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By G. L. F. Powell, L. M. Hogan

Samples of silver and Ag-O alloy, 0.12 wl pet, have been undercooled to a maximum of 250°C by melting in a slag of commercial soda-lime glass. Grain refinement occurred in undercooled silver samples whet the degree of undercooling exceeded a critical value which varied from 153° to 175°C, depending on the melting conditions. In under cooled Ag-0 samples, the grain structure was fine and equiaxed at degrees of undercooling larger than 50oC. The grain refinement in both silver and Ag-0 alloy samples was the result of recrystallization during or immediately following freezing. When the oxygen content was increased, the degree of undercooling at which recrys-tallization occurred decreased. INVESTIGATIONS of the grain structure of strongly undercooled melts of nickel and cobalt1-4 have revealed a change in room-temperature grain size as a function of undercooling. At large degrees of undercooling the grain size was fine, but samples undercooled by smaller amounts exhibited a much coarser grain structure. The reported undercooling at which this grain size transition occurred in nickel varied between 140° and 175'C undercooling. Walker1 attributed this grain size change to an enhanced nuclea-tion rate. Cavitation at a melt-solid interface was con sidered to generate a high-pressure wave. This, in turn, produced nucleation in the liquid as a result of an elevated freezing point increasing the undercooling in the melt. A similar grain size transition was observed in undercooled silver between 133° and 153°C undercooling.5 The grain size transition in undercooled silver was found to be dependent on oxygen con tent. More recently, the authors5 showed that a simila transition occurred in an undercooled Cu-O alloy, =0.08 pet, but not in oxygen-free copper undercooled more than 200°C. The transition was attributed to recrystallization of the solidified grains after freezing; the greater the degree of undercooling, the more complete the recrystallization. In view of this effect of oxygen content on grain size transition, additional undercooling experiments were carried out with silver. This paper reports the results of these studies on silver and Ag-O alloys undercooled more than 50°C EXPERIMENTAL 1) Materials. The silver was supplied by Matthey Garrett Pty., Ltd., as fine granulate, silver mini' mum 99.9 pet, and contained the trace impurities shown in Table I. The silver originated from two different refineries resulting in the difference in impurity content shown in Table I. However, the variation in composition had no discernible effect on the undercooling behavior. Samples of the silver were melted in an open-ended, vertical, cylindrical resistance furnace, wound with Kantha] wire, the temperature of which was adjusted by means of a variable transformer, 2) Melting and Undercooling. Bulk samples of silver were undercooled by preoxidation and removal of oxidized impurities in a glass slag, as previously described,5 but special precautions were necessary to allow comparison between oxygen-free silver and a Ag-O alloy. Samples free of oxygen were prepared by first melting weights of granulate varying from 200 to 400 g in clean vitreous silica crucibles in air for approximately 15 min. A sample was frozen and a complete glass cover formed over the top surface by adding granulated glass rod. The sample was then melted and frozen several times until no oxygen bubbles appeared in the glass slag. Subsequent thermal analysis indicated that the oxygen content was negligible. The glass acted as an efficient barrier to absorption of oxygen from the air. Ag-O alloys were prepared by melting silver in silica crucibles with only a partial cover of glass, so that the top surface of the melt was in contact with air. The silver was held molten in air for several hours at 1000° to 1050°C, in the expectation that the oxygen in solution would reach equilibrium with atmospheric oxygen. At T - 1000°C the equilibrium oxygen content was calculated as 0.12 wt pet O, using Sieverts law and Eq, [l], derived by Mizikar et a!.7 from the data of Sieverts and Hagenacker8 for 1 atm of oxygen:

Jan 1, 1970

By V. Koump, T. F. Perzak, R. G. Olsson

The rate of dissolution of recrystallized alumina in molten iron oxide in equilibrium with iron at 1450°C was studied by rotating alumina disks in the melt. The samples were rotated from 1 to 12 min at speeds from 54 to 3270 rpm. It is concluded from the experimental results that the process is limited by diffusion in the liquid boundary layer. The interdiffusion coefficient for the dissolution of alumina disks in pure molten iron oxide was estimated to be about 3 xl0-5= sq cm per sec. THE erosion of refractory materials is of considerable importance in iron- and steelmaking processes. Since many refractories contain alumina, this investigation was undertaken to develop a further understanding of the erosion mechanism. Shurygin et al.1-3 have investigated the rate of dissolution of alumina in molten silicates and molten fluorides by rotating alumina disks in various melts. The rates were reported to be limited by diffusion in the liquid boundary layer. In the present investigation the rate of dissolution of alumina in molten iron oxide in equilibrium with iron was studied in a similar manner; alumina disks were rotated at different speeds in molten iron oxide at 1450°C. The relationship between the rate of dissolution of the flat surface of the disk and the speed of rotation was used to determine whether the process is limited by diffusion in the liquid boundary layer or by some chemical reaction mechanism. APPARATUS AND EXPERIMENTAL PROCEDURE A schematic diagram of the experimental apparatus is shown in Fig. 1. The iron oxide melts were contained in 2-in.-diam high-purity iron crucibles that were centered in a cylindrical graphite susceptor. Prior to each sequence of experiments, sufficient reagent-grade ferric oxide was added to a crucible to form approximately 350 g of molten iron oxide; the depth of the melt was about 11/2 in. To prevent excessive erosion of the crucible, about 70 pct of the required iron was added to the crucible in the form of iron powder. The mixture was melted in an atmosphere of purified argon by induction heating with a 250-kc generator. After the oxide was molten, a Pt/Pt-10 pct Rh thermocouple at the exterior of the crucible was calibrated against an iron-sheathed thermocouple immersed in the melt and was thereafter used to regulate the melt temperature. All experiments were conducted at 1450°C with a probable error of 10°C. The alumina disks were constructed from pure re-crystallized alumina having a density of 3.7 g per cu cm (porosity approximately 6 pct). For each disk sample a 1-cm-diam alumina cylinder of known length (0.9 to 1.0 cm) was cemented into a ground seat at the end of a 1.27-cm-OD centerless ground alumina tube. The ends of the tube and cylinder were flush and served as the disk surface. When mounted in the apparatus the alumina tube was guided by tungsten carbide bearings and rotated by a variable-speed motor. The bearing and drive assembly could be raised and lowered in order to allow the sample to be quickly moved in and out of the melt. At the start of an experiment, a sample was preheated for several minutes directly above the melt, and then immersed in the middle of the melt to a depth of a in. and rotated for 1 to 12 min at speeds of 64 to 3270 rpm. The speed of rotation was measured by a hand tachometer. At the start and completion of each experiment, a melt sample was withdrawn with a cold copper rod and the alumina content was determined by chemical analysis. During individual experiments the alumina concentration of the melt changed by the order of 0.25 wt pct. Average melt compositions for individual experiments were used in the subsequent computations and were in the range between 0.12 and 4.3 wt pct alumina. The extent of dissolution was determined by sectioning the end of the sample and measuring the final length of the alumina cylinder with a microscope. Since the exposed end of the cylinder, i.e., the disk surface, was etched in a manner characteristic of this geometry, a number of measurements were made on

Jan 1, 1969

By F. Weinberg

Measurements have been made of solute segregation during dendrilic growth by using radioactive solute elements and ,measuring the activity of den(12-ites cut from decanted specimens. This has been done for both lead awl tin based binary alloys contaitzing the following solute additions: Ag, T1204, was dependet on ko, the equilibrium distribution coefficient in the following way Fay k &apos;c 0.1, C/C 0.6; for k0 >0.1. 0.6 <c,/c,< I. Qualitative obse?-vations were madc of dendritic segregation, by using autoradiographic techniques, for the Sn + Ag110 and Sn + Tlo4 systems. The observation were found to he in general agreement with the measurements ofCA/Co. Autoradiographic were also obtained of scctiolccl delzr11-iie stalks. These indicated that the stalks had a substructure, dclileated by solute corzetlt?atio?zs nlolg the substructure walls. A new dendrite growth direction <JI2> is reported for tila. SOLUTE segregation in dilute binary alloys has been investigated by Pfann,&apos; Smith, Tiller, and Rut-ter,&apos; and others. They considered the case of a slowly advancing plane solid interface, and derived expressions for the distribution of solute in both solid and liquid during solidification. To determine these expressions, they assumed no diffusion in the solid and either complete mixing in the liquid:&apos; or diffusion controlled solute movements in the liquid without any convective mixing.&apos; The present investigation considers solute segregation during dendritic growth, in which case the solid-liquid interface is not plane, and the growth rates are rapid. Segregation under these growth conditions has not been treated mathematically, because of the relative complexity of the system. It has been suggested by Chalmer, on the basis of preliminary results, that an alternative to the diffusion and heat flow controlled conditions during growth is &apos;diffusionless" dendritic growth in which solid is formed with the same composition as the liquid. He suggests this type of growth may depend upon a solvent-solute relationship that permits some solid solubility without excessive increase in internal energy, as is the case for solutions of tin in lead. On the other hand, Montariol,4 and others, have shown experimentally that some segregation does occur during dendritic growth in metals using etching and radioactive tracer techniques to indicate the concentrations of the solute. The present investigation was undertaken to determine, both qualitatively and quantitatively, the extent of solute segregation associated with dendritic growth in a series of binary alloys, as a function of solute concentration. PROCEDURE The solvent materials used were Vulcan Electrolytic tin (99.997 pct purity) and Tadanac lead (99.998 pct purity). The solute materials were Zn, Sn, and Sb (better than 99.998 pct purity), Ag and Co (99.5 pct purity), and T1 (Fisher "purified" metal sticks). Activation of the solute metals was carried out in the reactor at Chalk River, Canada. Master alloys were prepared by induction heating from the radioactive solute metal and the pure solvent, under argon, in graphite crucibles. Pieces of these alloys were then added to the solvent to give the required solute concentration. Dendrites were grown in essentially the same manner as that described by Weinberg and Chalmer, , in which controlled orientation single crystals were grown dendritically in horizontal graphite boats, and the liquid decanted. The crystals were grown and decanted in an atmosphere of tank argon. Before decanting, a sample of the liquid was drawn up in a glass tube and allowed to solidify rapidly. The orientations of the single crystals were such that <loo> was parallel to the growth direction, and (100) in the horizontal plane for lead, and [1101 and (110) respectively for tin. With these orientations long dendrite stalks formed along the bottom of the boat in the dendrite direction (<100> for lead and [I101 for tin) from which secondary branches grew. Only these secondary branches, which grew freely in the liquid from the dendrite stalk to the liquid surface, were used in the measurements. Accordingly, effects due to substrates and oxides on the surface of the liquid need not be considered. In order to measure the solute concentration C, of the dendrites, individual dendrite stalks were cut from the decanted specimens, remelted, and formed

Jan 1, 1962

By E. J. Dulis, R. M. Fisher, K. G. Carroll

IT is well known that ferritic steels containing more than 15 pct Cr when subjected to temperatures in the range of 700" to 1000°F exhibit increasing hardness and decreasing ductility. The phenomenon has been widely termed the "885 °F embrittlement," after the temperature of most marked effect.&apos;.&apos; In view of the excellent review articles available in the literature3-6 only a brief account of experimentally established facts need be given here. The extent of changes in physical characteristics during embrittlement depends on chromium concentration and time at temperature, higher alloy content and longer time both promoting more rapid and extensive changes. In a 27 pct Cr steel, changes in impact strength and in angle of fracture in bending can be detected after only a 1 hr exposure at 885°F; after 50 hr this steel becomes quite brittle. Hardness increases slowly with time during thousands of hours exposure and may attain a maximum hardness number twice as large as that of the unexposed steel. Microstructural changes accompanying embrittlement have been described as an initial widening of grain boundaries followed by eventual darkening of ferrite grains. Embrittled steels etch more readily, e.g., the weight loss of a 27 pct Cr steel in acid solution may occur at a rate one hundred fold greater following exposure at 885 OF. Marked changes which accompany embrittlement have been observed in electrical resistivity, specific gravity, and magnetic coercive force. Changes in physical properties may be readily removed by heating at temperatures above the embrittling range, such as a treatment at 1100°F for 1 hr. It has frequently been noted that the 885 °F embrittlement suggests precipitation on a submicro-scopic scale of a chromium-rich constituent, the nature of which has not been revealed by X-ray diffraction. Progressive broadening of the body-centered cubic diffraction lines during embrittlement has been observed," and recent observations by Lena and Hawkes&apos; upon single crystals have shown early asterism in X-ray photographs, disappearing within an hour at 900 °F. Many workers have ascribed8-13 the phenomenon to a precipitation of a phase (FeCr), which is known to cause embrittling effects at temperatures much higher than 885°F. Two general observations, however. suggest that a precipitation cannot be responsible for the 885°F phenomenon: 1—prior cold work greatly enhances a formation, whereas it scarcely affects the 885 °F embrittlement, and 2—the presence of an alloying element such as nickel or manganese may have an effect on the 885 °F embrittlement which is opposed to its effect upon a formation. The slight enhancement of a formation and 885°F embrittlement observed in the presence of elements with strong carbide and nitride forming tendencies&apos; is probably a consequence of lessened chromium depletion of the matrix. The bar graph in Fig. 1 shows a typical example, taken from two 27 pct Cr steels used in this work, of the hardness after exposure for 10,000 hr at 900°, 1050°, and 1200°F. Steel A (0.03 pct C, 3.13 pct Mn) showed marked hardening at 900" and 1200°F, whereas steel B (0.12 pct C, 0.63 pct Mn) exhibited only the 900°F hardening. The cr phase was found in steel A at the higher temperatures but not in steel B. Presumably a formation is enhanced by the low-carbon and high-manganese concentrations in A," Thus there are two distinctly different hardening phenomena present which cannot both be ascribed to v precipitation without invoking a transition phase possessing remarkable properties. Materials A number of chromium steels exposed for long periods (5000 to 34,000 hr) at 900°F, as well as unexposed samples of one of the steels, were available for this investigation. Table I gives the chemical compositions and aging treatments of these steels. In addition to these steels exposed in the elevated temperature test furnaces of the National Tube Division, a number of high-chromium steels were heated for short periods in small laboratory air furnaces and lead baths. Supplementing these commercial steels, a sample of high-purity (0.018 pct C, 0.002 pct N) 28 pct Cr iron, exposed 1000 hr at 887°F. was furnished by the Union Carbide and Carbon Corp. In addition, an alloy of iron and chromium of high purity containing 46 pct Cr was used. This

Jan 1, 1954

By H. Pew, J. H. Field

To alleviate pollution more restrictive legislation is being enacted, either limiting emission of pollutants or the type of fuel that can be utilized. The nature and magnitude of air pollution problems affecting the mining, preparation, coking and combustion of coal are described. Methods for combatting particulate emissions by use of mechanical separators and electrostatic pre-cipitors are discussed. Proposed methods to meet the problem of gaseous emissions currently receiving considerable attention are described, with special emphasis on methods to decrease pollution by sulfur oxides. Concern about air pollution goes back several centuries, but until very recently most effort has been aimed at coal smoke and other visible pollutants. The classic example of a &apos;successful&apos; campaign for smoke abatement and control is the fruitful combined effort of the city of Pittsburgh and its surrounding Allegheny County, which eventually led to the reconstruction of downtown Pittsburgh at an estimated cost of one billion dollars. Historically, the city&apos;s downtown Golden Triangle district had been afflicted by pollutants evolving from steel mills, from a variety of other industries, and from railroad locomotives. Efforts to alleviate the situation prior to 1943 were virtually ineffective. In 1945, however, a comprehensive redevelopment plan was prepared and backed by state authority. Within a few years a clean, modern metropolis has evolved where once stood America&apos;s famous &apos;smoky city.&apos; But the victory in Pittsburgh, as in various other American cities, has not solved the national problem. Current estimates indicate that 133 million tons annually of air pollutants from all sources still are emitted annually into the atmosphere above the United States. About 10% of this annual effluent is particulate matter so that most of the remaining pollution problems will be solved only when other effluents are reduced. Essentially, these are sulfur oxides, nitro- gen oxides, hydrocarbons, and carbon monoxide. Over the years, both states and local communities have tended to increase the restrictions on smoke and fly ash — problems mostly of concern in the combustion of coal. Prior to the middle 1950&apos;s, ordinances sometimes permitted emissions of smoke equivalent to as much as No. 3 on the Ringlemann scale. Since 1956, no ordinance has been passed which allows smoke of greater than No. 2. Under today&apos;s conditions of improved fuels, equipment and practice, a few communities have passed laws prohibiting emission of smoke of any density darker than Ringlemann No. 1. The majority of existing laws on fly ash emission in the U.S. limit emissions equivalent to 0.85 lb of fly ash per 1000 Ib of flue gas. In recent years, however, regulations which have been adopted give cognizance to the higher level of performance now obtainable with improved equipment. A comparison of the restrictions of five codes adopted since 1960 is given in Table I. The most stringent of these is the one for New York City which provides for a maximum emission of 0.6 lb fly ash per million Btu heat release (equivalent to roughly 0.51 lb/1000 lb of flue gas). The first comprehensive effort to restrict the emission of SO2 resulted from the passage of a 1937 law in St. Louis. This regulation stipulated that coal containing in excess of 23% volatile matter and 2% sulfur must be washed, thereby presumably producing some effective reduction in the input sulfur content. This was followed in 1949 by a Los Angeles County law which prohibited the emission of SO2 in concentrations greater than 0.2%. Most SO2-restrictive legislation passed since that date has been based on this limiting 0.2% SO2 by volume, although modifications are occasionally permitted under selected conditions, sometimes based on the fact that certain limiting ground concentrations are not exceeded — such as in the rules adopted by the San Francisco Bay area. To date, no legislation has been passed in the U.S. to limit the generation of nitrogen oxides from the combustion of fossil fuels. However, such oxides are considered to be of potential importance in air pollution control because of their possible detrimental effects on health and their reported role in the formation of photochemical smog. Interest in reducing oxides of nitrogen from powerplant and auto exhausts is increasing and regulations limiting their quantity can be expected in the future.

Jan 1, 1968

By M. J. Marcinkowski, E. N. Hopkins

An investigation has been made of the ß(fcc) — a(hexagonal) transformation which occurs in lanthanum using both electrical resistivity and transmission electron microscopy techniques. It has been shown that the ß phase can be retained below the transformation temperature by rapid quenching but that the sample immediately begins to transform to the a phase. The transformation is observed to nucleate in the vicinity of inclusions. Based on the above observations, a detailed model of the transformation has been advanced which involves the nucleation of an extrinsic stacking fault bounded by a pair of Shockley partial dislocations in the vicinity of some heterogeneity, i.e., an inclusion. The stress field of the resultant dislocation pair acts to nucleate extrinsic faults in adjacent planes and leads quite naturally to the B-a conversion with a minimum of strain energy induced in the crystal. LANTHANUM possesses an fcc structure ß) (8) upon cooling transforms to a hexagonal modification (a) in much the same way as cobalt. The one exception, however, is that the stacking sequence of the closest packed planes in a La is ABAC ABAC, and so forth,1 whereas in cobalt it is ABAB, and so forth, i.e., hep. Mainly on the basis of transmission electron microscopy techniques, there seems to be little doubt that the 0 — a transformation in cobalt involves a dislocation mechanism2 although its exact nature still remains obscure. Although the ß - a transformation in lanthanum is somewhat more complex than that occurring in cobalt, it was thought that its very uniqueness would be helpful in understanding fcc — hexagonal transformations in general. Such a general understanding of these transformations is important since they represent what are perhaps the simplest of the martensitic class of transformations. The experimental techniques used were those of electrical resistivity and transmission electron microscopy. EXPERIMENTAL PROCEDURE The lanthanum used in this investigation was prepared by the calcium reduction of lanthanum fluoride in a tantalum crucible under an argon atmosphere as described by Spedding et al. The residual calcium was removed by vacuum remelting in a tantalum container. Portions of the metal were then analyzed by emission spectroscopy as well as vacuum fusion. The amounts of the various impurities that were found are listed in Table I. A portion of the lanthanum ingot was swaged at room temperature into 0.030-in.-diam rod for the resistivity samples while the remainder was rolled into 0.010-in.-thick sheet for the transmission electron microscopy phase of this investigation. In order to eliminate the plastic deformation induced in the samples during fabrication, they were sealed in evacuated tantalum lined quartz capsules and annealed for 1 hr at 700° C. Specimen resistances were measured using the conventional "four-wire" technique described by MacDonald, 4 employing a type K-3 Universal Potentiometer and a standard 0.001-ohm resistor, both manufactured by Leeds and Northrup Co. By taking into account all of the possible errors in the apparatus, it was felt that the absolute resistivities of the approximately 0.87-in.-long specimens measured are reliable to 3 pct. Resistances from room temperature to 700°C were measured using a vacuum furnace. In order to avoid sample contamination by the thermocouple, chromel-alumel leads were spot-welded into a tantalum shield which in turn was welded to the lanthanum specimen. Resistances below room temperature were obtained by transferring the sample to a helium gas-filled quench tube and slowly dripping liquid nitrogen into a surrounding dewar flask. A steady reduction of temperature to about —190°C was completed in about 23 hr, and the resistances were measured at various temperature intervals. As will be shown shortly, rapid quenching from above about 350°C was sufficient to suppress the B -a transformation initially but subsequent annealing at lower temperatures leads to a partial ß --a conversion. To investigate this aspect of the transformation, the samples were suspended in the hot zone of a helium-filled furnace from which they could be rapidly dropped into the quench tube mentioned previously which was now

Jan 1, 1969

By G. R. St. Pierre, S. W. Gilby

The equilibrium vapors established over Fe-Cr-Ni alloys containing up to 50 pct Cr at 1600°C were determined by collecting the effusate from thoria Knud-sen cells. From these data, the activities of iron and nickel in the ternary alloys have been calculated. In a separate set of measurements, a mass spectrometer was used to determine the activities of the three components in a series of Fe-Cr-Ni alloys. In binary Ni-Cr melts, the activity coefficient of chromium referred to pure solid chromium as standard state is less than one at low chromium contents but is greater than one for melts containing more than 30 at. pct Cr. For Ni-Cr melts containing less than 10 at. pct Cr, In ?Cr = 0.15-0.92 N2. Ternary activity diagrams for Fe-Cr-Ni alloys are presented. In iron-rich alloys, the interactions between the three components are negligible; however, in nickel-rich alloys, the three principal interactions have significant positive values. THE activities of components in binary Fe-Ni alloys have been measured both in the solid solution1-3 and liquid solution4-6 temperature ranges. The determination for Fe-Ni melts have been made by three different methods: transpiration,4 Knudsen-cell effusion with analysis of an accumulated condensate,&apos; and Knudsen-cell effusion with direct measurements of relative component intensities by mass spectrometric methods.6 The results obtained by the three different methods in three different laboratories are in excellent agreement and considerable confidence is placed in the activity coefficients for Fe-Ni melts. At 1600°C, the activity coefficients in dilute solution are ?°Ni = 0.66 and ?oFe = 0.40 for the pure liquid components as standard states. Measurements made on solid solution Fe-Cr alloys7-10 indicate that the component activities are slightly greater than ideal. The measured activities in iron-rich Fe-Cr liquid alloy11-16 indicate almost ideal behavior for this system. The recommended activity coefficient of chromium in dilute solution in liquid iron at 1600o C, ?oCr, is 1.18 when referred to pure solid chromium as standard state and 1.04 when referred to pure undercooled liquid chromium as standard state." The conversion factor of 1.14 is based on a melting point of 1857°C and an enthalpy of melting of 4047 cal per g at. for chromium.21,37 The activity coefficient of iron in dilute solution in liquid chromium has not been measured. Activity measurements on solid Ni-Cr alloys have been made by several investigators.17-20 During the review of this manuscript, Fruehan12 published the results of a study on liquid Ni-Cr alloys at 1600°C. He used galvanic cells22 with ZrO2(CaO) as the electrolyte and Cr/Cr2O3 for the reference electrode. His results indicate that ?oCr in nickel at 1600°C is about 1.01 when referred to pure solid chromium and about 0.89 when referred to pure undercooled liquid chromium as standard state. The activity coefficient of nickel in liquid chromium is unknown. The compositions of vapors freely evaporating from some molten Fe-Cr-Ni alloys have been measured by Lyubimov, Granovskaya, and Berenstein.23 Their results, combined with known vapor pressures, indicate strong positive deviation for chromium in iron-rich alloys and strong negative deviation for chromium in nickel-rich alloys. EXPERIMENTAL PROCEDURES This research was undertaken to determine the activities of components in ternary Fe-Cr-Ni melts containing up to 50 at. pct Cr. At the initiation of the research, we decided to use a Knudsen-cell method wherein a portion of the effusing vapor is condensed on a cold plate.&apos; The composition of the equilibrium vapor for each alloy composition is determined in this method by holding the system under constant conditions for a period of time sufficient to collect enough condensate on the cold plate for chemical analysis. The activities of components in ternary alloys can be calculated from such vapor composition determinations by Eq. [1].24

Jan 1, 1970

By John J. McKetta. Jr., Louis B. Lesem, George H. Fancher

The proration of oil produced in the field frequently is based partially or entirely upon the gas-oil ratio of wells. The measurement of the gas-oil ratio is one of the more important field tests in regulatory and pro-ration work, and the test always should be conducted according to standardized methods and procedure. Obviously, the gas-oil ratio and the volume of gas produced by a well depend upon many factors but should be independent of the method of measurement and of the devices used to measure gas and oil. Consequently, the volume of gas accompanying a barrel of oil produced by a well may be measured by any reliable and accurate device or instrument. Frequently either a critical flow prover or an orifice well tester is used for this purpose, and for a particular well the same rate of flow of gas should be obtained regardless of whether a critical flow prover or an orifice well tester is employed in the test. In Texas, when using either instrument, either Capacity Table 1 or 5&apos; is employed in making the necessary computations. If the tables are used, a discrepancy always is found whenever the two instruments are compared by extrapolation to the same conditions of flow. Clearly, Tables 1 and 5 must be at fault in some respects. The orifice well tester and the Bureau of Mines type of critical flow prover are essentially the same instrument; both devices utilize a square-edged orifice 1/8 in. in di-ameter as the primary element, and both freely discharge gas to the atmosphere. Tables for the orifice well tester&apos; have been published in the ranges of 0 to 15 in, of water and 0 to 40 in. of mercury (Hg) differential in pressure. Coefficients for the critical flow prover have been published for differentials in pressure greater than 75 psia. An extrapolation of either differs from the other set of data as much as 18 per cent at some points. No immediately obvious reasons for the discrepancy was found, and data available from the literature were insufficient to effect reconciliation. Consequently, a series of experiments was performed to check the available data and to determine discharge coefficients for the two devices in an overlapping range of differential pressure. The correlating equation used in preparation of tables such as Tables 1 and 5&apos; is the so-called hydraulic equation, Q = C1 vh .....(1) The tables cover orifice sizes from 1/8 through 1 1/4 in. A second set of tables, for use with greater differentials in pressure, apply only to the 3/4, 1, and 1 1/4-in. orifice plates over a pressure range of 0 to 40 in. of Hg. The correlating equation used in preparation of the tables is Q = C2 vH(29.32 + 0.3H)/G . (2) Each of these equations is valid only for the range to which it has been applied, and neither equation is valid for extrapolation. Theoretical equations for flow through an orifice are based upon assumptions of fractionless flow and an emergent jet the size of the orifice. A multiplier, C,, called the discharge coefficient, is inserted into the theoretical expression to compensate for both frictional losses and the contraction of the jet experienced in the plane of the orifice. Buckingham3 shows that C, should vary with the ratio of upstream to downstream pressure for flow of a compressible fluid at any average linear velocity through the orifice less than the velocity of sound. Work published by the National Advisory Committee for Aeronautics4 indicates that the variation continues into the so-called critical flow region until the vena contracta coincides with the plane of the orifice. The NACA work, however, does not indicate a leveling off of the coefficient. The work of the Bureau of Mines5 for the critical flow prover was based on differential pressures greater than 75 psi and indicates that the discharge coefficient, Cd, is constant in this region at a value of about 0.86.

Jan 1, 1958

By J. L. Christie

A recent case of failure of a-low karat gold alloy by stress-corrosion cracking is of interest because it illustrates a principle frequently overlooked: the relation between high residual stress and high yield point. The alloy contained Gold, per cent 39.6 Copper, per cent 51.1 Silver, per cent 5.2 Nickel. per cent 3. 1 Zinc, per cent 1.0 Tension tests on the alloy showed the following:

Jan 1, 1945

By E. T. Turkdogan, W. L. Worrell

The rate of reaction of hydrogen sulfide with ferrous sulfide was studied by measuring the initial rates of sulfidation of iron strips in hydrogen sulfide-hydrogen-argon mixtures at 670°, 800°, and 900" C. The time-dependent surface sulfur activity is derived from the instantaneous rate of sulfidation with the assumption that diffusion in the sulfide layer is in a pseudo-steady state with the gas-sulfide chemical reaction. The rate of sulfur transfer from hydrogen sulfide to the surface of iron sulfide is proportional to the partial pressure of hydrogen sulfide and inversely proportional to the activity of sulfur at the surface of the sulfide layer. The derived rate equation is based on the assumption that most of the surface sites on the chemisorbed layer of iron sulfide are occupied by sulfur atoms and that the slow rate-controlling reaction is the dissociation of hydrogen sulfide on the chemisorbed layer. The experimental results are in reasonable accord with this reaction model. 1 HE slow approach to parabolic growth rate in the sulfidation of iron in hydrogen sulfide-hydrogen mixtures is a manifestation of slow approach to surface equilibrium between the gas and the surface of the sulfide layer. For example, at 800°C and Fig. 7 in Part I, the parabolic growth rate begins after approximately 1 day of sulfidation time. With decreasing temperature and decreasing partial pressure of hydrogen sulfide, the time necessary to reach gas-sulfide surface equilibrium is much longer. These observations are similar to those reported previously by Turkdogan et al.&apos; on the oxidation of iron to wustite in hydrogen-water vapor mixtures. The slow approach to gas-sulfide equilibrium is well-demonstrated by the results in Fig. 1 where the square of the weight gain per unit area, ( g S per sq cmI2, is plotted against time for gas mixtures having PH g//>H = 1.0 with 0, 25, 50, and 67 pct Ar at 800°C. The value of 1.5 X lo-&apos; (g S per sq cm)&apos; corresponds to almost complete sulfidation of the iron strip (- 0.05 cm thick) to iron sulfide. The points fall on S-shaped curves. If the inflection parts of the curves are considered to be linear, thus indicating parabolic growth rate and the establishment of gas-sulfide equilibrium, their slopes would have to be the same for a fixed ratio in the gas mixture. Such is not the case, and the slopes for the "linear" parts of the S-shaped curves are lower than the value when gas-sulfide surface equilibrium is established. The subject matter of this paper is the kinetics of the surface reaction of hydrogen sulfide with iron sulfide during the early stages of sulfidation of an iron strip. EXPERIMENTAL RESULTS The apparatus and materials used were the same as those discussed in Part I. In the present experiments of relatively short duration, a purified iron strip, 5 by 2 by 0.05 cm, which was suspended in the uniform hot zone of a vertical zircon tube, reacted with a flowing gas mixture of hydrogen sulfide-hydrogen-argon. The sample was suspended from a gold chain attached to a calibrated silica spring. The amount of sulfur picked up by the iron, which formed a layer of iron sulfide, was determined by measuring the displacement of a reference point on the silica spring, using a cathetom-eter. In all cases, the degreased sample was first annealed in a stream of oxygen-free dry hydrogen for several hours to remove any impurities such as oxygen, nitrogen, or carbon which might have been present on the surface of the sample.&apos; The rate measurements were carried out at 670°, 800°, and 900° C in hydrogen sulfide-hydrogen-argon mixtures with pH2 s/Ph2 ratios from 4 to 0.1. In a few experiments of short duration no sulfur deposition was observed in gas mixtures with (comments concerning sulfur deposition at high pH s/ph ratios were made in Part I.) Typical examples of the results obtained are shown in Fig. 2 for two temperatures and two gas compositions. With the present experimental technique, reproducible rate data could be obtained only after a uniform

Jan 1, 1969

By T. H. Alden

BaCKOFEN, Avery, and Turner have shown"2 that the large tensile elongation in superplastic metals is correlated with a high strain-rate sensitivity of the flow stress. At present, the reported superplastic materials are complex eutectic or eutectoid systems with a large solubility range in at least one of the terminal phases. This note describes the behavior of a solid-solution alloy (at 22°C)5,= Sn-1 wt pct Bi, which has been processed in order to produce a fine grain size.2,3,6 The chill-cast, 1/2-in.-diam ingot was homogenized for 1 week at 130°C, extruded at O°C to 0.083 in. diam, quenched within 15 sec, and stored at -196°C. Samples were warmed to room temperature 16 hr prior to mechanical testing or metallography. During tension testing at room temperature, it was found that the flow stress reached a steady-state value which depended on the imposed elongation rate and was independent of strain to an error of about ±5 pct. In Fig. 1 is plotted log a vs log e where a is the measured steady-state flow stress and i the imposed strain rate. These quantities were calculated using "instantaneous" area and length to an accuracy in the sample diameter of ±5 x 10-4 in. The log a-log 1 curve is sigmoidal. The strain-rate sensitivity, which may be given by d log old log 6,&apos; has a maximum value of about 0.48 between i = 5 x 10-4 and 5 x 10-3 per min. The value is well above frequently measured values for metals at creep temperatures (0.1 to 0.3) and comparable to the rate sensitivity for Pb-Sn solder: which has been shown to be superplastic.2,3 Note, however, that below or above the strain rate of 10-3 per min the rate sensitivity of Sn-1 pct Bi decreases and at high rates is more nearly like that of conventional metals. The grain size of this material is relatively uniform, Fig. 2, and may be specified by an average linear intercept i.&apos; Grain growth occurs at room temperature but after 16 hr, L = 5. Tensile elongation was measured on finer-grain-size specimens, which had been at room temperature 15 min prior to testing. After 15 min, the samples are fully recrystallized. Three strain rates, 5 x 10-3, 10-2, and 2.5 x 10-2 per min, were employed for which the measured elongations varied from 425 to 500 pct. These results would seem to justify the term superplastic, applied to this alloy, although the ductility is inferior to that of the Sn-Bi eutectic.3 The Sn-1 pct Bi alloy, which is a single-phase material with equiaxed, fine grains, is an excellent one on which to make an estimate of the Nabarro-Herring9 creep rate. According to the model of Avery and Backofen2 a substantial part of the deformation must be accomplished by N-H creep in the strain rate-temperature regime in which d log a/d log 1 = 0.5. The maximum N-H creep rate is given by9 L2kT where v is the atomic volume and D the self-diffusion coefficient. For this test v = 3.O x 10-23 cu cm, D = 1.4 e-23,300/RT = 2.5 x 10-l7 sq cm per sec,7 L = 5x 10+4 cm, T = 295°K, and a = 2000 lb per sq in. = 1.38 x 10&apos; dyne per sq cm. The calculated creep rate is E= 4 x 10-11 per sec which may be compared with the measured creep rate at this stress, Fig. 1, of 6.7 x 10-5 per sec. The N-H creep rate is seen to be negligible. Earlier consideration2 of this question was in error because incorrect diffusion coefficients were used. The assistance of R. R. Russell in metallography is acknowledged gratefully.

Jan 1, 1967

By Frank T. Sisco

SHORTAGES of metals and minerals and substitution of less critical materials for those in which a virtual famine exists received detailed and frank discussion at a recent conference in Washington called by Robert E. McConnell, director of the Conservation Section, Office of Production Management. Representatives of the large engineering societies in the United States and O.P.M. department heads participated in the discussion. To open the session Mr. McConnell called upon Zay Jeffries, chairman of the O.P.M. Advisory Committee on Metal Conservation and Substitution, to summarize the present situation. Dr. Jeffries touched briefly on the status of metal shortages, which has already received considerable attention in the popular press, and said that much of the present difficulty was caused by the too-hasty substitution-without sufficient con¬sideration of the immediate conse¬quences-of metals having ordinarily a small yearly consumption, for those having a much larger annual consumption. As a result, all the alloying metals used in alloy-steel manufacture have become scarce, with the exception of molybdenum; further¬more, despite a reserve of 11,000 tons of molybdenum in this country and despite a normal consumption of only about 3500 tons and a prospective annual production rate of 21,000 tons by October, a shortage of molybdenum is possible before this war is over.

Jan 1, 1941

By W. B. Morrison

The principal factors influencing the total percent elongation of a lead-tin eutectic and several low-alloy steels which exhibit superplasticity were investigated. These factors are: a) Strain-rate sensitivity of the stress (m). b) Specimen geometry; in particular, diameter-to-length ratio (d0/l0), and surface irregularity (e.g. notches). A feature of the tensile deformation of a superplastic alloy is the early appearance of a neck which continues to extend throughout the test. As a result the total elongation is directly related to the diameter-to-gage-length ratio. The strain-rate sensitivity of the stress varies during extension, mainly because of microstruc-tural changes. A satisfactory correlation exists between the minimum strain-rate sensitivity and total elongation, as indicated in the equation shown below. Notches cause a considerable reduction in total elongation and their effect is greatest at high values of strain-rate sensitivity. An initial notch depth exists below which there is no appreciable effect on elongation. An equation relating elongation to strain-rate sensitivity and specimen geometry is proposed, THE total percent elongation of a tensile specimen is usually regarded as a measure of superplasticity. However, little attention has been given to the fact that total elongation is also a function of specimen diameter and gage length, even though the importance of geometry has long been recognized in the tensile testing of materials in which elongations are relatively small.* In the present study, the effect on superplas- tic extension of the ratio of specimen diameter to length was investigated and an analysis was made of strain distribution along the gage length. Superplastic alloys are characterized by having a flow stress sensitive to strain rate but relatively insensitive to strain.2"&apos; In an earlier study,5 it was found that the strain-rate sensitivity of several low-alloy steels changed during the test and that the minimum value correlated best with total elongation. In the present study, the variation of strain-rate sensitivity during the straining of a lead-tin alloy was investigated. EXPERIMENTAL PROCEDURE Lead-tin eutectic alloy (32 pct Pb 68 pct Sn), made from high-purity lead and tin, was cast in the form of 1-in. diam rod. The rod was cold-rolled in stages to 0.25 in. diam and then swaged to 0.125 in. diam. Samples were cut from the rod at various stages of the mechanical working process. Tensile test specimens were then made by attaching threaded steel ends to the rods by an epoxy-base thermosetting adhesive which required a setting treatment of 2 days at room temperature. The rods were seated in the steel ends in holes approximately 0.01 in. oversize. The test pieces thus made were of various diameters and gage lengths. These were stored at a temperature of about -70°C (-94°F) until required for testing. The distribution of strain along the gage length of a specimen was determined at various stages of the tensile test. The lead-tin alloy was superplastic at room temperature and direct observations during testing were simple. The gage length was divided by reference marks into five initially equal units. These reference marks were made with ink and were reinforced during the test when the initial marks became faint. Photographs were taken, at intervals, of the deforming specimen and measurements of the elongation in each of the gage units were made on these photographs. Tests were also done on several superplastic low-alloy steels whose properties were described earlier.5 The chemical compositions of these steels are given in Table I. A diagram of the notched tensile specimen used in the steel tests is shown in Fig. 1. The specimens were tested at 900°C (1650°F) in a furnace through which a mixture of He + 2 pct H was continuously circulated.5 RESULTS Strain Distribution. Approximately 20 tests on lead-tin alloy specimens were conducted to determine the strain distribution during the test. An example is shown in Fig. 2 of the strain distribution in the gage

Jan 1, 1969

By Howard Martens, Pol Duwez

IN two papers published in 1949, alloys of chromium with the refractory metals tungsten, molybdenum, tantalum, and columbium were investigated in view of their possible use as high temperature resisting materials. For the Cr-Ta system, a partial phase diagram was presented and the only intermediate phase was identified at Ta2Cr3. A phase of the same composition was also observed in the Cb-Cr system. The X-ray diffraction data presented in these papers, however, were insufficient for crystal structure determination. It is shown in the present study that the only intermediate phase in both the Ta-Cr and the Cb-Cr systems corresponds to the ideal stoichiometric ratio TaCr2, or CbCr2. Both structures are cubic, MgCu, type. At high temperature, however, TaCr2 has a hexagonal MgZn, type structure, which can be retained at room temperature by fast cooling. The alloys were prepared by melting in a helium arc furnace on a water-cooled plate. The design of the furnace was essentially the same as that described in ref. 3. Some alloys were also obtained by sintering compacts made of the mixed powders pressed at 80,000 psi. The sintering was carried on for 4 hr at 1375°C. The tantalum and columbium powders were supplied by Fansteel Metallurgical Corp., North Chicago, 111. The tantalum powder was the reagent grade, with a particle size smaller than 400 mesh and a total impurity content less than 0.1 pct. The columbium powder was smaller than 325 mesh and contained approximately 0.1 pct C and traces of Fe, Ti, and Zr. The electrolytic chromium powder from Charles Hardy, Inc., New York, was smaller than 300 mesh and contained about 0.1 pct Na, 0.05 pct Ca, and traces of Cu, Al, Mg, Si, and Co. Powder diffraction patterns were obtained with a 14.32 cm camera, using copper Ka radiation filtered through nickel foil. The powder pattern of the TaCr2 alloy obtained by sintering at 1375&apos;C was different from that obtained on the same alloy rapidly cooled from the melt. Contrary to this result, the powder pattern of CbCr2 was the same, whether the alloy was made by sintering at 1375°C or by melting, and was similar to that of the TaCr, sintered. It was also found that the structure of the TaCr2 specimen obtained by melting was retained after heating for 4 hr at 1590°C, but transformed into the structure found in the sintered specimen after heating for 4 hr at 1375°C. Hence, the structural change of TaCr2, appears to be a reversible polymorphic transformation. CbCr2 and ToCr2 Structure, Low Temperature Form By using large scale Hull-Davey charts, the powder pattern of CbCr, and of the low temperature form of TaCr2 were readily interpreted on the basis of a face-centered cubic lattice with a parameter of approximately 6.95 kX. The indices of the reflections together with the values of sin&apos; 0 are given in Tables I and 11. From this list of observed reflections, it appears that the (200), (600), (024), (046), and (028) reflections are missing. The lack of (h00) reflections for h 4n indicates a four-fold screw axis. The missing (Okl) spectra for k + 1 An indicate the existence of a diamond glide d. The combination of these symmetry elements can be found in the O— Fd3m space group, which is therefore the most probable one. After having determined the approximate density of TaCr, by the immersion method, the number of molecules per unit cell was calculated and found to be nearly eight. This information, added to the fact that the most probable space group is O leads to the consideration of a structure of the MgCu2 type, in which the atoms have the following positions: 8 magnesium in a and 16 copper in d. On the basis of this structure, intensities were computed by means of the usual formula: 1 cos&apos;20 I a sin2 cos where F is the structure factor; 8, the Bragg angle: and p, the multiplicity factor. As shown in Tables I

Jan 1, 1953

By M. B. Kasen

eralized strain equation appear quite different they are really identical. This identity can be shown in a simple mathematical rearrangement. Referring to Eq. [I], the substitution of ln(1 +?E) for ? (where ?E is the engineering strain) yields: a =s8- (s8?so)e?1/?c) ln(1 +s8??E) [4] or s =s8?(s8?so)(l +?E)-1 [5] Replacing (1 + cE) with its equivalent, (l/lo): s=s8?(s8?so)l/€c [6] Adding and subtracting the quantity (s8 - so) yields: s=s8?(s8?so)+ (s8-so) l-(^)l/£c [7] = s0+(s8?so)ri-(^0)1/^l [8] Letting n = 1/?c it follows that: This equation will be seen to be identical to the generalized strain expression in Eq. [3], the coefficients a and b being equal to so and (s8?so)/?c or n(s8?so), respectively. Confirmation of the relationship noted above was provided by fitting experimental stress-strain data5 for 304 stainless steel (total strain from 0.072 to 1.15) at room temperature at a true total strain rate of 4 x 10-3 sec-1 to Eqs. [I] and [3]. In this study nonlinear regression analyses were used to evaluate the equation constants. The results of this evaluation are presented in Table I. When these constants are used in Eqs. [I] and [3] the curves obtained are coincident. It will also be noted that the value of n is equal to YE,, the value of a is equal to so, and the value of b is equal to (s8 - so)/Ec. These equalities are illustrated in Table I where the Voce constants in parentheses were calculated from a, b, and n values. Excellent agreement is seen to exist. It is concluded, therefore, that the Voce equation is, indeed, identical to the linear expression involving generalized strain. In applying the generalized strain concept to stress-strain data, it has been noted3 that data within what was termed the transition region (first few points in the plastic flow region) must be excluded from the analysis in order to obtain the linearity in Eq. [3]. A similar exclusion applies in the use of the Voce equation since an expression which would describe the strain-hardening region would usually be found to be not too effective in the elastic-plastic fillet2 which connects the elastic line to the strain-hardening region. These considerations further substantiate that Eqs. [I] and [3] do indeed describe identical behavior. Note added in proof. Since this manuscript was submitted for publication, correspondence with Professor S. R. Davies, University of Edinburgh, (of Ref. 3) indicated that his associates have also confirmed the identity between the Voce equation and the generalized strain concept. It is sometimes desirable to obtain a small grain size in super-purity metals. Unfortunately, the absence of impurities, the low recrystallization temperatures, and the low activation energies for boundary motion result in rapid growth of relatively few grains when conventional recrystallization practices are used. For example, Arajs et al.1 were unable to obtain a mean grain diameter less than about 0.4 mm in iron containing 38 at. ppm of impurity. The <1 at. ppm purity range in many modern superpurity metals results in yet larger limiting grain sizes. It is the purpose of this communication to describe a technique by which mean grain diameters as small as 30 µ have been produced in aluminum having an impurity content < 0.5 at. ppm. The technique should be applicable to superpurity metals other than aluminum and has potential applications to grain refinement of commercial alloys. The necessary and sufficient requirements for production of a small grain size are as follows: a) a large number of potential nucleating sites must be created randomly within the cold worked material, b) sufficient thermal energy must be provided at a rate that will simultaneously activate a large fraction of the potential nuclei, and c) the grain growth must be stopped shortly after grain impingement.&apos; A high density of potential grain nuclei may be cre-

Jan 1, 1970

By M. J. Duggin

Evidence is produced below to show that the carbide "Mn,C," reported by Kw,and ersson&apos; is really a mixture of two carbides. One is an Fe-Mn carbide which is probably isomorphous with hexagonal Mn15C4, reported by Bouchand and Pruchart.&apos; The othev is a tetragonul Fe-Mn carbide with cell dzmensions close to those reported previously by the author3 for an Fe0.1Mno. carbide. IN a recent investigation several Fe-Mn carbide specimens, each weighing about 5 g, were prepared in a manner previously reported.3 Debye-Scherrer X-ray photographs were taken of at least two samples from different parts of each specimen using a 114.6-mm-diam camera. The exposure time was 2 hr using filtered chromium radiation, a tube voltage of 35 kv and a tube current of 10 ma. In the case of one specimen, which contained nominal weight percentages of iron, manganese, and carbon of 9.40, 84. 6,, and 6.0, respectively, duplicate chemical analyses indicated f1.2 and 5.5 wt pct C to be present in the two samples taken of the specimen: no graphitic carbon was detected. One powder-diffraction pattern of this specimen indicated the presence of the tetragons: carbide reported earlier,3 (a = 6.772A, c = 9.427A), plus cementite (O), plus an unknown structure. The second powder-

Jan 1, 1970

(The following list contains the names of those members of the Institute of whose connection with military service we have only recently become acquainted; it also includes the names of a few who have recently been promoted or transferred, indicated by a *. A complete list was published in the Year Book, issued as a supplement of the Bulletin for March, 1918.) ANDERSON, AMIL A., Co. A, 309th Regiment Engineers, 84th Div., American E. F., Care Postmaster, New York. ARMSTRONG, HAROLD K., Pilot, Naval Aviation, U. S. Naval Air Station, Miami, Fla. BASSETT, ARTHUR F., Co. 6, -E. 0. T. S., Camp Humphreys, Va. BEDDALL, THOMAS H., Adjutant, 1st Battalion, 1st Gas Regiment, American E. F. BERNARD, CLINTON P., Candidate, 12th Observation Battery, F. A. C. 0. T. S., Camp Taylor, Ky. BROOKE, LIONEL, 50th Co., 13th Battalion, 5th Replacement Engineers, Camp Gordon, Ga. BROWNING, EDWARD, Corp., 1st Regiment Engineers, Co. D, American E. F. BUGBEE, EDWARD E., S. A. T. C. Headquarters, Raleigh N. C. BUTTERFIELD, G. BRUCE, CO. C, 605th Engineers, American E. F. CADOT, J. J., 1st Lieut., U. S. Air Service, American E. F., Care Post- master, New York. CARLOCK, J. B., Major, 2d Battalion, 1st Gas Regiment, American E. F., France.

Jan 12, 1918

Taking advantage of the lessons in administrative organizations which were taught by the war, Director Van H. Manning has put into effect a new form of organization in the Bureau of Mines. The Bureau is divided into two branches, one the investigations branch and the other the operations branch. F. G. Cottrell, who has been chief metallurgist of the Bureau, has been promoted to the position of assistant director and will be in charge of the investigations branch. F. J. Bailey, who has been chief clerk of the Bureau, has been placed in charge of the operations branch with the title of, assistant to the director. All research and technical control will be exercised through the investigations branch, while the non-technical work, administrative duties, and dissemination of the results of the Bureau&apos;s activities will be handled by the operations branch. Under the investigations branch comes the division of mineral technology, in charge of Dr. C. E. Parsons; the division of fuels, under the chief mechanical engineer, 0. P. Hood; the division of mining, under the direction of George S. Rice, the chief mining engineer; the division of petroleum and natural gas, under the direction of J. 0. Lewis, the chief petroleum technologist; and the division of experiment stations, under the direction of D. A. Lyon.

Jan 9, 1919

By Shun-ichi Satoh

A NEW method of treating steel by ammonia has recently been intro¬duced into the metallurgical industry, by means of which piston cylinders of internal combustion engines, shafts and other similar steel products can be made more resistant to wear than formerly. In 1929 the writer&apos; published the historical notes and some experiments upon nitrification. He believes that the corrosion of nitrified steel is a matter of importance and worth studying; and as there have been few studies thereof, he has conducted experiments in this respect. L. Guillet,2 in 1927, reported corrosion tests made on nitrified and cemented steels with acid and salt solutions. In 1928, 0. Hengstenberg3 reported the high resistance of nitrified steel to water vapor and air. V. 0. Homerberg4 stated that nitrified steel has high resistance to water, air and salt solution; and G. Hagg5 studied nitrified layers of iron by means of the X-ray. Though F. Hanaman6 in 1913, measured the single potential difference of steel treated with ammonia at above 650° C., no one, it seems, has ever measured the single potential difference of the so-called nitrified steels having extraordinary hardness. The writer measured the single potential differences of nitrified special steels in various solutions with the object of studying the corrosion problem and, if possible, of knowing whether any new physical property would be caused by nitrification.

Jan 1, 1932

By R. J. Wasilewski

Electrical resistivity variation with temperature was measured on a series of alloys containting up to 33 at. pct of oxygen over the range 77° to1500°K. The resistivity behavior is highly anomalous and itzconsistent with simple metallic conduction. Both composition and temperature-depended resistivity singularities were observed. A few experiments carried out on Ti-N and Zr-O alloys indicate the presence of similar anomalies. These observations, together with the published data on effects of substi-tutional alloying on the resistiuity of titanium, suggest that the anomalies are inhevent in the electron structure of this group oj metals. The existence of two-band conduction, and a significant shift of bands relative to each other with temperature and/or the electron concentration are suggested. CONSIDERABLE advances have been made in recent years in the alloy theory of simple metals. Very little, however, is known about the bonding in transition metals and their alloys.&apos; Titanium, with its relatively few electrons, may be expected to show simpler alloying behavior than the more complex transition elements. Its alloys with the interstitial elements appear particularly attractive in an investigation of bonding characteristics because of a) the simple nature of the solute elements, b) the remarkable similarity between the equiatomic structures Tic, TiN, and TiO, and c) the extensive solid solubility ranges of oxygen and nitrogen in a titanium reported.2,3 The Ti-O system was chosen for the most extensive investigation because of the relative ease of preparation of suitable specimens. Since the main object of the work was to obtain data on the bonding and its changes on alloying, electron-sensitive properties were primarily investigated. The present work describes the investigation on the electrical resistivity-temperature-oxygen content relationships. A few experiments were also carried out at selected compositions in the Ti-N and Zr-O systems. EXPERIMENTAL Materials and Method. Polycrystalline specimens were prepared in the form of hairpin strips some 50 by 5 by 0.15 to 0.50 mm by direct metal-gas reaction. This was carried out by controlled oxidation followed by a homogenizing anneal at a higher temperature. All the test specimens were fully homogenized as judged from the uniform microstructure and microhardness. To avoid preferred orientation, each strip specimen was annealed in the ß range prior to the oxidation, this procedure assuring random orientation in the strip;4 hence any texture resulting from the oxidation reaction itself affected all the specimens to a similar extent. Titanium used was of high purity (66 DPN, 10 Kg load; major impurities 0,-43G ppm, N,-70 ppm, C-25 ppm, Fe-14G ppm). The solute content of the alloys was determined by weighing, after the reaction with a known amount of oxygen. The specimens in which the discrepancy between the volumetric and gravimetric measurements exceeded 2 pct (or 0.2 mg for the low oxygen alloys) were rejected. The mean between the two measurements was then taken as the oxygen content of the alloy. Check analyses showed no measurable nitrogen contamination. All oxygen contents are given in atomic percent. Zr-O alloys were prepared in identical manner from hafnium-free crystal bar metal, cold-rolled to strip 0.25 mm thick. Ti-N alloys required very long reaction times at the maximum temperature available (1250°C). In order, therefore, to detect possible oxygen contamination, duplicate specimens were reacted in every experimental run, and one of these was analyzed both for oxygen (vacuum fusion) and for nitrogen (Kjeldahl). Only the specimens in which the check analysis showed < 1000 ppm O were then used for resistivity investigation. Since only relatively high nitrogen alloys (7.1 at. pct; i.e., 2 wt pct N,) were investigated, this oxygen contamination was considered permissible. Dc resistance was measured by the four-probe method as previously described.= The temperature was determined with a calibrated thermocouple placed in the center of the specimen hairpin. The errors in the specimen resistance values thus obtained were estimated at 1 pct due almost exclusively to the finite thickness of the potential wires and the consequent uncertainty as regards the true resistance length of the specimen. For the calculation of the specific resistance, however, no dimensional measurements could be carried out on most of the

Jan 1, 1962