Ternary-chalcopyrite structure ZnGeP2, CdGeAs2 (II-IV-V2) and AgGaS2, AgGaSe2 (I-III-VI2) compounds are currently of technological interest. They show the most promise for practical nonlinear optical applications in the areas of high-efficiency optical parametric oscillators and frequency up-converters for the infrared (ir) range as well as for widespectral-range optoelectronic devices. (See also the article by Schunemann, Schepler, and Budni in this issue.) However extensive realization of their potential has still not been achieved. One of the principal difficulties in the way to obtaining high-device-quality ZnGeP2, CdGeAs2, AgGaS2, and AgGaSe2 single crystals is undesired optical absorption in their transparency range near the fundamental band edge induced by lattice-related defects. This article summarizes selected aspects of dopant-incorporation techniques of these crystals including dopant choice of dopant material and monitoring of dopant incorporation as done in our laboratory.

In general for the ternary chalcopyrite compounds, doping-incorporation processes are more complicated in comparison to those of binary zinc-blende III-V compounds. The most common sources of dominant incorporation of acceptors and donors in as-grown chalcopyrites are believed to appear from (1) nonstoichiometric melts as well as by doping with different elements during the growth process and (2) incomplete removal of disorder on the cation sublattice during subsequent cooling. Furthermore the chalcopyrite structure II-IV-V2 undergoes a disorder-order phase transition upon cooling through approximately 1220 K for ZnGeP2 and 900 K for CdGeAs2. At these transition temperatures, solidification can be complicated also by supercooling phenomena, and the crystals transform from the cubic zinc-blende structure (where Zn atoms randomly fill cation sites) to the ordered chalcopyrite structure (e.g., when Zn and Ge occupy alternating cation sites in ZnGeP2).

The low-pressure synthesis of rather pure nanocrystalline diamond films from fullerene precursors suggests that for a small enough grain size the diamond structure may be energetically preferred over graphite. Because of the small grain size of typically about 15 nm in these films, a significant fraction of the carbon atoms is situated in the grain boundaries (GBs). The surprisingly high wear resistance of these films even after the substrate is removed and their high corrosion resistance suggest that the grains are strongly bonded. Grain-boundary carbon is also believed to be responsible for the absorption and scattering of light in these films, for their electrical conductivity, and for their electron-emission properties. In spite of all these indications of a critical role played by GB carbon in achieving the remarkable properties of nanocrystalline diamond films, to date the atomic structures of the GBs are essentially not known.

It is well-known that the electronic and optical properties of polycrystalline silicon films are significantly affected by the presence of GBs. For example GBs can provide active sites for the recombination of electron-hole pairs in photovoltaic applications. Also, in electronic devices such as thin-film transistors, GBs are known to play an important role. Because of silicon's strong energetic preference for sp3 hybridization over other electronic configurations, the structural disorder in silicon GBs is accommodated by a distortion of the tetrahedral nearestneighbor bonds and in the extreme by the creation of dangling bonds—that is, of three-coordinated Si atoms each having one unsaturated, bound electron in an otherwise more or less tetrahedrally coordinated environment.

Materials designed for rolling or sliding contact, like corrosion-resistant materials, can provide great cost savings to industry. So why can't such “tribomaterials” be designed based on materials properties that control friction and wear? In recent years, it has become clear that the properties we need to understand are not only those of the starting materials—whether bulk solids or engineered surfaces—but also those of the materials generated within the sliding (or rolling) contact, the so-called “third-body” materials. This article reviews third-body processes and their role in controlling friction and wear of practical surface treatments.

It may seem odd that properties like slipperiness and wear resistance can be easily described but are not scientifically understood. We can feel that polytetrafluoroethylene is slippery but know that it is not very wear-resistant because we can scratch it with a fork. We can sense low friction when we rub talcum powder between our fingers but know it cannot lubricate indefinitely because it gets ejected as we rub. Is it contradictory to ask for materials that are both slippery and wear-resistant? Slipperiness is associated with low shear strength whereas wear resistance is modeled on high hardness—that is, high shear strength. Before we can answer this seeming contradiction, it is useful to review some aspects of friction and wear.

More than 50 years ago, Bowden and Tabor explained how a low shearstrength film can reduce the friction coefficient between two higher shearstrength materials in sliding contact.

The phenomenon of electrorheology (ER), described in the guest editors' introduction to this issue, can largely be explained by electrostatic interactions between particles by an externally applied electric field. The purpose of this article is to review particle-polarization mechanisms active in ER suspensions and to highlight some poorly understood electrostatic phenomena that inhibit the commercialization of ER technology.

Figure 1 plots the apparent viscosity of a 20-wt% suspension of alumina particles in silicone oil as a function of shear rate for different electric field strengths. The viscosity changes are most pronounced at small deformation rates. At small shear rates the viscosity is proportional to E2 where E is the electric field strength.

Light emission in silicon has been intensively investigated since the 1950s when crystalline silicon (c-Si) was recognized as the dominant material in microelectronics. Silicon is an indirect-bandgap semiconductor and momentum conservation requires phonon assistance in radiative electron-hole recombination (Figure 1a, top left). Because phonons carry a momentum and an energy, the typical signature of phonon-assisted recombination is several peaks in the photoluminescence (PL) spectra at low temperature. These PL peaks are called “phonon replicas.” High-purity c-Si PL is caused by free-exciton self-annihilation with the exciton binding energy of ~11 meV. The TO-phonon contribution in conservation processes is most significant, and the main PL peak (~1.1 eV) is shifted from the bandgap value (~1.17 eV) by ~70 meV—that is, the exciton binding energy plus TO-phonon energy (Figure 1a).

Oxygen-implanted silicon-on-insulator (SOI) material, or SIMOX (separation by implantation of oxygen), is another chapter in the continuing development of new material technologies for use by the semiconductor industry. Building integrated circuits (ICs) in a thin layer of crystalline silicon on a layer of silicon oxide on a silicon substrate has benefits for radiationhard, high-temperature, high-speed, low-voltage, and low-power operation, and for future device designs. Historically the first interest in SIMOX was for radiation-hard electronics for space, but the major application of interest currently is low-power, high-speed, portable electronics. Silicon-on-insulator also avoids the disadvantage of a completely different substrate such as sapphire or gallium arsenide. Formation of a buried-oxide (BOX) layer by high-energy, high-dose, oxygen ion implantation has the advantage that the ion-implant dose can be made extremely precise and extremely uniform. However the silicon and oxide layers are highly damaged after the implant, so high-temperature annealing sequences are required to restore devicequality material. In fact SIMOX process development necessitated the development of new technologies for high-dose implantation and high-temperature annealing.

Progress made in the growth of “free-standing” (e.g., colloidal) quantum dots (see also articles in this issue by Nozik and Mićić, and by Alivisatos) and in the growth of semiconductor-embedded (“self-assembled”) dots (see also the article by Bimberg, Grundmann, and Ledentsov in this issue) has opened the door to new and exciting spectroscopic studies of quantum structures. These have revealed rich and sometimes unexpected features such as quantum-dot shape-dependent transitions, size-dependent (red) shifts between absorption and emission, emission from high excited levels, surface-mediated transitions, exchange splitting, strain-induced splitting, and Coulomb-blockade transitions. These new observations have created the need for developing appropriate theoretical tools capable of analyzing the electronic structure of 103–106-atom objects. The main challenge is to understand (a) the way the one-electron levels of the dot reflect quantum size, quantum shape, interfacial strain, and surface effects and (b) the nature of “many-particle” interactions such as electron-hole exchange (underlying the “red shift”), electron-hole Coulomb effects (underlying excitonic transitions), and electron-electron Coulomb (underlying Coulomb-blockade effects).

Interestingly, while the electronic structure theory of periodic solids has been characterized since its inception by a diversity of approaches (all-electron versus pseudopotentials; Hartree Fock versus density-functional; computational schemes creating a rich “alphabetic soup,” such as APW, LAPW, LMTO, KKR, OPW, LCAO, LCGO, plane waves, ASW, etc.), the theory of quantum nano-structures has been dominated mainly by a single approach so widely used that I refer to it as the “Standard Model”: the effective-mass approximation (EMA) and its extension to the “k · p” (where k is the wave vector and p is the momemtum). In fact, speakers at nanostructure conferences often refer to it as “theory” without having to specify what is being done. The audience knows.

The origins of the game of cricket are not known, but the first recorded game took place in London, England, on February 25, 1719. The basic rules of the game were formulated at a meeting of a self-appointed group of noblemen, held at the Star and Garter public house, Pall Mall, in 1774. These rules form the basis of the modern game.

Only under-arm bowling was allowed originally, and bats incorporated a curved blade to sweep the ball along the ground. Over-arm bowling where the ball was allowed to pitch or bounce was introduced late in the 18th century. This was quickly accepted and led to bats with straight blades to accommodate a variable bounce. Following a famous incident in 1793 in which a wide bat was used to completely cover the wickets, a new rule was quickly introduced to restrict the width of the blade to four and a quarter inches—a rule that survives to the present day.

The weight of the bat is not restricted by the rules, and the modern tendency is to use heavier bats. In the middle of the present century, a typical bat would weigh between 34 and 36 ounces. Currently bats of 40–44 ounces are widely available, and bats of 48 ounces and more are not uncommon among professional players. Although a heavier bat is widely believed to be capable of propelling the ball faster, there is a penalty. Heavier bats cause players to tire more quickly, and this affects their timing. Individual players tend to have a personal preferred weight of bat.

Over the past decade, we have been working to develop intelligent photonic-crystal materials with unique properties, which will be useful in a number of technological areas. These photonic-crystal materials utilize mesoscopically periodic arrays of spherical particles as their active optical elements and are easily fabricated chemically by the use of crystalline-colloidal-array (CCA) self-assembly techniques.

Crystalline colloidal arrays are mesoscopically periodic fluid materials, which efficiently diffract light meeting the Bragg condition. These photonic-crystal materials consist of arrays of colloidal particles that self-assemble in solution into either face-centered-cubic (fcc) or body-centered-cubic (bcc) crystalline arrays (Figure 1), with lattice constants in the mesoscale size range (50-500 nm). Just as atomic crystals diffract x-rays that meet the Bragg condition, CCAs diffract ultraviolet, visible, and near-infrared light, depending on the lattice spacing; the diffraction phenomena resemble those of opals, which are close-packed arrays of monodisperse silica spheres.

The CCA however can be prepared as macroscopically ordered arrays of non-close-packed spheres. This self-assembly is the result of electrostatic repulsions between colloidal particles, each of which has numerous charged surface functional groups. We have concentrated on the development of CCAs that diffract light in the visible spectral region and generally utilize colloidal particles of ~100-nm diameter. These particles have thousands of surface charges, which result from the ionization of surface sulfonate groups. The nearest-neighbor distances are often >200 nm.

Low-phonon energy glasses are desirable hosts for rare-earth (RE) ions because they enable emission from RE energy levels that would otherwise be quenched in high-phonon energy glasses. Such emissions are of interest for fiber amplifiers operating at telecommunications wavelength band s of 1.31, 1.46, and 1.55 μm, and for up-conversion lasers and three-dimensional displays.

Phonons are optical-frequency molecular vibrations in a material. If the RE energy level of interest lies only a few phonons in energy above the next lower lying level such as the 1G4 level of Pr3+, which is only 3,000 cm −1 above the 3F4, only three Si—O vibrational phonons (1,100 cm−1) are required to bridge the gap as shown in Figure 1. Thus any electrons excited to the 1G4 level via an external pump source will be deexcited to the 3F4 on down to the 3H4 ground state via phonons, and no radiation of usable light will be produced. This is why emission from the 1G4 level of Pr3+ is absent in silicates and why researchers have gone to great lengths to make low-phonon energy glasses.

Currently spectroscopists are studying many semiconductor quantum-dot (QD) systems in great detail because of their scientific and technological importance. However as in all nanostructure systems in which significant confinement energies exist, size fluctuations lead to inhomogeneous broadening of the spectral lines. This blurring of the spectra severely reduces the amount of information obtainable from spectroscopy. The finding has initiated an effort to isolate optically and study spectroscopically individual QDs. Studies involving individual QDs in most QD systems have been published. Here the results of a series of experiments are reviewed on GaAs QDs defined by interface fluctuations in narrow GaAs quantum wells. These experiments demonstrate the elegance and potential of single-QD spectroscopy.

An example of single-QD photoluminescence (PL) spectroscopy appears in Figure 1. The spectra shown were obtained at a temperature of 6 K by successively reducing the size of the laser spot on a GaAs quantum-well sample through the use of small apertures in a metal mask. The bottom trace is a PL spectrum obtained with a macroscopic laser spot diameter of 25μm. The spectrum shows two broad peaks corresponding to the recombination of excitons in parts of the quantum well that are either 10 or 11 monolayers wide (2.8 or 3.1 nm). The spectrum is strongly inhomogeneously broadened as shown most directly by a reduction in the aperture size. The relatively broad lines break up into a decreasing number of extraordinarily narrow PL spikes as the aperture is reduced to submicroscopic dimensions. These PL spikes arise from excitons localized in individual QD potentials. Remarkably the linewidth decreases from several meV in the ensemble-averaged spectrum (25-μm aperture) to 10s of μeV in the single QD spectra, corresponding to an effective improvement in resolution of two orders in magnitude. By probing individual QDs, it becomes possible to resolve directly a number of phenomena that previously were hidden in the inhomogeneous linewidth.

Phosphate glasses generally have lower glass-transition temperatures and greater thermal-expansion coefficients than silicate and borate glasses, and so are candidate materials for specialty glass-metal sealing applications. Phosphate glasses also have desirable optical properties (rare-earth stimulated-emission cross sections, low thermo-optical coefficients, ultraviolet transparency, etc.) and so are used as laser glasses. Recent developments of novel compositions for biomedical applications and as fast-ion conductors have also revived interest in the study of phosphate glass structures.

The structure/property relationships in simple phosphate glasses have not received nearly the attention paid to silicate and borate-based systems. This is partially due to the difficulty in preparing anhydrous compositions with more than 50-mol% P2O5. Ultraphosphate compositions (where the [O]/[P] ratio is less than 3) are highly reactive to ambient water. In addition because P2O5 sublimes at about 300°C, ultraphosphate glasses have not been studied to the extent that metaphosphate ([O]/[P] = 3) and polyphosphate ([O]/[P] > 3) compositions have.

On April 13, 1982, the Duke Power Company energized an experimental pad-mount distribution transformer in Hickory, North Carolina. The transformer, manufactured by General Electric, provided electric power to a local residence. That same month, the Georgia Power Company installed a similar transformer, made by Westinghouse Electric, atop a utility pole in Athens, Georgia. It supplied electricity for the exterior lights at the Westinghouse Newton Bridge Road plant. These devices shown in Figure 1 were unique among the nearly 40 million distribution transformers in service in the United States because their magnetic cores were made from an Fe–B–Si amorphous-metal alloy. This new material, produced by Allied-Signal (formerly Allied Chemical), was capable of magnetizing more efficiently than any electrical steel. By replacing grain-oriented silicon steel in the transformer cores, the amorphous metal reduced the core losses of the transformers by 75%.

Although distribution transformers are relatively efficient devices, often operating at efficiencies as high as 99% at full load, they lose a significant amount of energy in their use. Because of the number of units in service, coupled with the fact that the core material is continuously magnetized and demagnetized at line frequency, transformers account for the largest portion of the energy losses on electric power distribution systems. It is estimated that over 50 × 109 kWh are dissipated annually in the United States in the form of distribution transformer core losses. At today's average electricity generating cost of $0.035/kWh, that energy is worth over $1,500 million.

Diamond always has been considered an exciting material. In addition to its other outstanding properties, diamond's capacity for cold electron emission has become a “hot” topic of research in recent years. The electron emission from diamond films is important for both fundamental and applied purposes, which may be expressed by two questions as follows: “Why does diamond emit electrons?” and “How can an efficient cold cathode be made?”

The “diamond-emission era” started in 1991. That year three papers were published, reporting unexpectedly high electron emissivity from diamond. The first paper by Djubua and Chubun was very practical, having been written by device engineers. They tried different materials for the fabrication of pointed field emission cathodes and found that emitters made from diamondlike carbon (DLC) demonstrated a lower emission threshold compared to other materials. A second paper by Wang et al. reported a low emission threshold for chemical-vapor-deposition (CVD) diamond films, whereas a third paper by Geis et al. described the fabrication and operation of diamond cold cathodes.

Since 1991 hundreds of papers about electron emission from diamond and diamondlike materials have been published. The motivations for this increasing activity are prospective applications in vacuum microelectronics—that is, vacuum emission devices fabricated by modern microelectronic technologies. Experts expect that combining the physical advantages of emission devices and the technological progress in solid-state microelectronics will result in the development of an entirely new generation of high-performance electronic devices—among them, flat-panel displays and miniature microwave tubes.

In this article, we review the present status of experimental and theoretical work on friction at the interface between extended macroscopic bodies, rough on the micrometer scale. We show that systematic detailed studies of low-velocity friction using dynamical systems analysis, together with their shear response in the static state, provide a tool for investigating the physical processes taking place on the mesoscopic scale of real contacts between rough surfaces. This approach should shed light on the enduring question of the relationship between macroscopic friction and microscopic dissipative mechanisms. This still open issue has come back to the fore during the last decades, following considerable progress due to the development of “molecular tribometers.

Bowden and Tabor pointed out that, because nominally flat surfaces are in general rough on small scales, the real area of contact Ar (Figure 1) is only a small fraction ϕ of apparent contact area A0. On the other hand, they postulated the existence of a stressσs characteristic of the shear strength of the interface between a given couple of solids. Hence the friction force:

In this framework, the Amontons-Coulomb (AC) law F = μFN amounts to stating that Ar is proportional to the normal load FN where μ is the coefficient of friction.

When considering soft metals, Bowden and Tabor noticed that ϕ ≪ 1 entails that the nominal local pressure p on the real contacts—of the order of FN/(ϕA0)–generally overcomes the yield strength Y so that the contacting asperities flow plastically until p = H≈ 3Y, the “hardness” of the (softer) material. So, Ar = FN/H.

In recent years, chalcopyrites have distinguished themselves as the nonlinear optical materials of choice for mid- to far-infrared (ir) laser applications. In particular, AgGaSe2, ZnGeP2, and CdGeAs2 have demonstrated the highest conversion efficiencies and output powers in the wavelength range beyond 4 μm. The superior performance of these crystals arises from their high nonlinear optical coefficients (39 pm/V, 75 pm/V, and 236 pm/V, respectively), their relatively large birefringence (sufficient for phase matching), and advances in crystal growth and processing that have improved transparency and eliminated cracking.

The two most direct approaches to generating laser output in the mid-ir (in particular the 3–5-μm atmospheric transmission window) are (1) shifting the output of a solid-state laser to longer wavelengths via optical parametric oscillation (OPO), or (2) doubling the frequency of a CO2 laser (9–11 μm) via second-harmonic generation (SHG). The OPO approach offers the advantage of tunability combined with potentially more compact and efficient solid-state lasers, whereas the SHG approach benefits from the greater maturity of high-power CO2 laser technology. A third process, difference-frequency generation (DFG), is also a 3-wave interaction similar to OPO that can be used to mix two photons (from two lasers or from an OPO) to produce longer wavelength photons (corresponding to the small difference frequency) over a large spectral range. The optimum chalcopyrite crystal (AgGaSe2, ZnGeP2, or CdGeAs2) for a given approach depends on a complex combination of material parameters described in the sections that follow.