Nonepitaxial Growth of Hybrid Core-Shell Nanostructures with Large Lattice Mismatches

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Science  26 Mar 2010:
Vol. 327, Issue 5973, pp. 1634-1638
DOI: 10.1126/science.1184769

Perfect Mismatch

Heteroepitaxy, or the overgrowth of one crystalline material onto a second crystalline material, is a key fabrication method for making thin films and nanoparticles. But if the lattice mismatch between the two materials is too large or messy, fractured interfaces result. Zhang et al. (p. 1634) describe a synthesis strategy to obtain spherical nanoparticles with a core-shell architecture that does not depend on heteroepitaxy. Silver was deposited and converted to various semiconductors through a series of chemical transformations to yield structurally perfect single-crystal semiconductor shells on a gold core, despite mismatches approaching 50%.


We report a synthetic route to achieving nanoscale heterostructures consisting of a metal core and monocrystalline semiconductor shell with substantial lattice mismatches between them, which cannot be obtained by conventional epitaxial techniques. By controlling soft acid-base coordination reactions between molecular complexes and colloidal nanostructures, we show that chemical thermodynamics can drive nanoscale monocrystalline growth of the semiconductor shell with a lattice structure incommensurate with that of the core. More complex hybrid core-shell structures with azimuthal and radial nanotailoring of structures and compositions of the monocrystalline semiconductor shell are also demonstrated.

Growth of single-crystal semiconductor-based heterostructures with modulated composition is a prerequisite for exploring fundamental nanoscale semiconductor physics (1, 2) and can offer technological devices with optimum characteristics, including enhanced optical properties with high quantum yields (3), engineered electronic bandgaps (46), and various solid-state optoelectronic properties (79). Unintentional crystalline imperfections (such as polycrystallinity, dislocations, and other structural defects) lead to performance degradation or even premature failure of devices. For example, although the optical quality of semiconductor CdSe nanoparticles (NPs) could be improved by an overlayer of epitaxially grown CdS or ZnS, problems appear once the shell thickness becomes larger than the critical layer thickness (about two monolayers) due to the existence of strain-induced defects (3, 10, 11). Current methods that achieve high-quality monocrystalline heterostructures are all based on epitaxial growth, which requires moderate lattice mismatches (<2%) between the two different materials. This lattice-matching constraint is a severe obstacle, particularly for growth of core-shell nanostructures with (quasi-) spherical core NPs with highly curved surfaces that present many different crystallographic facets (12). In addition to such lattice-matching requirements, the issues related to differences in crystal structure, bonding, and other properties have been found to inhibit epitaxial growth of dissimilar hybrid materials such as monocrystalline semiconductors on metals (13).

Attempts to use epitaxy to achieve hybrid core-shell nanostructures have been unsuccessful, resulting in either polycrystalline semiconductor shells or anisotropic structures with segregation of the core and shell, thus limiting their usefulness (1418). We report a general nonepitaxial growth strategy that achieves precise control of the hybrid core-shell nanostructures, whereby the monocrystalline semiconductor shells are not dependent on the structure of the core NPs. In this approach, growth of the core-shell nanostructures is based on the Lewis acid-base reaction mechanism, where the entire nanostructure is spatially confined by an amorphous matrix (19). Because monocrystalline growth of the semiconductor shell is fully directed by chemical thermodynamic properties of reactions within the matrix, the shell’s lattice structure can be independent of that of the core NPs, thus circumventing the limitations imposed by epitaxial strategies.

Figure 1 highlights the results of Au-CdS growth, where the lattice mismatch between the two majority lattice planes of bulk Au and CdS is up to 43% (table S1). Large-scale transmission electron microscope (TEM) images (Fig. 1A) show uniform core-shell nanostructures. The monocrystalline feature of the CdS shell is evident in Fig. 1, B to E. Powder x-ray diffraction (XRD) patterns further confirm that this hybrid core-shell structure grows homogeneously as uniform crystalline domains, and the CdS shells form a hexagonal wurtzite lattice (Fig. 1F). The XRD features of the CdS shells do not show detectable strain-induced bond-length shifts when compared with bulk-indexed peaks, which is different from previous epitaxially grown core-shell nanostructures with much smaller lattice mismatches (20). The perfect crystallinity of the as-grown semiconductor shells is further revealed by angle-dependent TEM characterization under various viewing angles (Figs. 1, G to J). This hybrid Au-CdS nanostructure is stable for months without noticeable changes in the overall structure or degradation of the quality of the semiconductor shell.

Fig. 1

Au-CdS core-shell nanostructures with monocrystalline shell. (A) Typical TEM image showing uniform core-shell nanostructures. Scale bar, 20 nm. (B to E) High-resolution TEM images of core-shell nanostructures from (A). Whereas Au core NPs can manifest monocrystalline (B), single-fold twin (C), fivefold twin (D), and multiple-twin (E) lattice structures, all CdS shells are monocrystalline. The red lines highlight the lattice orientations within the Au core NPs. Scale bar, 5 nm. (F) XRD pattern of Au-CdS sample shown in (A). Bulk Au [red solid lines, Joint Committee on Powder Diffraction Standards (JCPDS) #04-0784] and wurtzite CdS (blue solid lines, JCPDS #41-1049) are also provided for reference and comparison. (Inset) A ball-and-stick molecular model of Au-CdS, illustrating a cubic core and wurtzite shell. (G to J) Angle-dependent high-resolution TEM characterization. The sample depicted has a larger shell thickness than the one in (A) to emphasize the extremely high-quality crystallinity of the shell. The CdS shell shows perfect monocrystalline features without detectable structural defects under a different viewing angle. Scale bar, 5 nm.

The steps of our synthesis protocol are outlined in the flowchart in Fig. 2A. We use Au-CdS as an illustrative example of the process to substantiate our proposed mechanism, in which each growth stage is characterized in detail by high-resolution TEM images (Fig. 2B) and XRD spectroscopy (Fig. 2C), as well as elemental analysis (fig. S1). Our nonepitaxial growth mechanism can be qualitatively understood on the basis of the thermodynamics and coordination chemistry of ionic transformation involved in the growth reactions. Controlling the thermodynamics associated with the chemical transformation processes can initiate and facilitate semiconductor monocrystalline growth in a well-defined amorphous matrix grown outside of the core NPs, and represents the key to our method.

Fig. 2

Nonepitaxial growth process and mechanism of hybrid core-shell nanostructures with substantial lattice mismatches. (A) Schematic of growth process. (B) Series of high-resolution TEM images highlighting different synthetic stages of Au-CdS growth. Scale bar, 5 nm. Red and yellow dashed lines are guides for the eye, distinguishing the core and shell boundaries, respectively. (C) Corresponding XRD patterns of different stages illustrated in (B). Bulk Au (red solid lines, JCPDS #04-0784), monoclinic Ag2S (green solid lines, JCPDS #14-0072), and wurtzite CdS (blue solid lines, JCPDS #41-1049) are provided for reference. Bulk Ag is not shown because its XRD pattern is very similar to that of Au.

Starting from the core NPs (stage S1), an overlayer of metal with soft Lewis acidity is grown onto the core (stage S2) (19, 21). For all hybrid core-shell structures, we choose a Ag metal overlayer based on the following considerations:

1) According to the theory of hard-soft acids and bases, silver cations behave as a strong acid (acid softness = +3.99) compared with many other common metal cations, such as Zn2+, Pb2+, and Cd2+ (fig. S2). Thus, silver cations can easily share their d electrons and coordinate with various soft bases via back-donating π bonds to form a rich family of organometallic complexes. The free energy of reaction (∆G) is qualitatively determined by the coordination stability of these complexes and can further govern the equilibria of reaction (22, 23). The high acid softness of silver therefore offers broad thermodynamic control of the synthetic process.

2) The silver layer can be grown onto a wide variety of core NPs (including metallic, magnetic, and semiconductor core NPs) with precise thickness control down to a single monolayer (21).

3) The electronegativity of silver is similar to that of many anions X (chalcogenides, As, P) (24). Under certain conditions (e.g., appropriate temperature and anion molecular complexes), the silver shells in stage S2 can be modified to form silver-compound shells (Ag2X) with an amorphous structure (stage S3) (the amorphous feature of Ag2S in stage S3 is confirmed from both high-resolution TEM and XRD patterns shown in Fig. 2, B and C), providing a crucial platform for the next chemical transformation stage, ultimately leading to monocrystalline growth. It has been demonstrated that nanoscale chemical transformations, such as cation exchange, represent a versatile route for converting one crystalline solid to another (25, 26). We show that this process can be harnessed to drive the single-crystal growth by carefully controlling the thermodynamic properties of the reaction (27):

Embedded Image

Tributylphosphine was selected because it is a soft base and can behave as a phase-transfer agent to transport metal ions (Mn+) to the surface of the core NPs by binding to free cations in solution (fig. S2). The high acid softness of Ag+ favors the exchange process between Ag+ in the amorphous matrix and Mn+ in solution as long as the softness of Mn+ is small enough to result in a positive ∆G. This in turn provides the impetus to initiate reorganization of the M2Xn crystalline lattice and to grow into a monocrystalline domain once Ag is completely expelled from the shell (stages S4a-S4b-S5) (fig. S2) (28). The processes from stages S3 to S5 can take from minutes to a few hours depending on the softness of the Mn+ in solution (19).

The effect of an amorphous versus crystalline phase of the Ag2X shell on the resulting crystalline quality, as well as geometry of the core-shell nanostructures, was also investigated. We observed that crystalline Ag2S shells typically led to phase segregation between the core and shell, forming nonconcentric anisotropic shapes (such as dumbbell nanostructures) (fig. S3). The CdS shells in such asymmetric nanostructures appeared as either polycrystalline or monocrystalline. By contrast, amorphous Ag2X shells not only provided a well-defined regime for cation exchange (thus defining the dimensions of the monocrystalline semiconductor shells in stage S5), but also promoted the motion of the ions inside the shells as well as the growth of the monocrystalline domain of M2Xn due to a reduction of interfacial and grain boundary energies between amorphous Ag2X and crystalline M2Xn (28, 29).

According to the above proposed growth mechanism, our technique should be readily applicable to other semiconductor hybrid systems as long as the softness of Mn+ is less than that of Ag+ to achieve positive ∆G (fig. S2). To demonstrate such versatility, Fig. 3 (and figs. S4 and S5) displays different combinations of uniformly grown hybrid systems using our technique (tables S1 and S2) (19). For all the systems the monocrystalline features of the semiconductor shell, whose lattice structure was determined from XRD measurements (fig. S4), are independent of the core NPs. Similar to the results of Au-CdS, XRD measurements of all hybrid core-shell nanostructures confirm that there is no evidence of strain-induced lattice changes in the semiconductor shell.

Fig. 3

Large-scale (left) and high-resolution (right) TEM images of different hybrid core-shell nanostructures with various combinations of the core and shell components. All semiconductor shells show monocrystalline features. Scale bars for large-scale and high-resolution TEM images are 20 and 5 nm, respectively. (A) Au-CdSe; (B) Au-CdTe; (C) FePt-CdS; (D) Au-PbS; (E) Au-ZnS; and (F) Pt-CdS.

One of the important merits of conventional epitaxial growth techniques is precise thickness control. In our nonepitaxial approach, similar precise control of the monocrystalline semiconductor shell layer is achievable because the preceding Ag growth (stage S2) is controllable down to a single monolayer (21). As an example, precise and independent control of the core and shell sizes in Au-CdS is shown in figs. S6 and S7. Because the optical properties of the semiconductor shell and metal core are dependent on their dimensions (due to quantum confinement and surface plasmon resonance effects, respectively), our independent control of both the shell and core dimensions can lead to tunable optical properties, as demonstrated in figs. S6 and S7. An additional advantage of our technique is the clear absence of a critical layer thickness intrinsic in expitaxial growth techniques (1). For instance, the monocrystalline CdS shells were grown up to 15 nm thick onto Au core NPs without detectable structural defects (fig. S8).

Our technique can be used to make more complex nanoscale heterostructures with precise structural and compositional tailoring. Figure 4 (and fig. S9) highlights three examples with independent azimuthal and radial engineering of hybrid core-shell nanostructures. In Fig. 4A, half of the amorphous Ag2S is first converted into monocrystalline CdS shells followed by sequential growth of PbS (this process can be confirmed by monitoring the compositional changes at each stage, as shown in fig. S10). Large-scale TEM images show that this controlled process can preserve the uniformity of the nanostructures (Fig. 4B). High-resolution TEM images reveal two distinct monocrystalline lattices split 50/50 with a Au core in the center, as evidenced by single-particle energy-dispersive x-ray spectroscopy (EDS) measurement (Fig. 4C). Enabled by such, multiple monocrystalline semiconductors can be seamlessly integrated into a single core-shell unit with a precisely tunable ratio of different components (fig. S11).

Fig. 4

Growth of complex hybrid core-shell nanostructures with tailored structures and compositions of the monocrystalline shells. (A to C) Control of the monocrystalline cation species within the shell: the case of Au-(CdS+PbS). (A) Schematic of the growth procedure. (B) Large-scale TEM image. Scale bar, 20 nm. (C) (Top) High-resolution TEM image. Blue and green dashed arc curves highlight the monocrystalline CdS and PbS regimes, respectively. CdS and PbS manifest distinct lattice planes that can be assigned to (100) and (220), respectively. Scale bar, 5 nm. (Bottom) Single-particle EDS measurements in the CdS and PbS regimes. Peaks from Cd, Pb, and S elements are highlighted. (D to F) Control of the monocrystalline anion species within the shell: the case of Au-CdS1-aSea. (D) Schematic growth procedure. (E) Large-scale TEM image. Scale bar, 20 nm. (Inset) High-resolution TEM image showing the monocrystalline alloy shell. Scale bar, 5 nm. (F) XRD patterns highlighting lattice evolution from CdSe to CdS with different ratio a.

Whereas Fig. 4, A to C, demonstrate integration of the monocrystalline cation species within the shell, Fig. 4, D to F, illustrate our fine control of anion species. Ternary single-crystal CdS1-aSea alloys represent an important semiconductor with a bandgap and lattice constant monotonically tunable by the ratio a. They can exhibit large nonlinear susceptibilities, as well as desirable photoconductive properties, and offer promising technological applications, such as a tunable laser (30). Figure 4D schematically shows the procedure for growing a monocrystalline CdS1-aSea alloy shell in a typical hybrid core-shell nanostructure, which begins with the reaction of the silver shell formed in stage S2 with a mixture of S and Se organocomplexes (with a predetermined ratio) developing an amorphous Ag2S1-aSea shell followed by sequential cation exchange with Cd2+. Figure 4E shows the uniformity as well as monocrystalline features of such an alloy shell grown onto Au core NPs. EDS measurements confirm that the atomic ratio of (S+Se)/Cd is very close to 1, which suggests formation of a ternary phase, but the ratio a is tunable in the monocrystalline shell layer (fig. S12). Powder XRD measurements reveal the lattice evolution of this ternary alloy as a continuous function of ratio a from pure wurtzite CdSe to wurtzite CdS; decreasing the S concentration increases the lattice constant of the monocrystalline alloy shell layer.

The excellent stability and monocrystalline quality of the as-synthesized core-shell nanostructures imply that the semiconductor shells can be further applied as a template for continual growth of different shell layers along the radial direction; one such example of hybrid Au-CdS-CdSe core-shell-shell nanostructure is presented in fig. S9. It has been demonstrated that through coupling with surface plasmons in metallic nanostructures, the luminescence intensity of the fluorophores can be significantly enhanced, depending on coupling strength (3133). Therefore, this radial engineering of the hybrid core-shell-shell nanostructures offers a precise and controllable way to explore such enhancements by tuning the thickness of the CdS shells, and may prove useful for interfacing with biological systems with enhanced bioimaging and biolabeling capability (fig. S13) (19).

Supporting Online Material

Materials and Methods

SOM Text

Figs. S1 to S13

Tables S1 and S2


References and Notes

  1. Materials and methods are available as supporting material on Science Online.
  2. This work was supported by Office of Naval Research YIP award (N000140710787), NSF CAREER award (DMR-0547194), and Beckman YIP grant (0609259093). Facility support was from Maryland Nanocenter and its Nanoscale Imaging Spectroscopy and Properties Laboratory (supported in part by the NSF as a Materials Research Science and Engineering Center shared experiment facility). We also thank G. Jenkins, T. Einstein, and J. R. Anderson for reading and polishing the manuscript.
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