Search Documents

By H. L. Gilbert, C. Travis Anderson, W. J. Kroll

IN a previous paper,1 the use of a split graphite tube resistor as a heater element for high-temperature furnaces has been described. The principal advantages of this type of construction are: I. The quadruple resistance or voltage compared with the top-bottom clamp resistor furnace. 2. Free expansion of the resistor. 3. The ease of protection from oxidation by air. 4. The firmness by which it may be held in the compact split clamp. This construction appeared to be well adapted for the fusion of high melting metals which require fusion in a vacuum. Zirconium metal had previously been melted at the Albany Laboratory of the Bureau of Mines in a vacuum arc furnace on a water-cooled plate, using a movable tungsten top electrode.2 When larger furnace capacities were desired, this type of construction proved to have its limitations as the largest tungsten rods available were about 7/8in. in diameter. The supply of direct current of high amperage for producing the . desired arc characteristics was also not available. Considerable difficulty was also found in compensating the atmospheric pressure on the flexible tubing of the movable electrode, at the same time keeping its movement sufficiently free to enable the striking of an arc on any part of the metal briquet. Among the various materials used for the electrodes in conducting experiments in the vacuum arc furnace was graphite. It was found that zirconium fused in a graphite cup with a carbon arc was contaminated to only 0.2 pct carbon. It was therefore decided to melt zirconium directly in graphite crucibles. A high-frequency furnace was first used in the attempt to melt the zirconium in graphite crucibles. The high-frequency heating coil was placed outside a silica tube which was evacuated by means of an oil diffusion pump backed up by a mechanical pump. The graphite crucible was thoroughly degassed at 1900°C before use. Considerable difficulty was experienced in operating at such temperatures in vacuo because, as shown by Zintl,3 carbon vapor formed and reacted with silica according to the reaction: Si02 + C = SiO + CO. The silicon monoxide which is volatile at these high temperatures reacted with the zirconium, thereby contaminating it with silicon and oxygen. The presence of carbon monoxide was detected by means of a gas discharge tube placed in the vacuum system as indicated by the pale blue glow of the tube as soon as carbon vapor started to react with the hot silica. The reaction between the carbon vapor and hot silica was avoided by the use of split cylindrical molybdenum shields which were inserted between the graphite crucible

Jan 1, 1948

By H. E. McCoy, C. J. McHargue

Single crystals of columbium containing various levels of oxygen, nitrogen, carbon, or hydrogen were deformed by slaw compression and impact loading at -196°C. For the slow deformation rates. 1500 to 1900ppm oxygen, 500 pprn nitrogen, or 150 to 200 ppm carbon completely suppressed mechanical twinning. Hydrogen in concentrations to 1000 ppm had no apparent effect on ease of twinning. Impact loading caused twins to form at all levels of contamination studied, and cleavage mucking on (100) planes occurred in specimens having higher interstitial contents. It is well established that, under some conditions, the bcc metals vanadium, columbium, tantalum, chromium, molybdenum, tungsten, and iron deform by formation of mechanical twins. In general it has been accepted that high strain rates, low temperatures, and large grain sizes are conducive to this kind of deformation.' Since interstitial impurity atoms have a relatively large effect on slip in bcc metals, one might sus- pect a similar effect on twinning. This could occur directly by the interstitials affecting the nucleation or growth of the twins or indirectly from their effect on slip. It has often been argued that the large increase in yield stress caused by interstitial elements enables the effective stress to exceed that necessary to cause twinning. The twinning stress has always been thought to be higher than the stress necessary to cause slip in pure metals.2 simonsens and Tipper and Hall4 found profuse twinning in very pure a iron even at room temperature, whereas Low and Fueste15 and Churchman and tottrell' observed twins in a iron containing carbon but none in highly decarburized specimens. McHargue7 reported mechanical twins to form readily in vanadium which contained 390 ppm interstitigs, while Clough and Pavlovic8 found twins only in the one or two grains adjacent to the fracture surface in vanadium which contained 1500 to 1700 ppm interstitials. In columbium, McHargue9 found high-purity single crystals to twin easily whereas Churchman,10 Adams, Roberts, and mallman,11 and Johnson12 studied less pure material and found far fewer twins. Although the results of Leadbetter and Argent13 showed oxygen to inhibit twinning, a study by sheely14 indicated the opposite to be true. Chromium exhibits this mode of deformation15,16 but there are no data regarding impurity effects. Lawley, Van den Sype, and Maddin17 observed twins in very pure molybdenum single crystals deformed at 77°K but none in less pure crystals. The work by Cahn18 and

Jan 1, 1963

By E. C. W. Perryman

The recovery and recrystallization characteristics of superpurity aluminum have been investigated using electrical resistivity, X-ray line breadth, and hardness measurements for the former and the micrographic method for the latter. The three different properties recover at different rates and have different activation energies. The recrystallization results agree well with Avrami&apos;s theory and furthermore indicate that the perfect subgrains formed during recovery are not the nuclei for re-crystallization. WHEN a metal is plastically deformed, its physical and mechanical properties generally undergo considerable changes and by subsequent annealing these changes are partly or wholly annihilated. Thus, a recovery process can be discussed, taking this term in its general sense. In practice, however, there is reason to discriminate between two apparently different processes, one most easily followed at low temperatures, in which the properties return to an almost constant value between that of the cold worked and fully annealed material, and a second process in which the properties return to their original values before cold working and which is accompanied by the formation and growth of new grains having an orientation different from that of the matrix. In this paper the word recovery will be taken to mean the changes in some property as a function of annealing time which occur either without the appearance of new grains or under conditions such that the new re-crystallized grains are very small (= 2 microns), are very few in number, and substantially do not affect the property being measured. This definition is rather abitrary, for it will depend upon the sensitivity of the technique used for the observation of new recrystallized grains, which in the present work was about 1 to 2 microns. However, it is helpful to use the term recovery in this sense and to reserve the term recrystallization for the processes of nucle-ation and growth of new grains in the cold worked matrix. Although considerable work has been done on recovery and recrystallization, most workers have based their study on the measurement of one or perhaps two parameters. Since very small amounts of impurities have such a profound effect on the recrystallization characteristics of a pure metal, it becomes extremely difficult to correlate one piece of work with another. With this in mind, the present work on recovery and recrystallization was done on the same material. Experimental Procedure Material Used and Fabrication: The composition of the superpurity aluminum used throughout this investigation was 0.002 pct Cu, 0.003 pct Fe, 0.003 pct Si, and <0.001 pct Mg. The ingot was hot rolled to 0.250 in., annealed, and cold rolled to 0.034 in. A large number of reductions and intermediate anneals were carried out so as to produce material with a minimum of preferred orientation and maximum homogeneity. For the recovery part of the investigation, the final cold reduction was 20 and 80 pct and for the recrystallization part, 20 pct. After each pass in the cold rolling process, the material was quenched in cold water in order to keep the rolling temperature as near room temperature as possible. Annealing Procedure: For the recrystallization work, specimens 1x1 in. were cut from the 0.034 in. cold rolled sheet and a hole was drilled in each through which a wire was threaded to support it in the salt bath. The temperature of the salt bath was controlled to +2°C and the time taken for a specimen to reach temperature was approximately 5 sec. These 1 in. squares were then divided into three groups, one of which was given 5 min at 318°C and another 2 hr at 244°C. These treatments were such that recovery was almost complete and a well defined subgrain structure produced. Separate specimens of each group were annealed for different times at 301°, 318°, 355°, and 373°C, i.e., three specimens for each annealing time. The delay between finish of cold working and start of annealing was about 1 hr. For the recovery work, strips 0.062 in. thick were cut from the cold worked sheet, annealed, and then given the last cold rolling operation. This was done for each annealing temperature. By this means it was possible to minimize the delay between cold working and annealing. In general, all measurements were carried out within 1 hr of the last cold rolling operation. Annealing at low temperatures was done in an oil bath the temperature of which was maintained constant to +1°C. Electrical Resistivity Measurements: Strips 20x0.5x0.05 in. were machined and the electrical resistance measured using a Kelvin double bridge. Measurements were made in an oil bath maintained at 20rt0.1°C. The same specimen was used for the complete isothermal annealing curve.

Jan 1, 1956

By R. I. Jaffee, R. W. Wood, H. R. Ogden, D. N. Williams

D. W. Morgan, D. H. Polonis,and R Taggart (Uuivevsity of Washington)—Dispersion hardening in titanium-copper alloys is of particular interest to us in view of our current research activities in phase transformations in this system. We should like to raise questions concerning several points in the paper. The flow sheet for fabrication and heat treatment (Authors&apos; Fig. 2) indicates that 0.045 in sheet is treated in air for a period of 30 min at 1625&apos; F. Were any surface coatings used to reduce oxygen contamination at this stage? One might ask the same question about the 8-hr annealing at 1400F. Contamination during fabrication and heat treatment of such thin sections could be considerable without some protection. Although the authors state that no analyses were made, they note that contamination masks the distinction between base strengths of the different alloys. In this connection, what fabrication and heat-treatment procedures were used in obtaining the comparison specimens of unalloyed titanium? It would seem that oxygen contamination must cast doubt on the reliability of the proposed strength-mean free path relationship deduced from the data (Authors&apos; Fig. 5). It is usual to evaluate dispersion hardening on the basis of yield strength rather than ultimate tensile strength, and accordingly we have replotted the authors&apos; Fig. 1 to include yield strength. The yield strength data do not justify a discontinuity of 3 p at either 70" or 1000 OF. Can the authors substantiate their conclusion 1) with additional evidence? Otherwise it seems that they have succeeded only in pointing out the possibility of a linear relationship between yield strength (or Uts) and the logarithm of the nleall free path, comparable with that demonstrated for iron-carbon alloys (Authors&apos; Ref. 2). Work in our own laboratories has shown that the hardness of the phase Ti,Cu is 235 Vpn, and that of a 1.5 pct (at.) (2 pct wt) supersaturated solid solution of copper in titanium is 190 Vpn. This small difference in hardness would suggest that Ti2Cu should not be a very effective dispersion-hardening phase. D. N. Williams, R. W. Wood, H. R Ogden, RI. L Jaffee (authors&apos; reply)— This discussion brings out some interesting points with respect to the interpretation of our results, and also raises some questions regarding the experimental technique. We shall consider the latter questions first. Contamination during heat treatment was prevented by performing all heat treatments in argon-filled Vycor capsules. Fabrication was done in air, with the usually observed scaling. Scale was removed

Jan 1, 1962

By A. J. Opinsky, J. L. Orehotsky, L. L. Seigle

TUNGSTEN incandescent lamp filaments possess a typical structure of elongated crystals generated upon heating the silica-alumina doped wire rapidly to 2200°C or above.&apos; It is known that these very long grains exhibit a preferred orientation, but there is disagreement in the literature about the direction. Rosi,2 and later Rieck; reported that the <135> direction of the large grains in doped wire tends to be parallel to the wire axis, but other investigators have reported that <110> is the preferred direction. Some further investigation of the orientations of large grains in tungsten wire is reported in the following. In the first series of experiments, 6-in. lengths of wire were heated rapidly in a bell jar. Current was passed through the wires and within 15 sec the wires had reached the grain coarsening temperature. 2 min was the annealing time. Wires were heated in hydrogen, nitrogen, and vacuum. Later, wires were heated in an evacuated tube furnace in which the tube was brought from 1000°C to the operating temperature within 25 sec and held 5 min at temperature. Temperatures were measured with an optical pyrometer. After heating, the wires were electro polished and etched, and the orientations of the large grains determined by the Laue back-reflection technique. Most of the experiments were carried out on a single spool of 5-mil silica-alumina doped wire, but a few tests were made with 5-mil thoriated tungsten wire and 10-mil undoped wire. 5-mil silica-alumina doped wires were reduced to smaller diameters by electropolishing for one series of experiments. Orientations of the large grains in silica-alumina doped wire heated in various atmospheres are presented in Figs. 1 and 2 as plots of the direction of the wire axis relative to the cube axes of the crystals. For most of the conditions studied, the preferred orientation cannot be described as a single direction, but the points might be said to group themselves about the arc passing through <110> and <113>. This arc is the trace of the (211) planes and contains the <135> and <124 as well as <011> directions previously observed by other investigators. The preferred orientation of large crystals in silica-alumina doped tungsten wire might be most generally described, therefore, by stating that a (211 ) tends to be parallel to the wire axis. This is true also for wires reduced in diameter by etching and furnace heated, Fig. 2, i.e., the recrystallization texture is apparently unchanged upon removal of the outer layers and substantially independent of the temperature gradient within the wire.

Jan 1, 1962

By R. H. Johnson

All strata not yielding oil or gas in commereial quantities or a corresponding amount of water may be called dry in a wide sense. In petroleum geology, however, we may exclude all sands of too low or fine porosity to yield gaseous or fluid contents to the hole drilled in the sand before any original pressure that its contents may be under is disturbed. Most rocks are of this class and they are not reservoirs in our definition; their "dryness" is wholly a matter of course. What are the contents of the pores or what is the exact porosity of such rocks is of almost no concern to us, for economically they are "dry." What does interest us is the content of a rock having sufficient porosity and the pores of sufficient size to yield oil or gas in commercial quantities, if they were present under original pressure. Dryness of these reservoirs is a matter of supreme practical importance. Three views current as to such dryness seem, to me, to apply in a few cases only. It is the purpose of this paper to give reasons for this position and for believing that, in ordinary sedimentary rocks, there is only rarely a reservoir of competent porosity and undisturbed pressure that is dry in the sense of not yielding water, oil, or gas when first penetrated. 1. Gardnerl writes of some Kentucky sands, "There has never been present any salt water or other water in the sand." Absence of water cannot demonstrate this position. It is necessary to show that the rocks were not laid down in water, but in air, and that they became so enclosed, while still above the water-table of the ground water, that water has not been able to enter since. Most of these sands, and certainly the productive limestones, were deposited in water; and such sands as have been commercially productive show no reason for believing that the overlying shale or limestone was not laid down progressively from one direction and in water that would have flooded it. NO adequate explanation has been offered for this hypothesis, which is so inherently improbable. 2, Reeves2 urges that "sands originally water filled may have been drained of their water and not filled when later covered." It is difficult I James H. Gardner: Kentucky as an Oil State. Science, N. S. (1917) 46, 279-280. Reeves: Origin of the Natural Brines of Oil Fields. Johns Hopkins Univ. Circ., N. S. (1917) N. 3; Absence of Water in Certain Sandstones of the Appalachian Field. Econ. Geol. (1917) la, 354-378.

Jan 1, 1921

By Axel Hultgren, B. Herrlander

THE original aim of the present investigation was to study the mechanism of cracking on hot-deforming red-short steels. During the microscopical examination of hot-deformed soft steels attention was directed to various patterns of so-called veining in the ferrite, as related to variations in the deformation procedure and the heat-treatment of the steels. Later, a hot-short steel of higher carbon content was also studied. Historical Storey1 suggested in 1914 that veining developed during the gamma-alpha transformation from the growing together of several alpha nuclei within the same gamma grain. Rawdon and Berglund2 found that in soft steel, forged somewhat below A3, the ferrite showed profuse veining, while usually there was very little veining after forging at very high temperatures. They emphasized the similarity between the pattern of veining formed by deformation about 600°C and that of slip lines formed by deformation at room temperature. They also suggested a relation between the gamma-alpha transformation and veining. They became convinced from etching and from the slip-line pattern in cold-deformed material, that the orientation within any ferrite grain showing veining was uniform. In recrystallized alpha grains veining was absent. Veining did not appear materially to affect the properties of ferrite. Independent of veining, continuous networks, attributed to preexisting delta and gamma grain boundaries were sometimes found, the former only in cast steel. Those networks were usually connected with small inclusions. Ammermann and Kornfeld3 confirmed that recrystallized ferrite grains showed no veining. In soft steel annealed between At and A3 there was veining only in the ferrite grains formed during cooling. Electrolytic iron deformed at 880° was free from veining. Bannister and Jones4 considered veining in ferrite to be a manifestation of microscopical and submicroscopical inclusions. Northcott,5,6 in studies on veining in steel and other metals and alloys, attributed veining to oxide inclusions. It could generally be removed by annealing in hydrogen, a suggestion earlier made by Rawdon and Berglund. Tritton (discussion in ref. 5) emphasized the importance of perfect polishing for bringing out veining by etching, and thought the veins or subboundaries were produced during the gamma-alpha transformation, when it was rapid, as a result of the volume change and contraction during cooling.

Jan 1, 1946

By Ralph Hultgren

IN solid solutions atoms of differing sizes occupy the same crystalline lattice, requiring that some of them be compressed and others expanded. The energy involved has been called misfit strain energy and is an important concept of crystal chemistry. If the atomic sizes and elastic constants of interatomic bonds are known, the misfit energy may be calculated,&apos; provided certain simplifying assumptions are allowable. Usually, isotropic crystals are assumed and interatomic distances are taken to be the statistical average determined from X-ray diffraction. Such calculations yield values of the misfit energy of the order of 1 or 2 kcal per atom in alloys such as Au-Cu at compositions of 50 atomic pct. However, evidence has accumulated in recent times that atoms change their sizes with composition of alloys, implying electronic rearrangement of the bonds. The size changes have been found particularly by application of the X-ray method developed by Warren, Averbach, and Roberts.&apos; Thus, Averbach, Flinn, and Cohen3 determined radii in Au-Cu alloys. Oriani4 showed that these new radii led to a calculated misfit energy in disordered AuCu, which was decreased from the values calculated by the usual theory more than twenty-fold, to only 80 cal per g atom. Thermodynamic calculations from the phase diagram5 also show misfit energy to be no more than a few hundred calories per g atom in this alloy. The question of what electronic rearrangements are possible therefore becomes compelling in estimating misfit energy. In the following pages the results of certain calculations on the AuCu tetragonal superlattice are submitted. Conclusions drawn from these should be applicable in large degree to disordered solid solutions. As in all ordered states, bonding distances in the superlattice are individually known, rather than being merely average distances as found from lattice constants of disordered states. Moreover, only the Au-Au and Cu-Cu distances are strained; the elastic constants of these are known in the elementary state. In the usual calculation it is necessary to assume elastic constants for Au-Cu bonds. Misfit energy has thus been calculable without the need of many simplifying assumptions usually made. It is still assumed that equilibrium bond lengths and elastic properties of the bonds are the same in the alloy as in the pure metals. As previously discussed, this is probably not correct. Also assumed is that the bonds are not affected by strain of neighboring bonds. A calculation of Young&apos;s modulus from compressibility data shows this to be far from true; extensive electronic rearrangements take place. It would seem that misfit energy cannot be calculated from elasticity data for the elements. The usual methods may, however, give an upper limit which is often much higher than the true value. The question of electronic rearrangement is, of course, a complex one. Pauling&apos;s theory gives a simple, approximate treatment of the relation between type of bond and bond distance. This has been applied with some success to the Au-Cu system, as will be shown in a later section. Misfit Energy in Au-Cu Alloys Hume-Rothery and Raynor6 discuss the Au-CU system as a type example of strain energy. The gold atom is 12.8 pct larger in diameter than the copper atom, near the size factor limit beyond which solid solubility is severely restricted. They therefore consider the misfit energy to be large, a conclusion for which they believe they find evidence in the phase diagram. Gold and copper are completely miscible in the solid state, but the alloy has a minimum melting point at an intermediate composition. From this Hume-Rothery and Raynor conclude that the strain energy is nearly large enough to prevent miscibility; the phase diagram tends toward a eutec-tic type. In Ag-Cu, which has almost identical size relationships, solid miscibility is quite limited; whereas in Au-Ag, where atomic sizes are nearly the same, there is complete miscibility without a minimum in the melting point. From their arguments the heat of formation of Au-Cu would be expected to be endothermic or only slightly exothermic, that of Ag-Cu to be endothermic, and that of Au-Ag to be exothermic. Deviations, from Ve-gard&apos;s law of additivity of atomic radii support these conclusions, since Au-Cu and Ag-Cu both have pronounced positive deviations, and Au-Ag has a negative deviation. Nevertheless, Au-Cu alloys form exothermically; indeed, considerably more exothermically than Au-Ag, Table I. Hence, strain energy must be much less important in this case than Hume-Rothery and Raynor have supposed.

Jan 1, 1958

By St. John, JOHN Ancel

AT THE New York meeting of the Institute of Metals Division in February, 1927, Jeffries and Tarasov presented a paper on Tungsten and Thoria,1 in which the experimental facts were interpreted in accordance with the prevailing point of view that the thorium is present as the oxide, thoria. In discussing that paper the writer of this article stated that his own investigations, which had been unavailable to Messrs. Jeffries and Tarasov, showed that the so-called "thoriated" tungsten filament failed to disclose any sign of the oxide when subjected to X-ray diffraction analysis, but did give plenty of evidence of the presence of metallic thorium and, in certain cases, of a material believed to bean alloy of tungsten in thorium. Since these results were completely at variance with the results reported in the paper under discussion and with other papers which had preceded it from time to time in the past, it was felt that further discussion should be withheld until the new evidence could be presented adequately. Arrangements were therefore made to present that evidence at this meeting. MATERIAL EXAMINED The material examined was high-emission filament such as is used in radio tubes, supposed to contain about 0.75 per cent. thoria. The actual analysis was not stated.

Jan 1, 1927

By George Holderer

THE Enos Coal Mining Co. owns an area underlaid by coal, of approximately 5000 acres. This property is situated 9 miles from Oakland City, in the southern part of Indiana. The mine has been in operation 5 years, during which time some 350 acres have been stripped, and stripping is now proceeding at the rate of about 100 acres per year. TOPOGRAPHY These coal lands consist of gently-rolling agricultural fields which in times past were farmed for wheat. There is no timber and but very little brush to contend with. The overburden varies from 10 to 60 ft. in depth. A thin bed of slate overlies the coal, followed by limestone, which varies from 3 to 15 ft.; this, in turn, is followed by sandstone, slate, shale and clay. There is a very wide variation in the thickness of all of these beds. The coal has an average thickness of 5 ft., and is a good grade of high-volatile bituminous. MINING OPERATIONS Briefly, the sequence of operations is: drilling holes with churn drills; blasting with L. 0. X.; stripping overburden to uncover coal and, finally taking up the coal, which is sent to a tipple for cleaning and sizing. In order to show what is being accomplished by the use of L. 0. X., results will be shown from May to November, 1926, inclusive, and compared to the same period in 1925. No work was done in No. 1 pit during May, 1925, hence all figures are unduly low for this month. Drilling There are in use 8 No. 44 Loomis clipper drills, caterpillar traction with gasoline engines; recently these were equipped with extra heavy stems, measuring 26 ft. overall and 4 ½ in. dia., to increase the speed of drilling. In addition, 4 No. 1 ½ G non-traction drills have been acquired, and are used to save the time of the larger drills when recleaning is necessary.

Jan 1, 1927

By E. A. von Hahn, T. R. lngraham

The rates of cementation of pd11 on electropolislzed copper cylinders were studied in aqueous perchloric acid solutions at pdII concentrations from 0.02 to 0.1 mM. The cylinders were rotated at various high speeds to reduce the diffusion layer to a minimum thickness. In 0.1 M HClO4 solutions, the cemented palladium is dense, adherent, and shiny. The rate data indicate that there are two stages in the deposition. In the initial stage, the rate of cementation is first order with respect to the pd11 concentration. The second stage is much slower. The first stage is consistent with rate control by the diffusion of pd11 ions to the copper surface, and/or chemical reaction at the surface, and the second stage is consistent with rate control by the diffusion of copper ions from the cop-per surface through the deposit of cemented palladium, out to the main body of solution. In 0.001 M HClO4 solution, only the first stage is evident and the rate is more rapid than at higher acidities. This rate enhancement is attributed to PdOH+ ions that predominate at low acidity and aye more reactive than unhydrolyzed pdII. The deposit is porous and loose. All of the cemented deposits are Pd-Cu alloys rather than pure palladium. Activation energies are 9.5 kcal per mole in 0.1 M HClO4 and 7.4 kcal per mole in 0.001 M HClO4. CEMENTATION or displacement reactions occur between aqueous solutions of metal salts and immersed metals according to the general equation: where M1 is electrochemically the more noble metal. Although these reactions find considerable application in metals recovery processes (e.g., the cementation of copper192 or gold3), in electrorefining processes (e.g., solution purification before electrolysis4,5), and in plating and metal finishing processes,6 few studies have been made of the kinetics of such reactions.7-9 The rates of cementation reactions will depend on one or more of the following factors: a) chemical reactions at the metal-solution interface; b) ionic transport to or away from the interface;&apos;-&apos; and c) the character and adherence of the cemented deposit on the substrate metal,6a,10 because the deposit will inhibit the rate of transport of dissolving ions (M2m+) into the solution. In this investigation, the kinetics of the early stages of cementation were studied to obtain an understanding of the reaction mechanism under conditions in which the deposit was sufficiently thin that its inhibiting effect could be disregarded. To achieve this, very dilute solutions in Mn+ were used. Cementation rates are often fast in the initial stages and transport-contr01led.9 To find conditions under which transport control would shift to chemical control, very high stirring velocities were used to minimize the diffusion-layer thickness. Of the many possible cementation reaction systems, the palladous perchlorate-copper system was chosen for this investigation because it was believed to involve simple ion-for-ion exchange. In addition, there are no interfering side reactions, such as the reduction of hydrogen ions, and anion effects are usually absent in perchlorate solutions. EXPERIMENTAL Materials. Reagent-grade chemicals and redistilled water were used in all experiments. Palladous perchlorate stock solutions (0.02 M) were prepared by dissolving 99.5 pct Pd sponge (Johnson Matthey and Mallory Ltd.) in concentrated nitric acid, adding concentrated perchloric acid, evaporating twice slowly (to prevent PdO precipitation) to a small volume, and diluting in a volumetric flask. To prevent gradual hydrolysis of PdII, the stock solutions were made 0.4 M in HC104. Tests for pdIV (Ref. 11) and C1- were negative. Copper strips were cut accurately from 0.025-in. electrolytic tough pitch sheet, Sample 1 (Mines Branch stock or Canadian Copper Refineries Ltd., Montreal), or from oxygen-free, high-conductivity sheet, Sample 2 (American Metal Climax O.F.H.C. brand, sold by Utility Brass and Copper Corp., Brooklyn). The strips were annealed for 1 hr at 335°C in a stream of purified nitrogen (Union Carbide, Linde Division) which had

Jan 1, 1967

ZAY JEFFRIES (communication to the Secretary*).-Having seen Mr. Ruder&apos;s micrographs of electrolytic iron, I am of the opinion that the tentative explanation offered verbally is correct. Mr. Ruder has informed me that both of the samples were cooled in the furnace. The structures as shown in Figs. Band C resulted from the last cooling through the temperature range near 900° C. Assuming that the rate of cooling through this range was the same in both cases, the differences in struc-ture of the gamma iron in the two samples must have ]:)]-ought about, in some way, the variations in grain size and grain shape shown in Mr. Ruder&apos;s micrographs. Even though there is practically no temperature gradient in muffle-furnace treatment during the "soaking" period, during both heating and cooling, there is a temperature gradient. Fig. B should normally have been fine-grained prior to the last cooling, due to the comparatively low maximum temperature to which it was heated (1,000° C.). On cooling through the allotropic change point at 900°, these small gamma grains, iii conjunction with the temperature gradient, should form an ideal condition for germination. Germinant grains should form near the surface of the sample (the-colder portions) and the small gamma grains would be progressively absorbed as they changed into non-gamma iron. The elongated grains in Fig. B, converging toward the axis of the piece of iron, indicate that germinant grains first formed near tine surface and advanced their boundaries toward the interior of the sample. The gamma grains in Fig. C, however, should have been large, prior to the last cooling, due to time high temperature treatment (1,250 to 1,300° C.). In other words, to use one of Prof. Howe&apos;s expressions, the crystalline organization should be very complete and the change from gamma to non-gamma iron on cooling would take place reluctantly. This resistance to change from gamma to non-gamma iron would probably be greater in the interior of the large grains than at their boundaries. Non-gamma iron centers or nuclei would probably form at a great many places, at the grain boundaries, thus masking the influence, to a large

Jan 6, 1917

By H. P. Reid

Portland cement is manufactured under either of two general processes, the wet or the dry. The raw materials in general consist of limestone, shells, marl, cement rock, clay, shale, blast-furnace slag, etc., so proportioned that when calcined the desired composition will be obtained. In the wet process the combined raw materials are ground to extreme fineness in the presence of water, to a resulting slurry of from 30 to 50 per cent moisture content. The moisture of the slurry may be reduced by thickening or filtering processes for feed to the cement kiln or the slurry may be fed to the kiln with its full percentage of water. The dry process differs from the wet in that the surface moisture of the raw materials is removed prior to grinding. For the drying of these materials heat is required and in most cases is produced by burning of raw fuel either on grates or stokers before the dryers or in suspension in the form of pulverized coal, oil or gaseous fuel. The power required for making the cement is sometimes purchased from outside sources; sometimes it is generated from the heat in the waste gases leaving the cement kilns, or generated at the plants direct from raw fuel. We have, then, as the possible requirements of fuel in the cement industry: (1) fuel for power generation, (2) fuel for drying raw materials, including coal before pulverizing, and (3) fuel for burning the raw materials to cement clinker. There is no essential difference in quality required in fuel for burning the materials by the wet and,dry processes. Fuel fob Power Generation As there is nothing peculiar to steam generation in a cement mill power plant as compared with other power plants this use classification of such fuel will be found in other reports of the Technical Committee. Fuel for Drying Raw Materials and Coal Any fuels capable of producing moderate temperature gases may be used for drying raw materials. There is no problem of material contamination from fuels at this part of the cement manufacture to be guarded against.

Jan 1, 1932

By T. Sterry Hunt

I have in a previous communication called the attention of the Institute to the geognostical relations of the crystalline iron ores belonging to the Eozoic racks of North America, at which time I noticed (1) those associated with the gneisses of the Laurentian series, (2) those of the crystalline schists of the Huronian, as seen in Northern Michigan and in many localities along the Appalachians; (3) those of Southeastern Missouri associated with petrosilex porphyries and schistose rocks, and presumed to be nearly of the same age with those of the Huronian schists; and (4) the highly titaniferous ores of the Norian or Labradorian series from Essex County, N. Y., and from various parts of Canada. It may here be remarked that, although the presence of titanium is most marked in the Norian ores, this objectionable element is sometimes found in the schistose ores of the Huronian, and to a lesser extent in certain Laurentian magnetites. The ores of each of these geological divisions have their individual peculiarities, which serve to distinguish them to such a degree that the magnetites and specular ores of the Laurentian are as dissimilar to those of the Huronian as are the crystalline silicated rocks of the two series. As it happens that these mineralogical dissimilarities correspond to diflerences in composition and in adaptability to different processes of manufacture, which, in the present condition of the iron industry, are objects of special consideration to metallurgists, it will be seen that these geognostical distinctions become of great practical importance. During the past year I was desired by the Geological Commissioners of Pennsylvania to devote a few weeks to certain geological studies in the eastern part of the State, in the course of which I was enabled to make, with regard to its crystalline iron ores, some observations to which I desire to call the attention of the Institute. I may mention, in passing, that the great South Mountain belt, from the Schuylkill to the Delaware, is composed of Laurentian rocks, in which are found the characteristic ores of the series like those of the Highlands of New Jersey and New York, and of the Adiron-

By R. L. Skaggs, R. A. Molsberger

DARKEN1 has pointed out that in most binary liquid metal systems the thermodynamic behavior is relatively simple in the terminal regions. Between the terminal regions. the thermodynamic behavior is more complex. To describe this behavior. he has defined the stability and the excess stability as shown: Qualitatively, an inflection in the free energy of mixing curve results in a change in sign, from positive to negative, of the stability function and is associated with a phase separation. Conversely, a large, positive stability over a narrow composition range is indicative of compound formation. Darken has shown a number of binary systems where the excess stability is constant in the terminal regions and varies sharply in between. The work described in this communication was undertaken to study the variation of stability and excess stability in a binary system as a function of temperature. In addition, this work provides means for the direct comparison of previously existing equilibrium (activity) and calorimetric (molar heat of mixing) data on the Hg-Sn system. Hildebrand2 measured the vapor pressure of mercury over Hg-Sn alloys at 595°K. Calorimetric studies were made by Kawakami3 at 523°K and by Kleppa4 at 423°K. In this work, the equilibrium vapor pressure of mercury was measured over a series of Hg-Sn alloys at 453" and 503°K by use of fused salt manometer with one arm exposed to mercury vapor. Deflections of the fused salt manometer were balanced by argon pressure on the other arm. Argon pressure was subsequently measured with a dibutyl phthalate manometer. Tin activities were determined by graphical integration of the Gibbs-Duhem equation making use of the familiar Darken5 a function. The activities of mercury and tin showed positive deviation from Raoult&apos;s law at 453 and 503°K as Hildebrand observed at 595 K. In Fig. 1 are plotted log activity of mercury and tin vs reciprocal temperature. The plot contains activities obtained from calorimetric studies. Hult-gren6 has calculated activities from Kleppa&apos;s4 data, and Kubaschewski7 has similarly treated Kawakami&apos;s7 data. Presumably. in each case a regular solution was assumed. For regions where a two-phase field

Jan 1, 1970

By R. W. Raymond

The " law of the apex," with its vague and impracticable " extra-lateral rights," is defended—if it is defended at all— on the ground that it is required by the peculiar conditions of the mineral deposits or of the mining industry in those States and Territories in which it is operative. As I have repeatedly attempted to show, in papers read before the Institute, and most recently in an article on " The Relation of the Mining Law of the United States to the Davelopment of its Mineral Resources," published in Mr. Rothwell&apos;s Mineral Industry for 1894, there is now (whatever may have been the case formerly) no peculiarity of conditions in the regions referred to which warrants the subjection of their mining industry to a burden so vexatious and injurious as this law imposes. I have called attention to the absolute freedom from expensive litigations over mining titles enjoyed by all other States in the Union and all other regions in the world where the absurd extra-lateral right does not exist.

Jan 1, 1896

By J. A. Pusateri, M. A. Orehoski, N. R. Arant

Plant studies on commercial-size ingots and laboratory experiments with induction furnace heats have demonstrated that the major source of ingot cracks is associated with two conditions arising during top-pouring practice: 1—solidification during pouring, and 2—turbulence created by the impact of the stream. Methods of controlling the two factors were effective in eliminating or significantly reducing ingot cracks. BECAUSE the process of removing surface defects from hot-rolled product is so costly, the steel industry is striving constantly to develop methods of preventing, or at least decreasing, the occurrence of surface defects. Investigations have revealed steel-making and processing variables related to major surface defects, and controlling these variables has led to improvements in surface quality. However, the fundamental causes of major surface defects, such as ingot cracks, have not been determined; consequently such defects persist. At the Research and Development Laboratory of the United States Steel Corp. in Pittsburgh, a seam research program was initiated to determine the fundamental causes of certain major defects. As a part of the program, ingot cracks in killed, finegrained C1020 steel were selected for study. A cost survey indicated that, of the steels produced in sizable tonnages, carbon steels in the range of 0.18 to 0.23 pct C content require the most conditioning. Since this is particularly true of C1020, any im-provement of surface quality that might be effected from the study would be beneficial. Also, since the frequency of ingot cracks is exceedingly high for this grade, the steel would provide an excellent opportunity for a thorough study of the ingot-crack-ing problem. Why steels in this carbon range tend to exhibit more ingot cracks than do other steels was not considered in the investigation; the mechanism of ingot-crack formation was of paramount importance. Also, only the top-pouring practice was considered in the investigation, because this pouring procedure is used more extensively than others. In this study, ingot cracks are defined. as deep surface defects that sometimes are observed in an ingot prior to rolling but usually are observed during the initial stages of rolling on the primary mills. These defects may occur at any angle to the rolling direction but are most prominent in the transverse or nearly transverse direction. The appearance of ingot cracks on the rolled product varies with respect to their angle of formation and with the extent of rolling after their first occurrence. Ingot cracks are also termed "deep seams," "arrowheads," "irregular cracks," and "transverse cracks." Fig. 1 shows ingot cracks in a rolled bloom. This report summarizes: l—the exploratory investigation of commercial-size ingots; 2—the labor-atory investigations related to determining the source of ingot cracks and developing corrective methods; and 3—the plant evaluations of laboratory methods of decreasing ingot cracks. Exploration of lngot Cracks in Commercial Ingots Materials and Experimental Work: At the Du-quesne Works of the United States Steel Corp., five heats of killed, fine-grained C1020 steel were selected for the exploratory phase of the seam-research program. The heats were made by open-hearth practices that are considered most conducive to good surface quality. They were top-poured into 22x25 in. big-end-up hot-topped molds. One as-cast ingot and four other ingots rolled to the following cross-sectional sizes were set aside from each heat: 16x20 in.; 9 1/2x10 in.; 5x5 in.; and 2x2 in. The processing of the heats, from the time of charging in the open hearth furnaces to the end of the rolling operation on the primary mill, was observed carefully and recorded. The cast ingot and the four rolled ingots from each heat were examined for surface defects. The cast ingots were split longitudinally near the vertical center plane to expose the ingot structure beneath badly cracked regions. Also, a series of transverse sections was cut at about 1 in. intervals in a badly cracked area of one ingot. All sections were ground, polished, examined by sulphur printing and deep etching, and photographed. The rolled ingots

Jan 1, 1955

By Maurice Cook

NORMALLY the grain size of cold-rolled and annealed copper sheet is of the order of 0.02 to 0.06 mm., and 0.1 min., for example, would, for many purposes, he regarded as undesirably large. The occurrence of very large grain sizes in a few isolated cases of copper sheet led the authors to inquire into the causes giving rise to such unusual results. During the course of this work a very large number of experiments have been carried out to determine the effect of varying the conditions of the final and penultimate rolling and annealing operations, as well as of earlier rolling and annealing treatments. A correspondingly large number of observations have been made, but within the space of one paper it is possible only to outline relatively briefly what appear to be the more important of these experiments and observations. The presentation of the results does not follow chronologically the order in which the experiments were carried out and they have been rearranged for the sake of simplicity. In recent years a considerable amount of work has been done in connection with the recrystallization and grain growth in metals and alloys that take place on annealing, and a number of generalizations have been reached. It is known that the crystal size in wrought material in the annealed condition depends on a number of factors, including the time and temperature of the annealing operations and the nature and amount of cold-work to which the material has been subjected. Results, however, that are not covered by these generalizations can be obtained, which, in some instances, may even appear to he contradictory to them. According to these generalizations gradual changes in grain size are effected by progressive variation of the factors involved in annealing or cold-working. The work recorded in the present paper, however, relates to the development of grain sizes which are far beyond those normally encountered in ordinary annealing operations and far in excess of grain sizes which, according to ordinary standards, would be classified as coarse or large.

Jan 1, 1938

By A. I. Krynitzky, Henry S. Rawdon

The most commonly used material for terminals in spark plugs is commercial nickel wire, because of its relatively high temperature of melting, excellent heat conductivity, and slow rate at which the metal is oxidized, even upon continued heating at high temperatures. The grade commonly used for this purpose averages 97 per cent. nickel, the remainder being manganese, cobalt, iron, copper, with traces of other impurities always found in commercial nickel. A peculiar and interesting type of deterioration that occurs in these nickel terminals during the service life of the spark plug was recently brought to the attention of the Bureau of Standards. Commercial spark plugs vary greatly in their size and shape and the terminals differ considerably as to their relative size, shape, number, and arrangement with respect to one another. Although this defect was studied in detail in only one form of spark plug, the deterioration is a characteristic of the material, i.e., of nickel wire, rather than of the particular type of spark plug in which it was noted and studied. The mechanical features of the plugs, however, have a considerable bearing on the time required for the deterioration to become serious. The results of this study of the deterioration of nickel terminals in service should be of value to the makers of all types of spark plugs in which the terminals are made of this metal. CharacteR of Deterioration of WIres During Service Macroscopic Appearance and Properties of Changed Wire.—The nickel terminals in which the deterioration was studied were taken from spark plugs having a central terminal with the ground or side terminals attached firmly at both ends to the shell of the plug. In some of the plugs, two side terminals were used, one on each side of the central one. Although both terminals were found to have deteriorated to some extent, the attack of the central one was quite negligible compared to that of the side ones. These latter wires had developed, in service, transverse cracks that, in many cases, were as sharp and definite as a knife-cut. After a separation occurred, the breach widened by loss of material from the

Jan 1, 1921

By R. P. Marshall

This note describes positive evidence that hydrogen in titanium alloys diffuses under the influence of a thermal gradient. The experiments confirmed the expected similarity of this system to the H-Zr system and produced qualitative information on the magnitude of the heat of transport for hydrogen in titanium alloys. This effect has practical importance in applications where titanium alloys operate in a thermal gradient under conditions such that hydrogen embrittlement may develop within the expected lifetime of the structure, for example in supersonic aircraft service. The theories of thermal-gradient diffusion may be found in recent reviews and experimentation in this field by Shewmon,1 Sawatzky,2 Markowitz,3 and Darken and Oriani.4 The basis for Soret thermal diffusion is the "heat of transport", denoted by Q*. This is an energy term for the amount of heat (energy) transferred by atomic diffusion, and can be either positive or negative signifying a driving force for atomic diffusion to the cooler or hotter region, respectively. Experimentally, Q* is determined by plotting the natural logarithm of the point-to-point concentrations after exposure to a thermal gradient vs the respective point-to-point temperatures in units of the inverse absolute temperature scale. The slope of the straight line is Q*/R, R being the gas constant. For the titanium experiments, specimens (3 by 1 by 1/16 in.) of a) commercially pure titanium, b) the 8 Al-1 Mo-1 V titanium alloy, and c) the 6 A1-4 V titanium alloy were exposed to thermal gradients of 6° to 115°C per in. for 18 days at 250° to 550°C. The temperature gradient was in the direction of the 3-in. dimension. The hydrogen concentrations existing after exposure were determined on sections of the specimens by vacuum extraction. Three sections of unexposed sheet of each alloy were analyzed as standards to allow a check for hydrogen

Jan 1, 1965