Crystalline materials are almost always stronger when plastic deformation is confined to smaller volumes. Such size effects on plasticity can be understood either in terms of the gradients of strain that can accompany deformation in small volumes or in terms of the microstructural constraints on dislocation motion. But recent experiments have shown that single crystals in the micrometer size range show significant size effects, with smaller being stronger, even when the crystals are tested in uniaxial compression. Plastic deformation in these tests is unconstrained in the sense that dislocations are free to leave the crystal during deformation and no significant gradients of strain are imposed. Other experiments involving both thin films on substrates and nearly free standing thin films show very similar size effects. These size effects suggest that plasticity is dislocation source-controlled, wherein smaller volumes are stronger because fewer sources of dislocations are available. The evidence for source-controlled plasticity is reviewed and recent attempts to describe the corresponding size effects on strength are described. One key idea is that for very small crystals dislocations leave the crystal more frequently than they multiply, forcing other, harder, sources of dislocations to be activated. In the extreme, dislocations have to be nucleated either at the surface or in the bulk of the crystal, which requires very high stresses; if the availability of these sources scales with the sample size, as statistical analyses have shown, then the observed size effects might be explained. Support for this picture is provided by recent in situ TEM observations showing that plastically deformed nickel pillars are completely free of dislocations during and after deformation. Also, gold crystals show no x-ray line broadening even after ~25% strain. Progress in developing these ideas will be described.

Quasi-one-dimensional (1D) nanostructures (nanowires, nanobelts and nanorods) are the forefront nanomateirals for nanotechnology. Oxide nanostructures have been synthesized for a wide range of semiconducting oxides [1] that are potential building blocks for constructing numerous nanodevices. Using the technique demonstrated for measuring the mechanical properties of nanotubes [2,3], the mechanical and field emission properties of the oxide nanobelts have been characterized. Field effect transistors [4], ultra-sensitive nano-size gas sensors [5], nanoresonators and nanocantilevers [6] have been fabricated using nanobelts. Among all of the oxide nanostructures we have investigated, ZnO is very unusual. The two important characteristics of the wurtzite structured ZnO are the non-central symmetry and the polar surfaces. The structure of ZnO can be described as a number of alternating planes composed of tetrahedrally coordinated O2- and Zn2+ ions, stacked alternatively along the c-axis. The oppositely charged ions produce positively charged (0001)-Zn and negatively charged (000-1)-O polar surfaces, resulting in a normal dipole moment and spontaneous polarization along the c-axis. The polar surfaces give raise a few interesting growth features, such as the formations of nanosprings [7], nanorings [8], nanobows [9] and nanohelices [10]. These nanostructure are semiconductive and piezoelectric and have potential applications as nano-scale sensors, traducers, and actuators. This presentation will be about the synthesis, characterization and potential applications of these novel nanostructures [11]. [1] Z.W. Pan, Z.R. Dai and Z.L. Wang, Science, 209 (2001) 1947.[2] P. Poncharal, Z.L. Wang, D. Ugarte and W.A. de Heer, Science, 283 (1999) 1513.[3] R.P. Gao, Z.L. Wang, Z.G. Bai, W. de Heer, L. Dai and M. Gao, Phys. Rev. Letts., 85 (2000) 622.[4] M. Arnold, P. Avouris, Z.L. Wang,. Phys. Chem. B, 107 (2002) 659. [5] E. Comini, G. Faglia, G. Sberveglieri, Zhengwei Pan, Z. L. Wang Applied Physics Letters, 81 (2002) 1869.[6] W. Hughes and Z.L. Wang, Appl. Phys. Letts., 82 (2003) 2886.[7] X.Y. Kong and Z.L. Wang, Nano Letters, 2 (2003) 1625 + cover.[8] X.Y. Kong, Y. Ding, R.S. Yang, Z.L. Wang, Science, 303 (2004) 1348.[9] W.L. Hughes and Z.L. Wang, J. Am. Chem. Soc., 126 (2004) 670[10] P.X. Gao, Y. Ding, W.J. Mai, W.L. Hughes, C.S. Lao and Z.L. Wang, Science, 309 (2005) 1700.[11] Thanks the support from NSF, DARPA, NASA and Airforce.[12] for details see: http://www.nanoscience.gatech.edu/zlwang/

Recent advances in miniaturized mechanical testing of materials with small length scales have created considerable interest on exploring the origins of size-effects at micron and submicron dimensions. Especially, the underlying dislocation mechanisms determining the strength of thin foils, whiskers and single-crystals, already mentioned in a short paragraph of the seminal book Theory of Crystal Dislocations by F.R.N. Nabarro, are now of central importance with the advent of nanomaterials, nanostructured coatings, miniaturized sensors, micro-electromechanical systems (MEMS), semiconductor devices, etc., in modern technological applications.In this talk in situ tensile testing of wires with diameters between 50µm and 0.5µm and of films with thicknesses down to 40 nm are reported. The in situ experiments performed in the scanning electron microscope revealed that the aspect ratio of gauge length to side length becomes very important at the micron length scales. At small aspect ratios dislocation pile-ups form and cause a significant increase in flow stress, while for a large aspect ratio the size dependence of the flow stress is weak and attributed to the size of the dislocation sources governing plastic deformation. The in situ tensile testing of 40 nm thick single crystal Au films revealed dislocation driven plasticity. However, in contrast to film thicknesses above ca. 100 nm deformation is carried by partial dislocations rather then perfect dislocations. At the late stages of deformation prior to fracture, dislocations start to glide on the (001) planes. Acknowledgements: Significant contributions from D. Kiener, R. Pippan, S.H. Oh (Leoben), M. Legros (Toulouse), and P. Gruber (Stuttgart), R. Spolenak (Stuttgart/Zurich) are acknowledged.

One of the clear impacts of Nabarro’s work is the understanding that the effect of dislocations on the properties of materials (electrical, mechanical, etc) depends on their interaction with both extended defects, such as surfaces and interfaces, and point defects, such as vacancies and impurities (either substitutional or interstitial). The difficulty for the experimentalist is that often these factors overlap, requiring characterization methods that can observe the atomic structure of the dislocations, quantify the elemental composition in and around the cores with ~single atom sensitivity, and determine the location of the dislocations within the overall material or device. The ability to characterize dislocations on this level is afforded by the combination of atomic resolution Z-contrast imaging and electron energy loss spectroscopy (EELS) in the scanning transmission electron microscope (STEM). The incoherent Z-contrast image allows the atomic structure of the dislocation core to be determined directly and compositional variations to be identified from contrast variations. Using the image as a guide to position the electron beam, the integrated intensities of the characteristic core-loss edges in the EEL spectrum permit vacancies and impurities to be identified and quantified, while the fine structure of these edges allows the local electronic structure to be observed. Both of these capabilities have improved significantly in recent years with the technical innovations of aberration correction (higher spatial resolution) and monochromation (higher energy resolution). Here, results will be presented from III-V and II-VI semiconductor systems that highlight the ability of the STEM techniques to identify the presence of impurities and vacancies in and around interfaces and dislocation cores and understand both the final electronic properties and how the structures developed during growth. In addition, the benefits and drawbacks of the improved sensitivity of aberration corrected and monochromated STEM methods will be described, with the potential directions for the future analysis of dislocations discussed.

Low angle grain boundary is known to consist of an array of dislocations. Recently, the studies concerning the low angle grain boundaries have been performed to study dislocation core structures in oxide ceramics. It has been also found using high-resolution transmission electron microscopy (HRTEM) that the lattice dislocations in oxide ceramics often dissociate into partial dislocations. In the case of alumina, we previously reported that the boundary dislocations on low angle tilt grain boundaries dissociated into two partials dislocations with the narrow separation distance of a few nanometers and the separation distance decreased with increasing tilt angle. In this study, alumina bicrystals with the low angle tilt boundaries including a slight twist component were fabricated using diffusion bonding technique, and then the resulting grain boundary structures were analyzed by HRTEM and WBDF to investigate the structure transitions of boundary dislocations by a slight twist component.The grain boundary with pure 2 degree tilt shows dot-like contrast arrays along the boundary with equal separation distance. Averaged separation distance in this case was about 13.2 nm. These contrasts were expected to appear due to the edge dislocations along the grain boundary since the tilt angle should be compensated. On the other hand, WBDF image of the boundary with 2 degree tilt component and a slight twist component showed that the contrasts in groups of five appear in addition to the array of pair contrasts with regular separation distance. HRTEM observations show that the contrasts in groups of five are due to the partial dislocations. From the Burgers circuit analyses, each contrast is reflected by a dislocation which has the lattice discontinuity of 1/6[11-20] in the [1-100] projection. Considering the [0001] view of alumina crystal with corundum structure, the grain boundary was found to consist of the array of partial dislocations, but two different structural regions were observed. One region was an area including the pairs of b=1/3<10-10> and b=1/3<01-10> partial dislocations with a stacking-fault in between. The other region was an area composed of the unique multiple-partial-structures. Systematic observations revealed that three to seventeen multiple-partial-structures were formed along the boundary in addition to the groups of five partials. The dark-field TEM image contrasts can be used to distinguish the Burgers vectors of the partials, and the structure of the multiple-partial-structures are consistently identified by the image contrast analysis. The multiple-partial-structures have n/6[11-20] edge component and the constant value of screw component, 1/2[-1100]. The screw angle of the bicrystal was estimated to be about 0.1°. It is concluded that the multiple-partial-structures should be introduced to compensate a very small amount of twist component of the boundary that was introduced during the fabrication of the bicrystal.