High Deposition Rate

Due to the very high deposition rate and use of single pass welding, material distortion is minimal.

From: Applied Welding Engineering, 2012

Chapters and Articles

Aerosol techniques for manufacturing piezoelectric materials

J. Akedo, in Advanced Piezoelectric Materials, 2010

14.4.1 Deposition ratio and influence of starting powder properties

High deposition rates can be achieved easily with AD because the source material is already in particle state form. These deposition rates are at least 30 times higher than other conventional thin film formation methods. Therefore, AD is an attractive manufacturing process due to high throughput. The particle velocity and concentration in the aerosol jet at the nozzle increase with increasing carrier gas flow rate, resulting in increased deposition efficiency. Starting particle properties, such as the average particle size and the size distribution, mechanical and surface properties can also strongly affect the deposition efficiency. [41]

The deposition rates for PZT AD layers using powders subjected to various milling procedures are indicated in Fig. 14.8. It can be seen that, by increasing the milling time, the deposition rate of the PZT layer significantly increased and reached a maximum of 73 μm/min for a 5 mm2 deposition area when powder was milled for 5 h. This value is 30 times higher than that for a starting powder without the milling procedure. An interesting fact is that the deposition rate decreased as a result of further milling to 30 h. It is assumed that particle surface properties (for example, surface activation, defects, gas absorption) will change by longtime milling making them less probable to be deposited in the same conditions.

14.8. Deposition rate for PZT film formation at room temperature using powder milled for different duration times with (black bar) / without (hatched bar) heat-treatment procedure at 800 °C for 4 h in air [41].

The milling procedure is also strongly influenced by the film density. The cross-sectional SEM images of the deposited layers vs. the milling time are shown in Fig. 14.9. With the increase of powder milling time, density and hardness of the deposited layer decreased. On the other hand, for 5 h and 30 h milling procedures, grain images of starting powders with the diameters from 100 to 300 nm and porous structures were distinct in the deposited layer images. At the same time, color and transmittance of layers were markedly changed from yellow to white due to the increase of optical scattering. [41] Thus the milling procedure is applicable to control the porosity of ceramic layers deposited by the AD process.

14.9. SEM images of PZT films deposited on glass substrates at room temperature using powders milled for different duration times [41].

The starting powder particle size and shape strongly influence the RTIC phenomenon in the AD process. If spherical α-Al2O3 ultra-fine particles with average particle size around 50 nm were used, after AD deposition the films have a pressed-like structure and the RTIC phenomenon could not be observed even if the ejecting particle velocity from the nozzle was over 400 m/sec and particle size was very small (Fig. 14.10(a)). In contrast, if non-spherical α-Al2O3 fine powder with average particle size around 1 μm was used, the deposited particles on the substrate were consolidated at room temperature and RTIC phenomenon was observed even for particle velocities around 200 m/sec, as shown in Fig. 14.10(b). As a result, high density and transparent ceramic layers were obtained. These results are explained by aerodynamic properties of the particle jet flow near the substrate. If particle size and weight are too small, the particle follows the carrier gas flow as shown schematically in Fig. 14.11. Therefore the particle velocity normal to the substrate is largely decreased and is not high enough to obtain RTIC phenomenon. Detailed further investigation about the particles’ aerodynamic properties in the AD process still needs to be conducted.

14.10. Influence of starting particle diameter for RTIC phenomenon on AD method: (a) average diameter: 50 nm, (b) average diameter: 700 nm.

14.11. Particles’ trajectories in an aerosol jet flow near substrate in AD method.

An advantage of AD over conventional thin film and thermal spray coating methods is that substrate surface does not need pre-cleaning to achieve good deposition. During the initial deposition stage, the particles impacting the substrate will act as cleaning agents in a similar way as in sand-blasting processes. Surface contaminants such as dirt and oils are removed by the initial particle collisions. The deposition automatically begins when the surface becomes sufficiently clean. The film adhesive strength to glass and metal substrates may be in excess of 30 MPa, because an anchoring layer having a thickness of about 100–200 nm was formed in the interface between the substrate and the deposited layer. To obtain maximum adhesive strength, a substrate with suitable hardness and elasticity is needed to allow the formation of the anchoring layer. A substrate that is very soft will be etched by the particle jet flow and the deposition will not occur. On the other hand, when a substrate with large hardness value is used, the adhesion strength between the deposited layer and substrate is weak and the film may easily peel off.

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Innovations in laser cladding and direct laser metal deposition

C. Leyens, E. Beyer, in Laser Surface Engineering, 2015

8.4.2 Induction-assisted laser cladding

The second approach for high deposition rates in laser cladding is based on hybrid processing. Therefore, a head was developed to integrate a module for additional inductive heating during the laser process (Figure 8.8). This hybrid technology concept further improves the deposition rate and the energy efficiency of the entire process; it is compatible with diode, fiber, and disc lasers. The localized inductive heating directly supports the laser beam and compensates heat losses, which leaves more laser energy to melt a larger amount of powder. As a result, deposition rates increase by a factor of 2-2.5. The coaxial nozzle designed for this application can handle a powder throughput of up to 18 kg/h. Typical practical deposition rates for INCONEL 625 are about 14-16 kg/h with the simultaneous application of 8 kW laser and 12 kW induction power.

Figure 8.8. Powder cladding head with full turn induction unit for cylindrical components.

Another potential application of this hybrid process is to use the induction coupling to modify coating strategies with tailored heat management. The additional heat leads to a significantly increased cooling time t8/5, as well as to decreased spatial temperature gradients (Figure 8.9). This configuration facilitates processing of crack free, ultrahard and wear-resistant metal alloys [10]. A good example is a protective coating made from Stellite 20 with a hardness exceeding 60 HRC.

Figure 8.9. Temperature profile at laser cladding without (a) and with (b) inductive preheating.

The hybrid process also leads to substantial cost reduction of the technology. The specific investment costs for the required energy sources are reduced by 50% while the overall energetic efficiency of the process is doubled.

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Cadmium telluride (CdTe) thin film solar cells

Kazi Sajedur Rahman, in Comprehensive Guide on Organic and Inorganic Solar Cells, 2022

3.1.1.6.4 Vapor transport deposition

Basically, VTD permits very high deposition rate at high substrate temperature and pressure forthcoming 0.1 atm onto moving substrates (McCandless & Sites, 2003). Whereas CSS is diffusion-restricted, VTD works by convective exchange of a vapor stream soaked with Cd and Te to the substrate, where supersaturation of the Cd and Te vapors creates condensation and reaction to form CdTe (McCandless & Sites, 2003). The CdTe source entails a heated chamber comprising solid CdTe in which the carrier gas combines with Cd and Te vapors and is exhausted through a slit over or under the moving substrate at a distance of ~1 cm (McCandless & Sites, 2003). The geometrical arrangement of the source impacts the consistency and use of the vapors in the carrier gas. The carrier gas composition can be changed, as with CSS, to comprise N2, Ar, He, and O2. As-deposited VTD films are like CSS films, with closely arbitrary orientation and normal grain size distribution (McCandless & Sites, 2003).

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Devices and Applications

S. Guha, ... B. Yan, in Comprehensive Semiconductor Science and Technology, 2011

6.08.4.5.1 High-rate a-Si:H and a-SiGe:H solar cells

Several deposition methods have been used for high-rate depositions of a-Si:H and a-SiGe:H solar cells, including the conventional rf glow discharge (Guha et al., 1992), vhf glow discharge (Curtins et al., 1987a, 1987b; Shah et al., 1992; Yue et al., 2007), microwave glow discharge (Guha et al., 1994), HWCVD (Mahan et al., 1991), thermal expansion PECVD (van de Sanden et al., 1998), and others.

RF-deposited materials shows a clear increase of defect density (Guha et al., 1992) and microvoid density (Williamson, 2003) as the deposition rate increases. As a result, the solar cell efficiency, especially the stable efficiency, decreases with the increase of the deposition rates. Figure 12 shows an example of initial and stable solar cell efficiency as a function of deposition rates. One can see that the light-degraded efficiency drops more steeply than the initial efficiency.

Figure 12. Initial and stable solar cell efficiency of a-Si:H solar cells deposited using conventional rf glow discharge at different rates. From Guha S, Yang J, Jones SJ, Chen Y, and Williamson DL (1992) Effect of microvoids on initial and light-degraded efficiencies of hydrogenated amorphous silicon alloy solar cells. Applied Physics Letters 61: 1444–1446.

In terms of the mechanisms of deterioration of material quality with deposition rates, three major effects have been identified. First, the high-rate deposited materials usually contain high dihydride and microvoid densities, which are responsible for the enhanced LID. The dihydrides are mainly caused by the SiH2* and SiH* species in the plasma, which are generated by high-energy electrons. Because the density of high-energy electrons increases with the increase of RF power, an effective and commonly used way to increase the deposition rates, the high-rate RF-deposited a-Si:H and a-SiGe:H solar cells show poorer efficiency than that for the low-rate deposited solar cells. Second, the ion energy increases with the increase of RF power, and high-energy ion bombardment is also believed to cause material degradation. Finally, during high-rate deposition, the impinging species do not have sufficient time on the growth surface to move and find favorable locations with low energy to form compact and low-defect structures. This results in poorer material.

The vhf glow discharge showed significant advantages to obtain high material quality at high deposition rates (Curtins et al., 1987a, 1987b; Shah et al., 1992; Yue et al., 2007). The deposition rate increases with the excitation frequency. In principle, one can obtain high-rate deposition with increased frequency without increasing the excitation power density. In this case, the electron temperature does not increase and the densities of SiH2* and SiH* in the plasma are not increased dramatically with deposition rates. Moreover, the high frequency reduces the ion energy and increases the ion flux intensity. The lower-energy and high-flux ion bombardments do not damage the growth surface but enhance the surface mobility of impinging species. The combination of these effects results in a much lower sensitivity of a-Si:H and a-SiGe:H properties on the deposition rates. Figure 13 shows the deposition rate dependence of initial and stable efficiency, as well as the ratio of LID, in a-Si:H solar cells deposited with a modified vhf glow discharge (Yue et al., 2007). One can see that both the initial and the stable efficiencies do not show a clear correlation to the deposition rate up to ∼14 Å s−1. The a-SiGe:H materials and solar cells deposited with vhf at high rates also show superior properties over the RF-deposited high-rate materials. With high-rate a-Si:H and a-SiGe:H deposited materials for triple-junction cells, initial and stable cell efficiencies of 11.3% and 10.1% have been achieved using vhf-deposited intrinsic layers, where the deposition time was less than 50% of the current corresponding deposition time in the production lines.

Figure 13. Deposition rate dependence of initial and stable efficiencies as well as the light-induced degradation of a-Si:H solar cells deposited with a modified vhf glow discharge. From Yue G, Yan B, Yang J, and Guha S (2007) High rate deposition of amorphous silicon based solar cells using modified very high frequency glow discharge. Materials Research Society Symposium Proceedings 989: 359–364.

Microwave glow discharge has also been demonstrated as an effective method for high-rate depositions of a-Si:H and a-SiGe:H solar cells. Guha et al. (1994) showed that high-efficiency a-Si:H and a-SiGe:H cells can be deposited at deposition rates up to 100 Å s−1. However, although the initial cell efficiency was reasonably good, the LID was very high. SAXS measurement showed that the microvoid density was high in the microwave-deposited materials (Guha et al., 1994; Williamson, 2003).

HWCVD is another method to demonstrate high deposition rate of a-Si:H and a-SiGe:H materials. The deposition rate can be increased to 100 Å s−1. However, reasonable cell efficiency was only achieved with a deposition rate around 10 Å s−1 (Mahan et al., 1998). In addition, some large-area issues and wire lifetime issues have delayed the use of HWCVD in large-scale solar panel production. In recent years, the thermal expansion glow discharge showed some promising results on high-rate a-Si:H solar cells. However, no high-efficiency solar cells have been demonstrated with this method (van de Sanden et al., 1998).

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Welding and Soldering

W Lucas PhD, DSc, CEng, FIM, FWeldl, EWE, S Westgate BSc(Hons), in Electrical Engineer's Reference Book (Sixteenth Edition), 2003

10.1.5.5 Series arc welding—single power source

As series arc welding is used for high deposition rate welding applications, specialised techniques have been designed specifically to increase the deposition rate and the welding speed. As the maximum welding current on a single wire is limited by a deterioration in weld quality through excessive arc forces, the techniques are based on the use of multi-arcs.

The simplest arrangement is two wires connected to the same power source to give d.c. electrode-positive and d.c. electrode-negative arcs or, alternatively, two a.c. arcs (Figure 10.31). When used in tandem (d.c. electrode-positive leading), a substantial increase in welding speed, typically 1.5 times the single-wire process, can be achieved with no deterioration in the weld-bead shape.

Figure 10.31. Series arc welding

The two arcs operate into a single weld pool and, because of this close proximity, there is significant arc interaction. With d.c., the arcs will diverge, i.e. be repelled by each other, whilst two a.c. arcs will be largely unaffected.

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Growth, properties, and applications of β-Ga2O3 nanostructures

Mukesh Kumar, ... R. Singh, in Gallium Oxide, 2019

5.2.1 β-Ga2O3 nanostructures using the CVD technique

The CVD method is attractive due to a high deposition rate, capable of producing pure materials, reproducibility of synthesis and ability to control the morphology of nanostructures by controlling process parameters, and capability of producing materials on an industrial scale. Many groups have studied CVD growth of β-Ga2O3 nanostructures including nanowires, nanosheets, and nanobelts [44–64]. The morphology of β-Ga2O3 nanostructures depends on various experimental parameters in CVD such as precursors, growth temperature, growth duration, separation distance between metal source and substrate (with and without catalyst nanoparticles), type of catalyst nanoparticles, and source gases flow rates. For example, metallic Ga and oxygen as source materials have been widely used to grow β-Ga2O3 nanostructures [53, 62]. Different precursors like a mixture of Ga and Ga2O3 have also been adopted to grow β-Ga2O3 nanostructures with and without catalysts [47, 48, 65]. Au nanoparticles are commonly used as a catalyst to grow β-Ga2O3 nanostructures. Other catalysts such as Ni [65], Fe [60], and spin-coated Ga2O3 films as a catalyst [58] have also been reported. Commonly used substrates for β-Ga2O3 nanostructures include silicon [44, 46] and alumina [45]. β-Ga2O3 nanostructures such as nanoribbons and nanorods have been shown by Auer et al. [53] and suggested that the vapor-solid (VS) mechanism is responsible for the obtained nanostructures. Chang et al. [5] have grown β-Ga2O3 nanowires using Ga vapors and H2O as precursors. The reaction temperature was kept in the range of 700–950°C. Grown nanowires exhibited a VLS growth mechanism and had diameters in the order of a few tens of nanometers with lengths of a few microns. β-Ga2O3 nanostructures such as nanosheets, nanowires, and nanobelts have been observed in self-catalytic nanostructure growth using the CVD technique [58], where spin-coated Ga2O3 films were used as a catalyst to initiate the growth of β-Ga2O3 nanostructures. The resulting rectangular nanosheets of β-Ga2O3 have widths and lengths in the range of 0.5–1.2 and 5–15 μm, respectively. On the other hand, nanowires grown by the self-catalytic method have diameters and lengths of about 50–150 nm and a few tens of micrometers, respectively. In various kinds of nanostructures, nanowires are one of the more promising nanostructures of β-Ga2O3, and show potential for nanoscale device applications.

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An introduction to welding processes

John Norrish, in Advanced Welding Processes, 2006

Submerged arc welding (SAW) {12}

Submerged arc welding is a consumable electrode arc welding process in which the arc is shielded by a molten slag and the arc atmosphere is generated by decomposition of certain slag constituents (Fig. 1.9). The filler material is a continuously fed wire and very high melting and deposition rates are achieved by using high currents (e.g. 1000 A) with relatively small-diameter wires (e.g. 4 mm).

1.9. Submerged arc welding.

The significant features of the process are:

high deposition rates;

automatic operation;

no visible arc radiation;

flexible range of flux/wire combinations;

difficult to use positionally;

normally used for thicknesses above 6 mm.

The main applications of submerged arc welding are on thick section plain carbon and low-alloy steels and it has been used on power generation plant, nuclear containment, heavy structural steelwork, offshore structures and shipbuilding. The process is also used for high-speed welding of simple geometric seams in thinner sections, for example in the fabrication of pressure containers for liquefied petroleum gas. Like shielded metal arc welding, with suitable wire/flux combinations, the process may also be used for surfacing.

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Radiation Effects in Structural and Functional Materials for Fission and Fusion Reactors

G. Federici, ... V. Barabash, in Comprehensive Nuclear Materials, 2012

4.19.4.3 Considerations on Plasma-Sprayed Beryllium

In the past, plasma spraying was considered as a high deposition rate coating method, which could offer the potential for in situ repair of eroded or damaged Be surfaces. Development work was launched during the early phase of the ITER R&D Program in the mid-1990s.136 In the plasma spray process, a powder of the material to be deposited is fed into a small arc-driven plasma jet, and the resulting molten droplets are sprayed onto the target surface. Upon impact, the droplets flow out and quickly solidify to form the coating. With recent process improvements, high quality beryllium coatings ranging up to more than 1 cm in thickness have been successfully produced. Beryllium deposition rates up to 450 g h−1 have been demonstrated with 98% of the theoretical density in the as-deposited material. Several papers on the subject have been published.136–138 A summary of the main achievements can be found in Table 4.

Table 4. Main achievements of ITER-relevant plasma-sprayed technology (summary of best results, not always achieved together)

ParameterValue/resultsComments
Residual porosity (%)˜2Could be more than 5%
Thermal conductivity (W mK−1)Up to 160 at RTDepends on temperature of substrate, maximum achieved at T  600–800 °C with addition of H
Bond strength (MPa)100–200Reasonable
Substrate temperature (°C)>450Very important for good strength, adhesion, and thermal conductivity. Keep in mind that CuCrZr temperature should not be higher than 500 °C for several hours due to overageing of CuCrZr
Substrate preparationNegative transfer arcNeeded, but very difficult to do in situ
Deposition rate (kg h−1)4.5Reasonable
Thickness (mm)>10Reasonable
Deposition efficiency (%)>90It means that more than 10% of powder will be lost in chamber
Thermal fatigue (MW m−2/number of cycles)5/680; 1/3000For first-wall conditions tested

However, based on the results available, the initial idea of using plasma-sprayed beryllium for in situ (in tokamak) repair was abandoned for several reasons. First was the complexity of the process and requirements to control a large number of parameters, which affect the quality of the plasma sprayed coatings. Some of the most important parameters include plasma spray parameters such as (1) power, gas composition, gas flow-rate, nozzle geometry, feed, and spray distance; (2) characteristics of the feedstock materials, namely, particle size distribution, morphology, and flow characteristics; (3) deposit formation dynamics, that is, wetting and spreading behavior, cooling and solidification rates, heat transfer coefficient, and degree of undercooling; (4) substrate conditions, where parameters such as roughness, temperature and thermal conductivity, and cleanliness play a strong role; (5) microstructure and properties of the deposit, namely, splat characteristics, grain morphology and texture, porosity, phase distribution, adhesion/cohesion, and physical and mechanical properties; and (6) process control, that is, particle velocity, gas velocity, particle and gas temperatures, and particle trajectories. Second, plasma-sprayed beryllium needs (1) inert gas pressure, (2) reclamation of the oversprayed powder (more than 10%), and (3) strict control of the substrate temperature. The higher the temperature the higher the quality of the plasma-sprayed coating, but unfortunately, an easy and reliable method to heat the first wall to allow in situ deposition was not found. Finally, tools to reliably measure the quality of the coating and its thickness are not available today and a strict control of the coating parameters is difficult to achieve.

Thus, it was concluded that plasma-sprayed beryllium for in situ repair is too speculative for ITER without further significant developments. Nevertheless, this method still remains attractive and could be used for refurbishment of damaged components in hot cell, albeit it may be cheaper to replace a damaged component with a new one.

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Film Structure

Milton Ohring, in Materials Science of Thin Films (Second Edition), 2002

9.6.5 AMORPHOUS METAL ALLOY SYSTEMS

Even though thin-film deposition techniques can independently facilitate high deposition rates and low substrate temperatures, not all metal alloys can be amorphized. Necessary, but not sufficient, for this purpose are atoms of different size, which alloys obviously provide. In addition to atomic size and achieving the required quench rates, alloy compositions are critical. Most of the presently known glass-forming binary alloys fall into one of four categories (Ref. 57):

1.

Transition metals and 10–30 at.% semimetals

2.

Noble metals (Au, Pd, Cu) and semimetals

3.

Early transition metals (Zr, Nb, Ta, Ti) and late transition metals (Fe, Ni, Co, Pd)

4.

Alloys consisting of IIA metals (Mg, Ca, Be)

In common, many of the actual glass compositions correspond to where deep or low-temperature eutectics are found on the phase diagram. Two different amorphous alloy–film system types will be considered. They shed different lights on aspects of these interesting materials and their properties.

9.6.5.1 Au–Co Amorphous Films

We first consider amorphous Co–30 Au films since they were among the first to be well characterized structurally and through electrical resistivity (ρ) measurements (Ref. 58). The Au– Co equilibrium phase diagram reveals simple eutectic behavior with critical eutectic composition and temperature values of 23.5 at.% Co and 996°C, respectively; the alloy composition selected thus bears little relation to these eutectic features. Amorphous films were evaporated from independently heated Co and Au sources onto substrates maintained at 80 K. Dark-field electron microscope images and corresponding diffraction patterns are shown side by side in Fig. 9-21. The as-deposited film is rather featureless with a smooth topography, while the broad-halo diffraction patterns cannot be easily and uniquely assigned to known lattice spacings of crystalline alloy phases in this system. Both pieces of evidence point to the existence of an amorphous phase. The question of whether so-called amorphous films are in reality microcrystalline 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.

Figure 9-21. Electron micrographs and diffraction patterns of Co–30 at.% Au. (Top) As deposited at 80 K, warmed to 300 K (amorphous); (middle) film warmed to 470 K (single-phase FCC structure); (bottom) film heated to 650 K (two-phase equilibrium).

(From Ref. 58.)

Resistivity changes accompanying the heating of Co–38 at.% Au (an alloy similar to Co–30 at.% Au) revealed a two-step transformation as shown in Fig. 9-22. 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. 9-21. Prior to this, the high resistivity of the amorphous films is due to enhanced electron scattering by the disordered solid solution. Crystallization to the FCC structure reduces the resistivity, and phase separation, further still.

Figure 9-22. Resistivity of a Co38 at.% Au film as a function of annealing temperature. Reversible values of /dT in various structural states of the film are shown together with changes in ρ during phase transformations.

(From Ref. 58.)

Both the amorphous and metastable phases are stable over a limited temperature range in which the resistivity of each can be cycled reversibly. Once the two-phase structure appears, it of course can never revert back to less thermodynamically stable forms. This amorphous–3crystalline transformation apparently proceeds in a manner first 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 or single-phase metastable matrices. For example, the equilibrium phase diagram for Ag–Cu is that of a simple eutectic with relatively pure terminal phases of Ag and Cu which dissolve less than 0.4 at.% Cu and 0.1 at.% Ag, respectively, at room temperature. These limits can be extended to 35 at.% on both sides by vapor quenching the alloy vapor. Similar solubility increases have been observed in the Cu–Mg, Au–Co, Cu–Fe, Co–Cu, and Au–Si alloy systems.

9.6.5.2 Ni–Zr and Co–Zr Amorphous Films

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 (Ref. 59). Consider Fig. 9-23, which shows the result of annealing a bilayer couple consisting of pure polycrystalline Ni and Zr films to 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 as kinetic considerations dominate 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 compounds. However, the driving force for either transformation is similar. Unlike other amorphous films extensive interdiffusion can be tolerated in Ni–Zr without triggering crystallization.

Figure 9-23. Cross-sectional electron micrograph of an amorphous Ni–Zr alloy film formed by annealing a crystalline bilayer film of Ni and Zr at 300°C for 4 h.

(Courtesy of K. N. Tu, IBM, T. J. Watson Research Center. From Ref. 59.)

The very earliest stage of the solid-state amorphization reaction in the Co–Zr system was detected with atomic resolution employing field ion microscopy (Ref. 60). It appears that the large negative heat of mixing plus the decrease of interface energy at the incoherent Co/Zr interface provide the driving forces for the reaction. Initially a 2.5 nm amorphous region forms at the interface even at room temperature. The process starts with the supersaturation of large-angle Zr grain boundaries by Co. Beyond a critical Co concentration the Zr lattice is destabilized and the amorphous phase grows spontaneously. Evidently the increase in supersaturation is due to the presence of both immobile Zr atoms and defects. However, point defects and dislocations in Zr help to nucleate the amorphous phase because in their absence no amorphization occurs.

In Section 11.4.2.1 the kinetics of the amorphous-to-crystalline transformation in CoSi films will be described.

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TECHNOLOGICAL CHALLENGES IN WELDING FOR CANADIAN OFFSHORE ENVIRONMENTS

J.C. Price, ... N.F. Eaton, in Welding for Challenging Environments, 1986

2.5.2 For Shipping, Caissons and Structural Applications

For shipping, caissons and structural applications it is envisaged that the high deposition rate processes, developed to utilize consumables with improved toughness properties will be used but with more stringent toughness requirements than those for the North Sea. It has been shown that high heat input submerged arc processes achieve adequate toughness in the HAZ when welding the new generation of structural steels. This is particularly true for high efficiency tandem or triple wire submerged arc welding which has already enhanced the development of the Japanese shipbuilding industry. The long electrode distance SAW process has already been used to weld parts of the mobile Arctic caisson. However, as the submerged arc process is limited in positional applications it may be anticipated that the flux cored arc and gas metal arc processes will be utilized more than in North Sea work where shielded metal arc welding has been normally utilized. This will be possible because of the satisfactory weld quality and properties that have been achieved for these processes in the last five years. In particular multi-headed high current GMAW with tandem or triple wire has also been used in mobile Arctic caisson construction. FCAW has been successfully applied to North American shipbuilding, particularly the self-shielded type, for out of position weld seams, when submerged arc cannot be applied. Consequently, the ratio of semi-automatic to manual process use would be expected to increase due to the higher deposition rates and improved quality of FCAW and the gas metal arc processes.

Finally, the development of pulsed GMAW welding, particularly the frequency modulated type, as developed at the Welding Institute of Canada would appear to have great potential for structural as well as piping fabrication in Canadian offshore development. BY this method the spray transfer mode may be applied to positions such as the vertical where conventional GMAW using dip transfer was difficult to apply because of problems in controlling the weld pool and hence producing defective welds. Furthermore, thinner materials may be welded in the spray transfer mode with the pulsed GMAW process that was previously difficult due to burn through with the higher heat inputs of conventional GMAW. Finally, the weld bead of pulsed GMAW has less spatter and a smoother profile than conventional GMAW in the spray transfer mode.

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