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By Ancel St. John

DURING a session on coal and coke at the February, 1926, meeting of the American Institute of Mining and Metallurgical Engineers, the writer called attention to the important work on the X-ray analysis of coal1,2 being carried on in England by C. Norman Kemp and his associates. This aroused so much interest that a paper on the subject was requested for the October meeting. It has seemed advisable to make this paper a preliminary survey of the present status in the application of X-ray methods to coal problems and to follow it from time to time with further articles as the art develops. A survey of this sort naturally divides itself into two parts, radioscopic analysis of the distribution of constituents through the body of the coal, and diffraction analysis of the state of association of the elementary components. In this article the subject matter concerning radioscopic analysis has been drawn almost entirely from the writings of Kemp and his associates and the illustrations have been specially furnished by him. It is a pleasure to record here a great indebtedness for this courtesy and cooperation. The subject matter and illustrations concerning diffraction analysis are taken from the author's unpublished researches.

Jan 10, 1926

By H. Doi, Y. Fujiwara, K. Miyake

In order to throw light on the mechanism of plastic deformation of WC-Co alloys, compressive tests of WC-(7 to 43) vol pct Co alloys have been carried out at room temperature, and stress-micro strain relation has been investigated in detail. The analysis of the factors affecting the yield stresses reveals that the yield stresses can be predicted by modified Oro-wan's theory if one properly estimates the planar in-terfiarticle spacings. Conzpressive straining of some of the alloys by 0.066 to 0.17pct increases the decrements by a factor of as much as 3.4 to 14, whereas the corresponding increase in the electrical resistivities is less than 10 pct. The analysis of the decrement data in terms of -Gramto and Lücke theory shows that the marked increase is attributed to increased dislocation darnping itt the binder (cobalt) phase. By cornbilling the decrement data and the conzjwession duta, one obtains the relation between flow stress in shear (?t) and increase in dislocation density (p): At = const . v6 . This is interHeted to mean that the mechanism of strain hardening of CirC-Co alloys is essentially sarne as the one for dispersion strengthened alloys. The possible effect of bridge formations between the carbide particles has also been examined. OWING to the combination of hardness, strength, and other physical and chemical properties, WC-Co alloys have opened the way for unique fields of applications, the recent ones being, for instance, anvils for super-high-pressure generation apparatuses. In such applications, the alloys are frequently subjected to very high compressive stresses: these stresses may cause the alloys to deform plastically and eventually to fail. However, much remains obscure regarding the nature of the plasticity of the alloys. Evidently, the alloys owe their high strength to the hard carbide particles which frequently occupy as much as 80 to 90 pct in volume fraction, whereas the ductility required for practical applications is provided by the small amount of the binder phase between the carbide particles. When the volume fraction of the carbide phase is not very large, deformation behavior of the alloys may be described by some of the current dispersion strengthening theories. However, greatly increasing the carbide phase is thought to lead to some carbide skeleton structure or bridge formations owing to the increased chances for direct contacts between the carbide particles;1,2 this may appreciably affect the plasticity of the alloys. Regarding the effect of formation of the carbide skeleton structure, it is interesting to note the work by Ivensen et al.3 on compression tests of the alloys containing somewhat large carbide particles; they observe extensive generation of slip bands in the carbide particles after application of some preliminary compressive stresses. They interpret the results in terms of plastic deformatiot: of the carbide particles which are supposed to have formed a skeleton structure; the binder phase plays only a passive role, at least in the early stages of the deformation. That carbide crystals exhibit microplasticity at room temperature is apparent from the work of Takahashi and Freise4 and French and Thomas5 on indentation of WC single crystals. On the other hand, Dawihl and coworkers6-10 maintain that even when volume fraction of the carbide phase is very large (for instance, more than 90 pet), a very thin binder layer generally exists between the carbide particles. They interpret the results of the extensive mechanical tests in terms of the plasticity of such a layer. Gurland and Bardzil11 point out that decrease in ductility of the alloys with increase in the carbide phase is caused by the effect of plastic constraint exerted by the dispersed carbide particles. Drucker12 further develops this concept from a continuum-mechanics approach on an assumption that a continuous thin binder layer separates the carbide particles. A common feature of the studies reported so far on the plasticity of the alloys is that the information deduced is invariably qualitative in nature. Thus, very few systematic experiments for obtaining reliable and sufficiently detailed stress-strain curves of the alloys varying widely in the microstructural features have been carried out. In particular, it may be of special interest to investigate in detail the early stages of the plastic deformation of the alloys in order to shed light on the strengthening mechanism. However, such work appears to be extremely rare. Doi et al.13 recently reported a first brief account of the results of some quantitative analysis of the plasticity of the alloys in terms of dislocation theory. Their experiment was rather limited in the composition range covered (volume fraction occupied by the carbide phase: 79 to 83 pct), and thus they could not necessarily elucidate the controlling mechanism of plastic deformation of the alloys of a more general composition range. Consequently, in the present investigation, deformation behavior and some other physical properties of the alloys were investigated and discussed in more detail over a much wider composition range. SPECIMEN PREPARATION WC-Co alloys used in this experiment were prepared in cylindrical or rectangular form by sintering in vacuo compressed mixtures of tungsten carbide and cobalt

Jan 1, 1970

By R. E. Ogilvie, A. R. Kaufmann, S. H. Gelles

The solid-solubility limits of iron in beryllium were determined between 850o and 1200oC by analysis of differential type multiphase diffusion couples, using an X-ray absorption technique. The maximum value of the solubility limit was found to be 0.92 ± 0.02 at. pct (5.46 wt pet) at the eutectic temperature 1225°C. The solubilities of nickel and beryllium were determined between 900°and 1200°C by the same technique and the maximum solubility was found to be 4.93 + 0.01 at. pct (25.2 wt pet) at the eutectoid temperature, 1065°C. A previously unreported high-temperature phase which decomposes eutectoidally at 1065 °C was found to exist in the beryllium-nickel system at a composition of approximately 8 at. pct Ni (36 wt pet) by diffision-couple analysis. The presence of this phase was confirmed by thermal analysis and metallo-graphic analysis of the structure resulting from the eutectoid decomposition. G. V. Raynor1 has treated the solid solubilities of some of the elements in beryllium on the basis of the "Hume-Rothery" rules2 which have been modified to include ionic size and ionic distortion effects. It was predicted that the solubility of iron and nickel in beryllium should be slightly less than that of copper. The lowering of the solubility, according to Raynor, is due to a more unfavorable relative valency effect and an ionic size effect. Kaufmann and corzine3 have compiled data on the solubilities of elements in beryllium and have discussed them in the light of the Raynor paper. These authors feel that, because the elements having the greatest solubility in beryllium systematically fall in the Group VIII and IB Columns of the periodic table, the electronic structure greatly influences the maximum solid solubility of elements in beryllium. The solubility of iron in beryllium was determined by Teitel and cohen4 as part of the study of the beryllium-iron phase diagram. The determination was carried out by X-ray and thermal analysis and according to the phase diagram presented, the maximum solubility of iron in beryllium is 0.41 at. pct (2.5 wt pct) at 1225oC. However, it is estimated that the uncertainty in the position of the a-beryllium primary solid-solution boundary is about 0.5 at. pct (3wtpct). Losana and Goria3 in studying the beryllium-nickel phase diagram, determined the solid solubility of nickel in beryllium by thermal analysis. They found the maximum solubility to be between 1.65 and 2.65 at. pct (10 to 15 wt pct) at 1240°C. This value decreased rapidly with decreasing temperature. In determining approximate ranges of solubilities for different elements in beryllium, Kaufmann, et al,8 reported a value of between 1.3 and 1.7 at. pct (7.9 to 10.1 wt pct) for the solubility of nickel in beryllium. The value was obtained by metallographic examination of quenched alloys and lattice-parameter measurements. However, the authors also noted a single-phase structure for a 1.7 at. pct Ni alloy (10 wt pct) on cooling from the liquid. This would indicate a higher solubility range than was reported. ~isch,' in his X-ray studies of beryllium-copper, beryllium-nickel, and beryllium-iron intermetallic compounds, reports the disappearance of a second phase (Ni,Be2) in the beryllium primary solid solution at approximately 4 at. pct (20 wt pct). THEORY The analysis of concentration gradients in diffusion couples has proven to be a useful tool in determining phase equilibria.8-14 In this particular study the diffusion couples were chosen to straddle the expected composition range of the phase boundary, then heat treated at a given temperature and the concentration gradient evaluated. The composition of the phase boundary for a given temperature appears at a point of discontinuity of the composition gradient. Examples of typical phase diagrams and the concentration gradients which should be found in such systems are shown in Fig. 1. In the present work, gradients of the form of Fig. l(c) were obtained in diffusion couples made of pure beryllium and two-phase alloys of beryllium with either iron or nickel. The composition at the point where the gradient becomes discontinuous, Cs, corresponds to the solubility limit of either iron or nickel in beryllium. The analysis of the concentration gradients was carried out by an X-ra absorption method developed and applied by Ogilvie and later used by Moll13 and Hilliard.l4 It depends on the fact that the absorption of X-rays by matter is determined by the concentration and type of the various atomic species present. The relationship for the intensity, I, of a monochro-

Jan 1, 1960

IN the early period of steelmaking, ingot structure and segregation were of no practical importance. Crucible melting required very small ingots that gave little segregation, and a small inserted hot top prevented pipe. As larger ingots were poured from bessemer converters and open-hearth furnaces, segrega- [ ] tion, piping, and gas evolution became increasingly important. The techniques of making semikilled and rimmed ingots were developed with little knowledge of the actual mechanism of the process. The first studies of solidification were made by the Brearleys1 and Gathmann,2- who used stearine to show the effects of shrinkage and mold shape. Fig 11-1 is a picture of a series of stearine ingots at various stages of solidification. These studies showed that the liquid solidifies at a nearly uniform rate from the sides and bottom of the mold and demonstrated the formation of

Jan 1, 1964

By A.J. PHILLIPS, Wm. B. Price

ADMIRALTY nickel is a new corrosion-resisting and heat-resisting white metal alloy composed of 70 per cent. copper, 29 per cent. nickel and 1 per cent. tin. It has been given the trade name "Adnic." In 1924, Win. B. Price1 described this alloy and its successful use as a diaphragm metal. The physical properties of Adnic at ordinary and elevated temperatures were given in Mining & Metallury.2 A patent3 was granted on Adnic, April 28, 1925. Work was first started on this alloy during December, 1921. An investigation of the patent situation in this country and abroad, and also a search of the literature, did not reveal any work on the ternary diagram, copper-nickel-tin; and the only wrought copper-nickel-tin alloy found was that known as "Newloy," composed of copper, 64 per cent.; nickel, 35 per cent.; tin, 1 per cent. This was being manufactured and sold by Wm.. Gallimore & Sons, Ltd., Sheffield, England, for unplated spoon and fork stock, and first came to our attention in January, 1922. The preliminary work in connection with the development of Adnic was started by casting four, binary copper-nickel alloys of the compositions shown in Table 1. Various amounts of tin were added to each of these mixtures, to produce the series of ternary copper-nickel-tin alloys also given in Table 1. A 200-gm. sample of each of these was cast by melting in a clay graphite crucible and allowing the melt to slowly cool with the furnace. A longitudinal section from each of these castings was polished and etched for-examination. The castings were not homogeneous solid solutions, but showed the usual cored structure associated with cast alloys. However, from an examination of these sections we roughly determined

Jan 1, 1928

By C. Baldwin Sawyer

During the last half century, much time has been devoted to investigations of the effect of nitrogen on the physical properties of steel, but in all discussions of results there is considerable doubt as to the accuracy and interpretation of analyses. No conclusions concerning the effects of nitrogen on the physical properties of steel can be more reliable than the methods used for determining the quantities of nitrogen present. The work described in this paper endeavors to establish some additional conditions required by the customary distillation method for reliable and nccurate determinations. Results by a combustion method of analysis for total nitrogen content are presented; these indicate that all of the nitrogen in steel, exclusive of that contained in blowholes, may be determined by the distillation method. Based on these more sound and more reliable analyses, results havc been obtained that bear on the following divisions of the general subject, nitrogen in steel: Nitrification of steel by melting in nitrogen, nitrification of steel by heating in ammonia, decomposition of nitrified steel by heating, and thermal analysis of nitrified steel. Consideration of Two Methods of Analysis and Some Results The details of the distillation and combustion methods are chemical in character and of little interest to most readers. They are therefore described in an appendix. The distillation method is essentially that described by Hurum and Fay1 and was generally used throughout this work. It is capable of the greatest accuracy and all results given are significant to the last, figure. Where analyses do not check, the samples

Jan 1, 1923

By J. Chipman, N. J. Grant, H. L. Bishop, H. N. Lander

Data of several studies on the equilibrium between molten iron and open hearth-type slags have been combined to determine some of the chemical reactions involved in steel-making. Effects of slag composition and temperature on the iron oxide activity, distribution of sulfur between slag and metal, and the carbon content of the metal are discussed. Also, an overall reaction for the equilibrium between sulfur and oxygen is presented. DURING the past 15 years, a number of significant equilibrium studies have been carried out between liquid iron and simple basic and acid slags in an effort to throw light on the chemical reactions involved in steelmaking. This paper is a summary of the experimental reports on investigations which have been conducted at the Massachusetts Institute of Technology, and of related work in other laboratories. The investigations summarized in this paper include the works of Fetters and Chipman1 and Taylor and Chipman,' which dealt with simple slags that are similar to those found in the open hearth. The sulfur determinations of the slag-metal tests of Fetters and Chipman as reported by Grant and Chip-man" are included. Also, the investigations by Win-kler and Chipman' and Grant and ChipmanQ f more complex slag systems are included in this summary. The experimental procedure is described in detail in each of these papers and will not be discussed. Since the previous reports have been somewhat independent, it is desirable to review and combine the data of the various investigations to determine, as accurately as possible, the reactions in the slag-metal systems of the open hearth. It is the object of this report to present the combined data in such a way as to enhance their utility and applicability. Oxidizing Power of Slags of the System (CaO+MgO)-Si0,-FeO in Equilibrium with Carbon-Free lron The slags selected to determine the iron oxide activity in equilibrium with pure liquid iron were those containing less than 2 pct each of P,O,, MnO, or A1,0, on a weight percentage basis. The method used to calculate activities of iron oxide in liquid slags was that developed by Taylor and Chipman.' The solubility of oxygen in molten carbon-free iron under pure iron oxide slags can be expressed by the following eauation —6320 log (pctO)-------------— + 2.734 [1] where T is the absolute temperature. The oxygen solubility values as determined from Eq. 1 were used throughout this study to calculate the iron oxide activity. The iron oxide activities of slags of various iron oxide contents are plotted in Fig. 1 as a function of the basicity, which is defined as the ratio (pct CaO + pct MgO)/pct SiO,, and are based on molecular percentage values of the components. The percent (FeO), is the total iron content in the slag calculated to FeO. Fig. 1 shows that as the basicity increases at constant mol pct (FeO),, the iron oxide activity increases rapidly and reaches a plateau which extends on the basicity scale from about 1.3 to 2.3; the iron oxide activity then decreases rapidly until a basicity of about 3.0 is reached, where it tends to level off. The solid portions of the curves represent the experimental data of Fetters and Chipman and Win-kler and Chipman. Their experimental heats were made in magnesia crucibles, and hence the slags were saturated with magnesia. The broken lines on the left represent the experimental data of Taylor and Chipman, who used a rotating crucible. In general, these slags contained less than 3 pct MgO. The absence of data for slags containing less than about 40 mol pct (FeO), and basicity ratios greater than about 2.5 made it necessary to calculate curves using ionic concepts of slags. The broken lines on the right, which disregard the solubility limits of di and tricalcium silicate and lime, were obtained by following the treatment of Flood and Grjotheim,"

Jan 1, 1957

By E. H. Bunce, H. M. Haslam

The "direct process" for the manufacture of pigment zinc oxide produces the oxide directly from ore. This is accomplished by reducing the zinc by means of carbonaceous fuels and immediately burning the evolved zinc vapor to zinc oxide. For many years this was done by smelting a loose charge of ore and fuel on perforated grates. Since this process was developed in America and was extensively used here, it has become known as the "American process." Alternative methods of producing pigment zinc oxide are the "indirect process " and the "wet process." In the indirect process the raw material is metallic zinc, which is volatilized and burned to the oxide. For a long time this has been done in externally fired horizontal refractory retorts. This process has become known as the "French process," from the country of its origin. The wet process precipitates the carbonate or hydrate from zinc-bearing liquors and then calcines the precipitate. This method has been but sparingly used for producing commercial quantities of zinc oxide. It appears to be advisable to restrict the use of the terms "American process" and "French process" to the methods of the earlier periods and to consider the terms "direct process" and "indirect process" as generic terms to designate, first, all methods of producing pigment zinc oxide directly from ores, and, second, to include all those methods of producing zinc oxide from metallic zinc and metallic zinc-bearing substances. American Process In 1796 John Atkinson was granted British Patent No. 2094/1796 entitled, "Manufacture of White Paint." This patent describes a method of producing zinc oxide in which roasted zinc blende or calcined calamine is mixed with charcoal, and the mixture heated in a muffle to liberate zinc vapor, which is then burned to obtain zinc oxide fume. About 1850 The New Jersey Zinc Co., at its plant in Newark, N. J., attempted to make zinc oxide by smelting, in a reverberatory furnace, a mixture of coal and ore from its Franklin mine. A considerable amount

Jan 1, 1937

By A. M. Gow

THE trend in ball milling has been toward mills of larger diameter, but without fundamental laws as a guide. The speeds at which mills are run have been a matter of cut-and-try. This paper deals with both diameter and speed and proposes to show the laws which apply to some laboratory ball mills, ranging in diameter from 18 to 42 in., run at various speeds. The correlation of power, speed, and grinding rate of the respective mills will be shown by curves of markedly significant similarity. Before presenting the results some consideration will be given to the mechanics of ball milling.

Jan 1, 1930

By P. K. Koh, H. A. Lewis, H. F. Graff

New data on the distribution of silicon between slag and carbon-saturated iron at 1600Oand 1700OC are presented which, in combination with previously published data, permit the determination of silica activities over a broad range of compositions in the CaO-Al2O3-SiO2 system. The distribution of silicon between graphite-saturated Fe-Si-C alloys and blast furnace-type slags in equilibrium with CO has been described in previous publications.1"3 In this past work the silica-silicon relation was established at temperatures of 1425" to 1'700°C for slags containing up to 20 pct A12O3. This paper presents the results of additional studies at 1600" and 1700° C which extend the silicon distribution data at these temperatures for CaO-A12O3-SiO, slags over a range from zero pct Al2O3 to saturation with Al2O3, or CaO.2Al2O3. The upper limit of SiO2 is set by the occurrence of Sic as a stable phase when the metal contains 23.0 or 23.7 pct Si at 1600" or 1700°C, respectively. The activity of silica over the expanded range is determined directly from the distribution data.3 Recently4-7 other investigators have studied the activities of SiO, and CaO, principally in the binary system, using different methods and obtaining somewhat different results. EXPERIMENTAL STUDY The experimental apparatus and procedure have been fully described in previous publications.1, 3 Six new series of experimental heats have been made, four at 1600° and two at 1700°C. Master slags of several fixed CaO/Al203 ratios were pre-melted in graphite crucibles, and these were used with additions of silica to prepare the initial slag for each experiment. Slag and metal were stirred at 100 rpm and CO was passed through the furnace at 150 cc per min. The initial sample was taken 1 hr after addition of slag at 1600°C or 1/2 hr after addition at 1700°C. The run was normally continued for 8 hr at 1600°C or 7 hr at 1700°C, and the final sample was taken at the end of this period. Changes in Si and SiO2 content indicate the direction of approach to equilibrium, and in a series of runs where the approach is from both sides this permits approximate location of the equilibrium line. Fig. 1 shows the results of such a series of 15 runs at 1600°C for slags of CaO/Al,O3 = 1.50 by weight. Figs. 2 and 3 record other series at 1600°C and Fig. 5 a series at 1700°C with fixed CaO/Al0 ratios. The results of the experiments at 162003°C have been reported in part in a preliminary note.3 In the experiments recorded in Figs. 4 and 6, the slags were saturated with A12O3 (or with CaO.2A12O3 within its field of stability) by suspending a pure alumina tube in the melt during the course of the run. The final slag analyses were used to establish the liquidus boundaries8 in the stability fields of CaO.2Al2O3 and of Al20,. ACTIVITY OF SILICA The free-energy change in the reaction has been calculated by Fulton and chipman2 from recent and trustworthy data including heats of formation, entropies, and heat capacities. The more recent determination by Olette of the high-temperature enthalpy of liquid silicon is in satisfactory agreement with the values used and therefore requires no revision of the result which is expressed in the equation: SiO2 (crist) + 2C (graph) = Si + 2CO(g.) [1] &F° = + 161,500 - 87.4T The standard state for silica is taken as pure cristobalite and that of Si as the pure liquid metal. Since the melts were made under 1 atm of CO and were graphite-saturated, the equilibrium constant for Eq. [I] reduces to K1 = asi /asio2. The value of this constant is 1.77 at 1600°C and 16.2 at 1700°C. Through K1, the activity of silica in the slag is directly related to the activity of silicon in the equilibrium metal.

Jan 1, 1960

By Francis Foley

Small cracks in a plane normal to the axis of steels are found to be prevalent around the water hole of drill steels that have been in service for an unknown period of time. Cracks are not found on the exterior surface of the steels nor to any noteworthy extent even in the water holes of new drill steels. Undulations and serrations are found in the surface of the water hole in new steels; it is in the bottom of such defects that cracks appear to develop in used steels. The occurrence of the more serious of these cracks at the point in the steel to which it has been heated for forging or hardening is thought to account to a great extent for failures in that region. DURING the past, year, we have investigated the cause of 88 per cent, of the failures by breakage near the bit end of some 1-in. hollow, hexagon, drill steel used in a metal mine. This breakage in the region of the bit end of the steel was peculiarly interesting because we have been led to believe that breakage is most prevalent at the shank end of drill steels near the place where the steel is gripped in the chuck of the machine. Records from the mine showed, also, that the breakage was much more frequent in the short lengths of large gage than in the longer steels; this is shown graphically in Fig. 1. Because the shorter lengths of steels are usually those that have been made from broken pieces of long steels or from steels that have worn down in service, rarely being made from new stock, those who are wont to suspect steel of tiring and crystallizing in its old age through fatigue are quite apt to see in this instance an argument in favor of their idea. As usual, a careful examination of these steels reveals defects of a character hardly to be overlooked in favor of such a cause. The steel in question was rolled with a sand core and was of the usual drill steel composition: Carbon 0.85 per cent., manganese 0.60 per cent., silicon 0.25 per cent., phosphorus 0.02 per cent., and sulfur 0.02 per cent., with the exception that in some specimens there was vanadium, not exceeding 0.20 per cent.

Jan 3, 1924

By S. D. Strauss

The marketing of nonferrous metals and of the ores and concentrates from which these metals are recovered is a fascinating trade, international in character, sensitive to every change in the economic tide, and complex in its organization. Each individual metal has its own marketing mechanism; no two are exactly alike. The differences sometimes are a matter of tradition rather than of necessity, but usually they have developed because of the very nature of the problems of producing or consuming the individual metal. This description, because of limitations of space, can do no more than indicate the major aspects of the trade and touch lightly on the marketing structure of each metal. The dictionary states that a market is a "meeting of people for selling and buying." In the case of nonferrous metals this meeting may occur on a half-dozen different occasions between the time the ore has been first mined and the time the refined metal is available for use by industry; or it may not occur at all until after the metal itself has been rolled, drawn, stamped, cast, or forged into a useful semifabricated shape. To illustrate: A Chilean miner may sell his small ore production to a custom mill, which produces concentrates that are sold to an ore dealer, who finances the shipments of the concentrates to a United States custom smelter, which buys these' concentrates for conversion into blister, which is subsequently refined into electrolytic copper and sold. Conversely, much of the production of aluminum is in the hands of integrated concerns that mine bauxite, convert bauxite to alumina, smelt the alumina to produce aluminum, and then fabricate the aluminum into semifinished or even finished products before title to the material changes hands for the first time. Some nonferrous ores are not converted into metal at all but are sold by the miners or dealers directly b the ultimate consumers. A good example would be chrome ores used by the chemical industry. And of course a substantial part of the commerce in nonferrous metals is in scrap, which is the reconversion into useful form of metals that have long since been mined and smelted. Thus, one-third of all the lead used in the United

Jan 1, 1959

By Morris E. Fine, Richard A. Kloske

The stress-strain beha1,ior in tension of Au-Ag alloy single crystals containing nominally 1,3, 5, 95, 97, and 9.9 at. pct Ag was studied uS strain role and lektlperalure down lo 4.2K. A slrain aging yield hob lc,rrs observed on aging under stress in the temperature range 30° lo 75°C. The species diffusing to the dislocation is thought to be a divacancy-solute complex wilh /he. solute then pinning the dislocation by short-Range ovdering and possibly Suzuki locking. At room tempevullcve the critical reso1ved shear stress .follo~c.s an empirical equation of the form tc = A +Bc' where A and B are 90 and 420 ,g per sy mm tor gold base alloys and 60 and 510 g p~r sq )11ttz .tor si1ver base alloys. This strengthening was attributed to a long-range size ejtect. The alloy slrenglliening at i.2'K is greafer than at room temperature. The additional atrloutzf was altribuled lo a uwak skovl-range i?~levacIion between the solute and dislocallon cove. The activation energy .tor deformation a/ 4.Z0h7 decreases on alloying. 11 increases irr other fee systetns. hi the 5 at. pcl Air-Ag- alloy in pmvlicular there is very extensive easy glide and lavge overshooting- of the synmetry line at 4.Z°K. There is also u sharp decrease in the vale of. strain hardenitlg near the widdle of stage II with the new rate bezng abold the sarne as tlzal In easy glide. This uws attributed to a sudden reduction of activity on the pvitrrarg sgstertr. In Au-Ag alloys the critical resolved shear stress Tc. at room temperature is essentially a parabolic function of atomic concentration:' the 50 at. pct Ag-50 at. pct Au alloy is about 10 times stronger than the pure metals. This strengthening occurs in spite of the similarity in valence and atomic size. Solid-solution strengthening in fcc metals is characteristically greater at low temperatures than at room temperature. In Ag-10 at. pct Au and Au-10 at. pet Ag; r, (20.4K)/Tc (300°K) is 2.3 and 2.5. respectively.- compared to 1.2 in pure silver.3 Suzuki' and Flinn5 explained the alloy strengthening at room temperature as a combination of Suzuki locking and short-range order hardening. The agreement between the computed and measured strengthening was very good: however, in Ag base-A1 alloys Hendrickson and Fine6 concluded that Suzuki locking resulted mainly in a yield drop effect. Recent reviews of the solid-solution strengthening in fcc metals by Fleischer2 and Haasen8, 9attributed the strengthening at room tenl-perature to the combined effect of atom nlisfit and change in shear modulus 011 the long-range stress field surrounding a dislocation. The parameter used by Fleischer was The shear tnoduli of Ag-0 to 6 at. pct Au alloys were measured by Pur-wins. Labusch. and Haasen.In While addition of 4 at. pct Au caused an appreciable increase in the C,, elastic constant. there was a decrease on increasing the solute to 6 at. pct Au: C,, and C :: were not measured in the 6 at. pct alloy. The greater solution hardening at low temperatures implies the presence of short-range interactions between the solute atoms and dislocations These include size and electronegativity effects. Whether the extra alloy strengthening at low temperatures is due to locking or friction hardening is still a controversy."".' In the present research the stress-strain behavior of Au-Ag single crystals containing 1 to 5 at. pct Ag and 1 to 5 at. pct Au was measured vs strain rate and temperature down to 4.2' K. Particular attention was paid to yield drop effects. These have not been previously reported. PROCEDURES The alloys were supplied by the Engelhard Industries. Inc.. as strips 0.050 by 0.250 by 12 in. The compositions are given in Table I. Single crystals were grown under static vacuum using a traveling molten zone technique. Crucibles shaped to the size of the strips were made of high-purity graphite prebaked for 24 hr at 1100°C in vacuum. The crucible and its charge were placed in

Jan 1, 1970

By H. A. Doerner

The process for smelting zinc developed several centuries ago is still in use. Through the experience accumulated over this long period of time, details of the process have been perfected until there is little opportunity to make any further improvement except by a radical change in method. Such a change was accomplished by the process recently developed by the New Jersey Zinc Co. The process to be herein discussed is a still more radical departure, which has been developed to large-scale laboratory operation by the Metallurgical Division of the U. 5. Bureau of Minesl. † The reduction of zinc oxide by carbon is a highly endothermic reaction, which requires temperatures in the range 1200° to 1400' C. for even a slow reaction rate. The transfer of the heat of reaction through thick fire-clay walls, the high radiation loss in the type of furnace used, and the practical impossibility of recovering or using the combustible products of the reduction (CO) cause a deplorable waste of fuel. With the one exception mentioned, all zinc smelting is done in small retorts manually charged and discharged. Attempts to mechanize these operations have been only partly successful, and there seems to be no way to avoid or improve the uncomfortable conditions in which the work is done. The reduction of zinc oxide by methane has a relatively high reaction rate at 1000' C. At this temperature most ore residues do not sinter, and it is possible to use metal retorts. These factors make gaseous reduction especially adapted to a continuous operation and compared to reduction by carbon make possible: (1) higher rate of reaction at lower temperature, (2) more nearly complete extraction of the zinc, and (3) much better utilization of fuel and reagent. Although the reaction between methane and zinc oxide requires 105,900 calories per mole, compared to 86,860 cal. in the carbon reduction, it is obvious that much less fuel will be required to supply 105,900 cal. through 36 in. of metal at 1000' C. than to transfer 86,860 cal. through 1 in. of fire clay at 1300° C. Moreover, the gaseous products of gaseous reduction (CO and H2) can supply 183,430 cal. by complete combustion.

Jan 1, 1937

By Henry M. Howe

Jan 1, 1908

By R. L. Templin

THE increasing use of aluminum and its alloys in commercial fields has demanded a better understanding of their machining properties. This fact is exemplified by problems that have arisen in the automotive and airplane industries, but many in other fields might be cited. As pure aluminum and its alloys in their various commercial conditions show appreciable differences in their machining properties, it is not surprising that quite divergent solutions have been offered for the machining problems encountered. However, if the fundamental requirements of the most suitable cutting tools for these metals are understood, these machining problems lend themselves more readily to satisfactory solutions. Since the machining of free cutting brass and mild steel is understood by most persons accustomed to working these metals, it may serve our purpose better to first make a general comparison of the tools more commonly used in machining these metals with the tools most suitable for machining aluminum, then proceed to a more specific discussion of the individual tools. COMPARISON OF CUTTING TOOLS Cutting tools commonly used for machining free cutting brass usually have little, if any, top and side rake; they are ground on a medium to coarse abrasive wheel and used without any cutting compound or with a cutting compound that has a paraffin base. Those ordinarily used for steel have some top and side rake, are usually ground on a medium to fine abrasive wheel, and are often used with soluble-oil cutting compounds. The proper tools for aluminum and its alloys should have appreciably more side and top rake than the tools for cutting steel; should have very keen edges obtained by grinding with fine or very fine abrasive wheels supplemented in many cases by hand stoning with an oil-stone; and should be used with suitable cutting compounds whenever possible. In many cases, tools suitable for machining aluminum and its alloys are not appreciably different from tools commonly used for cutting hardwoods. The front clearance of a tool most suitable for machining aluminum and its alloys should be about 6°, the top rake from 30° to 50°, making the total angle of the cutting edge of the tool from 35° to 55°. A side

Jan 1, 1928

By P. F. Gnirk, J. B. Cheatham

Single bit-tooth penetration experiments under static load were conducted on six rocks at confining pressures of O to 5,000 psi using sharp wedge-shaped teeth with included angles ranging from 30 to 120° In general, the force-displacement curves for all rocks exhibit an increasingly nonlinear and discontinuous behavior with decreasing confining pressure. The confining pressure at which a rock exhibits a macroscopic transition from predominantly ductile to predominantly brittle behavior during penetration vmies from about 500 to 1,000 psi for the limestones to greater than 5,000 psi for dolomite. The correlation between calculated values of force per unit penetration based on plasticity theory and experimental values is quite encouraging, even at confining pressures as low as 1,000 psi A qualitative correlation between volume of fragmented rock per unit energy input for a single bit-tooth and drilling rate for microbits appears to exist over a confining pressure range of 0 to 5,000 psi. INTRODUCTION Laboratory experiments utilizing a small-scale drilling apparatus have demonstrated that penetration rates are reduced considerably as a result of increasing the confining pressure from atmospheric to a few thousand psi?-3 This undesirable situation can, in general, be attributed to a combination of decreased efficiency of chip removal at the bottom of the borehole, increased rock-failure strength, and a possible change in the mechanism of chip generation and rock fragmentation with increasing confining pressure. To more fully understand the principles underlying the last circumstance, it is the purpose of this investigation to experimentally study the mechanism of single bit-tooth penetration into dry rock at low confining pressures and, in particular, to establish the confining pressure at which the penetration mechanism may undergo a brittle to ductile transition for various rock types commonly encountered in drilling. Confining pressure as used here refers to the differential pressure between the borehole fluid pressure and the formation-pore fluid pressure. EXPERIMENTAL PROCEDURE Using an experimental apparatus previously described: a single, sharp wedge-shaped tool was forced under a "statically" applied load into an effectively semi-in finite dry rock sample subjected to a prescribed confining pressure. To prevent the invasion of the con fining-pressure fluid into the pores of the rock sample during penetration, the exposed surface of the rock was jacketed with a layer of silicon putty.* Electrical instrumentation incorporated into the apparatus yielded a graphical plot of force on the tool as a function of penetration or displacement of the tool into the rock during an experiment. During the course of the experimentation the following conditions were maintained constant: (1) pore pressure — atmospheric (i.e., the rock was dry); (2) temperature — 75F; (3) rate of loading — essentially static (approximately 0.002 in./sec); (4) bit tooth — a sharp wedge-shaped tool loaded normal to the rock surface; (5) rock surface — smooth and flat; (6) drilling fluid — hydraulic oil; and (7) maximum depth of penetration — approximateiy 0.1 in. In addition, each experiment was performed on a different rock sample so the rock surface is free of a layer of cuttings and of any previous indentation craters. The influence of the corners of a borehole was neglected, since each rock sample was cemented inta a section of aLuminum tubing to simulate a semi-infinite body. Indentation experiments were made on a sandstone, two limestones, a marble, a dolomite and a schist with a 60' bit-tooth over a range of confining pressures from atmospheric to 5,000 psi and at a constant confining pressure of 1,000 psi for a variety of bit-teeth with angles ranging from 30 to

Jan 1, 1966

By J. B. Cheatham, C. E. Murphey

Presently, computer control of Borobolic direction cannot be obtained during drilling, and most straight-holc drilling methods attempt to resist hole deviation rather than control direction. Many of the theories advanced as possible explanations of the cause of hole deviation are Summarized berein. A new correlation of physical partables is introduced to indicate bow factors such iis drill collar stiffness clearance and hit weight influence borehole deviation, .A methord is proposed for predicting the rate of change of hole angle when drilling conditions are changed. INTRODUCTION Control of borehole direction during drilling can be difficult and costly. Unintentional crooked holes are often lrilleti in dipping formations and many times directional drilling is required when the surface location is not directly above the target area —- for example, at offshore and mountainous locations. Prilllng progress can be greatly hindered in either air or liquid drilling when it be comes necessary to use low bit weight to prevent excessive hole angle build-up. If hole inclination becomes too great, drill pipe drag becomes excessive and fishing risks are increased, logging is more difficult and problems of differential sticking, key seating and fatigue failures may be encountered. Dog-legs and key seats are more serious problems than steep inclination angles: therefore, reducing rate of direction change is preferred to attempting to maintain absolutely vertical holes. Consequently, a straight inclineti hole is preferable to a nearly vertical crooked hole containing numerous dog-legs. In this paper, theories of the cause of hole deviation and analyses of drill string behavior under down-hole conclitions are summarized. Methods for computing hole deviation are presented und systems for resisting deviation as well as neans for provic!ing control of holt- direction are iliscusseI. A new correlation of physical variables is introduced tu indicate how factors such as drill collar stiffness, clearance and bit weight influence borehole deviation. A method is oroposeti for predicting the rate of change of hole anFrle when drilling conditions are clianged. IIEVIEU' OF PREVIOUS WORK ON HOIaE DEVIATION Significant progress in the theoretical analysis of hole deviation problems has been made in the past 1 5 years. The pioneering work has been primarily a result of the efforts of Lubinski and Woods.l-5 In 1950, Lubinski 1 considered the buckling of a drill string in a straig:rt vertical hole, a problem also considered by Willers6 in 1941. It wns concluded that very low bit weights must be used to prevent hole deviation resulting from drill collar buckling. The use of conventional stabilizers was proposed2 in 1951 by Mac Donald and Lubinski as a method for permitting greater Sit weights to he carried without drill collar buckling. These authors pointed out that a 2° nearly vertical spiral hole can cause severe key seatinp, and drill pine near, wilereas n 3 straight inclined hole with deviation all in one direction, while not vertical, will not result in serious drillirig or producing problems. Studies were continued with an investigation of straight inclined holes by Lubinski and Woods3 in 1953. In this paper they concluded that perfectlv vertical holes cannot be drilled even in isotropic forriations unless extremely low bit P-eights are used. .l'l~ey postulated that constant drilling conditions produce holes of constant inclination angle and varying conditions cause the hole to drill at a neiv equilibrium angle. This analysis was not concerned with drill string buckling since it was based on an equilibrium solution in which the ~irill string Was presumed to lie along the lower side of the hole abcrve the point of tangency. Weight of the drill collars below the point of tangency tends to force the hole toyard the vertical, whereas the weight on hit tends to force the hole aurav from the vertica l. The concept of an anisotropic formation Lvas introduced as an empirical method for explaining actual drilling data and as a means for extrapolating known deviation data to otiier conditions of hit

Jan 1, 1967

By J. R. Long, C. E. Armantrout, A. H. Roberson, T. R. Graham, R. S. Dean

THE general program of the Federal Bureau of Mines on the study of alloys made with electrolytic manganese has been extended to copper-manganese-aluminum alloys. The initial results of the work are presented in this report and concern a determination of the mutual solubilities of these three metals to form the alpha solid solution. The principal studies of this system, reported in the literature, have been focused on the Heusler alloys containing 10 to 30 per cent Mn with 9 to 15 per cent Al. These investigations concern the ferromagnetic properties and the determination of the crystallization contours, with little emphasis on the phases present in the solid state outside of this area. Heusler1 recognized a breakdown of the beta solid solution of the magnetic alloy, which he says decomposes into the alpha, beta, and gamma phases of the copper-aluminum system, somewhat modified by the presence of manganese. Rosenhain and Lantsberry2 examined alloys ranging up to 10 per cent Mn and io per cent Al but were unable to reach definite conclusions concerning the phase boundaries in that area. Krings and Ostmann3 made magnetic studies and examined the system by thermal analysis. They present contours of the liquidus surface but do not place any of the phase boundaries. Corson4 [ ] has drawn a hypothetical boundary for the alpha field with few data to support such construction. The binary systems involved have received considerable attention by numerous investigators, and their general features are well established. The hopper-aluminum system (Fig, I) has been thoroughly examined by Stockdale,5 who places the limit of solubility of aluminum in copper at 7.4 per cent at the liquidus and 9.5 per cent at room temperature. Smiths established the eutectoid temperature at 570°C. by a study of isothermal decomposition of the beta phase. The copper-manganese system (Fig. 2) has been recently reviewed by Dean and co-workers.7 Using high-purity alloys, they place the solubility of alpha man-

Jan 1, 1946

By C. Baldwin Sawyer

During the last half century, much time has been devoted to investigations of the effect of nitrogen on the physical properties of steel, but in all discussions of results there is considerable doubt as to the accuracy and interpretation of analyses. No conclusions concerning the effects of nitrogen on the physical properties of steel can be more reliable than the methods used for determining the quantities of nitrogen present. The work described in this paper endeavors to establish some additional conditions required by the customary distillation method for reliable and nccurate determinations. Results by a combustion method of analysis for total nitrogen content are presented; these indicate that all of the nitrogen in steel, exclusive of that contained in blowholes, may be determined by the distillation method. Based on these more sound and more reliable analyses, results havc been obtained that bear on the following divisions of the general subject, nitrogen in steel: Nitrification of steel by melting in nitrogen, nitrification of steel by heating in ammonia, decomposition of nitrified steel by heating, and thermal analysis of nitrified steel. Consideration of Two Methods of Analysis and Some Results The details of the distillation and combustion methods are chemical in character and of little interest to most readers. They are therefore described in an appendix. The distillation method is essentially that described by Hurum and Fay1 and was generally used throughout this work. It is capable of the greatest accuracy and all results given are significant to the last, figure. Where analyses do not check, the samples

Jan 1, 1923