Search Documents

By P. C. Johnson, B. A. Stein

This paper describes a study of the effects of combined heat treatment and explosive loading on the mechanical properties of high-strength steels. nis program investigated two distinct areas: 1) the effect of shock waves, without gross irreversible defmmution, on a 3-Cr steel at various stages of heat treatment; and 2) the effect of rapid deformation (explosive forming) on H-11 and D6-AC steels in the metastable austenitic state. The mechanical properties of these steels were improved, in some cases markedly, as a result of these treatments. ,AUSFORMWG, which requires the plastic deformation of metastable austenite, is a process which can appreciably improve the properties of selected alloy stee1s.l,2 The Ausform process significantly increases the strength of these steels without decreasing their ductility. The properties at high temperatures are also improved through a change in the response of the steels to tempering. Although the mechanism by which ausforming alters the properties of these steels is not fully understood, it appears that the dislocation arrays produced by deformation of the metastable austenite influence the structure of the martensite on subsequent transformation. This, in turn, affects the strength, ductility, and tempering response of the martensite. This research used chemical explosives to deform steels at various stages in their heat treatment in order to improve the properties of these steels. The explosive energy is used in two ways; 1) high-pressure shock waves are propagated through the steel to produce extensive microscopic shear strain without causing a large irreversible change in shape, and 2) explosive energy is used to cause extensive macroscopic plastic strain in the metastable austenitic state (explosive forming). I) AUSFORMING WITH INTENSE SHOCK WAVES The steel used in this phase of the research was an alloy having a nominal composition 0.43 pct C, 3.0 pct Cr, 1.5 pct Ni, and 1.5 pct Si. The steel was subjected to intense shock waves in three conditions: 1) in the metastable austenitic state, 2) in the tempered mar tens itic state, and 3) in the tempered martensitic state after ausforming by conventional techniques. The specimens were in the form of disks 2.75 in. in diam and 5/16 in. thick. These were incorporated into a specimen assembly consisting of two disks pressed into a 5 by 5 by 1 in. block of stainless steel, Fig. 1. Spalling (or scabbing) is confined to the front disk. The specimen is protected from oxidation and decarburization by the surrounding metal. The temperature of the assembly is monitored by a thermocouple inserted into one side of the stainless steel block. The assembly is positioned over an oil reservoir which serves both as a means of catching the disks and as quenching medium for the disks shocked under ausforming conditions. Plane shock waves are introduced into the assembly by a metal driver plate impacting the top surface of the block. The driver plate is accelerated by a chemical explosive sheet supplied by E. I. du Pont de Nemours & Co. All the specimens were subjected to plane shock waves having a peak pressure of approximately 430 kbar. The pressure is that quoted by G. E. Dieter for the plane wave generator used in this work.' The driver plate used was 1/4 in. thick, so that the initial pulse was essentially a l/2-in.-wide square wave. The attenuation of the peak pressure during the subsequent 1/4 in. is estimated to be less than 5 pct. The shock front induces a temperature rise, a portion of which is irreversible. Rough estimates (+25 pct) of this temperature rise have been made for iron shocked at room temperature.4 For a 500-kbar shock wave, the temperature rise in the shock front is about 700°F, and is held for a time of the order of microseconds. The irreversible temperature rise, which remains after the shock wave passes, is about 450°F.4 The disks are quenched to room temperature within a few seconds of the shock treatment. It should be emphasized that the temperature rises given above are estimates for pure iron at room temperature, and are not necessarily true for the tests made in this work. The disks shocked at temperatures in the metastable austenitic range were austenitized in the stainless steel assembly in a furnace protected from the firing area. The assembly was removed from the furnace and placed over the recovery reservoir. The plane wave generator was then positioned and

Jan 1, 1963

By WM. E. Mlligan

IN the last year, much information had become available as to the extent that zinc participated in the war effort. The importance of foreign zinc in this program had been indicated by Bateman (M&M April 1945) as of the order of 400,000 tons annually in zinc or contained zinc. Pehrson (M&M April 1945) reviewed the domestic ore-reserve situation with respect to most minerals and metals and showed the extent to which commercial zinc ores had become depleted in this country. This picture was further darkened by O.P.A. figures appearing in E.&M.J. July 1945 in which figures for total production during the three previous years were broken down into percentages produced under ceiling and various quota price categories. True, factors other than depleted ore reserves contributed to this picture. Among the more important of these, were shortage and high cost of labor, both skilled and unskilled, and the general difficulty and cost of maintaining adequate equipment and supplies. The cost of production in the Tri-State district was most affected.

Jan 1, 1946

By G. B. Misra

While investigating on the aerodynamic resistance of airways supported with peripheral timber sets, at regular intervals, the following theoretical equations were developed by the author to estimate the resistance coefficient of such airways: [ ] for S < 1, where f is Darcy-Weisbach resistance coefficient of the airway, C is modified drag coefficient of the supporting member, D is equivalent diameter of the bare airway, 8 is ratio of the approach velocity over the sets to the average velocity of the bare airway, A is cross-sectional area of the bare airway, a is projected frontal area of the sets, A., is cross-sectional area of the air stream at the vena contracta inside the set, S is spacing of the sets, f, is resistance coefficient of the bare airway, l is length of aerodynamic influence of sets, p is perimeter of the bare airway, p, is setted portion of the perimeter of the bare airway, pe is unsetted portion of the perimeter of the bare airway, and P shielding factor. The equations were verified experimentally in a model rectangular airway supported with one- (bars), three-, and four-piece sets of square-section timber of three different sizes and were found to hold true. The work has been further extended to one-, three-, and four-piece sets of round timber of 2.6, 3.2, and 3.8 cm diam with the same experimental set up. Tests have been carried out for spacings of 25, 50, 75, 100, 150, and 200 cm over a regime of flow defined by the Reynolds number (with respect to the equivalent diameter of the bare duct) ranging from about 1.5 X 106 to 5 X 106 using the same experimental techniques. The values of f are calculated in the manner indicated in [Ref. 1]. Unlike with square-section timber, the resistance coefficient f of the airway setted with round timber shows a distinct variation with the Reynolds number of flow. This conforms to observations made by Sales and Hinsley.2 In order to have a comparable value of f for all types of sets with all sizes of timber, it was necessary to select the value of f at a fixed Reynolds number of flow. Since f is chiefly a function of the drag coefficient of the sets, the appropriate Reynolds number RE is that with respect to the diameter of timber in the set. Considering the diameters of timber used and the regime of flow over which measurements were made, f was chosen at a value of RE = 20,000 in all cases. The f vs. S curves are maximal in nature and in conformity with theory, the f vs. 1/S curves are straight lines up to a value of S = 1 beyond which they show a distinct flexure. The observed values of 1, the length of aerodynamic influence of sets, agree with the relation 1 = 42 e, developed for square-section timber sets, thus suggesting that the shape of timber has little influence on the length of aerodynamic influence. The value of the modified drag coefficient CD for round timber was calculated in the same way as for square timber in Ref. 1, taking the contraction factor Z = 1.5 for round-edged constrictions. CD has an average value of 0.96 with a standard deviation of 6.08% as compared to the free stream drag coefficient of 1.2 at RE = 20,000 for long cylindrical obstructions The shielding factor [ ] is plotted against S/1 in [Fig. 1]. The curves are more or less independent of the size of timber, but are different for the different types of sets, possibly due to their different degree of symmetry. Values of f calculated by the author&apos;s [Eqs. 1 and 2], using experimental values of CD&apos; and [ ] and taking I = 42 e, are plotted in [Fig. 2] against experimentally measured values of f for different types of sets with different sizes of round timber. The values agree closely with a standard deviation of only 5%, thus establishing the veracity of the theoretical equations developed by the author for round timber as well. A comparison was made between the Xenofontowa4 equations (the only other reasonable relations available for the estimation of the resistance coefficient of supported airways) and the author&apos;s [Eqs. 1 and 2] by comparing in [Fig. 3] the values of the resistance coefficient f computed by the Xenofontowa relations with those experimentally measured by the author. In order to make

Jan 1, 1975

By T. E. Leontis

The specific effect of various rare-earth metals on the room- and elevated-temperature properties of magnesium has been evaluated. Alloys containing didymium exhibit the highest tensile and compressive strengths at room and elevated temperatures. All the rare-earth metals increase the creep resistance of extruded magnesium at temperatures in the range of 400° to 600°F, but the degree of enhancement depends on the temperature and on the concentration of the added metal. THE effects of rare-earth metals on the properties of sand-cast magnesium were discussed in some detail in earlier paper by the author.&apos; The present paper deals with the effect of the same alloying elements on the properties of extruded magnesium. This investigation also had as its aim the development of a wrought alloy having a better combination of room-temperature strength and ductility and elevated-temperature strength and creep resistance than is found in magnesium-Mischmetal-manganese alloys, which have been reported earlier.2-5 The only known attempt to study wrought magnesium alloys containing pure cerium instead of Mischmetal was made by Mellor and Ridley.6 They found that in the form of rolled bars there is a definite, optimum cerium content for creep resistance at 200 °C and that the creep resistance of these alloys at 200 °C is significantly Improved by heat treatment at 550" to 580"C. In the present investigation the compositional variation in mechanical properties of the following alloy systems is presented: I—magnesium-Misch-metal. 2—magnesium-cerium-free Mischmetal. 3— magnesium-didymium. 4—magnesium-cerium. 5— magnesium-lanthanum. Alloys containing predominately praseodymium are not included in this series because of the lack of this material. Experimental Procedures The alloying ingredients used in preparing the alloys described herein are the same as those reported in the earlier paper.&apos; Cerium-free Mischmetal is the rare-earth mixture remaining when the cerium is removed from Mischmetal, which contains all the rare-earth metals as they occur naturally in mon-azite sand, the ore from which Mischmetal is produced. Removal of both cerium and lanthanum from Mischmetal leaves what is commonly called "didym-ium," consisting predominantly of neodymium and praseodymium. Although the composition of the particular batch of each metal used may differ somewhat from the analysis presented previously, these differences are not great enough to warrant repeat- ing the specific composition of each material. The electrolytic magnesium used as the starting material in these alloys has the same typical analysis as that given in the earlier paper.&apos; The alloys were prepared in small laboratory melts applying all the necessary precautions for alloying rare-earth metals with magnesium described by Marande.&apos; Most melts were large enough to cast one 3 in. diam billet 10 in. long. In a few cases, particularly the didymium-containing alloys, the lack of sufficient amounts of the rare-earth metal limited the size of the billet to 6 to 8 in. All billets were scalped to a diameter of 2 15/16 in. and faced to a length of 9 ¼ in. as limited by the size of the extrusion container. The alloys were extruded into ½ in. diam rod on a 500-ton direct-extrusion press using a 3 in. container. The details of the extrusion step are: billet preheat, 925°F (2 hr); container temperature, 900°F; die temperature, 900°F; extrusion speed, 10 ft per min; reduction ratio, 36:1; and percent reduction, 97.3. The lower melting point of alloys containing didymium&apos; necessitated reduction of the extrusion speed to 5 ft per min in order to prevent hot shorting during extrusion. Adequate lengths were cropped from both ends of each extruded rod to assure that all the material used for tests was extruded under uniform conditions. Tensile and compressive properties at room temperature are reported in the several conditions of heat treatment. The ASTM designations are used to distinguish these conditions as follows: T5—Direct age at 400°F (16 hr) T4—Heat treat at 950°F (4 hr) for alloys containing didymium T4—Heat treat at 1050°F (4 hr) for all other alloys T6—T4 + age at 400 °F (16 hr) The lower heat-treating temperature for alloys containing didymium is necessitated by the lower melting point of these alloys. All heat treatments were conducted in electrically heated, circulating-air furnaces. A protective atmosphere containing 0.5 to 1.0 pct sulphur dioxide was used for the high temperature heat treatments. Tension and creep specimens 6 Yz in. long and compression specimens 1½ in. long were cut from the extruded rod. A reduced section of ? in. diam was machined on the tension specimens, whereas on the creep specimens a reduced section of 0.450 in. diam

Jan 1, 1952

By W. L. Phillips

Stress-strain curves were obtained for single crystals of 70 pct Ag-30 pct Zn tested in tension and shear. Samples tested in tension and shear had comparable resolved shear stresses and stress-strain curves. The {111} <110> slip system was observed. It zoas found that the9.e is a barrier to slip in both latent close -packed directions and that the magnitude of these barriers is proportional to prior strain during easy glide. It was observed that cross-slip in tension and shear was most frequent in crystals with an initial orientation near <100> "Oershoot" zoas observed in tension. The amount of this "overshoot" was independent of initial orientation. AN idealized concept of plastic deformation indicates that a single crystal should yield at some stress that is dependent on crystal perfection and it should then continue to deform plastically by the process of easy glide which is characterized by a linear stress-strain curve and a low coefficient, d/dy, of work hardening. Hexagonal metal crystals generally conform to this ideal concept of laminar flow. In fcc metals the range of easy glide is always restricted in magnitude and it is strongly dependent on orientation, composition, crystal size, shape, surface preparation, and temperature. Since one of the principal differences between the two crystal systems, both of which deform by slip on close packed planes, is the existence of latent slip planes in the fcc crystals, it has been proposed that the transition from easy glide to turbulent flow, characterized by rapid linear hardening, is due to slip on secondary planes intersecting the primary plane.ls Several theories have been proposed to explain the linear hardening and parabolic stages of the stress -strain curve.6"10 The easy-glide region is the least understood of the three stages. The stress-strain characteristics of Cu-Zn, which shows a long easy-glide region, have been extensively investigated."-" In light of recent ideas on dislocations, cross-slip, effect of solute atoms, and stacking fault energy, it was felt that the certain features of this earlier work might be compared with another alloy, Ag-30 pct Zn, which also exhibits a long easy-glide region. Tension and shear stress at room temperatures were employed. The results obtained, together with some interpretation of the observations, are described below. EXPERIMENTAL PROCEDURE The silver and zinc used for mixing the alloys were 99.99 pct pure. The two components were weighed to within 0.1 pct of the weights required fo the alloy composition. They were then placed in a closed graphite mold and the mold and contents were heated in 100°C stages from 500&apos; to 900°C with sufficient time and vigorous agitation at each stage provided to dissolve the silver. The crucible was then heated to 1150°C and agitated violently before being quenched in oil. The resulting alloy rod was machined free of sur face defects and then placed in a graphite mold designed for growing single crystals. The graphite mold was closed with a graphite plug and was encased in a pyrex glass tube which was connected to a vacuum system. The tube and mold assembly were placed in a furnace; the tube was evacuated and the furnace was rapidly heated to a temperature sufficient for fusing and sealing the glass. The glass-encased evacuated mold and contents were then lowered through a vertical furnace. The top section of the furnace was held at 100 °C above the melting point of the alloy. The lowering rate was 1.5 in. per hr. The tension specimens were 1/4 in. diam; the shear specimens were 1/2 in. diam. These specimens were then removed from the mold, etched, and chemically polished with hot (60°C) Chase etch reagent (Crz03-4.0 g, NH4C1-7.5 g, NHOs-150 cc, HzS04-52 cc, and Hz0 to make 1 liter). In preparation for tensile testing, the specimens were carefully machined to a diameter of about 0.200 in. to permit a gage length of 6 in., annealed for 16 hr at 800&apos; to reduce coring, and then cleaned and polished. A modified Bausch-type shear apparatus which has been described previously18 were employed. The gage length was 1/8 in. This shear apparatus was placed in an Instron tensile testing machine. EXPERIMENTAL RESULTS A) Tension. Several specimens were extended at room temperature to determine the effect of initial orientation on the stress-strain curves of Ag-30 pct Zn. The initial orientation and the resolved shear stress supported by the active slip system at various total strains are plotted in Fig. 1. The critical resolved shear stress, t,, initial rate of work hardening, d/dy, and length of the easy-glide region are independent of orientation. The arrival at the symmetry line is shown by an arrow in Fig. 1. During the easy-glide region of the stress-strain

Jan 1, 1963

By J. A. Porter

In offering some data relative to this property, and the treatment of its ores, it is proper to say that a purely scientific article has not been attempted. It is hoped, however, that the economic conditions presented may be of interest to the members of the Institute, and that a discussion upon the concentration of the ore may throw some light upon a very unsatisfactory part of the work. The Smuggler-Union mines are located at the head of the San Miguel river, just south of the water-shed between Ouray and Telluride. The latter town, reached by the Denver and Rio Grande railway, is most picturesquely situated, at ail altitude of 8500 feet, in a narrow valley, the lower portion of which exposes only the sedimentary rocks. Along the sides of the valley, large masses of coarse alluvial deposits are left high above the present channel of the stream. Although much placer-work has been done, and some very rich material has been found, the limited extent of the isolated tracts has made placer-mining in this locality generally unprofitable. The valley terminates two miles above Telluride in an amphitheater around which walls rise abruptly several thousand feet. Here the company&apos;s mill is situated, being reached by a Bleichert tramway from the mines and also by a spur of the Rio Grande railway from Telluride. At this point, not only are the reddish sandstone and conglomerate beds exposed, but the contact between the conglomerate and the overlying gray volcanic breccia is plainly in view, even at a distance of many miles. A marked stratification is noticeable, indicating the successive layers of igneous breccia, locally termed trachyte. Throughout the entire region this is the favorite name among the miners for the rocks composing the igneous cap of the San Juan

Jan 1, 1897

By David Segura

THE Huitzuco mercury deposits have been described by C. W. Vaupell,1 therefore the general information such as location, history, production, etc. will not be repeated here. The present owners acquired the property when the deep levels of the Cruz mine were flooded, and it was believed that a relatively large tonnage of low-grade ore and fills would be encountered in the flooded section, because of the early mining methods used in the Cruz mine. The mine was partly unwatered and about 187,000 tons of fills was pit-sampled in the San Cayetano stope. These fills average 2.4 kg. Hg per ton. During the period of maximum production at Huitzuco, the ores were sorted in the mine and cobbed into 3-in. pieces, which were treated in Scott furnaces. Large pieces of raw ore are found in the sinter, and it is indicated that both recovery and costs were not very satisfactory. Recent practice has been to treat the ores in closed retorts made of 10-in. pipes. Because of the presence of combined antimony in the mineral livingstonite (HgS.2Sb2S3) and considerable native sulphur, little or no metallic mercury is recovered in the condensers and a second retorting of the condenser muds is necessary. This makes the treatment of the ores expensive and tends to low recoveries of mercury. The principal costs of retorting are, in their relative order: (1) fuel, (2) low recovery of mercury, (3) replacement of retort pipes, (4) labor. METHODS TESTED Three series of experiments have been made, with the object of reducing the costs of recovering mercury and antimony from the ores, by: (1) wet metallurgical methods, (2) bulk roasting in large mechanical furnaces, (3) concentration to reduce the bulk to be followed by roasting in some mechanical furnace. The wet metallurgical methods did not appear feasible because of the high cost of the alkaline solvents used, and investigations of these were not carried beyond the preliminary stage.

Jan 1, 1938

By E. Teghtsoonian, K. G. Davis

THE preparation of cobalt crystals offers problems: on cooling through 400°C a phase transformation takes place whereby the structure changes from face-centered cubic to the low temperature close-packed hexagonal form. There are a few references in the literature which describe the preparation of cobalt crystals for magnetic work1,&apos; where melt-solidification methods were used. Reasons will be given why such methods were found unsuitable for our purposes. The following techniques were tried: 1) Solid state methods. Any mechanical shaping of tensile specimens would lead to uncertainty with regard to residual strains. With this in mind, following Kaa, strain-anneal and grain growth techniques were extensively investigated. In no Case were useful single crystals obtained. The transformation boundaries, which are crystallographically similar to twin boundaries, are probably extremely stable and inhibit the grain growth. 2) Growth from the melt. Normal melt-solidification methods were tried, using recrystallised alumina crucibles, and also a osoft moldo technique in which cobalt rod was packed in alumina powder. Trouble was encountered with gas evolution on solidification of the melt, giving porosity. Some adhesion of cobalt to the alumina also took place. Only a small proportion of the runs gave single crystals. 3) Zone-refining. Much the best results have been obtained using an electron-beam floating zone refiner. The refiner, based on a design by Calverly et Al .5 had a single-turn tungsten filament. It was operated at 600 v with an emission current around 45 ma. Commercial cobalt rod 1/8 in. in diam was given one pass at a rate of zone travel normally 25 cm per hr, under a vacuum approximately 10-5mm of Hg. Over 50 pct of the runs gave long sections of single crystal. Crystals up to 2o Cm long with good uniformity of cross section and straightness have been grown in this way. Increase in the length of the molten zone resulted in higher single-crystal yields, but control of zone stability was reduced. Sub-boundaries in the zoned crystals were few, and Laue X-ray spots were sharp. The success of the method is believed to lie in two factors, namely i) the absence of crucible restraints during the transformation, and ii) extreme axiality of heat flow during cooling. The axes of the zoned crystals were clustered around the [ 1010] pole, the pole of the basal plane being at least 60 deg from the specimen axis. Crystals grown at lower rates of zone travel showed the same orientation range, Fig. 1. This orientation is similar to that for other hexagonal metals grown from the melt, which is a little surprising when the mode of formation of the hexagonal phase is considered. The preferred orientation is probably caused by anisotropy of heat conduction, i.e., that crystal which can most effectively conduct heat away from the transformation interface will be favored. Fig. 2 shows the area on a standard cubic stereo-gram in which there are no poles more than 60 deg from one of the four < 111> directions. This area

Jan 1, 1962

By D. T. Peterson, T. Carnahan

The diffusion coejTicients of carbon, nitrogen, and oxyget were determined in $ thorium over the tempernilcre range 1440" io 1715°C. The diffusion coyfiicir?zls are given by: D = 0.022 exp (-27,000/RT) jor carbo)~, D = 0,0032 exp(-l7,00Q/RTj for nitrogen, and D = u.0013 expt(-11,UOU/RT) for oxygen. Cavl~orz was found to increase the hardness of thoriunz nearly linearly with concentration over the range 100 to 1000Ppm carbon. ThORIUM has a fcc structure up to 1365°C and a bcc structure from this temperature to its melting point at 1740°C. Diffusion of carbon, oxygen, and nitrogen in bcc thorium was of interest in connection with the purification of thorium by electrotransport.&apos; In addition, it was possible to measure the diffusion of all three of these interstitial solutes in the same bcc metal. Only in niobium, tantalum, vanadium, and a iron have all three interstitial diffusion coefficients been measured in a given bcc metal. Diffusion coefficients have been measured for carbon and oxygen in a thorium by Peterson2, 3 and for nitrogen by Gerds and Mallett.4 Activation energies for diffusion are reported by the above authors to be 38 kcal per mole for carbon, 22.5 kcal per mole for nitrogen, and 49 kcal per mole for oxygen. Values of the diffusion coefficients of carbon and nitrogen in 3 thorium have been reported by Peterson et al.&apos; However, these were secondary results of their investigation of electrotransport phenomena in thorium and it was hoped that the present study could provide more precise data. EXPERIMENTAL PROCEDURE The specimens used in this study were the well-known pair of semi-infinite bar type. The couple was formed by resistance butt welding two 0.54-cm-diam by 3.0-cm-long bars of thorium together under pure helium, the concentration of the solute being greater in one cylinder than that in the other. The finished couple then contained a concentration step at the weld interface and diffusion proceeded only along the axis of the rod. The thorium used in this study was prepared by the magnesium intermediate alloy method.5 The total impurity content was less than 400 ppm. The major impurities were: carbon, 100 ppm: nitrogen, 50 ppm; and oxygen. 85 ppm. The total metallic impurity content was less than 150 ppm. The high solute concentration portions of the diffusion couples were prepared by adding the solute to the high-purity thorium in a non-consumable electrode arc melting procedure. Carbon and nitrogen were added in the form of spectroscopic graphite and nitrogen gas while a Tho2 layer was dissolved by arc melting to add oxygen. High-purity thorium formed the low concentration portions in the carbon and nitrogen couples. The low oxygen portions were obtained by deoxidizing high-purity thorium with calcium for 3 weeks at 1000°C according to a method reported by Peterson.3 The high C-Th contained 400 ppm C, the high N-Th contained 400 ppm N, the high 0-Th contained 220 ppm 0, and the low 0-Th contained 25 ppm O. The high O-Th was brine-quenched from 1500°C to retain most of the oxygen in solution at room temperature. These concentration levels were all below the solubility limits in 0 thorium at 1400°C. A resistance-heated high-vacuum furnace was used to heat the couples. The samples were mounted horizontally on a tantalum support which had small grooves near each end. Spacer rods of thorium, 0.4 cm in diam, were placed in these grooves to prevent contact between the sample and the tantalum support. This arrangement should have prevented contamination of the sample by contact with the support. In further effort to reduce contamination, the oxygen diffusion couples were sealed inside evacuated outgassed tantalum cylinders lined with thorium foil. Thorium rings around each end of the samples acted as spacers in this case. Pressure during diffusion runs was about 10-6 torr after an initial outgassing stage. Temperature measurements were made by sighting on black body holes in the sample support adjacent to the samples with a Leeds and Northrup disappearing-filament optical pyrometer. Temperatures were constant during a diffusion anneal to ±5C. The observed temperatures were corrected for sight glass absorption after each diffusion run. The pyrometer was checked against a calibrated electronic optical pyrometer and a calibrated tungsten strip lamp with the electronic pyrometer being taken as the standard. All temperature readings agreed to within ±3C over the temperature range 1450" to 1690°C. Time corrections due to diffusion during heating and cooling were necessary because of the short diffusion times. The diffusion times ranged from 6 min for the oxygen sample run at 1690°C to 90 min for the carbon sample run at 1500°C. A series of temperature vs time plots were made for heating and cooling of the samples to the various diffusion temperatures. This data was then used in a method according to shewmon6 to determine the time corrections. The corrections amounted to

Jan 1, 1970

By H. I. Aaronson, C. Wells

Configurations of ferrite crystals have been found in a plain carbon steel which appear to have resulted from the nucleation of new ferrite crystals at the interphase boundaries of previously formed crystals despite the high carbon concentrations which necessarily develop at these boundaries. This phenomenon has been termed sympathetic nucleation. An attempt has been made to reconcile the occurrence of sympathetic nu-cleation with current nucleation theory. THIS investigation is one of a series on the formation of proeutectoid ferrite from austenite. From the viewpoint of chemical composition, this reaction consists of the nucleation and diffusional growth of crystals of carbon-poor ferrite within a matrix of carbon-rich austenite. The austenite adjacent to the austenite-ferrite boundaries will be greatly enriched in carbon, approximately to the value of the y/(a + y) equilibrium curve or its metastable extrapolation at the temperature of transformation. Those areas of austenite appreciably farther removed from the growing ferrite, on the other hand, will be relatively unaltered in composition, especially at the earlier stages of transformation. Since rates of nucleation are considered to decrease exponentially with decreasing supersaturation,&apos; the frequency with which ferrite nuclei appear at austenite-ferrite boundaries should be negligible in relation to that at which they form in other regions of the austenite. During this investigation, however, many groupings of ferrite crystals have been found which appear to have resulted from the nucleation of ferrite at austenite-ferrite boundaries. This phenomenon has been given the name of sympathetic 71.1tcleation. A number of micrographs of morphological configurations caused by sympathetic nucleation will be presented, after which an explanation for this reaction will be proposed in terms of current nucleation theory. Some of the structures to be considered are composed of bainite, an aggregate of ferrite and carbide, rather than of ferrite. Since ferrite and bainite differ only in that bainite forms under conditions which result in the nucleation of carbides behind the advancing austenite-ferrite boundaries,&apos; it will usually be unnecessary, for the purpose of this paper, to distinguish between the two reaction products. All studies were performed on an electric furnace steel (obtained from the Vanadium Alloy Steel Co.) containing 0.29 pct C, 0.76 pct Mn, 0.25 pct Si, 0.005 pct P, and 0.007 pct S. The alloy was cast as a 150 Ib, 7x7 in. cross section ingot and forged into bars 2x2 in. in cross section. These bars were homogenized for 48 hr at 1250°C in an Endo-Gas atmosphere. The depth to which decarburization penetrated during this heat treatment was determined by chemical and microscopic analyses and the affected metal was removed by machining. Specimens for isothermal transformation studies were cut from the remaining material; most of these specimens were 1/2x1/4X1/16 in., though some with a thickness of 1/32 in. were prepared for use at the shorter reaction times and lower reaction temperatures. Specimens were austenitized for 30 min at 1300°C, isothermally reacted for various times at temperatures ranging from 775" to 475 "C, and then quenched in iced water. The austenite grain sizes within individual specimens ranged from ASTM Nos. 1 through —4. A commercial heat-treating salt which was continuously deoxidized by an immersed graphite crucible served to minimize the loss of carbon during austenitizing; thick covers of powdered graphite and immersed graphite rods effectively prevented decarburization in the lead pots employed for the isothermal reaction treatments. The heat-treated specimens were sectioned and mounted in Bakelite. Following the completion of standard grinding and mechanical polishing procedures, the specimens were electrolytically polished with a Buehler-Waisman apparatus and etched in 2 pct nital. Experimental Results Rules of Evidence for Sympathetic Nucleation—On the basis of observations made on a single plane of polish, one precipitate crystal may be considered to have been sympathetically nucleated at the inter-phase boundary of another precipitate crystal when the following conditions are fulfilled: 1) The sympathetically nucleated crystal is not in contact with a grain boundary or a subboundary in the matrix phase. 2) The shape, size, and location of the crystal at whose boundary sympathetic nucleation occurred (hereafter termed the base crystal) and the crystal formed by sympathetic nucleation substantially pre-clude the possibility that the plane of polish em-ployed may have concealed the fact that both crys-tals actually nucleated at a grain boundary or a sub-boundary in the matrix phase.

Jan 1, 1957

By Paul A. Farrar, Harold Margolin

The system Ti-Cr has been re investigated From 40 to 100 pct Cr. The high-temperature modification of TiCr, was found to exist in the temperature range from 1365 where it terminates in the eutectoid reaction TiCr,(TI)* — HE chromium-rich region of the Ti-Cr system was investigated by Vogel and Wenderott,&apos; Mc Quillan,&apos; Van Thyne et dl.,&apos; and Cuff Duwez and Taylor determined the crystal structure of the low-temeperature modification of TiCr, to be fcc a = 6.943A and Levinger&apos; discovered the existence of a second (high-temperature) modification of TiCr, and determited its crystal structure to be hexagonal a = 4.932A, c/a = 1.617i. No information was obtained as to the temperature range over which each modification of TiCr? existed; therefore the present investigation was initiated. I) Experimental Techniques Alloys employed for the delination of this system were prepared from iodide titanium and high-purity chromium. The purity of these materials is as follows: Iodide titanium: purity 99.99 pct; major impurities—carbon 0.0001 pct; oxygen 0 to 0.002 pct; nitrogen 0 to 0.002 pct; silicon 0 to 0.005 pct; manganese 0.001 pct; copper 0.001 pct (Foote Mineral Co., Lot 609-29). Chromium #1: metallic purity 99.997 min.; interstitial content—oxygen 0.27 to 0.46 pct; hydrogen 0.028 to 0.038 pct; nitrogen 0.0002 pct (Ave-quipo Inc.). Chromium #2: purity 99.9 pct; major impurities—carbon 0.01 pct; oxygen 0.013 pct; nitrogen 0.02 pct; silicon 0.01 pct; sulfur 0.03 pct; iron 0.01 pct (Fulmer Research Institute Ltd.). A total of sixteen alloys of thirteen different compositions was prepared in the range from 40 to 95&apos; *All percentages are given in weight percent. pct Cr. The compositions of the alloys prepared were as follows: Ti-40, 45, 50, 60, 62, 64, 66, 68, 70, 80, 85, 90, and 95 pct Cr. Two buttons were prepared of the 66, 90, and 95 pct Cr alloys. The 40 pct Cr and one of the 66 pct Cr alloys were prepared from the chromium # 2. All the other alloys were prepared from chromium #1. The alloys were arc-melted as 25-g buttons in a copper-hearth glass-walled furnace7 under an inert atmosphere of high-purity helium. The alloys were sintered at 250 amp and melted twice at 350 amp for 30 sec with the buttons being turned over between melts. Weight losses during melting were determined and in the range 45 to 68 pct Cr they were found, in all cases, to be less than 1.25 pct of the total alloy weight. A number of the high-chromium alloys shattered during cooling from either the first or second melt, which made the weight loss considerably larger than usual. The 70 and 80 pct alloys lost approximately 5.7 pct of their total weight, the 85 pct alloys 1.5 pct, the 90 pct alloy 2 pct, and the 95 pct alloy 3 pct. However, it is felt that the actual composition of the alloys in this region is probably within approximately 1 pct of the intended composition, as a large portion of the material lost was that due to shattering after melting. Nominal compositions were therefore used in plotting data and referring to the alloys. The alloys were cut into specimens, wrapped in molybdenum sheet, and encapsuled in argon-filled quartz or Vycor tubes for heat treatment. The heat-treatment times and temperatures are given in Table I. The specimens were quenched in iced brine (I.B.Q.) following heat treatment. The standard techniques of grinding, polishing., and etching were used in preparing the specimens for metallo-graphic examination. For alloys below 68 pct Cr

Jan 1, 1963

By E. A. Gulbransen, J. W. Hickman

DURING the past two decades considerable progress has been made in the art of manufacturing heater alloys. The conventional iron-chromium-nickel alloys have been improved by the addition of small quantities of other elements. Up to the present time, to our knowledge, there has not appeared a systematic study of these alloys where the newer tools of investigation have been brought to bear on the problem. It is not yet clear what roles the additional elements play in increasing life time in usage. As a first approximation one might guess that elements such as manganese, silicon, zirconium, calcium and aluminum would aid in the formation of a protective film on the alloy surface, thus retarding the oxidation rate and increasing lifetime. If these additional elements do form protective oxides in the outer surface of the oxide film, it may be possible to identify them by means of reflection electron diffraction. It is also possible that the additional elements may not diffuse to the surface but may remain at the interface between the oxide and the metal. Then, too, the major elements in the alloys undoubtedly form oxides which are stratified on the surface. If the minor elements concentrate in a layer in contact with the substrate where they undergo oxidation and if this oxide film offers resistance to the diffusion of iron, chromium and nickel ions, then the lifetime of the heater element might be expected to increase. It is also possible that preferential reduction in the diffusion rates of the several ions may Occur to such a degree that oxides may form in the body Of the film and On the outer surface which are different from those that would normally form if the additional elements were not present. The use Of transmisson electron diffraction may aid us in jnterpreting these phenomena. UP to the Present time all Of the factors which influence the lifetime of a given heater element are not definitely known. Among these are: (I) The rate of oxidation of the alloy. This may be dependent upon the composition and the physica1 state of the "alloy. (2) The stability of the structure of the coating which forms. Transitions from one structure to another or solid phase reactions resulting in a change of composition may occur. The tendency of the film to adhere to the substrate when heated and cooled successively will be important. (3) The evaporation of various constituents from the alloy such as the remova1 of carbon by Oxygen in a decar-burization reaction. With these possibilities in view, an electron diffraction study of the oxides occurring on six different nickel-chromium alloys was undertaken with the hope that the results of this study could be correlated with oxidation rate studies as determined by means of the vacuum microbalance method1 and with lifetime tests carried

Jan 1, 1949

By Edward Saibel

THE object of this study is to investigate the effect of prior tensile strain on the fracture stress of a metal. This is done in a theoretical manner starting from the point of view developed by the author in his thermodynamic theory of the fracture of metals. The results are compared with experimental findings and a rational interpretation is given to work that shows the variation of fracture stress with prior tensile strain. THEORY OF FRACTURE In the thermodynamic theory of the fracture of metals developed earlier,1 it was shown that a critical strain energy per unit volume exists. This is characteristic of the material and may be calculated from basic thermodynamic data. When the area under the flow curve reaches this characteristic value, fracture occurs. The general criterion for fracture is expressed in the form u ± uo = JLm?AV/V [1] where Lm, is the latent heat of melting of the unit volume, ?V/V is the ratio of the change in volume of the unit volume to the original unit volume on passing from the solid to the liquid state, [uo] is an initial energy density, and J is a conversion factor, which changes energy from heat to mechanical units. As a first approximation, it will be assumed that the stress-strain curve is the straight line that extends from the point of maximum load onward, extrapolated back to the axis of zero strain, for example, line AC of Fig. 1. If the point of intersection of the straight line with the axis of zero strain is represented by so, and the slope of the straight line is p, the analytic form of the flow curve is [s = s + pd [2]] and the strain energy per unit volume evaluated from [JvdS] becomes (vsod + pd2/ 2). When this latter term reaches the critical value JLm?V/V (designated hereafter by M), it is assumed that fracture occurs. This leads to the following expression for the fracture stress: [s1 = \/C r.2 + 2Mp [31] The strain at fracture may be calculated from Eqs 2 and 3. It is assumed in this paper that the initial energy density is either zero or that it may be neglected. Furthermore, the complicated states of stress and of strain existing in the necked region of the tensile specimen after "necking" has occurred will also be neglected for the present, so that average conditions will be assumed to prevail. FLOW AND FRACTURE CURVES Ludwig2 attributed to materials two fundamental properties: (I) a resistance to flow and (2) a resistance to fracture. He reasoned that at each instant the

Jan 1, 1947

By Harry H. Power

The Specific purposes of forma! engineering education include training in the basic sciences, the engincering-prob]em method, the rudimentary development of technical skills, an appreciation of values and costs, and an understanding of the art of engineering as distinguished from its science, Two major divisions of an educational program in engineering are recognized: (I) the scientific-technological, and (2) the humanistic-social. Hence, the entire curriculum for the four-year program in specialized petroleum engineering schools must be designed to give the student an initial impetus in the directions indicated. In accordance with suggestions from the Engineering Council for Professional Development, Inore attention is being given to the fundamental approach in the undergraduate curriculum. This paper presents a cross-sectional viewpoint of the various educators in the petroleum engineering schools concerning course contents in the specialized curricula An outline is suggested for the content of the specialized courses wherein the engineering problem method or quantitative approach is emphasized. Although the particular viewpoints of the author have been stressed, yet due acknowledgment is made to a considerable number of petroleum engineering educators and in in industry interested in educational work, for their contribution of course outlines or helpful suggestions with respect to programs that will be of maximum value to students, institutions, and the industry they serve. Although "standardization" of course material is not suggested, the advantage to be gained by cooperative determination of purpose and scope are apparent paiticular1y from &apos;he standpoint of the proper accreditment of petroleum engineering schools by petroleum engineering educetors who have laid a basis from which to measure educational values. Petroleum engineering has "come of age" and deserves the same careful attention with respect to. curricula that has been accorded other engineering branches. Introduction Engineering schools have a major responsibility to the public, to industry and to the professions they serve. The specific purposes of formal education include training in the basic sciences, the engineering-problem method, the rudimentary development of technical skills, an reciaciation of values and costs, and an understanding of the art of engineering as distinguished from its science. Other purposes implied include the ability to read, write, and speak the English language effectively, an understanding of social and human relationships, a knowledge of the duties of citizenship, a broad appreciation of cultural interests, and an indoctrination in professional standards and relations. Throughout these educational processes is the development in the student of habits of accuracy, thoroughness, powers of analysis, creative ability and integrity with respect to all phases of his work. Major Divisions of Program In the Report of the Society for the Promotion of Engineering Education,† two † Report of Committee on Engineering Education after the War, dated January 1944.

Jan 1, 1945

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

Alpha solutes increase the strengths of Ti-Mn alloys through solid-solution strengthening. The substitutional a addition, aluminum, decreases, and the interstitial solutes, carbon and nitrogen, increase the rate of nucleation and growth of a from ß. The best combinations of properties of a-ß alloys are obtained when there is a sufficient quantity of a phase in the structure to dissolve the a solutes. OF the many different titanium-base alloy systems, the predominant alloy type is the a-ß alloy. The properties of the a-ß alloys are dependent on solid-solution strengthening and heat-treatment effects involving the a-ß ratio and transformation reactions. Another variable which influences the mechanical properties of a-ß alloys is the a-stabilizer content of the alloy. An a solute may be present as an intentional addition, such as aluminum, or as an impurity element, such as carbon, oxygen, or nitrogen. It is known that these a stabilizers, when added to titanium, form single-phase alloys which are not heat treatable but which obtain their strength from solid-solution strengthening. Thus, it would be expected that a additions to a-ß alloys would increase the strength of the alloys by solid-solution strengthening of the a phase. In addition, they would affect the transformation kinetics of the ß-to-a reactions and other processes based on the instability of the ß phase. The effects of heat treatment and structure on the mechanical properties of Ti-Mn alloys have been shown in a previous paper.6 This system offers a good base to demonstrate the effects of typical a solutes on the properties of a-ß alloys. The three a solutes described in this work are aluminum, representative of a substitutional a solute, nitrogen, representative of an interstitial a solute, and carbon, representative of an interstitial compound-forming element. The effects of heat treatment and microstructure on the properties of a alloys containing these three elements are described in concurrent publications. Some of these data are used for base-line points in several of the curves used for illustration herein. Experimental Procedures Iodide titanium was used as the base for all of the alloys studied in this work. The alloys were prepared as ½ lb ingots by double arc melting in an argon atmosphere. The ingots were forged to ¾ in. rounds, vacuum annealed for 6 hr at 900°C at a pressure of 10 &apos; to 10-5 mm of Hg to remove hydrogen, and hot swaged to 1/4 in. diam rod. After me- chanical descaling, test specimens were prepared for heat treatment. The alloys used in this study together with the fabrication temperatures are given in Table I. Heat treatments were done in argon. For the most part, the specimens were sealed in Vyeor capsules under a partial pressure of argon. Quenching was accomplished by breaking the capsule under water. Other cooling methods used included oil quench, argon cool (simulated air cool in an argon atmosphere), and furnace cool. The times for the various heat-treating temperatures are given in Table 11. The tests performed on the alloys consisted of tensile tests on ? in. diam specimens, hardness tests, and microimpact tests. Specimen sizes have been adequately described in a previous publication.&apos; The micrographs presented in this paper were taken from specimens cut from the shoulders of broken tensile specimens. Final polishing was done with Linde B on a slow-speed wheel, and the specimens were etched with a 1½ HF — 3½ HNO, solution. Ti-N-Mn Alloys The transformation diagram and microstructures of the Ti-0.1 pct N-Mn alloys used in this investigation are given in Fig. 1. The effect of small nitrogen additions on the binary Ti-Mn diagram is to raise the ß-transus temperature with little effect on the a solubility of manganese. Also, as has been noted previously,&apos; high manganese-content alloys containing nitrogen, when quenched from temperatures high in the ß field, contain a subgrain boundary phase which appears to be nitrogen-rich a. Marten-site is formed when alloys containing less than about

Jan 1, 1956

By Albert L. Toenges

Alaska requires an adequate fuel supply for its development, and has large potential coal reserves ranging from lignite to subbituminous and anthracite. Coal production in the Territory now is less than the requirements. In L947, production was 361,000 tons, divided about equally between bituminous coal from the Matanuska field and subbituminous coal from the Nenana field. There is need for development of modern mechanized mines, which should produce the required output with a minimum of workers. These mines should not be thought of in the light of large potential capacity, as in the States. However, these new mines should be developed by modern methods, which should result in lower costs than at Present. Because of the diverse physical conditions in the coal fields, intensive investigation by diamond drilling is necessary to properly plan the development and equipment of mine sites at minimum cost. The Bureau of Mines is doing this now in the Wishbone Hill area of the Matanuska field. All coal land in Alaska is government-owned and subject to the Coal Leasing Act. Coal mining in Alaska has been handicapped by an inadequate supply of dependable workers. Wage scales have been very high in recent years, and "floaters" have been attracted to the coal mines. This type of labor is not conducive to efficient operation. Modernization of mines will require skilled and dependable workers and to induce such men to become interested in coal mining in Alaska, modern facilities, such as modern houses, schools, churches, and recreation halls, must be provided near the mines. Alaska is still a frontier and improvements in living standards must accompany development of mines in order that a stable community may be established. Coal is known to have been mined in Alaska at Port Graham, on Cook Inlet, by the Russians in 1854. During the Klondike gold rush in 1898, coal was mined in scattered areas for use on steamers plying the Yukon River and its navigable tributaries. After the Alaska Railroad was built, two important fields were developed—the Nenana and the Matanuska (Fig 1). The Bering River field was explored by the Guggenheim interests at one time, and a railroad to the area was constructed. However, the field was abandoned because of litigation and reported unsatisfactory physical condition of the beds. There are areas that today are unimportant but which may be developed in the future. Possibly the most important of these is the deposits of subbituminous coal adjacent to Cook Inlet, near Homer, on the Kenai Peninsula (see Fig 1). It is expected that the demand for coal in the Kenai Peninsula will increase because the population in that area has increased as a result of agricultural development. Numerous outcrops of coal are exposed on the peninsula and development of a mine there would obviate dependence on supplies of coal from the Matanuska field. Coal from this field must be shipped approximately 160 miles by rail to Seward and transported thence by boat about 180 miles to Homer. There has been no regular boat service from Anchorage to Homer. There are deposits of coal in the Arctic region but a description of them will not be given in this paper, which will describe the more important deposits in the Alaska Railroad belt. The results of a reconnaissance in the Arctic have been described.&apos; Description of Fields MATANUSKA FIELD The Matanuska field lies in the valleys of the Matanuska River and its tributaries and their separating ridges. This field has been interpreted by geologists as a sunken fault block that separated the Talkeetna Range of mountains from the Chugach Range. The Matanuska River flows west along the south side of the field and the Glenn highway, which extends from Anchorage via Palmer to the Richardson highway near Glen Allan, follows the Matanuska Valley. This road to the coal field can be traveled the year around. Although the snow-

Jan 1, 1950

By Glenn S. Wyman

CHUQUICAMATA, located in northern Chile in the Province of Antofagasta, is on the western slope of the Andes at an elevation of 9500 ft. Because of its position on the eastern edge of the Atacama Desert, the climate is extremely arid with practically no precipitation, either rain or snow. All primary blasting in the open-pit mine at Chuquicamata is done by the churn drill, blasthole method. Since 1915; when the first tonnages of importance were removed from the open pit, there have been many changes in the blasting practice, but no clear-cut rules of method and procedure have been devised for application to the mine as a whole. One general fact stands out: both the ore and waste rock at Chuquicamata are difficult to break satisfactorily for the most efficient operation of power shovels. Numerous experiments have been made in an effort to improve the breakage and thereby increase the shovel efficiency. Holes of different diameter have been drilled, the length of toe and spacing of holes have been varied, and several types of explosives have been used. Early blasting was done by the tunnel method. The banks were high, generally 30 m, requiring the use of large charges of black powder, detonated by electric blasting caps: Large tonnages were broken at comparatively low cost, but the method left such a large proportion of oversize material for secondary blasting that satisfactory shovel operation was practically impossible: Railroad-type steam and electric shovels then in service proved unequal to the task of efficiently handling the large proportion of oversize material produced. The clean-up of high banks proved to be dangerous and expensive as large quantities of explosive were consumed in dressing these banks, and from time to time the shovels were damaged by rock slides. As early as 1923 the high benches were divided, and a standard height of 12 m was selected for the development of new benches. The recently acquired Bucyrus-Erie 550-B shovel, with its greater radius of operation compared to the Bucyrus-Erie 320-B formerly used for bench development, allowed the bench height to be increased to 16 m. Churn drill, blasthole shooting proved to be successful, and tunnel blasts were limited to certain locations where development existed or natural ground conditions made the method more attractive than the use of churn-drill holes. Liquid oxygen explosive and black powder were used along with dynamite of various grades in blasthole loading up to early 1937. Liquid oxygen and black powder were discontinued because they were more difficult to handle due to their sensitivity to fire or sparks in the extremely dry climate. At present ammonium nitrate dynamite is favored because of its superior handling qualities and its adaptability to the dry condition found in 90 pct of the mine. In wet holes, which are found only in the lowest bench of the pit and account for the remaining 10 pct of the ground to be broken, Nitramon in 8x24-in. cans, or ammonium nitrate dynamite packed in 8x24-in. paper cartridges, is being used. This latter explosive, which is protected by a special antiwetting agent that makes the cartridges resistant to water for about 24 hr, currently is considered the best available for the work and is preferred over Nitramon. Early churn drill hole shots detonated&apos; by electric blasting caps, one in each hole, gave trouble because of misfires caused by the improper balance of resistance in the electrical circuits. Primarily, it was of vital importance to effect an absolute balance of resistance in these circuits, the undertaking and completion of which invariably caused delays in the shooting schedule. Misfires resulting from the improper balance of electrical circuits, or from any other cause, were extremely hazardous, since holes had to be unloaded or fired by the insertion of another detonator. The advent of cordeau, later followed by primacord, corrected this particular difficulty and therefore reduced the possibility of missed holes. After much experimentation, the blasting practice evolved into single row, multihole shots, with the holes spaced 4.5 to 5 m center to center in a row 7.5 to 8 m back from the toe. Such shots were fired from either end .by electric blasting caps attached to the main trunk lines of cordeau or primacord. The detonating speed of cordeau or primacord gave the practical effect of firing all holes instantaneously. Double row and multirow blasts, fired instantaneously with cordeau or primacord, proved to be unsatisfactory in the type of rock found at Chuquica-

Jan 1, 1952

By Sidney Fischler

Large striation-free single crystals of bismuth have been grown from the melt by rapid freezing. Zone-refined bismuth, together with doping impurities if desired, is placed in a shallow flat-bottomed graphite boat and melted in air with a propane hand torch. The torch is then withdrawn in a manner which causes the melt to freeze direction-ally. Crystallization, which resuires only a few minutes, usually results in the formation of a single crystal even when a seed crystal is not used. Crys -tals of any desired orientation may be grown by using oriented seeds. Undoped crystals grown by this method have residual resistivity ratios greater than 200. THE growth of large single crystals of bismuth by either the Czochralski or horizontal zoning technique is not entirely satisfactory. Specifically, difficulties are encountered in producing single crystals of the required dimensions in all desired orientations, and striations caused by low-angle polysynthetic twins are frequently present in the crystals. In addition, both methods are time-consuming and require special apparatus of some complexity. A simpler method has now been developed for growing large striation-free bismuth single crystals of desired orientation in a short time. Fig. 1 shows a typical setup consisting of a rectangular graphite boat which contains zone-refined bismuth, a 1/8-in.-thick flat quartz plate which covers the entire inner bottom of the boat, and three additional quartz plates about 1/4 in. thick which are used to separate the bismuth from the graphite everywhere except at a small area at the left of the boat. The graphite boat is 1 in. high, and its sides and bottom are about 1/8 in. thick. The quartz plates should be smooth and clean. The graphite boat is heated from the right with a propane torch, as shown in Fig. 1, until the bismuth is completely melted. The melt has the shape of a triangle with a narrow neck at the apex farthest from the torch. The melt is frozen direc-tionally by gradually moving the torch toward the right, away from the boat. The bismuth in contact with the graphite, at the left end of the neck, freezes first. The freezing interface then moves down the neck into the main bulk of material, where it develops a convex shape ideal for the continuation of single-crystal growth. The interface continues to move through the melt until the entire bulk is solid. The entire procedure may be completed, in air, in a matter of minutes. The technique described almost always yields a single crystal whose basal plane is nearly perpendi,cular to the bottom of the graphite boat. In earlier experiments, in which the bottom of the melt was in direct contact with the graphite boat, single crystals were grown with basal planes parallel, perpendicular, or at some intermediate angle to the bottom of the boat. At times the orientation of the bulk of the material differed from the orientation of the material in the narrow neck. In these cases, a nucleation site initiated the growth of a differently oriented crystal, and the thermal conditions favored the new orientation over the initial one. The thermal conditions depend on a number of factors, including the heating technique, the placement, shape, and thickness of the quartz plates, the thickness of the walls and bottom of the graphite boat, and the quantity of bulk bismuth employed. All of these factors, plus the initial orientation and the presence and effectiveness of nucleation sites, will determine the orientation of the final large single-crystal slab. When a crystal of specific orientation is desired, an oriented section of a rapid-freeze crystal is shaped by spark cutting and grinding for use as a seed. To grow a doped crystal, the desired impurity is placed in the graphite boat together with the bismuth chunks and seed. Crystals doped with mercury, cadmium, lead, and selenium have been grown. The rate of freezing is so great that the distribution coefficient of any impurity approximates unity. On a gross scale, therefore, impurities should be more homogeneously distributed in rapid-freeze crystals than in Czochralski or zoned crystals. Because of the possibility of constitutional supercooling, however, it is quite possible that impurities are not homogeneously distributed on a microscopic scale in the rapid-freeze crystals. Generally the single crystal slabs which have been prepared are initially 5 to 7 mm thick. Thicker crystals may be obtained by using one of these slabs as a seed. The slab is placed in a graphite boat resting on a large aluminum block, either air- or

Jan 1, 1964

The movement for Federal endowment of engineering and industrial research was revived early in the present Congress by the introduction of two bills into the Senate. The first bill, introduced by Senator Smith of Georgia, provides for the establishment of an engineering experiment station in each state under the direction of and in connection with some university, engineering school, or land-grant college. The state legislature of each state is authorized to designate the institution best equipped to conduct the work but wherever the land-grant college has facilities approximately equal to those of the other institutions of the state, it shall be designated. All designations of institutions are made subject to the approval of the Secretary of Commerce. The Secretary of Commerce is authorized to secure practical uniformity of methods, and efficiency and economical expenditure of funds by indicating such lines of work as shall seem to him important from a military; naval, industrial or national standpoint, but the responsibility for the initiation and conduct of research shall rest with the individual engineering experiment stations. Uniform standards of research are further provided for by making each experiment station a depository of the Department of Commerce and the Bureau of Standards, and the work of each station shall conform to the standards established by the U. S. Government through said bureau, which may whenever practicable utilize the facilities of the station for the prosecution of research.

Jan 8, 1919

By C. B. Alcock, L. I. Cheng

By the use of radiochemical methods for the study of the gas-liquid equilibria at low temperature, and for the determination of the sulfur contents of metal beads which had been equilibrated with H2S/H2 mixtures of known sulfur potential, it has been possible to obtain the liquid solubility and the free energy of solution of sulfur in liquid tin and lead at temperatures between 500°and 680°C. THE gas-liquid equilibrium method has proved in the past to be most successful in the determination of the thermodynamic behavior of dilute solutions of sulfur in liquid metals.1,2 One of the basic requirements for success with this method is that the volatility of both the metal and its lowest sulfide should be small, otherwise sulfide will be deposited at the cool end of the furnace, where it may react with the outgoing gases to form either sulfur-rich lowest sulfide or higher sulfides. The resultant value of the apparent equilibrium constant will then be lower than the correct one. This argument applies even at sulfur potentials below that in equilibrium with a separate condensed phase of the lowest sulfide at the reaction temperature, T. The mass of sulfide which is deposited at the cold end of the furnace, and hence the extent to which further reaction occurs with the outgoing gases, depends on the time taken for equilibrium to be reached between metal and gas. Since this will depend principally on the bulk of the metal phase which is used, one should clearly attempt to uie as small metal samples as possible. These considerations are important in the study of dilute solutions of sulfur dissolved in liquid tin and lead which both have moderately high vapor pressures as metals and form volatile sulfides. The limit on the size of the metal samples which may be used is set chiefly by the difficulties of analysis for very small amounts of sulfur. The oxygen or carbon dioxide combustion method, followed by iodimetric determination of the sulfur dioxide which is formed,has been found to be successful for the determination of small amounts of sulfur in copper, iron, cobalt and nickel.4 This method was unsatisfactory for sulfur dissolved in tin and lead, mainly because the sulfur dioxide was to some extent absorbed by the copious tin or lead oxide deposits which were formed on the walls of the combustion tube. Furthermore some of the sulfur was found to segregate on the surface of the beads as flaky sulfide crystals which would easily be lost in the transfer of a bead from a boat in the gas equilibration apparatus to one in the combustion apparatus. Oxidation in aqueous media to sulfate ion followed by precipitation as barium sulfate was, therefore, adopted as the analytical procedure. The gas-metal equilibrium experiments were all carried out with radioactive sulfur and thus the analysis involved the counting of barium radiosulfate. Furthermore the use of the radioisotope meant that the approach to the gas-metal equilibrium could be followed continuously by gas counting.&apos; The metal beads were held separately in glass crucibles during equilibration and were transferred from the furnace to the beaker for dissolution in nitric acid still in the crucibles, and thus the possibility of sulfur loss by detachment of the sulfide segregates was eliminated. The temperature range of this investigation was 500° to 680°C. EXPERIMENTAL APPARATUS AND METHOD The apparatus consisted of two furnaces placed in series in a gas recirculation system, Fig. 1. One furnace F1, which was vertical was used to heat the alumina crucible, A, holding six metal beads in separate glass crucibles. The beads weighed between 300 and 700 mg each. The crucible assembly was introduced and removed from the furnace mechanically under a stream of oxygen-free argon. The other furnace, F2, was horizontal and was used to heat a cobalt Co9S8 mixture, held in an alumina boat, and made with radiosulfur containing about 1/2 millicurie per g of sulphur. This mixture, which was finely powdered, was used as a source of known H2S/H2 mixtures6 for a given furnace temperature. The recirculation system also contained a gas re-circulation pump (P), an end window Geiger-Miiller counter (N)—placed downstream of F1 so as to monitor the H2S pressure in the gas leaving this furnace— a sample volume for chemical analysis of the gas phase (G), gas drying tubes (D), filling taps and other standard ancillary equipment. The gas sampling volume was principally used in the cali-

Jan 1, 1962