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By Harry A. Wahl

Precent fracture design procedures are bared on the total fracture area created. A method to distinguish beI,,.ecn total area and [he propped or effective fracture area has not been available. This paper presents a solution to this problem, applicable to horizontal fractures. The difierences between effective fracture area and torn] area are demonstrared in example calculations. This work is hayed on experimentally determined transport efficiencier of solids in sand-liquid slurries. Newtonian and non-Ne~vtoninn systems are considered. INTRODUCTION Fifteen years after commercial introduction, hydraulic fracturing remains the most successful stimulation technique in the oil field. This success is primarily due to ability of induced fractures to penetrate and alter permeabilities deep within formations. Many fields producing today could not have been developed without the hydraulic fracturing process. Because of wide usage, fracture-treatment design has received a great deal of engineering and research effort. This work, resulting in improved equipment and materials, has increased the benefits from fracture treatments as well as the applicability of the process. A major contribution was the development of fluid-loss additives. Necessarily, the number of parameters to be considered in treatment design has steadily increased, resulting in more complicated design techniques. Almost all present design procedures are based on the precepts set forth by Howard and Fast. Relating the fluid volume lost into the formation, the volume required in extending the fracture, and the total slurry volume injected, they developed an expression for the total fracture area created in terms of pertinent treatment parameters. Fluid loss during treatment was expressed as a function of time for three flow mechanisms. Although modifications of fluid loss equations have been made, the total fracture area concept has remained unaltered. A vast amount of field data indicate that induced fractures must be propped and held open to be effective. A notable exception is the Mesa Verde formation in the San Juan basin. However, analysis of these treatments shows that improved well productivities are obtained when propping agents are incorporated in the treating fluid. Although propped fracture area has been recognized as an important design parameter, a method to distinguish between total area and effective fracture area has not been available. The necessary information on slurry-sand transport in fractures has been lacking. Interest in the propped region of induced fractures is not restricted to areal extent alone. The distribution of sand within fractures is important from the standpoint of fracture flow capacities. Flow capacity affects the increase in well productivity after stimulation. The work of Huitt and Darin4 hows that partial monolayers of sand have large flow capacities compared to thick* dense sand packs. It has been postulated that gelled fluids have the ability to transport sand within the fractures at the deired low concentrations. An early contribution in the area of sand placement in fractures was made by Kern et al.' They studied sand movement in a transparent vertical fracture model. It was observed that the sand tended to settle out in the bottom of the model before moving very far. When the fluid velocity exceeded a certain critical value, all of the sand injected began moving through the crack even though it settled to the bottom. This critical velocity was determined under several flow conditions. Some work on sand movement in horizontal fractures has been reported in Russian publications. Sand movement was studied by Izyumova and Shan'gin' using a transparent "pie-shaped" flow model to simulate a horizontal radial flow system. However, the data were limited, especially in a quantitative sense. Dorozhkin, Zheltov and Zheltow studied the behavior of sand-liquid slurries in a horizontal linear flow model. The quantitative data were restricted primarily to the thickness of sand deposits formed at the bottom of the fracture. An earlier paper provided basic data on the flow of sand in horizontal fractures. This study was designed to yield specific quantitative information on rate of advance of sand particles and pressure behavior under various flow conditions. A comprehensive photographic study was undertaken in a 10-ft windowed flow cell to provide the necessary qualitative and quantitative data. Since the number of potential variables far exceeded the capacity of the initial study, emphasis was placed on the effects of sand concentration, oil viscosity and oil flow rate. A detailed description of these experiments and the results are described in Ref. 9. However. the implications of this work on the fracture design calculations were not discussed. An analysis of these data as well as new data is provided in the following sections. EXPERIMENTAL RESULTS The primary objective of the experimental investigation was to provide information on the rate of advance of the solids in sand-liquid sturries. A 10%-ft long transparent

Jan 1, 1966

By G. P. Airey, R. F. Mehl, T. A. Hughes

The coarsening of cementite particles in a ferrite matrix has been studied in a series of steels with 0.15 pct C only and 0.15 pct C plus 1 pct Ni, Mn, and Cr, respectively. Two initial states were employed: quenched nartensite, and quenched and cold-rolled martensite. A series of tempering temperatures between 500' and 700" and tempering times of up to 190 hr were used. The structures were studied by replica and transmission electron microscopy. Particle size distribution curves were determined. From the average size value coarsening curves were obtained. These were plotted in accordance with the Wagner analysis assuming diffusion control. A discussion of the significance of the results is given. L HE reactions occuring upon the tempering of martensite have long engaged the attention of metallurgists. The latter stages, when cementite particles coarsen in a ferrite matrix, have been studied both qualitatively and quantitatively. Studies of such coarsening processes have recently been spurred by the publication of the Lifshitz-Wagner theory1, and the extension of this to the a Fe-Fe3C system by Oriani3 and by Li, Blakely, and einold. Following Wagner the coarsening process is often designated as "Ostwald Ripening". The only quantitative data on the rate of coarsening, except for the work of Hyam and uttin,' in the a Fe-Fe3C system are those of Bannyh, Modin, and odin' for a commercial eutectoid steel and those of Heckel and ereorio" using a pure eutectoid steel. The data of Bannyh, Modin, and Modin have been employed by 0riani3 to derive the a Fe-FeE interface energy. The reaction is one of the most important ones in steel and is worthy of detailed study. This is the purpose of the present study. Laboratory heats were prepared; these were steels with approximately 0.15 pct C, selected so that the num ber of carbide particles would be relatively small and thus so that the overlapping diffusion fluxes would be minimized, presumably a desirable circumstance.'-3 In addition to Fe-C alloys, comparable heats containing 1 pct of Ni, Mn, and Cr, respectively, were included with a view of appraising the effect of alloying elements. This report includes an account of the micro-structures observed, primarily with the electron microscope, and of kinetic data and their interpretation. MATERIALS AND TREATMENT The alloys were prepared from electrolytic iron ("Plastiron") and high-purity graphite; these were melted in a zirconia crucible using a vacuum furnace. The alloy steels were made by adding electrolytic nickel, electrolytic manganese, and "vacuum grade" chromium, respectively, under a partial pressure of argon. Each melt was poured into a mold within the vacuum furnace and cooled in the mold. The ingots were 2 in. in diam. and 8 to 10 in. long. The analysis of the alloys is given in Table I. These ingots were hot-rolled to strip 0.1 in. thick, then cold-rolled to 0.05 in. and each alloy split into two batches. One batch was austenitized at 1200 for 1 min, quenched in cold brine, then cold-rolled to 0.02 in.; samples given this treatment are hereinafter designated as "worked". The other batch was cold-rolled to 0.025 in., austenitized at 1200" for 1 min, and quenched in cold brine; such samples are hereinafter designated as "quenched". These two batches were then tempered, as below. The purpose of the treatment given the first batch was to provide an initial structure of cold-worked martensite, with the expectation that the additional defect structure created by cold work would encourage a higher rate of nucleation of cementite on tempering and hence a more uniform distribution of cementite particles. Individual specimens were sealed in evacuated quartz or Pyrex tubes, then tempered in a muffle furnace. The temperature control was better than 3'C at 700. Tempering treatments wer: performed at 400°, 500°, 550°, 600°, 65o°, and 700C for time periods between 15 min and 190 hr. PREPARATION OF SPECIMENS Specimens for optical and replicalelectron microscopy were mounted, polished conventionally, and etched with 2 pct nital. For electron microscopy, single-stage "formvar" replicas were made, dry-stripped and rotary-shadowed with chromium at an angle of 30 deg. Carbide extraction replicas were prepared from electropolished specimens usirig the method described by Smith and uttin.' Thin foils for electron transmission microscopy were prepared by chemical thinning in an H202-HF bath prior to electropolishing in a chromium trioxide-acetic acid solution. The most

Jan 1, 1969

By R. W. Regehr, R. A. Burmeister

Epitaxial growth of GaAs 1-xPx on germanium substrates was achieved using an open tube vapor transport system. The compositional range of 0.3 < x < 0.4 was examined. The best results were obtained with (311) orientation of the germanium substrate. The physical and chemical properties of the resulting layers were investigated using several techniques. Spectrographic analyses of the layers indicate substantial incorporation of germanium into the GaAs t-X Px layer. Evidence is presented which indicates that this incorporation occurs via a vapor phase transport process rather than by solid phase dijfu-sion. Electrical measurements suggest that the germanium thus incorporated behaves predominantly as a deep donor in the compositional range of 0.33 < x * 0.40 and has a deleterious effect upon the luminescent properties of GaAs1-x Px. The increasing technological importance of GaAs1-xPx for use in light-emitting devices has led to an evaluation of several aspects of existing growth processes. The method most commonly used to prepare GaAs1-xPx for electroluminescent device applications is vapor phase epitaxial growth on GaAs substrates.&apos;-4 In a typical electroluminescent diode structure the active region of the diode is entirely within the epitaxial layer and thus the electrical properties of the substrate are relatively unimportant since it is effectively a simple series resistance (assuming hetero-junction effects to be negligible). The use of germanium rather than GaAs as the substrate material is of interest for several reasons. First, GaAs of reasonable structural quality has been epitaxially grown on germanium4-2 and it is reasonable to expect that GaAs1-xPx could subsequently be deposited on the GaAs layer. Second, germanium substrates are readily available with both lower dislocation densities and larger areas than GaAs. Finally, single crystals of germanium are more economical than GaAs single crystals. The principal objective of the present investigation was to test the feasibility of growing GaAs1-xPx epi-taxially on germanium substrates, and to evaluate the properties of such layers with regard to electroluminescent device requirements. The approach used was to a) demonstrate epitaxial growth of GaAs1-xPx on germanium, and b) characterize the relevant structural, electrical, and optical properties of the GaAs1-xPx layers. The possibility of germanium incorporation into the grown layers was of special interest since there was some indication of this in previous studies of GaAs growth on germanium.5&apos;11,12 Although a study of the electrical properties of germanium in GaAs1-xPx was not an intent of this investigation, several features of the electrical properties of the layers grown in the present study which appear to be due to germanium are described. EXPERIMENTAL PROCEDURE The open-tube vapor transport system used for the epitaxial growth of GaAs1-xPx is illustrated in Fig. 1. This system utilizes the GaC1-GaC13 transport reaction and is similar in most respects to the larger system described elsewhere.&apos; The germanium substrates were n-type, with a resistivity of 40 ohm-cm (Eagle-Picher Co.). These were cut to the orientations of {100), {111), and (3111, and were mechanically polished and chemically etched in CP-4 (5 min at 0°C) prior to growth. In some cases, a GaAs substrate was employed in addition to the germanium. The orientation of the latter was {loo}, and they were also mechanically polished and chemically etched prior to growth. The initial composition of the deposited layer was pure GaAs. After approximately 10 microns of GaAs was deposited on the germanium substrate, the phosphorus content of the layer was gradually increased over a distance of approximately 15 microns to the desired concentration and maintained at this value throughout the remainder of the growth. Typical operating parameters used during growth are given in Table I. Selenium was used as a n-type dopant in several runs to facilitate comparison of the electrical properties of the layers grown on germanium with those of layers grown on GaAs substrates, which are usually doped with selenium. The concentration of H2Se in the gas phase was adjusted to a value which would normally yield a carrier density of 1 to 5 x 101 7 at room temperature in layers grown on GaAs substrates. The terminal surfaces of the epitaxial layers were examined by optical microscopy for structural characteristics. Laue back-reflection photographs (Cu radi-ation) were also made on the terminal surface to verify the epitaxial nature of the deposit. After these steps

Jan 1, 1970

By R. O. Williams

It is shown that Ni3Al precipitates homogeneously from nickel-rich alwminum alloys as plates on the (100) planes. Prior to actual precipitation a process occurs which is believed to be one of increasing short-range order. After precipitating the Ni3Al plates enlarge through competitive growth. Discontinuous precipitation can occur simultaneously with the above processes. Recent ideas of the origin of precipitation strengthening appear adequate to explain the hardness changes. REMARKABLY little appears to be known about the precipitation process in Ni-Al alloys in spite of their technical importance. This investigation originated to supply additional information about precipitation in general, this system in particular. Information on the structures and kinetics have been obtained through the use of hardness, X-rays, microscopy, calorimetry, and resistivity on high-purity alloys. PROCEDURES Six alloys, Table I, were prepared by melting carbonyl nickel and high-purity aluminum in alumina crucibles in vacuum and casting into 1-in. graphite molds. All rods were homogenized at least once at 1300°C for 24 hr prior to swaging and this was repeated on the first three alloys after 75 pct reduction. Alloy 4 could be reduced only 10 pct at 1000°C (probably in two-phase field) prior to fracture but 1/4-in. samples quenched from 1100°C were readily reduced cold. Alloy 5 was reduced 15 pct cold but failed on the next pass while alloy 6 of essentially the same aluminum content failed inter-granularly without apparent flow up to 1000°C. The alloys were heated in hydrogen at the elevated temperatures and formed thin, coherent aluminum oxide coatings which provided excellent oxidation resistance at lower temperatures. However, freshly prepared surfaces showed considerably less resistance at 500"to 700°C in air and apparently resulted in internal oxidation. As a consequence, low-temperature agings were carried out in evacuated tubes. RESULTS The isothermal hardening behavior of these alloys at 500"and 565C is given in Figs. 1 and 2. These results were obtained from samples cold worked 75 pct, recrystallized at 1000°C (1100°C for the 7.8 pct Al) and quenched in water. This recrystallization was used to give smaller grain sizes so as to obtain more uniform hardness values and the points represent an average of five readings. The electrical resistivity was measured on 1/16-in. wires quenched from 1000°C during aging at 495°C to give Fig. 3. The energy release and its rate are given in Fig. 4 for the 6.9 pct Al alloy during aging around 500°C. Inasmuch as this was a single run, its accuracy is not known but certainly the general shape and magnitudes are correct. The method used to obtain these results is described elsewhere.&apos; Data for the aging at 600°, 700°, and 800°C of these alloys cold worked 50 pct are given in Fig. 5. Supplementary information from microscopy and X-ray diffraction have been included to indicate recrystallization, discontinuous precipitation and the appearance of superlattice lines from the Ni3Al. The hardness of these alloys as annealed, aged, cold worked, and cold worked and aged is given vs composition in Fig. 6. Those samples which were isothermally aged, Figs. 1 and 2, were reaged at 532°C and at successively higher temperatures for the indicated times to give the data of Fig. 7. These results as well as certain others, support the idea that the level of hardness reached for temperatures above 600°C are equilibrium values more or less independent of path. This being the case, the breaks in the curves would be the complete solution of the Ni,Al. The electrical resistivity versus temperatures for some of these alloys, both aged and unaged, is given in Fig. 8 along with those data from heating slowly (10 deg per day) to high temperatures. Interesting points include the lowering of the Curie temperature (the change in slope), the lack of any indications of a solubility limit and the large temperature coefficient for the Ni3Al. A slight break for Ni3Al around 100C shows up but this is not a Curie temperature as Ni3Al is not ferromagnetic down to -190°C. Metallographically both the nickel-rich solid solution and the Ni3Al appear very much like pure nickel. Profuse twin boundaries are present both

Jan 1, 1960

By J. G. Richardson, J. W. Graham

The purpose of this paper is to present the results of theoretical and experimental studies of water imbibition. The imbibition processes are involved in recovery of oil from stratified and fractured-matrix formations in natural water drives and water flooding. An understanding of the role of inhibition in implementing the recovery of oil from such formations is deemed essential to proper control of these reservoirs to achieve maximum recovery. The theoretical studies involved development of the differential equations which describe the spontaneous imbibition of water by an oil-saturated rock. The dependence of the rate of water intake by the rock on the permeability, interfacial tension, contact angles, fluid viscosities and fluid saturatiorls is discussed. A few experiments were performed using core samples to determine the effects of core length and presence of a free gas suturation. The role of water imbibition in recovery of oil from a fractured-matrix reservoir by water flooding was investigated by use of a laboratory model. This model was scaled to represent one element of a frac-tured-matrix formation. Water floods were made at various rates with several fracture widths. Interpretations were made of the behavior expected in a system containing many matrix blocks. The presence of a free gas sntu.ration was found to reduce the rate of water imbibition. In the reservoir prototype of the fractured-matrix model, water imbibition rather than direct displacement by water was the dominant mechanism in the recovery of oil at low rates. INTRODUCTION Imbibition may be defined as the spontaneous taking up of a liquid by a porous solid. The spontaneous process of imbibition occurs when the fuid-filled solid is immersed or brought in contact with another fluid which preferentially wets the solid. In the process of wetting and flowing into the solid, the imbibing fluid displaces the non-wetting resident fluid. Common examples of this phenomenon are dry bricks soaking up water and expelling air, a blotter soaking up ink and expelling air and reservoir rock soaking up water and expelling oil. As increasingly better lithological descriptions have been made of the characteristics of petroleum-bearing formations, it has become obvious that imbibition phenomena which were once considered laboratory curiosities are of practical importance. For instance, in reservoirs composed of water-wet sand strata of different permeability in intimate contact, the tendency of water to channel through the more permeable stratum is offset by the tendency for water to imbibe into the tight sand and expel oil into the coarse sand. Also, in fractured-matrix formations the tendency of water to channel through the fractures is offset by water-wet matrix blocks. As some imbibition of the water into the of the largest fields in the world are fractured-matrix reservoirs, it has become increasingly important to understand all the factors involved in the imbibition process. Examples of fractured-matrix reservoirs are the Spraberry field in West Texas which produces from a fractured sandstone&apos;, the giant Kirkuk field in Iran&apos;, the Dukhan field in Qatar, Persian Gulf2, and the Masjid-I-Sula-main and the Haft-Kel fields in Southwestern Iran, which produce from fissured limestone3. Research into recovery of oil from fractured-matrix formations was stimulated by the rapid decline of oil productivity of wells in the Spraberry formation. One result of this research was the water imbibition process developed by the Atlantic Refining Co.4 Another idea was that much of the Spraberry oil could be recovered by conventional water-flooding procedures5. Subsequently, pilot floods were conducted in this field to test the feasibility of these ideas. It was felt that an understanding of the role played by imbibition processes in displacement of oil from a fractured-matrix reservoir could not be obtained from field data alone because of the many complicating factors and uncertainties involved. Therefore, theoretical and laboratory studies were undertaken to provide this understanding. Study of the equations which describe the linear, countercurrent imbibition process provided an insight into the role of various factors in the process, such as the permeability of rock and inter-facial tension. In addition to the theoretical studies, imbibition experiments were conducted with core samples to determine the effect on the rate of imbibition of such variables as core length and free gas saturation. The principal experimental studies were conducted by water flooding a scaled model of an clement of a frac-tu red-matrix reservoir to evaluate

By T. Toda, R. Maddin

A study has been made of the microstructure and mechanical properties of splat-cooled aluminum-base gold alloys with gold concentration from 0.25 to 5.0 wt pct. These alloys have been quenched from the liquid state by a torsion-catapult technique, which has made it possible to pepare specimens suitable for mechanical property measwements. From the electron micrographs it has been shown that the solid solubility of gold in aluminum can be extended to 2.5 wt pct (0.35 at. pct) by splat-cooling, while the maximum equilibrium solubility is known to be less than 0.3 wt pct (0.04 at. pct). The very fine grain size (several tenths of a micron), the extended solid solubility, and the fine dispersion of a second phase (AuAl2) contribute concurrently to a substantial strengthening effect. In Al-5 wt pct Au splat-cooled specimens of less than 50 thickness, the yield strength is 17 kg per sq mm or 6 times as large as the strength of bulk specimens. For the Al-1.0 to 2.5 wt pct Au solid solution obtained by splat-cooling, the yield strength reaches 7.5 kg per sq mm after an aging treatment (for 10 hr at 200°C), while it is 3.7 kg per sq mm for the corresponding bulk specimens. A great deal of research has been done in recent years on the structure and the properties of metals and alloys rapidly quenched from the liquid state.&apos; The term "splat-cool" has been used with the meaning of a rapid quenching from the liquid state., The splat-cooling techniques have produced large numbers of new structures, which are expressed in terms of metastable phases,3 concentrated solid solutions,4 amorphous phases,5&apos;6 new phases,7 and so forth. Nearly all previous studies have concentrated on the physical properties; i.e., crystallography, structure, electrical resistivity, magnetism, and so forth, of the splat-cooled metals and alloys. The mechanical strength of splat-cooled metals and alloys has hardly been investigated except for some recent work by MOSS&apos; on A1-V alloys. The principle common to all experimental techniques developed to obtain very rapid quenching rates is based on the heat transfer by conduction. Liquid must be in good thermal contact with a substrate of high heat conductivity. Both of the published devices known as the "gun" and the "piston and anvil" techniques suffer from certain shortcomings. For example, the specimen obtained by the gun technique is very small and flaky, and hence inadequate for mechanical properties measurements. On the other hand if the material is forced to yield a continuous speci- men by the piston and anvil technique, it is probable that some plastic deformation occurs during the quench. A novel method for rapid quenching of a liquid metal or alloy, the "torsion-catapult", has been devised by Roberge and Herman9 at the University of Pennsylvania. In the apparatus the melt is thrown out of a curved furnace by a catapult and impinges on a copper substrate. The apparatus has the advantage of producing a continuous foil which is relatively large in size and of a quality suitable for the measurements of mechanical properties. The quenching rate is estimated to be of the order of l05 to l06 ºC per sec, (comparable to rates achieved by the piston and anvil technique). In selecting an alloy to be studied we were made aware of the fact that gold was believed to be "insoluble" in in and consequently age hardening in the A1-Au system appeared to be interesting. Quite recently Heirnendahl13-15 revealed that the solid solubility, as determined by transmission electron microscopy, was 0.3 wt pct Au at 640°C and 0.25 wt pct Au at 600°C, decreasing with decreasing temperature. In an A1-0.2 pct Au alloy after quenching from a solution treating temperature of 600°C the yield stress was 2 kg per sq mm, and it increased up to 6 kg per sq mm after aging for 1 to 10 hr at 200°C. The precipitation occurred in the form of platelike particles mainly on (100) matrix planes. The intermediate phase n&apos;, the equilibrium phase n (AuAl2), and lattice relationships between both precipitates and the matrix were also investigated by electron microscopy. One of the purposes of the present research is to determine whether or not the solid solubility in this system, in which gold has a very small solubility in

Jan 1, 1970

By G. L. Fisher, R. H. Geils

The effect of small amounts of columbium (<0.01 to 0.10 pct) on the ?-a transformation occurring during the continuous cooling of a low carbon Ni-Cu steel was investigated. Dilatometer specimens were aus-tenitized at 950" and 1068?C and cooled at 17? and 375 C° per min. Columbium caused a marked depression in the ?-a transformation temperature except when cooling at the slower rate from 950°C. The effec of columbium on the transformation temperature was greater the higher the austenitizing temperature and rate of cooling. A maximum depression of 92 C" was observed. Metallographic examination of specimens of <0.01 and 0.07pct Cb steels heated at 1200°C for 1 hr and cooled at various rates showed that columbium had a major effect on the ferrite morphology. The fer rite in the columbium -free steel remained equiaxed at cooling rates as high as 440 C? per min while the columbium-bearing steel exhibited mixed structures o equiaxed and bainitic ferrite at cooling rates as low as 130 C° per min. The ? grain boundaries in the columbium -free steel provided the ferrite nucleation sites in rapidly cooled specimens. There was a complete absence of nucleation at these sites in the colum bium-bearing steel. It is concluded that columbium depresses the transformation temperature by suppressing ferrite nucleation at the austenite grain bound-aries. In this respect the effects of columbium are analogous to those of boron in low C-Mo steels. It is well known that small columbium additions can substantially strengthen plain carbon steels. As little as 0.02 pct Cb can increase the yield strength of mild steels by 10,000 psi.1 A fine precipitate of CbC has been observed in columbium-bearing steels2 and is generally thought to be responsible for the strengthening. Little attention has been devoted to the effect of columbium on the ?-a transformation. Webster and woodhead3 have studied the effect of columbium on the isothermal proeutectoid ferrite reaction in mild steels. They found similar transformation behavior in steels both with and without columbium additions. However, as the austenitizing temperature increased, the incubation time for the start of the ferrite transformation became longer in the columbium-containing steel. Morrison1 found that the addition of 0.03 pct Cb to a C-Mn steel lowered the transformation temperature by 50 C° during cooling from 1200°C at a rate of 80 C" per min. The strengthening effect of columbium has recently been utilized in an age-hardenable, low-alloy steel containing copper and nickel.4 A small amount of columbium has a substantial effect on the as-rolled strength of this steel. By increasing the columbium level from <0.01 to 0.13 pct the as-rolled yield strength is increased by 15,000 psi. Columbium also significantly lowers the ?-to-a transformation temperature of this steel during continuous cooling from the austenitizing temperature. Because of the low carbon level in this steel (0.05 pct max), it is almost entirely ferritic. Thus, it offers the opportunity of studying the effect of small columbium additions on the proeutectoid ferrite reaction. Of particular interest in this study was the reason for the marked lowering of the transformation temperature by columbium during continuous cooling. EXPERIMENTAL PROCEDURE Materials. The compositions of the steels used in this investigation are shown in Table I. The steels were 30-lb air induction melts. They were forged to 4 by 8 by 1 in. plate at 1230°C, air cooled, and then reheated to 1230°C and cross-rolled in two passes to in. plates. Dilatometry. A Leitz Bollenrath dilatometer was used to record the transformation during continuous cooling from two different austenitizing temperatures. The dilatometer specimens were + in. in diam and 2 in. long. Oxidation and decarburization of the specimens was prevented by maintaining a small positive pressure of dry argon in the dilatometer furnace and by plating the specimens with 1 mil of Cu. For the lowest cooling rate, 17 C" per min, the temperature of the specimen was measured with a Pt-Pt 10 pct Rh thermocouple placed in a & in. diam well in the center of the specimen. During air cooling, 375 C° per min, this method of measuring the temperature interfered with the operation of the dilatometer. However, it was found that the temperature of the specimen could be measured accurately by placing a thermocouple in an identical specimen in a holder adjacent to the one being used to operate the dilatometer mechanism. The dilation-temperature curves were recorded on photographic film and then converted to volume percent ferrite-vs-temper-ature curves. The cooling rates obtained with the dilatometer are shown in Table 11. Cooling rates of

Jan 1, 1970

By E. Miller, J. M. Eldridge, K. L. Komarek

The liquidus curve of the Mg-Pb system was accurately redetermined. The compound Mg2Pb decomposes peritectically at 538.2° ± 0.3°C to liquid and to a compound p&apos; which melts congruently at 35.0 at. pct Pb and 549.0° ± 0.3°C. The solidus curve of ß&apos; was determined. X-ray diffraction studies indicate that 4&apos; has an orthorhombic structure. Activity values of magnesium calculated from the phase diagram agree with those published in the literature. EXPERIMENTAL thermodynamic properties of binary metallic systems have to be consistent with values calculated from the phase diagram. In systems forming intermetallic compounds the shape of the liquidus curve near a compound is determined by the thermodynamic properties of the coexisting solid and liquid phases. Hauffe and Wagner&apos; neglected the temperature dependence of the chemical potentials and obtained the potential differences of the components of the liquid alloys, relative to stoichiometric liquid. Their calculations were based on the liquidus curve and on the heat of fusion of the compound, and were only valid near the congruent melting point. Steiner, Miller, and Komarek2 developed equations which account for the temperature dependence and obtained the chemical potentials of liquid Mg-Sn alloys over the entire phase diagram from the liquidus and solidus curves and from enthalpy values with the pure components as the standard states. The Mg-Pb phase diagram has been studied by several investigators whose results have been compiled and critically evaluated by Hansen.3 Although the liquidus curve was poorly defined, the general features of the diagram, i.e., one congruent melting compound, Mg2Pb, of essentially stoichiometric composition, two eutectics, and limited terminal solid solubilities, seemed to be suitable for a similar thermodynamic analysis. A careful redeter-mination of the liquidus by thermal analysis revealed, however, the existence of another compound. The liquidus curve between the two eutectics was precisely delineated and the structure and solidus curve of the new compound were investigated. The revised phase diagram was thermodynamic ally analyzed to evaluate the activity of magnesium in the liquid alloys. EXPERIMENTAL PROCEDURE The magnesium metal (Dominion Magnesium Ltd., Toronto, Canada) had a purity of 99.99+ pct; lead (American Smelting and Refining Co.) contained 99.999 pct Pb. Most experiments were carried out in graphite crucibles. Several experiments were made in high-purity alumina (Triangle R.R., Mor-ganite, Inc.) and in Armco iron crucibles to test the inertness of the graphite crucibles. Chemical analysis of magnesium and detailed description of the procedure for thermal analysis have been given previously. For the determination of the solidus curve of the compounds, specimens of initial composition Mg2Pb were equilibrated in a closed isothermal system with magnesium vapor. The source of the magnesium vapor was an alloy which had a gross composition lying in the 0&apos; + L field at the temperature of equilibration. As equilibrium was approached, the specimens lost magnesium to the two-phase reservoir thereby lowering the activity of magnesium in the specimens until activity and composition equaled that of the ß&apos;/ß&apos; + L boundary. Crucibles (1.9 cm ID by 2.2 cm OD by 4.1 cm high) and tightly fitting lids were machined from a molybdenum rod; small, shallow trays were fashioned from thin (0.005 in.) molybdenum sheet, and all the molybdenum components were degreased in hot carbon tetrachloride and then dried. The pieces were then degassed in vacuum at 950°C for about 6 hr. The two-phase alloy was placed at the bottom of the crucible and small specimens of the Mg2Pb compound, weighed on an analytical balance, were placed in two molybdenum trays above the two-phase alloy. The crucible was closed by forcing its lid on and then inserted in a titanium crucible. This crucible was evacuated, flushed twice with argon, and welded under argon. The specimens were equilibrated for about 1 week in a resistance furnace regulated by a Celectray controller, and the runs were terminated by water quenching. The specimens were again weighed and the equilibrium compositions were calculated on the basis that the weight losses were solely due to a loss of magnesium to the two-phase alloy. The structure of the B&apos; phase was investigated by the Debye-Scherrer X-ray diffraction technique. Selected ingots from thermal-analysis experiments containing about 35 at. pct Pb were re-melted, slowly cooled, and crushed in an argon-filled glovebox until the entire ingot passed through a 50-mesh sieve. The powder was thoroughly

Jan 1, 1965

By W. S. Hyler, L. W. Berger, R. I. Jaffee

Hydrogen additions of 390 ppm to A-55 titanium and 368 ppm to Ti-8 pet Mn have no deleterious Hydrogenadditionseffect on the unnotched and notched rotating-beam fatigue properties of these materials. &apos;These amounts of hydrogen, however, are sufficient to cause severe notch-impact thesematerials.embrittlement in A-55 titanium and pronounced loss of tensile ductility in Ti-8 pet Mn. The lack of embrittling effect in fatigue in the latter alloy is consistent with the postulated strain-aging mechanism of hydrogen embrittlement in a-ß alloys. There is a significant strain-agingincrease in the unnotched endurance limit of A-55 titanium with the addition of hydrogen. This increase may be explained as the result of internal heating effects which would dissolve the hydride and cause solid-solution strengthening. TITANIUM and its alloys may be seriously embrittled by relatively small amounts of hydrogen. The form which this embrittlement takes has been shown to vary with alloy type. The a alloys, for example, suffer most strongly from loss of notch-bend impact toughness&apos; when sufficient hydrogen is added, and this effect has generally been associated with the presence of hydride phase in the micro-structure. In a-ß alloys, on the other hand, hydrogen is most detrimental to tensile ductility in slow-speed tests,2-1 and the embrittlement may be detected in a most convincing manner by means of rupture tests at room temperature. This particular kind of embrittlement has not been associated with a change in microstructure, but has been classified rather generally as associated with a strain-aging type of mechanism.&apos; In the present paper, the effect of an embrittling amount of hydrogen on the rotating-beam fatigue properties of both an a and an a-ß titanium alloy is covered. For this study, annealed commercially pure (A-55) titanium was chosen as an a alloy, and equilibrated and stabilized Ti-8 pet Mn as representative of a typical a-ß alloy. Nominal hydrogen levels of 20 and 400 ppm were evaluated, the latter amount having been shown previously to be severely detrimental to the impact toughness of commercially pure titanium and to cause pronounced strain-aging embrittlement in the Ti-8 pet Mn alloy. The only report of the effect of hydrogen on the fatigue properties of titanium is given by Anderson et al.,° in which a push-pull type of fatigue test was conducted on as-received commercial-purity titanium sheet. Much scatter was found in the results, but generally the presence of hydrides slightly decreased the fatigue strength of unnotched specimens in the longitudinal direction. The results of notched tests were masked too greatly by scatter to be significant. Experimental Procedure Preparation of Materials—Analyses of the A-55 titanium and the Ti-8 pet Mn alloy used in this investigation are given in Table I, which indicates the 8 pet Mn alloy to be more nearly a 6 pet Mn alloy. This alloy will be referred to as Ti-8 pet Mn, however, since this is the commercially designated composition. Both alloys were received in the form of5/8-in. diam rod and, after suitable surface preparation, 5-in. lengths were vacuum annealed at 820°C. Half of the rods for each material were then hydrogenated at 820°C to a nominal hydrogen content of 400 ppm. The hydrogenated and vacuum-annealed A-55 rods were hot swaged at 700°C from 5/8-in. diam to 1/4-in. diam, and then annealed 1 hr at 800°C and air cooled prior to preparation into test specimens. Fabrication of the Ti-8 pet Mn alloy was by hot swaging to 3/8-in. diam at 760" and then 1/4-in. diam at 704°C. This material was then annealed 1 hr at 704", followed by furnace cooling to 593"C, and finally air cooling to room temperature. Evaluation—In order to examine more completely the effects of hydrogen on the particular materials studied, slow-speed tensile and notch-bend impact properties were determined in addition to fatigue data. Tensile specimens were of the standard ASTM type with a reduced section of 1/8-in. diam and a gage length of 1/2 in. A subsize cylindrical Izod specimen was used for impact tests. These specimens had a 45" notch with a 0.005-in. radius and a 0.150-in. root diam, and the stress concentration factor of this notch in bending was Kr = 3. Both the ten-

Jan 1, 1959

By Tuncel M. Yegulalp, Malcolm T. Wane

In general, many problems relating to the exploitation of mineral deposits are probabilistic in nature. This derives from the fact that the geologic universe is inherently random. Probability theory and statistics have been found useful for forecasting the behavior of natural events that occur in the geologic universe. The objective of this paper is to illustrate the application of the theory of extremes to this fore-casting problem. For example, it is customary for design purposes to determine the rupture strength of geologic materials. The theory of extremes is exceedingly useful in describing that portion of the frequency distribution of rupture strength which contains the least strengths. Parameters describing the distribution of the least strengths are more important to the designer of mining excavations than parameters describing the total distribution. The basic principles of the theory of extremes will be detailed and illustrated. Any person required to work in the laboratory of nature is aware that uncertainty is a salient feature of all mining enterprises. A mining engineer required to plan the most efficient, practicable, profitable, and safe mine finds himself face to face with numerous ill-understood and often unquantifiable states of nature. Basic information necessary for adequate planning is often lacking or derived from incomplete tests on samples or experience of doubtful validity. The planning procedure usually takes the form of determining a feasible layout with the intent of determining an optimal layout when and if the necessary details and information become available. The crux of the entire procedure is the choosing of numbers to put into the operational and structural models which encompass the plan. Many times these numbers must be assigned qualitatively from past experiences and are called the "most probable ones." At other times, load records, performance records and material tests provide a basis for extrapolation. In any event, the numbers are chosen from a distribution or set of all numbers. Since each number in the distribution represents a possible state, the choice of any particular value is based upon a decision rule. To illustrate, consider the design of an underground structure or the design of a rock slope. The initial step is the formulation of the various possible structural actions which result from the geometry of the layout. For a given structural model various intensities of behavior are possible depending upon the load, deformation, and material characteristic spec-trums, respectively. Of particular interest to mining people is the failure behavior or condition, i.e., when there is a complete collapse of structural resistance by either structural instability or fracture. A necessary feature of the analysis is the "rupture strength" of the material. Information on the rupture strength is derived from testing either in situ or in the laboratory and the usual outcome is a variation in the test results. The methodology used to overcome this variation is to construct a frequency distribution of rupture strengths, and then determine a measure of central tendency and variability. The main idea involved is that the central tendency number will be used in the failure calculations and the measure of dispersion will be used to estimate the probability of failure. In particular if the distribution of rupture strength is normal, the mean rupture strength is the central tendency number and the standard deviation of the rupture strength is the measure of variability. Suppose the mean value of rupture strength is 1000 psi and the standard deviation is 200 psi. Insertion of 1000 psi into the failure calculation produces results that are unsafe, hence a common decision rule is to reduce the mean value by a "factor of ignorance" so that the failure calculation will produce a "safe result." If two is chosen as a factor of ignorance, this means the value inserted in the calculation is 500 psi or 2.5 times the standard deviation. The next step is to determine the percentage chance that failure will occur from a design created on this basis. Tables on the normal distribution function show that this percentage chance is 0.621% or approximately 7 times out of 1000. In practice, however, the situation is more complicated than represented by the foregoing illustration. The laboratory or field testing program usually constitutes a pathetically small sample of the geologic universe of interest and not enough testing is carried out to determine the exact form of the distribution of the test results. The normal, Cauchy and Student&apos;s T distributions are strikingly similar, and it becomes a matter of mathematical convenience to assume the normal law for phenomena which follow other laws.

Jan 1, 1969

By W. C. Hagel, H. J. Beattie

Seven austenitic alloys of varions base compositions and minor-alloy additions were solution-treated, aged systematically between 1200oand 1800oF, and examined by X-ray and electron metallography. Intragranular preczpitations of µ, Laves, s, ?&apos;, Ni3Ti, and x phases were observed as a function of composition and aging time and temperatwre. Phase solubility limits were detevtnitzed within 100Fo intervals. These inter metallic compounds fall into two distinct general classes, and whichever class predomznates depends on base composition. It has become increasingly evident that multicom-ponent austenitic alloys are well characterized by their precipitation processes. Since certain groups of elements act as one, the relationships among these processes are reasonably simple; complete identification of such processes is usually attainable by a systematic aging study with a combination of techniques centered on microscopy and diffraction. Several nickel- and cobalt-base alloys illustrating cellular precipitation and its interaction with general precipitation were reported previously.1 The group of alloys covered in the present paper demonstrates precipitation-hardening reactions involving two distinct classes of intermetallic compounds where the predominating class appears to depend on base composition. This dependency ties in with a crystal-chemistry regularity first observed some twenty years ago by Laves and Wallbaum but never amplified to our knowledge. Results of electron-microscope and X-ray diffraction studies on systematically aged hot-rolled alloys known commercially as S-816, S-590, Rene-41, Incoloy-901, M-308, and M-647 are reported here. Some of these alloys have previously undergone minor-phase analyses by other investiators. Alloy S-816 was investigated by Rosenbaum, Lane and Grant,3 and Weeton and Signorelli.4 Rosenbaum found only CbC in hot-rolled bars. Lane and Grant found CbC and a small amount of M6C in the cast structure and stated that both carbides form during aging, most of the precipitation being CbC. Weeton and Signorelli found CbC, M23C6 and a weak indication of a phase after a slow step-down cooling cycle from 2250°F. Rosenbaum also investigated hot-rolled samples of S-590 and identified CbC and M6C. Preliminary information on Rene-41, gained partly from the present work, was reported by Morris.5 Long-time precipitation phenomena in Incoloy-901 at 1350°Fwere investigated by Clark and Iwanski.B heir raw data re- semble those of our present heat with 0.1 pct B, while their interpretation of these data resembles our interpretation of data from another heat with only 0.001 pct B; they made no statement as to boron content. No previous minor-phase studies of alloys M-308 or M-647 have been reported. EXPERIMENTAL METHODS Table I gives alloy compositions in both weight and atomic percent. Specimens were solution-treated from 1700º to 2200ºF, aged at logarithmic-time intervals up to 1000 hours between 1200 and 1800 F, and examined in accordance with procedures previously described in detail. &apos; &apos; Phase extractions were carried out in electrolytic cells containing 800 ml of either 7 pct HC1 in denatured ethanol or 20 pct H3PO4 in water. After electrolysis for 48 hr at 0.1 to 0.2 amp per sq inch, residues were separated by filtration or centrifuging. X-ray powder patterns of residues were recorded on a diffractometer for accuracy and on film for sensitivity. Lattice parameters were calculated by least-squares analyses of indexed sin 8 values, and relative abundances were estimated from intensities of strongest lines of each phase. These phase abundances denote relative amounts with respect to each other rather than to the alloy. Mechanically polished specimens were etched in a freshly mixed solution of 92 pct HC1, 5 pct H2SO4, and 3 pct HNO3. Parlodion replicas for the electron microscope were chromium-shadowed in high vacuum at a glancing angle of 20deg. All electron micrographs are reproduced here with the shadowing source above. The correspondence betweenelectronmicrostructures and phases identified by X-rays was established by a high redundancy of correlation between relative amounts at different stages of aging and examination above and below critical transformation or solubility temperatures. EXPERIMENTAL RESULTS S-816 and S-590—The phases found in S-816 and S-590 after various aging and solutioning treatments are listed in Table 11. These data and the observed

Jan 1, 1962

By F. C. Langenberg, J. Chipman

New data on the distribution of silicon between slag and carbon-saturated iron at 1600oand 1700oC are presented which, in combination with previously published data, permit the determination of silica activities over a broad range of compositions in the CaO-Al2O3-SiO2 system. The distribution of silicon between graphite-saturated Fe-Si-C alloys and blast furnace-type slags in equilibrium with CO has been described in previous publications.1"3 In this past work the silica-silicon relation was established at temperatures of 1425" to 1700°C for slags containing up to 20 pct Al2O3. This paper presents the results of additional studies at 1600" and 1700° C which extend the silicon distribution data at these temperatures for CaO-A1203-SiO2 slags over a range from zero pct A12O3 to saturation with A12O3, or CaO.2A12O3. The upper limit of SiO, is set by the occurrence of Sic as a stable phase when the metal contains 23.0 or 23.7 pct Si at 1600" or 1700°C, respectively. The activity of silica over the expanded range is determined directly from the distribution data.3 Recently, 4-7 other investigators have studied the activities of SiO, and CaO, principally in the binary system, using different methods and obtaining somewhat different results. EXPERIMENTAL STUDY The experimental apparatus and procedure have been fully described in previous publications.1, 3 Six new series of experimental heats have been made, four at 1600° and two at 1700°C. Master slags of several fixed CaO/A12O3 ratios were pre-melted in graphite crucibles, and these were used with additions of silica to prepare the initial slag for each experiment. Slag and metal were stirred at 100 rpm and CO was passed through the furnace at 150 cc per min. The initial sample was taken 1 hr after addition of slag at 1600°C or 1/2 hr after addition at 1700°C. The run was normally continued for 8 hr at 1600°C or 7 hr at 1700°C, and the final sample was taken at the end of this period. Changes in Si and SiO2 content indicate the direction of approach to equilibrium, and in a series of runs where the approach is from both sides this permits approximate location of the equilibrium line. Fig. 1 shows the results of such a series of 15 runs at 1600°C for slags of CaO/Al2O3 = 1.50 by weight. Figs. 2 and 3 record other series at 1600°C and Fig. 5 a series at 1700°C with fixed CaO/Al2O3 ratios. The results of the experiments at 162003°C have been reported in part in a preliminary note.3 In the experiments recorded in Figs. 4 and 6, the slags were saturated with A12O3 (or with CaO.2A12O3 within its field of stability) by suspending a pure alumina tube in the melt during the course of the run. The final slag analyses were used to establish the liquidus boundaries8 in the stability fields of CaO.2Al,O3 and of A120,. ACTIVITY OF SILICA The free-energy change in the reaction has been calculated by Fulton and chipman2 from recent and trustworthy data including heats of formation, entropies, and heat capacities. The more recent determination by Olette of the high-temperature enthalpy of liquid silicon is in satisfactory agreement with the values used and therefore requires no revision of the result which is expressed in the equation: SiO, (crist) + 2C (graph) = Si + 2CO(g.) [1] &F° = + 161,500 - 87.4T The standard state for silica is taken as pure cristobalite and that of Si as the pure liquid metal. Since the melts were made under 1 atm of CO and were graphite-saturated, the equilibrium constant for Eq. [I] reduces to K1 = asi /asio2 The value of this constant is 1.77 at 1600°C and 16.2 at 1700°C. Through K1, the activity of silica in the slag is directly related to the activity of silicon in the equilibrium metal.

Jan 1, 1960

By R. D. Nelson, J. C. Shyne

The ß and a ß transformations in plutonium were studied with particular emphasis on the transformation kinetics and microstructure. Significant observations are: 1) The kinetic data show conclusively that the ß — a transformation in high-purity plutonium can proceed isothermally with no athermal component. 2) Plastic deformation of the stable (3 phase retards the subsequent (3 — a transformation. 3) Plastic deformation of the stable a phase accelerates the a — ß transformation; the acceleration is attributed only to residual stresses. 4) The a and a?a volume changes occur anisotroPically in textured plutonium. 5) An apparent crystallogvaphic relationship exists between the parent and the product phases of the and (3 — a transformations. 6) Both applied uniaxial compressive stresses and uniaxial tensile stresses raise the starting temperature for the ß — a transformation. 7) A given uniaxial tensile stress favors the a — ß transformation more than an equivalent applied uniaxial compressive stress opposes the transformation. These observations of the (ß —a and a — ß phase changes in plutonium are consistent with known mar-tensitic transformations. ThIS paper elucidates some of the characteristics of the a— ß and ß —a transformations in plutonium. Because considerable conjecture exists about the mechanisms by which the phase transformations occur in plutonium, experiments have been performed to provide indirect information concerning the mechanisms responsible for the a —ß and ß -a transformations. Indirect information is of particular value in the study of plutonium because of the experimental difficulties presented by the metal. Single crystals have not been produced in any of the allotropes. The large density results in high X-ray and electron-absorption factors and consequently complicating X-ray and electron diffraction. The kinetics of ß — a and a — ß transformations of plutonium and the behavior of the transformations under a variety of conditions have been investigated in detail. Information about the mechanisms of the allo-tropic transformations of plutonium was obtained indirectly from three sources: 1) the effect of plastic deformation of the stable parent phase upon the transformation kinetics; 2) the behavior of the metal transforming under applied stresses; and 3) the microstruc-tural and crystallographic features between parent and product phases. PHASE-TRANSFORMATION CHARACTERISTICS In characterizing solid-state phase transformations in metals and alloys, it is useful to define several types of transformations. An aim of the present work was to identify the low-temperature transformations in plutonium by type, i.e., as martensitic or nonmar-tensitic. Practical definitions for these terms follow. The terms commonly used to categorize phase transformations lack universally accepted definitions. This confusion arises doubtlessly because some terms specify crystallographic or morphological character while other words have a kinetic or a thermodynamic connotation. For example, martensitic specifies certain definite crystallographic restrictions. Unfortunately, martensitic is sometimes used in an ill-defined way to imply kinetic characteristics. Further confusion attends the use of such expressions as nucleation and growth, diffusional, and massive. From time to time new systems of phase-transformation nomenclature are suggested; unfortunately none of these has gained general acceptance.1,2 The authors of the present paper have no intention of entering the controversy. We recognize that some readers may object to the nomencliture used here. For exampie, the terms military and civilian have recently been used in much the same way as martensitic and non-martensitic are used in this paper. This paper is intended to describe several specific details of the low-temperature phase transformations in plutonium. The authors have found it useful to identify these transformations as martensitic; the term was chosen as the best available to describe the experimentally observed features of the transformations studied. A martensitic transformation is one that occurs by the cooperative movement of many atoms; the rearrangement of atoms from parent to product crystal structure occurs by the passage of a mobile semico-herent growth interface. The geometric features characteristic of a martensitic transformation are a specific orientation relationship between the product and parent phase lattices, a specific habit-plane orientation for the growth interface, and a shape change with a specifically oriented shear component. There is no alloy partition between the parent and product phases in a martensitic transformation. Martensitic transformations may display either athermal kinetic behavior or thermally activated isothermal kinetic behavior. Some martensitic transformations occur isothermally, although more commonly martensitic transformations are athermal. Isothermal martensitic transformations are suppressible by rapid cooling. In athermal martensitic transformations, nucleation and growth are not thermally activated and the transformations are essentially time-independent. Nucleation, growth, or both can be thermally activated in isothermal martensitic reactions. Transformation of the parent phase into a marten-

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 H. S. Chiang, D. F. Lu, S. S. Peng

INTRODUCTION The most important element in a successful lingual mining is a good roof control. The modern longwall mining employs hydraulic powered supports for roof control at the face area. The application of hydrau¬lic powered support requires the knowledge of over¬burden strata behavior for proper selection of sup¬port type and capacity. Failure to do so could lead so serious loss. There are several methods available for determining the required support capacity (1-3). While these methods are simple for application, they do not include the complicated roof behavior observed in longwall mining. As research progresses and operational experience accumulates (4,5), the concept about the designing and selection of powered support improves. The design of a longwall powered support consists of three major phases: 1. structural integrity and stability of the powered support, 2. external loadings induced by the movements of the overburden strata, and 3. interaction between the support, roof and floor. Phase 1 involves structural analysis (5) and full-sized testing (6) of the supports. Its validity is limited by the accuracy of the assumed external loading because of the uncertainty about the actual loading underground. The third phase includes the reaction of the support and the floor to the movements of the overburden strata and vice versa. Among the three phases, the second phase concer¬ning the external loading seems to be the least known because of the complicated behavior of the roof strata. There are many unresolved problems. For example, does the main roof break periodically and cause periodic roof weighting in the face area? If so, are there any rules governing its behavior? How does the roof load on the support canopy! Finally, how can one determine the required support capacity and select a proper type of support to meet a certain roof behavior? In order to answer those questions, underground instrumentation and observations were performed at 4 longwall panels in 3 separate mines for the past two years. This paper summarizes the current findings. PANEL LAYOUTS AND EQUIPMENT EMPLOYED The three mines selected are all located in West Virginia; two in northern and one in southern West Virginia. As shown in Table 1, seam conditions (i.e. seam, depth and thickness) and panel layouts are different among the three mines. The most significant difference in equipment is the face powered supports. Three mines used three different types of shield; 2-leg caliper, 2-leg lemniscate, and 4-leg lemniscate chock-shield. (Fig. 1) UNDERGROUND INSTRUMENTATION AND OBSERVATION PROGRAM Two events were instrumented in each observed longwall face: one was the hydraulic pressure (resistance) of the powered supports and the other was the canopy load distributions. In addition, the gob caving conditions were visually observed and recorded. Leg and Support Resistances One or two automatic Weksler Pressure Recorders were installed at the designated shield support,. In most cases, the daily charts were used to record the pressure variations in both the front or the rear legs (for the 4-leg shield), or in both the leg and the fore-pole ram (for the 2-leg shield). The recorded pressure w a s then converted to load or resistance by multiplying it by the cross-sectional area of the hydraulic leg or canopy ram piston. Fig. 2 shows the typical pressure-recorded charts for the 4-leg and 2-leg shields in a 23-24 hour period. The support resistance is the summation of the resistance in each of all the legs for that support. Generally, the resistance of the fore-pole ram will not be considered in determining the capacity of the support because of its rather small vertical compo¬nent force at the tip of the fore-pole. Canopy Load Distribution External load distribution on the canopy as exer¬ted by the roof was monitored. The measurements employed 12-14 pieces of pressure cells (6-inch square) that were uniformly arranged in two rows on the canopy. After support setting, the pressure changes in the cells were monitored at various stages of the mining (supporting) cycle while the support leg pressures were recorded continuously by the pressure recorders. Based on the calibration chara¬cteristics of each pressure cell as performed in the laboratory before and after each underground test, the cell pressures were converted to actual loadings. From these load measurements the canopy load distri¬butions and the relations between measured canopy loadings and support leg resistances were determined. Accordingly, the supporting efficiency of the shield support can be determined.

Jan 1, 1982

By Charles F. Hickey, Eric B. Kula

Thermomechanical treatments applied to the maraging steels include a) cold working in the austenitic condition at 650°F, followed by transformation to martensite and aging, b) cold working in the murtensitic condition and aging, and c) cold working in the aged condition with and without subsequent reaging. The strength increases in these steels are very small compared to the increases observed in conventional carbon and alloy steels. The changes that are observed are compatible with the strengthening mechanisms operative during thermomechanical treatment of conventional steels, however. Differences are caused by the absence of a carbide precipitate and the low work-hardening rate in both the solution-treated and the aged conditions. ThE 18 pct Ni maraging steels represent a class of steels which are finding great interest for high-strength applications.1~2 They are essentially carbon-free, and contain 7 to 9 pct Co, 3 to 5 pct Mo, and 0.2 to 0.8 pct Ti. Although austenitic at elevated temperatures, they can be air-cooled to room temperature to form a martensite, which because of the absence of carbon is relatively soft. On subsequent reheating age hardening occurs and strength levels of 250 to 300 ksi yield strength can be attained. These steels appear to be particularly suitable for studying the response to various thermome-chanical treatments for additional reasons other than the obvious one of attempting to improve their already attractive properties. Thermomechanical treatments can be defined as treatments whereby plastic deformation, generally below the recrystal-lization temperature, is introduced into the heat-treatment cycle of a steel in order to improve the properties. With an absence of intermediate transformation products on air cooling the maraging steels have good hardenability and hence can readily be cold-worked in the austenitic condition prior to transformation to martensite. Further, they can be worked in the martensitic condition prior to aging, and even can be deformed in the fully aged condition. Finally, it is of interest to compare their re- sponse to that of the more conventional alloy and carbon steels, where the role of carbides is important in the strength increase by thermomechani-cal treatments. The thermomechanical treatment of conventional steels has been the subject of a recent review.&apos; I) MATERIALS AND PROCEDURE The steel used in this investigation was a commercially produced vacuum-melt heat, which had been rolled to 0.090 in. and mill-annealed. The composition of the alloy was as follows: 0.02 C, 0.08 Mn, 0.10 Si, 0.009 P, 0.009 S, 18.96 Ni, 7.34 Co, 5.04 Mo, 0.29 Ti, 0.05 Al, 0.004 B, 0.01 Zr, and 0.05 Ca. Unless otherwise stated the heat treatments used were the standai-d solution treatment at 1500°F for 1 hr, air cool, followed by a 900°F, 3 hr age. In this condition, the material exhibited 232 ksi yield strength and 239 ksi tensile strength. Mechanical properties were determined by Vicker&apos;s hardness measurement (20 kg) and by tensile tests on standard 1/2-in.-wide, 2-in.-gage-length sheet tensile specimens. Notch tensile tests were run using the 1-in.-wide NASA type, edge-notched specimen.4 Fracture-toughness determinations were made on 3-in.-wide, center-notched, fatigue-cracked specimens, following the recommendations of the ASTM Committee on Fracture-Toughness Testing.4 An electric-potential technique was used for measuring the crack size at the onset of rapid crack propagation5 which is necessary for calculations of Kc, the critical stress-intensity factor under plane-stress conditions. The critical stress-intensity factor under plane-strain conditions KI, was also calculated, using the stress at which the first observable crack growth occurred. 11) RESULTS A) Cold-Worked in the Austenitic Condition. The reported M, temperature for the 18 pct Ni maraging steel is about 310°F.1 Therefore, a temperature of 650°F was selected as suitable for rolling in the austenitic condition. Specimens were solution-treated at 1500°F for 1 hr, air-cooled to 650°F, and rolled varying percentages from 0 to 60 pct, at 20 pct reduction per pass. Tensile and hardness properties after aging at 900°F for 3 hr are shown in Fig. 1. The tensile strength increases from 253 to 271 ksi and the yield strength from 247 to 265 ksi as a result of a reduc-

Jan 1, 1964

By R. G. Agarwal, Ramey Jr. H. J., Al-Hussainy R.

Although it has long been realized that gas recovery from a water-drive gas reservoir may be poor because of high residual saturations under water drive, it appears that only limited infomlation on the subject has been available until recently. This study was performed to show the qiiantitative potential importance of water influx. Results indicate that gas recovery may be very low in some cases: perhaps as low as 45 per cent of the initial gas in place. Gas recovery under water drive appear to depend in an important was on: (I) the prodirction rate and manner of production; (2) the residual gas saturation; (3) aquifer propertie.); and (4) the volumetric displacement effciency of water invading the gas reservoir. The manner of estimating water-drive gas reservoir recovery can vary considerably. Examples are: the steady-state tnethorl. the Hurst modified steady-state method, and various unsteady-state methods such ac. those of van Ever-dingen-Hurst, Hurst, and Carter-Tracy. The Carter-Tracy water influx expression was used in this study. In certain cases, it appears that gas recovery can be increased significantly by controlling the production rate and manner of production. For this reason, the potential importance of water influx in particular gas reservoirs should he investigated early to permit adequtrtr planning lo optirtize the pay reserves. INTRODUCTION In recent years, the economic importance of natural gas production has become increasingly apparent. This has been evidenced by more intensive exploration efforts aimed at gas production, and exploitation of both deep, as well as low-permeability gas reservoirs. Technical developments such as deep-penetration fracturing have made development of such formations economically feasible. Unfortunately, water influx has forced abandonment of a number of gas reservoirs at extraordinarily high pressures. Although reservoir engineering methods for estimating water influx have long been available, it appears that application of these methods to the water-drive gas reservoir has been sporadic.&apos;a Available methods for estimating water influx which can be applied to the water-drive gas reservoir problem include the steady-state method,1 the Hurst modified steady-state method and various unsteady-state methods such as those of van Everdingen-Hurst. Hurst, and Carter-Tracy. Interesting applications of these solu- tions to gas reservoir and the aquifer gas-storage problems have appeared recently.3,12,14 The experimental study of residual gas saturations under water drive by Geffen et al. in 1952 indicated that residual gas saturations could be extremely high. A value of 35 per cent of pore volume is often used in field practice when specific information is not available. The study of Geffen et al. showed that residual gas saturation might be much higher in some cases. Naar and Henderson concluded that the residual non-wetting phase saturation under imbibition should be about half of the initial non-wetting phase saturation. The Naar and Henderson result that residual gas saturation under water influx should be about half the original gas saturation is recommended as an estimate if laboratory measurements are not available. Thus, it is clear that a considerable portion of the initial gas in place might be trapped in a water-drive gas reservoir as residual gas at high pressure. A full water-drive would result in loss of residual gas trapped at initial reser.voir pressure. Consideration of transient aquifer behavior leads to the conclusion that high-rate production of water-drive gas reservoirs could result in improved gas recovery by reduction of the abandonment pressure. However, there appears to be little quantitative information on this possibility. One of the few advantages of water-drive gas production appears to be improved deliverability through water-drive support of the reservoir pressure. There may also be an advantage in higher condensate recovery caused by pressure maintenance for gas-condensate water-drive reservoirs. In view of the preceding, this study was made to assess the potential importance of water-drive in gas reservoir engineering. The Carter-Tracy approximate water-influx expression was used because this equation offers some advantages in hand-calculation which do not appear to have been generally recognized.&apos; However, calculations were performed in the main with a high-speed digital computer to permit evaluation of the effect of water-drive under a large variety of conditions. CALCULATION METHOD Water-drive gas reservoir performance can be estimated in a manner completely analogous to oil reservoir calculations: a materials balance is written for the reservoir, and a water influx equation is written for the aquifer. Siniltaneous solution provides the cun~ulative water influx and the reservoir pressure. When reservoir performance data (gas produced and reservoir pressures) are available, it is usually possible to match performance data to determine the initial gas in place and the water influx parameters —

Jan 1, 1966

By J. H. Moran, E. E. Finklea

The pressure build-up technique is a recognized method of determining permeability from conventional drillstem tests. In this paper an effort is made to extend such techniques to the interpretation of data obtained from the wireline formation tester. Such a study is necessary because of the differences, for this case, in the magnitude of the flow parameters (rate of flow, amount of recovered fluids) and in the flow geometry (flow through a perforation vs flow across the face of the wellbore, etc.) involved in the solution of the equations of flow for compressible fluids. The perforation is replaced by a spherical hole, and the effect of the borehole is neglected, so that the flow can be considered to be radial in a spherical co-ordinate system. Arguments are presented to justify this idealization. Assuming single-phase flow, general relations between pressure and flow rate are developed for a homogeneous medium. The study is then extended to permeable beds of finite thickness. It is shown that the early stages of pressure build-up tend towards spherical flow, while the later stages tend towards cylindrical flow. The thinner the bed, the more quickly flow approaches the cylindrical model. The prevalence of thin beds in practical work makes this analysis quite important. Cases involving permeability anisotropy are treated. INTRODUCTION From wireline formation tester operation, two types of data are obtained: (1) the nature and amount of recovered fluids, and (2) the pressure history recorded during the test. A number of papers have been written dealing with the interpretation of formation production on the basis of the recovered fluids.&apos;.&apos; In general, the methods described have been quite accurate for both high- and low-permeability formations. The present paper will deal with an analysis of the pressures observed. An analysis of the pressure build-up curves obtained in hard-rock country has already been attempted on the basis of the formula proposed by Hor-ner. Although this approach has met with success in many instances, some questions have been raised as to its validity. It is the aim of the present study to place the analysis of pressure build-up in the formation tester on a firmer basis, from which more detailed methods of interpretation can evolve. Because of the great differences between the operation of the wireline formation tester and the conventional drillstem test, modifications are necessary in the interpretation. The major difference relates to the flow geometry. Once the flow geometry has been established other features such as multiphase flow, skin effect, afterflow, etc., well described in the literature, can be introduced. It will be assumed that the mechanical operation of the formation tester is already known to the reader.6 t will suffice here merely to state that the tester provides the means for taking a relatively small sample of the fluid immediately adjacent to the borehole, and for recording the subsequent pressure response. In comparison with conventional drillstem tests, the time required for a satisfactory pressure build-up response is much shorter, because of the relatively small quantity of fluid withdrawn by the wireline tester. This feature is highly desirable in the case of low-permeability formations. For an analysis of the pressure response within the formation, three simple flow geometries are considered— linear, cylindrical and spherical. The spherical and cylindrical flow geometries are most pertinent to the formation tester; therefore, they will receive the major emphasis. Since the configuration of the borehole and the perforation made by the tester complicate the flow geometry, it is necessary to allow for them in the drawdown response. However, because of the volume of formations contributing to the pressure-response, the details of the perforation shape are unimportant in the build-up period. Since relatively small amounts of fluid are withdrawn from the formation, in contrast to a conventional drill-stem test, a study of the "depth of investigation" and the significance of drawdown as well as build-up data will be included. Because the "depth of investigation" will be shown to be rather large, the effect on the build-up curves of the finite thickness of the permeable bed is considered. It is this consideration that leads to the importance of cylindrical flow geometry. Also included is a discussion of permeability anisotropy and its effect on the interpretation of the tester results. The pressure curves recorded by the formation tester will follow two general patterns, depending upon whether the formation is of high or low permeability. Fig. I (a and b) schematically illustrates these two responses. In Fig. 1(a), the high pressure recorded during fill-up of the tool is essentially the pressure differential across the choke in the system. In Fig. l(b), the flow rate is

By N. A. Gokcen, John Chipman

SILICON is the most commonly used deoxidizer and an important alloying element in steelmak-ing; hence a detailed study of this element in liquid iron containing oxygen is of considerable interest. The equilibrium between silicon and oxygen in liquid iron has been studied by a number of investigators but generally with inconclusive or incomplete results. The variation of the activity coefficients of silicon and oxygen with composition is entirely unknown. Published investigations deal with the reaction of dissolved oxygen with silicon in liquid iron and the results are expressed in terms of a deoxidation product. For consistency and convenience in comparison of the published information, the deoxidation product as referred to the following reaction is expressed in terms of the percentage by weight of silicon and oxygen in the melt in equilibrium with solid silica: SiO (s) = Si + 2 O; K&apos;l = [% Si] [% 012 [I] Theoretical attempts to calculate the deoxidation constant for silicon in liquid iron from the free energies of various reactions yielded results which were invariably lower than the experimental values. Thus, the deoxidation "constants" calculated by McCance,1,2 Feild,3 Schenck, and Chipman were of the order of 10, which is below the experimental values by a factor of more than 10. Experiments of Herty and coworkers" in the laboratory and steel plant resulted in an average deoxidation constant of 0.82x10 &apos; at about 1600°C. The technique employed in their investigation was crude and the reported temperature was quite uncertain. The concentration of silicon was obtained by subtracting silicon in the inclusions from the total. Since at least some of the inclusions resulting from chilling must represent a fraction of the silicon in solution at high temperatures, such a subtraction is not justifiable. Results of Schenck4 for K&apos;1 from acid open-hearth plant data yielded a value of 2.8x10-5, which was later revised as 1.24x10 at 1600°C. Similarly Schenck and Bruggemann7 obtained 1.76x10-5 at 1600OC. The discrepancies and errors involved in the acid open-hearth plant data as compared with the results of more reliable laboratory techniques were attributed by these authors to the lack of equilibrium and the impurities in liquid metal and slag, and are sufficiently discussed elsewhere." Korber and Oelsen" investigated the relation between dissolved oxygen and silicon in liquid iron covered with silica-saturated slags containing varying concentrations of MnO and FeO. The deoxidation products obtained by their method scatter considerably, and their chosen average values of 1.34x10, 3.6x10-5, and 10.6x10-5 1550°, 1600°, and 1650°C, respectively, represent the best experimental results which were available until quite recently. Darken&apos;s10 plant data from a steel bath agree approximately with their data at 1575° to 1625°C. Zapffe and Sims" investigated the reaction of H2O and H2 with liquid iron containing less than 1 pct Si and obtained deoxidation products varying by a factor of more than 20. Inadequate gas-metal contact and lack of stirring in the metal bath should require a longer period of time than the 1 to 5.5 hr which they allowed for the attainment of equilibrium. Furthermore, their oxygen analyses were incomplete and irregular and confined to a few unsatisfactory preliminary samples. Their results did indeed indicate that the activity coefficient of oxygen is decreased by the presence of silicon, although they made no such simple statement. They chose to attempt to account for their anomalous data by the unlikely hypothesis that SiO is dissolved in the melt. Hilty and Crafts" investigated the reaction of liquid iron with acid slags under an atmosphere of argon, making careful determinations of silicon and oxygen contents at several temperatures. Despite erroneous interpretation of the data at very low silicon concentrations, their data represent the most dependable information on this equilibrium that has been published. In the range 0.1 to 1.0 pct Si, their data yield the following values for the deoxidation product: 1.6x10-5, 3.0x10- &apos;, and 5.3x10 at 1550°, 1600°, and 1650°C, respectively. The purpose of the work described herein was to study the equilibrium represented by eq 1 as well as the following reactions, all in the presence of solid silica: SiO2 (s) + 2H2 (g) = Si + 2H2O (g);

Jan 1, 1953

By D. V. Edmonds, C. J. Beevers

Tensile tests were carried out on a! titanium containing 850, 1250, and 2700 ppm 0, and up to -500 ppm H. The tests were performed at -196", -78", 20°, 150°, and 300°C at a strain rate of -1.0 x 10??3 sec-1. Increasing oxygen content, increasing grain size, and decreasing test temperature resulted in enhanced embrittlement of the a titanium by the hydrogen additions. Metallographic observations showed that this can be correlated with the influence of these parameters on the introduction of cracks into the a! titanium by fracture of titanium hydride precipitates. CRAIGHEAD et al.1 reported that the hydrogen content normally found in commercial-purity a! titanium (60 to 100 ppm) was sufficient to cause a substantial lowering of the impact strength, and they attributed this embrittling effect of hydrogen to the precipitation of titanium hydride. Lenning et al.&apos; found that in commercial-purity a titanium there is an almost complete loss of impact strength at about 200 pprn H, which is approximately half the value needed to eliminate the impact strength of high-purity a titanium. They also showed that the presence of 3000 ppm hydrogen reduces the room-temperature tensile ductility of commercial-purity material to a value of the order of 10 pct; the corresponding hydrogen concentration for high-purity titanium is over 9000 ppm. It thus appears that the detrimental effect of hydrogen on the mechanical properties of commercial-purity titanium becomes evident at much lower hydrogen contents than for high-purity titanium. The main difference between the two types of a titanium might be expected to be the higher level of interstitial impurity in the commercial-purity grade. Jaffee et a1.3 studied the influence of temperature and strain rate on the hydrogen embrittlement of high-purity and commercial-purity ! titanium. In general, the behavior was the same for both materials; embrittlement was enhanced by decreasing temperature and increasing strain rate. Recent results from tests on commercial-purity a titanium containing 850 ppm O and varying amounts of hydrogen up to -500 ppm showed that the degree of embrittlement by hydrogen is intimately related to the fracture characteristics of titanium hydride precipitates.4 The present paper considers the interrelationship between the mechanical properties and micro-structural features of commercial-purity a! titanium containing 850, 1250, and 2700 ppm 0 and varying amounts of hydrogen up to -500 ppm. 1. EXPERIMENTAL PROCEDURE Three types of commercial-purity titanium supplied by IMI* were used in the investigation, and for the *Address: Witton, Birmingham 6, United Kingdom. purpose of this paper are designated Ti 115, Ti 130, and Ti 160. The principal impurity elements are given in Table I. The material was received in the form of 12.7 mm diam bars having a fully recrystallized structure. Tensile specimens with a round cross-section of 4.5 mm diam and a gage length of 15.2 mm were machined from the bars. In order to develop the same grain size (mean linear intercept of grain boundaries) in each of the three types the specimens were annealed under a dynamic vacuum of <10?5 mm Hg, Table 11. Specimen hydriding was carried out in a modified Sieverts apparatus;&apos; hydrogen was taken into solution at 450°C and after holding the specimens at this temperature for 24 hr they were furnace-cooled to room temperature at an average rate of -100 C deg per hr. By this method nominal hydrogen contents of 0, 50, 100, 250, and 500 ppm were introduced into specimens of Ti 115, Ti 130, and Ti 160 (100 ppm (wt) -0.5 at. pct). The actual hydrogen contents were calculated from the weight differences obtained by weighing the specimens before and after the hydriding treatment. Tensile tests were carried out at temperatures of -196", -78", 20°, 150°, and 300°C on a 10,000 kg In-stron machine at a nominal strain rate of -1.0 x 10-3 sec-1. Fractured specimens were sectioned in planes parallel to the tensile axis, mechanically polished to 0.25 µm grade of diamond paste, and then attack polished using a solution containing by volume 99 parts H2O, 1 part HF, and 1 part HNO3. Although the latter treatment unavoidably opened out cracks and voids visible after mechanical polishing, it did reveal the grain structure, titanium hydride morphology, and deformation twinning structure.

Jan 1, 1970