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Fig. 3. (A) Stress displacement curves calculated from first principles. The first point of each line indicates the GB expansion with respect to the bulk interlayer distance in the (012) direction (d_bulk = 0.785 A). The relative GB expansion from the first point of each line indicates the displacement between fracture surfaces. The curves are plotted for the following cases: bulk (open squares), clean GB (solid squares), GB0 4/4 (open circles), GB0 4/4, GB2 1/4 (solid circles), GB0 4/4, GB2 2/4 (open upward triangles), GB0 4/4, GB2 3/4 (solid upward triangles), and GB0 4/4, GB2 4/4 (open downward triangles). (B) These tensile tests were performed in such a way that two crystals divided by a fracture plane were gradually separated in the direction of the c axis, with no relaxation for atomic positions and ab axes' lengths. The tensile stress (GPa) perpendicular to the GB(012) plane is calculated from the derivative of total energies. Here, the fracture plane lies between the GB0(GB1) and GB2 plane in order to measure the tensile stress between the two fractured surfaces.

Tensile stress

Tensile stress

GB0,GB11 Displacement | of fracture - surfaces

( GB expansion + displacement ) / d_bulk

Fig. 3. (A) Stress displacement curves calculated from first principles. The first point of each line indicates the GB expansion with respect to the bulk interlayer distance in the (012) direction (d_bulk = 0.785 A). The relative GB expansion from the first point of each line indicates the displacement between fracture surfaces. The curves are plotted for the following cases: bulk (open squares), clean GB (solid squares), GB0 4/4 (open circles), GB0 4/4, GB2 1/4 (solid circles), GB0 4/4, GB2 2/4 (open upward triangles), GB0 4/4, GB2 3/4 (solid upward triangles), and GB0 4/4, GB2 4/4 (open downward triangles). (B) These tensile tests were performed in such a way that two crystals divided by a fracture plane were gradually separated in the direction of the c axis, with no relaxation for atomic positions and ab axes' lengths. The tensile stress (GPa) perpendicular to the GB(012) plane is calculated from the derivative of total energies. Here, the fracture plane lies between the GB0(GB1) and GB2 plane in order to measure the tensile stress between the two fractured surfaces.

tensile strength for the bulk case in the (012) direction is 30 GPa, which is close to the ideal tensile strength in the (100) direction (37 GPa) that is estimated from experimental data (14). Second, we see the result for a clean GB case. In this case, the c-axis length is already expanded by 0.4 A, which is the GB expansion (SOM). The resulting tensile strength of a clean GB is 26 GPa, which is slightly smaller than that of the bulk case. For the GB that includes four S atoms at four GB0 sites (GB0 4/4), the GB is already expanded by 0.9 A. The resulting tensile strength reduces to 22 GPa (15). Further tensile tests were performed by substituting S atoms one by one for Ni atoms at the four GB2 sites. The GB expansion further increased and the resulting tensile strength decreased. The decrease of the tensile strength was clearly in proportion to the increase in GB expansion. The final tensile strength was 2.5 GPa, which is only 10% of the clean GB case. This is comparable to the ideal shear strength of Ni,

2.4 to 7.1 GPa, estimated by Mackenzie's method (14) (SOM), and is only five times larger than the experimental tensile strength of Ni, 493 MPa (16). Although these tensile tests are within the limits of only one kind of S5(012) GB, the strong decohesion can occur for any kind of GB that includes a sufficient number of segregation sites having large segregation (binding) energies and neighboring each other within a distance about

2.5 A. As for the other impurity elements, the calculated results for Ni-P (phosphorous) and Ni-B (boron) systems are in agreement with experimental facts (SOM).

The estimated critical concentration of S for the strong GB decohesion agrees well with the experimental data. In Heuer et al. (1, 2), the GB region is assumed to be within about 5 A depth from the GB plane, from analysis of Auger electron spectroscopy. Similarly, we assume the GB region from the GB0 site plane to the GB±7 or GB±8 site plane within about 5 A from the GB, in which 56 or 64 atomic sites are included. Then we assume that this strong decohesion begins in a range of S occupations from (GB0 4/4, GB2 2/4) to (GB0 4/4, GB2 4/4). In this region, the tensile strength of GB is less than half of that for the clean GB case. Thus, the critical S concentration is estimated to be 9 to 14 atomic % S (6 out of 64 to 8 out of 56), which is in agreement with the experimentally determined concentration (10 to 16 atomic % S) (1, 2).

References and Notes

1. J. K. Heuer, P. R. Okamoto, N. Q. Lam, J. F. Stubbins,

2. J. K. Heuer, thesis, Univ. of Illinois at Urbana-

Champaign (2000).

3. R. P. Messmer, C. L. Briant, Acta Metall. 30, 457 (1982).

4. M. E. Eberhart, K. H. Johnson, R. M. Latanision, Acta

5. J. R. Rice, J.-S. Wang, Mater. Sci. Eng. A 107, 23 (1989).

6. R. Wu, A. J. Freeman, G. B. Olson, Science 265, 376 (1994).

7. D. McLean, Grain Boundaries in Metals (Oxford Univ. Press, London, 1957).

8. Materials and methods are available as supporting material on Science Online.

9. G. Kresse, J. Hafner, Phys. Rev. B 47, R558 (1993).

10. G. Kresse, J. Furthmüller, Phys. Rev. B 54, 11169 (1996).

11. G. Kresse, D. Joubert, Phys. Rev. B 59, 1758 (1999).

12. A. Larere, M. Guttmann, P. Dumoulin, C. Roques-Carmes, Acta Metall. 30, 685 (1982).

14. A. Kelly, N. H. Macmillan, Strong Solids (Clarendon Press, Oxford, ed. 3, 1986).

Porous solids, such as zeolites, ordered mesoporous materials (M41S type), aerogels, and photonic crystals, find applications in such technologies as catalysis, sorption, filtration, thermal and acoustic insulation, and optical switches. Their distinctive properties are a function of the high internal surface area; the size, shape, and degree of interconnectedness of the pores; and the chemical nature of the framework. The pore structure can act as either a thoroughfare or a trap, or can even modulate the optical properties (as in photonic crystals), whereas the chemical interface is where catalysis, sensing, and sorption take place. The vast majority of porous solids are oxide based, thereby limiting the range of chemical functionality and physical properties that can be engineered into these materials. Recently, there has been sustained effort to prepare metal chalcogenide (sulfide, selenide, and telluride) analogs that are semiconducting substitutes for traditional oxide materials, the majority of which are electrically insulating.

There has been considerable success with the creation of microporous (<2 nm pores) zeolitic chalcogenide structures (1-3), as well as macroporous (>50 nm) materials created

Department of Chemistry, Wayne State University, Detroit, MI 48202, USA.

*To whom correspondence should be addressed. E-mail: [email protected]

15. G. S. Painter, F. W. Averill, Phys. Rev. Lett. 58, 234 (1987). (They pointed out the importance of lattice expansion for the embrittlement of Ni due to expanded valence orbitals of S.)

16. A. Buch, Pure Metals Properties (ASM International, Materials Park, OH, 1999).

17. T. Miyahara, K. Stolt, D. A. Reed, H. K. Birnbaum, Scr. Metall. 19, 117 (1985).

18. N. Barbouth, J. Oudar, C. R. Acad. Sci. Ser. C 269, 1618 (1969).

20. G. M. Mehrotra, V. B. Tare, J. B. Wagner Jr., J. Electrochem. Soc. 132, 247 (1985).

21. We thank J. K. Heuer, P. R. Okamoto, N. Q. Lam, and J. F. Stubbins for their fundamental and precise ex-

by sphere-templating strategies (photonic crystals) (4-6). In contrast, attempts to prepare similarly ordered mesoporous (2 to 50 nm) materials of metal chalcogenides have met with less success. Efforts have largely emphasized surfactant templating methodologies, originally developed for M41S-type materials, that employ metal chalcogenide building blocks and transition metal ion linkers. The resulting materials demonstrate longrange ordering of the surfactant/chalcogenide framework; however, attempts to remove the surfactants to access the pores invariably result in structure collapse (7-12).

We have developed a strategy for the production of mesoporous nanostructured metal chalcogenides based on the assembly of discrete nanoparticles to produce aerogel-type frameworks. Aerogels are highly porous architectures that arise when the solvent in a wet polymeric gel is replaced by air while the structural integrity of the framework is maintained (13). They are defined by a nanonet-work of particles, each with a low degree of connectivity, resulting in porosity that is intrinsic to an aerogel (regardless of whether a monolith is produced) and not to a precipitate. Aerogels are typically produced by supercritical drying because other routine techniques for drying gels (e.g., vacuum extraction or heating) result in collapse of the pore structure due to capillary forces and, ultimately, in the formation of dense aggre-

periments for S-induced embrittlement of Ni. Part of this research is supported by ACT-JST.

Supporting Online Material

www.sciencemag.org/cgi/content/full/1104624/DC1

Materials and Methods

SOM Text

Figs. S1 to S3

Tables S1 to S3

References and Notes

30 August 2004; accepted 07 December 2004

Published online 6 January 2005;

10.1126/science.1104624

Include this information when citing this paper.

gates (xerogels). Unlike the mesostructured materials reported to date, there is no longrange order in aerogels; however, the range of pore sizes and interconnectedness should be optimal for properties dependent on facile transport of small molecules through the interior space, such as photocatalysis and sensing (14). Such properties have only begun to be exploited in aerogels and are usually achieved through the incorporation of conductive or fluorescent moieties into conventional silica aerogels (15, 16). In contrast, the semiconducting and luminescent nature of the aerogels presented herein is a consequence of the intrinsic properties of the metal chalcogenide aerogel framework.

The methodology for chalcogenide aerogel formation is simple, comprising three steps: (i) Nanoparticle formation/thiolate capping, (ii) gelation through controlled surface-group loss, and (iii) supercritical CO2 drying to maintain the pore architecture (17-19). Both room-temperature reverse-micellar strategies (CdS, ZnS, PbS, and CdSe) and high-temperature arrested-precipitation techniques (CdSe) can be employed for nanoparticle production. The controlled loss of surface groups to reveal reactive sites for nanoparticle condensation is achieved by chemical oxidation of thiolate capping groups (17) or, in the case of CdSe, by photo-oxidation of the capping groups (20). This step is pivotal for the preparation of porous gel frameworks, because rapid surface-group loss leads to the formation of dense aggregates and precipitation. The resultant gels are exchanged multiple times with acetone to remove the oxidized disulfide by-product and then are dried at ~40°C by using supercritical CO2. For comparison, samples were also permitted to air dry on the benchtop, producing xerogels.

The chalcogenide aerogels appear to be morphologically similar to base-catalyzed silica aerogels, as depicted for CdS and CdSe (Fig. 1, A and B). The nanoparticle building blocks that make up the pearl-necklace morphology are visible, as is the presence of mesopores (2 to 50 nm). Unlike traditional silica aerogels, the building blocks appear to be crystalline, as evidenced by the presence of lattice fringes (Fig. 1C), which

Porous Semiconductor Chalcogenide Aerogels

Jaya L. Mohanan, Indika U. Arachchige, Stephanie L. Brock*

Chalcogenide aerogels based entirely on semiconducting II-VI or IV-VI frameworks have been prepared from a general strategy that involves oxidative aggregation of metal chalcogenide nanoparticle building blocks followed by supercritical solvent removal. The resultant materials are mesoporous, exhibit high surface areas, can be prepared as monoliths, and demonstrate the characteristic quantum-confined optical properties of their nanoparticle components. These materials can be synthesized from a variety of building blocks by chemical or photochemical oxidation, and the properties can be further tuned by heat treatment. Aerogel formation represents a powerful yet facile method for metal chalcogenide nanoparticle assembly and the creation of mesoporous semiconductors.

correspond to the (102) planes of the hexagonal wurtzite (CdS) structure. This is also in contrast to what is observed in the surfactant-templated mesostructured materials where the framework itself was amorphous, despite the presence of long-range order due to the surfactant substructure. Electron diffraction patterns of CdS aerogels confirm the crystal-linity, revealing four rings. These could likewise be indexed to the wurtzite structure (fig. S1).

Notably, monoliths can be achieved for CdS and ZnS with bulk densities as low as 0.07 g/cm3 and 0.35 g/cm3, respectively, representing 1.4 to 8.7% of the density of a single crystal (CdS, 4.83 g/cm3; ZnS, 4.04 g/cm3). Figure 1D shows a representative CdS aerogel monolith (right) compared with the wet gel (center). Only a small volume loss is observed upon aerogel formation, in contrast to a sample dried on the benchtop (xerogel) (Fig. 1D, left). If the wet gels are washed before benchtop drying, xerogel monoliths cannot be obtained at all, because fragmentation occurs upon drying.

An assessment of the porosity of the chalcogenide aerogels and xerogels is achieved by using N2 adsorption/desorption analysis (Table 1). All samples exhibit a type IV isotherm with H3-type hysteresis loop attributed to an interconnected mesoporous system with a broad pore-size distribution

Metal Chalcogenide Aerogels Xerogels

(fig. S2) (19, 21, 22). The small upward hysteresis exhibited by the isotherms suggests a cylindrical pore geometry (23). Pore-size distribution plots of the aerogels confirm the presence of a range of pores extending from the meso to the macro regime, with averages of 15 to 45 nm (fig. S2 and Table 1). The xerogels (washed) have a considerably narrower size distribution, with average pore diameters in the lower mesoporous range (CdS is 5 to 7 nm for a xerogel versus 29 to 30 nm for an aerogel).

The surface areas achieved for the chalco-genide aerogels range from 120 to 250 m2/g and are comparable to those obtained for many oxide-based aerogels. For example, aerogels of vanadia (24) are reported to have surface areas of 150 to 280 m2/g, whereas those of manganese oxide (25) have surface areas of up to 210 m2/g. Silica sets the benchmark, with surface areas of up to 1600 m2/g (600 m2/g is a typical value) (13). When compared on a per mole basis, our CdS aerogels with mean surface area of 245 m2/g are equivalent to a silica aerogel of surface area 590 m2/g. In contrast, the washed xerogels exhibit considerably lower surface area (47 m2/g), not unlike values obtained for precipitated CdS nanoparticles (56 m2/g) (26).

Despite these aerogels having a three-dimensional connected chalcogenide network, they nevertheless retain the optical

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