 Accepted date: 12 April 2017 Cite this article as: Gokul Muthupandi, Ka Ram Lim, Young-Sang Na, Jieun Park, Dongyun Lee, Hanjong Kim, Seonghun Park and Yoon Suk Choi, Pile-up and sink-in nanoindentation behaviors in AlCoCrFeNi multi-phase high entropy a l l o y , Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2017.04.045 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Pile-up and sink-in nanoindentation behaviors in AlCoCrFeNi multi-phase high entropy alloy
Gokul Muthupandi1, Ka Ram Lim2, Young-Sang Na2, Jieun Park3, Dongyun Lee4,5, Hanjong Kim6, Seonghun Park6, Yoon Suk Choi1*
1
School of Materials Science and Engineering, Pusan National University, Busan, Korea
2
Titanium Department, Korea Institute of Materials Science, Changwon, Gyeongnam, Korea
3
Department of Cogno-Mechatronics Engineering, Pusan National University, Busan, Korea
4
Department of Nano Fusion Technology, Pusan National University, Busan, Korea
5
Department of Nanoenergy Engineering, Pusan National University, Busan, Korea
6
School of Mechanical Engineering, Pusan National University, Busan, Korea
*
Corresponding author: [email protected]
Abstract Microstructures and nanoindentation behaviors were studied on annealed AlCoCrFeNi high entropy alloy. Both pile-up and sink-in characteristics were found in the grain boundary and grain regions, respectively. The multiple phases present in the AlCoCrFeNi high entropy alloy are the reasons behind the different nanoindentation behaviors, which were identified using electron microscopy. The identified phases showed the grain boundary segregation to have A1 lattice, viz., FCC structure while the grain was distributed with A2 and B2 lattices, 1
viz., BCC and ordered BCC structures, forming the matrix with nano-precipitates of the other. The reason for the pile-up and sink-in is attributed to the dislocation activity in the individual crystal structure: large dislocation activities were found under the pile-up and little dislocation activities under the sink-in, only limited to the indenter tip. Results from a finite element analysis under an isotropic elasto-plastic condition by varying the hardness-tomodulus ratio show that high hardness-to-modulus ratio results in pile-up and the lower ratio results in sink-in. This was associated with the susceptibility to plasticity and the elastic recovery for individual phases of the AlCoCrFeNi high entropy alloy.
Keywords: High entropy alloy, Microstructure, Nanoindentation, Pile-up, Sink-in
1. Introduction Traditionally, alloys have base elements where small quantities of additional elements helped obtain desired properties. The amounts of these alloying additions have been increasing to design alloys for advanced applications. In the past few decades, high entropy alloys (HEAs) with multiple metallic elements of nearly equimolar compositions are being developed [1,2]. The idea is to stabilize solid solution phases and prevent the formation of intermetallics by increasing configurational entropy ΔSconf [3], intended to retain the ductility while possessing the high strength due to distorted lattice structures predominantly of single-phase FCC or BCC. HEA holds many properties which are better than most of the conventional alloys: good thermal stability [4,5], very high hardness [6,7], superior strength [7,8], excellent anticorrosion properties [9–11], and special electrical and magnetic properties [12]. Studies have
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been extensively carried out under different metallurgical conditions on HEAs based on transition elements [13–16]. However, the tests carried out were to assess mostly mechanical properties, little effort has been made to identify deformation micro-mechanisms associated with microstructures of HEAs. The nanoindentation technique has been often used to study micro-mechanical properties of HEAs. This is partially attributed to the limitation in size of the sample and the convenience in ‘location-specific’ property measurement [17–26]. In this study, in-depth investigation was made to understand the pile-up and sink-in behaviors that we recently observed in AlCoCrFeNi multiphase HEA during the nanoindentation. A self-similarity of the microstructure at different length scales was also studied for AlCoCrFeNi multiphase HEA. In particular, an extensive effort has been made to understand nanoindentation behaviors of multiphase AlCoCrFeNi in line with its responsive microstructural features. This also enabled us to experimentally provide evidences for different dislocation behaviors under sink-in and pile-up.
2. Experimental The multiphase AlCoCrFeNi high entropy alloy used in the current study was produced in a vacuum arc melting furnace (purity of the raw materials used was 99.99%). The alloy was remelted multiple times to ensure chemical homogeneity. The alloy was suction-cast into rods of 6 mm diameter and 50 mm long using a copper mould. Specimens of 5 mm thickness were cut and annealed at 1100 °C for 2 hours. Microstructural characterization was carried out using a light optical microscope (Nikon Eclipse MA200) under the polarized light. Further characterization of the microstructure and nanostructure was carried out using a scanning electron microscope, SEM (TESCAN MIRA3, 15kV) and a transmission electron microscope, TEM (FEI Tecnai F20, 200kV). The phase analysis was carried out by studying the selected area diffraction patterns, SADP from the TEM. The compositional analysis was 3
carried out using energy dispersive spectroscopy, EDS in both the SEM and TEM. The mechanical behavior of the HEA was investigated using nanoindentation (Agilent G200) with a Berkovich tip (tip radius 50 nm). The maximum load was set to 3mN. The load was optimized after extensive testing under different loading conditions to obtain an indentation size where the distance between the neighboring phases was considerably away. The loaddisplacement data was analyzed using Oliver-Pharr method [17,18] to determine the hardness and modulus. A minimum of 10 tests were carried out in each phase to understand the behavior. The indentations were imaged using SEM and the residual surface topology of the indentation was mapped using atomic force microscopy, AFM. Indentation marks with pileup and sink-in topologies were separated. Focused ion beam (FIB) cut specimens of 100 nm thick were made across the indents for the TEM analysis to study the dislocation behavior beneath the pile-up and sink-in. Numerical analysis using finite element method (ABAQUSTM) was also carried out to interpret the pile-up and sink-in behaviors of materials with different stress-strain responses. A 2D isotropic model was used with a sharp conical indenter with a half angle of 70.3° which simulates the area function of a Berkovich tip of the nanoindentation.
3. Results and Discussion Earlier studies on the equiatomic AlCoCrFeNi HEA found that the as-cast microstructure has a dendritic structure [3]. Fig. 1a shows the annealed (1100 °C for 2 hours) macrostructures (under the as-polished condition) of the specimen, which shows no change in macrostructure. This is attributed to the distorted lattice and the sluggish diffusion in the high entropy alloys [2]. The as-cast structure was studied to have dendritic structure, which in turn have substructures with the matrix phase rich in Al-Ni and precipitates rich in Fe-Cr [8,24,25,27– 30]. This phase separation was known to be due to the spinodal decomposition of the phases 4
[27–30]. The annealed microstructure (Fig. 1a) shows further decomposition of the alloy phase during the heat treatment resulting in precipitation and dispersion of the phases at different length scales, as shown in Fig. 1b to 1d. At the micro-scale shown in Fig. 1b and 1c, the main features are the precipitates distributed along the grain boundary (hereafter shall be denoted as “GB micro-precipitate”) and the matrix (hereafter shall be denoted as “micromatrix”). Here, the micro-matrix consists of topological separations of engraved and embossed regions (more clearly seen in Fig. 1c). These regions were identified to posses different chemical compositions (The details will be discussed later). Fig. 1d is the submicron scale microstructure taken from the engraved region (indicated by a thick arrow) of the micro-matrix in Fig. 1c. This higher-magnification microstructure shows the fine dispersion of the nanoscaled precipitates (will be noted as “nano-precipitates”) in the matrix (will be referred to as “nano-matrix” hereafter). The embossed region of the micro-matrix in Fig. 1c also showed similar microstructure is the sub-micron scale (will be discussed in the following section). The TEM analysis was performed for annealed AlCoCrFeNi in order to identify different phases at different length scales (observed in Fig. 1), and the results are shown in Fig. 2. Fig. 2a shows the microstructure taken from the TEM specimen. In the figure, the micro-matrix and GB micro-precipitate regions were indicated using the notation in Fig. 1b and 1c. Fig. 2b to 2e are the SADPs taken from local areas A to D in Fig. 2a. In Fig. 1c, the micro-matrix consisted of topologically engraved and embossed regions. These two different regions are marked as A and B in Fig. 2a. The SADP analysis from areas A and B (Fig. 2b and 2c, respectively) indicated that these two regions in the micro-matrix have mixtures of A2 and B2 lattice structures with different mixing ratios such that the one consists of A2 nanoprecipitates in the B2 nano-matrix while the other with B2 nano-precipitates in the A2 nanomatrix, resulting in the sub-micron scale microstructures shown in Fig. 1d (through SEM) 5
and Fig. 2f (through TEM). The EDS line scan result (in Fig. 3) for areas A and B in Fig. 2a and 2f shows that the A2 phase is rich in Fe-Cr and the B2 phase is rich in Ni-Al with a homogenous distribution of Co. The GB micro-precipitate (area C) of Fig. 2a was identified to have an A1 (disordered FCC) lattice structure (from the SADP in Fig. 2d). Interestingly, the micro-precipitate located at area D in Fig. 2a also showed the same lattice structure (identified from the SADP in Fig. 2e). Perhaps, the GB micro-precipitate (area C) is threedimensionally interconnected with the micro-precipitate in D (although this needs further clarification). The A1 phase in areas C and D was found to be rich in Fe-Cr-Co. In order to understand the mechanical behavior of the individual phases observed, nanoindentation tests were carried out. Figs. 4 and 5 show nanoindentation responses of the GB micro-precipitate and the micro-matrix (particularly for a region of a B2 nano-matrix with A2 precipitates), respectively. Fig. 4a identifies the location of the indentation on the GB micro-precipitate observed through SEM. Fig. 4b shows the SEM image of a representative nanoindentation (identified with a box in Fig. 4a) on the GB micro-precipitate. Here, slip traces near the edges of the indentation are observed mainly in the pile-up regions. The topology of the area of Fig. 4b measured by the AFM clearly shows the pile-up along the two edges of the indentation, as indicated by arrows in Fig. 4c. The pile-up along only two edges seems to be attributed to the anisotropy of the elastic-plastic behavior of the GB microprecipitate, which has the disordered FCC (A1) structure as confirmed in Fig. 2e [31,34]. TEM analysis to understand the dislocation behavior underneath the indentation of the GB micro-precipitate was carried out. The sample was cut across the indentation as indicated in Fig. 4b (with dashed lines). The TEM bright field image for the through-thickness section of the indentation was then taken and shown in Fig. 4d. From Fig. 4d, the indentation was confirmed to be through the GB micro-precipitate (FCC phase). It was also found from Fig. 4e that the micro-matrix (having the B2 nano-matrix with A2 nano-precipitates) is located 6
right underneath the GB micro-precipitate. Fig. 4f and 4g are SADPs taken from area A, which is away from the edge of the indenter, and area B, which is adjacent to the face of the indenter (showing pile-up above), in Fig. 4d, respectively. These SADPs indicate the different amount of dislocation activities between areas A and B underneath the indentation. Area B, which is near the face of the indenter, shows larger dislocation activities than those in area A, which is away from the edge of the indenter. This seems to be due to differences in magnitude and state of stresses between the two areas under the indentation. In particular, heavy dislocation activities in area B seem to result in the pile-up on the surface near the edge of the indentation (Fig. 4b and 4c). Fig. 4h is the TEM bright field image showing the dislocation interaction at the interface between the GB micro-precipitate (FCC phase) and the micro-matrix (B2 nano-matrix with A2 nano-precipitates), as indicated by a dotted circle in Fig. 4d. Here, the interface seems to act as a barrier against dislocation gliding from the GB micro-precipitate to the micro-matrix, as it appears that only selected (perhaps, energetically favorable) dislocations are activated in the micro-matrix region near the interface. Fig. 5a is the SEM micrograph showing nanoindentations carried out on the micro-matrix (Fig. 1c). Among these indentations one indentation was selected (indicated with a box in Fig. 5a) in an engraved region of the micro-matrix for the further analysis of the microstructure and topology using AFM and TEM, respectively. The SEM image of the selected indentation in Fig. 5b and the corresponding topology map in Fig. 5c show the sinkin near the indentation. It was also confirmed that this micro-matrix region consists of the nano-precipitates in the nano-matrix, as previously observed in Fig. 1d. Again, the throughdepth TEM analysis was carried out across the indentation (the FIB cut indicated by the dashed lines in Fig. 5b), and the resulting TEM bright field image is shown in Fig. 5d. Here, the indented tip location is indicated by an arrow in the figure. The SADP analysis (Fig. 5e) 7
for the region adjacent to the face of the indenter (area A in Fig. 5d) indicated that the indented micro-matrix region consists of the B2 nano-matrix and A2 nano-precipitates. The further SADP analysis (Figs. 5f and 5g) for areas B and C in Fig. 5d showed that another micro-matrix region having the A2 nano-matrix with B2 nano-precipitates is located underneath the indented micro-matrix region (having the B2 nano-matrix with A2 nanoprecipitates). The boundary between the two micro-matrix regions is indicated by the dashed line in Fig. 5d. The TEM image in Fig. 5d and the SADPs in Figs. 5e and 5f indicate that the indentation in the micro-matrix region induces the lower level of dislocation activities underneath the indentation, compared to the indentation in the GB micro-precipitate (Figs. 4d to f). This limited plasticity for the micro-matrix is probably due to the limited availability of active slip systems for the B2 phase nano-matrix and the presence of A2 phase nano-precipitates, which act as dislocation barriers. Also, the formation of antiphase boundaries (APB) and stacking faults in the B2 (ordered) phase is believed to limit the slip activity. Under this limited plasticity in the micro-matrix, dislocation activities seem to be localized underneath the indented tip (area B in Fig. 5d), while little dislocation activities were found adjacent to the face of the indenter (area A in Fig. 5d). In particular, the boundary between the two micromatrix regions does not seem to be a strong barrier for the dislocation transmittance, as confirmed by no contrast change across the boundary in the TEM bright field image (area B in Fig. 5d). This seems to be associated with the similar microstructure between the two micro-matrix regions (the B2 nano-matrix with A2 nano-precipitates and the other way around) and lattice dimensions of the B2 and A2 phases. It was confirmed from Figs. 4 and 5 that the GB micro-precipitate and the micro-matrix have different crystal structures and microstructural features (A1 (FCC) phase for the former and the B2 nano-matrix with A2 nano-precipitates (or the other way around) for the latter), and 8
revealed the pile-up and sink-in responses upon nanoindenation, respectively. Dislocation activities near and underneath the indentation were also different between the two cases. In order to further investigate this difference, load-displacement curves after nanoindentation tests were systematically analyzed and compared between the two cases. Fig. 6 shows the representative load-displacement curves (corresponding to nanoindentations in Figs. 4b and 5b) for the GB micro-precipitate and the micro-matrix after the load-controlled (3 mN) nanoindentation tests. The serrations (or pseudo pop-in) observed in loading curves for both cases are signs of the evolution of dislocation barriers, which impede the dislocation motion and increase the energy required for deformation to continue. This seems to be associated with solid solution strengthening for the GB micro-precipitate and dispersion strengthening of nano-precipitates (in addition to APB strengthening) for the micro-matrix. The unloading curve shows the similar recovery behavior between the two cases. In Fig. 6, however, the indentation depth was deeper in the GB micro-precipitate than in the micro-matrix for the same applied load (3mN). In order to perform statistically more reliable analysis for this nanoindentation response, a series of nanoindentation tests was performed over the different regions of the specimen, and the resulting (final) indentation depth (hfinal) values were compared between the two cases in Fig. 7. Different symbols in Fig. 7 mean different regions for nanoindentation tests. Here, hfinal values were plotted as a function of the hardness since the hardness is inclusive of plasticity and elasticity, also the anisotropy does not significantly affect the hardness [21,35]. The hfinal-hardness plot in Fig. 7 shows an inverse relation for both cases, which is understandable because the hardness is the resistance to the deformation under the indentation or abrasion. In particular, the GB micro-precipitate shows lower hardnesses and higher hfinal values, compared to the micro-matrix. This indicates that the GB micro-precipitate is more susceptible to the plastic deformation than the micro-matrix. Using the same data in Fig. 7, the amount of the elastic recovery (Dh) was 9
determined from load-displacement curves and plotted as a function of the hardness in Fig. 8. The result showed higher elastic recovery in the micro-matrix than in the GB microprecipitate. Results from Figs. 7 and 8 imply that the nanoindentation deformation of the micro-matrix is relatively elastic in nature, compared to the GB micro-precipitate. Combining results from Figs. 4 to 8, distinctive nanoindentation characteristics were identified between the GB micro-precipitate (FCC phase) and the micro-matrix (the B2 nanomatrix with A2 nano-precipitates) for the heat-treated AlCoCrFeNi multiphase HEA. The pile-up observed in the GB micro-precipitate was accompanied by heavy plasticity (underneath the indentation) particularly near the face of the indenter, which was confirmed by high plasticity susceptibility in load-displacement curves. However, the sink-in was observed in the micro-matrix and associated with relatively low level of plasticity particularly confined only near the tip of the indenter and large elastic recovery. A number of studies have been carried out to understand pile-up and sink-in behaviors by analyzing the nanoindentation data in terms of the ratio of hardness (H) and indentation modulus (E*) [3133,36–38]. According to those results, for hard materials with high H/E* ratio, greater than 0.1, irrespective of the strain rate sensitivity and work hardening, the material will show sinkin. Perhaps, this is the case for the low susceptibility of plasticity and the high elastic recovery. Also, in this case the plastic strain is expected to be relatively low and highly concentrated near the indenter tip. In order to confirm such a tendency, nanoindentation simulations were performed using the 2D finite element analysis (ABAQUSTM) with the two different material models, one with H/E* = 0.001 and the other with H/E* = 0.1 for the same work hardening. Here, no anisotropy in elasticity and plasticity was considered in the simulation. Fig. 9 shows the simulated plastic strain fields for the two different H/E* ratios. Simulation results predicted the pile-up for H/E* = 0.001 and the sink-in for H/E* = 0.1. Low H/E* ratio is associated with 10
low hardness and/or high modulus. Both low hardness and high modulus promote plasticity even in the early stage of deformation, and the plastic strain is easily spread along the face of the indenter (with low elastic recovery upon unloading), resulting in the pile-up (Fig. 9a). However, high H/E* ratio, which is high hardness and/or low modulus, tends to suppress plasticity and confine the plastic region only near the indenter tip (with high elastic recovery upon unloading), leading to the sink-in (Fig. 9b). Strain fields predicted for the pile-up and sink-in in Fig. 9 seem to be in accordance with nanoindentation deformation behaviors for the GB micro-precipitate (Figs. 4d to 4f, and Figs. 7 and 8) and the micro-matrix (Fig. 5d to 5f, and Figs. 7 and 8), respectively. In fact, elastic moduli measured from nanoindentation tests were not much different between the two regions (225 GPa ± 45 GPa for the GB micro-precipitate and 236 GPa ± 38 GPa for the micro-matrix). According to the H/E* ratio rule, this means that only the hardness difference between the two regions leads to the pile-up and sink-in responses [19,20,39]. Relatively low hardness for the GB micro-precipitate is due to the high susceptibility of plasticity for the A1 (disordered FCC) phase (Fig. 7). High work hardening is also expected for the GB microprecipitate due to solid-solution strengthening by equimolar elements (Fe-Cr-Co rich in Fig. 3). In this case, the deformation constraint induced by the indentation is easily accommodated by local plasticity spread along the faces of the indenter (Figs. 4d, 4g, and Fig. 9a), leading to the formation of selective slip traces (Fig. 4b) near the surface of the pile-up. On the other hand, high hardness for the micro-matrix (Fig. 7) seems to be attributed to the limited slip activity for the B2 (ordered) phase (due to APB strengthening) along with dispersion strengthening of A2 phases. In this case, the plastic strain field is confined near the tip of the indenter (Figs. 5d to f, and Fig. 9b), resulting in the sink-in (Figs. 5b and 5c) with high elastic recovery (Fig. 8).
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4. Conclusions Multiscale microstructural characterization and nanoindentation tests of selected microstructures were carried out on an annealed AlCoCrFeNi multiphase HEA to understand the micromechanics of the deformation involved in the material. The experimental results were analyzed extensively and the following conclusions were drawn. 1. Multiphase microstructures were observed at different length scales. A1 (disordered FCC) phase precipitates (GB micro-precipitate) were distributed along the grain boundary, and the micro-matrix showed topological separations of engraved and embossed regions, which consisted of the B2 (ordered BCC) phase nano-matrix with A2 (disordered BCC) phase nanoprecipitates and the other way around over the region. 2. Nanoindentations on the GB micro-precipitate and the micro-matrix (the B2 nano-matrix with A2 nano-precipitates) showed the pile-up and sink-in morphologies, respectively. The TEM analysis confirmed that the pile-up was associated with higher degree of dislocation activities upon indentation, particularly near the face of the indenter, while the sink-in was due to relatively limited dislocation activities, only concentrated near the indented tip. This seems to be associated with the limited active slip systems for the B2 phase nano-matrix (including APB strengthening) and the presence of A2 phase nano-precipitates as dislocation barriers. 3. Load-displacement curves from nanoindentation tests showed high susceptibility of plasticity for the GB micro-precipitate and larger elastic recovery with relatively low plasticity for the micro-matrix. 4. Nanoindentation simulations using the 2D finite element analysis with the two different material models, high and low hardness-to-indentation modulus ratios, captured the pile-up and sink-in behaviors observed in experiments (even for the plastic strain distribution). Considering comparable indentation modulus values between the GB micro-precipitate and 12
the micro-matrix, the susceptibility of plasticity (viz., readiness of slip) seems to play a key role in determining the pile-up and sink-in in the current case.
Acknowledgements This study was supported financially by Fundamental Research Program of the Korean Institute of Materials Science (KIMS).
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[28] Y.F. Kao, T.J. Chen, S.K. Chen, J.W. Yeh, Microstructure and mechanical property of as-cast, -homogenized, and -deformed AlxCoCrFeNi (0 ??? x ??? 2) high-entropy alloys, J. Alloys Compd. 488 (2009) 57–64. doi:10.1016/j.jallcom.2009.08.090. [29] A. Manzoni, H. Daoud, R. Völkl, U. Glatzel, N. Wanderka, Phase separation in equiatomic AlCoCrFeNi high-entropy alloy, Ultramicroscopy. 132 (2013) 212–215. doi:10.1016/j.ultramic.2012.12.015. [30] W.-L.L.W.-R.R. Wang, W.-L.L.W.-R.R. Wang, J.-W.W. Yeh, Phases, microstructure and mechanical properties of AlxCoCrFeNi high-entropy alloys at elevated temperatures, J. Alloys Compd. 589 (2014) 143–152. doi:10.1016/j.jallcom.2013.11.084. [31] Y. Wang, D. Raabe, C. Klüber, F. Roters, Orientation dependence of nanoindentation pile-up patterns and of nanoindentation microtextures in copper single crystals, Acta Mater. 52 (2004) 2229–2238. doi:10.1016/j.actamat.2004.01.016. [32] T. Csanádi, M. Bl’anda, N.Q. Chinh, P. Hvizdoš, J. Dusza, Orientation-dependent hardness and nanoindentation-induced deformation mechanisms of WC crystals, Acta Mater. 83 (2015) 397–407. doi:10.1016/j.actamat.2014.09.048. [33] K.K. McLaughlin, W.J. Clegg, Deformation underneath low-load indentations in copper, J. Phys. D. Appl. Phys. 41 (2008) 74007. doi:10.1088/0022-3727/41/7/074007. [34] S.B. Luyckx, F.R.N. Nabarro, S.-W. Wai, M.N. James, The anisotropic workhardening of WC crystals, Acta Metall. Mater. 40 (1992) 1623–1627. doi:10.1016/0956-7151(92)90104-M. [35] D. Tabor, A Simple Theory of Static and Dynamic Hardness, Proc. R. Soc. Lond. A. Math. Phys. Sci. 192 (1948) 247–274. http://www.jstor.org/stable/98140. 17
[36] J.-M. Song, Y.-L. Shen, C.-W. Su, Y.-S. Lai, Y.-T. Chiu, Strain Rate Dependence on Nanoindentation Responses of Interfacial Intermetallic Compounds in Electronic Solder Joints with Cu and Ag Substrates, Mater. Trans. 50 (2009) 1231–1234. doi:10.2320/matertrans.M2009016. [37] J. Jin, S.A. Shevlin, Z.X. Guo, Multiscale simulation of onset plasticity during nanoindentation of Al (001) surface, Acta Mater. 56 (2008) 4358–4368. doi:10.1016/j.actamat.2008.04.064. [38] A. Bolshakov, G.M. Pharr, Influences of pileup on the measurement of mechanical properties by load and depth sensing indentation techniques, J. Mater. Res. 13 (1998) 1049–1058. doi:10.1557/JMR.1998.0146. [39] S. Mandal, S. Kose, A. Frank, A.A. Elmustafa, A numerical study on pile-up in nanoindentation creep, Int. J. Surf. Sci. Eng. 2 (2008) 41. doi:10.1504/IJSURFSE.2008.018967.
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Fig 1. (a) As-polished macrostructure of annealed AlCoCrFeNi under polarized light microscopy; (b) SEM micrographs at 1,000X and (c) at 5,000X taken from the central area of (a); (d) SEM micrograph at 100,000X taken from the engraved region of (c), indicated by a thick arrow. Fig 2. (a) Micrograph of an FIB-cut TEM specimen; SADPs showing (b) B2 with A2 peaks at area A, (c) A2 with B2 peaks at area B, (d) FCC peaks at area C and (e) FCC peaks at area D; (f) TEM dark field image taken from areas A and B. Fig 3. (a) TEM bright field image for area B in Fig. 2a and 2f and (b) EDS profiles for the line scan shown in (a); (c) TEM bright field image for area A in Fig. 2a and 2f and (d) EDS profiles for the line scan shown in (c).
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Fig 4. SEM micrographs showing (a) nanoindentations along the GB micro-precipitate and (b) a representative nanoindentation taken from the box indicated in (a); (c) the topology of the representative nanoindentation measured by AFM; (d) TEM bright field image of the through-depth indentation section (FIB cut by dashed lines in (b)); SADPs taken from (e) the GB micro-precipitate (far away from the indentation), (f) area A and (g) area B in (d); (h) TEM birght field image showing the interface between the GB micro-precipitate and the micro-matrix (magnified from a dotted circle (d)). Fig 5. SEM micrographs showing (a) nanoindentations in the micro-matrix and (b) the selected indentation in the engraved region of the micro-matrix taken from the box indicated in (a); (c) the topology of the selected indentation measured by AFM; (d) TEM bright field image of the through-depth indentation section (FIB cut by dashed lines in (b)); SADPs taken from (e) area A, (f) area B and (g) area C in (d). Fig 6. Representative load-displacement curves for the GB micro-precipitate and the micromatrix region having the B2 nano-matrix with A2 nano-precipitate (indents shown in Fig. 4b and Fig. 5b, respetively). Fig 7. The final indentation depth (hfinal, after unloading) versus hardness from nanoindentations on GB micro-precipitates and the micro-matrix. Open and closed symbols indicate indents on the nano-matrix and GB micro-precipitates, respectively. Fig 8. The elastic recovery, Δh, versus hardness from nanoindentations on GB microprecipitates and the micro-matrix. Open and closed symbols indicate indents on the micromatrix and GB micro-precipitates, respectively. Fig 9. Plastic strain distribution predicted from 2D finite element analysis of the nanoindentation with two different material models: (a) H/E* = 0.001 and (b) H/E* = 0.1.
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PII: DOI: Reference:
S0921-5093(17)30497-5 http://dx.doi.org/10.1016/j.msea.2017.04.045 MSA34946
To appear in: Materials Science & Engineering A Received date: 8 March 2017 Accepted date: 12 April 2017 Cite this article as: Gokul Muthupandi, Ka Ram Lim, Young-Sang Na, Jieun Park, Dongyun Lee, Hanjong Kim, Seonghun Park and Yoon Suk Choi, Pile-up and sink-in nanoindentation behaviors in AlCoCrFeNi multi-phase high entropy a l l o y , Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2017.04.045 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Pile-up and sink-in nanoindentation behaviors in AlCoCrFeNi multi-phase high entropy alloy
Gokul Muthupandi1, Ka Ram Lim2, Young-Sang Na2, Jieun Park3, Dongyun Lee4,5, Hanjong Kim6, Seonghun Park6, Yoon Suk Choi1*
1
School of Materials Science and Engineering, Pusan National University, Busan, Korea
2
Titanium Department, Korea Institute of Materials Science, Changwon, Gyeongnam, Korea
3
Department of Cogno-Mechatronics Engineering, Pusan National University, Busan, Korea
4
Department of Nano Fusion Technology, Pusan National University, Busan, Korea
5
Department of Nanoenergy Engineering, Pusan National University, Busan, Korea
6
School of Mechanical Engineering, Pusan National University, Busan, Korea
*
Corresponding author: [email protected]
Abstract Microstructures and nanoindentation behaviors were studied on annealed AlCoCrFeNi high entropy alloy. Both pile-up and sink-in characteristics were found in the grain boundary and grain regions, respectively. The multiple phases present in the AlCoCrFeNi high entropy alloy are the reasons behind the different nanoindentation behaviors, which were identified using electron microscopy. The identified phases showed the grain boundary segregation to have A1 lattice, viz., FCC structure while the grain was distributed with A2 and B2 lattices, 1
viz., BCC and ordered BCC structures, forming the matrix with nano-precipitates of the other. The reason for the pile-up and sink-in is attributed to the dislocation activity in the individual crystal structure: large dislocation activities were found under the pile-up and little dislocation activities under the sink-in, only limited to the indenter tip. Results from a finite element analysis under an isotropic elasto-plastic condition by varying the hardness-tomodulus ratio show that high hardness-to-modulus ratio results in pile-up and the lower ratio results in sink-in. This was associated with the susceptibility to plasticity and the elastic recovery for individual phases of the AlCoCrFeNi high entropy alloy.
Keywords: High entropy alloy, Microstructure, Nanoindentation, Pile-up, Sink-in
1. Introduction Traditionally, alloys have base elements where small quantities of additional elements helped obtain desired properties. The amounts of these alloying additions have been increasing to design alloys for advanced applications. In the past few decades, high entropy alloys (HEAs) with multiple metallic elements of nearly equimolar compositions are being developed [1,2]. The idea is to stabilize solid solution phases and prevent the formation of intermetallics by increasing configurational entropy ΔSconf [3], intended to retain the ductility while possessing the high strength due to distorted lattice structures predominantly of single-phase FCC or BCC. HEA holds many properties which are better than most of the conventional alloys: good thermal stability [4,5], very high hardness [6,7], superior strength [7,8], excellent anticorrosion properties [9–11], and special electrical and magnetic properties [12]. Studies have
2
been extensively carried out under different metallurgical conditions on HEAs based on transition elements [13–16]. However, the tests carried out were to assess mostly mechanical properties, little effort has been made to identify deformation micro-mechanisms associated with microstructures of HEAs. The nanoindentation technique has been often used to study micro-mechanical properties of HEAs. This is partially attributed to the limitation in size of the sample and the convenience in ‘location-specific’ property measurement [17–26]. In this study, in-depth investigation was made to understand the pile-up and sink-in behaviors that we recently observed in AlCoCrFeNi multiphase HEA during the nanoindentation. A self-similarity of the microstructure at different length scales was also studied for AlCoCrFeNi multiphase HEA. In particular, an extensive effort has been made to understand nanoindentation behaviors of multiphase AlCoCrFeNi in line with its responsive microstructural features. This also enabled us to experimentally provide evidences for different dislocation behaviors under sink-in and pile-up.
2. Experimental The multiphase AlCoCrFeNi high entropy alloy used in the current study was produced in a vacuum arc melting furnace (purity of the raw materials used was 99.99%). The alloy was remelted multiple times to ensure chemical homogeneity. The alloy was suction-cast into rods of 6 mm diameter and 50 mm long using a copper mould. Specimens of 5 mm thickness were cut and annealed at 1100 °C for 2 hours. Microstructural characterization was carried out using a light optical microscope (Nikon Eclipse MA200) under the polarized light. Further characterization of the microstructure and nanostructure was carried out using a scanning electron microscope, SEM (TESCAN MIRA3, 15kV) and a transmission electron microscope, TEM (FEI Tecnai F20, 200kV). The phase analysis was carried out by studying the selected area diffraction patterns, SADP from the TEM. The compositional analysis was 3
carried out using energy dispersive spectroscopy, EDS in both the SEM and TEM. The mechanical behavior of the HEA was investigated using nanoindentation (Agilent G200) with a Berkovich tip (tip radius 50 nm). The maximum load was set to 3mN. The load was optimized after extensive testing under different loading conditions to obtain an indentation size where the distance between the neighboring phases was considerably away. The loaddisplacement data was analyzed using Oliver-Pharr method [17,18] to determine the hardness and modulus. A minimum of 10 tests were carried out in each phase to understand the behavior. The indentations were imaged using SEM and the residual surface topology of the indentation was mapped using atomic force microscopy, AFM. Indentation marks with pileup and sink-in topologies were separated. Focused ion beam (FIB) cut specimens of 100 nm thick were made across the indents for the TEM analysis to study the dislocation behavior beneath the pile-up and sink-in. Numerical analysis using finite element method (ABAQUSTM) was also carried out to interpret the pile-up and sink-in behaviors of materials with different stress-strain responses. A 2D isotropic model was used with a sharp conical indenter with a half angle of 70.3° which simulates the area function of a Berkovich tip of the nanoindentation.
3. Results and Discussion Earlier studies on the equiatomic AlCoCrFeNi HEA found that the as-cast microstructure has a dendritic structure [3]. Fig. 1a shows the annealed (1100 °C for 2 hours) macrostructures (under the as-polished condition) of the specimen, which shows no change in macrostructure. This is attributed to the distorted lattice and the sluggish diffusion in the high entropy alloys [2]. The as-cast structure was studied to have dendritic structure, which in turn have substructures with the matrix phase rich in Al-Ni and precipitates rich in Fe-Cr [8,24,25,27– 30]. This phase separation was known to be due to the spinodal decomposition of the phases 4
[27–30]. The annealed microstructure (Fig. 1a) shows further decomposition of the alloy phase during the heat treatment resulting in precipitation and dispersion of the phases at different length scales, as shown in Fig. 1b to 1d. At the micro-scale shown in Fig. 1b and 1c, the main features are the precipitates distributed along the grain boundary (hereafter shall be denoted as “GB micro-precipitate”) and the matrix (hereafter shall be denoted as “micromatrix”). Here, the micro-matrix consists of topological separations of engraved and embossed regions (more clearly seen in Fig. 1c). These regions were identified to posses different chemical compositions (The details will be discussed later). Fig. 1d is the submicron scale microstructure taken from the engraved region (indicated by a thick arrow) of the micro-matrix in Fig. 1c. This higher-magnification microstructure shows the fine dispersion of the nanoscaled precipitates (will be noted as “nano-precipitates”) in the matrix (will be referred to as “nano-matrix” hereafter). The embossed region of the micro-matrix in Fig. 1c also showed similar microstructure is the sub-micron scale (will be discussed in the following section). The TEM analysis was performed for annealed AlCoCrFeNi in order to identify different phases at different length scales (observed in Fig. 1), and the results are shown in Fig. 2. Fig. 2a shows the microstructure taken from the TEM specimen. In the figure, the micro-matrix and GB micro-precipitate regions were indicated using the notation in Fig. 1b and 1c. Fig. 2b to 2e are the SADPs taken from local areas A to D in Fig. 2a. In Fig. 1c, the micro-matrix consisted of topologically engraved and embossed regions. These two different regions are marked as A and B in Fig. 2a. The SADP analysis from areas A and B (Fig. 2b and 2c, respectively) indicated that these two regions in the micro-matrix have mixtures of A2 and B2 lattice structures with different mixing ratios such that the one consists of A2 nanoprecipitates in the B2 nano-matrix while the other with B2 nano-precipitates in the A2 nanomatrix, resulting in the sub-micron scale microstructures shown in Fig. 1d (through SEM) 5
and Fig. 2f (through TEM). The EDS line scan result (in Fig. 3) for areas A and B in Fig. 2a and 2f shows that the A2 phase is rich in Fe-Cr and the B2 phase is rich in Ni-Al with a homogenous distribution of Co. The GB micro-precipitate (area C) of Fig. 2a was identified to have an A1 (disordered FCC) lattice structure (from the SADP in Fig. 2d). Interestingly, the micro-precipitate located at area D in Fig. 2a also showed the same lattice structure (identified from the SADP in Fig. 2e). Perhaps, the GB micro-precipitate (area C) is threedimensionally interconnected with the micro-precipitate in D (although this needs further clarification). The A1 phase in areas C and D was found to be rich in Fe-Cr-Co. In order to understand the mechanical behavior of the individual phases observed, nanoindentation tests were carried out. Figs. 4 and 5 show nanoindentation responses of the GB micro-precipitate and the micro-matrix (particularly for a region of a B2 nano-matrix with A2 precipitates), respectively. Fig. 4a identifies the location of the indentation on the GB micro-precipitate observed through SEM. Fig. 4b shows the SEM image of a representative nanoindentation (identified with a box in Fig. 4a) on the GB micro-precipitate. Here, slip traces near the edges of the indentation are observed mainly in the pile-up regions. The topology of the area of Fig. 4b measured by the AFM clearly shows the pile-up along the two edges of the indentation, as indicated by arrows in Fig. 4c. The pile-up along only two edges seems to be attributed to the anisotropy of the elastic-plastic behavior of the GB microprecipitate, which has the disordered FCC (A1) structure as confirmed in Fig. 2e [31,34]. TEM analysis to understand the dislocation behavior underneath the indentation of the GB micro-precipitate was carried out. The sample was cut across the indentation as indicated in Fig. 4b (with dashed lines). The TEM bright field image for the through-thickness section of the indentation was then taken and shown in Fig. 4d. From Fig. 4d, the indentation was confirmed to be through the GB micro-precipitate (FCC phase). It was also found from Fig. 4e that the micro-matrix (having the B2 nano-matrix with A2 nano-precipitates) is located 6
right underneath the GB micro-precipitate. Fig. 4f and 4g are SADPs taken from area A, which is away from the edge of the indenter, and area B, which is adjacent to the face of the indenter (showing pile-up above), in Fig. 4d, respectively. These SADPs indicate the different amount of dislocation activities between areas A and B underneath the indentation. Area B, which is near the face of the indenter, shows larger dislocation activities than those in area A, which is away from the edge of the indenter. This seems to be due to differences in magnitude and state of stresses between the two areas under the indentation. In particular, heavy dislocation activities in area B seem to result in the pile-up on the surface near the edge of the indentation (Fig. 4b and 4c). Fig. 4h is the TEM bright field image showing the dislocation interaction at the interface between the GB micro-precipitate (FCC phase) and the micro-matrix (B2 nano-matrix with A2 nano-precipitates), as indicated by a dotted circle in Fig. 4d. Here, the interface seems to act as a barrier against dislocation gliding from the GB micro-precipitate to the micro-matrix, as it appears that only selected (perhaps, energetically favorable) dislocations are activated in the micro-matrix region near the interface. Fig. 5a is the SEM micrograph showing nanoindentations carried out on the micro-matrix (Fig. 1c). Among these indentations one indentation was selected (indicated with a box in Fig. 5a) in an engraved region of the micro-matrix for the further analysis of the microstructure and topology using AFM and TEM, respectively. The SEM image of the selected indentation in Fig. 5b and the corresponding topology map in Fig. 5c show the sinkin near the indentation. It was also confirmed that this micro-matrix region consists of the nano-precipitates in the nano-matrix, as previously observed in Fig. 1d. Again, the throughdepth TEM analysis was carried out across the indentation (the FIB cut indicated by the dashed lines in Fig. 5b), and the resulting TEM bright field image is shown in Fig. 5d. Here, the indented tip location is indicated by an arrow in the figure. The SADP analysis (Fig. 5e) 7
for the region adjacent to the face of the indenter (area A in Fig. 5d) indicated that the indented micro-matrix region consists of the B2 nano-matrix and A2 nano-precipitates. The further SADP analysis (Figs. 5f and 5g) for areas B and C in Fig. 5d showed that another micro-matrix region having the A2 nano-matrix with B2 nano-precipitates is located underneath the indented micro-matrix region (having the B2 nano-matrix with A2 nanoprecipitates). The boundary between the two micro-matrix regions is indicated by the dashed line in Fig. 5d. The TEM image in Fig. 5d and the SADPs in Figs. 5e and 5f indicate that the indentation in the micro-matrix region induces the lower level of dislocation activities underneath the indentation, compared to the indentation in the GB micro-precipitate (Figs. 4d to f). This limited plasticity for the micro-matrix is probably due to the limited availability of active slip systems for the B2 phase nano-matrix and the presence of A2 phase nano-precipitates, which act as dislocation barriers. Also, the formation of antiphase boundaries (APB) and stacking faults in the B2 (ordered) phase is believed to limit the slip activity. Under this limited plasticity in the micro-matrix, dislocation activities seem to be localized underneath the indented tip (area B in Fig. 5d), while little dislocation activities were found adjacent to the face of the indenter (area A in Fig. 5d). In particular, the boundary between the two micromatrix regions does not seem to be a strong barrier for the dislocation transmittance, as confirmed by no contrast change across the boundary in the TEM bright field image (area B in Fig. 5d). This seems to be associated with the similar microstructure between the two micro-matrix regions (the B2 nano-matrix with A2 nano-precipitates and the other way around) and lattice dimensions of the B2 and A2 phases. It was confirmed from Figs. 4 and 5 that the GB micro-precipitate and the micro-matrix have different crystal structures and microstructural features (A1 (FCC) phase for the former and the B2 nano-matrix with A2 nano-precipitates (or the other way around) for the latter), and 8
revealed the pile-up and sink-in responses upon nanoindenation, respectively. Dislocation activities near and underneath the indentation were also different between the two cases. In order to further investigate this difference, load-displacement curves after nanoindentation tests were systematically analyzed and compared between the two cases. Fig. 6 shows the representative load-displacement curves (corresponding to nanoindentations in Figs. 4b and 5b) for the GB micro-precipitate and the micro-matrix after the load-controlled (3 mN) nanoindentation tests. The serrations (or pseudo pop-in) observed in loading curves for both cases are signs of the evolution of dislocation barriers, which impede the dislocation motion and increase the energy required for deformation to continue. This seems to be associated with solid solution strengthening for the GB micro-precipitate and dispersion strengthening of nano-precipitates (in addition to APB strengthening) for the micro-matrix. The unloading curve shows the similar recovery behavior between the two cases. In Fig. 6, however, the indentation depth was deeper in the GB micro-precipitate than in the micro-matrix for the same applied load (3mN). In order to perform statistically more reliable analysis for this nanoindentation response, a series of nanoindentation tests was performed over the different regions of the specimen, and the resulting (final) indentation depth (hfinal) values were compared between the two cases in Fig. 7. Different symbols in Fig. 7 mean different regions for nanoindentation tests. Here, hfinal values were plotted as a function of the hardness since the hardness is inclusive of plasticity and elasticity, also the anisotropy does not significantly affect the hardness [21,35]. The hfinal-hardness plot in Fig. 7 shows an inverse relation for both cases, which is understandable because the hardness is the resistance to the deformation under the indentation or abrasion. In particular, the GB micro-precipitate shows lower hardnesses and higher hfinal values, compared to the micro-matrix. This indicates that the GB micro-precipitate is more susceptible to the plastic deformation than the micro-matrix. Using the same data in Fig. 7, the amount of the elastic recovery (Dh) was 9
determined from load-displacement curves and plotted as a function of the hardness in Fig. 8. The result showed higher elastic recovery in the micro-matrix than in the GB microprecipitate. Results from Figs. 7 and 8 imply that the nanoindentation deformation of the micro-matrix is relatively elastic in nature, compared to the GB micro-precipitate. Combining results from Figs. 4 to 8, distinctive nanoindentation characteristics were identified between the GB micro-precipitate (FCC phase) and the micro-matrix (the B2 nanomatrix with A2 nano-precipitates) for the heat-treated AlCoCrFeNi multiphase HEA. The pile-up observed in the GB micro-precipitate was accompanied by heavy plasticity (underneath the indentation) particularly near the face of the indenter, which was confirmed by high plasticity susceptibility in load-displacement curves. However, the sink-in was observed in the micro-matrix and associated with relatively low level of plasticity particularly confined only near the tip of the indenter and large elastic recovery. A number of studies have been carried out to understand pile-up and sink-in behaviors by analyzing the nanoindentation data in terms of the ratio of hardness (H) and indentation modulus (E*) [3133,36–38]. According to those results, for hard materials with high H/E* ratio, greater than 0.1, irrespective of the strain rate sensitivity and work hardening, the material will show sinkin. Perhaps, this is the case for the low susceptibility of plasticity and the high elastic recovery. Also, in this case the plastic strain is expected to be relatively low and highly concentrated near the indenter tip. In order to confirm such a tendency, nanoindentation simulations were performed using the 2D finite element analysis (ABAQUSTM) with the two different material models, one with H/E* = 0.001 and the other with H/E* = 0.1 for the same work hardening. Here, no anisotropy in elasticity and plasticity was considered in the simulation. Fig. 9 shows the simulated plastic strain fields for the two different H/E* ratios. Simulation results predicted the pile-up for H/E* = 0.001 and the sink-in for H/E* = 0.1. Low H/E* ratio is associated with 10
low hardness and/or high modulus. Both low hardness and high modulus promote plasticity even in the early stage of deformation, and the plastic strain is easily spread along the face of the indenter (with low elastic recovery upon unloading), resulting in the pile-up (Fig. 9a). However, high H/E* ratio, which is high hardness and/or low modulus, tends to suppress plasticity and confine the plastic region only near the indenter tip (with high elastic recovery upon unloading), leading to the sink-in (Fig. 9b). Strain fields predicted for the pile-up and sink-in in Fig. 9 seem to be in accordance with nanoindentation deformation behaviors for the GB micro-precipitate (Figs. 4d to 4f, and Figs. 7 and 8) and the micro-matrix (Fig. 5d to 5f, and Figs. 7 and 8), respectively. In fact, elastic moduli measured from nanoindentation tests were not much different between the two regions (225 GPa ± 45 GPa for the GB micro-precipitate and 236 GPa ± 38 GPa for the micro-matrix). According to the H/E* ratio rule, this means that only the hardness difference between the two regions leads to the pile-up and sink-in responses [19,20,39]. Relatively low hardness for the GB micro-precipitate is due to the high susceptibility of plasticity for the A1 (disordered FCC) phase (Fig. 7). High work hardening is also expected for the GB microprecipitate due to solid-solution strengthening by equimolar elements (Fe-Cr-Co rich in Fig. 3). In this case, the deformation constraint induced by the indentation is easily accommodated by local plasticity spread along the faces of the indenter (Figs. 4d, 4g, and Fig. 9a), leading to the formation of selective slip traces (Fig. 4b) near the surface of the pile-up. On the other hand, high hardness for the micro-matrix (Fig. 7) seems to be attributed to the limited slip activity for the B2 (ordered) phase (due to APB strengthening) along with dispersion strengthening of A2 phases. In this case, the plastic strain field is confined near the tip of the indenter (Figs. 5d to f, and Fig. 9b), resulting in the sink-in (Figs. 5b and 5c) with high elastic recovery (Fig. 8).
11
4. Conclusions Multiscale microstructural characterization and nanoindentation tests of selected microstructures were carried out on an annealed AlCoCrFeNi multiphase HEA to understand the micromechanics of the deformation involved in the material. The experimental results were analyzed extensively and the following conclusions were drawn. 1. Multiphase microstructures were observed at different length scales. A1 (disordered FCC) phase precipitates (GB micro-precipitate) were distributed along the grain boundary, and the micro-matrix showed topological separations of engraved and embossed regions, which consisted of the B2 (ordered BCC) phase nano-matrix with A2 (disordered BCC) phase nanoprecipitates and the other way around over the region. 2. Nanoindentations on the GB micro-precipitate and the micro-matrix (the B2 nano-matrix with A2 nano-precipitates) showed the pile-up and sink-in morphologies, respectively. The TEM analysis confirmed that the pile-up was associated with higher degree of dislocation activities upon indentation, particularly near the face of the indenter, while the sink-in was due to relatively limited dislocation activities, only concentrated near the indented tip. This seems to be associated with the limited active slip systems for the B2 phase nano-matrix (including APB strengthening) and the presence of A2 phase nano-precipitates as dislocation barriers. 3. Load-displacement curves from nanoindentation tests showed high susceptibility of plasticity for the GB micro-precipitate and larger elastic recovery with relatively low plasticity for the micro-matrix. 4. Nanoindentation simulations using the 2D finite element analysis with the two different material models, high and low hardness-to-indentation modulus ratios, captured the pile-up and sink-in behaviors observed in experiments (even for the plastic strain distribution). Considering comparable indentation modulus values between the GB micro-precipitate and 12
the micro-matrix, the susceptibility of plasticity (viz., readiness of slip) seems to play a key role in determining the pile-up and sink-in in the current case.
Acknowledgements This study was supported financially by Fundamental Research Program of the Korean Institute of Materials Science (KIMS).
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Fig 1. (a) As-polished macrostructure of annealed AlCoCrFeNi under polarized light microscopy; (b) SEM micrographs at 1,000X and (c) at 5,000X taken from the central area of (a); (d) SEM micrograph at 100,000X taken from the engraved region of (c), indicated by a thick arrow. Fig 2. (a) Micrograph of an FIB-cut TEM specimen; SADPs showing (b) B2 with A2 peaks at area A, (c) A2 with B2 peaks at area B, (d) FCC peaks at area C and (e) FCC peaks at area D; (f) TEM dark field image taken from areas A and B. Fig 3. (a) TEM bright field image for area B in Fig. 2a and 2f and (b) EDS profiles for the line scan shown in (a); (c) TEM bright field image for area A in Fig. 2a and 2f and (d) EDS profiles for the line scan shown in (c).
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Fig 4. SEM micrographs showing (a) nanoindentations along the GB micro-precipitate and (b) a representative nanoindentation taken from the box indicated in (a); (c) the topology of the representative nanoindentation measured by AFM; (d) TEM bright field image of the through-depth indentation section (FIB cut by dashed lines in (b)); SADPs taken from (e) the GB micro-precipitate (far away from the indentation), (f) area A and (g) area B in (d); (h) TEM birght field image showing the interface between the GB micro-precipitate and the micro-matrix (magnified from a dotted circle (d)). Fig 5. SEM micrographs showing (a) nanoindentations in the micro-matrix and (b) the selected indentation in the engraved region of the micro-matrix taken from the box indicated in (a); (c) the topology of the selected indentation measured by AFM; (d) TEM bright field image of the through-depth indentation section (FIB cut by dashed lines in (b)); SADPs taken from (e) area A, (f) area B and (g) area C in (d). Fig 6. Representative load-displacement curves for the GB micro-precipitate and the micromatrix region having the B2 nano-matrix with A2 nano-precipitate (indents shown in Fig. 4b and Fig. 5b, respetively). Fig 7. The final indentation depth (hfinal, after unloading) versus hardness from nanoindentations on GB micro-precipitates and the micro-matrix. Open and closed symbols indicate indents on the nano-matrix and GB micro-precipitates, respectively. Fig 8. The elastic recovery, Δh, versus hardness from nanoindentations on GB microprecipitates and the micro-matrix. Open and closed symbols indicate indents on the micromatrix and GB micro-precipitates, respectively. Fig 9. Plastic strain distribution predicted from 2D finite element analysis of the nanoindentation with two different material models: (a) H/E* = 0.001 and (b) H/E* = 0.1.
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GB Microprecipitates
Micro-matrix
c d
c
e
b
d
f
A B2 matrix + A2 particles
B
A2 matrix + B2 particles
c
a
Counts Counts
b
100
200
300
400
500
600
200
400
600
800
1000
100
100
200
200
300
Position (nm)
Position (nm)
300
400
400
500
Ni
Co
Fe
Cr
Al
Ni
Co
Fe
Cr
Al
a Slip traces
(μm)
c
(μm)
Pile-up
(nm)
A
B
GB micro-ppt
Micro-matrix (B2 matrix + A2 ppt.)
d
Pile-up
e
f
g
100 nm
h
a A2 nano-precipitates
B2 nano-matrix
b
C
B
Micro-matrix (B2 matrix + A2 ppt.) Micro-matrixx (A2 matrix + B2 ppt.)
d
A
e
f
a
b