pubs.acs.org/NanoLett
Screw Dislocation-Driven Epitaxial Solution Growth of ZnO Nanowires Seeded by Dislocations in GaN Substrates Stephen A. Morin and Song Jin* Department of Chemistry, University of Wisconsin-Madison, 1101 University Avenue, Madison, Wisconsin 53706 ABSTRACT In the current examples of dislocation-driven nanowire growth, the screw dislocations that propagate one-dimensional growth originate from spontaneously formed highly defective “seed” crystals. Here we intentionally utilize screw dislocations from defect-rich gallium nitride (GaN) thin films to propagate dislocation-driven growth, demonstrating epitaxial growth of zinc oxide (ZnO) nanowires directly from aqueous solution. Atomic force microscopy confirms screw dislocations are present on the native GaN surface and ZnO nanowires grow directly from dislocation etch pits of heavily etched GaN surfaces. Furthermore, transmission electron microscopy confirms the existence of axial dislocations. Eshelby twist in the resulting ZnO nanowires was confirmed using bright-/ dark-field imaging and twist contour analysis. These results further confirm the connection between dislocation source and nanowire growth. This may eventually lead to defect engineering strategies for rationally designed catalyst-free dislocation-driven nanowire growth for specific applications. KEYWORDS Zinc oxide, gallium nitride, dislocations, nanowires, epitaxy, and Eshelby twist
N
anowire (NW) growth can be driven by dislocations, where an axial screw dislocation spiral at the tip of the NW propagates the anisotropic crystal growth.1-4 Under low supersaturations, layer-by-layer (LBL) growth is prohibited due to the energy penalty for creating a new surface layer and growth occurs only at the dislocation spiral causing highly anisotropic one-dimensional (1D) growth.3 So far this NW growth mechanism is less reported than the wellknown vapor-liquid-solid (VLS) growth mechanism.5,6 Because dislocation-driven NW growth does not require catalysts and thus consideration of the catalyst/NW material eutectics, this mechanism could eventually become more versatile. In principle, dislocation-driven NW growth could be applied to any material system or growth phase simply by rationally designing the supersaturation conditions of the system.3,7 This could open the door to large-scale low cost NW synthesis of a variety of different materials for various applications. Originally discovered and elucidated in the vapor phase growth of lead sulfide NW pine trees,1 we recently expanded dislocation-driven growth to the aqueous solution synthesis of ZnO NWs, nanorods (NRs), and nanotubes (NTs).3 ZnO NWs have found applications in solar8,9 and piezoelectric energy conversion10 as well as in photonics11 and have been widely studied and reported in literature.12-15 Until recently, there was no consensus on the aqueous growth mechanism of ZnO nanomaterials, but now it is clear that the growth of these materials is dislocation-driven.3 In these cases, the
dislocation sources, the “seeds”, which propagate dislocation-driven NW growth, come from highly defective nanoparticles synthesized under high supersaturations. However, this method of dislocation seeding is difficult to control and most useful for generating randomly oriented NWs. Since many applications demand control over NW orientation and location, more control over the dislocation “seed” is critical to advancing the utility of dislocation-driven NW growth. As with VLS grown NWs, which can be organized spatially by controlling the location of the catalysts,11,16 precisely engineering dislocation “seeds” on a surface can allow control over the arrangement of dislocation-driven NWs. Here we demonstrate the first step of such intentionally initiated dislocation-driven growth by utilizing dislocations that are contained within GaN thin films and exposed on the surface to propagate the growth of epitaxial ZnO NWs directly from aqueous solutions. There have been many examples of catalystfree epitaxial growth of ZnO NWs on GaN using metal-organic chemical vapor deposition (MOCVD)17,18 or from aqueous solutions19,20 but no clear explanation for 1D growth. In fact, the favorable lattice match between the basal planes of GaN and ZnO (both wurtzite structure with aZnO ) 3.249 Å and aGaN ) 3.186 Å)21 would tend to promote two-dimensional epitaxial film growth over 1D NW growth. We realized that GaN substrates were so effective at enabling epitaxial 1D growth of ZnO NWs probably because GaN thin films grown on sapphire substrates contain high densities of dislocations22 that can actually act as dislocation sources to propagate dislocationdriven 1D growth. We show that the resulting ZnO NWs contain axial dislocations and that their growth originates from the screw dislocations on the GaN surface with comparable densities. The Eshelby twist of these NWs was measured by
* To whom correspondence should be addressed. E-mail:
[email protected]. Received for review: 04/30/2010 Published on Web: 08/18/2010 © 2010 American Chemical Society
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small etch pits at the core of many dislocations with pure or partial screw character as revealed by AFM (Figure 1e,f). The dislocation density of these substrates is approximately ∼3 screw dislocations per 10 µm2 of surface, or ∼0.3/µm2, which is typical for MOCVD grown GaN.22 Optimizing etching conditions to be just enough to expose the pristine GaN surface and the continuous dislocations, usually ∼5 min at 165 °C, followed by flow reactor growth of ZnO using the optimized aqueous precursor concentration, which was 55 µM Zn(NO3)2/hexamethylenetetramine (see experimental details in Supporting Information), yields vertical epitaxial ZnO NWs over a large area (Figure 2a). The NW density of ∼0.5/µm2 agrees reasonably well with the estimated density of screw dislocations. Furthermore, if we employ GaN substrates that have been more heavily etched to open large etch pits, visible under scanning electron microscopy (SEM, Figure 2b) and AFM (Figure 2c), it is apparent that virtually every ZnO NW grown originates from a dislocation etch pit. Interestingly, small spiral protrusions are often visible at the bases of some removed NWs that were examined using TEM (Figure 2d). We speculate this is where the dislocation etch pits were “filled in” during the initial stages of NW growth. It is interesting to note this picture matches our intuition of how these NWs would begin to grow (Figure 1b). The degree of epitaxy in the optimized NW sample (Figure 2a) is easily confirmed by the X-ray diffraction (XRD) rocking curve around the ZnO (002) reflection (Figure 2e) which has a full width at half max (fwhm) of 0.3°.26 TEM characterization of these ZnO NWs grown on GaN further shows that they are dislocation-driven. Select area electron diffraction (SAED) (Figure 3a inset) confirms that the NWs are single crystalline with the wurtzite structure, which is in agreement with previous findings. Using highresolution TEM (HRTEM, Figure 3b), we have determined the growth direction of these NWs is along the [001] direction (the c-axis), which is the same axis reported previously for dislocation-driven ZnO NWs3 and most often observed. TEM observation of axial dislocations in these NWs directly is not trivial because dislocations can move or be worked out during sample preparation. It is well-known that dislocations in small volumes are especially mobile and instable.27-30 By employing the most noninvasive and gentle dry transfer of NWs to lacey carbon TEM grids, we were able to observe axial dislocations in many NWs. The representative example shown in Figure 3c was imaged using zero beam bright-field TEM and shows a very distinct axial dislocation that is also visible in the corresponding dark-field TEM image (Figure 3c inset). We further note that dislocation contrast is not clearly observed in all objects examined due to the mobility of dislocation under the inevitable mechanical perturbation, in fact, sometimes dislocations can even move during the course of TEM observation. The observation of dislocations in these epitaxial ZnO NWs and on the GaN substrates and the presence of spiraling protrusions at their bases strongly support our hypothesis that these NWs grow via a disloca-
FIGURE 1. Schematic of dislocation-driven epitaxial nanowires grown from dislocated substrates. (a) GaN substrate terminated with (001) Ga surface on basal face of sapphire is cleaned and etched mildly to expose dislocations that then seed dislocation-driven growth of ZnO NWs directly from aqueous solutions. (b) Illustration of a dislocation-driven epitaxial NW that results from this process. (c) Diffraction contrast cross sectional TEM shows the high density of dislocations in thin film GaN grown on sapphire (reproduced from ref 22 with copyright permission from The Japan Society of Applied Physics). (d) Deflection mode AFM image showing top down view of dislocations on GaN. (e) Lowmagnification deflection AFM image post cleaning and etching. (f) Highmagnification height mode AFM image of the area marked in (e). An etch pit is marked by a black arrow.
indexing twist contours, further confirming the dislocationdriven NW growth mechanism. Our strategy is illustrated in Figure 1a. First, GaN substrates are degreased to remove surface contaminants then mildly etched to remove any oxide layer and expose the dislocations on the surface. When subjected to low-supersaturation ZnO growth solutions, these dislocations “seed” or propagate the dislocation-driven growth of ZnO NWs generating epitaxial arrays of dislocation-driven NWs directly from solution (Figure 1a,b). It is well-known that GaN thin films grown on sapphire have high dislocation densities apparent from previous cross-sectional transmission electron microscopy (TEM) studies (Figure 1c).22 They are defect rich because of the slight lattice mismatch between GaN and the sapphire substrates onto which the films are epitaxially grown, as previous reports have clearly shown.21,23 The resulting dislocation-driven epitaxial ZnO NWs (Figure 1a,b) are seeded by the spiral dislocations from the GaN substrates and grow via their propagation along the NW growth axis. In this study, GaN thin films terminated with the (001) Ga face were grown epitaxially on the c-plane of sapphire using well established MOCVD techniques.23,24 Contact mode atomic force microscopy (AFM) of the as grown GaN surface prior to etching confirms this high dislocation density (Figure 1d). Degreasing clears organic surface contamination and etching using a 10 wt % solution of potassium hydroxide in ethylene glycol removes surface oxide layers as well as the top layers of GaN.25 These treatments also ensure consistent surfaces for each experiment. It can also open up © 2010 American Chemical Society
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FIGURE 2. Dislocation-driven epitaxial ZnO nanowires on GaN. (a) Tilted (18°) SEM image of epitaxial ZnO NWs over a large area. The insets show progressively higher-magnification views of these NWs. (b) ZnO growth on an overetched GaN surface shows NWs grow directly from dislocation etch pits. (c) AFM height image showing the dislocation etch pit features on the substrate used in (b). (d) TEM images of the base of a ZnO NW show a spiraling protrusion that presumably “filled in” the spiral dislocation etch pit. (e) XRD rocking curve, centered around the ZnO (002) reflection, for sample in (a) verifies the high degree of epitaxy.
where b is the Burger’s vector magnitude and r is the NW radius, or by hollowing out of the dislocation core to create a tube, or by both mechanisms together when the tube wall is thin creating twisting tubes.3 The final morphology therefore depends on the energy balance between these two mechanisms, which is dependent on the properties of the material and the dislocation itself. For the ZnO grown here, only solid core NWs have been observed and not NTs, therefore the strain energy created by axial dislocations in these NWs must go into the Eshelby twist mechanism. Verifying this twist provides further proof for the dislocationdriven mechanism. Eshelby twist in 1D nanomaterials can be observed using several common microscopy or diffraction techniques. Direct SEM observation can reveal the twist in surface-faceted 1D crystals, however the crystals must be large enough to exhibit faceting and the twist sufficiently dramatic for easy observation. Because the Eshelby twist equation (eq 1) says larger diameter NWs have smaller twist, such conflicting requirements are hard to satisfy simultaneously. As a hypothetical estimate, a NW with a 100 nm diameter and a Burger’s vector magnitude of 1 nm should have an Eshelby twist of about 7.3°/µm. In the occasional fortunate cases where visible reference points other than facets are available, such as the epitaxial side branches on the NW pine trees of PbS1,2 and PbSe31 previously reported, SEM observation and tracking of small amounts of twist over long distance becomes much easier. Most directly, SAED patterns can be collected along the growth axis of a NW whose lattice is twisting using TEM and electron diffraction (ED).31 If two different known zone axes orthogonal to the NW growth direction can be observed at different locations of the same NW, the angle between these zone axes and the distance between the collection points can be used to calculate the twist. This may not be trivial since it is rare for a NW to be perfectly oriented on a TEM grid for this type of analysis and
FIGURE 3. Transmission electron microscopy characterization of ZnO NWs. (a) Low-resolution bright-field TEM image of a single NW confirmed to have the wurtzite structure by the inset SAED pattern, where the area of analysis is indicated by the red dotted circle. (b) HRTEM image of the same area confirming growth along the [001] direction (inset is the corresponding fast Fourier transform). (c) Zero beam brightfield TEM image showing an axial dislocation in another ZnO NW (inset is the corresponding dark field image for the (12¯0) g vector).
tion-driven mechanism and that the dislocations that initiate growth come from highly dislocated GaN substrates. Axial dislocations create stress and strain in the NWs that contain them. This strain may be alleviated by a torque around the NW, known as the Eshelby twist (R)27,28
α)
b πr2
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respectively (Figure 4b top and bottom diagrams, respectively), the difference between the two angles (∆θ) can be calculated geometrically using the relation
∆θ ) (θ2 - θ1) )
2
- g1 |
(2)
where λ is the electron wavelength (the inverse of k). Note the approximation θ ∼ λg/2 is made because of the large radius of the Ewald sphere. For a twisting NW, the upper and lower conditions in Figure 4b can be simultaneously met at different points along the NW and the resulting twist contours can be imaged using zero beam bright-field TEM and indexed directly using displaced aperture dark-field TEM. Unlike bend contours, each of these twist contours is associated with one set of planes and can be easily differentiated. The measured real space distance (L) between two indexed twist contours can be used to calculate the real space twist (R), effectively the Eshelby twist, of the NW using the following relation
FIGURE 4. Observation and determination of Eshelby twist using twist contours. (a) Schematic illustration of a NW showing g vectors orthogonal to the growth axis and the k vector of the electron beam. (b) Conditions where -g and +g vectors (top and bottom diagrams, respectively) meet the Laue condition. (c) Zero beam bright-field TEM image of a representative NW showing two indexed twist contours and the physical separation (L). Inset is the convergent beam electron diffraction pattern. (d,e) Displaced aperture dark-field TEM used to index the two labeled contours as (1¯22¯) and (12¯2). For this object, the g vectors were not perfectly orthogonal and the calculated twist is 9°/µm.
α)
( ∆θL ) ) ( 2Lλ )|g
2
- g1 |
(3)
We have implemented this procedure and estimated the twist in our ZnO NWs (Figure 4c-e) to be typically between 5 and 10°/µm, which is consistent with the Eshelby prediction (eq 1) for a reasonable Burgers vector magnitude of ∼1 nm. Specifically, for the NW shown in Figure 4c that has a measured Eshelby twist of 9°/µm and radius of 46 nm, the estimated b is 1.0 nm. While the schematic (Figure 4a,b) illustrates the simplest case of g vectors perfectly orthogonal to the NW growth axis, nonorthogonal g vectors may be used so long as the correct geometric correction is applied. In practice, using twist contours to quantify Eshelby twist is much more easily implemented on NW samples than the other techniques discussed above and in general more accessible than X-ray analysis methods not discussed here. Moreover, twist contour analysis can quantify twist of ∼1°/µm or less, which is a much higher resolution than the other techniques. This technique may make observation and confirmation of the dislocation-driven NW mechanism more routine. Here we have deliberately generated dislocation-driven epitaxial ZnO NWs from aqueous solutions using dislocations contained in GaN substrates to “seed” and propagate dislocation growth. We have observed the dislocations in both the GaN substrates and the ZnO NWs with NW growth densities similar to the dislocation density observed for such GaN substrates. Furthermore, the direct connection between NW growth and dislocation etch pits plus the presence of protrusions at the bases of many NWs supports the conclusion that dislocations in GaN are “seeding” and propagating the NW growth. Additionally, we have described a technique
often twist is not dramatic enough to allow the observation of two different zone axes across the length of a NW. Under the best case scenario using standard TEM instrumentation, a twist magnitude of ∼5°/µm is the smallest twist measurable using this technique.3 Since many cases will involve twist magnitudes below this limit, we introduce a more precise technique below for estimating NW twist based on twist contours observed in TEM. Originally described by Drum in the 1960s,32 the quantification of lattice twist using twist contours can be implemented as follows. First twist contours must be differentiated from the bend contours commonly observed in TEM of nanomaterials. Bend contours, as the name suggests, occur when a NW is bent about its axis. This distorts many atomic planes locally at the bend causing the diffraction of electrons from this region of the crystal to be different than everywhere else. Therefore, when a zero beam bright-field TEM image is taken of the bended object, because electrons are diffracted in such a way that virtually none of them remain in the zero beam the bend contour appears very dark compared to everywhere else. In contrast, when the real space crystal lattice twists along the growth axis causing the reciprocal space lattice to twist simultaneously, the various reciprocal g vectors that are orthogonal to the NW growth axis (Figure 4a) come on and off of the Laue condition at different angles, creating several twist contour bands (Figure 4b). If we generically choose -θ for -g and +θ for +g, which we will call θ1, g1 and θ2, g2, © 2010 American Chemical Society
( 2λ )|g
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for characterizing small Eshelby twist in 1D nanomaterials by indexing twist contours. More sensitive and high throughput, this technique could facilitate the identification of more 1D materials whose growth follows the dislocation-driven mechanism. In retrospect, we now believe it is possible that many previously reported catalyst-free epitaxial ZnO NW syntheses using either solution growth19,20 or MOCVD17,18,26 on dislocation-prone substrates inadvertently follow the same mechanism we outline here but perhaps the critical role of dislocations from the substrates to the growth of the 1D ZnO nanomaterials (NWs and NTs) was not recognized. Moreover, it was often commented that “seeding layers” of ZnO film or particles initially grown on sapphire26,33,34 or GaN35,36 substrates were crucial for the successful epitaxial growth of ZnO NWs. We suspect such “seeding layers” could in fact be the highly defective substrates that contain numerous screw dislocations that then initiate the dislocationdriven growth of ZnO NWs, that is, they likely fill the role of the GaN films intentionally used here. Generally speaking, GaN is not essential to our strategy, since dislocations exist in all crystals. Rather it is really a question of how many dislocations there are and can we “engineer” them to grow 1D architectures suited for specific purposes. In the future, such rational engineering of substrates or seeds to contain dislocations of a specific density, arrangement, character, and magnitude may enable the use of the dislocation-driven NW growth mechanism for producing 1D materials with controlled location and morphology.
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Acknowledgment. This research is supported by NSF (CAREER DMR-0548232). S.J. also thanks research Corporation Cottrell Scholar Award, DuPont Young Professor Grant, and Sloan Research Fellowship for support. S.A.M. was partially supported by a 3M Graduate Research Fellowship and by UW-Madison NSEC (NSF DMR 0832760). We thank Professor T. F. Kuech, Department of Chemical and Biological Engineering, UW-Madison for donating the samples of GaN thin films on sapphire used in this research. We thank J. M. Higgins for assistance in X-ray diffraction studies.
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Supporting Information Available. Experimental details. This material is available free of charge via the Internet at http://pubs.acs.org.
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