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By Louis Moyd

In the preparation of illustrations to accompany reports of investigations concerning particle shapes of various natural and manufactured materials proposed for use as he aggretates in concrete structures, it was found that stereoscopic views of such materials presented the information much more effectively than the usual two-dimensional photographs. With the whole-hearted cooperation of the Photographic Section, a technique was developed in the petro-graphic laboratory of the Concrete Research Division, U.S. Corps of Engineers, at Clinton, Miss., for simple and rapid preparation of paired stereoscopic photographs and micrographs. The method requires only equipment which is usually part of the general stock of any laboratory, and it is adaptable for all purposes in which three-dimensional views would be found superior to ordinary photographs. Stereoscopic Photographs Where the subject or subjects to be photographed are small but not of microscopic size, an ordinary camera having a ground-glass focusing screen is set up vertically over the field. The material to be photographed is placed on a card, the surface of which is suitable as a background. The card is mounted so that it can be tilted about 5' above and below the horizontal plane, on an axis that coincides with the north-south diameter of the viewing screen. A simple way of achieving this is to mount the card on an object having a convex lower surface. In our laboratory, a watch-glass has been used, with plasticene as the mounting medium. The convex surface should be mounted on a support having a round opening on top and a flat bottom. A roll of tape, a petri dish, a beaker, or any other piece of equipment having the proper shape to give stability to the setup will be satisfactory. The arrangement of the required equipment is shown in Fig 1. For the best results, the subject to be photographed should be mounted in the plane of the axis of rotation. To achieve this effect, the depth of the mass of plasticene between the background and the watch-glass is varied so that when the watch-glass is tilted back and forth on the axis, the subject to be photographed does not appear to move back and forth across the viewing screen. A permanent piece of equipment, embodying an adjustable vertical screw, with lock-nuts, could be made for this purpose. The subject may be illuminated in any manner desired. It was found that the presence of shadows intensifies the effect of depth. In taking the photographs, the concave surface is mounted so that one side of the background card is about 5' below the horizontal, while the axis remains in the horizontal plane. The subject is focused carefully on the ground glass and a picture is taken. The card is then tilted so that the opposite side is about 5' below the horizontal, while the axis still remains in the horizontal plane. The subject is again put in focus and another picture is taken. After both pictures are developed and printed, they are placed side by side in the same orientation and examined with a stereoscopic viewer. If depressions appear where there ought

Jan 1, 1950

BASIC open-hearth slags have no obviously unique features when compared with slags from other metallurgical operations. Open-hearth slags form and exist at temperatures ranging from 2500 to 3100 F (1370 to 1700 C). In the course of an open-hearth heat, the liquid slag changes slowly and continuously in chemical analysis-the changes reflecting changes in temperature, changes in metal analysis, solution of charged materials, solution of feed additions, and solution of refractories. Through- out the process, the slag performs specific primary functions that are responsible for a large measure of the technical and economic importance of basic open-hearth steelmaking. It is intended in this chapter to present the chemical reactions by which the primary functions of basic open-hearth slags are carried out, to describe the physical and chemical principles governing the extent to which these reactions proceed and can be regulated, and finally to indicate the limitations imposed on other features of the process by the fulfillment of the primary functions. Basic open-hearth slags perform three primary functions. First and foremost, the slag provides a means of removing phosphorus from the liquid metal. The ability to remove phosphorus is largely responsible for the widespread use of the basic open- hearth process. Second, the slag serves as a controlling factor in transferring the oxygen required for the other oxidizing reactions. Third, the slag provides a means of eliminating some sulfur from the open-hearth bath. While the extent of sulfur removal is not comparable to that of phosphorus, the ability to remove some sulfur contributes much to the value of the process. The effectiveness with which the slag fulfills its functions varies with temperature and composition. Since operating temperatures are fixed largely by the specified nature of the product and the equipment, the speed and extent to which the functions of the slag are carried out depend largely on its composition. The necessity for regulating slag composition to accomplish the re-

Jan 1, 1951

By D. P. Pope, T. Vreeland

Observations of dislocations in zinc using Berg-Barrett X-ray micrography confirm the validity of a dislocation etch for (1010) surfaces. A technique for measurement of the depth in which dislocations can be imaged in X-ray micrographs is given. This depth on (0001) surfaces of zinc was found to be 2.5 µ using a (1013) reflection and CoKa radiation. BUCHANAN and Reed-Hill (B & RH) have recently questioned the ability of a dislocation etch to reveal all of the basal dislocations which intersect (1010) surfaces in annealed zinc crystals.' This etch was developed by Brandt, Adams, and Vreeland who conducted a number of different experiments to check its ability to reveal dislocations.2,3 B & RH prepared (0001) foil specimens for transmission electron microscopy from annealed crystals and observed dislocation densities of about l08 cm per cu cm in the foils, while the etch indicated densities of the order of l04 cm per cu cm in their annealed crystals. As this etch has been used in a number of studies of dislocations in zinc, it is of considerable importance to reassess its validity in the light of the B & RH results. The X-ray work reported here was undertaken to check the ability of the etch to reveal dislocation intersections on (1070) surfaces of zinc. The X-ray technique was chosen for this check because it could be applied to the as-grown crystals with a relatively small amount of specimen preparation. We believe that the possibility of accidental deformation in preparation of the bulk specimens is considerably less than that for thin foil specimens suitable for transmission electron microscopy. Unfortunately, basal dislocations are not visible on Berg-Barrett topo-graphs of (1010) surfaces, which are the surfaces on which the etch is most effective. Therefore, a one-to-one correspondence between the etch and X-ray observations could not be made. Basal dislocations near (0001) surfaces have been observed by Schultz and Armstrong4 using the Berg-Barrett technique, but they did not report the as-grown dislocation density observed in their crystals. We have applied the X-ray technique in this study to surfaces oriented from 1 to 2 deg of the (0001) to determine the basal dislocation density, and have compared this density with that observed using the etch on a (1070) plane of the same crystal. The X-ray observations permit determination of the depth in which basal dislocations can be observed under the diffracting conditions used. SPECIMEN PREPARATION High purity zinc crystals are very soft, so a good deal of care must be exercised in the preparation of observation surfaces. As-grown crystals approximately 2.5 cm in diam and 20 cm long were acid cut into 1.25 cm cubes. A thin slab was cleaved from an (0001) surface to produce an accurately oriented reference surface on the specimen. Some of the cubes were examined in the as-machined condition while some were annealed in argon at 410°C for 2 hr. Heating and cooling rates were less than 2°C per min. Some of the specimens were scratched on a (0001) surface with a razor blade to produce fresh dislocations. Approximately 2 mm of material was acid lapped from one face of a cube to produce a surface oriented between 1 and 2 deg from the basal plane and parallel to the [1210] direction. A (1070) surface was also acid lapped. The lap used a 1 to 3 pct solution of HN03 in water to saturate a soft cloth which was backed by a stainless steel plate. The cloth was moved over the crystal surface at a rate of 20 cm per sec while a normal force of about 4 g was maintained between the cloth and the specimen. As-lapped surfaces were examined as were surfaces which were chemically and electrolytically polished after lapping. The small angle between a lapped surface and the (0001) plane was measured to 0.1 deg using a Unitron microgoniometer microscope (the cleaved surface was used as a reference in this measurement). The microscope was modified so that the intensity of reflected light could be continuously monitored on a meter. This modification produced nearly a ten-fold increase in the reproduceability of orientation readings. OBSERVATIONS The Unitron Microgoniometer observations indicated that the lapped surfaces had a terraced structure with the terraces quite rounded and spaced about 0.1 mm in the [1010] direction. The maximum change in slope between terraces was 0.25 deg, indicating a terrace height of about 0.1 µ. A Unitron measurement of the average angle between (0001) and a lapped surface was checked by micrometer measurement of the specimen and found to agree within 0.1 deg. The Berg-Barrett micrographs using (1013) reflections and CoKa radiation5 revealed subboundaries, short dislocation segments, spirals, and loops near the surfaces which were oriented from 1 to 2 deg of the (0001). Micrographs of surfaces prepared by lapping appeared very similar to those of the chemically and electrolytically polished surfaces. The loops and spirals were not extinct in (1013) or (0002) reflections, indicating that they have a nonbasal Burgers vector. Extinctions of the short, straight dislocations indicated that they belonged to an (0001)(1210) system. Fig. 1 is an example of a micrograph which shows a subboundary, and dislocation segments which are pre-

Jan 1, 1970

BASIC open-hearth slags have no obviously unique features when compared with slags from other metallurgical operations. Open-hearth slags form and exist at temperatures ranging from 2500 to 3100 F (1370 to 1700 C). In the course of an open-hearth heat, the liquid slag changes slowly and continuously in chemical analysis-the changes reflecting changes in temperature, changes in metal analysis, solution of charged materials, solution of feed additions, and solution of refractories. Throughout the process, the slag performs specific primary functions that are responsible for a large measure of the technical and economic importance of basic open-hearth steelmaking. It is intended in this chapter to present the chemical reactions by which the primary functions of basic open-hearth slags are carried out, to describe the physical and chemical principles governing the extent to which these reactions proceed and can be regulated, and finally to indicate the limitations imposed on other features of the process by the fulfillment of the primary functions. Basic open-hearth slags perform three primary functions. First and foremost, the slag provides a means of removing phosphorus from the liquid metal. The ability to remove phosphorus is largely responsible for the widespread use of the basic open-hearth process. Second, the slag serves as a controlling factor in transferring the oxygen required for the other oxidizing reactions. Third, the slag provides a means of eliminating some sulfur from the open-hearth bath. While the extent of sulfur removal is not comparable to that of phosphorus, the ability to remove some sulfur contributes much to the value of the process. The effectiveness with which the slag fulfills its functions varies with temperature and composition. Since operating temperatures are fixed largely by the specified nature of the product and the equipment, .the speed and extent to which the functions of the slag are carried out depend largely on its composition. The necessity for regulating slag composition to accomplish the re-

Jan 1, 1964

By F. A. Olson, J. S. Cho, M. E. Wadworth

An infrared study of the states of hydration of MgSO4 revealed a hitherto unreported metustable dehydration state in the temperature range just below that of the stable anhydrous salt. Infrared, thermogvavimetric, and X-ray studies indicated that the state had bisulfate and hydroxide characteristics. A formula fitting all of the experimental evidence is hydrogen-bonded Mg(OH)HSO4, or possibly the chemically equivalent 1/2 Mg(HSO4)2. Mg(OH)2 Infrared spectra of the 7, 6, 1, and OH2O states of hydration are given as well as the spectra of the newly found state which is called the bisulfate state in this paper. Thermogvavimetric studies as a function of water-vapor pressure gave an enthalpy value of 16 cal per mole, and 25 eu for the conversion of the bisulfate state to that of the anhydrous salt. The free energies of the transition varied from 1.76 to 0.18 kcal per mole as the transition temperature varied from 279° to 336°C. The bisulfate state is believed to be of importance in catalytic dehydration phenomena using MgSO4 and may be of interest because of properties in sulfating reactions differing from those of the monohydrate. Also, an analogous state may exist as a hydration state of other sulfates. ANHYDROUS magnesium sulfate is a catalyst with a high degree of specificity for the dehydration of secondary alcohols.1'2 The only product obtained at the usual operating temperature range of 360" to 400°C other than the corresponding olefin is water. Thus, the bonding of water to the anhydrous salt as well as an understanding of the hydration states of MgSO4 became of interest in attempting to understand the catalytic mechanism better. In the process of this study using infrared, X-ray, and thermogravimetric analyses a dehydration state not previously reported was found. Proof of its existence, the nature of the state, and its role, if any, in the catalytic mechanism then became of interest. This paper is concerned principally with the first two aspects of this problem. Magnesium sulfate forms many hydrates, of which the monohydrate and heptahydrate (Epsom salts)3 are most important. Other hydrates4 are known with 1-1/4,5 l-1/2, 5 2, 4, 5, 6, and 12 H2O's. The hydrates with 1-1/4, 1-1/2, 2, 4, and 5 HzO's are unstable. The anhydrous salt is very stable thermally and has a strong affinity for water. It can be heated without additional decomposition to about 800°C.' At 900°C noticeable decomposition takes

Jan 1, 1964

By D. W. Mitchell

While pure magnesium does not corrode rapidly the presence of even very small quantities of certain other metals accelerates corrosion remarkably. Because magnesium is such an electropositive metal (E° = +2.34 volts)' the presence of substances that permit galvanic activity is extremely undesirable. It is stated that as little as 0.017 pct of iron in magnesium causes a marked increase in the corrosion rate.2 Magnesium and its alloys are commonly melted in iron or steel pots; consequently the commercial metal has the opportunity to pick up iron if it has that tendency. Quantitative data on the solubility of iron in magnesium are of practical as well as of academic interest. Some magnesium alloy castings tend to have large grain size unless suitable treatment is given to the liquid metal just prior to casting. One method of grain refinement consists of "superheating" the metal just prior to casting. "Superhcating" is heating to a temperature 100 to 200°C above the casting temperature and holding for a short period of time. Superheated metal has a much finer grain than unsuperheated metal. A satisfactory explanation of this phenomenon has not been forthcoming. One hypothesis attributes the grain refinement in the superheated metal to precipitation of some phase that is mis-cible at the temperature of superheat but not miscible just below the casting temperature and which is thought to provide nucleating centers for crystallization. This phase, if the foregoing hypothesis is correct, would probably be a metal or an intermetal-lic compound. Of the metals likely to be found in magnesium alloys, iron is one of the most likely candidates for the role of grain refiner for iron occurs in small quantities in all commercial magnesium and magnesium alloys. A. Beck³ states that magnesium, in either the liquid or the solid state, cannot dissolve iron and that iron exists as a finely divided suspension which does not settle out because of its extreme fineness. Beck also presents a micrograph showing a large iron dendrite in magnesium but does not comment on the crystallization of iron from a solution supposed to contain none. It is doubtful that dendrite formation from suspended iron -would be possible. This suspension theory is the basis of patents4 covering the removal of iron from magnesium by allowing the iron particles to act as nuclei for the formation of primary magnesium crystals which can be separated by segregation. Also covered by patents is the removal of iron by the addition of another metal to form an intermetallic compound with the iron which, because of its increased particle size, is able to settle out of the molten metal. Here, too, quantitative information on the extent of iron solubility in liquid magnesium is of more than theoretical interest. M. Hansen,5 in his critical collection of the data on binary alloys, states that iron is not soluble in liquid magnesium. W. R. D. Jones% ho studied the physical properties of low iron magnesium alloys observed a 57°

Jan 1, 1949

By John Chipman

THE equilibrium between titanium in liquid iron and titanium oxides has been studied by Hadley and Derge.' They have shown that a minimum occurs in the oxygen content of the metal between 0.1 and 1.0 pct Ti, but its exact level was not fixed. The oxide phase in this minimum range was principally TiO, containing some iron oxide. On the other hand, for concentrations above 5 pct Ti the oxide phase was TiO and the oxygen content of the metal increased rapidly with increasing titanium content. From their observations and the free energy of Ti0 and TiOz reported by Kelley and ah,' it is possible to derive approximate values for the activity coefficients of titanium and oxygen in liquid iron-titanium alloys and the conditions for equilibrium in steel deoxidation. Neglecting possible departures from the stoichio-metric composition, the condition of equilibrium with Ti0 may be expressed as follows: TiO= Ti_ + C) log K = log NTl + log [% O] + log yTi + log fQ Here as a matter of convenience the titanium concentration is expressed in mol fraction while that of oxygen is in weight percent. The activity coefficients are yTi and fo. The value of log fo approaches zero as the concentration of titanium approaches zero, while that of log yTi approaches a limiting value of In Fig. 1 the sum of log NTi + log [% 0] is plotted for mol fractions of titanium between 0.05 and 0.50. The smooth extrapolation gives a value of —4.15 which corresponds to log K — log y:i . We therefore have the equation: According to the data of Kelley and Mah,' the free energy of formation of Ti0 at 1900°K is -82,400 cal. The free energy of solution of half a mol of oxygen in iron, taking 1 pct 0 as the standard state, is -29,300 cal. From these we write: From these equations the activity coefficient of titanium at infinite dilution in iron is found to have a value of 0.011 or log y°. = -1.96. This is only an approximate value but is certainly better than the estimated 0.05 in Basic Open Hearth Steelmaking To obtain approximate values for the activity coefficient at intermediate concentrations, it is as- sumed that log yTi(l is constant and equal to -1.96. This corresponds to a value of E;: = 3 In yTi/aNTi = 9.0 which is not unreasonable We now have the value for the activity of titanium and the percentage of oxygen in equilibrium with TiO. It is therefore a simple matter to calculate the activity coefficient of oxygen which is shown in Fig. 2. The effect of titanium is rather large, being comparable to that of vanadium, but not so large as has been reported for aluminum. The slope of the line at low concentrations corresponds to E: = aln yo/ aNTi=-0.37. At concentrations below 1 pct Ti the deoxidation constant and oxygen content for equilibrium with TiO, can be calculated from the foregoing data and the free energy of formation of TiO,, which according to Kelley and Mah is -144,600 cal at 1900°K. The free energy of solution of oxygen has been given; that of Ti to form a dilute solution with 1 pct Ti as the standard state is RT ln (0.5585 y:i/47.9) = -34,800 cal. From these data we find (for concentrations in wt pct): TiO, = Ti + 2 -0; hFTgO0= + 51,200 cal log if =-5.89 The interaction coefficients, concentrations being expressed as wt pct, are e;: = alogfTi/a[% Ti] = +0.048 and e:' = a logfo,/a[ojTi] = -0.187. The calculation gives the following results which, in view of the impurity of the oxide phase, must be regarded as approximate:

Jan 1, 1961

By Victor A. Phillips

A series of strips of a Ni-12.7 at. pct A1 alloy were Prepared containing cubical y'(NisAl) precipitates with edge lengths from 60 to 500A. A particle-free solution-tveated strip was included for cornparison. They weve cold-rolled 95 pct and the effects of particle size on the isochronal (1/2 hr) annealing behavior between 300° and 950°C studied (by hardness and light and electron microscopy). It ulas inferred that the particles deformed with the lnatvix becoming lamellae which remained coherent. Comparison with published data fov pure nickel showed that aluminum greatly re-tavded softening and recrystallization, but it made little difference whether or not particles were present. The presence of pakticles led to a heterogeneous distribution of precipitates after annealing at 700" to 750°C. Recovery was not detected. Recrystallization occurred by the growth of new grains into unrecrys-tallized material. In a previous study by the author,' the growth of Ll2-type ordered yl(Ni3Al) precipitates was followed in Ni-12.7 at. pct A1 alloy as a function of aging at 600" and 700°C. The particles were showo to be cubical in shape in all sizes from 50 to 3000A and remained coherent. This work was used as a guide in preparing the starting structures for the present study of the effect of these particles on the annealing behavior of heavily cold-rolled strip. Another question of present interest was whether dislocation and particle hardening were additive, since the structures before rolling ranged from solution-treated to peak-aged to overaged. Also, precipitation might occur on annealing after cold-rolling. Reference may be made to other papers2"5 for previous work in this relatively unexplored field and only some recent work will be mentioned her:. phillips2 studied the effect of deformable 0 to 590A-diam cobalt particles on a Cu-3.23 pct Co alloy rolled 95 pct and found that the particles, which rolled out into thin lamellae, impeded softening and recrystallization. Tanner and servi3 likewise studied the annealing of cold-swaged Cu-2 pct Co alloy containing 150A-diam particles and found impeding effects. Haessner et a1.,4 on the other hand, found that incoherent 2-p-diam non-deformable particles of B4C (0.04 vol pct) tended to increase the rate of recrystallization of copper rolled up to 95 pct reduction. They attributed this to the formation of new grains at the particle interfaces. Humphreys and artin' found that nondeformable silica particles in copper rolled to 30 pct reduction accelerated recrystallization if the particle spacing was large and retarded it if the spacing was smaller. Haessner et a1 4 also studied a rolled Ni-Cr-A1 alloy; however, the particles of y'(Ni3Al)-type precipitate were not put in before rolling, but separated during the isothermal annealing at 750°C. No previous work appears to have been carried out on the effect of y' (Ni3A1) particles on the annealing of Ni-A1 alloy. Hornbogen and ICreye7 redetermined the solubility c of aluminum in nickel as a function of temperature T and showed that it was given by c = 32.6 exp(-1940/RT). This relation gives aluminum solubilities of 15.1, 14.2, 12.0, and 10.7 at. pct at 1000°, 900°, 700°, and 600°C, respectively. The phase precipitated from the nickel-rich solid solution is fcc y1 (Ni3A1) which has a Cu3Au -type ordered structure8 and remains ordered up to 1000°C.B EXPERIMENTAL PROCEDURE The alloy used was identical with that used before. Chemical analysis showed 6.27 wt pct (12.71 at. pct) Al, the principle impurities being 0.065 pct Fe, 0.022 pct Co, 0.020 pct Cu, and 0.004 pct C. Bar stock of 1 in. diam was cold-swaged to % in. diam, cold-rolled to 0.300-in.-thick strip, and annealed at 900°C in dry hydrogen. It was cold-rolled to 0.100-in. thickness and solution-treated for 1 hr at 1000°C while sealed in a quartz tube in argon, quenching in iced brine with the aid of a device to snap off the nose of the tube. Lineal analysis gave an average grain size of 0.055 mm. Pieces of strip were aged at 700°C in vacuo for 30 min, 51/4 hr, and 1 week to produce nominal average particle widths of 60, 150, and 500A, respectively, as known from the previous work.' The average diamond pyramid hardness was determined. The heat-treated strips were rolled from 0.100 to 0.005 in., a reduction of 95 pct, and the rolled strips stored at about -5°C. Small pieces were annealed within 1 week for 30 min at temperatures from 300° to 950° ±2°C in a horizontal vacuum furnace. Strips were withdrawn into a cooling zone, giving an estimated initial cooling rate from 950°C of about 50°C per sec. Average diamond pyramid hardnesses were determined on a lightly electropolished spot on the surface of each strip using 300-g load. Each point on the softening curves represents a separate annealed specimen. Sections containing the rolling direction were examined by optical metallography. Selected specimens were electrothinned to the center plane' and examined by transmission at 100 kv in a Siemens Elmiskop I electron microscope. It is well-known that changes in the structure tend to occur when a deformed strip is electrothinned below a thickness of a few hundred angstroms, although this is less serious with a material such as nickel which has a high melting point, and also is apt to be less serious when particles are present. Observations were nevertheless confined to thicker regions of the foils with estimated thicknesses over 1000A. No changes were observed due to beam exposure.

Jan 1, 1968

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 R. M. Brown, W. J. Murphy, B. M. Kapadia

An investigatiott was conducted to study the influence of nitrogen, titanium, and zirconium on the boron llardenabilzty effect in a low-carbon constructiona2 alloy steel. The experimental steels investigated exhibited a significant variation in hardenability, the variation being dependent on the interactions of boron, titanium, and zirconium with the nitrogen. Only the boron not combined with nitrogen was effective in increasing hardenability. Titanium, and with lesser effectiveness zirconium, combined with available nitrogen, thereby protecting the boron. The hardenabil-ity effect mas related to an empirical expression for the "effective" boron content, P, deduced from experimental evidence of these interactions. The hardenabzlity effect reached a maximum at about 0.001 wt pct 0, and decreased somewhat as P increased further. The physical understanding of this relationship is discussed. FOR many years boron has been added to steels to obtain high hardenability. Although a great deal of research has been conducted on boron-treated steels, certain aspects of the boron hardenability effect have not been fully understood. For instance, the magnitude of the hardenability effect has been observed to vary markedly, depending on the steelmaking technique, even when the amount of boron in the steel was essentially constant. Furthermore, the optimum amount of this element to be added has not been definitely established. A better understanding of the boron hardenability effect is essential because too small an addition of boron is likely to be ineffective, while an excessive amount can cause brittleness'' and hot shortness. The findings of earlier investigations have shown that the hardenability effect cannot be consistently related to the amount of boron added or retained in the steel. Grossmann observed that in a 0.60 pct C steel the hardenability increased to a maximum with mold additions up to about 0.0025 pct B and then decreased with larger additions. Other investigators5 likewise reported a maximum in the hardenability at about 0.003 pct B. Crafts and Lamont, however, found that in commercial open-hearth heats of medium-carbon steel the hardenability increased linearly with boron up to 0.001 pct and remained essentially unchanged with larger percentages up to 0.006 pct. Other investigators7,' also observed a rather constant hardenability effect in the range about 0.0005 to 0.0035 pct B. These observations and other evidence suggest that the effectiveness of boron in increasing hardenability probably depends, in addition to the amount, on the form of boron retained in the steel, this form being influenced by the presence of other elements. Both oxygen and nitrogen apparently exert the strongest influence on the hardenability behavior, since, at the temperature of liquid steel, boron readily combines with these elements, thereby losing its effectiveness as most experimental evidence seems to indicate. For consistent recovery of the boron effective in increasing hardenability, it is necessary that the oxygen and nitrogen in the steel be either reduced to extremely small amounts by the steelmaking practice or neutralized by combination with other elements before the addition of boron. The importance of achieving adequate deoxidation prior to the addition of boron in order to realize the full hardenability effect of boron has been sufficiently emphasized by earlier investigators. Digges and Reinhart' and others have investigated the role of nitrogen and have shown that nitrogen also interacts with boron and reduces or nullifies altogether its effect on hardenability. Moreover, their work also demonstrated that the addition of strong nitride formers such as titanium and zirconium reduce the deleterious effect of nitrogen on boron hardenability by combining with nitrogen to form stable nitrides. Another element which has a pronounced influence on the boron hardenability effect is carbon. It has been shown7'10 that the hardenability effect of boron diminishes with increasing carbon content, and becomes almost negligible at the eutectoid composition. This observation is useful in comparing the potential increase in hardenability from boron of steels with different carbon contents, but is not relevant to a study of the effects of normal steelmaking variables. The amounts of oxygen and nitrogen in steel vary with the steel composition and steelmaking practice employed. Most commercia1 low-alloy steels are fully deoxidized by the addition of silicon and aluminum, or other strong deoxidizers, which adequately protect the boron from oxidation. In addition, one or more of the elements such as titanium or zirconium are usually added, either separately or in combination with boron, in the form of complex ferroalloys, to protect boron from combination with nitrogen in the steel. However, the actual amount and type of addition employed for a given processing requirement are usually selected by trial and error, and have a rather limited range of applicability. As a result, substantial variations in the hardenability of boron-treated steels are often observed in practice, particularly when the nitrogen content of the steel is a significant processing variable. These variations might therefore be reasonably attributed to the interactions between boron, nitrogen, and titanium or zirconium present in the

Jan 1, 1969

By H. Biloni, R. Morando

The surface structure and substructure of Al-Cu alloys solidified as conventional ingots and under particular conditions such as those used by Bower and Flemings are studied. The influence of lampblack coating on the mold walls is especially considered and the results compared with those obtained in copper and graphite molds where no coatings exist. When high heat extraction conditions exist the observations show that mechanism of copious nucleation is responsible for most of the chill zone. When the heat extraction through the mold walls is low, a coarse grain structure with dendritic morphology arises, with a size that depends on the degree of convection present, analogous to that analyzed by Bower and Flemings. In both cases the effect of the convection on the macroscopic and microscopic appearance is discussed. The ingot macrostructure consists of one or more of three zones: "chill zone", "columnar zone", and central "equiaxed zone". The mechanism of the columnar-equiaxed transition has been subject of considerable interest and at present at least three theories exist about the formation of the equiaxed region: 1) the constitutional supercooling theory1 maintains that the equiaxed crystals nucleate after the columnar zone has formed, as a result of the constitutional supercooling of the remaining liquid; 2) chalmers2 pointed out, however, that there were several objections to this proposal, and that consideration should be given to the possibility that all the crystals, equiaxed as well as columnar, originated during the initial chilling of the liquid layer in contact with the mold; 3) Jackson et aL3 and O'Hara and ~iller~ suggested that a remelting mechanism of the dendrite arms is responsible for the formation of the equiaxed region. After the work of Cole and Bolling and other authors6 it became evident that convection (natural, reduced, or forced) plays a very important role in the transition from columnar to equiaxed and on the size of the resultant equiaxed structure. Until recently the accepted explanation of the chill zone was that it occurs as a result of copious nucleation in the liquid layer in contact with the mold walls.798 The columnar region is a subsequent result of the growth of favorably oriented grains and, as a result of a selection mechanism studied by Walton and Chalmers,9 elongated grains with marked texture are formed. Recently, however, Bower and Flemings" using an ingenious laboratory experiment introduced the idea that the "copious nucleation" mechanism is not responsible for the formation of the chill zone and that the presence of convection, introducing some form of "crystal multiplication", plays a decisive role in the formation of the chill zone. Unfortunately, it is important to consider that for their conclusions Bower and Flemings extrapolated the results obtained in their special experiments to the case of conventional ingots, and that these authors only analyzed the macrostructures of the specimens. Let us consider the work by Biloni and chalmers" concerning predendritic solidification. These authors were able to show that a study of the segregation substructure of A1-Cu gives information about the nucleation and growth of crystals formed in contact with a cold surface. A spherical predendritic region characterizes the first part of every grain nucleated in contact with the surface as a result of the chill effect. The aim of this paper is to elucidate through the observation of the segregation substructure the conditions under which (in the Bower and Flemings type of experiments and in conventional ingots) either the nucleation or the multiplication mechanism gives rise to the structure in contact with the mold walls. I) EXPERIMENTAL TECHNIQUES The experiments were performed on two alloys: Al-1 wt pct Cu and A1-5 wt pct Cu. The purity of the aluminum was 99.99 pct and the copper 99.999 pct. The results obtained with both alloys were similar. In the Bower and Flemings type of experiments the apparatus employed to obtain rapid solidification against a surface was similar to that used by those authors. The liquid was drawn by partial vacuum into the thin section mold cavity. Plate casts were 5 cm wide and usually 7.5 cm high. The thicknesses of the cast were 0.1 and 0.3 cm. Two different materials were used for the mold, copper and nuclear-grade graphite. The internal mold surfaces were polished and left uncoated for some experiments. In other experiments, the copper or graphite surface was coated with a thin film of lampblack material. In some of these particular experiments one of the mold walls was left with an uncoated region (usually in the form of a cross). The conventional ingots were cast in graphite or copper molds. In different experiments the mold walls were sometimes uncoated or coated with lampblack material. The results obtained in conventional and Bower and Flemings copper molds were compared with those obtained with copper molds coated with a very thin film of graphite; the results obtained were essentially similar. The size of the conventional ingots was 5 cm diam and 7 cm high in all cases. The cast surfaces produced by the Bower and Flemings type of experiments and conventional methods were observed macroscopically and microscopically without any metallographic preparation. As Biloni and Chalmers showed," the observation of the chill surface can give considerable information about the structure and segregation substructure.

Jan 1, 1969

By Joseph Kelly

DEPLETION of anthracite resources in Lackawanna County, Pennsylvania, has forced the recovery of coal tracts formerly considered unminable. Chief among these are the large areas of outcrop coal lying on a heavy pitch on the mountain sides, abandoned thirty or forty years ago as too costly and too hazardous to mine. The old companies, in many instances, stopped their chambers 80 or 100 ft. hack from the actual outcrop, and then dropped back another 100 ft. before starting to rob the pillars, thus leaving large areas of coal. Recovery of this coal naturally leads to mining hazards not encountered years ago. The first step is to sink a slope from the surface with the pitch of the vein, and then drive counters across the pitch as close to the robbed area as proper protection for the lower side of the -counter will permit. Next, the old chambers are advanced or new ones driven in the solid, as the conditions require. As these chambers approach the outcrop the real safety work begins-that of preventing the miner from being buried or locked in by an avalanche of mud.

Jan 1, 1936

By R. J. Schaefer, M. E. Glicksman

Direct, in situ observation of metallic alloys during solidification reveals that in some systems crystal multiplication occurs by means of a dendrite remelt-ing process, without the necessity of independent nu-cleation events. The solidification is observed on optically clean surfaces of supercooled molten alloys. In one system (Sn-3 wt pct Bi) , a fine-grained equi-axed structure is produced by the complete disintegration, via remelting, of transient initial dendrites. In contrast, several other metallic systems show little or no evidence of remelting activity. The behavior of these metallic systems is consistent with that expected on the basis of published observations of dendrite remelting in transparent organic alloy systems. It is concluded that dendrite remelting is an important crystal multiplication process in some metallic alloys, but is velatively unimportant in some other alloys. The detachment of dendrite branches from the primary stem by a remelting process has been observed by Jackson, Hunt, Uhlmann, and seward1 during the solidification of some transparent organic alloy systems. The branches released by this remelting process can drift free of the primary stem and serve as seeds for the formation of numerous independent crystals. This proliferation of new crystals does not require repeated independent nucleation events, but only modification of the initial solidification structure. The mechanism of dendrite remelting, as described by Jackson et a1.1 and by Tiller and O''Hara,2 is dependent on the transport of both heat and solute and is, therefore, expected to operate only in alloys, and not in pure substances. It was found in the transparent organic alloy systems that remelting occurred primarily in the constricted regions near the base of dendrite branches, where the branches had grown through the solute-rich liquid layer surrounding the primary dendrite stem. It was suggested that impure solid had formed in these constricted regions at a relatively low temperature and had later become unstable as further solidification of purer material raised the temperature. Jackson et al. proposed that the equiaxed zone of metal alloy castings originates from a similar remelting process, and they demonstrated the operation of this mechanism in "ingots" poured from saturated salt solutions. Although these substances were especially selected to show solidification behavior similar to that of metals,3 the remelting mechanism has, until now, been subject to the objection that its operation has not been directly observed in a metallic system. Considering the fact that thermal properties such as diffusivity may differ by as much as two or three orders of magnitude between metals and organic compounds, this objection was considered to be possibly serious. A method has now been developed by which it is possible to make direct optical observations of metals during solidification. This method has revealed transient solidification processes which would not be detected by metallographic examination of totally solidified specimens. In particular, a grain-refinement process in a metallic alloy has been observed to be a direct consequence of the total disintegration of an initial dendritic structure. The influence of an initial, transient solidification structure on the final solidification structure is demonstrated here for the first time in a metallic system. I) EXPERIMENTAL The technique of using a transparent flux to maintain an optically clean metal surface during solidification has been described by us elsewhere.* The

Jan 1, 1968

By H. G. Doll

A new logging method, called induction logging, is described; it measures the conductivity, or resistivity, of the strata traversed by a bore hole. The apparatus, which is briefly described, comprises a coil system coupled with the ground by induction, so that there is no need for direct contact with the mud, or with the formations. For that reason, the method is particularly useful in oil base mud, and it was first applied to that case. A discussion concerning the respective contribution of each region of ground in the vicinity of the coil system is given. Through the concept of a geometrical factor, it is demonstrated that the total signal, which measures the apparent conductivity, is the summation of the partial products — geometrical factor by conductivity — for the different regions of the ground under consideration. It is thus possible to establish, for each type of coil arangement, investigation characteristic as well as departure curves and correction charts for beds of various thicknesses. This should greatly facilitate the quantative interpretation of the field logs, even in cases where such interpretation would have been difficult on electric logs taken with electrodes. It is also shown that focussing combinations of coils can be used, which give, directly on the recorded log, a value close to the true conductivity, even for fairly thin beds, thus eliminating in most cases the need for corrections. These focussing systems also give a sharper definition of the boundaries between successive strain.

Jan 1, 1949

By H. G. Doll

A new logging method, called induction logging, is described; it measures the conductivity, or resistivity, of the strata traversed by a bore hole. The apparatus, which is briefly described, comprises a coil system coupled with the ground by induction, so that there is no need for direct contact with the mud, or with the formations. For that reason, the method is particularly useful in oil base mud, and it was first applied to that case. A discussion concerning the respective contribution of each region of ground in the vicinity of the coil system is given. Through the concept of a geometrical factor, it is demonstrated that the total signal, which measures the apparent conductivity, is the summation of the partial products — geometrical factor by conductivity — for the different regions of the ground under consideration. It is thus possible to establish, for each type of coil arangement, investigation characteristic as well as departure curves and correction charts for beds of various thicknesses. This should greatly facilitate the quantative interpretation of the field logs, even in cases where such interpretation would have been difficult on electric logs taken with electrodes. It is also shown that focussing combinations of coils can be used, which give, directly on the recorded log, a value close to the true conductivity, even for fairly thin beds, thus eliminating in most cases the need for corrections. These focussing systems also give a sharper definition of the boundaries between successive strain.

Jan 1, 1949

By A. U. Seybolt, R. A. Oriani

The pressure and composition of the solid solution of nitrogen in chromium has been investigated as a function of temperature. The partial molal heat of solution, the activity of nitrogen in solid solution, and the partial molal free energy have been computed. The enthalpy of three reactions is summarized: N + 2Cr(3) = Cr2N(4), H° = —28,500 cal per mol 1/2N2(p) + 2Cr(8) = Cr2N(3)1 ?HO = —23,500 cal per mol 1/2N2(p) = N, ?H° = 5,000 cal per mol. WORKERS who have investigated the problem of the preparation of ductile chromium have been concerned with the effects of minor amounts of impurities upon this refractory metal. In particular, gaseous impurities such as oxygen and nitrogen have been suspected of causing trouble because of many known cases where brittleness in transition elements has been definitely traced to these elements. In order to evaluate the possible effects of the nitrogen impurity in chromium, it appeared desirable to learn something of the dependence of the nitrogen solubility upon temperature and pressure. Previous Work Maier3 in 1942 in his treatise on sponge chromium reviewed the thermodynamics of the chromium nitrides Cr2N and CrN, both of which are well-estab- lished compounds, but which apparently have not been studied in much detail. Maier presented some free energy-temperature expressions for the reactions 2 Cr,N = 4 Cr + N2 [1] and 4 CrN = 2Cr2N + N2 PI based mainly on the work of Satoh2 and Sano." Pearson and Ende3 in 1953 briefly reviewed the thermodynamics of Cr-N using the same data available to Maier, and also derived a free energy-temperature expression for reaction 1. Of the two nitrides mentioned previously, the present investigation was only concerned with Cr,N, since this nitride is in equilibrium with the metal phase, at least over most of the temperature range, Blix4 in 1929 established the structure of Cr2N (close-packed-hexagonal) by X-ray diffraction methods, and indicated that its homogeneity limits extend from 11.2 pet to 15.1 wt pet N. At a temperature of 1200°C, 38 mm Hg from the Maier expression and 43 mm Hg from the Pearson

Jan 1, 1957

By D. O. LIVINGTON

(Presented by invitation at a meeting of the Spokane Local Section of the institute, Feb. 17, 1912.) THE Pilares de Nacozari mine is located in Sonora, 75 miles south of Douglas, Ariz. The town of Douglas is on the International Boundary and is the place at which the ores from the Bisbee mines are smelted. The Moctezuma Copper Co. owns the mines at Nacozari, and the copper-concentrates shipped to the Copper Queen smelter, at Douglas, make a good smelting-mixture with the Bisbee ores. The Moctezuma Copper Co. and the Copper Queen smelter at Douglas, as well as the railroad from Douglas to Nacozari, are owned by the Phelps Dodge Co. The ore-deposit occurs in the form of a large ellipse with a major axis of approximately 2,000 ft. and a minor axis of about 600 ft., the major axis bearing about 90 W. of N. The whole of this ellipse is more or less mineralized; the surface being principally an iron gossan with some occasional copper-stains. Below the oxidized zone, which is not more than 50 ft. deep, the minerals are pyrite and chalcopyrite, with occasional seams of chalcocite, the latter being rare. There appears to be no well-marked zone of secondary enrichment, the oxidized gossan over the greater part of the deposit changing suddenly to what is apparently the original unaltered ore. The copper-values are concentrated around the perimeter of the ellipse, and it is principally around this perimeter that the mining is done. The ore mined averages a little more than 3 per cent. of copper, with a small amount of silver, less than an ounce per ton. Some ore of considerably higher grade than this has been shipped, however, but the above average is of the mill-run. The ore is wider near the two ends of the ellipse than along the sides, and is mined in some cases to a width exceeding 200 ft. The country-rock for the first vertical 500 ft is a volcanic acid breccia, probably rhyolitic; below

Sep 1, 1912

Impact Transition Temperatures of Some Pearlite-Free Mild Steels As Affected by Heat Treatments in the Alpha Range H. P. Tardif (Canadian Armament Research and Development Establishment, Quebec)—Mr. Josefsson is to be congratulated for his study of a most difficult subject and for the theories proposed to explain the re- sults. One point which, however, might not have received proper attention is that of the shape of the stress-strain curve at the yield point of quenched and furnace-cooled specimens and of specimens strain aged

Jan 1, 1956

By George V. Smith, Daniel E. Sonon

Various mechanical properties of the Ni-Co-C alloy system were investigated to delineate the strengthening effect of carbon. Carbon concentration, cobalt concentration, vain size, temperature, and strain rate were varied so that thermal activation analysis and the Hall-Petch analysis could be used to evaluate the strengthening effect of carbon. Increasing carbon increased the strength of nickel and a Ni-60 pct Co alloy , with the effect becoming more pronounced at lower temperatures. Yield stress depended linearly on carbon concentration in nickel, but it depended on the square root of carbon concentration in the Ni-60 pct Co alloy. The Hall-Petch slope of nickel increased with carbon concentration; however, that of the Ni-60 pct Co alloy did not. The yielding behavior of these alloys was sensitive to composition, grain size, and temperature. Cobalt eliminated serrations in the flow curve of carbon-containing nickel at 300' and weakened them severely at higher temperatures. Pairs, or clusters, of carbon atoms appear to be responsible for the observed strengthening behavior. FLINN' conducted several experiments with carbon in nickel in an effort to provide information on the strengthening effect of interstitial impurities in solid solution in fcc metals and alloys. Strengthening which increased with decreasing temperature led him to conclude that carbon causes Cottrell locking in nickel. Fleischer2 analyzed Flinn's data and calculated that the strengthening effect of carbon in nickel was smaller by a factor of fifty than the strengthening effect of carbon in a! iron. Fleischer2 termed the magnitude of strengthening of carbon in nickel "gradual" and that of carbon in a! iron "rapid". He attributed "gradual" hardening to hydrostatic strains and localized changes in modulus of elasticity around solute atoms, whereas he attributed "rapid" hardening to tetragonal strains around solute atoms. Sukhovarov et a1.3-7 reported strain aging and serrated plastic flow in nickel, both of which they attributed to the presence of carbon. Serrated plastic flow has been rationalized by a process involving a series of dislocation pinning and multiplication steps.8, This process is more probable when screw dislocations are strongly pinned. Screw dislocations cannot be pinned by pure hydrostatic forces from the symmetrical strains of an interstitial impurity in an fcc lattice, except for small, second-order effects. However, they might be pinned by localized changes in modulus of elasticity around solute atoms,' by the pinning of the edge components of the partial dislocations of an extended screw dislo~ation,'~ or by clustered groups of solute atoms whose net elastic stress field is unsymmetric. The purpose of the present work was to investigate various mechanical properties of the Ni-Co-C a1loy system which are sensitive to pinning effects in order to delineate the specific pinning mechanism of carbon. Carbon concentration, grain size, temperature, and strain rate were varied so that thermal-activation analysis and the Hall-Petch analysis could be used to evaluate the pinning mechanism. Cobalt was added to lower stacking fault energy so that the number and extension of split, screw dislocations would be increased in order to test the possibility of pinning by carbon at extended screw dislocations. EXPERIMENTAL PROCEDURE Nickel and cobalt (both 99.98 pct-. pure) were melted with graphite in stabilized zirconia crucibles and cast at lo-' Torr to form Ni-C and Ni-60 pctCo-C alloys. Two ingots were heated to 1250°C and were forged to 1-in.-sq bars. These bars were machined to 4-in.-round bars, and then swaged cold to 0.144-in. -diam rods. Reductions in area of approximately 75pct were used with intermediate anneals at 900°C for 1 hr. The carbon content of batches of 0.144-in.-diam rods from each ingot was reduced to two levels by annealing 5-in. lengths in palladium-purified, dry hydrogen at 1100°C for 25 and 100 hr. The remaining material from each ingot was annealed at 10"5 Torr for 1 hr at 1100"~. These treatments gave a total of three carbon levels for both the nickel and the Ni-60 pct Co alloy. The 0.144-in.-diam rods were swaged to 70-mil wire, cut into test specimens, and then re crystallized at lom5 Torr in capsules for 1 hr at temperatures ranging from 760" to 1050" ~. The capsules were broken and the specimens were immediately quenched into water. Average grain size was measured using Hilliard's method of circular intercepts." Annealing twin boundary intercepts were counted in addition to grain boundary intercepts to establish an average grain size. Average grain sizes ranged from 5 to 140 p depending on the cobalt concentration and re-crystallization temperature. Tension tests were made in duplicate at various temperatures at a crosshead speed of 8.34 x 10"4 in. per sec with an Instron Universal Testing Machine. Specimens of 1-in. gage length with soldered ball ends were used at atmospheric and cryogenic temperatures. Pinch grips were used on specimens at elevated tem-

Jan 1, 1969

By R. F. Hehemann, D. J. Cometto, G. L. Houze

The w transformation in the Zr-Nb system was studied using X-ray diffraction, dilatometric, re-sistornetric, hardness, and metallographic techniques. w forms in a diffusionless, completely reaersible manner on quenching and in a diffusion-controlled manner- on aging. The temperature at which w begins to form on quenching was determined as a Junction of composition and was found to decrease with increasing solute content. w formed bv aging establishes a metastable equilibrium with an enriched ß phase. The ß/w + ß transus has been determined for this metastable equilibrium and employed to rationalize retrogression and reversion phenomena observed in these alloys. The decomposition mechanism is discussed in terms of a gradual or continuous transformation from ß to the w state. BETA- stabilizing alloying elements lower the MS temperature of the martensitic bcc (ß) to hcp (a') transformation in zirconium and titanium alloys. In certain titanium alloys, a lower symmetry modification of the martensitic structure, termed (a"), also has been reported.1,2 These martensitic transformations are suppressed when the amount of the ß-stabilizing element exceeds a critical level. However, the high-temperature ß phase cannot be quenched to room temperature without change. At alloy contents just above the critical level, |ß trans-forms to the w phase when cooled rapidly.3-6 The amount of w in quenched alloys is reduced by increasing alloy content, and this phase virtually disappears above a critical composition.5 In addition to the transformation during cooling, w also can be formed by aging ß at temperatures below approximately 500°C, and significant alloy enrichment of retained 13 accompanies the isothermal transformation.1-8 The structure of LC is closely related to that of the ß from which it forms.9-15 The bcc (ß) lattice can be generated using a hexagonal cell with three atoms located at (000) and ±(2/3, 1/3, 1/3). This cell has an axial ratio of 0.612 and is oriented with respect to the cubic cell such that (0001)H (111)C and [1120]H [101]C. Consequently, there are four possible orientations of the hexagonal cell, associated with the four (111) planes of the bee lattice. Formally w can be obtained from 0 by allowing the two atoms at the ±(2/3, 1/3, 1/3) positions to approach the coordinates ±(2/3, 1/3, 1/2) and there are four equally probable orientations of w for each 0 grain. In titanium alloys w retains the cubic axial ratio of 0.612" and hence also can be indexed as triply cubic, but this is not the case for aged Zr-Nb alloys where w is clearly hexagonal with an axial ratio of 0.622." The lattice parameters and atomic positions of w depend on thermal history and alloy content. The atomic positions (000), ±(2/3, 1/3, 1/2) provide reasonable agreement between calculated and observed diffracted intensities for w in the fully aged condition.10,14 In the quenched condition, on the other hand, the atoms appear to be displaced from the central plane12,13,15 and assume positions ±(2/3, 1/3, Z = 0.42-0.48) rather than ±(2/3, 1/3, 1/2). This results in a structure with trigonal rather than hexagonal symmetry. The readily detectable parameter and axial-ratio changes of w in Zr-Nb alloys make this system especially attractive for studying the structural changes that occur in the formation and aging of w. Particular attention in this investigation has been directed to the relationship between the structure of w in the quenched and aged conditions, and to certain aspects of the reaction kinetics. MATERIALS AND PROCEDURE Zr-5, 12, 17.5, and 25 pct Nb alloys were prepared by a double arc-melting practice, encased in stainless-steel cans and hot-rolled to 1/8-in.-thick sheet.* Charged weights have been employed for the niobium contents and interstitial analyses are reported in Table I. Dilatometric, X-ray diffraction (filtered CuK, and MoK, radiation), metallographic, and hardness techniques have been employed to follow the transformations during isothermal and quench-aging heat treatments. In quenched specimens, the electrical resistivity also was studied. Betatizing was conducted for 1 hr at 900°C. With the exception of the isothermal dilatometric studies, samples were

Jan 1, 1965