The bulk materials were prepared by using standard metallographic methods, so that the samples could be glued on SEM specimen mounts. Bulk samples were carefully cut by using a diamond wire saw (Walter Ebner GmbH, Germany). They were then hand ground with silicon carbide particles of decreasing grit size down to 5 µm using a custom-made stainless steel cylindrical holder that was employed to produce samples. Subsequently, the samples were electropolished to reduce the damage produced by the harsh cutting and grinding procedures, and to obtain a flat surface needed for the later microcompression experiments. The electropolishing conditions used for each of the samples are summarized in the following table (Table 3.1). Table .: Electropolishing conditions of the different samples. Sample Parameter W Solution Voltage Time 97% distilled and 3% NaOH 20 V 40 sec Ta Solution Voltage Time 95% CH3OH and 5% H2SO4 34 V 30 s β-CuZn Solution Voltage Time 60 wt% and distilled water 1.5 V 4 min NiAl Solution Voltage Time Struers electrolyte A2-I and A2-II (C2H6O, C6H14O2, HClO4, H2O) 33 V 90 s The W and Ta samples were mounted onto SEM using silver-based electrically and thermally conductive paste (Pyro-Duct 597-A, Aremco Products, New York, NY), which is suitable for elevated-temperature usage. The bonded arrangement was then cured at 70 °C and degassed in vacuum in a furnace (Gero Hochtemperaturöfen, Neuhausen, Germany) to avoid additional hydrocarbon contamination of the SEM used to carry out further in situ compression tests at high temperatures. The β-CuZn and NiAl samples were glued onto an aluminum holder using conventional silver paste (Acheson Henkel Silver DAG 1415, Germany), which is widely used for electron microscopy purposes.
(only system observed for 200 nm diam. pillars) 0.5 0.47 0.28 0.44 0.42 46 43 12 47 47 2 47 16 37 40 Ill-defined slip traces 0.3 18 18 33 21 Ill-defined slip traces 0.5 48 1 Stress-strain response Stress-strain behavior of representative β-CuZn compression pillars with diameters ranging from 5 µm to 200 nm are shown in Figure 5.5 as a function of crystal orientation. These stress-strain curves have distinctive features typical of micropillar deformation. During the initial loading (well under 1% strain), relatively low and inconsistent stiffness values are recorded. This is thought to be caused by pillar surface features and initial seating of the indenter on the pillar [87]. Quickly the stiffness increases and the pillar undergoes elastic loading, until a critical point is reached and strain increases sharply under a constant applied load. Interestingly, such strain bursts remained prevalent for relatively large diameters. This is in contrast with previous studies on various FCC and BCC metals that have shown decreasing strain bursts for pillar diameters of 2 µm or larger [6,7,84,97,135]. Nevertheless, bulk deformation of β-CuZn has demonstrated ‘jerky’ or ‘serrated’ flow stress, somewhat resembling the strain bursts observed in micropillar testing [48,49]. Figure .: Representative stress-strain behavior of β-CuZn compression pillars. The curves fundamentally change shape with decreasing diameter. Figure 5.6a shows the critical resolved shear stress (CRSS) values at 2.5% strain for all samples tested. A 2.5% strain value was chosen because it is beyond the first observed strain burst, yet still low enough to be minimally affected by strain hardening. Although
We studied the compressive behavior of single crystalline NiAl micropillars at room temperature in to determine the impact of sample size on the plastic behavior of such a semi-brittle material. Overall, more than 200 pillars with different crystal orientations and aspect ratios were tested at different compression rates. The following was found: The -, - and -oriented micropillars exhibited very different plastic response as a consequence of the activation of different slip systems: for -oriented micropillars and - and -oriented micropillars. In contrast to larger pillars, 200 nm diameter pillars deformed in and slip steps irrespective of crystal orientation. This might be caused by surface dislocation source activation promoted at this small size scale. Constant strain rate tests on pillars with different aspect ratios compressed along the and axes showed that kinking also takes place at the microscale as a result of crystal misorientation; -oriented micropillars deformed by slip in while micropillars deformed by kink banding. Very high ductility was obtained for the compressed micropillars in contrast to the semi-brittle bulk specimens. This might be due to surface effects such as Ga ion damage as well as the fact that, at the (sub-) micron scale, plasticity is dominated by dislocation source activation rather than dislocation-dislocation interaction (thus minimizing deformation induced embrittlement effects). The strain-stress response of the pillars evidenced the plastic behavior of highly anisotropic materials, showing the effect that crystal orientation has in the overall mechanical response. More importantly, plastic anisotropy, defined as the ratio of the critical resolved shear stresses between hard and soft crystals, strongly decreased with decreasing sample size. The reduced plastic anisotropy might be attributed to a decrease in the intrinsic resistance of the specimens at the sub-micron scale (surface effects). Very dissimilar size effects were observed as a function of crystal orientation. Soft crystals exhibited a strong size effect, similar to that observed in FCC metals. By contrast, hard crystals exhibited a diminished size effect, similar to