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By H. D. Hodges, J. K. Godbey

In order to better understand the fracturing process, bottom-hole pressures were measured during a number of typical fracturing operations. A recently developed system was used that allows simultaneous surface recording of both the bottom-hole and wellhead pressures on the same chart. The results from six fracruring treatments are summarized on the basis of the pressure data obtained. Al-though no complete analysis is attempted, the value of accurate pressure measurements is emphasized. Important characteristics of the bottom-hole pressure record do not appear at the wellhead because of the damping effect of the fluid-filled column. In four of the six treatments described, the formations apparently fractured during the initial surge of pressure with only crude oil in the well. The properties of the fluids used during the treatments are given and the fluid friction losses are obtained directly from the pressure records. This technique is also shown to be adequate for determining when various fluids, used during the process, enter the formation. INTRODUCTION Hydraulic fracturing for the purpose of increasing well productivity is now accepted in many areas as a regular completion and workover practice. Numerous articles have appeared in the literature discussing the various techniques and theories of hydraulic fracturing'. In general, three basic types of formation fractures are recognized today. These are the horizontal fracture, the vertical fracture, and fractures along natural planes of weakness in the formation'. Any one or all three of these fracture types may be present in a fracturing operation. However, with only the wellhead pressure record as a guide, it is difficult at best to determine if the formation actually fractured, and is almost impossible to determine the type of fracture induced. These difficulties arise in part because the wellhead pressure record, especially when fracturing through tubing, does not accurately reflect the pressure variations occurring at the formation. Several factors contribute to this effect and preclude the possibility of using the wellhead pressure as a basis for accurately calculating the bottom-hole pressure. These factors are: 1. The compressibilities of the fluids which damp the pressure variations. 2. The changes in the densities of the fluids or apparent densities of the sand-laden fluids. 3. The flowing friction of the various fluids and mixtures, which is dependent on the flow rates and the condition of the tubing, casing, or wellbore. 4. The non-Newtonian characteristics of a sand-oil mixture and its dependence upon the fluid properties, the concentration of sand, and the mesh size used. 5. The unknown and variable temperatures throughout the fluid column. Because of these reasons it was determined that in order to obtain a more accurate knowledge of the nature of fracturing, the bottom-hole pressure must be measured along with the pressure at the surface during a fracturing treatment. Even with accurate pressure data, a reliable estimate of the nature of fracturing is still dependent upon knowledge of the tectonic conditions. However, the hydraulic pressure on the formation is basic to any approach to a complete analysis. In order to accomplish this objective a system was developed to record the wellhead and bottom-hole pressures simultaneously at the surface. By recording both pressures on a dual pen strip-chart recorder, it was possible to greatly expand the time scale so that rapid pressure variations would be faithfully recorded. By such simultaneous recording, time discrepancies inherent in separate records are eliminated, thus overcoming one of the most difficult problems associated with bottom-hole recording systems. This paper illustrates the results obtained by using this system during six typical fracturing operations. All of these tests were taken in wells that were treated through tubing. By a direct comparison of the wellhead and bottom-hole pressures, the importance of obtaining complete pressure information during a fracturing treatment is emphasized. THE INSTRUMENTATION AND PROCEDURES The bottom-hole pressure measuring instrument consisted of a pressure-sensing element, a telemetering section, and a lead-filled weight or sinker bar. The pressure-sensing element used was an isoelastic Amerada pressure-gauge element. By using an isoelastic element, no temperature compensation was necessary in the tests described, since the temperature was believed to be well below the maximum temperature limit of 270°F. The rotary output shaft of this helical Bourdon tube element was coupled to a precision miniature potentiometer. The rotation of the pressure-gauge shaft thus changed the resistance presented by the potentiometer

By J. Alfred Berger, O. M. Katz

Selected samples of hydrided Zr-Hf alloys were rapidly quenched to voom temperature and exrtrnined metallographically, by X-ray diffraction, and through micro hardness studies to confirm high-temperutuve data Confirming experiments sllowed that there were five phases in this Lernary system: 1) hextrgonal with lattice parameters similar to that of the initia1 Zr-Hf alloy but slightly enlarged due to dissolved hydrogen; 2) fee with properties of a brittle, intermediate, hydride compound; 3) fct with c/a crvoltnd 1.07 and which appeared as a neetilelike precipitale; 4) hexagonal, designated ?, with c/a ratio of 2.37; and 5) orthorhombic, designated X, with a = 4.67, b = 4.49, and c = 5.093 and whose tnicro-st?ruct~ival nppetrl-nnce depcncled o/i, heat lvecrt~r~ent. The tetragonrrl phase never crppeal-erl witkorct the cubic hydricle. Abpecrrtrnce of 0 and A also tlependet on the hafnium content of the zirconium. A previous paper' on the Zr-Hf-H system described the thermochemical data obtained with a high-vacuum, high-sensitivity mirrogravimetric apparatus. This data presented a fairly complete picture of the phase relationships at elevated temperatures. However, it could not establish the actual crystal structures, lattice parameters, or metallographic disposition of the hydride phases. The present complementary study utilizes X-ray powder patterns along with light and electron microscopy to characterize completely the five hydrided phases found in Zr-Hf-H alloys quenched to room temperature. Crystallographic features of the zr-Hf,2,4 zr-H,5-7 and Hf-H8 systems have been summarized in Table I. Designations of a, ß, and ? were retained in the Zr-Hf-H system for the phase regions through which the pressure-composition isotherms always sloped. However, it was not firmly agreed that these were single-phase regions.' In fact, the region designated y always contained a cubic as well as a tetragonal phase after quenching to -196°C. MATERIALS Preparation of the high-purity Zr-Hf alloys has been described.' The four zirconium alloys which were hydrided contained 37 wt pct Hf (23 at. pct), 51 wt pct Hf (37 at. pct), 73 wt pct Hf (58 at. pct), and 91 wt pct Hf (82 at. pct), respectively. These were designated B-2, B-4, B-6, and B-8. Photomicrographs of the initial alloys showed the material to be quite clean as would be expected from the precautions exercised in producing them. However, there were a number of annealing twins but no other subgrain structure. In addition to the four original alloys, fifteen hydrided samples were observed at room temperature. Hydrogen compositions are given at the top of Tables I1 to V. APPARATUS The phases present at elevated temperatures were studied by quenching hydrided samples to room temperature by two different methods, both under vacuum: 1) fast cooling of the sample tubes of the microgravimetric apparatus1'9 with flowing air and 2) rapid quenching into liquid nitrogen. The cooling rate for 1) was 750° to 250°C in 30 sec. Since the microbalance chamber was not designed to permit very rapid cooling of a hydride sample, all liquid-nitrogen quenching was done in an auxiliary experiment. The auxiliary quenching apparatus consisted of a small-bore, high-temperature furnace, a sealed SiO2 tube containing the sample, and a dewar quenching flask filled with liquid nitrogen. The hydrided sample, previously quenched in the microgravimetric reaction chamber, was placed in a platinum boat in a vacuum-degassed SiO2 tube. A zirconium wire getter and degassed SiO2 rod, to reduce the internal volume, were also in the tube. After sealing the tube under vacuum the zirconium getter was heated to absorb the last traces of gas. Only the sample was heated at the reaction temperature for the desired length of time, and then the tube dropped through the opposite end of the furnace into the dewar. A quenching rate of 200" to 400° C per sec was estimated. Analyses of samples after the auxiliary experiment also showed practically no increase in oxygen or nitrogen content from heating in the SiO2 tube. All of the samples were examined at room temperature by the X-ray powder method. The majority of the powder patterns were obtained with double nickel-filtered CuKa radiation after 8- and 16-hr exposures in an 11.48-cm-diam camera. Cobalt and chromium radiation were also used to spread out the high d value end of the Pattern. Such patterns readily identified the minor phases. NO oxide or nitride lines were found. Where sharp back-reflection lines existed it was possible to reduce the

Jan 1, 1965

By H. F. Bishop, F. A. Brandt, W. S. Pellini

The solidification mechanism of experimental steel ingots (7x7x20 in.) was studied by thermal analysis. It was determined that solidification proceeds in wave-like fashion at rates which are determined by the carbon level, superheat, and mold thickness. The thermal cycles of the mold walls were related to the course of solidification. ESPITE marked advances in the field of solid state transformation, metallurgical research has contributed comparatively little exact quantitative data on the mechanism of solidification of metals. There is, therefore, a great need for such data in the various metallurgical industries. The mechanics of solidification of ingots have been investigated in the past primarily by studies of the rate of skin formation as indicated by bleeding or "pour out" tests. The "pour out" method, however, is a tool which gives only approximate information. In the case of alloys with wide solidification ranges, such as irons and certain nonferrous alloys, the method will not work at all; in the case of alloys of intermediate solidification ranges, such as commercial steels, the information may be misleading. Thus, the general adoption of this method has resulted in divergent conclusions regarding the solidification process. Chipman and Fondersmith1 by means of bleeding tests have shown that the linear growth of a solidifying ingot wall follows a parabola of the general form, D = K C, with the start of solidification delayed until superheat is exhausted, as indicated by the constant C. These tests were carried only to a wall thickness of about 5 in. using an ingot of approximately 17x39 in. in cross-section; hence the latter stages of solidification were not studied. Matuschka2-3 indicated that linear solidification of ingots is rapid at first, then slow, but toward the end of solidification the rate becomes extremely rapid again. Spretnak's4 bleeding studies indicated that, wall growth is expressed more rigorously by two parabolas, and that their point of intersection corresponds to a change of solidification mode from columnar to equiaxed. Spretnak also showed that the K values of the first parabola are always the same regardless of superheat. Nelson bled ingots of square cross-section and found that linear wall growth is initially rapid but decreases continually until the end of solidification. He also concluded that rate of solidification in ingots of square cross-section increases 2.15 pct for every 10 pct increase in cross-sectional area of the mold. The mold ratios considered (ratio of cross-sectional area of the mold to cross-sectional area of the ingot) were all less than 2 to 1. The subject of solidification has also been treated mathematically in many cases, but because of the lack of accurate thermal constants and the simplifying assumptions required, as their authors generally acknowledge, they represent only approaches to the actual conditions of ingot solidification. A third method of studying solidification is the electrical analogue method promulgated by Pasch-kis6-7 and by Jackson and coworkers.8 This method treats solidification as a heat transfer problem with the solidification cycle synthesized on an electrical circuit. Paschkis in his treatment of solidification considered the fact, which was generally ignored, that solidification of steel is not simply the growth of a plane solid wall but a more complex process occurring over a temperature range as indicated by the constitution diagram. Undoubtedly, the anomalous results obtained by bleeding tests arise from the inability to measure quantitatively this mushy condition. The shape of Paschkis' solidification curves are more nearly in accord with those of Matuschka, in that they indicate rapid linear solidification at the beginning and end of solidification with intermediate solidification occurring at a slower rate. Paschkis indicates a definite lengthening of solidification time with increasing superheat. Thermal analysis is a direct method providing exact information for all types of metals regardless of solidification range and was thus adopted in the present program to follow the entire course of solidification from the surface to the centerline of the ingots. The method has the added advantage of being adaptable to following the thermal cycle of the ingot mold during the course of solidification. Test Methods The ingots studied were of square cross-section, 20 in. long, tapered from 71/4 in. at the top to 63/4 in. at the bottom, and fed with a hot top 7 in. in diam and 12 in. high. The molds were uniform in wall

Jan 1, 1953

By A. U. Seybolt

In part 111' of this series it was shown that during the selective oxidation of a 3 1/4 pct Si-Fe alloy in damp hydrogen, only silica, (observed at room temperature) as low cristobalite or low tridy-mite or both, was formed as an oxidation product. In some in- „ stances where the film was fairly thin (probably well under 100A) there was some suggestion of an amorphous form of SiO2. The present investigation of oxidation rate showed that the selective oxidation of silicon-iron can be rather complicated, and apparently impossible to rationalize in an unequivocal manner. In some temperature regions, notably near 800" and 1000°C, the data seem to obey the familiar parabolic rate law. However, at intermediate temperatures complications were noted, some of which are possibly due to the order-disorder reaction in the silicon-iron solid solution. IN an earlier report' it was shown that during the oxidation of 3 1/4 pct Si-Fe alloys in H2O-H2 atmospheres only silica films were formed in the temperature range from 400° to 1000°C in hydrogen nearly saturated with water at room temperatures, or at dew points as low as -45°C. In the work to be reported here, some observations are made on the rate of oxide film formation. As in the earlier investigation, electron diffraction patterns generally showed either low tridymite or low cristobalite or both, except for some very thin films. These sometimes showed diffuse rings, presumably due to a very small crystallite size, or in a few cases, diffuse bands probably caused by an amorphous film. EXPERIMENTAL PROCEDURE Vacuum-melted silicon iron made of high-purity materials was rolled into strips 0.014 in. thick, and cut into samples 1/2 in. wide by 1 in. long. Chemical analysis showed 3.2 pct Si and 0.002 pct C. All samples were surface abraded with 600-grit paper, were solvent cleaned, and then placed in an paper,apparatus containing a "Gulbransen type"2 micro-balance. Here the gain in weight of the samples of about 5 sq cm area could be followed as a function of time during the oxidation caused by the water in atmospheres of various controlled water-hydrogen ratios. The water-hydrogen ratios can most easily be described as varying from a dew point of 0°C (PH2O-p^2 = 6.2 x 10-3 , to K (P j -40°C (PH2O/PH^= 1.3 X 10-* Most of the experiments were conducted at the 0°C dew-point atmosphere because drier atmospheres caused so little gain in weight that the accuracy of measurement was poor. Because of this, only the data obtained at PH2O,/P,,,= 6.2 x X3 will be reported. The temperature range extended from 800" to 1000°C; and most of the oxidation runs lasted for about 24 hr. The reproducibility of any reading was about ± 1 ?, but the sensitivity of the balance was about 0.2 ?. The atmosphere, flowing at 200 cm per-min, was preheated to the furnace temperature before contacting the specimen. While the gas flow caused a measurable lift on the sample, it was ordinarily sufficiently constant so that it was not an appreciable source of error. X-ray and electron diffraction checks of the samples before and after oxidation showed no evidence of preferred orientation, either on the metal samples or on the silica films formed. EXPERIMENTAL RESULTS The data obtained are summarized in Table I, and some are given in detail in Figs. 1 to 4. In the fourth column of Table I, kp refers to the parabolic rate constant in the expression (?/cm2)2 = kpt + c [1] where ? = micrograms gain in weight kp = parabolic rate constant in units r2 /cm4 t = time in minutes c = constant It will be noted that in many cases no value for kp is given; this is because in these instances the data did not obey the parabolic rate law. The silica film thicknesses given in the last columns are values calculated from the weight gain, an average tridy-mite-crystobalite density, and by assuming a perfectly plane surface. Fig. 1 shows the data plotted in the form of Eq. [I], hence a linear plot indicates parabolic behavior. It has been frequently observed in the literature that

Jan 1, 1960

By C. S. Matthews, H. C. Lefkovits

Drilling progress is often delayed by sticking of the drill string. The development of preventive and remedial methods has been hampered by incomplete understanding of the sticking mechanism. A recent lahorntory investigation hns indicated that one type of sticking may be attributed to the difference in pressure between the borehole and formation. This paper shows, by means of soil mechanics, that the primary cause for differential pressure sticking is cessation of pipe movement, whereas diflerential pressre and stanrtding time determine the severity of the sticking. The analysis stresses the importance of using low-weight muds with low solids content and low water loss to alleviate diflerential pressure sticking and describes why packed hole drilling, long strings of drill collars, and a large deviation from the vertical are conducive to sticking. Finally, preventrve and remedial methods ore evaluated, and a theory is presented on the release of stuck pipe by spotting oil. INTRODUCTION Since drilling with long strings of oversize drill collars has become standard practice in many areas, the incidence and severity of the stuck pipe problem has increased. It has been noticed that in the majority of these cases the sticking could not possibly be attributed to key seating or caving of shales. It appeared that, due to the differential pressure between the mud column and the formation fluid, the collars were pressed into the wall and so became "wall stuck". Points to note about differential pressure sticking are: (1) sticking is restricted to the drill collars, (2) the collars become stuck opposite a permeable formation, (3) the sticking occurs after an interruption of pipe movement, (4) circulation, if interrupted, can be restarted after the sticking is noticed, and (5) no large amounts of cuttings are circulated out after restarting circulation. Helmick and Longleyl investigated pipe sticking by differential pressure in the laboratory and found an empirical relationship between the differential pressure, the sticking time and the required pull-out force. In this paper an explanation of the mechanism is given based on Terzaghi's theory of clay consolidation. A qualitative description is given in the following paragraphs while the derivation of fonnulas is given in Appendices. This paper is a first attempt to explain pressure differential sticking and many points will require additional theoretical and practical investigation before the problem can be fully understood. PRESSURE DIFFERENTIAL STICKING AS A CONSOLIDATION PROBLEM In any borehole, where the mud pressure is higher than that exerted by the formation fluids, a mud cake is formed opposite the permeable sections of the hole and a continuous flow of filtrate takes place from the mud, through the cake and into the formation. This radial flow pattern requires a certain distribution of the hydraulic and the effective (grain-to-grain) stresses inside the mud cake. Any quantitative or qualitative change in the external pressure conditions will produce a change in the flow pattern and, consequently, also in the internal stress distribution inside the cake. In view of the low permeability and the high compressibility of a clay mud cake, the adjustment of the internal stress distribution is slow and is accompanied by a change in volume. Time dependent stresses are thus created which gradually diminish as the new state of equilibrium between internal and external pressures is approached. Some 30 years ago, Terzaghi developed his "Theory of Consolidation" to account for the time-dependent stresses and settling of clay formations under the influence of external loads. He derived a differential equation by which the time-dependent hydraulic stress and the consolidation can be computed for any point inside the layer during the consolidation process. His theory is based on the assumption that the change in stress is solely due to a change in water content and it may only be applied to one-dimensional consolidation phenomena. Other investiga-tors5,10 have expanded his theory to include processes of more than one dimension. The difference between the external pressures on the mud cake before and after sticking is a qualitative one (isolation of part of the cake by the static contact with the drill collars after pipe movement has been stopped)', and the time-dependent stresses thus created may be investigated by means of Terzaghi's theory. By this analysis the changes in the nature of the contact surface between the drill collars and the mud cake during the sticking can be explained; and the friction force between the two may be computed as a function of the sticking time, the borehole dimensions and the mud cake characteristics.

By R. F. Lynch, J. D. Wood

The Al-1.0 pct Mg alloy 565 7 was studied using optical microscopy and electron microprobe X-ray analysis. Constituent particles were found to exist inter-dendritically in the as-cast material in a region of precipitate free a -aluminum. Five phases besides a fine precipitate and a-Al were identified in the cast structure: Fel3, Fe2Al7, Mg2Al3, CUMgAl2, and Cu2FeAl7. Thermal treatments conducted for 100 hr at 1180°, 1130°, 1080°, 1030°, 980°, and 880° F revealed a general dissolution and spheroidization of the in-terdendritic constituent network observed in the cast structure. The principal constituents present in the thermally treated structures were FeAl3 and Fe2Al7 with the relative amount of Fe2A17 to FeAl3 increasing with a decrease in the treatment temperature. The phases Present in the wrought structure were identical to those observed after the thermal treatments, with the constituent particles strung out in the direction of rolling. ALLOY 5657 is a nonheat-treatable commercial purity Al-1.0 pct Mg alloy utilized extensively because of its bright finishing characteristics. This investigation was conducted to determine the constituents present in 5657 alloy, and to study the effect of extended thermal treatments on morphology. Numerous studies have been carried out to establish the equilibrium diagrams for various aluminum systems,1-3 with phase identification based on X-ray analysis, morphology, and the etching response of relatively large particles. Phragmen4 conducted a study of the phases in aluminum eutectic systems and compiled a "corrected" table of etching responses, drawing on his work plus that of Schrader,5 Keller and Wilcox,6 and Mondolfo.7 A review of the original work of Keller and Wilcox, and Mondolfo, which was concerned with the constituents found in commercial alloys, reveals that in numerous cases their etching responses differ from those reported by Phragmen and from each other. These inconsistencies may occur because a specific constituent will react to a given etch in a varying manner depending upon its size, the elements dissolved in the phase, the other constituents surrounding a phase, and the solute content of the matrix. Work with a commercial orientation was conducted on alloys 2024 and 3003 by Sperry8,9 and on alloy 3003 by Barker,10 where the relationship between the phase diagram and the nonequilibrium structure of an alloy was examined. Backerud11 investigated the A1-Fe binary system and found that at high cooling rates the equilibrium eutectic reaction forming a-A1 and FeA13 is replaced by another lower temperature eutectic reaction forming a-A1 and metastable FeAl6, a constituent first identified by Hollingsworth et al.12 Most of the above mentioned studies were conducted on materials having a significantly greater alloy content than 5657 alloy, where the relatively small size and sparse distribution of second phase particles hinders the process of identifying constituents. EXPERIMENTAL PROCEDURE Material with a composition as given in Table I was examined in the as-direct chill cast, hot rolled and cold rolled conditions, and after thermal treatment of the cast structure. Thermal treatments were terminated by a water quench. Microscopic examination was conducted under various lighting conditions following the application of standard etchants as specified in Table 11. A semi-quantitative electron microprobe X-ray analysis was conducted for Al, Mg, Fe, Cu, Si, Zn, and Ti. RESULTS AND DISCUSSION Microstructure of the As-Cast Material. Particles of second phase material were found to exist inter-dendritically, principally in regions of precipitate free a-Al, as illustrated in Fig. 1. Adjacent to the ingot edge was a region of inverse segregation, resulting in an increased amount of second phase material containing large sized particles which aided in phase identification. Phase Identification. Cast Structure. Five phases besides a-Al and a fine precipitate were identified using optical microscopy and electron microprobe X-ray analysis, as presented in Tables II and III, respectively. FeA13 and Fe2A17 are often found with Fe2A17 forming a sheath around the core of FeAl3, resulting from an incomplete peritectoid reaction. These phases have a nearly identical appearance under white light, although they are easily differentiated under crossed polarizers, as characteristically illustrated in Figs. 2(a) and 2(b), respectively. Microprobe analysis con-

Jan 1, 1970

By J. M. Woodall

GuAs crystals which have been grown in quartz boats by the horizontal Bridgman method in the pvesence of Ga20 vapov have beetz found to have carrier and donor distributions which do not correspound to those expected from simple dopant seg-vegation during directional freezing; Instead, the carrier distribution is determined by the heat-trentnzent history of the crystal, while the donor distribution, zohiclz is principally due to silicon, is fixed by the pozuth rate, the geo)tlet.ry of the crystal growth vessel, and tlze initial Ga20 pressure. WHEN semiconducting materials are made into doped crystals by the normal freezing method,&apos; they usually exhibit doping variations along the growth axis. If a) there is no dopant diffusion in the solid, b) the dopant is distributed uniformly in the melt, and c) the distribution coefficient, k, does not vary with composition, then the doping variation along the growth axis is represented by the equation: where C is the doping level in the solid at a point where a fraction g of the original liquid has frozen, and Co is the mean concentration. When the dopant is either a singly ionized shallow donor or shallow acceptor, C also represents the carrier concentration. Even though this equation accurately describes the dcping profiles of a large number of normal freeze systems, there are several special systems for which Eq. [I] does not apply. One such system is the horizontal Bridgman method for preparing oxygen-grown GaAs crystals using quartz vessels. Several workers2"* have shown that GaAs crystals grown by the horizontal Bridglnan method using quartz vessels are generally contaminated with silicon in concentrations in excess of 5 < 10lG atoms per cu cm. This contamination is ascribed to a reaction: occurring at the walls of the crystal-growth vessel which liberates silicon into the melt and Ga2O vapor. It has been shown4 that this reaction, and, hence, the silicon contamination, can be suppressed by the addition of oxygen to the crystal-growth apparatus. It is the purpose of this paper to describe a special apparatus capable of yielding single crystals of GaAs grown in the presence of oxygen and to describe both the kinetics of silicon suppression in this system and the relationship between the carrier concentration profile and the silicon concentration profile. EXPERIMENTAL-CRYSTAL GROWTH A schematic diagram of the apparatus used in the oxygen addition experiments is shown in Fig. 1. The most important features of this apparatus are: a) the use of a sand blasted quartz boat, h) a quartz rod of length 1, with a hole of cross section A, that is placed near the boat to limit the free space volume V, over the melt during the growth, and c) the temperature gradient at and near the solid-liquid interface. Sand blasting the boat is necessary to prevent wetting of the melt. The quartz rod retards the diffusion of the GazO vapor away from the melt to the colder portions of the ampoule. A crystal is prepared by first loading a 5.5-in. boat, which has been cleaned in aqua regia, with 40 g of 99.9999 pct Ga along with 1 to 8 mg of Ga203 powder. GazO3 is a convenient source of oxygen since it reacts with gallium at the melt temperature to form Ga20 vapor, the species which apparently controls the suppression of silicon contamination. The loaded boat, the quartz rod, and the 99.9999 pct As are placed in the ampoule as shown in Fig. 1. Generally, GaAs seeds were not used since most of the unseeded growths resulted in monocrys-tals. The ampoule is evacuated to 10-5 Torr and sealed off. The GaAs melt is synthesized by placing the ampoule into a two-zone horizontal Bridgman furnace. The two zones are separated by several sandwiched layers of -in. Fibre-Fax board drilled with holes slightly larger than the OD of the ampoule. This causes a very large temperature gradient between the two zones, which is necessary for single-crystal growth. The two-zone furnace is mounted on a stand fitted with a roller bearing which allows uniform motion of the furnace in a horizontal direction. A uniform velocity of the stand is achieved by the use of a windlass device which winds up a wire attached to the stand. Movement of the solid-liquid interface is accomplished by fixing the position of the ampoule and moving the stand. Growth rates investigated in this experiment were between 0.4 and 1.2 in. per hr. The composition of the melt is fixed by maintaining an arsenic reservoir at 618°C. The stand is moved away

Jan 1, 1968

By Charles H. Grant

Some of the limiting factors relative to explosive crushing of rock and ways to overcome a few of these problems are presented. Relationships between borehole diameters, bench heights, and spacings, along with a review of the influence geometry has on energy as these are changed, are discussed. Efficiency in use of explosives and the decay of energy as it moves through rock and is absorbed and dissipated, is described, along with fragmentation as a function of spacings and energy zoning, etc. Communications are one of the major problems encountered. In an effort to provide a better understanding of the use of explosives, it is necessary to take a little different view of what explosives are, how to look at them as tools to fragment rock, and some of the problems encountered in doing so. First, take the explosive: although there are many factors involved, consider these as being reduced to only two — shock-strain imparted to the rock by the high early development of energy, and the gas effect which is a combination of heat, moles of gas formed, rate of formation of these gases which develop pressures, etc. First, consider shock energy by itself and assume there is no gas effect in the reaction. Fig. 1 illustrates a block or cube of rock, in the center of which is detonated an explosive charge which is 100% shock energy. Tensile slabbing would be seen on the surface and probably the cube of rock would generally hang together even though microcracks were formed. If the situation is reversed and an explosive whch has no shock energy and only gas effect (Fig. 2) is considered, the cube of rock would act as a pressure vessel and contain the pressure from the gas effect until it exceeded the rock-vessel strength; then the rock would break in a few large pieces. If these two kinds of energy are put together and the area of shock-strain around the explosive (Fig. 3) is considered, the two energies will be seen working together to furnish broken rock. The gas effect applies pressure to the microcracks formed from the shock energy to weaken the rock-pressure vessel and propagate these cracks to break the rock apart. It not only will be broken more finely, but will break apart at a lower pressure than the gaseffect case, since the shock energy has first weakened the rock vessel. Although tensile spalling from the shock-strain imparts momentum to the rock, the main source of displacement comes from the gas effect. The term "rock" is being used to mean any material to be blasted. These energies are absorbed by the rock in different ways. First, classify rock into two main categories: "elastic" and "plastic-acting." Elastic rock should be thought of as rock which can transmit a shock wave and is high in compressive strength, such as granite or quartzite. Since this elastic rock transmits a shock wave well, it makes good use of the shock energy from the explosive-forming cracks, etc., for the gas effect to work on. Plastic-acting rocks are rock masses which are relatively low in compressive strength and absorb shock energy at a much faster rate, thereby making poor use of the shock energy by not developing as extensive a cracked zone for the gas effect to work on. Rocks of this type are generally softer materials such as some limestones, sandstones, and porphyries. For the most part, the shockenergy part of the explosive reaction is wasted in plastic-acting rock, leaving most of the work to the gas effect. Since the ratio of gas effect to shock energy is different in different explosives, it is easy to understand why some explosives perform well in elastic rock and poorly in plastic-acting rock, and vice versa. Some of the most difficult blasting situations arise when mixtures of plastic-acting and elastic rock are encountered (Fig. 4). Fig. 4 shows an example of granite boulders cemented together with something like a decomposed quartz monzonite which is plastic-acting. The elastic granite boulders will transmit the shock-strain within itself, but when this shock tries to move through the monzonite to the next boulder, its intensity is absorbed by the monzonite and little shock-strain is placed on the adjoining boulder. In addition to this loss by absorbtion, shock reflection at the surface of the boulder will effect tensile spalling. The net effect is poor breakage of the boulders which do not have drillholes in them as they simply will be popped out with the muck. The same is true (Fig. 5) when layers and joints make

Jan 1, 1970

By O. F. Kammerer, W. E. Miller, D. H. Gurinsky, J. Sadofsky, J. R. Weeks

Zirconium and titanium inhibit solution mass transfer of steels by liquid bismuth, mercury, and lead. It is shown that in bismuth and mercury, these adsorb on the surface of the steels and subsequently react with nitrogen and possibly carbon from the steels to form inert, adherent surface layers of ZrN, TiN, or TiN + Tic. Data are presented which describe the condition under which thase deposits form. These inhibitors decrease the solution rate of iron into bismuth, and require a higher supersaturation for precipitation of iron from bismuth. USE of the low-melting heavy metals (bismuth, lead, mercury, and their alloys) as coolants has been limited because solution mass transfer of steels occurs in these liquids; i. e., iron dissolves in the hot sections of the heat transfer circuit and deposits in the colder sections. The rate of solution of iron and the temperature coefficient of solubility are sufficiently great to cause complete or partial stoppage by the deposition in the coldest section of a closed circuit in finite time, even though the actual solubilities are extremely low. In the development of the mercury vapor turbine by the General Electric Co., Nerad and his associates1 discovered that the addition of as little as 1 ppm Ti or Zr to magnesium-deoxidized mercury reduced the mass transfer of ferrous alloys by mercury to a negligible amount. Reid2 reported that titanium was detected chemically on the surface of steels contacted with this mercury alloy in amounts varying from 2.0 to 2.6 mg per sq in., the greatest amount being found in the hottest portion of the circuit. Reid stated that the titanium forms the intermetallic compound Fe2Ti by reaction with iron on the surface of the steels. This compound was presumed to be highly insoluble in mercury. More recently, El-gert and Egan3 have reported a greater than 100-fold reduction in the rate of mass transfer of a 5 pet Cr steel by liquid bismuth upon the addition of titanium (in excess of 50 ppm) and magnesium (350 ppm) in the liquid metal, during experiments performed in thermal convection loops* over the temperature differential 700° to 615° C. Also, Shep-ard and his associates&apos; have reported that the addition of titanium to liquid bismuth and Pb-Bi eutec-tic produced a marked decrease in the rates of solution of both iron and chromium from type 410 steel capsules under static conditions. This inhibiting effect increased with repeated reuse of the capsules. Tests performed in this laboratory under carefully controlled conditions have shown that the addition of zirconium and magnesium, or titanium and magnesium, to liquid bismuth or lead greatly reduces the rate of mass transfer of chromium alloy steels and carbon steels in thermal convection loops with a maximum temperature of 550°C.5-9 The present paper will review the data obtained to date at this laboratory on the behavior of iron and steels in contact with liquid bismuth alloys containing titanium or zirconium, and will attempt to explain the role of the above additives in reducing solution mass transfer. Reaction between the Zirconium or Titanium Dissolved In Liquid Bismuth and an Iron or Steel Surface Reaction between Zirconium Dissolved in Bismuth and the Surface of Pure Iron-—A small pure iron crucible (analyzed by the supplier to contain 0.8 ppm N was contacted with bismuth containing approximately 0.1 pet Mg and varying amounts of a radioactive zirconium tracer. The crucible was then inverted at the temperature of contact. The thin residual layer of adherent bismuth was dissolved in cold, concentrated nitric acid. The crucible surface and the solidified bismuth were then analyzed for radioactive zirconium. An analysis of the activity loss on the crucible surface and the weight loss of the crucible during the nitric acid treatment showed that the acid treatment removed the zirconium that had originally been dissolved in the adherent bismuth, but not any zirconium that may have reacted with the crucible surface. The crucible was then pickled in warm aqua regia to remove all surface activity, hydrogen-fired at 600°C, and recontacted with a new liquid alloy. The results of the experiments contacted 1 hr at 450°C show, Fig. 1, a Langmuir-type adsorption with an adsorption free energy of approximately 17 keal per g atom Zr.5 This deposit was estimated to contain 1 atom of zirconium for each 7 to 8 iron atoms on the crucible surface, assuming a surface roughness factor of the pickled crucibles to be five. Increasing the temperature to 520°C caused consi-

Jan 1, 1959

By H. I. Aaronson, C. Laird

Although the past results of X-ray experiments indicate that the broad faces of 0&apos; plates are coherent with their matrix, dislocations lying in arrays have frequently been observed at these boundaries by transmission electron microscopy. Critical experiments employing the latter technique have been carried out in order to determine the origin of these dislocations. It is concluded that 8&apos; plates are essentially coherent with the matrix at their broad faces throughout the aging temperature/time envelope studied. Virtually all of the dislocation arrays observed are deduced to have been formed by plastic deformation accompanying transformation. The proportion of dislocations arising from convexity of the plates is shown to be negligible by comparison with that from plastic deformation. At the higher aging temperatures, a[001] dislocations appeared in moderate numbers. These dislocations were traced directly, however, to the ledgewise dissolution of 0&apos; occasioned by the formation nearby of 0 crystals. On the other hand, since there is a parametric difference normal to the broad faces of the ?&apos; plates, mismatch dislocations do form at their edges. A previous conclusion that these dislocations have Burgers vectors of type a[001] was confirmed directly. The edges of 0&apos; plates were observed to develop octagonal shapes when growing, but circular shapes during dissolution. 1 HIS paper presents the results of an investigation of the interfacial structures of plates of the transitional phase, 8&apos;, formed in an A1-4 pct Cu alloy. In a companion paper, Part 11, the effects of these structures upon the migration kinetics of a:?f boundaries are reported. This work is pa.rt of a general program designed to establish the basis of precipitate morphology. The present authors in Al-Ag,1 and whitton2 previously in U-C alloys, have used transmission electron microscopy to examine directly the vander Merwe3-6 networks of dislocations anticipated7 to compensate the small amount of lattice misfit normally founda at the broad faces of Widmanstatten plates. Since the broad faces of 0&apos; plates are considered to be perfectly coherent with the corresponding habit planes in the a matrix,&apos; no dislocations should be present at these faces. Many reports have been published, however, giving evidence to the contrary.10-18 The primary objective of this investigation was therefore to ascertain the nature of these dislocation structures. An attempt to do this is described in the first three sections of this paper. Inspection of the matching of the a and 8 &apos; lattices at the orientations of the 0:0&apos; boundary corresponding to the edges of 0&apos; plates raises the possibility that these edges may be made up of rather closely spaced edge- type misfit dislocations oriented so as to be sessile with respect to the lengthening or shortening of the plates. Since this structure should severely inhibit migration of the plate edges (Ref. 7, Part II), a situation not originally anticipated,&apos; an experimental determination of the interfacial structure of the edges of 8&apos; plates was clearly in order, and is reported in Section III. Those aspects of the experimental procedure applicable to both Parts I and I1 are presented in the next section. Specific procedures applicable to individual aspects of each investigation, and also the relevant surveys of the literature, are then individually reported in the appropriate sections. I) GENERAL EXPERIMENTAL PROCEDURE The material used in both parts of these studies was the same as that of a previous investigation:" strips of A1-3.93 pct Cu, 0.009 in. thick, prepared as before, solution-annealed at 548°C for 6 hr, and quenched. Details of subsequent aging, and in some cases deformation treatments, are given in the Experimental Procedure sections of the individual parts of both papers. Specimens of the heat-treated strips were electro-thinned as beforeLg and examined in a Philips EM 200 microscope equipped with a goniometer stage. A commercial hot stage, of the grid-heater type and capable of * 30-deg tilt about one axis in the plane of the specimen, was also used for kinetic studies. The usual precaution of calibrating for the additional heat supplied by the electron beam was taken.19 A 16-mm cine cam-I era mounted outside the viewing window was frequently used to record the transformations. Conventional selected-area diffraction and dark-field viewing techniques were used to identify the precipitates in the foils. Normal bright-field images corresponding to two-beam diffracting conditions or dark-field images were employed to characterize the dislocations observed at the interfaces of the precipitates. The application of these techniques to the study of an interphase boundary, and the interpretation of the images,20&apos;21 has been fully described in a previous paper.&apos;

Jan 1, 1969

By R. C. Ku, A. J. McEvily, T. L. Johnston

The microplastic response of a series ofas-quenched Fe-Ni-C martensites has been measured at 77°K. At strains less than JO&apos;3 the flow stress is governed primarily by the transformation-induced dislocation structure of the martensite. Only at strains in excess of 10-3 is the influence of carbon manifested in the flow stress. At these macroscopic strains, typically 10-2, the solid-solution hardening is proportional to (wt pct C)1/3, and, in an alloy containing 0.39 wt pct C, amounts to 50 pct of the flow stress. THE technological significance of high-strength ferrous martensite has stimulated many investigations of its structure and properties. Although our knowledge of the characteristics of martensite has increased immensely, especially with the advent of high-resolution techniques, an understanding of the basic strengthening mechanism still remains elusive. The purpose of the present paper is to consider certain aspects of micro-plastic behavior of Fe-Ni-C martensite which we feel can help to resolve this important problem. Such alloys are particularly suitable for experimental investigation because their compositions can be adjusted to reduce the M, to a temperature low enough essentially to eliminate the diffusion of carbon in the freshly formed martensite.1 The mechanical properties in this condition are of interest inasmuch as they reflect a state that is free of the important but complicating influence of precipitation processes. In this virgin martensite the carbon is distributed as it was inherited from the parent austenite; i.e., it is present interstitially, and gives rise to tetragonality through strain-induced ordering.&apos; In order to determine the source of strength of such alloys, Winchell and Cohen1 investigated the low-temperature macroscopic stress-strain behavior of a series of virgin martensites of increasing carbon content but of common M, temperature (-35°C). They found that the flow stress increased rapidly with carbon content up to 0.4 wt pct; beyond this point the flow stress increased at a much slower rate. It was concluded that martensite is inherently strong. To account quantitatively for the strength of virgin or as- quenched martensite in terms of the role of carbon, Winchell and cohen3 suggested that the carbon atoms, trapped in their original positions by the diffusionless martensite transformation, interfere with dislocation motion according to a model akin to that of Mott and Nabarro. 4 In this treatment, individual carbon atoms are considered to constitute centers of elastic strain and thereby generate an average stress resisting the motion of dislocations throughout the lattice. The additional stress necessary to move dislocations, over and above that necessary for motion in a carbon-free martensite, is given by where L is an effective length of dislocation capable of motion. L was assumed to be limited to the spacing between the twins that are an essential structural element of Fe-Ni-C martensites. They assumtd the spacing to be invariant and of the order of 100A. However, recent work5 has shown that L is variable and can be in excess of 1000Å, so that the assignment of an appropriate value of L is not straightforward. In contrast to the above conclusion that there is an intrinsically high resistance to plastic flow, it has been suggested by Polakowski6 that freshly quenched martensite is in fact "soft" in the sense that dislocations are initially free to move upon application of stress. The high indentation hardness and macroscopic yield stress of ferrous martensites are then a consequence of rapid strain hardening that depends upon carbon in solution. Consistent with this point of view are the results of Beau lieu and Dubé who measured the rate of recovery of internal friction as a function of aging (tempering) temperature in a freshly quenched steel containing 0.90 wt pct C, 0.37 wt pct Mn, 0.1 wt pct Cr, and 0.07 wt pct Ni. The kinetics were clearly consistent with the idea that many dislocations are unpinned in the as-quenched state and that during aging they become progressively pinned by carbon at a rate controlled by carbon diffusion in the body-centered martensite lattice. In order to provide a basis upon which to distinguish between the "hard" and "soft" interpretations indicated above, we have made studies of the initial stages of plastic deformation in Fe-Ni-C martensites similar to those&apos;used by Winchell and Cohen. It will be shown that the results support the contention that dislocation segments in as-quenched material are indeed

Jan 1, 1967

By M. Khorouzan, L. Thomas

Designers of semiconductor devices have been strivi,ng to resolve problems associated with Au-A1 alloys in bonded in.tercomzeclions. One approach now being- used is that of waintaining a physical seyav-atioz between the two metals in bond areas. This is accolrzplished by alunzincnz-plating a bonding area on the tips oJ the kovar leads and using alcminurn wires to join the senzicondictor device to the leads. The portion of the kovar lead which is on the externul side of the sealed package is gold-plated to provide an oxide-free surface for soldering or welding. A discoloration condition originally thought to be sinilar to purple plague, occuving in the yluled uluninur bonding area after package sealing, has been investigated to determine its efiects ipm bond integrity. Electron-micro-probe analysis determined that no1 only gold, but lead, zinc, and silicon were also present in the discolored area. A series of samples conlaining&apos; conkrolled umonts of these inzpitrities weve prepared and subjected to a sil.zuluted sealing process. The investigations swcued that, of the contawiinants, only zinc toas detrinenlul to Lhe bond integily. The discoloration condition itself was found not to be detrimental to the bond integrity. DESIGNERS of semiconductor devices have been striving to resolve problems associated with Au-A1 alloys in bonded interconnections. One approach now being used is that of maintaining a physical separation between the two metals in bond areas. This is accomplished by aluminum plating a bonding area on the tips of the kovar leads and using aluminum wires to join the semiconductor device to the kovar leads. The portion of the kovar lead which is on the external side of the sealed package is gold-plated to provide an oxide-free surface for soldering or welding. Contamination as evidenced by discoloration of the aluminum-plated area was observed in a number of integrated circuits undergoing examination for defect characteristics which cause electrical failures.&apos; This paper contains the results of an investigation to determine the nature of this discoloration, its cause, and its effect upon the integrity of the interconnection bond. I) THE NATURE AND EXTENT OF ALUMINUM-BOND CONTAMINATION The initial hypothesis in the investigation was that the discoloration was caused by reaction of the aluminum film with some unknown contaminants during the sealing of the hermetically sealed integrated-circuit flat package. The package is a rectangular ceramic container sealed with glass which surrounds the kovar leads as well as joining the top to the bottom. The seal is made hermetic by heating and cooling the package to devitrify the glass. In the case of the packages under investigation, the hermetic sealing had been accomplished with dry air as internal atmosphere. The apparent effect of contaminations as observed by microscopic examination was the formation of surface oxides having variations in color encompassing the whole spectrum of visible light. The contamination appeared to be related to one of the more notorious examples of these colorations, the so called purple plague.&apos; In addition to purple plague, Fig. 1 shows the tarnish in the luster of the aluminized surface in the bond area which had been observed in many of the integrated circuits. To identify the contaminant in the bond area electron-probe microanalysis techniques were used.3 Fig. 2 shows the result of this analysis. The contaminants identified were gold, aluminum, zinc, lead, silicon, and cobalt. Fig. 2(a) is a back-scatter display of the area under study. The back-scattered electrons provide a general indication of the distribution of elements in the specimen surface. Elements with higher atomic number scatter more electrons back from the surface and are seen as light areas in the picture. The sample current, Fig. 2(b), is the amount of current conducted by the specimen as a result of electron-beam striking it and is an indication of element distribution. The Sample current is the reverse of back-scatter and complements it. Other pictures in Fig. 2 are produced by characteristic X-rays generated by the elements, allowing the isolation of the element of interest. The isolated element appears white and all other elements are dark. In this manner a comparative study provides a correlation between different surface areas and the elements which are in these areas. The area covered by the gold film, Fig. 2(c), shows that the boundary between the gold film and the kovar is not sharp as expected and that some sort of diffusion has taken place. Fig. 2(c) shows that some gold particles have been carried to the bond area and are in the proximity of the bonded wire in spite of the presence of a physical barrier in the form of the un-

Jan 1, 1967

By A. Ghosh, T. B. King

The kinetics of the discharge reaction: 20&apos;- (in silicate melt) = O,(g) + 4e- at a platinum anode in lithium silicate melts have been studied al 1350°C by galvanostatic methods. Plots of the sleady-state overpotential, q, as a function of the logarithm of the current density, i, showed injlections and were linear only at high current densities. The value of the overpotential was influenced by bubbling gas through the electrolyte. The ocer potential was also studied as a function of time. The rise and decay of overpotential were very slow processes. At low current densities transport is the likely rate-controlling process but at high current densities passivation of the electrode, Presumably by an oxide film on the surface, seems to be a contributory functor. IT is well-established that molten silicates behave as electrolytes&apos;5 and, except in a few cases,6 conduction is entirely ionic. Moreover, it is supposed that a possible, and perhaps predominant, mechanism for phase boundary reactions between metals and slags is similar to that in corrosion whereby anodic and cathodic processes occur at unrelated sites, the metal serving to conduct electrons.1&apos;8 Thus electrochemical studies of some slag-metal reactions would seem to be a useful way to diagnose the rate-controlling steps in the overall reaction. The electrochemical method is, in principle, a better diagnostic tool than the direct chemical method for the following reasons: 1) The partial electrochemical reactions, which are simpler than the overall reaction, may be studied individually. 2) The rate of reaction can be controlled at will and independently of the concentrations of reactants. 3) Fast reactions can be studied by relaxation methods.&apos; Esin and his coworkers5&apos;10"12 have pioneered such studies in silicates and have deviloped some ingenious techniques. Not all of their findings, however. can be accepted without a good deal of further work. In this investigation, the kinetics of the oxygen discharge reaction: 202- (in silicate melt) = Oz(g) + 4e- [I] at a platinum electrode were studied by both steady-state and transient galvanostatic techniques. Interest in this reaction was first developed as a result of the findings of Fulton and chipman13 that the reduction of silica, in a silicate slag, by carbon, dissolved in liquid iron, is a very slow reaction. Subsequent work, for example, by Rawling and ~lliott,&apos;~ has demonstrated that the reaction under these conditions must be slow, because the rate is limited by diffusion of oxygen in the iron to the metal-crucible phase boundary at which a CO bubble may be nucleated. Further work by Tarassof,&apos;~ in which the reduction of silica by aluminum dissolved in copper was studied, has shown that under these conditions, where carbon monoxide evolution is not involved, control of the reaction rate resides in diffusion of silica in the slag phase. However, there is no practical way of inducing sufficient convection in the system to make it clear that the phase boundary reaction is indeed fast. The overall reaction of silica reduction involves the discharge of silicon ions at cathodic sites and oxygen ions at anodic sites. In the examples cited, the discharged ions are dissolved in a liquid metal. In the present study of oxygen ion discharge, gaseous oxygen may be evolved at high current densities or oxygen may simply dissolve, possibly as oxygen molecules, in the silicate at very low current densities. The discharge of an oxygen ion at an anode must, in silicates less basic than the orthosilicate composition, be preceded by a reaction in the vicinity of the electrode, such as: which makes oxygen ions available. Platinum was chosen as the working electrode since it is comparatively inert to oxygen and is, therefore, expected to come rapidly into equilibrium with the electrolyte and with gaseous oxygen. Minenko, Petrov, and Ivanova16 have measured the electromotive force at a platinum electrode in molten silicates as a function of the partial pressure of oxygen in the atmosphere, the concentration of oxide ions in the melt, and the temperature. They found platinum to behave as a reversible oxygen electrode. At two different oxygen pressures, Po2 (I) and Pq (11). the electromotive force is given by: where F is the Faraday constant, equal to 23,060 cal per v equivalent, indicating that the electrode reaction is as written in Eq. [I.]. Platinum has been similarly used in molten silicates by other inve~ti~ators. "&apos;~~ In this investigation platinum was used only as an anode, since a current deposits other elements on its surface and changes its characteristics.

Jan 1, 1970

By Arne Bergholm

Australian rutile was chlorinated in the presence of CO or carbon. The chlorination velocity in CO was found to be strongly influenced by temperature and proportional to the CO concentration, but independent of the Cl, concentration. In the presence of carbon, the reaction velocity is much higher. The reactivity of the carbon and the distance between the carbon and the rutile surfaces are important variables. The reaction velocity is approximately proportional to the Cl, concentration and independent of the CO concentration of the surrounding atmosphere. Experiments with fluidized-bed chlorination of carbon-rutile mixtures indicate that the motion of the bed has little influence on the reaction velocity. At low temperatures, the chlorination velocity of dense tablets is much greater than that of TiO, coke mixtures suitable for fluidization. The reaction mechanism is discussed. In the industrial production of TiCl, rutile is chlorinated in the presence of carbon. Disregarding intermediate steps, the reaction may be expressed by the following equations: The ultimate object of this study was to find out whether the reaction velocity was higher in fluidized bed operation compared with chlorination of pelle-tized carbon-rutile mixtures. From a literature survey and preliminary experiments it was learned that some basic knowledge about the reaction mechanism was needed for a good experimental design. Therefore the following sets of experiments were carried out: 1) Studies on the reaction velocity in the chlorination of rutile with CO as the only reducing agent. 2) Chlorination of separate rutile-carbon tablets. 3) Chlorination of rutile-carbon tablets at different temperatures with various kinds of carbon, various grain sizes, and various tablet-making techniques. This series of experiments was carried out not only with pure chlorine but also with mixtures of chlorine with argon, CO and CO,. 4) Chlorination of rutile-carbon tablets made in a strictly standardized manner. 5) Chlorination of static-bed rutile-carbon mixtures. 6) Fluidized-bed chlorination of the same mixture. The experiments 4 to 6 formed the final and direct test of the main question: Static bed vs fluidized bedo. Although there are many patents and papers dealing with the general aspects of chlorination, only few experiments from which detailed information can be obtained have been reported in the literature. Pamfilov and coworkersl3 have studied chlorination of TiO, with CO or carbon as the reducing agent. They found that the weight decrease per hour at 600o to 800°C amounted to 11 to 14 pct with CO and 45 to 51 pct with carbon. They suggest that phosgene might be an intermediate in the chlorination of TiO,. Takimoto and Hattori have chlorinated reduced titanium oxide (TiO) and found a very high rate of TiC1, -production. They proved that the composition of the gas from the chlorination of rutile-carbon mixtures contained mainly CO. They reported for instance 74.7 pct CO, and 5.6 pct CO at 800°C at which temperature the Boudouard equilibrium composition is 12 pct CO and 88 pct CO. Seligman and Segerchano6 have studied the chlorination of TiO. They proved that Ti0 and chlorine react rather rapidly at temperatures as low as 300°C. Above 400°C, the reaction was complete. At 500°C the velocity of this reaction was much higher than was chlorination of TiO, + C. McTaggart,7 Nishimura, et al. and Wilskam have chlorinated mixtures of carbon and various kinds of rutile or beneficated ilmenite. Nishimura, et al. also report the results from reduction of TiO, with H2, CO, or carbon. At 900°C only 1 pct Ti O is formed in 2 hrs. At 1100°C 23.1 pct Ti, 0, was formed if elementary carbon was present. No carbide formation occurred below 1400°C. McIntosh and Cofferll have observed that the CO, content of the exit gases from chlorination of rutile and calcined petroleum coke is appreciably higher than found in the Boudouard equilibrium. At 900°C the ratio (CO, /CO + CO) is about 80 pct, whereas the equilibrium value is about 2 pct. W. E. Dunn12 has studied the chlorination rates of several TiO,-bearing minerals with CO + Cl, or COCl . Chlorinations were carried out either in a fluidized reactor or a fixed-bed reactor, both having a 30-mm diameter. The results obtained in both reactors were comparable. It was proved that benefi-ciated ilmenite (i.e., ilmenite from which the iron oxide had been removed by chlorination) was chlorinated 10 times faster thanrutile. Sore1 slag showed an intermediate rate. When phosgene is used, the

Jan 1, 1962

By W. C. Maurer

A study of bit-tooth penetration, or crater forniation. under simulated borehole condirions has been made. Pressure conditions existing when drilling with air, water and mud have been sirnulated for depths of 0 to 20.000 ft. These crater tests showed that a threshold bit-tooth force must he exceeded before a crater is .formed. This thresh old force increased with both tooth dullness and diflerenrial pressure between the borehole and formalion fluids. At low differential pressures, the craters formed in a brittle manner and the cuttings were easily removed. At high differenlial pressures, the cunings were firmly held in the craters and the craters were formed by a pseudoplas-tic mechanism. With constant farce of 6,500 16 applied to the bit reeth, an increase in differential pressure (sitnulated mud drilling) from 0 to 5,000 psi reduced the crater volumes by 90 per cent. A comparable increase in hydrostatic fluid pressure (simulated water drilling) produced only a 50 per cent decrease in volutne while changes in overburden pressure (simulated air drilling) had no detectable effect on crater volume. Crater tests in unconsolidated sand subjected to differential pressure showed that high friction was present in the sand at high pressures. Similar friction belween the cuttings in craters produces the transition from brittle to pseudo plastic craters. INTRODUCTION The number of wells drilled below 15,000 ft increased from 5 in 1950 to 308 in 1964. Associated with these deep wells are low drilling rates and high costs. High bottom-hole pressures produce low drilling rates by increasing rock strength and by creating bottom-hole cleaning problems. This paper describes an experimental investigation of crater formation under bottom-hole conditions simulating air, water and mud drilling. Although numerous investigators have studied bit-tooth penetration (cratering) at atmospheric pressure conditions, only limited work has been done on cratering in rocks subjected to pressures existing in oil wells. Payne and Chippendale2 have studied cratering in rocks subjected to hydrostatic pressure using spherical penetrators. Garner et aLJ conducted crater tests in dry limestone by varying overburden pressure and borehole fluid pressure independently and using atmospheric formation-fluid pressure Gnirk and Cheathem4,5 have studied crater formation in several dry rocks subjected to equal overburden and borehole pressure and atmospheric formotion pressure. Podio and Gray studied the effect of pore fluid viscosity on crater formation using atmospheric borehole and formation-fluid pressurc and varying overburden pressure. Although these studies have provided useful information on crater formation under pressure, they were limited in that the three bottom-hole pressures could not be varied independently and, therefore, that many drilling conditions could not be simulated. The prersure chamber used in this study allowed visual observation of the cratering mechanism and independent control of the borehole, formation and confining pressures. By using different fluids in the chamber, pressure conditions existing in air, water and mud drilling to depths of 20,000 ft were simulated. The mechanisms involved in cratering at these different pressure conditions were studied for teeth of varying dullness and at different loadins rates. High-speed movies (8,000 frames/sec) and closed-circuit television were used to visually study the crater mechanism under pressure. EXPERIMENTAT PROCEDURE PRESSURE CHAMBER The Pressure chamber in Fig. I was used to simulate bottom-hole pressure conditions. This chamber has been pressure-tested to 22,500 psi and is normally operated at pressures up to 15,000 psi. The chamber contains four lucite windows&apos; used for illuminating and observing the crater mechanism under pressurc. A closed-circuit television and a Fastax camera (8,000 frames/sec) have been used in these studies. Cylindrical rock specimens (8-in. diameter X 6-in. long) were subjected to three independently controlled pressures simulating overburden, borehole fluid and formation-fluid pressures. Overburden pressure, which corresponds to the stress induced by the overlying earth mass, was applied by exerting fluid pressure against a rubber sleeve surrounding the rock. Borehole pressure, which is the pressure exerted by the column of mud in the wellbore, was simulated by applying pressure to the fluid overlying the rock in the chamber. Formation pressure was simulated by applying pressure to the water saturating the rock. The borehole and formation pressures were equal except when mud was used in the chamber, in which case the differential pressure between these fluids acted across the mud filter cake.

Jan 1, 1966

By C. H. Li, R. J. Stokes, T. L. Johnson

The effect of alumina particles on the nucleation and growth of cracks through a silver- chloride matrzx has been investigated. It has been found possible to promote fibrous cracking in dispersion-strengthened silver, chloride under notch-impact conditions at temperatures at which silver chloride alone cleaves brittlely The modification) of fracture beIzavzor is thought to be due to the relaxation of hydrostatrc stress beneath a notch by the nucleation of cavities near alumina particles. In recent years, composite or dispersion-strengthened materials have been studied primarily to understand their high resistance to plastic flow particularly at elevated temperatures. Dislocation models have been developed with which it is possible to deduce with fair success the effects of interparticle distance, particle size, temperature, upon yielding and creep behavior.l-4 Much less attention has been paid to the fracture behavior of these materials (with the notable exception of common structural steels) and little is known experimentally about the manner in which inclusions affect the nucleation and growth of cracks through a matrix. Nevertheless a beginning has been made in connection with fibrous cracking in ductile matrices where inclusions appear to play an essential ro1e.5-7 During the severe localized plastic deformation which accompanies necking in a tensile test, cavities are believed to develop at inclusions; these cavities subsequently grow and coalesce by plastic flow until separation is complete. It is of interest to consider whether inclusions can affect fracture behavior under loading conditions which restrict the plasticity of the matrix itself (for example, cleavage under conditions of a high imposed strain rate at low temperatures). It is particularly interesting to study these effects in a solid which shows a spectrum of behavior ranging from fully ductile to semi-brittle behavior. Such a solid is silver chloride whose mechanical behavior depends sensitively upon temperature and strain rate.&apos;," The present paper is concerned with a study of the influence of inclusions (in the form of alumina particles) on the fracture behavior of silver chloride loaded uniaxially at low strain rates at room temper- ature and also under notch impact conditions over a wide range of temperature. In particular, it will be shown that the alumina particles can exert a startling effect on the ductile-brittle transition temperature of notched silver chloride and that the magnitude and nature of the effect depends upon both the quantity of alumina and the shape of the alumina particles. 1. EXPERIMENTAL PROCEDURE 1.1 Materials Used. Silver chloride powder of analytical reagent (AR) quality having an average particle size 6 was supplied by the Mallinckrodt Chemical Works (st. Louis, MO.). Acid washed 900 mesh alumina powder, designated A38-900, was supplied by the Norton Co. (Cambridge, Mass.). This powder was added to silver chloride in two forms: a) the as-received condition in which the individual particles were of random irregular shape; their statistical average size was 7; b) in a condition in which each particle was spherically shaped by a fusion technique. In this case, the statistical average particle size determined with the optical microscope was approximately the same (about 5p) but electron micrographic evidence indicated that many ultra fine particles were present in the spheroidized powder. 1.2 Preparation of Composite Materials. Silver chloride-alumina composites containing 2.5, 5, and 15 pct by volume of alumina were produced by the extrusion of mechanically mixed powders blended in a ball mill for 24 hr at room temperature. The mixtures were compacted at 50,000 psi at room temperature in the form of billets 3/4 in. in diameter and 1 in. long which were then extruded with a 16:l reduction ratio at 370°C through a radius-type steel die having a 5 deg lead-in angle. An extrusion temperature of 370°C was selected to ensure that all composites had sufficient plasticity to be extruded. Apart from this general requirement, the choice was arbitrary. 1.3 Microstructure of Composites. Attempts were made to check the distribution of alumina metallo-graphically by polishing transverse and longitudinal sections of the extruded rod. Specimens were wet-ground to 600 emery paper and lapped successively with 5 and 0.25-p grades of diamond paste. They were etched for 10 sec in 10 pct sodium thiosulfate solution and lightly polished in concentrated ammonium hydroxide. The most effective way to render the alumina particles and grain boundaries visible was to radiate the surface with intense white light to decorate the grain boundaries and the particle-matrix interfaces photolytically.

Jan 1, 1962

By Francis R. Conley, John A. Sievert, John N. Dew

A brief review is presented of the past performance of a number of large, thin, highly permeable reservoirs with low dips in the Bolivar Coastal fields of Venezuela. The performance of these reservoirs indicates that the fluids are segregated vertically within the sand section by gravity. With this assumption, equations are developed which describe the performance under pressure maintenance operations. Methods of solving these equations and results of simple example calculations are presented. Example calculations indicate that under pressure maintenance conditions injected gas tends to {tow preferentially along the top of the sands and that encroaching water has a tendency to flow preferentially along the bottom. The expected performance of segregated fluids is discussed and compared with that of fluids which are uniformly distributed in any sand section. INTRODUCTION The field performance of a number of reservoirs described herein indicates that the fluids are segregated vertically within the sand section by the force of gravity. Thus, it was felt that any method used to predict the future performance of these reservoirs should consider the effects of segregation in the sand sections. A study of the available information on some of the reservoirs suggests that increased recoveries may be expected if the reservoir pressure is maintained by either crestal gas injection or flank water injection. To predict future performance of reservoirs under pressure maintenance operations, a method of analysis was needed which would account for segregation of fluids in the sand sections. Several methods of analysis have been developed to take into account the segregation of fluids in the reservoir as a whole1,3 To our knowledge no method considers segregation of the fluids within sand sections in the manner indicated by the past performance of several reservoirs in the Mara-caibo Basin. This paper outlines part of the work done in studying some of these reservoirs and contains a description of their performance characteristics. The analysis presented is restricted to pressure maintenance conditions, since space limitations prevent a full discus- sion of the development of the mathematical relations and the various methods of solving them. Only a sketch of the development of the mathematical relations is given. The various methods of solving these relations are pointed out, but the actual determination of solutions to various problems are omitted except for one example. It is felt that these results may aid in the study of reservoirs outside the Maracaibo Basin. Some concepts on which the present analysis is based are outlined by D. N. Dietz. However, the analysis presented herein includes several factors not considered by Dietz: (1) variations in permeability and in the cross section, (2) various shapes of the cross section, and (3) the production of fluids. Also, the mathematical development presented by Dietz differs considerably from our corresponding analysis. RESERVOIR CHARACTERISTICS The reservoirs under consideration, large and thin, with low dip and high permeability, are large to the extent that they contain from 0.5 to over four billion bbl of oil initially in place. The section thicknesses vary from 100 to 400 ft and the lengths from 10,000 to 40,000 ft. Thus, a length-wise cross section of the producing formation appears to be long and thin. The dip angles vary from 0 to 6 degrees and the average permeabilities vary from 0.5 to over three darcies. The reservoirs contain from 20 to 50 per cent shale inter-laminated with the productive sands. Most shale breaks within the major sands do not correlate from one well to the next, and correlation is often difficult, even between wells drilled from the same location. However, it is often possible to correlate the shale breaks between major sands for some distance. Herein we are concerned with the performance of reservoirs containing oil of gravities over 20" API. Past reservoir performance indicates good pressure communication and, in cases where pressure sinks have developed, large amounts of fluid migration have occurred. Few of the reservoirs have had initial gas caps, but also, few have been found to be highly under-saturated. Due to gravity segregation, secondary gas caps usually form before the pressure has been reduced more than 25 per cent of its initial value. The effect of gravity has not only been apparent on a reservoir-wide basis, but has also caused segregation of the reservoir fluids in productive sections. Proof of this is found from the results of selective well tests, workovers, and electric and various other types of logs. Evidence that

By F. O. Garrabrant

Power is furnished to the Climax Molybdenum Co. by the Public Service Co. of Colorado over two 100,000-volt lines to a bank of three 3333-kva. transformers 100/13.8 kv. These transformers are so designed that with the addition of forced cooling their capacity may be increased 33? per cent at 12,000-ft. altitude. Under present conditions, these transformers are operated at 140 Per cent of their rating with forced air cooling. On the secondary side of this bank there is a 15-kv., 1200-amp. Westinghouse oil circuit breaker provided with automatic reclosing relays allowing three successive short circuits before lockout. From this breaker 13.8-kv. power is transmitted to a central switching station (Fig. I) where all circuits radiate to the different operations of the property. Each circuit is protected by a 23-kv., 600-amp. oil circuit breaker with by-pass and isolating disconnecting switches. Mine Power at 13.8 kv. enters the mine through a ventilation raise over a three-conductor, size 1-0 varnished cambric insulated, leaded, steel-wire armored cable, insulated from 17 kv. This power is received and distributed through one switching station and four underground substations, shown in Fig. 2. No. 370 substation (Fig. 3) contains three a 13,800/440-volt, single- phase Pyranol transformers protected on the high-voltage side by a 600-amp. manually operated oil circuit breaker. This station serves 20 slusher-hoist motors, a 45-kw., three-phase, 220-volt electric furnace, a Massco hot miller for tempering and milling Timken detachable bits and a hoist with 175-hp motor on the No, 3 development shaft. Also at this shaft are three pumps. One two-stage pump, 1200 gal, per min., 500-ft. head, 3600 r.p.m., driven by a 150-hp. motor, is at the bottom of the shaft, Two deep-well pumps, 500 gal. per min., 500-ft. head, are at the collar of the shaft. under normal operation the one pumping 1200 gal. per min, is controlled by a float switch to take care of the water. If this pump fails, or if its capacity is insufficient, the first 500-gal. pump is automatically started1 and if still more capacity is needed the second 500-gal. pump automatically starts. Time delay relays are provided on the deep-well pumps, so that if power fails the column will have time to drain before the Pumps restart. On failure of any Of these, an alarm is sounded in the train dispatcher&apos;s office, so that mechanics can be sent to ascertain the cause of failure and make repairs. The cable from No. 37&apos; substation is run to the switching station, in which are four 15-kv., 600-amp. direct-current sole-noid-operated oil circuit breakers. The current for Operating and tripping is furnished by a 125-volt storage battery. All breakers are operated from the switchboard in the compressor room. One circuit from this switching station goes to the

Jan 1, 1946

By Harold Margolin

The Ti-Al system certainly merits the investigative attention it has been receiving and this latest contribution by Blackburn is therefore to be welcomed. The titanium-rich end of the phase diagram shown here, as well as the diagram by crossley,17 Clark et a1.,18 and Tsujimoto and Adachi,29 reveal no other intermediate phase than Ti3A1, although a previous work by Ence and Margolin20 had reported an additional intermediate phase. Since then, Buerger Precession pictures of single crystals of 10.5, 11.5, and 12.5 wt pct A1 were made at New York university21 and the extinctions were found to be typical of a D6h a Ti structure. Some superlat-tice spots and/or streaks suggesting the presence of an Mg3Cd structure, i.e., Ti3Al, were also found. Since these specimens had been quenched from the y region20 and had revealed a Ti, it was concludedz1 that y does not exist and the structures interpreted as m + y could not be so considered. The absence of normal two-phase behavior in these structures had been indicated in the paper. Accordingly, the two-phase region y + 6 becomes a + Ti3Al and there is some agreement between the results of Blackburn and of Ence and Margolin20 as to both shape and existence above 900°C of an a + Ti3Al field. Crossley 17 proposed that Ti3Al is a line compound which decomposes congruently at 900°C. It is difficult to discount entirely the fact that clear two-phase structures, varying in amount with position in the two-phase field, could be detected above 900°C by Ence and Margolin.20 Precipitation during quenching could not have been so regular and the twin-phase syndrome17 does not apply. The heat treatment and quenching technique was essentially the same as was used by Crossley, whereas Blackburn, who used a rapid quenching procedure, also showed a two-phase a + Ti3A1 field above 900°C. If the two-phase field exists above 900° C, it is easy to picture a two-phase field which could account for the data of Crossley below 900°C without invoking a line compound existence of Ti3A1. All that would be necessary would be an 0 + Ti3Al/Ti3Al boundary which curved sharply to higher aluminum contents below 900° C. It is clear that the controversy must be solved by techniques which obtain data at elevated temperatures without quenching. It is conceivable that the difference in width of the a + Ti3N fields reported by Blackburn and by Ence and Margolin is due to the longer annealing times used by the latter investigators. The criterion for establishment of equilibrium in our work20 was the

Jan 1, 1969

By E. Clarence Peterson, Rip V. Thompson

ALTHOUGH the metallurgical treatment for the Ophir Hill tailings was developed more than 20 years ago and in the interim a vast amount of experimental work was done with later-developed flotation reagents, it was found necessary to revert to the original oil reagents and their combinations with specific chemical reagents in order to obtain results of economic value. This is of especial interest because in present ore-treatment investigations later-developed flotation reagents are used almost exclusively and the coal-tar derivatives, except in certain patented reagents, are virtually neglected. These coal-tar oils and derivatives still have a definite field of application, especially in the treatment of some oxidized ores when modified sulphidization is desirable. The Ophir Hill tailings deposit is in Ophir Canyon, Tooele County, Utah, about 55 miles southwest of Salt Lake City and 25 miles south of the Tooele smelter of the International Smelting and Refining Co. A road good for trucks leads directly from the Ophir mill to the smelter. TAILINGS These tailings were produced in the old Ophir Hill gravity concentrator previous to 1921. The original mill ore was a heavy lead-zinc-copper-silver-iron sulphide ore formed by limestone replacement. The gangue is a highly siliceous lime. An assay of the original milling ore would approximate 1.0 per cent Cu, 6 to 8 per cent Pb, 6 to 8 per cent Zn, 8 to 12 oz. Ag, 18 to 25 per cent Fe, 2.0 per cent CaO, 20 to 5o per cent insoluble. The gold content was 0.005 oz. or less. Tailing was impounded in a fairly wide area in a canyon about ½ mile below the mill. A number of dirt embankments were thrown up to form settling ponds and the tailing was allowed to flow through them in series. Coarser sand deposited at the upper end of the ponds and the finer slimes settled gradually as they passed through the successive settling ponds. As soon as the coarser tailings filled a pond, the tailings were by-passed to the next pond and teams and scrapers were used to haul the deposited sand onto the banks to make them higher. The final dump was approximately 2500 ft. long and 250 ft. wide at the bottoms, with banks sloping up at 45°. The coarser sands predominated at the upper ends and sides of the ponds while the proportion of slimes increased steadily at the lower ends. The lower end of the pond contained over 50,000 tons of slime tailings, 90 per cent of which passed through a 200mesh screen. SAMPLING The entire tailings deposit was sampled by drilling vertical holes so placed as to give data concerning size segregation and the contour of the underlying ground as well as to secure an assay sample. Different types of auger and post-hole drills were tried. None was found satisfactory because the stratified layers of

Jan 1, 1940