Direct Atomic-Scale Imaging about the Mechanisms of Ultralarge Bent

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Direct Atomic-Scale Imaging about the Mechanisms of Ultralarge Bent Straining in Si Nanowires Lihua Wang,† Kun Zheng,† Ze Zhang,†,‡,§ and Xiaodong Han*,† † ‡

Institute of Microstructure and Properties of Advanced Materials, Beijing University of Technology, Beijing, 100022, China Department of Materials Science and §State Key Laboratory of Silicon Materials, Zhejiang University, Hangzhou, 310008, China

bS Supporting Information ABSTRACT: To safely and reliably use nanowires (NWs) for exploring new functions for different nanodevices, the mechanical properties and structural evolution of the nanowires under external stress become highly important. Large strain (up to 14%) bending experiments of Si NWs were conducted in a highresolution transmission electron microscope at atomic resolution. The direct dynamic atomic-scale observations revealed that partial and full dislocation nucleation, motion, escape, and interaction were responsible for absorbing the ultralarge strain of up to 14% in bent Si nanowires. The prevalent full dislocation movement and interactions induced the formation of Lomer lock dislocations in the Si NWs. Finally, in contrast to the unlock process of Lomer dislocations that can happen in metallic materials, we revealed that the continuous straining on the Lomer dislocations induced a crystal
amorphous (c-a) transition in Si NWs. Our results provide direct explanation about the ultralarge straining ability of Si at the nanometer scale. KEYWORDS: Si nanowires, ultralarge bent strain, atomic scale, Lomer dislocation, crystal
amorphous transition

S

emiconductor nanowires (NWs) have attracted intensive attention because of their many applications in nanotechnology, such as in nanoscale transistors,1
3 sensors,4,5 actuators,6 nanogenerators,7,8 power sources,9 and battery anodes.10 Si NWs have been one of the most important types of NWs since their successful fabrication in 1998.11,12 To safely and reliably use NWs and to explore new functions for different kinds of nano devices, the mechanical properties and structural evolution of these NWs when under external stress is very important. In particular, NWs display novel properties and functions when they are under bending strain.7,8,13
16 For the past decade, the mechanical properties of Si NWs have been studied by many newly developed methods in atomic force microscopy (AFM), scanning electron microscopy (SEM), and transmission electron microscopy (TEM). Experimental results show that at the nanometer scale, Si exhibits different elastic and plastic properties compared to its bulk forms. Using an AFM tip, large fracture strain was achieved by bending single Si NWs.17
19 Using a nanomanipulator as the actuator and an AFM cantilever as the load sensor inside an SEM instrument, a large tensile strain of 12% was observed,20 and later, an elastic bending strain of 24% was claimed for Si NWs.21 However, no structural evolution details were reported to these bending processes. Some of these bending processes17
22 should involve highly active plasticdislocation events because the fracture strength in those studies should not be higher than the theoretical one.23,24 Direct TEM experimental results have shown that nanoscale Si possess r 2011 American Chemical Society

room-temperature plasticity,25
27 but several important issues remain unsolved for large straining in Si NWs. It is unknown if the nucleated dislocation is movable under bent strain, whether these dislocations interact with each other, and what is the consequence of these interactions. One of the key issues regarding Si NWs’ lowtemperature plasticity is that “does the crystalline Si transform to be amorphous by dislocation reactions?” We tend to answer these questions in the following by the investigations on bending Si NWs using the in situ atomic-scale experimental method.28,29 High-quality Si NWs 20
50 nm in diameter and several tens of micrometers in length were synthesized with thermal evaporation technique. The synthesized NWs were defect free before the elasticity limit was reached (Supporting Information Figure S1). The Si NWs were scattered on the prebroken colloidal film of the TEM grid. The scattered NWs could then be bridged on the cracked colloidal thin films. The colloidal thin films shrank under electron beam irradiation (see Supporting Information, Movie S1) and drove the nanowires to be bent or be conducted a tensile force. The in situ bending experiments of individual NWs can then be carried out with HRTEM, and the experimental processes can be recorded at the atomic scale. Figure 1 provides serial TEM images showing the continuous bending process of a single Si NW. According to the traditional Received: March 4, 2011 Revised: April 17, 2011 Published: May 05, 2011 2382

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Figure 1. A series of low-magnification TEM images that show the bending process of a single Si NW. (a
h) The strain sustained on the Si NW increased from 4.9 to 14.3%. No cracks were observed during the bending process.

formula εbent = r/(r þ R)%,30(where R is radius of the bending curvature of the NW and r is the radius of the Si NW), the maximum bent strain in the Si NW can be calculated and it increased from about 4.9% (Figure 1a) to be approximately 14.3% (Figure 1h). For a bent NW, the upper surface sustains the maximum tensile strain, while the bottom surface sustains the maximum compressive strain. The strain noted in the figures and the text indicates the maximum strain in the NW. All of the bending processes were conducted by showering the electron beam on the colloidal thin film to avoid overirradiation of the Si NWs. The colloidal thin film shrank and bent the Si NWs continuously at a strain rate about 5.4  10
4/s. The electron beam was moved onto the Si NWs only to capture the images of the bent Si NWs. The electron beam dose was kept at the level of 5  1019 e cm
2 s
1, which led to the local temperature being just a few degrees above room temperature27 (see Supporting Information). The electron irradiation energy for displacement of Si atoms is about 440 keV (see Supporting Information). Thus, both the knock-on effect and the heating from the electron beam irradiation can be neglected. Figure 2 shows the in situ observation of the full dislocation nucleation, movement and escaping process. Figure 2a is a lowmagnification TEM image that was captured at a bent strain of 4.9%. Figure 2b,c is the HRTEM images of the same region taken 30 s apart. Figure 2d,e is the enlarged HRTEM images corresponding to the framed regions in Figure 2b,c, respectively. Dislocation “1” nucleated at the bottom surface of the bent NW and is noted in Figure 2d. The local Burgers vector circuit indicates that the dislocation is a 60° full dislocation. This nucleated dislocation escaped during the deformation process, as shown in Figure 2e, and a second 60° dislocation “2” was observed in the NW (the same region was dislocation free in Figure 2d). In the same manner, dislocation “2” also escaped as the strain increased and a third dislocation “3” was nucleated afterward (Supporting Information Figure S2). This indicates that both the dislocation nucleation and movement indeed happened in the elastic
plastic transition when bending Si NWs to be large strains. Partial dislocations were captured in bent Si NWs when the maximum bent strain surpassed 6.9%. Those partial dislocation emissions from up-surface (under tensile strain) of the bent NW

Figure 2. In situ observation of the quick nucleation and motion of the dislocations during the bending process. (a) A low-magnification TEM image that was captured with a strain of 4.9%. (b,c) HRTEM images of the same region that were taken 30 s apart. (d,e) The enlarged HRTEM images corresponding to the framed regions in (b,c), respectively.

Figure 3. In situ observation of partial dislocation nucleation in the Si NWs when the bent strain surpasses 6.9%. (a) The Si nanowire with 4.9% bent strain and (b) the HREM image of the same region when the bent strain has been increased to approximately 6.9%.

left debris of the energy-cost stacking faults. Figures 3a (no partials) and 3b (emission of the partial from the same area) provide two typical atomic-scale images showing the partial dislocation emission process at the atomic scale directly. Besides the partial dislocation emission, we also observed that the length of the stacking faults changed from 13 atomic layers to 18 atomic layers of the (111) plane as the strain increased (Supporting Information Figure S3). It is interesting to observe partial dislocation nucleation at room temperature. Partial dislocations belong to the glide set31,32 with higher activation energy compared to the shuffle set.33 The low temperature activities of partial dislocation in bent Si NWs can be explained as the following: (a) The low dimensionality and defect-free structure of Si nanowires make it possible to sustain high strains and high stress. The high stress can supply enough energy to activate the dislocations belonging to the glide sets. (b) Full dislocations glided first in the shuffle planes and formed steps on the surface of the nanowire. The partials dislocations can then nucleate on the steppedsurfaces even at 0 K.34 2383

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Figure 4. The obvious movement of dislocations and the in situ observation of a Lomer lock dislocation formed by a dislocation reaction at the atomic scale. (a) A dislocation with a Burgers vector of a/2[101] (b) after the former dislocation escaped, another dislocation (b = 1/2 [011]) was nucleated; (c) the Lomer lock dislocation formed by a dislocation reaction; (d) the enlarged HRTEM image showing the Lomer lock structure.

Full dislocation activities were prevalent in the bent Si NWs under large strains. These full dislocations were observed to be mobile and, therefore, have chances to interact. Figure 4 shows the obvious movement of dislocations and the in situ observation of Lomer lock dislocation formed by a dislocation reaction at the atomic scale. In Figure 4a, a dislocation with a Burgers vector of 1/2 [101] was nucleated at the bottom surface of the bent Si NW and then escaped. Figure 4b was recorded from the same area of Figure 4a and shows another dislocation (b = 1/2 [011]) that was nucleated and then formed a new structural dislocation by reaction (as shown in Figure 4c). Local Burgers circuits were drawn to identify the Burgers vector b. The enlarged HRTEM image is given in Figure 4d. Extra planes are seen for both the (111) and the (111) planes. This configuration represents a Lomer dislocation35 exhibiting the Burgers vector of a/2[110]. The Lomer dislocation was formed by the interaction of two full dislocations with Burgers vectors a/2[101] and a/2[011], respectively, moving under applied strain on two intersecting slip planes, (111) and (111). The dislocation reaction can be written as 1 1 1 ½101 þ ½011 f ½110 2 2 2

ð1Þ

The two reacting dislocations mutually knit with each other to reach a low energy configuration.36 The b2 criterion for dislocation energy per unit length indicates a considerable energy reduction for the reaction. This is the first in situ demonstration of the Lomer lock dislocation formation in semiconductor NWs. With continuously increasing the bent strain, these locks could not be unzipped, and the locked lattices became disordered, leading to amorphization. Figure 5a
d was taken from a same area and shows the in situ observation of Lomer dislocation formation and the subsequent crystalline to amorphous transition process at the atomic scale. Through Figure 5a
d, the corresponding maximum bent strains were 10.3, 12.8, 13.3, and 14.3%, respectively. Figure 5a,b indicates the dislocation nucleation and motion. These activated dislocations were driven to meet under high strain/stress and then formed a Lomer

Figure 5. (a
d) The enlarged HRTEM images taken from the same region. The corresponding strains are 10.3, 12.8, 13.3, 13.8, and 14.3%, respectively. The c-a transition process was directly observed at the atomic scale.

dislocation by reaction in Figure 5c. The Lomer dislocation junction was expected to be sessile because it can glide in neither of the slip planes containing the reactant dislocations. High stresses are needed to break the Lomer lock, and its destruction process has only been observed in metals.36,37 This sessile dislocation will make the local stress to be very high as the strain increases. As shown in Figure 5d, unlike the destruction process of Lomer dislocation locks in metals,36,37 the lattices of the Lomer dislocation lock became highly distorted and finally completely disordered and transformed to be the amorphous state as the strain increased. This is the first direct atomic-scale observation of the crystal
amorphous (c-a) transition related to the Lomer dislocation. The deformation mechanisms presented in this paper can be extended to explain the large deformation ability of those semiconductor nanostructures, such as the high-strain bending ability of B4C NWs,38 the high bending strength achieved in Si nanobeams39 and the c-a transition observed in Ge NWs.40 In summary, more than 10 Si NWs were bent and in situ investigated at the atomic scale for exploring their mechanism of plastic deformation. The ultralarge bent straining of Si nanowires involves dislocation nucleation, movement, and interaction events. The ultrahigh straining process and the activated full dislocation movement led to the formation of Lomer locks. The strain-induced c-a transition was through the continuous increased straining on sessile Lomer dislocations. These results help to clarify the ultralarge straining ability of Si on the nanometer scale and to design nanowire-based flexible micro- and nanoelectronics.

’ ASSOCIATED CONTENT

bS

Supporting Information. Movie, additional information, and additional figures. This material is available free of charge via the Internet at http://pubs.acs.org.

’ AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected].

’ ACKNOWLEDGMENT This work was supported by the National Outstanding Young Investigator Grant of China (10825419), the Key Project of 2384

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Nano Letters C-NSF (50831001), the National 973 Program of China (2009CB623700), Beijing high-level talents (PHR20100503), Beijing PXM2011_014204_09_000053, the NSF(11004004) and Beijing Municipal Natural Science Foundation (1112004). The authors thank Dr. Xiaona Zhang and Mrs. Yan Qing for providing Si NWs.

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