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By I. R. Kramer

The apparent activation energy, U, JOY plastic deformation of aluminium, copper, and gold single crystals and polycrystalline aluminum was found to he dependent upon surface condition. At low strains, the value for U for aluminum single crystals was 4200 cal per mole ; however, this value decreased to 920 cal per mole when the specimens were polished during the test at a rate of 50 X 10-5 win-1 The activated volume also was found to increase with the rate of metal removal. That the surface plays an important role in plastic-deformation processes has been adequately demonstrated.1-4 Of particular interest at this time are the observations4 that the creep rate of single crystals of fee metals is increased markedly when specimens are polished during the deformation process. These observations were interpreted in terms of a dislocation debris layer at the surface regions of the crystal. This layer, which appears to be a region containing a high concentration of dislocations or segments of dislocations, not only will impede the egress of dislocations on the primary slip system, but will also provide a barrier against which dislocations may pile up. By this means, a stress which opposes the motion of dislocations will be generated. It is considered that this stress must be taken into account in any theory which describes the plastic-deformation process. Although a detailed dislocation model cannot be given at this time, it may be stated that the stresses opposing the motion of dislocations arise from the dislocations in the debris layer, Ts, and ti due to the resistance to motion imposed by internal obstacles such as Lomer-Cottrell barriers, subgrain and grain boundaries, forest dislocations, and so forth, as well as the resistance due to the presence of jogs on dislocations and the interaction of the stress fields of dislocations, and so forth. In essence, we have added an additional term, ts, to the usual form of the equation T = Ta- Tb, where Tb is the back stress due to the various types of obstacles which impede the motion of dislocation. Accordingly, the net stress, t, acting on a dislocation may be written as where Ta is the applied shear stress. For deformation which occurs by a thermally activated process5,8 ?=Ae-u/kT = A e(u,-VT)/kT [2] and U = Uo- Vt PI where U = activation energy for plastic flow, y = shear strain rate, Uo = activated energy for plastic deformation in the absence of stress and depends upon the particular deformation process, V = activated volume, k = Boltzmann constant, T = temperature, OK, and A = frequency factor. From Eqs. [I] and [3], U = Uo- V(Ta - Ti - TS) [4] and it is seen that if Ts decreases, the activation energy, U, must decrease if V remains constant. However, as will be shown, V is a function of the surface conditions. In their work on aluminum, Lytton and coworkers7 reported that at low strains the activation energy was 3400 cal per mole and at high strains it was 28,000 cal per mole. The lower value was interpreted in terms of the energy necessary to overcome the Peierls force, while the higher value was associated with the energy necessary for dislocations to move by a cross-slip process. It will be shown that the lower activation energy is associated with the surface debris layer rather than with the Peierls force. EXPERIMENTAL PROCEDURE Single crystals having a nominal cross section of 1/8 in. and a 3-in. gage length were prepared by a modified Bridgman technique using a multiple-cavity graphite mold. The initial purity of the metals was 99.997, 99.999, and 99.999 pet for aluminum, copper, and gold, respectively. The orientation of the specimens is given in Fig. 1. The measurement of the activation energy as a

Jan 1, 1964

By J. B. Cheatham, C. E. Murphey

Presently, computer control of Borobolic direction cannot be obtained during drilling, and most straight-holc drilling methods attempt to resist hole deviation rather than control direction. Many of the theories advanced as possible explanations of the cause of hole deviation are Summarized berein. A new correlation of physical partables is introduced to indicate bow factors such iis drill collar stiffness clearance and hit weight influence borehole deviation, .A methord is proposed for predicting the rate of change of hole angle when drilling conditions are changed. INTRODUCTION Control of borehole direction during drilling can be difficult and costly. Unintentional crooked holes are often lrilleti in dipping formations and many times directional drilling is required when the surface location is not directly above the target area —- for example, at offshore and mountainous locations. Prilllng progress can be greatly hindered in either air or liquid drilling when it be comes necessary to use low bit weight to prevent excessive hole angle build-up. If hole inclination becomes too great, drill pipe drag becomes excessive and fishing risks are increased, logging is more difficult and problems of differential sticking, key seating and fatigue failures may be encountered. Dog-legs and key seats are more serious problems than steep inclination angles: therefore, reducing rate of direction change is preferred to attempting to maintain absolutely vertical holes. Consequently, a straight inclineti hole is preferable to a nearly vertical crooked hole containing numerous dog-legs. In this paper, theories of the cause of hole deviation and analyses of drill string behavior under down-hole conclitions are summarized. Methods for computing hole deviation are presented und systems for resisting deviation as well as neans for provic!ing control of holt- direction are iliscusseI. A new correlation of physical variables is introduced tu indicate how factors such as drill collar stiffness, clearance and bit weight influence borehole deviation. A method is oroposeti for predicting the rate of change of hole anFrle when drilling conditions are clianged. IIEVIEU' OF PREVIOUS WORK ON HOIaE DEVIATION Significant progress in the theoretical analysis of hole deviation problems has been made in the past 1 5 years. The pioneering work has been primarily a result of the efforts of Lubinski and Woods.l-5 In 1950, Lubinski 1 considered the buckling of a drill string in a straig:rt vertical hole, a problem also considered by Willers6 in 1941. It wns concluded that very low bit weights must be used to prevent hole deviation resulting from drill collar buckling. The use of conventional stabilizers was proposed2 in 1951 by Mac Donald and Lubinski as a method for permitting greater Sit weights to he carried without drill collar buckling. These authors pointed out that a 2° nearly vertical spiral hole can cause severe key seatinp, and drill pine near, wilereas n 3 straight inclined hole with deviation all in one direction, while not vertical, will not result in serious drillirig or producing problems. Studies were continued with an investigation of straight inclined holes by Lubinski and Woods3 in 1953. In this paper they concluded that perfectlv vertical holes cannot be drilled even in isotropic forriations unless extremely low bit P-eights are used. .l'l~ey postulated that constant drilling conditions produce holes of constant inclination angle and varying conditions cause the hole to drill at a neiv equilibrium angle. This analysis was not concerned with drill string buckling since it was based on an equilibrium solution in which the ~irill string Was presumed to lie along the lower side of the hole abcrve the point of tangency. Weight of the drill collars below the point of tangency tends to force the hole toyard the vertical, whereas the weight on hit tends to force the hole aurav from the vertica l. The concept of an anisotropic formation Lvas introduced as an empirical method for explaining actual drilling data and as a means for extrapolating known deviation data to otiier conditions of hit

Jan 1, 1967

By E. A. Gulbransen, A. Pebler

Pressure-composition isotherms have been determinedfor the systems ZrMo2 -H2 between 0" and 900°C at hydrogen pressures between 10-4 and 760 Torr. Tkese studies plus X-ray diffraction analyses of selected compositions show that these intermetallic compounds form a complete series of hydrogen solid solutions within this temperature and pressure range. The partial molar enthalpies and entropies of solution of hydrogen in the three intermetallic compounds were calculated. The solubility of hydrogen in these Laves phases can be related to their free electron concentration. ThE reaction of hydrogen with inter metallic compounds of zirconium with Laves phase-type structure has both theoretical and practical interest. The in-termetallic compounds have structures different from any of the pure metals. In addition the electron to atom ratio of the intermetallics varies appreciably. In the nuclear materials field, these materials are of interest since intermetallic compounds may exist as precipitates in commercial zirconium alloys with transition metals. From studies on the hydrogen reaction, one can establish the stability of the intermetallic compounds relative to the hydrides of the component metals. The intermetallic compounds may also be used as getters for hydrogen and as materials for hydrogen storage. This paper will present the results of equilibrium and structural investigations of the ZrCr2-H2, the ZrV2-H2, and the ZrMo2-H2 systems for the temperature range of 0" to 900°C and the pressure range of Torr to 1 atm. A complete thermochemical analysis will be made of the measurements. Some of the results have been referred to in a survey paper' covering the types of hydrogen behavior observed in zirconium alloys. LITERATURE ZrCr2 is the only intermetallic phase in the Zr-Cr system.' It exists in two allotropic modifications. Between 900" and 1000°C the hexagonal MgZa (C14) structure transforms into the cubic MgCu2 (C15) structure. Beck4 has previously shown that the intermetallic ZrCrz takes up hydrogen into solid solution up to a composition H/ZrCr2 = 1.14 at room temperature and 1 atm hydrogen. A ZrCr2 hydride was stated to have a fcc structure. The binary Zr-V system has been investigated by williams5 although Rostoker and Yamamoto 6 had given earlier a partial phase diagram. ZrV2 was the only intermediate phase found in the system. wallbaum7 reported earlier that the compound ZrV2 hp the MgZm (CJ4) type structure with a = 5.288A and c = 8.664A. This was not comfirmed by Elliot8 nor by Matthias, Compton, and corenzwit9 who reported that ZrV2 has a cubic MgCua LC15) type structure. The latter authors gave a = 7.429A. Also, several authors do not agree on the temperature of the peri-tectic horizontal. It appears that ZrV2 may exist in several modifications. A high capacity for occluding hydrogen was reported by Beck4 who found H/ZrV2 = 1.38 at room temperature and 1 atm of hydrogen. ZrMoz is the only intermetallic phase in the Zr-Mo system.10 ZrMoa is formed by a peritectic reaction of the melt with a molybdenum-rich solid solution. ZrMo2 has the C15-type structure with a = 7.59A. MATERIALS High-purity chromium, vanadium, and molybdenum were used together with grade 1 crystal bar zirconium to prepare the alloys. Table I lists the chemical analysis and the sources for the metals. Ten-gram alloy compacts were levitated and melted in an inert argon

Jan 1, 1968

By L. F. Mondolfo, B. E. Sundquist

The crystallographic orientation relationships resulting when lead is nucleated from the liquid by Ni, Cu, Ag, and Ge were determined. For each nucleating agent several definite orientatioz relationships were found. These relationships seemed to be controlled by good symmetry relations and low crystallographic disregistry between mating planes. For any given nucleating agent the under colling for nucleation was found fairly constant and independent of the orientation relationship and consequent disregistry. It was also found that, upon re melting and refreezing the Pb, the orientation relationship was changed. These findings prove that crystallographic disregistry is not the controlling factor in heterogeneous nucleation from the liquid. The results of this investigation tend to confirm the theory presented in a preceding paper that heterogeneous nucleation starts with the formation of an adsorbed layer of nucleated metal on the nucleat-ing impurity. Evidence is given that cavities in the nucleating agent act as centers of nucleation. IT has long been known' that solid extraneous particles are active in catalyzing phase transformations that occur in a system, particularly condensation and crystallization. It is well established that these heterogeneities act as catalysts by providing surfaces upon which nuclei of the precipitating phase can form with activation energies smaller than those required for homogeneous nucleation. Numerous investigations have shown that in this process of heterogeneous nucleation: a) the nucleus forms with one, or several, definite crystallographic orientation relationships with the nucleating phase2-4 and b) that there is a small range of undercoolings or super saturations characteristic of the nucleation of a given solid on a given Substrate.5-10 Turnbull and vonnegut11 have developed a theory based on theories developed by Volmer12 and Turn-bull and Fisher1= for heterogeneous nucleation from gases and liquids, that relates the super saturation or undercooling required for nucleation to the dis-registry between the lattices of the nucleus and the nucleating agent. This theory predicts that nucleation should occur with the orientation relationship between the nucleus and nucleating agent that minimizes the disregistry. Further, it predicts that the undercooling or super saturation necessary for nucleation should be a function of the disregistry. Numerous investigations have dealt with the orientation relationships resulting from the condensation of vapors onto crystalline solid substrates2,3 and a few with the nucleation of one phase by a second phase in solidification4,14. Others have dealt with the supersaturation8-10 and undercooling5-7 associated with nucleation in condensation and solidification respectively. However, there is virtually no report that gives both of these factors for the same system. In this investigation a study was made of the undercoolings and orientation relationships resulting when Pb is nucleated from the liquid by Ni, Cu, Ag, and Ge. It was the purpose of this investigation to check the Turnbull-Vonnegut theory, i.e., the importance of crystallographic disregistry between nucleating catalyst and nucleated metal. The results indicate that disregistry is not an important factor in nucleation and that the nucleation process is probably somewhat more complex than current theories suggest. EXPERMENTAL PROCEDURE Small single crystals of nickel, copper, silver, and germanium were prepared from materials of four to five nines purity, and the Pb used was also 99.999+ pet pure. Cu and Ag single crystals were prepared by sealing small chips of Cu or Ag in an evacuated quartz capsule and heating the capsule at 2000°F for 1 hr before cooling. Nickel crystals of 200 diam were also prepared in evacuated quartz capsules, but melting was done by heating the capsules in an oxy-acetylene flame for a few minutes. These spheres were invariably polycrystalline so

Jan 1, 1962

By C. W. Covington, H. C. Cook, J. F. Libsch

Sputtered aluminum thin films were prepared in each of two conventional bell-jar vacuum systems. One system utilized an inner "getter sputtering" enclosure; the second system was a standard diode sputterirlg arrangement. Consistent and repeatable film properties were obtained in both systems provided sufficient cleanup and presputter time was allowed. The various problems associated with aluminum sputtering are discussed, with pavticular attention to the vzinimization of film contamination by outgassing and oxidation. The growth and structure of the aluminum thin films were studied with the aid of electron wzicroscopy and related to deposition rate, substrate temperature, and film thickness. The resistivity of the films was correlated to film structure and surface roughness. Resistiuities as low as 3.5 microhm-cm (2.3 x bulk resistiuity) were measured for relatively thick films (20,000). THIS investigation was undertaken to explore the problems and film characteristics associated with the sputtering of aluminum. Earlier work reported on aluminum sputtering has stressed the contamination of both the cathode and film as well as abnormally low deposition rates.' Data on sputtering yields,3 however, indicate that aluminum sputtering could be practical. Recently, it was shown by Theuerer and Hauser3 that aluminum sputtering was feasible in a conventional oil-diffusion pumped vacuum system provided an inner "getter sputtering" enclosure was employed to minimize gaseous contamination. This technique employs the getter action of aluminum to produce extremely low reactive gas partial pressures in the deposition zone. Aluminum films have already found use in solid-state electronic devices such as transistors and integrated microcircuits. To the present time, most aluminum films have been prepared by evaporation due to the ease with which this metal can be evaporated. Controlled sputtering of aluminum offers a number of advantages over evaporation, however, such as 1) manufacturing compatibility with other metals which must be sputtered, 2) better control of process variables resulting in more repeatable and uniform films, 3) generally better adhesion resulting from higher-energy metal atoms, 4) built-in material supply for continuous operation, and 5) the ability to deposit alloys. The data presented in this paper are the result of two separate studies. One study utilized a small "getter sputtering" enclosure (System A) and was concerned with the structure and properties of thin aluminum films up to about 1600A thick. In this thickness range, study of the film structure was possible with transmission electron microscopy. The second study was conducted with a conventional diode sputtering apparatus (System B) and covered $ much wider range of film thickness up to about 34,000A. Emphasis was placed on developing techniques which would extend aluminum sputtering to an existing in-line continuous deposition process.4 I) EQUIPMENT AND PROCEDURE System A. Fig. 1 shows a schematic of the "getter sputtering" inner chamber used and Fig. 2 is a photograph of this chamber mounted on an 18-in. access ring of a standard liquid-nitrogen trapped bell-jar vacuum system. This chamber was designed after a similar system used by Theuerer and auser.3 The configuration allowed the glow to emanate in all directions from the cathode. Argon (99.99') was allowed to enter only at the top of the chamber such that most of the reactive gases entering with the argon were reacted with the aluminum vapor and deposited on the chilled walls of the chamber before reaching the deposition zone. Provision was also made as shown, Fig. 1, for heating and cooling the substrate from about 20" to 400°C. Two substrate materials were used: Corning 7059 glass (12 by 12 by 0.048 in.) and carbon films about

Jan 1, 1967

By R. H. Oliver

The purposes, types, preparation, and testing of flocculants are discussed. A flocculation compendium is included, indicating choice of flocculant for a given set of conditions. An economic evaluation of the process is presented. Solid-liquid separation is a major expense in most mineral processing flowsheets today. Proper use of flocculation can often lower the overall cost of a solid-liquid separation by reducing the size of sedimentation or filtration equipment. PURPOSE OF FLOCCULATION When is use of a flocculant justified in terms of overall process economy? 1) when it results in process improvement, such as producing a clear liquor for electrolysis, precipitation, ion exchange, or solvent extraction; 2) when it results in satisfying requirements of pollution abatement ordinances; 3) when it results in increased recovery of values that would otherwise be lost; or 4) when it results in considerable savings in capital expenditure due to use of smaller equipment. When can the use of a flocculant adversely effect overall process economy? 1) when it results in increased filter cake moisture, 2) when it results in decreased filtration rates, 3) when it results in considerable bulking of sedimentation underflow, 4) when it contaminates either the supernatant or sludge, or 5) when the cost of using the flocculant is greater than the savings due to its use. First step in consideration of flocculants for a given application is establishment of desired goals in terms of increased clarity of effluent, increased recovery of solids, or decreased size of equipment. The most economical solution to a solid-liquid separation problem is that combination of equipment and flocculant costs which meets the established goals at the lowest total annual expenditure. This economic decision can be made only after technical data has been obtained as outlined in next four sections. TYPES OF FLOCCULATION Flocculation is a process wherein individual particles are united into more or less tightly bound agglomerates or flocs, thereby increasing effective particle size of solids suspended in a liquid. Degree of flocculation of a suspension of finely divided solids in a liquid is controlled by a combination of probability of collision between particles and probability of adhesion after the collision has occurred.' Probability of collision can be increased commercially through use of a paddle-type flocculator, combination flocculator, and clarifier or flocculating-type feedwell. Probability of adhesion usually can be increased by addition of a reagent known as a flocculant. Reagents act as flocculants through one or a combination of three possible mechanisms. The first is electrolytic neutralization of inter molecular repulsive force due to Zeta potential. This neutralization enables Van der Waal's cohesive force to hold the particles together after they collide.' The second is the precipitation, within a definite pH range, of voluminous metallic hydroxide flocs which entrap fine particles. This process is known as coagulation.14 Third is the bridging of two or more particles by either natural or synthetic long-chain, high-molecular-weight organic polymers. This bridging is accomplished by adsorption of two particles at different sites on the same molecule or by bonding of two molecules, each adsorbed to a different particle. Considerable effort has been expended in explaining the action of these polyelec-trolytes and detailed accounts are available in the literature. CHOICE OF FLOCCULANTS It is impossible, at the present time, to specify the optimum type or amount of flocculant for a given application from a knowledge of the material to be treated. The standard procedure is to choose reagents or combination of reagents most likely to be effective on the slurry; then test them.' The accompanying Flocculation Compendium (Table I) contains details of all currently available reagents sold as flocculants and it should be used when selecting the group of reagents for tests. This selection can be based on information listed in columns 4 and 5. Column 4 presents industries in which each floc-

Jan 1, 1961

By R. I. Jaffee, F. C. Holden, H. R. Ogden

The mechanical properties of Ti-C and Ti-C-0 alloys can be altered by heat treatments to dissolve or reject carbon from solid solutions. The maximum strength is obtained by annealing just below the peritectoid temperature. Quenching from the ß-carbide field results in softening. Impact behavior is influenced by the extent of solution of interstitials CARBON is not as important in titanium alloys as it was in earlier years when graphite induction melting was more prevalent. With induction melting, carbon pickup is about 0.5 to 0.8 pct. Cold-mold arc melting with graphite electrodes, which is among the more current methods, results in carbon pickup of 0.1 to 0.2 pct. The titanium-sponge melting stock itself contains about 0.02 to 0.07 pct C, picked up from the magnesium used in the reduction process. Thus, carbon is still one of the more important contaminants in titanium and needs careful study and evaluation in order to understand titanium and titanium alloys. Moreover, it is equally proper to consider carbon as an alloying addition because, in controlled amounts in titanium alloys, it may serve a most useful function. This is particularly the case where the ultimate in toughness is not required. The addition of carbon to titanium results in a peritectoid-type system as illustrated in Fig. 1. The solubility of carbon is seen to be higher in the a phase than in the ß phase in the temperature range of interest. Below the peritectoid temperature, the solubility of carbon in a titanium decreases with decreasing temperature. This type of system permits many microstructural and heat-treatment variations which may influence the properties of the alloys. The study of how the microstructure variations affect mechanical properties is described in this paper. Experimental Procedures The experimental techniques and testing methods used in this work were the same as those described in detail in a previous paper.' No further descriptions will be given except where changes have been made. High purity iodide titanium, produced by the New Jersey Zinc Co., was used in all of the alloys. Carbon additions were made using high purity spectro-graphic carbon electrodes in order to obtain the highest purity carbon available. The alloys were prepared by double arc melting to reduce the possibility of segregation. The alloys were forged at 875°C to ¾ in. round, descaled, vacuum annealed for 5 hr at 900°C to remove the residual hydrogen, and swaged to ¼ in. round at 750°C. All heat treatments were made on material in this initial condition. Two of the early Ti-C ingots were found to have been contaminated in melting. Vacuum-fusion analyses showed that the two contaminated alloys had picked up oxygen, so the test results on these two alloys permitted an evaluation of ternary Ti-C-0 alloys. A list of the alloys and their compositions are given in Table I. Properties of Ti-C Alloys In a peritectoid-type system, such as the Ti-C system, the chief microstructural variable which influences the properties is solid-solution strengthening. For a given composition, it is possible to control the disposition of carbon, either in solid solution or as the compound Tic, by suitable heat treatments. When in solid solution, the carbon present in the alloy makes a definite contribution to the strength of the alloy. However, when present as Tic, it has very little effect on the strength. Other microstructural variables considered in this investigation were transformation structures and grain size. As discussed in subsequent sections, these variables have only a minor effect on properties as compared to the solid-solution effect. Solid-Solution Strengthening: The principal strengthening mechanism for Ti-C alloys has been found to be solid-solution strengthening. Carbon in interstitial solid solution strengthens titanium, but has very little strengthening effect if present as the compound. Since carbon has a higher solubility in

Jan 1, 1956

By B. J. Reid, J. N. Greenwood

It has been shown1 that cavities are formed in the grain-boundaries of copper and 70:30 brass (as well as in magnesium) by the application of tensile stresses at elevated temperatures. For a given rate of strain, cavities are formed in brass at a lower temperature than in copper, and they are larger; it was thought that this earlier onset of cavitation in the alloy might be caused by the distortion produced in the metal lattice by the alloying element. If the cavitation is caused by the vacancy-migration mechanism, ' the distortion of the lattice would reduce the energy required for vacancy formation, and thus provide vacancies more readily to assist in the cavitation process. It was decided to study a range of alloys to find out whether increasing amounts of an alloying element would cause any progressive increase in, or modification of, the cavitation in copper. Nickel was chosen for this purpose, since it is soluble in copper throughout the whole range of alloys. A further advantage was its low vapor pressure, which enabled the preparation of the alloys in vacuum and the consequent reduction of their hydrogen content. It has been suggested that cavities may be nucleated by the migration of hydrogen atoms to the less regular atomic spacing at the boundary, and their combination to form molecules. The latter being too large to diffuse back into the lattice, form stable cavities into which vacancies could condense. PREPARATION OF ALLOYS All alloys, ranging from pure OFHC copper to 43.7 at. pct of nickel (carbonyl), were prepared in a vacuum furnace employing high-frequency induction heating. When a charge was melted, vigorous gas evolution occurred initially, after which the pressure was reduced to 10-4 mm Hg, using an oil-diffusion pump. After remaining molten for a further 1/2 to 1 hr to aid degassing, the melt was allowed to solidify in the crucible. In this way contamination by gases evolved from the surface of a separate mold was avoided. The list of alloys prepared is shown in Table I. The percentage of hydrogen present in the OFHC copper was reduced from 0.0054 at. pct to 0.0012 at. pct, as measured by a vacuum-fusion method. The alloys were cold-rolled to 0.03-in.-thick strip, and were homogenized in vacuum for 1 hr at 900°C after every 30 pct reduction. At a thickness of 0.05 in., the strip was held at 450°C for 24 hr under high vacuum (10-5 mm Hg) in an attempt to reduce the hydrogen content still further. Tensile specimens were prepared with a gage length of 1 in., and were finally annealed for 2 hr at 800°C to give an average grain diameter of 0.1 mm. TENSILE DEFORMATION The specimens were deformed in a tensile straining machine operated by a synchronous motor at a constant rate, equivalent to 1.7 pct per hr on a 1-in. gage length. The specimen and grips were contained in a vacuum jacket. A pressure of 10-2 mm Hg was maintained, which was sufficient to prevent surface tarnishing of a specimen held at 650°C for 24 hr. Tests were made between 20. and 650°C and the temperature was controlled to within ±3° over the gage length. The load was measured by the deflection of a cantilever spring. a) Mechanical Properties-The ductility curves shown in Fig. 1 fall into two groups. Figs. 1(a) and (b) for copper and the 5 pct nickel alloy, respectively, show a continuous fall in ductility with a steep fall between 200° and 400°C. Figs. l(c) to (f) show the same sudden loss of ductility between 200° and 400°C. With higher temperatures, however, there is a recovery in ductility values to a maxi-

Jan 1, 1959

By K. D. Sheffler, R. W. Hertzberg, R. W. Kraft

The Ni-Ni,Ti lamellar eutectic alloy responds to unidirectional solidification by alig-nment of the platelets of the two phases roilghly perpendicular to the moving solid-liquid interface. X-ray diffraction and zetallographic experimenls show that the following cryslallographic relationships exist in the directionally solidified ingots: lamellar interface w [0001] Ni3Ti 11 (111 }Ni local growth direction (1120) Ni3Ti n (110) Ni These relationships are rationalized on the basis of the crystallography of the constituent phases. In recent years researchers in several laboratories have shown that an aligned or "controlled" micro-structure can be produced in many binary eutectic alloys by forcing the alloy to solidify unidirectionally in an elongated crucible.'-3 It also has been found that a unique crystallographic relationship often develops between the two phases in such a structure.4, 5 Since a parallel arrangement of phase particles in an alloy should, and, in many cases, does, produce unique and highly anisotropic properties, the alloys so produced are of potential interest in a variety of applications. Studies of these alloys are also of interest at the academic level because they help to improve our basic understanding of the way in which metals solidify. This paper reports the metallographic and crystallographic structure of another of these alloys—that between nickel and Ni3Ti. In a separate paper, hot tensile and stress rupture properties as well as deformation and fracture behavior of the same alloy are discussed.19 EXPERIMENTAL DETAILS AND RESULTS General Characteristics of the Eutectic Reaction. Several small exploratory Ni-Ti melts were made to resolve a conflict in Hansen6 over the eutectic composition. The alloy containing 16.4 at. pct Ti, Fig. 1, appeared metallographically to be closest to the eutectic composition, although small dendritic patches of proeutectic nickel are present. The eutectic morphology is lamellar and appears quite sharp and angular, indicating a strong influence of crystallographic factors upon the microstructure. Some of the eutectic nickel platelets are continuous with the primary phase, whereas in hyper eutectic melts a "halo" of nickel generally surrounded the proeutectic inter-metallic platelets, indicating that nickel is the nucleating phase in this eutectic reaction. Extensive precipitation of Ni3Ti has occurred in the primary nickel phase in a Widmanstatten pattern which makes the dendrites difficult to differentiate from the eutectic mixture. This phenomenon may account for discrepancies among the eutectic compositions reported in the literature. The sloping solvus line bounding the nickel primary solid solution explains the presence of the intermetallic precipitate. Poole and Hume-Rothery7 report that the solubility of titanium in nickel drops from 13.8 at. pct at the eutectic temperature (1304°C) to approximately 9 at. pct at 700°C (Ni3Ti, on the other hand, melts con-gruently and exhibits essentially no solubility for nickel or titanium). Lever rule calculations based upon Poole and Hume-Rothery's phase diagram indicate an increase from 21 vol pct Ni3Ti in the eutectic mixture at the eutectic temperature to 61 pct at room temperature. Selective point count measurements of the volume fraction of eutectic Ni3Ti platelets in the eutectic microstructure yielded a value of 46 i 2 vol pct, which required extensive precipitation of Ni3Ti from the nickel solid solution onto the eutectic intermetallic platelets during cooling from the solidification temperature. Influence of Directional Solidification. Nickel bar and titanium sponge of 99.9 pct purity were induction-melted in zirconia crucibles under argon in the proportion of 13.8 wt pct (16.4 at. pct) Ti to produce 14-lb master heats which were cast into investment molds yielding 5/8-in.-diam pins approximately 10 in. long. These pins were remelted under argon in horizontal zirconia boats and unidirectionally solidified to produce 1/2-in.-sq ingots approximately 44 in. long. Details of the experimental techniques and apparatus are described elsewhere.'

Jan 1, 1970

By Donald A. Corrigan, John Chipman

It is shown that the heat of solution of nitrogen in liquid-iron alloys is Proportional to the interaction coefficient. This proportionality forms the basis for a method of predicting nilrogen solubility at any temperature when the interaction coefficient is knoun at 1600°C ov when the solubility is known at some one temperature. A method for predicting the solubility of nitrogen in liquid steels, including alloy steels, was proposed some years ago by Langenberg.' Since the data available at that time were restricted to temperatures of about 1600°C, his predictions were similarly restricted. More recently, Nelssonz devised an approximate method for estimating nitrogen solubilities at higher temperatures. It is the purpose of the present paper to present a method by which the solubility of nitrogen at any temperature within the steelmaking range may be predicted from the known interaction coefficients at 1600°C. Nitrogen in iron and in all of the alloys to be considered here is known to obey Sieverts' Law; hence the activity of nitrogen in an alloy of fixed composition is proportional to its concentration, here expressed in weight percent. The activity of nitrogen is defined as equal to its weight percent in the binary Fe-N solution at any temperature. Its activity coefficient in any alloy is the comparison being made at equal values of T and PN2. In all alloys the activity coefficient changes with composition, log fN being a linear function of the percent of alloying element at low concentrations. In the following calculations the alloy compositions are restricted to those in which this linear relation is exhibited. For most systems this is a fairly wide range, for example up to 5 pct V and to 10 pct Cr or Ni. For this composition range. complex alloy is the same as that in the Fe-N-j ternary solution. Langenberg's method of prediction was based on Eq. [3], the values of log fN for various additions being taken from a graph. Instead of using the simple addition indicated in Eq. [3], he used a more complex method suggested for Fe-S-j alloys by Morris and Buehl.3 For the composition ranges in which Eq. [2] is valid the two methods are equivalent. Nelson extended the prediction to other temperatures on the assumption that logfN is inversely proportional to absolute temperature. It may be shown that inverse proportionality would result if the partial molal entropy of nitrogen were unaffected by addition of the alloy element, an assumption which Nelson designated as a "regular" solution. It has been shown recently by the authors4 that this assumption is not a good approximation. The heat of solution of nitrogen (N2) in an alloy may be obtained if the solubility is known as a function of temperature. Measurements of this nature have been reported recently by Beer5 for Fe-Mn alloys, by Nelsson2 for Fe-C, by El Tayeb and parlee6 for Fe-V, and by Pehlke and Elliott7 for many alloys embracing a wide range of compositions. In connection with the work of the last named authors it is noted that their Fig. 22 gives the heat effect per half mole rather than per mole of N2. In using their data we have shifted some of their values slightly by using the best straight line in preference to their curve. Their data, supplemented by others from the literature, are summarized in Table I. In Fig. 1 the heat of solution of nitrogen is shown as a function of composition for the several alloy systems. The difference between the heat effect in the alloy. AH; and that in pure

Jan 1, 1965

By J. A. Elias, A. J. Heckler

A simplified method is described for the prepauatior~ of comnposite samples for determination of one quadrant of an X-ray pole figure of a sheet material by the Schulz reflection method. Other special conzposite sampling methods were used to investigate specific areas of the pole figures. Plotting nets for these sampling plans were drawn by a computer. Step-scanning techniques were used to obtain detailed information about specific regions of a pole figure. Satisfactory pole figures were obtained for large-grained (0.5 to 1 in. diam) sheet materials. COMPOSITE sampling techniques for preparing one quadrant of an X-ray pole figure have been reported by a number of authors.'-5 All of these methods are similar in that they use a composite sample prepared by stacking and holding together coupons cut from a sheet material. The schulze reflection technique is applicable to the composite sample so that only minor corrections are needed for defocusing and absorption effects. Also, a relatively large area of a sheet is sampled through the use of multiple cross sectional pieces. In general, sampling methods necessary to obtain a composite-type specimen have been formidable enough to discourage their use in routine pole figure determinations. This paper presents a versatile sampling technique which is simpler than those previously described along with an accurate method for obtaining stereographic plotting nets for any orientation of a composite specimen surface. A step-scanning technique is used to obtain more quantitative determinations of specific regions in pole figures, and the composite technique is satisfactorily applied to determine pole figures of very large grained sheet materials. REVIEW OF PREVIOUS TECHNIQUES Mueller and ~nott' first proposed a technique for obtaining one complete quadrant of an X-ray pole figure by using a sheet sample and six different composite samples in which the individual coupons were aligned and held together by two types of jigs. Data from the seven different samples were overlapped to obtain a single quadrant. ~eieran~ and Lopata and ~ula~ obtained a complete quadrant of a pole figure from a single sample prepared by forming a cube from coupons cut from a sheet. These coupons were bonded together using an epoxy cement, and a cut was made at 54.7 deg to the normal, transverse, and rolling directions of the sheet, Fig. I. The normal direction of the cut face projects into one quadrant of a stereographic projection at equal angles to the three principal orthogonal reference directions mentioned above. X-ray reflection data obtained along a spiral path in physical space could then be plotted on a stereographic net offset 54.7 deg from the normal, as shown in Fig. 2. A rotation about an axis representing the surface normal is designated as a, and a rotation about an axis lying in the plane of the surface is 4. From Fig. 2 it is apparent that a tilting angle of no more than 55 deg @ is necessary to obtain the complete quadrant. ~eber~ devised an improved sampling technique by cutting coupons 45 deg to the rolling direction of a sheet, stacking and bonding these with an epoxy cement at an angle of 54.7 deg to a reference plane, and finally grinding and lapping a flat specimen surface parallel to this plane. An important consideration of this technique and that of Lopata and ~ula~ is variation in the diffracting power of the specimen surface due to the bonding medium. While any composite-type sample will have some such variation due to boundaries between coupons, the problem can be minimized by elimination of the bonding medium and by avoiding prolonged etching times which could lead to significant inter lame llar attack of the coupon faces. EXPERIMENTAL TECHNIQUE We were working concurrently though independently on a technique similar to ~eber's,' but which does not

Jan 1, 1968

By Keith R. Bock, Norman Parlee, Albert M. Barloga

Coils of pure iron and iron-carbon alloy wire (0.05 to 0.80 pct C) and sufficient sulfur to saturate the solid phase were equilibrated in evacuated or argon filled tubes. After rapid cooling, and removal of the outside nonmetallic layer, the wires were analyzed for carbon and sulfur and the data used to construct an Fe-S binary and isotherms of the Fe-C-S ternary in the range 1149° to 1427°C. THE solid solubility of sulfur in steel is of interest in connection with such phenomena as hot shortness, burning, and so forth. "Burning," the more or less permanent damage that some steels suffer when heated for forging or rolling, has been shown to be related closely to the behavior of sulfur and less closely to carbon and oxygen.' Attempts to interpret burning phenomena in steels fail because of lack of data on the Fe-C-S and the more complex systems in this family. Rosenqvist and Dunicz 2 and Turkdogan, Ignatowicz, and pearson3 have largely elucidated the Fe-S diagram in the region of interest but no information on the Fe-C-S diagram in this region appears to be available in the literature. This paper deals with the elucidation of the Fe-C-S diagram in these interesting ranges. The method employed is different from those used by previous workers2'3 on the Fe-S system. EXPERIMENTAL METHOD Pure iron wires (Ferrovac E) or iron carbon alloy wires of 1 mm in diam were cleaned with acid and acetone, coiled, and placed in silica tubes (7 mm OD and 5 mm ID) previously closed at one end. Enough sulfur was added to assure saturation of the solid iron phase. The filled tubes were either simply evacuated and sealed, or filled with argon at a reduced pressure and sealed. The argon was required at the higher temperatures to prevent collapse of the tubes. The filled and sealed tubes were placed in the middle uniform temperature zone of a Globar tube furnace and equilibrated at temperatures ranging from 2100°F (1149°C) to 2720°F (1493°C). After equilibration the tubes were removed from the furnace and quenched in air or water, the form of quenching being found to have no effect on the results. The tubes were broken open and the coils were placed in a 1 : 1 HC1 solution to remove the sulfide-rich layer. The coils were then cut into small pieces and analyzed for sulfur and carbon. In the early stages of the investigation different equilibration times ranging up to 17 hr were tried and the cores of the wires were analyzed to test for saturation. One hour appeared to be sufficient to reach maximum sulfur content at 1454°C and 2 hr sufficient at 1149°C. The practice adopted was to use at least 3 hr at the higher temperatures and at least 5 hr at the lower temperatures. The iron-carbon alloy wires used were made by carburizing pure iron wire with carbon monoxide gas in a one inch diameter ceramic tube at about 1204°C. Differing carbon contents were obtained by allowing fairly large coils to react with the gas for varying lengths of time at a flow rate of 800 cc per min. The reacting times ranged from 15 min to 4 hr depending upon the amount of carbon desired. The furnaces used were controlled by means of Leeds and Northrup Speedomax Type H instruments with D.A.T. Control attachment. Thermocouples were calibrated against the melting points of gold and copper. The temperatures recorded appear to be accurate within i2.7OC. Starting with run No. 85 and continuing with the

Jan 1, 1962

By J. F. Butler

Boron nitride, BN, has been identified in boron low-carbon steels by means of light microscopy, electron microscopy and diffraction, and chemical analysis. This boron nitride is responsible for strain aging suppression in these steels through the removal of nitrogen from solution; however, small oxidizing potentials which result in deboronization In the austenitic temperature range also result in boron nitride decomposition and the subsequent occurrence of nitrogen strain aging. THE addition of boron to low-carbon steels is an effective means of eliminating the detrimental effects of strain aging;',' however, the mechanism by which boron suppresses strain aging has not been established. In studying the effect of heat treatment on the strain aging of boron low-carbon steels, Butler3 found that no detectable nitrogen was present in the ferrite solid solution after various heat treatments even though the total nitrogen analyses of the steels were 0.003 pct N. Apparently the function of boron in suppressing strain aging is the removal of nitrogen from solution to form a relatively stable second phase. A clue to the identity of this phase was found by Speight4 who extracted from an aluminum-killed boron steel an acid insoluble residue which contained boron and nitrogen in proportions corresponding to the compound, BN. Shyne and Morgan5 found that the addition of nitrogen to a hardenable boron steel lowered its hardenability in a manner which suggest-d the presence of nucleating particles (presumably BN) even at high austenitizing temperatures. Although this indirect evidence for the presence of boron nitride in boron steels has been found, the compound was not identified metallographically as a separate phase in these steels. Because boron has an affinity for oxygen comparable to that of silicon,5 boron steels must be thoroughly killed with aluminum in order to prevent the reaction of boron with oxygen.7 The affinity of boron for oxygen also is demonstrated through the ease of deboronization of a boron steel in an oxidizing atmosphere. Digges, Irish, and Carwile8 found that a mixture of wet natural gas and air caused both deboronixation and decarburization of a boron steel at 1900°F. The boron content rose at the surface of the steel due to the formation of boron oxide. Shyne and Morgan9 encountered deboronization at much lower oxygen potentials; loss of boron occurred in steels treated at 1750o F under argon purified to 0.0014 pct 0, even though decarburization did not occur. Boron oxide, B2O3, is much more stable than the nitride, BN, as is indicated by the standard free energies of formation at 980oC: -235 kcal per mol of B2O3 (liquid) and -28 kcal per mol of BN (crystal). It is evident that the oxygen potential must be kept low in the presence of BN in order to prevent its decomposition and the subsequent formation of B2O3' The purpose of the work described in this paper was to identify the nitride phase responsible for the removal of nitrogen from solution and to study its stability under oxidizing conditions. IDENTIFICATION OF BN In order to identify the phase containing boron and nitrogen, several alloys high in boron and nitrogen were made up in a small vacuum furnace. In alloys 3 and 4, Table I, electrolytic iron was used as a charge material. Boron additions were made in the form of 18 pct B ferro-boron after the iron was molten. In the case of alloy 3, the charge was solidified at this point; for alloy 4, nitrogen at 1 atm pressure was admitted over the melt before it was

Jan 1, 1962

By R. G. Schaal

The Magma Mine at Superior, Arizona had an unimpressive beginning as a. worked out silver mine that was purchased in 1910 for $130,000 and then incorporated into the Magma Copper Company which has been continuously mining the deposit since 1915. Copper is the principal metal produced. However, zinc, gold, silver and iron have also contributed to the total metal product. Prior to 1973, the production rate of 1500 tons per day was maintained by hoisting 4 and 5 ton pay-loads at speeds up to 1100 FPM through four small hoisting shafts. The hoists used included 8' and 7' diameter double drum, double clutched doc. and a.c. motor driven geared units built as early as 1917. One old timer that did yoeman's service for many years was an 8' Bollen Pulley with a.c. geared drive that utilized an endless rope system with mechanical grip and release mechanism built into the pulley's throat. Primary power for all of these old production hoists was 25 Hertz. Exploration that was completed in 1967 developed a reserve of 10 Million tons of 5 per cent copper which was sufficient to justify new capital expenditures for enlargement of productive facilities to 3300 tons per day. To reach the new reserves a 22' I.D. concrete lined vertical shaft was sunk to an initial depth of 48431. A 12 ½ diameter adit was bored from the mill side of the mountain to intercept the shaft at a point about 1100' below the collar. This 9200' long adit provides the right-of-way for the mine to mill railway and egress for heavy supplies from rail head as well as a conduit for potable water, electric power, compressed air and mine drainage. Mine operating levels begin 3582' below collar and thereafter have been cut at 100' vertical intervals. Skip loading stations have been installed 4282' and 46821 below collar with an in-shaft skip dump station located 800' below collar. This provides for 4500 tons of storage above the mine to mill train loading station located 1100' below collar. Ore and waste are moved horizontally from stope chutes and drift headings by rail haulage to transfer raises which channel all material to the two mentioned skip loading stations. The 22' diameter shaft crossection is arranged with the two skip compartments on the south side separated from the manway and utility compartments by the cage compartment. The cage counterweight operates in a portion of the southwest quadrant of the shaft. The hoisting plant consists of two conventional Nordberg double drum, single clutch, geared, d.c. motor driven units that are identical except for the drives and automatic controls. The drums are 15' diameter x 105" face with LeBus grooving for 2-1/8" diameter wire rope wound in three layers. Mechanical design emp1o::ed is the same as has been used for the past twenty- five years. One drum is attached to the main shaft by 2 large bolted integral shaft flenge. The loose drum is engaged by a 4-arm positive tooth clutch meshing with radially cut teeth of the clutch ring which is bolted to the drum. The drums are one-piece steel plate weldments having a finished weight of 60 tons each and are mounted upon a 36" diameter shaft that also weighs 60 tons. With the addition of brake rings, clutch, bull gear and rope, the total dead weight adds up to nearly 250 tons that is supported by just three double row spherical roller pillow block bearings. There is no bearing between the drums. All pinion shafts as veil as all head sheave journals are supported in the same type of antifriction bearings.

Jan 1, 1975

By A. J. Griest, A. P. Young, P. D. Frost

A study was made of the effect of ß grain size on the tensile ductility of a-ß titanium alloys haet treated to strengths in the range 165,000 to 180,000 psi. It was concluded that the primary cause of $ embrittlement is the large grain size which is obtained so rapidly in many a-ß titanium alloys when heated into the ß field. It was also found that grain-boundary a and acicular intragranular a in the microstructure are not necessarily detrimental to the tensile ductility if the prior ß grain size is sufficientlv fine. BETA embrittlement occurs in a-ß alloys when the finish hot-working or stages of the subsequent heat treatment have been done at temperatures in the ß field. In order to avoid this embrittlement, for-gings must be finished at temperatures under the ß transus. For better workability, it would be desirable to forge at temperatures in the ß field. For this reason, research has been conducted to determine the cause of ß embrittlement. Although ß grain size has been considered as a possible contributing factor to ß embrittlement, only limited research has been done to evaluate its effect. The effect of addition elements on ß grain growth in unalloyed titanium has been studied by Liu.' Complete mechanical property data, however, were not obtained for correlation with grain size. Ogden and coworkers2 have investigated the effect of both a and ß grain size and grain shape in Ti-Cr-Mo alloys. Their data indicated that tensile ductility was relatively insensitive to ß grain size at annealed strength levels (140,000 psi ultimate). No study has apparently been made to determine the relationship between ß grain size and ß embrittlement in a-ß alloys heat treated to high strength (180,000 psi ultimate). The high-strength condition is that in which ß embrittlement is particularly severe. Beta grain size is not easily varied through the use of a range of annealing times and temperatures in the ß field, because of the rapidity with which ß grain growth occurs. Short-time heating techniques are needed to obtain the smallest ß grains. EXPERIMENTAL PROCEDURE In this research, the experimental procedure was as follows: 1) Tensile and microimpact specimen blanks were resistance heated to produce a range of ß grain sizes. 2) The tensile and microimpact blanks were then solution-treated in the a-ß field and aged to produce predetermined ultimate tensile strengths in the range 165,000 to 180,000 psi. Tensile and impact ductility were then correlated with a grain size for constant strength levels. Materials—The alloys Ti-6A1-4V, Ti-4A1-4Mn, and Ti-16V-2.5Al are representative of the heat-treatable a-ß types in which ß embrittlement can be severe. The compositions and ß transi of the heats studied are given in Table I. The Ti-16V-2.5Al heat was melted and fabricated to 1/4 in. round at Battelle; the other alloys were obtained as 1/4 in. round rolled from commercial heats. The micro-structures of the as-rolled alloys were fine a-ß mixtures. This structure is typical of titanium alloys hot-worked extensively in the a-ß field. Resistance-Heating Treatments to Establish ß Grain Size—For direit resistance heating, a 1/4-in.-diam s~ecimen was heated with a 3000-watt transformer rated for 1000 amp with adjustable secondary voltage. The apparatus included a spray quench fixture. Both heating and quenching periods could be automatically controlled. For temperature measurements, 28-gage Chromel-Alumel wires were tack-welded to the specimen at its midpoint. A high-speed recorder-controller was used to control the heating cycle.

Jan 1, 1960

By T. O. Anderson, E. J. Stahl

Tests to increase knowledge concerning the process of hydraulic fracturing have been performed on three wells in the Allegheny field in western New York. The fracture orientation and its physical appearance at the well-bore was determined by impression packers run at appropriate times throughout the program. Fract~tre azi171uth and its dimensions thus determined were compared to the rock properties and treatment conditions. Logs to provide data concerning rock properties and to improve identification of vertical fractures were run before each step of the fracture treatment. Pressure osc~illu~iot~s were studied after each phase of tile treatments surface pressures during treatment. These measurements show that subsurface pressures exhibit many of the same characteristics as surface pressures during treatment of a well; however, transient pressures are dampened and deelayed in time at the bottom of the well when COmpared with surface pressures. Data obtained from this study provides a means to help locale and determine orientation of vertical fractures made in the open wellbore. Although these rests were performed in only one area, much of the information gained may be applied to fracturing programs in orher areas. INTRODUCTION In this study, various tests were made to provide data to better understand formation fracturing in situ. The reservoir studied is the Richburg oil sand which is a part of a series of predominantly grey sandstones and shales with occasional brown sandstone layers, known as the Caneadea unit of the Canadaway group of Upper Devonian age. Average physical characteristics of 12 cored wells in this area are 13-ft thick gross sand formation, 12-ft thick net sand. 1.7 md permeability, 11 percent porosity, 44 percent oil saturation (before secondary recovery) and 16 percent water saturation (before secondary recovery). This particular area of the Allegheny oil field was originally drilled prior to 1890. The first waterflood secondary recovery attempt was made in 1929. The present sustained waterflood development began in 1951 and is being carried out on a 4.5 to 5-acre spacing. Objectives of these tests were to (I) study fracture azimuth and dimension in relation to rock properties and treating conditions, (2) improve fracture identification techniques, (3) compare surface and subsurface pressure during all phases of treatment and (4) compare the periods of the fluid column oscillations occurring after each phase of treatment. Although the areas of investigation listed appear divergent in purpose, the program was designed primarily to increase knowledge of down-hole conditions leading to rock failure and extension of the fracture thus created. TEST PROCEDURES Wells upon which tests were conducted are designated as Wells 1 through 3. Well 2 was located 440 ft north of Well 1, and Well 3 was located approximately 880 ft west of Well 1. Wells 1 and 2 were completed with cable tools, while the producing interval of Well 3 was rotary cored using fresh water as the drilling fluid to obtain oriented cores. Tests were divided into separate phases on each well as the fracturing program progressed and consisted of measurements performed prior to treatment, during and after the formation breakdown operation. during fracture treatment and after treatment. To verify location and azimuth of the fracture created, and to obtain an impression of wellbore irregularities, a packer was used (Fig. 1). This packer was lowered on tubing into the wellbore opposite the interval under study and inflated in stages to allow the impression material to conform to the wall of the wellbore. A magnetic compass was attached to the tool for orientation reference. Impressions made prior to treatment showed no existing fractures in the wells under study. Impressions made after both formation breakdown and treatment pro-

By L. Luini, E. Lee

An exploratory survey of the heat treatment response of commercial titanium alloys (Ti-150A, RC-130B, and MST 3AI-5Cr al-loys) shows a wide range of possible hardness and microstructural characteristics. The hardening is primarily dependent on the solid state transformation of the phase. An age hardening reaction which apparently is associated with ß decomposition and precipitation has been shown. Brittleness and notch sensitivity appear to be characteristic of age hardening to high hardness. Ductility, tensile strengths, and elongations are given. A NUMBER of commercially produced titanium-base alloys have been made available by the titanium metals industry. These alloys were developed to combine the light weight of titanium with increased strength through alloy additions. Usually the properties of these alloys are reported for material in the hot rolled and annealed condition. Since the production of commercial titanium has preceded the investigation of titanium-base alloy systems, the potential of these alloys is not well known. Although heat treatability has been indicated for these alloys, data on the response to heat treatment are quite limited. The heat treatability of titanium is based on the allotropic transformation of the high temperature, body-centered cubic ß structure to the low temperature, hexagonal close-packed a structure at approximately 1625 °F. This transformation in the pure metal is of little practical significance in heat treatment since the high temperature phase is not amenable to controlled transformation. Alloy additions serve to alter the transformation kinetics and permit control of the transformation by heat treatment variables. The effects of different alloys on the allotropic transformation in titanium classify them in two general categories,1 a stabilizing and ß stabilizing elements. The a stabilizers, represented by aluminum and oxygen, raise the temperature limits of the low temperature phase and introduce a two-phase a + ß temperature range. The ß stabilizers, represented by molybdenum and columbium, lower the equilibrium temperature of ß and give rise to an extended temperature range where both a and ß phases are stable. Another type of ß stabilizer, represented by iron, chromium, and manganese, introduces a eutectoid reaction. Some commercial titanium alloys belong to complex multiple element systems. Based upon the phase diagrams of the binary systems, the general effect of commercial compositions is to vary the stability of the high temperature phase and significantly alter its transformation characteristics. Therefore it is expected that the characteristics of alloyed titanium can be altered by heat treatment. The present investigation was undertaken to survey the variations in hardness, microstructure, and mechanical properties as affected by heat treatment variables. Commercial Alloys lnvestigated The commercially available titanium-base alloys selected for this investigation and obtained as 3/4 in. diameter hot rolled and annealed bars are listed in Table I. The producer did not determine the chemical composition of the Ti-150A heats X1052 and L685. The early heats of Ti-150A, prefixed with the letter X, were melted with tungsten-tipped electrodes and cast in approximately 500 lb ingots 12 in. in diameter. The heats prefixed with the letter L were melted under a revised melting practice yielding approximately 1300 lb ingots 16 in. in diameter. The final annealing treatment was at 1200" to 1250°F for 1 hr. The MST-3A1-5Cr alloy was melted in an induction furnace and cast in 7 in. diameter ingots. This practice produces high carbon material and some variation in the amount of carbides was found in different sections of the same bar. The RC-130B alloy was melted with a carbon arc. This alloy showed approximately 2 pct carbides. No large segregation effects were observed.

Jan 1, 1955

By P. T. Carter, A. B. Murad, H. B. Bell

The distribution of manganese and oxygen between molten iron and FeO-MnO-SiO2 slags not saturated with SiO2 has been determined and used to calculate activities of MnO and SiO2 in MnO-SiO2 slags. These were found to be intermediate between the corresponding activities in FeO-SiO2 and CaO-SiO2 slags. THE equilibrium constant for the reaction between molten iron containing manganese and slags containing FeO and MnO, namely: FeO(slag) + Mn = MnO (slag) + Fe(1iquid) is most conveniently represented by: (MnO) Kmn== (FeOHMn] [Fe] is omitted, since for dilute solutions of manganese in iron, afe is approximately constant. When activities are used for (MnO), (FeO) and [Mn], Kmn is a true equilibrium constant and will apply to all types of slags, but when weight percentages are used, the values obtained for Kmn for different slags vary considerably. The equilibrium between iron containing manganese and relatively pure FeO-MnO slags has been investigated by several workers;1-4 the most recent investigation reported is that of Chipman, Gero, and Winkler.4 Their results, corresponding to slags containing up to 4 pct SiO2, were represented by the equation: 6440 log KMn = +2.95 In this case both slag and metal may be assumed to be ideal solutions, therefore the values of Kmn so derived represent the true constant for the manganese reaction. As the molecular weights of MnO and FeO are almost equal, weight concentration of these oxides may also be used in the above expression. Acidic oxides such as SiO2 and P2O5 affect the activities of FeO and MnO differently, so that when total concentrations by weight of FeO and MnO are used in the above expression for Kmn the values obtained are no longer constant. These will be denoted by K'mn. Körber and Oelsen5 investigated the manganese reaction using silica-saturated iron-manganese-silicate slags and found: 7940 logK'mn =3.172 This corresponds to Kmn values much higher than those obtained for pure oxide slags. The results of Körber and Oelsen, and Chipman, Gero, and Winkler are summarized in Table I. In the present investigation it was decided to study the manganese reaction using iron-manganese-silicate slags not saturated with silica, and which fell mainly between the orthosilicate join 2FeO . SiO2 - 2MnO.SiO2 and the FeO-MnO side of the ternary system FeO-MnO-SiO2. It has been shown" that the phases present in solidified slags of this part of the diagram are (FeO, MnO) and (2Fe0.Si02, 2Mn0-SiO,) solid solutions. Such an investigation would be of interest in deoxidation and also gives information on activities in liquid slags. This equilibrium has not as yet been investigated successfully. Hilty and Crafts7 attempted to examine this reaction in a rotating crucible furnace, but were unable to obtain suitable slag samples. In the present investigation one series of experiments was carried out in a rotating crucible furnace, because the success of Taylor and Chipman8 with such a furnace in their study of oxygen distribution between molten iron and (CaO, MgO) — FeO-Si0, slags suggested that it would minimize contamination of the slag by the crucible material. A second series was carried out using a wire-wound resistance furnace. Rotating Crucible Furnace Experiments In all cases the metallic charge consisted of 800 to 1000 g Armco iron containing 0.02 pct C, 0.03 pct Mn, and traces of Si. Manganese and silicon additions consisted of electrolytic manganese (99.9 pct)

Jan 1, 1953

By N. L. Carwile, G. W. Geil

A metallographic study was made of deformation twinning (Neumann lamellae) in ingot iron slowly deformed in tension at —196O and —150°C. The results showed that twinning is initiated mainly during the initial stages of plastic deformation. However, the breadth of the twins generally increases with increase in strain up to fracture of the specimen. HE results of an investigation of the effect of low I temperatures on the flow, fracture, strain hardening, and strain aging characteristics, and the true stress-strain relationship for specimens of ingot iron deformed in tension were reported recently.1,2 In that study specimens of ingot iron, in different initial conditions, were slowly deformed in tension at various temperatures between —196" and +100°C. A phenomenon observed, but only briefly reported, was the formation of deformation twins (Neumann lamellae) in the specimens deformed at the lower temperatures. Preliminary metallographic examination of these specimens indicated that there was apparently some correlation between the deformation twinning and the strain of the specimen and that this factor merited further investigation. Some recent investigations have indicated a direct association of Neumann lamellae and the fracture of the metal. Bruckner".' * postulated that the twinning immediately precedes the fracture and the latter is initiated by the stopping of the propagation of twins, either at subgrain boundaries or at other structural features. Tipper and Sullivan5 postulated that Neumann lamellae are initiated by the shock or a shock wave set up at fracture in the still highly stressed unbroken metal, and that a high rate of deformation is a necessary condition for their formation. An extensive metallographic examination, therefore, was made of many of the specimens used in the study of ingot iron at this Bureau and of additional specimens deformed to selected strains under various conditions. The present paper summarizes the effects of deformation under these conditions upon the breadth and scope of the deformation twins. Other features of twinning will be discussed in another paper. Material and Testing Procedure All of the specimens were prepared from one lot of ingot iron, as hot-rolled, of chemical composition other than iron as follows: C, 0.02 pct; Mn, 0.02; P, 0.005; S, 0.018; Si, 0.002; Cu, 0.10; 0, 0.058; and N, 0.002. Cylindrical tensile specimens with a 2 in. gage length were used. The reduced section was slightly tapered from each end; the diameter (0.438 in.) at the midsection of the gage length was about 0.003 in. less than at the ends. The specimens were finished to the final dimensions by grinding and polishing in the axial direction. The ends of the specimens were machined with 3/4 in. x 10 threads and the shoulder fillets were machined to a radius of 0.75 in. The specimens were extended slowly in tension (deformation rate less than 0.5 pct reduction in area per minute) while immersed in an appropriate cryogenic liquid maintained at the desired temperature. The minimum diameter of the specimen was measured during the test by means of a specially designed reduction of area gage accurate to 0.0001 in." A detailed description of the testing equipment and the method of maintaining the desired temperature is given in a previous paper.' Each specimen, after deformation to a specified strain at a selected temperature, was sectioned longitudinally and examined metallographically. The preparation of specimens for metallographic examination was carried out in the usual manner by mechanical polishing and etching. Results and Discussion Deformation twins in specimens of ingot iron extended slowly in tension at —196" and —150°C are shown by the representative micrographs of Figs. 1

Jan 1, 1954

By H. Bernstein

Grain coarsening tests were carried out on AI-4.5 pct Cu and AI-4.5 pct Si alloys. The effects of three variables, melt composition, pour temperature, and mold temperature, were determined. It was found that the macrostructure generally coarsened with increased pour and mold temperatures. Coarsening was extreme in the unrefined alloys but was retarded by the active grain refiners like titanium and columbium. The effect of boron was spectacular in suppressing coarsening tendencies. The results of the investigation support the carbide theory of nucleation as opposed to the peritectic theory. A REVIEW of prior work on the subject of grain refinement in the light metals indicated that, for the most part, it had been concerned with the positive effects of addition agents upon structure. This work produced various theories of refinement. Prominent among them were the peritectic theory and the transition element carbide theory. Certain inadequacies in the former theory appeared to be explained satisfactorily by the transition element carbide theory. In order to appraise these theories further, a study was undertaken of the permanence of grain refinement effects when subjected to thermal or chemical variations. It was considered that a study of this kind would be of practical value, since the light metals are exposed to thermal and chemical variations in the normal course of foundry operations. Accordingly, two alloys of commercial importance were selected for this investigation, the A1-Cu and A1-Si alloys. Experimental Procedure and Results Raw Materials: The base alloys, A1-Cu and A1-Si, were prepared from virgin aluminum and hardeners, induction melted, and pigged for remelting. The grain refiners, including titanium, columbium, zirconium, and boron, were added as master alloys. Analyses of the raw materials and base alloys are listed in Table I. Melting and Casting Practice: The metals were melted in clay-graphite crucibles in a Hevi-Duty pit-type resistance furnace, equipped with a Leeds and Northrup Micromax Recorder and Controller. After melt down of the base alloys, the additions were made. For each casting, approximately 80 grams of metal were poured into a preheated graphite mold held in a transite flask. The mold (Fig. 1) weighed 320 grams. To determine the effect of superheat upon structure, a set of melts was heated successively to four temperature levels, 705°, 775°, 845°, and 915°C, and test specimens were cast at each level. Prior to each pour, the furnace temperature was held constant for 20 min. Next, a set of melts was put through a superheat cycle and test specimens were cast at melt temperatures of 705" and 915°C successively. The crucibles then were cooled in air to 705°C as deter-mined by immersion thermocouple and additional test specimens were poured. The cooling interval was approximately 1 min. Another set of melts was furnace cooled with the cover slightly ajar and test specimens cast as before. The interval in this case was 15 to 20 min, varying with the superheat temperature. The mold temperature for all these tests was 400°C. To determine the effect of cooling rate without superheat upon structure, a set of melts was heated to 705°C, and castings were poured at mold temperatures of 205", 400°, and 540°C. An interval of 15 min elapsed between each pour. Mold temperatures were measured by a direct reading millivolt-meter, using a chromel-alumel thermocouple located at midwall thickness of the mold. To demonstrate the phenomenon of chemical coarsening, agents such as chromium and beryllium, which had acted as coarseners in a previous investigation,' were added to the A1-Cu alloy. Test speci-

Jan 1, 1955