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By S. E. Maddigan, B. R. Zimmerman

Most metallurgists are well acquainted with the contributions already made to the study of metals by the use of X-rays. On the one hand, the radiographic method is constantly becoming of increasing importance as a nondestructive testing procedure, while on the other hand the techniques of X-ray diffraction have made important contributions to the fundamental knowledge in studies of phase diagrams, stress conditions, and so forth. A new, and as yet insufficiently appreciated, tool exists in the application of microradiography to alloy structures. This new procedure in many respects bears the same relation to microscopic investigations as does macroradiography to the gross inspection of metal surfaces. However, where macroradiography is used ordinarily to determine the soundness of the metal structure, this new method can be used not only to investigate micro-un-soundness but also to examine the distribution of the constituent elements within the body of the alloy. When combined with the normal procedures of microscopy, this offers a technique of great potential power. It is intended in this paper to present a few examples of copper-base alloys demonstrating the additional knowledge obtained from the microradiographic method. The detection of micro-unsoundness is demonstrated in Fig. I for a sample of cartridge brass as cast. The sample ap- peared quite sound when viewed under the microscope, but in reality was thickly populated with minute voids. This application of the method conforms in every way to conventional large-scale radiography and has been used to facilitate production of metal for vacuum devices.' The true potentialities of the technique are revealed by comparison of Figs, 2a and 2b, which are respectively a photomicrograph and a microradiograph of a cast alloy of 80 per cent Cu, 10 per cent Sn, 10 per cent Pb. The remarkable difference in appearance provides some concept of the additional information this method may yield when used as an auxiliary to microscopic studies. Fig. 2a shows two phases of the copper-tin system plus segregated_ lumps of lead. Fig. 2b, on the other hand, shows with exceptional clarity the actual dendritic growth of the metal crystals during solidification. The practice of microradiography depends upon the laws of absorption of X-rays. This subject has been reviewed thoroughly in recent papers,l,2,3 but for clarity a brief resume will be given here. As is well known, when a beam of X-rays is transmitted through a, layer of homogeneous material, the intensity is reduced according to the relation: where Il is the transmitted intensity, I0 is the incident intensity, µ1 is the linear absorption coefficient of the material, t1 is the thickness of the material.

Jan 1, 1944

By S. E. Maddigan, B. R. Zimmerman

Most metallurgists are well acquainted with the contributions already made to the study of metals by the use of X-rays. On the one hand, the radiographic method is constantly becoming of increasing importance as a nondestructive testing procedure, while on the other hand the techniques of X-ray diffraction have made important contributions to the fundamental knowledge in studies of phase diagrams, stress conditions, and so forth. A new, and as yet insufficiently appreciated, tool exists in the application of microradiography to alloy structures. This new procedure in many respects bears the same relation to microscopic investigations as does macroradiography to the gross inspection of metal surfaces. However, where macroradiography is used ordinarily to determine the soundness of the metal structure, this new method can be used not only to investigate micro-un-soundness but also to examine the distribution of the constituent elements within the body of the alloy. When combined with the normal procedures of microscopy, this offers a technique of great potential power. It is intended in this paper to present a few examples of copper-base alloys demonstrating the additional knowledge obtained from the microradiographic method. The detection of micro-unsoundness is demonstrated in Fig. I for a sample of cartridge brass as cast. The sample ap- peared quite sound when viewed under the microscope, but in reality was thickly populated with minute voids. This application of the method conforms in every way to conventional large-scale radiography and has been used to facilitate production of metal for vacuum devices.' The true potentialities of the technique are revealed by comparison of Figs, 2a and 2b, which are respectively a photomicrograph and a microradiograph of a cast alloy of 80 per cent Cu, 10 per cent Sn, 10 per cent Pb. The remarkable difference in appearance provides some concept of the additional information this method may yield when used as an auxiliary to microscopic studies. Fig. 2a shows two phases of the copper-tin system plus segregated_ lumps of lead. Fig. 2b, on the other hand, shows with exceptional clarity the actual dendritic growth of the metal crystals during solidification. The practice of microradiography depends upon the laws of absorption of X-rays. This subject has been reviewed thoroughly in recent papers,l,2,3 but for clarity a brief resume will be given here. As is well known, when a beam of X-rays is transmitted through a, layer of homogeneous material, the intensity is reduced according to the relation: where Il is the transmitted intensity, I0 is the incident intensity, µ1 is the linear absorption coefficient of the material, t1 is the thickness of the material.

Jan 1, 1944

By Earl Bardwell

THE determination of suitable and safe annealing temperatures is one of the most important problems arising in the operation of a copper rolling mill. Certain of the larger mills have worked this problem out for themselves and in these mills the operation has become more or less standardized. In some of the mills pyrometers are used to measure the temperature of the annealing furnaces and careful watch is kept to see that overheating does not take place. In a considerable number of the mills, however, dependence is placed entirely on the judgment of the operator. The almost total lack of literature bearing on this subject seems sufficient warrant for the publication of this paper. The discussion of this subject naturally divides itself into two parts: (1) The effect of annealing temperature on physical properties; and (2) changes in the microstructure of refined copper due to cold rolling and annealing. I. THE EFFECT OF ANNEALING TEMPERATURE ON PHYSICAL PROPERTIES For the purpose of investigating this phase of the subject six test bars were selected, varying in oxygen content between the limits of 0.036 and 0.070 per cent. The chemical composition of each of these test bars was as follows:

Jan 8, 1914

By Stephen Royce

[Stephen Royce.. I .......................... E. S. Moore.. 9 .......................... E. L. Bruce.. 10 ......................... J. E. Hawley.. 11 ........................ L. M. Scofield.. I 2] STEPHEN ROYCE, * Crystal Falls, Mich.Having examined and reported upon the Steep Rock Lake property and vicinity on several occasions for clients, and having kept in touch with the operation from 1938 on, the writer offers the following discussion of the valuable paper by Messrs. Roberts and Bartley, which was so kindly prepared in response to the writer's earnest request, made as a member of the Committee on Mining Geology. ORIGIN OF THE ORES The writer wishes to express complete agreement with the authors that the origin of these deposits was replacement action by hot ascending sulphide solutions of igneous origin, gaining access via the limestone and dolomite breccia channels, and acting mainly to replace the dolomite and limestone breccia with hematite and hydrated hematite. Some replacement in adjoining greenstones and tuffs by the same solutions may also have taken place, but the latter occurrences may represent replaced iron formations originally bedded in with the volcanics. Iron formation in the same relationship to volcanics is extensively developed on the western Menominee Range of Michigan, in the Middle Huronian belts; also appears in similar relations on the Vermilion Range; and is abundantly exposed in the same stratigraphic position no farther away than the southeast arm of Steep Rock Lake itself. Furthermore, the writer's classification of drill cores, made June 11, 1938, on the prop¬erty, contains the following, for hole 8, in the original strike of the A ore body: "332-342 ore; 342-350 much-leached sugary chert with thin ferruginous slate bands." This is a typical iron formation. The ore in this first group of holes (Numbers 6 to 13, incl.) is described in the writer's 1938 notes as follows: "The ore is hard, massive hematite with some limonite, and shot through with many small vugs." The breccia forms in the float ore and in the drill samples were also noted. As a field classification this does not check badly with Professor Schwartz's microscopic work. Notes supplied by Mr. Neeland, then geologist at the exploration, show: "Hole 23 has some iron formation, not logged yet." It appears therefore that there is some iron formation at the exploration, much iron formation in similar stratigraphic position in the vicinity, and iron formation similarly located as to volcanics on the Vermilion and western Menominee Ranges. Certain of the holes, notably Holes 25 and 27, showed some specular hematites resembling those known to be formed from cherty iron formation elsewhere. Hence the writer concludes that at least a small part of the Steep Rock ores results from replacement of iron formation. By far the greater part of it is, as the authors say, clearly a replaced limestone and dolomite breccia, with the breccia outlines still preserved in much of the ore, and all of the replacement is by hot ascending waters. The authors' citation of the neighboring Atikokan pyrite, pyrrhotite and magnetite

Jan 1, 1943

By R. V. Smith, E. J. Dewees, R. H. Williams

Before this paper was presented, the terms "f" and "v1/f in flow equations had been called "friction coefficient" and "friction factor," respectively. However, many authors have used the names for the terms interchangeably, with the result that considerable confusion has existed in the literature on flow of fluids. In an effort to establish a clear and concise terminology for these terms, the authors have designated f as "resistance coefficient" and dm as "transmission factor." Transmission factors for the flow strings of gas wells were determined by tests on wells in six fields of Oklahoma, Texas, and Louisiana. Data were collected on 14 gas wells while gas was flowing through 1 1/4, 2, 21/2, and 3-in. tubing and 51/2 and 7-in. casing. The transmission factors developed from these data were found to follow the resistance equation published by Nikuradse1 for turbulent flow in "rough" pipes with an average absolute roughness of 0.000,65 in. This average absolute roughness is believed to be representative generally of the flow strings of gas wells with the exception of those that have been exposed to corrosion or other conditions causing undue roughness. Also, use of an absolute roughness of 0.000,65 in. probably will result in transmission factors that are more representative of gas-well conditions than those in common use today, provided that the transmission factors are calculated by the methods outlined in this report. The authors believe that use of these transmission factors will permit calculation of more accurate subsurface pressures in gas wells than those in use today. Engineers use transmission factors in flow equations to compute subsurface pressures in flowing gas wells from measured wellhead pressures. These subsurface pressures are necessary for determining the productivity of gas wells from "back-pressure" tests. The methods of calculating pressures in flowing gas wells, as outlined by Rawlins and Schellhardt2 in Bureau of Mines Monograph 7, are based on the transmission factors in the well-known Weymouth formula" Natural gas engineers have realized that errors introduced by using transmission factors as given by the Weymouth formula are relatively unimportant in testing low-capacity gas wells; but they also know that such factors are important considerations in testing large-capacity wells. Accordingly, the research described in this report was undertaken to fulfill the need for more accurate means of calculating subsurface pressures in gas wells. This report is part of the Bureau's study of the flow of natural gas through pipe and commercial pipelines. Several of the conclusions regarding flow in the flow strings of gas wells were influenced by the results of the study of flow in horizontal pipelines. These instances are mentioned in the text of this report. METHODS OF COMPUTING TRANSMISSION FACTORS Methods and procedures for computing transmission factors for turbulent flow in circular pipes from test data taken on gas wells have been derived and explained in detail in a paper by Smith4. As these methods were used to calculate the transmission factors presented in this report, a summary is included. Starting with

Jan 1, 1955

By W. T. Cardwell

Physically absurd, triple-valued saturations appear in the straight-forward solution of the Buckley-Leverell equations for the displacement of oil by water or gas. From an engineering viewpoint, the triple value causen no difficulty. It is well known how it may be compensnted in order to obtain physically meaningful, numerical results. From a scientific viewpoint, the question still arises: What did the triple value mean? This paper explains how and why the triple value arose in non-capillary Buckley-Leverett theory. The discussion should serve as a background for the understanding and use of the modern method of characteristics in di.~pkzc(>ment theory. INTRODUCTION Displacement theory was introduced to petroleum technologists in 1 941 by Buckley and Leverett.' Their first equation, which they wrote down without derivation, was equivalent to the following: ?s/?t= - qr/F ?f/?r (1) where S = saturation of displacing fluid (a fraction) r = time q,. = total volumetric velocity of both the displacing phase and the phase being displaced (usually, but not necessarily, assumed to be a constant) F = porosity (a fraction) f - fraction of the total volumetric velocity that is the volumetric velocity of the displacing phase (assumed to he a function of S only x = distance. Buckley and Leverett called Eq. I "a material balance equation". Actually it is derived from both a continuity equation, or material balance equation, and Darcy's law. Buckley and Leverett jumped from their Eq. 1 immediately to their Eq. 2, which was the equivalent of the following: (?x/?t)x = qr/F df/ds = qr/F f'(s) (2) The jump from Eq. 1 to Eq. 2 involves apparently simple mathematics, but in that apparently simple mathematics lies a subtle point that is part of the key to the meaning of the triple value to be discussed here. Eq. 2 may be readily interpreted as saying that a point of constant saturation (on a saturation-vs-dis-tance curve) moves with a constant velocity that is proportional to the total volumetric rate, inversely proportional to the porosity, and is otherwise a function of the saturation itself. So that if one knows the derivativc with respect to saturation of the fractional flow function, f(S), for each saturation, one knows the velocity of each point of the moving saturation-vs-distance curve. The function f(S) can he determined experimentally and its derivative f'(S) can be calculated. It turns out that the experimentally derived f'(S) is not a monotonic function of S, but instead it has a definite maximum and declines away at both high and low saturations. This in turn means, in accordance with Eq. 2, that on a moving saturation-vs-distance curve, ccrtain intermediate saturations will travel faster than the saturations either higher or lower. The result is indicated in Fig. 1. In Fig. 1, the abscissas represent distances along a column of porous medium, which column is assumed to be uniform in cross-section perpendicular to the ab-scissal direction. The ordinates represent fractions of pore space occupied by the displacing phase, which may be either water or gas. It is assumed that the saturation is uniform over all planes perpendicular to the abscissal direction. The saturation is a function only of time and one space variable. Time variation of the saturation is represented by change in shape of a curve such as curve AGJ, which represents the initial saturation-vs-

By D. J. Carney, E. C. Rudolphy

IT has been demonstrated that inclusion size, distribution, and composition affect the machin-ability of resulphurized steels. Merchant and Zlatinl concluded that large sulphide inclusions aided machining by forming a (lubricating) coating on the tool face. Boulger et al.² and Van Vlack³ noted that the size, distribution, and composition of the inclusions in the steel affected the machinability. Steel specimens containing large globular sulphide inclusions usually exhibited excellent cutting properties, while machinability was adversely affected by the presence of numbers of oxide-type inclusions. Consequently a thorough knowledge of all the factors which affect the inclusions in the final product is desirable. Since almost all the inclusions have their origin in liquid steel, it was necessary to begin a study of inclusions in free-machining steels by studying the inclusions and chemical segregation in the as-cast ingot. Very little information is available on the size, distribution, shape, and composition of inclusions in large, capped, free-machining steel ingots, particularly the B1113 grade. Gregory and Whiteley4 made a general study of the inclusions in a small, high sulphur, free-machining steel ingot. Also, numerous authors have described the solidification and segregation characteristics of the four basic types of steel ingots, namely, rimmed, capped,7 semikilled,7,8 and killed7,9,10 ingots. Most of these studies were made with plain carbon or low alloy, low sulphur steel. It was desirable to study not only the ingot but also the change in size, shape, and number of inclusions on rolling an ingot to a bloom and thence to a billet. This procedure was followed and it is hoped that this study may serve the dual purpose of adding to the general knowledge of ingot solidification as well as contributing to the knowledge of the size, shape, distribution, and composition of inclusions from the ingot to the billet in a high sulphur, free-machining steel. Procedure A 12.000-lb 23x35x75 in. slab ingot of the B1113 grade was cast, sectioned, and studied both macro-scopically and microscopically. An adjacent ingot from the same heat and of the same size was rolled to a 77/8x77/8 in. bloom and thence to a 21/2x21/2 in. billet. These various bloom and billet sections were also sectioned and studied macroscopically and microscopically. Sectioning the Ingot: The ingot herein described was obtained from the United States Steel Corp.'s South Works Bessemer Blow No. 0193, a B1113 mechanically capped heat. The 23x35 in. ingot (No. 2) was teemed according to normal procedures and after stripping and transportation to the rolling mill was not placed in the soaking pit but allowed to air cool in an upright position. When completely solidified, the ingot was cut into sections by means of a powder scarfing torch and further sectioned by saw cutting as indicated in Fig. 1. Cut No. 2 (1x10x12 in.) from sections A through H was cleaned thoroughly, macroetched in a solution of 50-50 water and hot muriatic acid and used to obtain a macrograph of a horizontal section from the surface to slightly beyond the center of the 23-in. ingot dimension. Cuts No. 5 and 3 (lx81/2xl0 in. each) from sections A through H were treated in a similar manner to obtain a macrograph of a horizontal section from the surface to slightly beyond the center of the 35-in. ingot dimension. The composite macrograph of these horizontal ingot sections, which shows a vertical section of the ingot from top to bottom, is shown in Fig. 2. It should be noted that sections No. 2 are normal to sections 5 and 3 in the composite. Drillings for chemical analyses were obtained from selected positions within the above-mentioned ingot sections as noted in Fig. 3. The oxygen content was determined by the vacuum-fusion method. Samples for microscopic examination were cut from

Jan 1, 1954

By M. L. Holzworth, Philip R. Sperry, Paul A. Beck

INTRODUCTION THE basic work of Z. Jeffries1,2,3 has long ago established the main features of grain growth in the presence of a dispersed second phase. Working with sintered specimens of initially fine grained tungsten, to which various amounts of thoria have been added, Jeffries found that grain growth was inhabited, that is, practically prevented, up to a certain annealing temperature. When this temperature, which increased with the thoria content, was exceeded, extremely large grains developed abruptly from the fine grained matrix. Such a "coarsening temperature" was later found by Gross mann13,14 in certain types of steel, while Bain showed" that in other steels grain growth was gradual, without inhibition and abrupt coarsening. Later work16,17,18,19,21 demonstrated the connection between coarsening and aluminum additions. The viewpoint that the direct cause of inhibition and coarsening in aluminum killed steels is a fine dispersion of aluminum oxide21 appears to be held quite generally, although some doubts have been voiced even very recently.20 In steels with Ti additions the titanium carbide phase is considered responsible for the inhibition effects found.20 The coarsening temperature increases, with the titanium content. That coarsening may occur, as a result of certain heat treatments, even in low carbon rimmed sheet steel, was shown by Samuels,22 He also detected coarsening in a steel ingot of similar composition, after giving it the same heat treatment as that used by him for the sheet material, In this instance the identity of the inhibiting phase has not been determined, However, Tangerding found 23 that even the few small carbide particles, which occur in carbonyl iron, have a very marked inhibitive effect. As a result of an oxidizing anneal at 850°C, large grains begin to form at the surface of the carbonyl iron specimen, and gradually grow inward as the subsurface oxidation of the carbide particles progresses and eliminates the obstruction, The carbide particles can be removed also by annealing in hydrogen, with similar results. But if the oxidizing anneal of 10 hr at 850°C in air, is followed by annealing in hydrogen for 75 hr at the same temperature, grain growth is more spectacular; usually the whole specimen is transformed into a single crystal (approx. 1,5 cm X 5 cm). The reversed procedure, that is, hydrogen annealing followed by oxidizing annealing leads to smaller grains. Tangerding arrived at the logical conclusion that there must have been in his specimens some additional obstruction to grain growth resulting from a minor impurity, other than carbon. It seems likely that this obstruction was caused by an oxide formed during the oxidizing anneal and gradually reduced during the subsequent hydrogen-annealing. Other instances of grain growth inhibition of varying severity, associated with a dispersed second phase, have been described by several authors. Notable examples: lead with 0.06 pct Cu or Ni,24 cartridge brass with 0.12 pct Cr,25 or with up to 0.15

Jan 1, 1948

By Robert D. Pehlke, Robert G. Blossey

The solubility of nitrogen in liquid iron alloys containing chromium and nickel has been measured in the temperature range 1550° to 1700°C at nitrogen pressures to 1 ah. The solubility surface has been determined for the iron corner of the liquid metallic ternary to 21 pct Cr and 11 pct Ni. The nitrogen solubility increases markedly with increasing chromium concentration. The effect of nickel and the bilinear effect of chromium and nickel are less pronounced. The solubility data are presented, and interaction coefficients are determined from them. THE solubility of nitrogen has been determined in liquid Fe-Cr-Ni binary and ternary alloys. The work in this laboratory has been directed toward a relatively limited composition range, that is to 20 pct Cr and 10 pct Ni. Humbert and Elliott1 studied the entire ternary; but their investigation was, understandably, limited in this region. The region considered here is of the greatest commercial interest. Knowledge of the solubility in these alloys is important for accurate control of the nitrogen levels in various stainless alloys. Also, the solution behavior is of theoretical interest because of its deviations from ideal dilute solution rules at moderate chromium levels. The nitrogen solubility in the iron binaries has been reviewed extensively by Pehlke and Elliott,2 particularly with respect to the dilute solution interaction parameters. More recently, there has been some interest in accurate descriptions of solute interactions in non-dilute solutions. Most of the investigators have adopted a power series fit to accommodate any non-linearity of the data. The Taylor series expansion is adopted here, using Scimar's notation3 for the interaction coefficients. An alternate method, instead of requiring higher-order expansions, is to restrict the range of the composition variables. A planar solubility surface may be fitted to the data and its range of validity determined. If the axes of this region are displaced from the usual origin—the solvent pure iron—a change in standard state is required for the solution reaction. The standard state for nitrogen in iron is commonly taken as the hypothetical 1 wt pct solution of nitrogen in pure iron. However, for an alloy of appreciable solute concentration, e.g., 13 pct Cr or 18 pct Cr-8 pct Ni, it would be convenient to use that base composition as the solvent and redefine the solution reaction in terms of that particular solvent. A brief derivation follows, defining the free energy of solution in an alloy solvent. Following the treatment by Darken and Gurry,4 for a nitrogen pressure of 1 atm: N (pure component) = N (1 wt pct in solvent) [l] AG=RT ln100/ wt pct N in solvent Taking this reaction for an alloy solvent less the same reaction for a pure iron solvent the relation between nitrogen in pure iron and an alloy is: N (1 wt pct in Fe) = N (1 wt pct in alloy) [2] AG = RT[ln(wt pct N in Fe) - ln(wt pct N in alloy)] Adding ½ N2(g) to each side and subtracting N (1 wt pct in Fe) = ½N2(g), the total free energy for solution of nitrogen in an alloy, referred to nitrogen at infinite dilution, is given by: ½N2 (g) = N (wt pct in alloy) L3] AG°3 = -RT ln(wt pct N in alloy )PN =1 atm Interaction parameters have their usual meaning when ?G°/RT is differentiated with respect to the proper composition coordinate. If desired, the dilute alloy treatment may be utilized with the new reference composition as the solvent. This has. indeed, been done by Small and pehlke5 in their analysis of nitrogen solution in liquid 18-8 alloys. The Sieverts' technique was used for this investigation on an apparatus previously described.6 The materials used in this program were: Ferrovac-E iron, 99.95 pct, Crucible Steel Co.: chromium, 99.95 pct, Union Carbide Metals Co.; nickel, 99.9 pct, International Nickel Co. Recrystallized alumina crucibles were used with no evidence of crucible-melt reaction. RESULTS The locations of the ternary compositions investigated are shown in Fig. 1. The data of Turnock7 and small8 have been alluded to in other papers.5'6 There were no departures from Sieverts' law observed for

Jan 1, 1969

By Alexander L. Field

This paper contains an account of experimental results obtained in an extensive investigation of steels containing zirconium. There recently appeared' an announcement by F. M. Becket, at whose suggestion and under whose general direction this investigation was made, a statement of the outstanding features of the work and a brief explanation of the factors that led to its inception. It is not possible, in a paper of this character, to record all of the operating and test data collected, which includes the production of more than 350 heats of steel. The difficulty of a complete presentation is increased by the fact that several lines of experimentation are too incomplete to warrant publication of the results at the present time. In addition to the methods of manufacture and test employed in the work, there is described experimental evidence on which are based certain definite conclusions regarding the remarkable effects brought about by the presence in steel of relatively small percentages of zirconium. When added to a molten steel bath, zirconium combines chemically with dissolved oxygen, nitrogen, and sulfur, in the order named, and is capable of neutralizing the embrittling effect of phosphorus in whole or in part; although in the case of phosphorus there is no evidence so far of actual chemical combination. These effects arc not in the nature of suppositions supported by theory, nor is their existence deduced merely on the basis of analytical evidence. In each phase of its relation to or reaction with the four impurities mentioned, zirconium affects beneficiauy the mechanical properties of the steel to which it is added, in a manner which it is one of the purposes of this paper to explain in terms of a tentative working theory.

Jan 1, 1923

By Alexander L. Field

This paper contains an account of experimental results obtained in an extensive investigation of steels containing zirconium. There recently appeared' an announcement by F. M. Becket, at whose suggestion and under whose general direction this investigation was made, a statement of the outstanding features of the work and a brief explanation of the factors that led to its inception. It is not possible, in a paper of this character, to record all of the operating and test data collected, which includes the production of more than 350 heats of steel. The difficulty of a complete presentation is increased by the fact that several lines of experimentation are too incomplete to warrant publication of the results at the present time. In addition to the methods of manufacture and test employed in the work, there is described experimental evidence on which are based certain definite conclusions regarding the remarkable effects brought about by the presence in steel of relatively small percentages of zirconium. When added to a molten steel bath, zirconium combines chemically with dissolved oxygen, nitrogen, and sulfur, in the order named, and is capable of neutralizing the embrittling effect of phosphorus in whole or in part; although in the case of phosphorus there is no evidence so far of actual chemical combination. These effects arc not in the nature of suppositions supported by theory, nor is their existence deduced merely on the basis of analytical evidence. In each phase of its relation to or reaction with the four impurities mentioned, zirconium affects beneficiauy the mechanical properties of the steel to which it is added, in a manner which it is one of the purposes of this paper to explain in terms of a tentative working theory.

Jan 1, 1923

By T. B. Massalski, Horace Pops

The occurrence and the temperature dependence of the athermal martensitic transformation in bcc Cu-Zn ß-phase alloys have been studied by cold-state microscopy, differential thermal analysis, and electrical -resistivity measurements. CinenzatograpJzic studies have shown that in the majority of alloys the transformation occurs in two stages. A thermoelastic phase, having a 'needlelike" uppearance, forms initially at the MS temperature and grows relatively slowly, chiefly in the edge (lengthwise) direction. With further cooling, an additional martensitic phase forms in rapid bursts at a lower temperature, MB, Both martensitic phases appear to have the same cry stallographic habit hut their modes of formation and growth are different. For the composition range studied, between 38.55 and 41.49 wt pct Zn, the habit planes determined are close to the (2.17 ,12) plane of the matrix, in general agreement with the prediction of the phenomenological theory of martensite formation. Thermal cycling through the MB temperature produces permanent plastic deformation in the 0 matrix and gradutally eliminates the typical burst phenomena. Thermal cycling only through the MS temperature produces no noticeable deformation effects. THE bcc Cu-Zn ß-phase alloys tend to become unstable at low temperatures and undergo a martensitic transformation,'-3 but the details of the transformation have not been completely determined. Following earlier metal log raphic and X-ray work1 Tichener and ever' determined by means of electrical-resistivity measurements the temperature range of the martensite transformation during continuous cooling. The transformation temperature was found to decrease with increasing zinc content, becoming about -130°C for an alloy containing 40.04 wt pct Zn. Hull and Garwood3 made a metallo-graphic study of Cu-Zn 0 brasses at temperatures down to -160°C, and determined the habit plane of the transformation product for an alloy containing 38.88 wt pct Zn. No metallographic studies below liquid-nitrogen temperatures have been reported. The martensitic transformation in ß-brass alloys can also be induced by deformation.1,4 The MD temperature corresponding to the appearance of the strain-induced martensite is some 300°C higher than that observed without strain, and hence the transformation can be strain-induced in alloys containing a higher percentage of zinc. Massalski and Barrett4 found that alloys containing as much as 51.89 wt pct Zn could be transformed into a faulted hcp structure by cold working at liquid-helium temperature. Because the details of the MS temperature and morphology reported in the literature are somewhat incomplete, a further study has been made of some of the above features. A design of a cryogenic stage for optical metallography5 has permitted the examination of specimens at temperatures down to 4.2oK, and the use of a cinecamera has permitted a study of the morphology on cooling and heating. The particular aim of the present work was to examine the transformation temperatures, the habit planes, and the growth morphology over a wide composition range in the ß-phase field. The effect of thermal cycling upon the subsequent transformation was also examined.

Jan 1, 1964

By H. E. Cline

Mechanical properties of the eutectic between CoAl and cobalt were measured over a range of- temnperatures and strain rates for a variety of microstructures produced by directional solidification and by thermo-mechanical processing. Directional solidification led to rodlike, lamellar, and irregular microstructures. The unusually high volume fraction of the cobalt-rich rods and the lurge spacing of the rods were explained by the phase diagram. The hot-worked structure consisted of fibers of COAL in a cobalt-rich matrix. The roonl- tevlperature strength of the uvrought material increased with decreasing grain size but the 1000°C strength decreased with decreasing pain size. At high temperatures the directionally solidified tnaterial was stronger and less strain-rate-sensitice than the hot-rolled material. The fine-grained hot-worked ma-terial became superplastic at high temperatures, with tensile elongations greater than 850 pct, while at root temperature this material was ductile and impact-resistant because of the ductile matrix. Fracture occurred in the directionally solidified material at elevated temperatures by inter phase separation and at roorn temperature by cracks in the intermetallic phase. Growth of these cracks was impeded by the ductile cobalt-rich phase. It was found that the CoAl intermetallic remains ordered at elevated temperatures. DIRECTIONAL solidification of eutectic alloys has been used to produce structures consisting of parallel rods or lamellae. Although the spacing of the phases depends on the growth rate, it has been observed to be of the order of 1 µ in many eutectics. A strong intermetallic phase has been used to reinforce a ductile matrix in the rod eutectic A1-Al3Ni and the lamellar eutectic Al-Al2cu.1 The microstructure of the eutectics A1-A13Ni and A1-A12Cu remained aligned after samples of these alloys were exposed to temperatures just below their melting points for long times. Stability of the microstructure at elevated temperatures makes directionally solidified eutectics a candidate for high-temperature applications. Recently Thompson4 has developed the Ni-NiMo eutectic containing 40 pct NiMo lamellae, giving a tensile strength superior to nickel-base superalloys. Unfortunately, this eutectic oxidizes rapidly at elevated temperatures. By mechanically processing a two-phase structure, one may reach strengths as high as 700,000 psi, as demonstrated by drawn pearlite.5 In this case the strengthening mechanism was thought to be related to the fine substructure formed during deformation.5 However, a heavily worked structure may recrystallize and coarsen at elevated temperatures. Some worked eutectics such as Al-Al2Cu,6 Pb-Sn,7 and Sn-Bi7 have shown unusually large elongations in tension when tested at elevated temperatures. This phenomenon of large extension, called "superplasticity", was related to the fine grain size of wrought two-phase alloys at elevated temperatures. A mode of deformation of "superplastic" material has been shown to be grain boundary sliding,' which has also been observed during creep of polycrystalline materials.9 The Co-A1 system was chosen for this investigation after examining the phase diagrams of binary eutectics." The high melting point of 1400°C, the aluminum content which was expected to give oxidation resistance, and the properties of the intermetallic CoAl were the chief factors influencing this choice. This eutectic consists of a mixture of the intermetallic CoAl and a cobalt-rich solid solution. The intermetallic CoAl has a large range of solubility and a CsCl structure." From the similarity between CoAl and NiAl one would expect CoAl to be ordered at elevated temperatures; however, specific heat measurements" show a transition at 800°C. West-brookL2 has measured the hot hardness of CoAl and found a rapid decrease in hardness above 600°C. The tensile properties and stress rupture properties of as-cast CoAl-Co eutectic were measured by Ashbrook and wallace13 who report a room-temperature tensile elongation of less than 1 pct. I) EXPERIMENTAL A) Sample Preparation. Experimental materials were prepared by directional solidification, by hot working, and by powder processing. The starting materials were electrolytic 99.9 pct pure cobalt and 99.99 pct pure aluminum. 1) Directional Solidification. A vertical Bridgman apparatus heated by induction was used to directionally solidify cylindrical ingots $ in. in diam and 5 in. long. The apparatus was first evacuated and then an argon atmosphere was introduced to retard evaporation. The ingots were contained in an alumina crucible inside a graphite susceptor that rested on a movable water-cooled base. The base was lowered out of the induction coil at a constant rate. Eight of the ingots were solidified using a drive rate of 2.5 cm per hr, two of the ingots at 1.2 cm per hr, and one ingot at 10 cm per hr. The eutectic composition Co 10 wt pct A1 was used in all but one ingot which had the composition of the stoichiometric CoAl intermetallic, Co 32 wt pct Al. 2) Mechanical Processing. At 1000°C the eutectic Co 10 wt pct A1 is a mixture of an intermetallic phase, Co 21 wt pct Al, and a cobalt-rich phase, Co 5 wt pct Al.10 These phases differ from the stoichiometric CO 32 wt pct A1 and pure cobalt because of the large solubility of this eutectic at elevated temperatures." Four rectangular slab ingots, 4 by 1 by 12 in. high, of Co 10 wt pct Al, Co 5 wt pct Al, Co 21 wt pct Al, and Co 32 wt pct A1 were cast for hot rolling. The cobalt was first vacuum-melted, H2-treated for 20 min,

Jan 1, 1968

By P. A. C. Wilson

Instantaneous outbursts of carbon dioxide in coal mines have occurred in Germany only in one part of the Waldenburg-Neurode mining district.' This mining region comprises the northeastern fold of the "Innersudetian" coal basin. The basin (Fig. 1) is in the form of a geologic trough approximately 60 km. (37 miles) long and 30 to 35 km. (19 to 22 miles) broad and is enclosed on three sides by ancient truncated mountain ranges—the Riesengebirge, the Bober-Katzbachgebirge and the Eulenge-birge. Only the southeastern part of the trough is not enclosed by mountains of older rock formation. The floor of the basin before the building up of the coal measures was formed on the western, northwestern and northeastern borders by the material from Archean rocks (gneiss and mica schists); in the northern part over this was deposited the Silurian formation. The carboniferous period is represented both by the Lower ("kulm") and Upper Carboniferous, with an unconformity between the stages; that is, the kulm was tilted up and partly eroded before the Upper Carboniferous was deposited. The kulm is a salt-water formation (with marine fauna) but Upper Carboniferous rocks are nonsaline water deposits, with fresh-water flora only. The Lower Silesian coal measures were thus formed in a lime basin fed by the running water coming down from the mountains. As already indicated, the bottom series of the Upper Carboniferous formation, termed "Waldenburg strata" were laid down unconformably on the Lower Carboniferous. Over the Waldenburg strata there are the "Weisstein strata," an intermediate barren stage without coal beds, 400 to 500 m. (1300 to 1600 ft.) thick. Another unconformity followed, which is overlain by the '(Saarbrucken" series, which contain the upper coal measures, and the ('Ottweil" strata. In the last named, it is only on the southwestern fold that beds of coal are developed.

Jan 1, 1931

By John T. Eash, Clair Upthegrove

During recent years nickel has had an increasingly important role as an alloying element in the copper-tin bronzes. Nickel additions not only produce better casting alloys but also make alloys whose properties can be varied by precipitation hardening; thus leading to increased commercial application of copper-nickel-tin alloys. Recent investigations have indicated a decrease in the solubility of tin in the alpha phase as the annealing temperature is lowered and as the nickel content is increased within certain limits. A further change caused by the nickel additions is produced in the excess phase of the copper-tin bronzes. The alpha plus delta eutectoid is replaced by a clear precipitated constituent in the alloys containing up to 15 per cent tin and not less than 1 per cent nickel. This paper has a dual purpose. The first is to present a new determination of the alpha-phase boundary in alloys containing from 0 to 20 per cent nickel, together with the liquidus and solidus temperatures above that field. The second purpose is to give the equilibrium conditions existing in alloys exceeding the alpha phase containing up to 31 per cent tin and 5 per cent nickel, so that the disappearance of the alpha plus delta eutectoid constituent may be explained. After considering the earlier investigations of the related alloy systems and the experimental procedure used in this investigation, the alloys will be discussed in groups as separate quasibinary systems containing constant amounts of nickel. The systems considered contain 0, 1, 2, 3, 5, 10 and 20 per cent nickel. The straight copper-tin system was reinvesti-gated in order to establish a foundation on which to build the remaining material. The 1 per cent nickel alloys studied give the relations existing in the region of the eutectoid transformation between the copper-tin system and the 2 per cent nickel copper-tin system which has been investigated in detail. The thorough examination of the latter system was made since 2 per cent of nickel is representative of the commercial additions to the cast bronzes and also because it is sufficient to produce

Jan 1, 1933

By David J. Mack

The structures and properties of the copper-aluminum alloys have been the subject of much study since the classic investigation of Carpenter and Edwards1 focused attention on. the engineering utility of these alloys. It was recognized at an early date that the metallographic structures developed in the aluminum bronzes were similar to those developed in steels, both in the annealed and the rapidly cooled state. Most investigators have been concerned with the acicular, martensitic-like structure formed from the 0 solid solution upon rapid quenching; the transformation of the ß under equilibrium cooling to a lamellar eutectoid having been relatively neglected. With the introduction in 1930 by Davenport and Bain2 of the isothermal transformation technique for studying eutectoidal decompositions, a new field was opened. Although an enormous amount of work has since been done on the isothermal transformation of steel, study of structurally analogous systems has been almost totally overlooked. The outstanding exception was the prize-winning paper† of Smith and Lind1ief³ ho investigated the decomposition of the 0 phase in copper-aluminum alloys by isothermal methods. This was followed in 1934 by Wasserman's review' of available information on analogous transformations in eutectoid alloys. Since that time no comprehensive study of isothermally transformed eutectoids analogous to steel has appeared, although many valuable contributions have been made to an understanding of the structures developed in such systems. Important papers have been published by Kurdjumow, Gridnew and co-workers, Obinata,18.'9 Greninger,20.21 and Others. The work to be described in this paper was an outgrowth of preliminary studies on the isothermal transformation of a entectoid aluminum bronze, after it became apparent that the alloy under study was reacting somewhat differently than the similar alloy used by Smith and Lindlief. Preliminary Experiments The material used in this study was a high purity aluminum bronze especially prepared by Ampco Metal, Inc. It analyzed: Copper—88.07 pet, Aluminum—11.89 pct, Iron-—0.02 pet, Manganese—-.OI pct, Others—Balance. Although some disagreement exists on the exact composition of the eutectoid, this alloy was believed to be of essentially eutectoid composition even though some pro-eutectoid* delta particles existed in the microstructure of furnace cooled specimens. Specimens transformed isothermally at temperatures slightly below the eutectoid showed relatively large amounts of pre-eutectoid delta, but as will be shown later, this results

Jan 1, 1949

By AIME AIME

0 N March 5, at Carnegie Hall, Cleveland, the Ohio Section held a joint meeting with the Cleveland Engineering Society, and the local sections of the American Chemical Society, American Society of Mechanical Engineers, and the American Society for Steel Treating. Prior to the meeting there was an informal dinner attended by 50 men. Then followed an address, "The Principles of Corrosion Protection," by Ulick R. Evans, the well-known British specialist in this problem. There were 266 present, presided over by Zay Jeffries, president of the American Society for Steel Treating and a member of the Cleveland Engineering Society. The address was well received and some discussion followed

Jan 1, 1929

By Charles Bohdanowicz

The accompanying tables' show that in 1936 production of crude oil in Poland decreased very slightly (Tables 1 and 5), by about 1 per cent, and that gas production increased by about 11 per cent. As in preceding years, drilling (Tables 1, 2 and 3) was devoted principally to extension of oil fields in the Jaslo district (fields 1, 5 and 7) and, in a lesser degree, of those in the Stanislaw6w district (field 27). Production of crude oil is being sustained in a large measure by the deepening of old wells; besides a total number of 183 new wells completed (Table 3) about 152 wells were deepened, and in a general way the results of deepening were better than those of new wells drilled. Reserves of "proved land" have not increased as far as crude oil is concerned, but a notable extension of gas fieldsZ was attained in Daszawa (field 25) and in the Roztoki area (field 5). In the sub-carpathian zone gas fields have been proved in the vicinity of Kalusz, Kos6wJ and in the northwestern sector of that zone, between Tarnow and Lezajsk. The gas horizons of Daszawa occur in the Upper Tortonian, those of Kos6w in the Lower Tortonian (Middle Miocene). Exploratory drilling has been performed near Kalusz, and has been commenced near Kos6w and in the Tarn6w area. The Kos6w gas approaches the "wet" type. A possible gas reserve with a volume of 4,000,000,000 cu. m. (141,000,-000,000 cu. ft.) has been calculated, for an area of about 5 sq. km. of the Daszawa field, on the basis of production obtained and of decreases in pressure and their distribution. The fields of Roztoki and Strachocina (No. 12) have not yet supplied sufficient data for a calculation of reserves, which might possibly be estimated at 2,000,000,000 cu, m. (70,000,-000,000 cu. ft.).

Jan 1, 1937

By Henry Knox

AMIDST the tumult and the shouting of the times, there are perceptible certain tendencies, quiet though serious, to take stock of the situation with a view to putting our house in order, and of these none is more encouraging than the awakening, interest in the higher education, and especially in technical education. Various committees of the professional organizations have cast much light on the matter and have pointed the way to progress and reform, but their scope is necessarily restricted by the requirements of precision and any excursion beyond these bounds must be undertaken by the unfettered individual who is free to rush in where responsible committees fear to tread. A discussion of the bearing of the higher education on character scarcely lends itself to objective treatment and yet, fraught as it is with imponderables and abstractions, the theme cannot be escaped in any com-prehensive scheme of education. The question is being pressed with growing insist-ence, why these United States, relatively to Europe, are deficient in intellectual output, and the ready reply, "Because America is a new country and Americans are too intent on developing their natural resources," is too facile and too false to carry conviction.

Jan 12, 1922

By James Gayley

The development of blast-furnace practice in America in the direction of large yields is mainly the history of our working since the year 1880, as the advancement that has been made in the last decade is greater than that in the third of a century previous. A new era in the manufactnre of pig-iron began in 1880 with the putting in blast of the Edgar Thomson furnaces. These furnaces at once leaped to the front as pig-iron producers, and have maintained that position—with but one brief interruption—ever since. I shall, therefore, confine myself mainly to a description of these works, showing the changes in design and practice by which these results have been achieved. In order to show more clearly the progress that has been made since 1880, I shall refer briefly to the best work that was done in the ten years previous to that time. The Struthers furnace, in Ohio, was one of the first to attract attention in the matter of large outputs. This furnace was 55 feet high, 16 feet in diameter of bosh, 9 feet in the hearth, and about 8 feet 6 inches at the stock-line. The fuel was raw coal. This furnace, in December, 1871, made 1602

Jan 1, 1891