NANO LETTERS
Microstructure-Hardened Silver Nanowires
2006 Vol. 6, No. 3 468-472
Bin Wu, Andreas Heidelberg, and John J. Boland* School of Chemistry and the Centre for Research on AdaptiVe Nanostructures and NanodeVices (CRANN), Trinity College Dublin, Dublin 2, Ireland
John E. Sader Department of Mathematics and Statistics, The UniVersity of Melbourne, Australia
XiaoMing Sun and YaDong Li Chemistry Department, Tsinghua UniVersity, Beijing, 100871, P R China Received December 8, 2005; Revised Manuscript Received January 20, 2006
ABSTRACT To exploit the novel size-dependent mechanical properties of nanowires, it is necessary for one to develop strategies to control the strength and toughness of these materials. Here, we report on the mechanical properties of silver nanowires with a unique fivefold twin structure using a lateral force atomic force microscopy (AFM) method in which wires are held in a double-clamped beam configuration. Force-displacement curves exhibit super elastic behavior followed by unexpected brittle failure without significant plastic deformation. Thermal annealing resulted in a gradual transition to weaker, more ductile materials associated with the elimination of the twinned boundary structure. These results point to the critical roles of microstructure and confinement in engineering the mechanical properties of nanoscale materials.
Metal nanowires have stimulated great interest as potential building blocks in nanoelectronic and nanoelectromechanical devices because of their high conductivities and high strengths. Although a variety of metal nanowires such as gold, copper, and silver have been synthesized successfully via wet chemistry or template-directed approaches, much less is known about the chemical or physical properties of freestanding nanomaterials at a single-object level. One of the main issues involved in the assembly of functional structures using these nano building blocks, however, is the relationship between material properties and structural elements such as size, geometry, and microstructure. Many earlier studies have focused on size-dependent properties.1-7 For example, micrometer-sized whiskers and nanowires have ultrahigh strength compared with that of their bulk counterparts.1-3 The exceptional strength is believed to originate from a decreased defect density resulting from reduced wire size. However, it is well established that microstructures such as crystallinity, defects, and grain boundaries play important roles in bulk material properties. Metals in bulk form can be engineered to have high strengths by cold-work hardening, grain-boundary hardening, and precipitation hardening, techniques that rely on restricting and hindering dislocation motion.8 However, on the nanometer scale strengthening methods that rely on the incorporation of impurities may be * To whom correspondence should be addressed. E-mail:
[email protected]. 10.1021/nl052427f CCC: $33.50 Published on Web 02/21/2006
© 2006 American Chemical Society
ineffective because of facile surface segregation and expulsion. This leaves microstructure as the best candidate route for controlling material strength. To date, single crystalline, polycrystalline, and twin structured nanowires9-12 have been synthesized, but the correlation between the microstructures and mechanical properties is poorly understood, principally because of the difficulties in performing standard tensile or bending tests on individual wires. Here in this work we demonstrate that microstructure control is a particularly effective means for controlling mechanical properties in nanowire systems. Using a fivefold twinned Ag nanowire system, we show that the precisely controlled grain orientation and grain-boundary organization within these wires is responsible for their anomalous strength and brittle failure. The principle slip directions in these grains intersect with the twinning boundaries that extend along the entire wire length to produce a uniformly hardened structure. These results demonstrate that, in the case of free-standing nanoscale materials, grain-boundary hardening is extraordinarily effective because of the ability to completely control the orientation and boundaries of the limited numbers of grains in these materials. Nanowire mechanical properties were measured using a AFM lateral bending technique developed recently,2 which in contrast to earlier methods,3,13-23 allows the full spectrum of mechanical properties to be measured, ranging from elasticity to plasticity and failure. Figure 1a shows a low-
Figure 1. (a) TEM image of pentagonal silver nanowires. (b) Cross-sectional high-resolution TEM image of pentagonal silver nanowires. (c) Schematic of slip systems for faced-centered cubic fivefold twinned Ag nanowires, showing that one of the slip planes (111) together with three possible slip directions 〈110〉 inevitably intersect the grain boundaries. Shaded planes indicate fivefold twinned boundaries.
magnification transmission electron microscopy (TEM) image of pentagonal silver nanowires prepared at a lower temperature (140 °C). Details of the preparation of these pentagonal silver nanowires can be found elsewhere.9 These wires have typical lengths of many micrometers and diameters ranging from 16 to 35 nm. The cross-sectional TEM image in Figure 1b shows the remarkable fivefold twinned grain-boundary structure that exists along the entire wire length. Note that there is a ∼2 nm thick carbon coating on the surface as determined from high-resolution TEM observation9 (data is not shown here). To perform three-point bending tests on 16-35 nm diameter pentagonal silver nanowires, we fabricated welldefined trench patterns on a substrate that was coated with a 10-20 nm TiN film. Typically, trenches having widths of 650 nm and depths of 475 nm were used in these experiments. This selection provides a reasonable ratio between the pinned nanowire length and its diameter while eliminating problems associated with wire droop. The pentagonal silver nanowires were dispersed in ethanol and then deposited on the prepatterned substrates after solvent evaporation. Silver nanowires found to bridge well-defined trenches were located by scanning electron microscopy (SEM) in a dual-beam (electron/focused ion beam) system and subsequently doubleclamped at the trench edges by electron-beam-induced deposition of Pt lines. Details of the experimental conditions and the pinning procedure can be found elsewhere.2 A wellcalibrated rectangular cantilever (Budgetsensors) with an average normal force constant of 1-3 N/m (75 kHz) was used in the bending experiments. AFM lateral manipulations were carried out using a Digital Instruments Nanoman System with closed-loop x-y-z scanner. By positioning the Nano Lett., Vol. 6, No. 3, 2006
AFM tip 450 nm into the trench, that is, 20-30 nm above the trench floor, frictional forces between the trench floor and tip were eliminated completely. Lateral bending measurements were then performed either as a single-shot experiment (in which the tip engaged the wire, elastically and then plastically deformed it, and finally broke the wire in a single manipulation) or in a series of loading-unloading cycles (in which the wire was increasingly loaded and unloaded in a series of manipulations so that progressive elastic and then plastic deformation occurred, followed by wire failure). The normal and lateral force signals were recorded using a Labview-based program. However, here we focus on the lateral force because in this geometry the normal force on the cantilever is less than 5% of the total force.2 Tip velocities were 20 nm/s throughout and all details relating to the manipulation and tip calibration procedures can be found elsewhere.2 Figure 2 shows a typical set of experimental data including AFM images before bending and after failure (Figure 2a and b) together with the F-d curves recorded during a series of loading and unloading cycles (Figure 2c). The two curves labeled 1 and 2 in Figure 2c (which are displaced relative to each other for viewing purposes) are essentially nonlinear but symmetric about the vertical dashed lines that identify the starting point for unloading, that is, the cantilever’s turning point. This symmetry reflects the full elastic recovery of the wire and is evident also from AFM images (not shown, but identical to Figure 2a) recorded both before and afterward, which reveal no permanent deformation of the wire. However, after reloading again the wire was subject to a large manipulation that resulted in F-d curve 3 in Figure 2c. In this curve, a sharp force-drop was observed and is associated with wire failure, which is confirmed immediately by the subsequent AFM image (Figure 2b). The F-d curves are nonlinear and reproducible; both curves 1 and 2 can be shifted to completely overlap curve 3 in Figure 2. These F-d curves were analyzed in terms of a generalized model that includes contributions from wire bending and tensile stretching.24 This approach has the advantage that it provides an accurate description of the mechanical properties over the entire range of elastic deformation. The exact analytical solution is expressed as24 192EI f(R)∆zcenter L3 R where f(R) ) xR 192 tanh 4 48 xR Fcenter )
(1)
( )
R is related to by (the detailed description of this approach can be found elsewhere24) ) ∆z2center R)
(AI)
6(140 + ) 350 + 3
(2) (3) 469
Figure 3. Plot of Young’s modulus versus the nanowire radius in the range of 10-15 nm for pentagonal silver nanowires. The Young’s modulus remains essentially the same before (circle) and after (star) thermal annealing experiments. The dashed line shows the average value of the Young’s modulus for bulk silver.
Figure 2. (a and b) Tapping-mode AFM images of a 23.6 nm diameter pentagonal silver nanowire before bending and after brittle failure. All scale bars are 250 nm. (c) F-d curves recorded during the consecutive manipulation by AFM tip-induced lateral bending of a 23.6 nm pentagonal silver nanowire. Curves 1 and 2 show that the wire was elastically loaded and unloaded. The unloading points are identified as vertical dashed lines. Curve 3 is a singleshot experiment and shows nonlinear elastic behavior of silver nanowire, followed by limited plastic deformation and then brittle failure. Note that F-d curves are shifted for clarity. Inset: schematic of bending test showing that the bending angle defined as the angle between the deformed wire and its original direction. (d) Fit of F-d curve 3 to the generalized formula, which yields a Young’s modulus of 90 GPa.
where Fcenter is the measured lateral force, E is Young’s modulus, I is the moment of inertia, and ∆zcenter is the displacement of the wire of suspended length L and crosssectional area A. Because the wires have a circular cross section as seen from the TEM image in Figure 1b, the moment of inertia is I ) πr4/4. AFM was used to determine the diameter and the suspended length of pinned wires. The diameter was measured at several points along the wires’ length and average values based on a typical scatter of (1 nm were used for fitting the curves. A complete analysis 470
also requires a detailed calibration of the AFM tip dimensions and a determination of cantilever lateral spring constant, the procedures for which are described elsewhere.2 The F-d curves in Figure 2 were fitted using eqs 1-3 to determine the mechanical properties of the nanowire: Young’s modulus, elastic deformation, yield point, and failure. The data analysis was performed using a MATHEMATICA 5.2based algorithm that inputted all measured parameters (wire diameter, the suspended length of wires, lateral force constant of the cantilever, etc.) and in which the Young’s modulus, E, was the only adjustable parameter. Analysis involved initially fitting data at small displacement and then extending the range to include increasingly larger displacements. The value of the modulus obtained by this method remained constant until the yield point associated with the onset of plastic deformation was reached and after which the apparent value of the modulus dropped and the fit became increasingly poorer. This approach is possible only because the generalized formula provides a complete description of the elastic properties over the entire range of displacements. Curve 3 in Figure 2c is replotted in Figure 2d together with the corresponding fit to the generalized theory. The excellent fit demonstrates that these pentagonal silver nanowires exhibit very good elasticity up to displacements of ∼80 nm; that is, these wires can be bent elastically up to an angle of ∼13.2° from the wire original direction (see the inset in Figure 2c). The axial strain induced by bending is related to the bending angle by ) (1/cos θ) - 1. After the yield point, Figure 2d indicates that the wire undergoes only 30 nm of plastic deformation prior to failure. Significantly, we measure ∼30 nm permanent displacement of wire in the postfailure AFM image Figure 2b, which underscores the validity of this analysis in the determination of both the modulus and yield point. We have measured the mechanical properties of 24 single wires, and Figure 3 shows values of the Young’s modulus determined for a range of pentagonal silver nanowires with diameters from 22 to 35 nm. The data scatter falls within that typically found using this technique, and details of the Nano Lett., Vol. 6, No. 3, 2006
error propagation can be found elsewhere.2 The average value of the modulus is 102 ( 23 GPa and is higher than that of bulk silver (83 GPa). Oxygen plasma experiments that removed the 2 nm carbon coating indicated that the coating is not responsible for the high modulus values of the pentagonal silver nanowires (see the Supporting Information). At present, the physical origin of the increased Young’s modulus is unknown, but similar observations have been reported for other nanoscale systems.5-7 The observation in Figure 2d that the extent of plastic deformation at failure is less than 40% of the elastic deformation is exceptional for a pure metal. In the case of gold nanowires, the total plastic deformation during bending is up to 450% of the elastic deformation, an order of magnitude greater than that observed for the present silver wires.2 For bulk metals, strengthening typically involves the incorporation of impurities or the modification of the microstructure. Although the former method is effective for engineering bulk materials, it has limited application on the nanoscale because impurities can be expelled easily from the material. The effectiveness of the microstructure modification, however, is limited by physical size in bulk specimens because there will always be grains with favorable orientations for plastic deformation. In contrast, the finite physical dimensions and limited numbers of highly orientated nanoscale grains within the present Ag nanowires make them particularly amenable to strengthening by microstructure modification. We hypothesize here that the novel fivefold twin microstructure of the nanowires is responsible for the unique mechanical properties seen in Figure 2. Wire growth occurs along the direction of the face-centered cubic (FCC) crystal structure, which is characterized by a slip system of four (111) slip planes, along the three directions (see Figure 1c). Consequently, the fivefold grain boundaries in these wires necessarily intersect with all of the possible slip systems and so the motion of dislocations associated with the initiation of plastic deformation along any slip direction is restricted by the twinning boundaries that extend into the center of the wire. In this manner, the fivefold twinned silver nanowires are effectively grainboundary hardened materials, which sacrifices ductility for strength. Because the twin boundaries exist along the entire length of wire, the whole wire is uniformly hardened and there are no defects that limit the strength of the wire. To test this hypothesis, we carried out nanowire annealing experiments at 240-250 °C in nitrogen gas and then let them cool slowly to room temperature before removing and mechanically testing the samples. These annealing conditions are sufficient to induce recrystallization even for bulk silver. As a result, atoms at the grain boundaries diffuse rapidly to reduce the interfacial energy, resulting in the gradual elimination of the unique pentagonal grain-boundary structure. Figure 4 shows single-shot mechanical measurements for pentagonal silver nanowires that were annealed for 17 and 48 h, respectively. These data, which are representative of a study involving 15 nanowires, illustrate several important points, which are summarized in Table 1. First, the Young’s modulus remains essentially unchanged before and after Nano Lett., Vol. 6, No. 3, 2006
Figure 4. (a) A typical single-shot F-d curve fit for a 23.2 nm diameter pentagonal silver nanowire following 17 h thermal annealing treatment in N2. The red curve is the best fit of the elastic region of F-d data showing three regimes: the elastic bending, grain-boundary hardened plastic bending and failure. The point where the red curve departs from the data is identified as the yield point. The curve fit gives a modulus of 121 GPa. (b) A typical single-shot F-d curve fit for a 17 nm diameter silver nanowire after 48 h thermal annealing. The modulus is 99 GPa in this case. (c) AFM images of a 16.5 nm diameter silver nanowire after 48 h thermal annealing in N2 during consecutive lateral bending tests. All scale bars are 250 nm.
annealing (see Figure 3). This can also be clearly seen from the solid curves in Figure 4a and b, which represents a best fit to the generalized formula over the elastic region of the data (see above). Second, the F-d curve fit for the 17 h annealed sample in Figure 4a shows that the wire is 471
Table 1. Summary of the Bending Angles in Different Samples
original Ag nanowire 17 h annealed Ag nanowire 48 h annealed Ag nanowire 200-nm Au nanowirea
maximum elastic bending angles
bending angle at failure
ratio between plastic and elastic displacements
13.2°
17.9°
35%
8.2°
21.9°
169%
6.9°
27.1°
304%
6.8°
34.0°
432%
a Typical bending angles for a 200-nm Au nanowire were shown for comparison.2
elastically deformed up to a displacement of 50 nm, which corresponds to a maximum elastic bending angle of ∼8.2°, followed by 90 nm of plastic deformation. The extended 48 h anneal resulted in an elastic displacement of 40 nm, a reduced elastic bending angle of 6.9°, and 130 nm of plastic deformation (Figure 4b). When compared with the mechanical properties of original unannealed wire, the annealed wires have much reduced elasticity with lower yield strengths (see Table 1). For these annealed wires, yielding at small deformation is induced by mechanical pure bending2 and the yield strength can be estimated as σy ) FyL/2πr3, resulting in values (∼7.3 GPa from Figure 4c) that are substantially larger than that for bulk silver (55 MPa). This analysis is not valid for the original wires because yielding involves both stretching and bending, but is indicative of even stronger materials. Finally, the data in Figure 4 also demonstrate that the transition from a brittle to ductile metal is a gradual one. Note that the yield point (the deviation point between the fitted curves and the data) in Figures 2d, 4a, and 4b progressively moves to smaller displacements and the smaller the elastic recovery following plastic deformation. This is consistent with reduced hardening as the grain boundaries disappear during annealing and is also evident from the variation in the bending angle at failure, which range from 17.9° for the original wire, to 21.9° and 27.1°, for the 17 and 48 h annealed wires, respectively (see Table 1). This dramatic increase in ductility can also be seen from the AFM images in Figure 4c recorded following consecutive manipulation of the 48 h annealed wire. These results demonstrate that the mechanical properties of nanowires can be uniquely tailored by controlling their microstructure. Although microstructure is a recognized means to engineering the strength of materials, it is exceptionally effective on the nanoscale level because it is possible to assemble materials with oriented, interlocking grains that are both grain-boundary hardened and oriented so as to eliminate all favorable slip orientations.
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Acknowledgment. This work was supported by Science Foundation Ireland under grant 00/PI.1/C077A.2. Note Added after ASAP Publication. A coauthor was not included in the version published ASAP February 21, 2006; the correct version was published March 3, 2006. Supporting Information Available: The possible effect of carbon coating on the elastic properties of pentagonal silver wires was performed by oxygen plasma etching experiments. AFM images, height analysis, and corresponding F-d curve fitting are shown. This material is available free of charge via the Internet at http://pubs.acs.org. References (1) Levitt, A. P. Whisker Technology; Levitt, A. P., Ed.; WileyInterscience: New York, 1970. (2) Wu, B.; Heidelberg, A.; Boland, J. J. Nat. Mater. 2005, 4, 525. (3) Wong, E. W.; Sheehan, P. E.; Lieber, C. M. Science 1997, 277, 1971. (4) Uchic, M. D.; Dimiduk, D. M.; Florando, J. N.; Nix, W. D. Science 2004, 305, 986-989. (5) Cuenot, S.; Champagne, S. D.; Nysten, B. Phys. ReV. Lett. 2000, 85, 1690. (6) Cuenot, S.; Fretigny, C.; Demoustier-Champagne, S.; Nysten, B. Phys. ReV. B. 2004, 69, 165410. (7) Liang, H. Y.; Upmanyu, M. Phys. ReV. B. 2005, 71, 241403(R). (8) Callister, W. D., Jr. Materials Science and Engineering: An Introduction, third edition; John Wiley & Sons: New York, 1994. (9) Sun, X. M.; Li, Y. D. AdV. Mater. 2005, 17, 2626. (10) Chen, H. Y.; Cao, Y.; Zhang, H. R.; Liu, L. B.; Yu, H. C.; Tian, H. F.; Xie, S. S.; Li, J. Q. J. Phys. Chem. B. 2004, 108, 12038. (11) Jiang, P.; Li, S. Y.; Xie, S. S.; Gao, Y.; Song, L. Chem.sEur. J. 2004, 10, 4817. (12) Wiley, B.; Sun, Y. G.; Mayers, B.; Xia, Y. N. Chem.sEur. J. 2005, 11, 454. (13) Yu, M. F.; Lourie, O.; Dyer, M. J.; Moloni, K.; Kelly, T. F.; Ruoff, R. S. Science 2000, 287, 637-640. (14) Poncharal, P.; Wang, Z. L.; Ugarte, D.; De Heer, W. A. Science 1999, 283, 1513-1516. (15) Gao, R. P.; Wang, Z. L.; Bai, Z. G.; De Heer, W. A.; Dai, L. M.; Gao, M. Phys. ReV. Lett. 2000, 85, 622-625. (16) Salvetat, J. P.; Briggs, G. A. D.; Bonard, J. M.; Bacsa, R. R.; Eulik, A. J.; Stockli, T.; Burnham, N. A.; Forro, L. Phys. ReV. Lett. 1999, 82, 944. (17) Salvetat, J. P.; Kulik, A. J.; Bonard, J. M.; Briggs, G. A. D.; Stockli, T.; Metenier, K.; Bonnamy, S.; Beguin, F.; Burnham, N. A.; Forro, L. AdV. Mater. 1999, 11, 161. (18) Kis, A.; Mihailovic, D.; Remskar, M.; Mrzel, A.; Jesih, A.; Piwonski, I.; Kulik, A. J.; Benoit, W.; Forro, L. AdV. Mater. 2003, 15, 733. (19) Walters, D. A.; Ericson, L. M.; Casavant, M. J.; Liu, J.; Colbert, D. T.; Smith, K. A.; Smalley, R. E. Appl. Phys. Lett. 1999, 74, 3803. (20) Tombler, T. W.; Zhou, C. W.; Alexseyev, L.; Kong, J.; Dai, H. J.; Liu, L.; Jayanthi, C. S.; Tang, M. J.; Wu, S. Y. Nature 2000, 405, 769. (21) Minot, E. D.; Yaish, Y.; Sazonova, V.; Park, J. Y.; Brink, M.; McEuen, P. L. Phys. ReV. Lett. 2003, 90, 1564011. (22) Falvo, M. R.; Clary, G. J.; Taylor, R. M.; Chi, V.; Brooks, F. P.; Washburn, S.; Superfine, R. Nature 1997, 389, 582. (23) Li, X. D.; Gao, H. S.; Murphy, C. J.; Caswell, K. K. Nano Lett. 2003, 3, 1495. (24) Heidelberg, A.; Ngo, L. T.; Wu, B.; Phillips, M. A.; Sharma, S.; Kamins, T. I.; Sader, J. E.; Boland, J. J., submitted for publication.
NL052427F
Nano Lett., Vol. 6, No. 3, 2006
