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By AIME AIME

ANOTHER meeting of the Institute has passed into history and it fully sustained the reputation of the Institute as a live organization of the men, and nowadays the women, concerned with the mineral . ? industries and those working therein. The field meetings of, the 1nstitute are cine of its pleasantest and most valuable features. As Dr. Raymond used to say, ((It is good for a mining engineer once in a while to see a mine" and these meetings bring back to the mines those who, through progress into higher executive positions, have come to visit them infrequently, at the same time that they give opportunity for the .younger men to see regions and mines aside from those in which they are then working. This, and the formal and' informal interchange of information resulting,, is of the greatest professional value to all, while the meeting of old and making of new friends causes each meeting to be notable in turn. This year's meeting at Salt Lake City was one of the largest that has been held outside of. New York City; 470 members and guests registered while many more were present at the sessions without attending to this formality. '

Jan 1, 1925

America.-Foreman, J. T. Ashland.-Sturdevant, J. C. AItalIa. Henderson, R. M. Bessemer.-Ball, E. M. Maschmeyer, W. L. McKenzie, W. C., Jr. Birmingham.-Abbott, C. E. Aldrich, T. H. Aldrich, T. H., Jr Aufderhaar, H. C. Berg, H. A. Blair, C. S. Bush, M. W. Clark, G. H.

Jan 1, 1923

Jan 1, 1968

Aamot, Olav Crone, (M'29) Chem. Engr., Elektro¬kcmisk, Raadhusgt. 32, Oslo. Norway. . Aase. Glenn D., (J'40) Jr. Met., American Smelt. & Refin. Co., Selby, Calif. Abadie, Henry G., (J'43) Petr. Engrg., Avenal, Calif. Abadilla, Quirico A., (M'38) Min. Engr., Dir., Bureau of Mines, Manila, P. I. Abbott, Agatin Townsend, (J'40) Jr. Engr., Phelps Dodge Corp., Morenci Branch. Morenci. Ariz. Abbott, Clarence E., (M'04) Civ. Engr. V.P., Tennessee Coal Iron & R.R. Co., 1542. Brown Mars Bldg. Birmingham. Ala. Abel, Raymond L.. (M'23) Chem. Engr., Com¬manding Officer, Pittsburgh Chem. Warfare urement Dist., American Bank Bldg.. Pittsburgh. Pa. Abel, Walter D., (M'32) Min. Engr., Supvr. Engr.. R.F.C., Los Angeles, Calif. Abele, Louis T., (M'31) 423 3rd Ave., Decatur. Ga. Abell, Leo M.: (M'31) Mill Supt. & Met.. Em¬ peror Gold Min. Co., Ltd., Vatukoula. Fiji Islands. Abernathy, G. E., (M'38) Geol., Kansas State Geol. Survey, Pittsburg, Fan. Abernathy, Jack H., (M'40) Cons. Petr. Engr., First National Bldg.. Oklahoma City, Okla.

Jan 1, 1943

By Alfred Boyles

"Heredity"1 in cast iron has been a subject of much discussion. Numerous experimenters have found that the properties of gray iron may vary greatly without corresponding variations in composition or thermal treatment. The main structural difference in such cases is in the size and distribution of the graphite flakes and the explanations given usually attribute this difference to some chemical element not shown by ordinary analysis or to some physical factor, such as inclusions in the melt. It cannot be said that any of these theories is completely satisfactory. Much remains to be learned about cast iron and for this reason a study of the material was undertaken as a part of the program of fundamental research at Battelle Memorial Institute. A review of the literature indicated that very little systematic work had been done on the mechanism of freezing in cast iron. The experiments described below are largely concerned with that problem. Normal and "Dendritic" Cast Iron As examples of variation in type of structure, photomicrographs of two test bars of cast iron are shown (Figs. 1 and 2) and the composition and properties are given in Table 1. Both are common, unalloyed gray iron. Test bar No. 13 has normal graphite flakes while No. 49 has very fine, lacy graphite arranged in a dendritic pattern, the matrix in both being pearlite. These two types of cast iron are well known to foundry-men and need no detailed description.

Jan 1, 1937

By Hansheinz Portisch, Harold Margolin

A surface expansion was found to occur sometime after etching in Ti-A1 alloys containing 9.5 to 12.5 wt pct Al. The structure formed, grew, and disappeared with tzrrze. The surface expansion was followed by microscope observations and interferometric and lattice parameter measurements. Activation energy measurements for the growth of the "expansion structure" and chemical analysis indicated that the phenomenon occurred as a result of hydrogen pzckup during etching. It is proposed that hydrogen initially enters octahedral sites of Ti3Al coherent with a Ti and later shifts to the tetrahedral sites. It is postulated that expansion occurs when hydrogen enters the tetrahedral sites. The expansion structure disappeared, it is proposed, because of diffusion of hydrogen from the surface into the body of the alloy and because of loss of coherency of Ti3Al. In examining Ti-A1 alloys, Ence and arolinl observed markings on the surface of specimens. These markings did not appear after electropolishing only, but rather appeared only after etching. The markings appeared to grow as a function of time after etching and later seemed to disappear. Although the markings had some similarity to precipitates, showing, frequently, a Widmanstatten type of arrangement, the observation that other microstructural markings continue to be seen within the new structure suggested that it was actually not a precipitate. The source of this structure was unknown and an attempt was made in the present investigation to develop some understanding of its nature. It has been labeled expansion structure. 1) EXPERIMENTAL PROCEDURE A) Alloys. Alloys in the range 6.5 to 14.5 wt. pct A1 were studied. Bars, 4 in. sq, forged from Bureau of Mines (73 Bhn) titanium consumably arc-melted 4-lb ingots,' as well as 50-g arc-melted buttons of desired aluminum contents were used. The 50-g buttons also used Bureau of Mines titanium (73 Bhn) and aluminum of 99.99 pct purity. Buttons containing up to 8.5 wt pct A1 were hot-rolled from a furnace at 900' . Those with higher aluminum contents were hand-forged on a titanium anvil, and heated with an oxygen-hydrogen torch in the region of 1200" to 1300°c. Frequent reheating kept the samples at the desired temperature range. After a reduction of about 30 pct, the samples were water-quenched. To eliminate any contamination picked up either during hot rolling or forging, at least 1 mil of the surface of the sample was taken off. B) Heat Treatment. Prior to heat treatment all alloys were vacuum-annealed to remove hydrogen. Samples were annealed at 900' until a vacuum of 10'5 mm Hg was established. After this treatment, the samples were wrapped in molybdenum sheet and heat-treated in argon-filled quartz capsules, which were broken under water or iced brine at the conclusion of the heat treatment. All heat treatments under dynamic vacuum were performed in a rapid-quench furnace. This consisted of a molybdenum-lined quartz tube attached to a vacuum system and through a stopcock to a beaker of water. At the completion of the heat treatment the vacuum stopcock was closed, the furnace shut off, and dropped below the quartz tube. Then immediately the inlet stopcock was opened and water admitted until the tube was filled. The steam formed was allowed to escape through the inlet stopcock. This method was used in heat treatments up to llOO°C. C) Metallography. Specimens for metallographic examination were ground, then electropolished using a Disa Electropol machine with a perchloric acid electr01te. Specimens were etched with R-etch.3 A standard etching time of 3 min was used, with the specimen being agitated during immersion. D) Sample Preparation for X-ray Analysis. Samples 0.2 to 0.5 mm in diam were produced from heat-treated rods which were turned to 2 mm diam on a lathe and then rotation-etched. An etchant consisting of 1 pt HF, 7 pt HN03, and 12 pt HzO was satisfactory. With the rod rotating in a vertical position at 50 to 100 rpm, a needle with uniform dimensions could be obtained. Care had to be taken to insure that the rod was at the center of rotation, otherwise cavitation developed. If a small unevenness developed, it was possible to grind it off with a fine emery paper. E) X-ray Diffraction. All X-ray work was done on a North American Phillips X-ray diffraction apparatus and a Jarrell Ash microfocus X-ray unit. The Phillips unit was used with a copper target and nickel filter. The Jarrell Ash unit was fitted with a cobalt target and iron filter. For the Phillips unit 35 kva and 20 ma were used, whereas for the microfocus unit with the 100-p fixed-focus gun 40 kv and 1.5 ma were used. It was found that alignment of cameras on the Jarre11 Ash unit was very critical. The X-ray beam contains an intense area which is not the beam center. The cameras were aligned with the intense region by monitoring the beam coming out of the camera with a Geiger counter. Adjustments were made until a maximum intensity was obtained. A Phillips diffractometer with a Brown chart recorder furnished some of the lattice parameter data. The divergence slit up to 80 deg 28 was I deg; above 80 deg, a 4-deg opening was used, while the scatter slit was1 deg and the receiving slit had a 0.003-in. opening. The general scanning rate was 1 deg per min, while peaks of special interest were rescanned at -j deg per min. For elevated-temperature X-ray diffraction a Uni-cam High Temperature Camera with a film diameter

Jan 1, 1969

By B. A. Wilcox, A. H. Clauer

The steady-state creep of TD Nickel NL + 2 001 pct TltOz) has been studied orer the telirperatve range 325&apos; to 1100O and the stress range 15,000 to 36,000 psi. At high temperatures (aboue 0.5 T& gran-boundary slzding is the )nost znportant )node of creep deformation, and the steady-state creep rate, is, can be related to stress and temperature by: where Q = 190 kcal pev mole and n has an unusually high value of 40. A creep mechanism based on cross slip of dislocations around The O2 particles can satisfactovily explain the low-temperature (T < 0.5 T,) cveep behavior, and the follo wing relation is applicable: Q, (a) is found to decrease from 57 to 46 kcal per mole as the stress is increased from 32,000 to 36,000 psi. THERE have been a variety of theories proposed to explain the influence of dispersed second-phase particles on the yield strength and flow stress of metals, and these have been reviewed recently by Kelly and icholson.&apos; However, only several attempts2"4 have been made to develop mechanistic treatments which characterize the creep behavior of dispersion-strengthened metals, and to date these have not been fully evaluated experimentally. weertman2 and Ansell and weertman3 proposed a quantitative creep theory for coarse-grairzed dispersion-strengthened metals, based on the concept that the rate-controlling process for steady-state creep was the climb of dislocations over second-phase particles, as suggested by choeck. The theory predicted that the steady-state creep rate, <,, was proportional to the applied stress, a, for low stresses and that is a4 o for high stresses. The activation energy for creep, Q,, was equivalent to that for self-diffusion, Qs.d., in the matrix. Some limited experimental evidence in support of this theory was obtained on a recrystallized Al-Alz03 S.A.P.-type alloy by Ansell and Lenel.6 Ansell and weertman3 also developed a semiquanti-tative theory for high-temperature creep of lineg-rained dispersion-strengthened metals in order to explain their results on an extruded S:A.P.-type alloy, which had a fine-grained fibrous structure. They suggested that the rate of dislocation generation from grain boundaries was the rate-controlling process, and fitted their results to the equation: where Q, was found to be 150 kcal per mole, i.e., QC- 4Q,.d. in aluminum. Similar high activation energies for creep7-&apos;&apos; and tensile deformation" of dispersion-strengthened alloys have been observed by other investigators for S.A.P.,&apos;" indium-glass bead omosites, and Ni + A1203 alls.&apos; There is no general agreement regarding the mechanisms involved in the creep of dispersion-strengthened metals, and this is due in part to the lack of detailed studies relating the structures of crept specimens to the mechanical behavior. The present investigation on thoriated nickel was undertaken with the aim of studying the structural changes which occur during creep of a dispersion-strengthened alloy and rationalizing the observed mechanical behavior in terms of the creep structures. EXPERIMENTAL METHODS The material used in this investigation was 1/2-in.-diam TD Nickel bar, which contained 2.3 vol pct Tho,. Obtained from E. I. duPont de Nemours & Co., Inc. The final fabrication treatment by DuPont consisted of -95 pct reduction by swaging followed by a 1-hr anneal at 1000°C. Transmission and replica electron microscopy revealed that the material had a fine-grained fibered structure with an average transverse grain size of -1 p and a longitudinal grain size of 10 to 15 p. Selected-area diffraction indicated that the fiber axis was parallel to (OOl), in agreement with the results of Inman eta1." All creep specimens were vacuum-annealed at 1300°C for 3 hr prior to testing. Transmission electron microscopy showed that the only structural change due to annealing was a slight decrease in dislocation density, confirming the reported high degree of structural stability.13 Furthermore, recrys-tallization or grain growth during creep was never observed. The structure typical of uncrept material (after the 1300 C, 3-hr anneal) is shown in Fig. 1. The grain boundaries are predominantly high angle and. although some areas show a tangled cell structure, the grain interiors are relatively dislocation-free. Individual dislocations are strongly pinned by the Tho2 particles; i.e., very rarely did dislocations move within a thin foil. The grey "halos" around some of the larger particles which protrude out of the foil surface arise from contamination in the electron microscoge. The Tho, particle size ranged from -100 to IOOOA, and the distribution is shown in Fig. 2. The technique used to obtain the data in Fig. 2 consisted of dissolving the nickel matrix in acid, collecting the Tho2 particles on cellulose acetate, and measuring about 1000 particle diameters in the electron microscope. Similar results were obtained by measuring about 600 particles in thin foils, an; the average particle size was found to be 2r, = 370A. Using the data in Fig. 2 (annealed structure), the mean planar center-to-center particle

Jan 1, 1967

By Donald T. Hawkins, Ralph Hultgren

The lorsion effusion method has been used to measure the vapor pressure of lead over pure lead and eight Pb-Sn alloys ranging from 9.1 to 87.9 at. pct Pb in the temperature range 950° to 1125°K. The resulting ?H°y 298 = 46,630 ± 200 cal per g-atow for pure lead is in excellent agreement with previous work. Deriued ?Gph values for the alloys were smoothed and correlated with previous heat of formation data to obtain entropies for the liquid alloys. values for the tin component were calculated by Gibbs-Duhem integmtion. Surfhce depletion of these liquid alloys was found to be absent, as shown by the fact that the measured pressure remained substantially constant with time in contrast with previous measurements on iron-rich solid Fe-Mn alloys. The activities and Gibbs (free) energies of the liquid Pb-Sn system are not well known although the heats of formation have been accurately determined by K1eppa.l Vapor pressure determinations of the activity of lead by Voronin and Evseev2 disagree with those of Predel.3 Elliott and chipman4 measured activities by the electromotive -force technique in the ternary Pb-Sn-Cd system. From the resulting values of the activity of cadmium, they were able to calculate the activities of lead and tin for the binary Pb-Sn system at the compositions xsn = 0.33 and 0.67. Atarashiya, Uta, Shimoji, and Niwa5 determined the activity of tin from Hz-H2O equilibrium measurements, but their results are subject to question due to uncertainty in the final state of their samples. To resolve these differences, measurements of the activity of lead were made using a torsion-effusion apparatus. Among other advantages, this apparatus rapidly detects errors caused by surface depletion, which tends to reduce the apparent vapor pressure of the more volatile component. If the measured torques remain nearly constant with time, only bulk depletion is occurring. This contrasts with the rapid decrease in torque found in solid Fe-Mn alloys at high temperatures and low manganese contents.6 EXPERIMENTAL 99.999+ pct Pb was obtained from the American Smelting and Refining Co. and 99.999 pct Sn from Vulcan Detinning Co. Eight Pb-Sn alloys were prepared by melting together weighed amounts of the pure metals in evacuated Pyrex tubes and quenching the melts in water. Loss of weight in preparation was less than 0.02 pct in all cases. The effusion cells were made from high-purity, nonporous graphite. Graphite was chosen because of its resistance to reaction with lead and tin and because it can be accurately machined with precision orifices. Two cells were used with overall dimensions as shown in Fig. 1. The orifices of cells 1 and 2 were, respectively, 1 and 1.5 mm in diam.

Jan 1, 1968

?AARONSON, ALFRED E , Vice-Pres , Mid-Co Petroleum Co, Mid-Co Bldg, Tulsa, Okla &apos;18 ABADIE, EMILE R, Min Engr Address wanted &apos;76 || ABADILLA, QUIRICO A, Geol Dept, Cia Mexicana de Petroleo "El Aguila," S A, Edifico "El Aguila," Esq, Calles Aurora y Comercio, Tampico, Tamaulipas, Mexico &apos;20 ABBOTT, CLARENCE E, Gen&apos;l Supt, Iron Mines Div, Tenn Coal, Iron & R R Co, 1405 Minnesota Ave, Bessemer, Ala &apos;04 ABBOTT, LOUIS BENJAMIN, Chief Engr, Consolidation Coal Co, Elkhorn Div, Jenkins, Ky &apos;15 ABBOTT, ROBERT R Met Engr, Peerless Motor Car Co, Cleveland, Ohio &apos;12 ABE, ASAKA Fujita Min Co, Dojima, Osaka, Japan &apos;14 ABEEL, GEORGE H, JR, Cons Min Engr, 635 St Paul Ave, Los Angeles, Cal &apos;18 ABEL, ALBERT E, Asst to Pres, Valley Mould & Iron Corpn, Sharpsville, Pea &apos;19 ABEL, W D, Min Engr Forum Bldg, Sacramento, Cal &apos;19 ABELL, O J, Pres & Treas, Abell-Howe Co, 332 South Michigan Ave, Chicago, Ill &apos;19 ABERNATHY G E Willow Springs, Mo &apos;20

Jan 1, 1923

By A. R. Troiano, A. B. Greninger

There is need for an adequate working hypothesis that would describe at least qualitatively the crystallographic mechanism for the transformation from austenite to martensite in steel. A general theory would not only be of great assistance in the correlation of existing data and in planning future experimentation, but would also form the basis for an eventual explanation of the property changes that accompany lattice transformations. The mechanism of martensite formation described in the following pages is the outcome of two new experimental determinations: (1) the accurate evaluation of the lattice relationship between austenite and individual crystals of martensite, and (2) the measurement and analysis of the change in position that a volume of austenite undergoes when it transforms into a crystal of martensite. For this latter study, the stereographic analysis of shear was employed; this method greatly simplifies the solution of lattice-shear problems, and some space is devoted to its description. Lattiice Relationships and Habits PREVIOUS WORK The lattice relationships between austenite and martensite in carbon steels have been studied by Kurdjurnow-and Sachs,&apos; and by Wassermann;2 and in iron-nickel alloys by Nishiyama.3 Wassermann,2 and Mehl and Derge.4 The studies by Kurdjumow and Sachs and by Wassermann were made on single (austenite) crystals of quenched 1.4 pct carbon steel, and the results of these two independent studies agree closely with the postulated lattice relationships: (III)r||(0ll)a [101]r [111]a [1] Both Nishiyama and Wassermann worked with an alloy of iron plus 30 pct nickel (all austenite at room temperature), and "martensitic alpha" was formed by cooling in liquid air. They agree that the lattice relationships for these conditions are: (lll)r||(011)a [112]r[011]a [2] Mehl and Derge worked with a range of iron-nickel alloy compositions centered about 30 pct nickel. They concluded that when the transformation takes place near room temperature or above, relationship [l] holds, but if the same specimen is made to transform at -195°C, relationship [2] obtains;intermediate temperatures evidently produced a combination of these two relationships. Sachs and co-worker~,1,7 as well as Nishiyama,3 have proposed transformation mechanisms for martensite formation.* Both of the mechanisms picture the transformation as initiated by shearing movements along the octahedral [Ill] plane of austenite, followed by suitable expansion and contraction to complete the transformation. These mechanisms have been described and commented upon by Mehl, Barrett, and Smith,8 Mehl and Derge,4 Wassermann,2 and others, and generally accepted by investigators in this field. Greninger and Troiano9 have recently published the results of a detailed study of the shapes and orientation habits of martensite crystals in plain-carbon and nickel steels. They found that, contrary to previous assumptions, martensite plates form not along the octahedral planes but along austenite planes of very high indices; these results were in line with those of previous studies of martensite reactions in nonferrous alloy systems by Greninger and coworkers.6,10 Greninger and Troiano concluded that the transformation theories of Kurdjumow and Sachsl and of Nishiyama3 are untenable. Phragmén11 reached a similar conclusion. EXPERlMENTAL METHODS The object of this particular study was the evaluation of (1) lattice relationship between austenite and martensite, and (2) relationship between the martensite lattice and the martensite plate. The technique used for solving these problems was similar to that used by Greninger6 on martensitic structures in 8 copper-aluminum alloys. Briefly, the method consists of grinding and polishing a specimen on a surface parallel to an individual martensite plate and thus exposing a single martensite crystal for an area of l to 3 mm. The orientation of this exposed martensite crystal is then determined by means of a back-reflection Laue X ray pattern.12 Another pattern is obtained from the matrix crystal, and a stereo-graphic plot of data from these two patterns then provides the solution to the problem. It is necessary to repeat this process on several crystals in order to determine whether or not the solution is unique.

Jan 1, 1950

By Larry Kaufman, Joan B. Berkowitz-Mattuck, Edward V. Claugherty

The oxidation characteristics of hafnium and zirconiunr diboride were measured between 1200 and 2200&apos;K by a thermal- conductivity method which continuously ttzeasures the rate of reaction of oxygen with the diboride and by a metallographic air oridation method zuhich provides a measure of the total arr7ount of bovide conuerted to oxide for a given time interval. The oxidized specimens obtained from the tl~eritaal-coi~ductitrity method were also examined by quantitatire metal-lographic procedares. The significant results obtained in this investigation reveal that metal-rich compositions of lzafi~iutil diboride proride the most oxidation-resistatzt material up to 2000°K; hafnium diboride is tmre oxidation- resistant titan zirconium diboride at all tempevatures examined; the morphology of the oxide formed on H/B2 and Llie temperature coefficient qi. the oxidation rate constants change at the temperature of ttze monoclinic to tetragonal phase transition] in HfO2; the oxidation of neither HfB2 nor ZrB2, results iN catastrophic Jazlure at lorc. oxygen pressures; and pvefevetztial gvaLti boundary oxidation was not obsevued for either HfBi or ZYB, A comprehensive study of the high-temperature characteristics of refractory transition-metal di-borides is currently in progress. This program has included investigation of the physical, thermal, and thermodynamic properties of TiB2, ZrB2, HfB2, NbB2, and TaB2. In addition, aspects of the synthesis and fabrication of such materials have been studied. In view of the diverse nature of this research, a number of other laboratories have actively participated and contributed specific capabilities for analysis and characterization of these materials. As a consequence, an extensive description of the relevant properties of these compounds has emerged which is central in evaluating their high-temperature (1200" to 2500°K) performance. To date, information on thermodynamic stability, specific heat, and vaporization characteristics,1 hot hardness and electrical resistivity,1, 3 therma1 expansion:&apos;4 and thermal conductivity 1, 5 has been presented. This information has been generated on materials of the highest purity (98.5 to 99.9 wt pct Me + B) and density currently available. Samples fabricated by zone melting6 and high-pressure hot pressing"3&apos;7 techniques have been used to generate suitable specimens for all of the aforementioned studies. dation characteristics of the most oxidation-resistant of these materials, hafnium and zirconium diboride, is presented and a description of the synthesis and the experimental procedures used to prepare and characterize specimens is given. The high-temperature range under consideration (1200" to 2200°K) and the known dependence of oxidation characteristics on sample chemistry, density, and oxidation conditions required a close coupling of the synthesis, fabrication, and evaluation procedures.8 This was accomplished by continual surveillance of chemical composition of starting materials before and after specimen fabrication and by evaluation of density, phase constitution, and microstructural features prior to and after oxidation exposure. I) PROCUREMENT AND CHARACTERIZATION OF STARTING MATERIALS In view of the current state of the art in fabricating refractory boride materials, the methods used in preparing samples for the present study are given in detail as follows: starting materials were purchased in high-purity powder form and fabricated by high-pressure hot pressing into 0.40 by 1.00 in. bars from which oxidation specimens were obtained. The hafnium diboride used in this study was purchased from Wah Chang Corp.; the zirconium diboride from U.S. Borax and Chemical Co. These powders were routinely characterized by quantitative chemical analyses for metal, boron, carbon, oxygen, nitrogen, and iron, by qualitative emission-spectrographic analysis for trace impurities, by X-ray procedures for extraneous phase identification, and by powder densitometry for comparison with X-ray (theoretical) density. Hafnium and zirconium metal and elemental boron were also purchased as high-purity powders and characterized for impurities by emission-spectrographic analyses. The hafnium diboride was procured in three shipments which were designated as HfBl.g7(1), HfB1.88(2A), and HfB2.12(2). The indicated stoichiometry is based on the atomic ratio of total boron to total hafnium; the number in parentheses identifies the shipment number. Shipment 1 was 5 1b, shipment 2A, 1 1b, and shipment 2, 8 1b. The zirconium diboride was procured as a 20-1b shipment and designated as ZrB1.89(1). A small quantity of purified zirconium diboride was also supplied and designated ZrB1, 9(P). The averaged results for chemical analyses which were generally performed according to the procedures set forth in the compilation by KrieGe9 are presented in Table I. Qualitative spectrographic analyses indicated that Ca, Cr, Ti, Si, Zr (in H~B~), and A1 were present at levels between 0.01 and 0.10 wt pct. Other metallic elements were found to be less than 0.01 wt pct. Since it is virtually impossible to purchase these materials in the desired quantities (5 to 20 lb) as single-phase compounds it is necessary to obtain

Jan 1, 1968

ALABAMA Alabama City.-Eddy, L. Altoona.-Cain, J. America.-Foreman, J. T. Anniston.-Foster, R. N. Ashland.-Sturdevant, J. C. Bessemer.-Ball, E. M. Calhoun, F. W. Maschmeyer, W. L. McKenzle, W. C., Jr. Birmingham.-Abbott, C. E. Aldrich, T. H. Aldrich, T. H, Jr. Ashton, J J.

Jan 1, 1923

By C. M. Wayman, J. F. Breedis

The crystallography of the martensitic transformation in the Fe-30 Ni alloy was reinvestigated. The scatter in the habit Plane as determined from measurements on the mid-rib plane was smaller than reported previously, and the orientation relationship was determined from a martensite plate having known direction cosines. From agreement with crystallographic theory and from observations of the morphnlogy of the martensite plates, it is suggested that the mid-rib plane represents the starting region of the transformation. THE crystallography of the martensitic transformation in the iron-nickel system has been the subject of many investigations.&apos;-7 However, for iron alloys containing approximately 30 wt pct Ni, a wide scatter in the reported habit planes exists. In only one instance has a Laue determination of the orientation relationship been made, but the habit plane of the martensite plate studied was not determined experimentally. Because the most successful thermodynamic and kinetic analyses8 of the martensite transformation have been based on Fe-30 Ni alloys, it was considered desirable to reinvestigate the crystallographic features of the transformation in this alloy. Lieberman9 suggested that the most significant information for use in a theoretical treatment involves the determination of all the crystallographic observables from a single plate of the product phase. This was first done in essence by Greninger and Troiano10 for Fe-22 wt pct Ni-0.8 wt pct C, and is the method followed in this investigation. Some theoretical approaches11-13 in describing the martensitic transformation consider the overall transformation distortion to be comprised of a Bain distortion, or lattice deformation (PI, which produces the structure change, and a lattice invariant deformation (G), such as twinning, which together leave the interface or habit plane un-distorted. The incorporation of a rigid body rotation (R) results in a transformation distortion (E) such that the interface plane is left both undistorted and unrotated. In matrix notation, E = RPG where the lattice invariant deformation has "taken place" in the austenite. Bowles and Mackenzie14-16 in formulating a theoretical approach proposed an adjustable scaler quantity, denoted as 6, which represents a uniform dilation in the interface plane. This dilation acts to adjust the lattice spacings of the parent and product structures at the plane of contact. Using the dilation parameter as a variable, these authors were able to account for a continuous variation of habit planes in steels along a curve in the stereogram between {225}A and {259}A. she invariant line strain14 of the transformation is 6 RP, and the corresponding transformation distortion is E/6.) With the dilation parameter known, the invariant line is computed as the triple coincidence of the habit plane, the shear plane of the lattice invariant deformation, and the initial Bain cone.14 The contraction axis for the Bain distortion can be taken9&apos;17 as [010]A, and the correct permutation of the habit plane indices used. Kelly and Nutting 18 have established the shear plane (of twinning) in the tetragonal martensite of Fe-22 Ni-0.8C to be {112)M by means of transmission electron microscopy. The reasonable shear system is then {112}M which is equivalent to {011}A <01l>A. The as-

Jan 1, 1962

By F. V. Lene, E. J. Westerman

The recrystallization behaviors of two extruded and cold-drawn experimental sintered aluminum powder alloys, containing 1.75 and 3.0 pct Al2O3 by weight, were compared with that of extruded and cold-drawn commercially pure alumirzum. The kinetics of recrystallization of the alloys are described semiquantitatively. For the alloy containing 1.75 pct A l203 the rates of nucleation and of growth were also semiquantitatively determined. THE most striking property of aluminum alloys strengthened by a dispersion of Al2O3, the so-called SAP alloys, is their stability at elevated temperatures. One of the manifestations of this stability is their resistance to recrystallization after they have been cold worked. Most of the commercial grades of either the Swiss SAP or of Alcoa&apos;s Aluminum Powder Metallurgy Products have not been recrys-tallized after cold working, even when they are heated for a long time at a temperature near the aluminum melting point. Lenel, however, observed that the dispersion strengthened aluminum alloys with a larger spacing between the oxide particles than that of most commercial grades would recrys-tallize.1 It appeared to be of interest to further investigate the mode and kinetics of recrystallization of these alloys, and to compare their recrystallization behavior with that of commercially pure aluminum. Because homogeneous deformation of these SAP alloys in tension did not provide sufficient cold work to induce recrystallization, they were cold worked by wire drawing; the nonuniformity of this deformation unavoidably complicated the interpretation of the recrystallization studies. EXPERIMENTAL DETAILS Extrusions—Two types of sintered aluminum powder extrusions were used in this study. One type, designated AT-400, was produced from Reynolds atomized aluminum powder consisting of spherical particles averaging 3µ in diam and containing 1.75 wt pct of Al2O3. This powder was very similar to the R3M powder from which extrusions were previously prepared with an average spacing of 0.9µ between oxide particles.2 The second type, designated MD 2100, was produced from Metals Disintegrating Co. flake powder containing 3.0 wt pct of Al2O3, with an average flake thickness of 0.8µ. The average spacing between oxide platelets in MD 2100 extru- sions was 0.45µ.2 Powder compacts of 3/4-in. diam were extruded at 1000°F into 0.097-in. diam (AT-400) and 0.093-in. diam (MD 2100) wires by methods previously described.3 In order to compare the recrystallization behavior of sintered aluminum powder extrusions with that of wrought commercially pure aluminum 3/4 in. rod stock of 1100 F aluminum was extruded at 1000°F into 0.102-in. diam wire. Wire Drawing—Tungsten carbide dies were used for the AT-400 and 1100 F alloys. They had an included angle of about 15 deg and reduced the wire area approximately 7 pct per pass. Steel dies with an included angle of 11 to 13 deg and an average reduction per pass of 10 pct were used for drawing the MD 2100 alloy, because drawing this alloy through the carbide dies produced overdrawing defects. Heat Treatment—The cold-drawn wires were cut into small samples, and the deformed ends were etched off. The samples were each wrapped tightly in a single layer of aluminum foil, and individually isothermally annealed in a lead bath. Metallography—The modes and kinetics of recrystallization were determined by metallography. Mounted and polished specimens were anodized in a solution of 1.8 pct HBF4;4 examination under polarized light clearly revealed their grain structures. The recrystallized grains were generally much larger than those of the unrecrystallized matrix, and could clearly be distinguished because they alternated between maximum and minimum light reflection when the microscope stage was rotated, while the unrecrystallized matrix had a comparatively homogeneous "salt and pepper" structure. The fractional recrystallized volumes of the dispersion hardened alloy wires were determined by cutting and weighing of recrystallized and total transverse areas on photomicrographs. The recrystallized grains in the 1100 F alloy were too small to be cut out individually; therefore a combination of cutting and lineal analysis was used in this case. RESULTS AND DISCUSSION Modes of Recrystallization—The modes of recrystallization of the three alloys varied widely. In the 1100 F alloy nucleation and growth started in the region midway between the center and the surface;

Jan 1, 1961

Coal Division Mineral Industry Education Division Industrial Minerals Division Minerals Beneficiation Division Mineral Economics Division Mining, Geology and Geophysics Division Council R C Stephenson, Chairman (Ind Min Div), M D Cooper (Coal), H F Yancey (Coal), S D Michaelson (MBD), Will Mitchell, Jr (MBD), J R Van Pelt, Jr (MIED), S S Cole (Ind Min Div), J D Forrester (MGGD), L E Shaffer (MIED), E L Clark (MGGD), C W Merrill (MED), E H Robie, Acting Secretary, 29 West 39 Street, New York 18, New York

Jan 1, 1955

By F. J. Meringer, C. W. MacGregor

Various investigations have been carried out to determine the effects of cold-rolling on the common physical properties as represented by the yield strength, tensile strength, percentage of elongation and ordinary reduction of area. Very little information is available in the literature with reference to the effects of cold-rolling on the true stress-strain properties. Since the ordinary physical properties base both the stresses and the strains on the original dimensions of the test piece, which are greatly changed during the test, they are essentially empirical values and do not represent physically either true stresses or true strains. Consequently, they do not give an accurate picture of the effects of different conditions, both mechanical and metallurgical, on the strength and ductility of the material. It is not, however, the purpose of this paper to review the various advantages accruing from the use of true stress and true strain values, since these have been shown elsewhere,1-l2 but to utilize them in representing the various effects considered. The true stress S and true strain (e = q&apos;) values mentioned above for the tension test are defined by1 S =1/A and where P, A, Ao are the axial load, the instantaneous area and the original area of cross section, respectively. From simultaneous readings of the axial load and the dimension or dimensions determining the instantaneous area of cross section, the true stress-strain curve is constructed by using the relations given in Eq. I. It is then the object of this paper to discuss the general effects of cold-rolling on the true stress and strain values for a low-carbon steel together with the effects of age-hardening and directionality on these properties. Preliminary Investigations The material considered throughout the paper is a low-carbon steel of the following chemical analysis: C, 0.19 per cent; S, 0.118; P, 0.092; Mn, 0.85; Si, 0.01. The stock was received in bars, 7 by 3/4 in., and was first annealed for one hour at 1650°F. followed by a furnace cool. Photomicrographs of this material showed considerable directionality even in the annealed condition. In view of the possible effects of the nonhomogeneity and directionality in this material, it was considered important to investigate their influence on the true stress-strain diagrams. In addition, the effect of strain-aging in this steel was also made the subject of a preliminary study. Effect of Aging on True Stress-strain Properties Since a certain time element would elapse between the cold-rolling operation and the final tension tests, it was determined to investigate the aging character-

Jan 1, 1945

Organization Place Date 1918 American Society of Mechanical Engineers...:.. Worcester, Mass. June 4-7 American Institute of Chemical Engineers Berlin, N. H. June 19-22 American Concrete Institute Atlantic City, N. J. June 24-26 American Society for Testing Materials Atlantic City, N. J. June 25 28 American Institute of Electrical Engineers Atlantic City, N. J. June 26-28 New York Electrical Society, Annual Meeting... New York, N. Y. June American Institute of Mining Engineers Colorado Sept. 2-7 National Association of Stationary Engineers.... Cincinnati, 0. Sept. 9-13 British Iron and Steel Institute, Autumn Meeting London, Eng. Sept. 12-13 National Exposition of Chemical Industries New York, N. Y. Sept. 23-28 American Chemical Society Cleveland, 0. Sept. National Petroleum Association.: Atlantic City, N. J. Sept.

Jan 6, 1918

By Bedrand L. Johnson

THE phosphate-rock industry is built upon natural deposits of rocks and minerals in which the element phosphorus is present as a phoshate. The term ?phosphate rock? is a general one, applied to certain rocks of different degrees of consolidation and of diverese origin, charater and occurrence, composed of intimate mixtures of several phosphate minerals, largely submicrocrystalline, usually calcium phosphates, in commerical quantities, with various impurities. This term includes phosphatized limestones, sandstone, shales, and igneous rocks; nodular phosphates ; residual weathered phosphatic limestones ; vein phosphates ; and phosphatic sediments, both consolidated and unconsolidated.

Jan 1, 1944

By H. W. Immolz

The Palo Pinto limestone pools of Jones and Shackelford Counties were defined and almost fully developed during the year. A new Palo Pinto limestone pool was apparently discovered when the K. B. Nowels NO. 1 Pardue, S.F.I. and W. survey, south central Haskell County, was completed for a small well in Palo Pinto limestonc at 3395 to 3421 feet. In Haskell County, Amerada Petroleum Corporation&apos;s C. J. Kleiner, J. M. Cass survey, drilled to the Ellenburger with a total depth of 5607 ft., where it was abandoned. A Bluff Creek sand field of undetermined size was discovered by the Fain-McGaha No. 1 Dawson-Conway, P. G. Holcomb survey, which was chompleted for 37 bbl. at 1496 feet. The Cook field was extended by the Roeser-Pendleton B-113 Cook in sec. 113, E.T.R.R. Co. survey, the well being completed for 510 bbl. from Cook sand encountered at 1545 to 1568 feet. Several shallow pools of minor importancc mere also discovered in Shackelford County during the year. In Jones County the Lewis and the Steffens fields, in block 15, T. and P.R.R. Co. survey, were further developed during the year but the limits of the Steffens field were not definitely defined. The most, important discovery in Jones County was the Guitar field, where most of the wells that had been producing from the Cook sand were deepened to the Hope limestone, a new producing horizon. In Taylor County, the Bowles field, see. 28, B.A.L. survey, was defined arid proved to he of minor importance. The Goldsboro and Novice area, ill northwest Coleman County, continued to hold considerable interest because of the developments near the town of Novice, where new Ieases continued to be opened for production. Nolan County was placed in the ranks of the producing counties when the S. B. Owens No. 1 Tipton, sec. 43, block 19, T. and P.R.R. Co. survey,

Jan 1, 1940

By H. W. Immolz

The Palo Pinto limestone pools of Jones and Shackelford Counties were defined and almost fully developed during the year. A new Palo Pinto limestone pool was apparently discovered when the K. B. Nowels NO. 1 Pardue, S.F.I. and W. survey, south central Haskell County, was completed for a small well in Palo Pinto limestonc at 3395 to 3421 feet. In Haskell County, Amerada Petroleum Corporation&apos;s C. J. Kleiner, J. M. Cass survey, drilled to the Ellenburger with a total depth of 5607 ft., where it was abandoned. A Bluff Creek sand field of undetermined size was discovered by the Fain-McGaha No. 1 Dawson-Conway, P. G. Holcomb survey, which was chompleted for 37 bbl. at 1496 feet. The Cook field was extended by the Roeser-Pendleton B-113 Cook in sec. 113, E.T.R.R. Co. survey, the well being completed for 510 bbl. from Cook sand encountered at 1545 to 1568 feet. Several shallow pools of minor importancc mere also discovered in Shackelford County during the year. In Jones County the Lewis and the Steffens fields, in block 15, T. and P.R.R. Co. survey, were further developed during the year but the limits of the Steffens field were not definitely defined. The most, important discovery in Jones County was the Guitar field, where most of the wells that had been producing from the Cook sand were deepened to the Hope limestone, a new producing horizon. In Taylor County, the Bowles field, see. 28, B.A.L. survey, was defined arid proved to he of minor importance. The Goldsboro and Novice area, ill northwest Coleman County, continued to hold considerable interest because of the developments near the town of Novice, where new Ieases continued to be opened for production. Nolan County was placed in the ranks of the producing counties when the S. B. Owens No. 1 Tipton, sec. 43, block 19, T. and P.R.R. Co. survey,

Jan 1, 1940