that is liquidlike in character especially at high substrate temperatures. Coales-cence decreases the island density, resulting in local denuding of the substrate where rurther nucleation can then occur. Crystallographic facets and orienta-tions are frequently preserved on islands and at interfaces between initially disoriented, coalesced particles. Coalescence continúes until a connected net-work with unfilled channels in between develops. With rurther deposition, the channels fill in and shrink, leaving isolated voids behind. Finally, even the voids fill in completely, and the film is said to be continuous. This collective set of events occurs during the early stages of deposition, typically accounting for the first few hundred angstroms of film thickness.
The many observations of film formation have pointed to three basic growth modes: (1) island (or Volmer-Weber), (2) layer (or Frank-van der Merwe), and (3) Stranski-Krastanov, which are illustrated schematically in Fig. 5-2. Island growth occurs when the smallest stable clusters nucleate on the substrate and grow in three dimensions to form islands. This happens when atoms or molecules in the deposit are more strongly bound to each other than to the substrate. Many systems of metáis on insulators, alkali halide crystals, graphite, and mica substrates display this mode of growth.
The opposite characteristics are displayed during layer growth. Here the extensión of the smallest stable nucleus occurs overwhelmingly in two dimen sions resulting in the formation of planar sheets. In this growth mode the atoms are more strongly bound to the substrate than to each other. The first complete monolayer is then covered with a somewhat less tightly bound second layer. Providing the decrease in bonding energy is continuous toward the bulk crystal valué, the layer growth mode is sustained. The most important example of this growth mode involves single-crystal epitaxial growth of semiconductor films.
The layer plus island or Stranski-Krastanov (S.K.) growth mechanism is an intermedíate combination of the aforementioned modes. In this case, after forming one or more monolayers, subsequent layer growth becomes unfavor-able and islands form. The transition from two- to three-dimensional growth is not completely understood, but any factor that disturbs the monotonic decrease in binding energy characteristic of layer growth may be the cause. For example, due to film-substrate lattice mismatch, strain energy accumulates in the growing film. When released, the high energy at the deposit-intermediate-layer interface may trigger island formation. This growth mode is fairly common and has been observed in metal-metal and metal-semiconductor systems.
At an extreme far removed from early film formation phenomena is a regime of structural effects related to the actual grain morphology of polycrys-talline films and coatings. This external grain structure together with the internal defect, void, or porosity distributions frequently determines many of the engineering propcrties of films. For example, columnar structures, which interestingly develop in amorphous as well as polycrystalline films, have a profound effect on magnetic, optical, electrical, and mechanical properties. In this chapter we discuss how different grain and dcposit morphologies evolve as a function of deposition variables and how some measure of structural control can be exercised.
Capillarity Theory
Atomistic Nucleation Processes
Cluster Coalescence and Depletion
Experimental Studies of Nucleation and Growth
Grain Structure of Films and Coatings
Amorphous Thin Films
Amorphous Thin Films
Systems, Structures, and Transformations
Amorphous or glassy materials have a structure that exhibits only short-range order or regions where a predictable placement of atoms occurs. However, within a very few atom spacings, this order breaks down, and no long-range correlation in the geometric positioning of atoms is preserved. Although bulk amorphous materials such as silica glasses, slags, and polymers are well known, amorphous metáis were originally not thought to exist. An interesting aspect of thin-film deposition techniques is that they facilitate the formation of amorphous metal and semiconductor structures relative to bulk preparation methods.
As noted, production of amorphous films requires very high deposition rates and low substrate temperatures. The latter immobilizes or freezes adatoms on the substrate where they impinge and prevents them from diffusing and seeking out equilibrium lattice sites. By the mid-1950s Buckel (Ref. 25) producedamorphous films of puré metáis such as Ga and Bi by thermal evaporation onto substrates maintained at liquid helium temperatures. Alloy metal films proved easier to deposit in amorphous form because each component effectively inhibits the atomic mobility of the other. This meant that higher substrate temperatures (~ 77 K) could be tolerated and that vapor quench rates did not have to be as high as those required to produce puré amorphous metal films. Although they are virtually impossible to measure, vapor quench rates in excess of 10l0 °C/sec have been estimated. From laboratory curiosities, amorphous Si, Se, GdCo, and GeSe thin films have been exploited for such applications as solar cells, xerography, magnetic bubble memories, and high-resolution optical lithography, respectively.
Important fruits of the early thin-film work were realized in the later research and development activities surrounding the synthesis of bulk amorphous metáis by quenching melts. Today continuously cast ribbon and strip of metallic glasses (Metglas) are commercially produced for such applications as soft magnetic transformer cores and brazing materials. Cooling rates of ~ 106 °C/sec are required to prevent appreciable rates of nucleation and growth of crystals. Heat transfer limitations restrict the thiclcness of these metal glasses to less than 0.1 mm. In addition to achieving the required quench rates, the alloy compositions are critical. Most of the presently known glass-forming binary alloys fall into one of four categories (Ref. 26):
1. Transition metáis and 10-30 at% semimetals
2. Noble metáis (Au, Pd, Cu) and semimetals
3. Early transition metáis (Zr, Nb, Ta, Ti) and late transition metáis (Fe, Ni, Co, Pd)
4. Alloys consisting of IIA metáis (Mg, Ca, Be)
In common, many of the actual glass compositions correspond to where "deep" (low-temperature) eutectics are found on the phase diagram.
Amorphous thin films of some of these alloys as well as other metal alloysand virtually all elemental and compound semiconductors, semimetals, oxides, and chalcogenide (i.e., S-, Se-, Te-containing) glasses have been prepared by a variety of techniques. Amorphous Si films, for example, have been deposited by evaporation, sputtering, and chemical vapor deposition techniques. In addition, large doses of ion-implanted Ar or Si ions will amorphize surface layers of crystalline Si. Even during ion implantation of conventional dopants,local amorphous regions are created where the Si matrix is suffíciently damaged, much to the detriment of device behavior. Lastly, pulsed láser surface melting followed by rapid freezing has produced amorphous films in Si as well as other materials.
Au - Co and Ni - Zr Amorphous Films
It is instructive to consider amorphous Co-30Au films since they have been well characterized structurally and through resistivity measurements (Ref. 27). The films were prepared by evaporation from independently heated Co and Au sources onto substrates maintained at 80 K. Dark-field electrón microscope images and corresponding diffraction patterns are shown side by side in Fig. 5-19. The as-deposited film is rather featureless with a smooth topography, and the broad halos in the diffraction pattern cannot be easily and uniquely assigned to the known lattice spacings of the crystalline alloy phases in this system. Both pieces of evidence point to the existence of an amorphous phase whose structural order does not extend beyond the next-nearest-neighbor distance. The question of whether so-called amorphous films are in reality microcrys-talline is not always easy to resolve. In this case, however, the subsequent annealing behavior of these films was quite different from what is expected of fine-grained crystalline films. Heating to 470 K resulted in the face-centered cubic diffraction pattern of a single metastable phase, whereas at 650 K, lines corresponding to the equilibrium Co and Au phases appeared. Resistivity changes accompanying the heating of Co-38Au (an alloy similar to Co-30Au) revealed a two-step transformation as shown in Fig. 5-20. Beyond 420 K there is an irreversible change from the amorphous structure to a metastable FCC crystalline phase, which subsequently decomposes into equilibrium phases above 550 K. The final two-phase structure is clearly seen in Fig. 5-19. The high resistivity of the amorphous films is due to the enhanced electrón scattering by the disordered solid solution. Crystallization to the FCC structure reduces the resistivity, and phase separation, further still.
Both the amorphous and metastable phases are stable over a limited tempera-ture range in which the resistivity of each can be cycled reversibly. Once the two-phase structure appears, it, of course, can never revert to lessthermody-
namically stable forms. This amorphous-crystalline transformation apparently proceeds in a manner fírst suggested by Ostwald in 1897. According to the so-called Ostwald rule, a system undergoing a reaction proceeds from a less stable to a final equilibrium state through a succession of intermediate metastable states of increasing stability. In this sense, the amorphous phase is akin to a quenched liquid phase. Quenched films exhibit other manifestations of thermodynamic instability. One is increased atomic solubility in amorphous Confounding the notion that rapid quenching of liquids or vapors is required to produce amorphous alloy films is the startling finding that they can also be formed by solid-state reaction. Consider Fig. 5-21, which shows the result of annealing a bilayer couple consisting of pure polycrystalline Ni and Zr films at 300 °C for 4 h. The phase diagram predicts negligible mutual solid solubility and extensive intermetallic compound formation; surprisingly, an amorphous NiZr alloy film is observed to form. Clearly, equilibrium compound phases have been bypassed in favor of amorphous phase nucleation and growth, askinetic considerations domínate the transformation. The effect, also observed in Rh-Si, Si—Ti, Au-La, and Co-Zr systems, is not well understood. Apparently the initial bilayer film passes to the metastable amorphous state via a lower energy barrier than that required to nucleate stable crystalline com-pounds. However, the driving forcé for either transformation is similar. Unlike other amorphous films, extensive interdiffusion can be tolerated in NiZr without triggering crystallization.
No hay comentarios:
Publicar un comentario