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By N. Pessall, A. D. McQuillan

The phase relalionsllips in an isothernzal section of the Ti-Zr-H system at 750°Chave been investigated by measurement of hydrogen equilibrium pressures. The resulting diagram indicates that the hydride phase has an anomalously wide range of homogencity extending from 66.6 to 30 at. pct H in alloys containing about 15 al. pet Zr. Measurement of the heat of solution of hydrogen as a function of alloy composition in ß-phase alloys indicates no anomalous behavior in the region of 15 at. pet Zr, and suggests that the extensive range of homogeneity in the hydride phase at about this composition must he attributed to an enhanced stability of the, hydride phase. THE manner in which titanium and zirconium react with hydrogen and the constitutional diagrams of the two metal-hydrogen systems are qualitatively very similar,'-6 as can be seen from Fig. 1 which shows the constitutional diagrams of the two systems. Quantitatively, however, there are some differences. For example, the exothermic heat of solution of hydrogen in ß titanium is smaller than that in ß zirconium, and consequently the equilibrium hydrogen pressure of a Ti-H alloy is higher than a Zr-H alloy at the same temperature and hydrogen concentration. Other differences manifest themselves in the hydride phases which exist at high hydrogen concentrations. In the Ti-H system, above room temperature, there exists only a single hydride phase in which the metal atoms are arranged on a fee lattice, but in the zirconium system there are two hydrides, and whereas the hydride of lower hydrogen concentration has a fee structure, the higher hydride has a face-centered tetragonal structure. The tetragonal hydride retains its stability up to temperatures in excess of 450oC.7 From the foregoing, it would be expected that the constitutional diagram of the ternary Ti-Zr-H system would have a very simple form, especially as titanium and zirconium are completely miscible in both the a and ß forms.8 The present work, which consists of a study of an isothermal section of the ternary system, was carried out in order to check the assumption frequently made that the electronic structure of the two metals and the mechanism by which hydrogen is held in solution in the metals is the same for each. Although this proved largely to be true, certain inter esting and unexpected features of the ternary system emerged which must be explained before we can really say that the reactions of titanium and zirconium with hydrogen are fully understood. EXPERIMENTAL PROCEDURE The most convenient method of studying hydrogen-metal systems is by the measurement of the hydrogen pressure in equilibrium with the hydrogen dissolved in the metal as a function of hydrogen concentration and temperature. In binary systems, the results are easily interpreted in terms of alloy constitution, since two-phase regions can be recognized by the fact that the hydrogen pressure at constant temperature does not vary with the hydrogen concentration in the alloy when it is in the two-phase condition. This criterion does not apply in ternary systems, but there is usually a marked change of slope in the isothermal pressure-composition curves when a phase boundary is crossed, and with care, and a knowledge of the limiting binary systems, the phase structure of the alloys can usually be deduced. In the

Jan 1, 1962

By C. K. Leith

CERTAIN ideas on iron and steel sanctions to follow originated in a series of conferences held under the joint auspices of the War Department and Brookings Institute in Washington last spring. The views presented may be regarded as those of substantial segments of the steel industry. In view of the fact that steelmaking capacity can be so quickly converted to the production of armament, any effective control must necessarily include the entire steel industry and not solely the parts devoted to the production of implements of war. It must also be exercised all the way from production and shipment of raw materials through their ultimate conversion into finished products and use. Control will be dis- cussed under two main headings: (I) control of raw materials at the source; (2) control of steel capacity within the Axis countries.

Jan 1, 1944

By Jack R. Elenbaas, Donald L. Katz

A radial turbulent flow formula has been developed which permits the computation of the pressure drop for radial flow in gas wells whether the flow is laminar, turbulent, or partially laminar and partially turbulent. Using the formula a complete back-pressure curve has been calculated and analyzed by a comparison with existing back-pressure curves. A procedure is presented for computing the permeability of the porous media when the porosity, sphericity and average particle diameter are known. Introduction Recent studies on the flow of fluids through porous media have provided new methods for computing flow under turbulent conditions. The present study was initiated as an investigation of pressure -drop computations for flow through porous sands in gas wells and as an analysis of present-day back-pressure tests. Although laminar flow exists in the producing formation of gas wells under normal flow rates, turbulent flow does take place adjacent to the well bore. As the flow rate is increased, turbulent flow exists further and further into the producing formation and under open-flow conditions in relatively deep wells a considerable portion of the flow through the sand will be turbulent. In order to calculate a complete backpressure curve for relatively deep wells without resorting to extrapolation of curves drawn from data obtained in back-pressure tests or calculations made in the completely laminar region, it was necessary to have some means of calculating pressure drop for flow through porous sands under turbulent conditions.4.6 As a result, a radial turbulent flow formula was derived from the equations recently presented by Brownell and Katz.1 A rigorous solution of this equation involves a trial and error graphical integration. By making the assumptions of average viscosity, temperature and compressibility factor of the flowing gas, it was possible to graphically integrate the relation for the general case of radial flow of gases through the porous sand of gas wells. A chart has been prepared relating the value of the integral to the radius of the well bore and to the Reynolds number of the flowing gas. The chart actually performs the graphical integration and the solution of the formula is then reduced to a simple trial and error calculation. A typical back-pressure curve is calculated to show the transition from laminar to turbulent flow. This paper will briefly describe the pro-cedurcs for computing flow through porous media, including a method for predicting sand permeability from porosity and grain size and shape. Flow through Porous Media Brownell and Katzl recently reported a correlation for computing flow of fluids through porous media by the use of an enlarged Reynolds number and friction factor in which the porosity of the bed is included as an additional prime variable. The enlargcd Reynolds number com-

Jan 1, 1948

By Donald L. Katz, Jack R. Elenbaas

A radial turbulent flow formula has been developed which permits the computation of the pressure drop for radial flow in gas wells whether the flow is laminar, turbulent, or partially laminar and partially turbulent. Using the formula a complete back-pressure curve has been calculated and analyzed by a comparison with existing back-pressure curves. A procedure is presented for computing the permeability of the porous media when the porosity, sphericity and average particle diameter are known. Introduction Recent studies on the flow of fluids through porous media have provided new methods for computing flow under turbulent conditions. The present study was initiated as an investigation of pressure -drop computations for flow through porous sands in gas wells and as an analysis of present-day back-pressure tests. Although laminar flow exists in the producing formation of gas wells under normal flow rates, turbulent flow does take place adjacent to the well bore. As the flow rate is increased, turbulent flow exists further and further into the producing formation and under open-flow conditions in relatively deep wells a considerable portion of the flow through the sand will be turbulent. In order to calculate a complete backpressure curve for relatively deep wells without resorting to extrapolation of curves drawn from data obtained in back-pressure tests or calculations made in the completely laminar region, it was necessary to have some means of calculating pressure drop for flow through porous sands under turbulent conditions.4.6 As a result, a radial turbulent flow formula was derived from the equations recently presented by Brownell and Katz.1 A rigorous solution of this equation involves a trial and error graphical integration. By making the assumptions of average viscosity, temperature and compressibility factor of the flowing gas, it was possible to graphically integrate the relation for the general case of radial flow of gases through the porous sand of gas wells. A chart has been prepared relating the value of the integral to the radius of the well bore and to the Reynolds number of the flowing gas. The chart actually performs the graphical integration and the solution of the formula is then reduced to a simple trial and error calculation. A typical back-pressure curve is calculated to show the transition from laminar to turbulent flow. This paper will briefly describe the pro-cedurcs for computing flow through porous media, including a method for predicting sand permeability from porosity and grain size and shape. Flow through Porous Media Brownell and Katzl recently reported a correlation for computing flow of fluids through porous media by the use of an enlarged Reynolds number and friction factor in which the porosity of the bed is included as an additional prime variable. The enlargcd Reynolds number com-

Jan 1, 1948

By Joseph J. Becker

RECENTLY the "oriented growth" hypothesis of recrystallization textures has been receiving considerable attention. According to this view, "whenever a new generation of grains is growing in a highly oriented matrix . . . one may assume . . . that from among available nuclei of a great variety of orientations those with suitable orientation relationship with respect to the matrix have by far the highest rate of growth (oriented growth) and will, therefore, predominate."' This statement includes primary recrystallization of single crystals and of polycrystalline materials showing strong deformation textures as particular cases. While the phenomenon of oriented growth undoubtedly does operate, and may be of importance in coarsening or secondary recrystallization, the writer feels that its extension to primary recrystallization is a different matter. Several aspects of this phenomenon may be pointed out which are difficult to rationalize on the basis of oriented growth. In recrystallizing deformed single crystals of alpha brass, Maddin, Mathewson, and Hibbard2 found that the twin composition planes of the recrystallized grains were parallel to the three operative slip planes they observed, and not to the fourth octahedral plane. The writer's current work seems to indicate further that when only one or two slip systems can be seen to operate, the recrystallized grains are aligned only with those slip planes actually observed. This behavior could hardly be accounted for by oriented growth, because on the latter basis the recrystallized texture of a deformed single crystal should certainly have the same symmetry as the crystal itself. Growth rates could not be expected to distinguish between crystallogra- phically equivalent directions, unless the presence of a visible slip system vastly increases the growth rate in the appropriate direction. This seems unlikely. These relationships were observed by the writer only when considerable care had been exercised before annealing to etch away the effects of cutting out the centers of tensile specimens with a jeweler's saw and to etch off the ends of compression specimens. When these.precautions were not taken, nu-cleation inevitably appeared to occur at the saw cuts or at the ends, and the resulting orientations never showed a [111] rotation but appeared quite random. The provision of many nuclei of many orientations did not enhance the correlation, but destroyed it. Recrystallization textures are sometimes strongly dependent on annealing temperature, which would imply that the anisotropy of growth rates is a suddenly varying function of temperature. This might be regarded as unlikely. Recrystallization textures can be strongly dependent on the degree of deformation. If textures are to be explained by selective growth from among nuclei of all orientations, it is difficult to see how the recrystallization process could have any dependence whatever on the manner or degree of deformation. It is felt that the above points should be considered before the oriented growth hypothesis is applied to primary recrystallization, with the necessary complete divorce of the mechanism of recrystallization from the details of the deformation process. References 1P. A. Beck and H. Hu: Reply to discussion of their paper on Recrystallization Texture and Coarsening Texture in High Purity Aluminum. Trans. AIME, 188, 1057; Journal oF Metals (Aug. 1950) TP 29083. 'R. Maddin, C. H. Mathewson, and W. R. Hibbard, Jr.: The Origin of Annealing Twins in Brass. Trans. AIME, 185, 655; Journal of Metals (Sept. 1949) TP 26763.

Jan 1, 1952

By AIME AIME

EARLE EDWARD SCHUMACHER, the new Chairman of the Institute of Metal, Division. is well known to the metallurgical profession. His election a, Chairman is the culmination of fifteen years' service to the Institute, beginning in 1930 when he was elected Member. Since that time he has served as member of the I.M.D. Data Sheet Committee. 1935-40; Papers and Publication Committee. 1936; Executive Committee. 1937-1910: Papers and Programs Committee, 1937-1938. (Vice-Chairman, 1939-40, Chairman. 1941-12): Annual Award Committee. 194142. (Vice-Chairman, 19431: Vice-Chairman, LM.D., 1943-44; Chairman Publication Committee, 1943-44; member Powder Metallurgy Committee), 1944. Pew other men have contributed so Much time anti effort in so many way- to the interests of the metals branch o4 the A.I.M.E.

Jan 1, 1945

The following list contains the names of members of whose deaths notices have been printed in MINING AND METALLURGY during the year ending April 1, 1937. Biographical sketches published in MINING AND METALLURGY are indicated in the last two columns. YEAR OF ELECTION - NAME - DATE OF DEATH - ISSUE CONTAINING BIOGRAPHY - PAGE I892 ALLDERDICE, TAYLOR May 20, I934 1919 ALTMEYER, MAURICE E June 9, I936 1917 ANDERSON, JAMES P Jan. I4, I937 I936 ANDERSON, O RAN M Jan. 12,1937 I908 AUSTIN, E. A I935 1914 BACON, MAURICE W Jan. 4, I936 Apr. I936 22I 1914 BARTHOLOMEW, WILLIAM F Oct. 7, I936 Dec. I936 602 I930 BERRY, HAROLD R July 1, I936 1915 Betts, Robert M Jan. 9, I937 HON

Jan 1, 1937

By Wm. E. Milligan

LEAD stocks at the beginning of 1943 were comfortable when compared with those of other base metals such as copper, zinc and tin. This situation was early recognized by W.P.B. and other Governmental agencies and every effort was made to encourage the substitution of lead for these latter metals in all essential uses. Despite this apparently healthy condition, lead was not denied participation in the various measures inaugurated by the Government to relieve metal shortages. Thus in the Mexican trade agreement, effective Jan. 20, the tariff on two important classes of lead importss was halved. Similarly, in the setting up of B quota premiums, a B base was allotted to lead, with all pro¬duction above this figure receiving an additional payment of 2.75c. per lb., or a total payment of 11.75c. per lb. At about the same time it was announced that all quota transactions would be based upon the lead content as shown by the smelter settlement sheets. This basis was 86 percent of the lead content for all ores shipped to custom milling plants and 95 per cent for ores and concentrates shipped to custom smelting plants. The importance of this action can be appreciated when in the Coeur d'Alene it was estimated as equivalent to a 5 per cent increase compared to the previous base.

Jan 1, 1944

By P. K. Foh, C. G. Dunn

IN an investigation of the rolling and recrystalliz-ing textures obtained from single crystals of Si-Fe, Koh and Dunn included specimens in the (111) [lie] orientation.1,2 This note reports some additional results obtained on samples of this orientation set aside for further study. After, primary recrystallization longer annealing treatments at 980°C produced only normal grain growth; after 16 hr at 980°C the largest grains were approximately 0.010 to 0.020 in. in diam. The orientations of 13 of the larger grains were found by the transmission Laue method. The unidentified (110) poles of these grains have been added to the (110) pole figure obtained after primary recrystallization2 to give the plot shown in Fig. 1. When grain identifications are also included, such a plot reveals the following. Two orientations agree almost exactly with the peak pole ,density positions of component B of the primary recrystallization texture, one with component A, and one with component C. All the remaining orientations except one have (110) poles near weaker and unidentified pole density peaks. In fact, the correlation is good enough for an identification of eight components. Previously only components A, B, and C were certain. The results are perhaps even more informative when combined with the (110) pole figure of the cold-rolled crystal to give the plot shown in Fig. 2. Each of the 13 orientations is identified and uniquely determined by selecting two (110) poles per crystal, which are 90" apart. The plot shows that every grain of the 13, except grain 6, has a (110) pole in common with a (110) pole concentration of the cold-rolled crystal. The orientation relationship with the cold-rolled orientation may be expressed as approximately 25" rotations about <110> axes, which is in agreement with the results found for coarse-grained primary recrystallization.2 If primary recrystallization had occurred according to the <110> rule without preference, there would have been 12 components. Actually our limited data reveals nine of the possible 12. The integrated pole figure also shows preferences among them, since three are considerably stronger than the rest. In addition to the above direct information, the following conclusions may be drawn. Primary recrystallization of the (111) [110] cold-rolled crystal produces a limited number of weak and strong components with the larger than average primary grains from these components related to the cold-rolled crystal by 25° <110> relationships. In arriving at this conclusion of limited orientations the (111) [110] orientation has to be ruled out as a source of large primaries for two reasons. First, growth in this orientation was not observed, despite an expected high grain boundary mobility (due to its 25" <110> relationship with every component of the texture). Secondly, primary recrystallization usually produces grains in new orientations, particularly in a single orientation matrix; but (111) [110] is not new. Other orientations besides those of the <110> type may be considered unlikely, of course, because the <110> relationship holds quite well for primary recrystallization2 and should produce the large primaries. There are exceptions to this rule, but probably not enough to provide large

Jan 1, 1958

By AIME AIME

ON Thursday, Oct. 9, oil men from far and wide gathered at Beaumont, Texas to participate in a three-day celebration of the fortieth anniversary of the completion of the famous Lucas gusher well at Spindletop field just south of Beaumont, and to pay tribute to Anthony F. Lucas, for whom the Lucas Medal of the A.I.M.E. is named. In the afternoon the gathered host drove to the site of the discovery well to dedicate a 58-ft. granite monu¬ment erected on the spot where the original derrick stood. Representatives of the A.I.M.E. took a prominent part in the dedication ceremonies, giving the two principal addresses of the afternoon&apos;s program. Everette DeGolyer, past President, paid a personal tribute to Capt. Lucas, as a pioneer producer, an engineer, and a friend and well wisher of all young engineers. He gave a splendid picture of the useful and full life of Capt. Lucas who always took a great interest in the activities of the Petroleum Division of the Institute, from its inception to the time of his death.

Jan 1, 1941

H. Van F. Furman, Denver, Colo. (communication to the Secretary): There are some statements in Dr. Carpenter&apos;s interesting paper which appear to require explanation, if not modification.

Jan 1, 1901

By AIME AIME

ON April 7 the twentieth national Open-hearth Conference of the A.I.M.E. will be held in Birming¬ham, Ala., in conjunction with a meeting of the Committee on Blast Furnaces and Raw Materials. At least 300 engineers and operators are expected to attend. Birmingham, the industrial city and iron, and steel center of the South, has been called the "Magic City." Men actively engaged in business in Birmingham today have seen practically every building in the business district erected, have witnessed the establishment of every existing industry, and have hunted birds cm the site now one of the busiest corners of the business district. Incorporated as a town in 1871, the city today has a population of more than 380,000 in its metropolitan area. It is

Jan 1, 1937

By AIME AIME

AT A meeting of the Boston Local Section of the Institute, on Oct. 3, approval was voted to the work done by its representatives on the Committee opposing the passage of a bill by the, Massachusetts Legislature to license engineers. The report is as follows: The Joint Committee on Licensing of Engineers, which included a Committee of the Boston Society of Civil Engineers and representatives of seven other societies,* submitted to the members of the respective societies, on Jan. 8, 1921, a report in which it was stated that on Jan. 4, 1921, a petition was filed in the Legislature of Massachusetts by the Boston Chapter of the American Association of Engineers accompanied by a bill to provide for the registration of professional engineers and land surveyors. This report also stated that at a meeting of this committee, held on Jan. 7, 1921, it was voted:

Jan 1, 1921

By Pradeep K. Rohatgi, Clyde M. Adams

Structures produced upon solidification of the eu-tectic composition (33 wt pct Cu) aluminum copper alloy have been examined as a function of freezing rate dfs /d? , the rate of change of fraction solid (fs) with time (8). Slow (dfs/d? = 0.0016 sec-1), intermediate (dfs/d? = 0.02 sec-1) and rapid (dfs/d? = 0.4 to 7.30 sec-1) freezing rates were used. The lamellar Al-Cual2 eutectic is arranged in the form of rod-shaped colonies at rapid freezing rates. The colonies are aligned parallel to the direction of heat flow, whereas the lamellae within the colonies are aligned at various angles, as high as 90 deg, to the direction of heat flow. The colony spacing (C) is proportional to the square root of inverse freezihg rate. The relationship is C = 15.5(dfs/d?)-1/2 where C is in µ and 8 is in sec. The ratio of colony spacing to lamellar spacing is greater than 20.0 and increases with a decrease in the freezing rate. A duplex dendritic structure is produced at intermediate freezing rates. A fine lamellar eutectic is arranged within the dendrites (exhibiting side branches at an angle close to 60 deg from the main stem) and a coarse irregular eutectic appears in the interdendritic regions. The duplex eutectic structure is also produced at slow freezing rates. However, at slow freezing rates there is a Platelat of CuAl2, along the center of the main stem of each dendrite and the other lamellae are arranged perpendicular to the central platelet. THE eutectic between CuA12 and a! aluminum has been reported to freeze in a lamellar form by several workers.&apos;-3 chadwick4 has measured the interlamel-lar spacing as a function of growth rate. Kraft and Albright2 have reported on irregularities in the lamellar structures, and have proposed growth models which account for the formation of faults during solidification. In certain instances the lamellar eutectic has been found to exist in colonies. The colony formation315 has been attributed to the breakdown of a planar liquid-solid interface due to rejection of impurities. The aim of the present work is to study the structures produced from the eutectic aluminum-copper alloy under relatively fast solidification rates, such as encountered in casting and welding operations. The solid-liquid interface presumably remains planar under conditions of slow unidirectional freezing which produce lamellae aligned parallel to the direction of heat flow. The local growth velocities are the same over the entire interface and are equal to the rate of growth of the all-solid region. The spacing between the eutectic lamellae is inversely proportional to the square root of the growth rate of the all-solid region. Under the freezing conditions used in the present study, the solid-liquid interface is cellular or dendritic and the local growth velocities are different in the different regions of the interface. The relationship between the growth rate of the all solid region and the local growth velocities varies with the location and the shape of the interface. The growth rate of the all-solid region is, therefore, an inadequate parameter to describe the eutectic micro-structures which depend upon the local growth velocities. For this reason the structures have been examined as a function of freezing rate, dfs/d?, where fs is the fraction solidified at time 0. The freezing rate was varied by a factor of 4000. The relationship between the freezing rate, dfs/d?, and the growth velocit of the all solid region depends upon the specimen geometry and the shape of the interface. EXPERIMENTAL PROCEDURES The A1-33 pct Cu alloy used throughout this study was made in an induction furnace, using electrolytic copper and aluminum of commercial purity (99.7 pct), the primary impurities being silicon (0.12 pct), iron (0.14 pct), and zinc (0.02 pct). Three ranges of freezing rates were investigated: 1) A spectrum of rapid freezing rates (ranging from 0.40 to 7.30 sec-1) was obtained in arc deposits made on 2-in. thick cast plates of the eutectic alloy. The arc was operated at constant power and was made to travel at constant velocity on the surface of the plate that was in contact with the chill surface during solidification. The pool of liquid metal formed under the moving tungsten arc solidified rapidly by heat extraction through the unmelted plate. Conditions of unidirectional heat flow were achieved near the fusion zone interface, especially in the center of the arc deposits. The great advantage of the arc technique is that rapid cooling and freezing rates can be varied in a qualitative way. The correlation between the arc parameters and the solidification rate is given by the following relationship:6-8

Jan 1, 1970

By M. J. Marcinkowski, K. S. Raghavan, A. S. Sastri

A detailed transmission electron microscopy investigation was carried out in connection with a manganese Hadfield Steel. At small plastic strains, numerous individual intrinsic stacking faults are observed. With increased plastic deformation, the stacking faults thicken into twin lamellae which in turn subdivide the original austenite matrix into smaller domains. The twin boundaries act as strong barriers to subsequent dislocation motion and is in a sense equivalent to grain refinement. It is this "grain refinement" which is believed to be the cause of the very high work hardening rates in the Hadfield Steels. In many cases, especially where an hcp phase is the stable one at low temperatures, the stacking fault energy in fcc metals and alloys decreases with decreasing temperature.&apos; Since stacking faults of the intrinsic type are precursors of both twins as well as the hexagonal close packed structure, both of these entities should become more frequent as the temperature of a fcc crystal is lowered. In the case of the twin, there is no chemical driving force for its formation and it is generally necessary to provide the required driving force by an applied stress, i.e., strain energy. In the case of the hcp structure the transformation from the fcc modification can occur spontaneously (marten-sitically) since a decrease in chemical energy does in fact occur; however, an applied stress will provide an even greater driving force toward complete transformation. Since the transformation products mentioned above occur in an inhomogeneous manner throughout the crystal and since these can act as potential barriers to further plastic deformation2 marked strengthening effects can be anticipated. Also because metal and alloy strenghening is in general proportional to the shear modulus, these effects should be greatest in steels of the austenitic type (y), i.e., the fcc types. Perhaps the two most important steels in this category are the austenitic stainless steels and the Hadfield manganese steels. Both may be quenched from elevated temperatures so as to retain the austenitic states characteristic of those temperatures. The effect of subsequent deformation at lower tem- peratures has a profound effect on the stress-strain curves of these alloys. In particular Fig. 1 shows the compressive stress-strain curves obtained with an 18-8 stainless steel which was quenched from 1850°C after annealing for 1 hr so as to produce all y. As the temperature is lowered, the work hardening rate increases markedly. Although some hcp or c mar-tensite can be generated by plastic deformation as the temperature is lowered,~ it is believed to be a transition phase4 and most of the martensite produced is of the bcc or a variety.3 It is this stress induced martensite which gives rise to the very low initial work hardening at 77°K as can be seen in the stress-strain curve in Fig. 1. Similar low initial work hardening rates have been observed in the stress induced Ni-Ti martensites.5 Fig. 2 shows that an even more rapid rate of work hardening occurs in the Hadfield steels treated in the same way as that described for the 18-8 stainless steels a; the temperature is lowered. It is this ability to work harden to such high stress levels that makes the Hadfield steels particularly suitable for armor plate and heavy construction equipment. However, unlike the case of Fig. 1, no initial low rate of work hardening is observed in any of the curves in Fig. 2. Thus the stress induced formation of any low energy martensite phase in any significant quantity must be ruled out. This observation is in accord with the X-ray findings of Otte.~ On the other hand, small quantities of the E phase have been observed by other investigators using transmission electron microscopy (TEM) above Even more significant was the fact that large numbers of deformation twins were observed in the deformed Hadfield steels,678 which were postulated to be one of the reasons for the high work hardening ability of this class of steels.8 It is the purpose in what follows to discuss a series of experimental observations pertaining to the stress induced transformation in a Hadfield steel and to formulate a dislocation mechanism which adequately accounts for the observed results. EXPERIMENTAL PROCEDURE The stainless steel used to obtain the curves shown in Fig. 1 was of the AISI Type 303 containing approximately 18.0 pct Cr and 8 pct Ni. On the other hand, the Hadfield manganese steel used to obtain the curves shown in Fig. 2 contained between 1.00 to 1.25 pct and 11.5 to 13.5 pct Mn. In all cases the samples were in the form of compression cylinders 0.220 in. in diam and 0.370 in. long. Prior to testing the samples were annealed for a hr at 1050°C and rapidly quenched into a brine solution. This treatment was sufficient to preserve the y phase for subsequent testing at lower temperatures. All samples were compressed in an Instron testing machine using a cross head speed of 0.02 in.

Jan 1, 1970

By A. S. Nowick, C. Y. Li

It is deMonstrated that quenching from an elevated temperataupe accelerates the grain boundary relaxation in two solid solutions (aAg-Zn and a Cu-Al). This result is consistent with the proposal that, in solid solutions, grain boundary relaxation occurs by a mechanism of&apos; self diffusion. Nevertheless, an alternative possibilitg, that quenching introduces vacancies into the boundary itself, must also be considered. THE phcnomenon of grain boundary relaxation has been well known for many years,1,2 yet the mechanism of this process is very poorly understood. One of the most interesting suggestions which relates to the mechanism of grain boundary relaxation was that of Ke,3 who claimed that the activation energy for grain boundary relaxation and for lattice self diffusion were essentially the same. The implication is therefore that the elementary step in the two processes is the same. This suggestion is particularly startling in view of the fact that activation energy for self diffusion along a grain boundary is very significantly lower than that for volume self-diffusion. Later evidence5-7 showed that there really are two grain boundary peaks, one which appears in high-purity metals, and the other (which develops at a higher temperature than the first) which appears in solid solutions beginning at solute concentrations in the range of 0.1 pct. Data for silver6 show that Kg&apos;s hypothesis is surely incorrect for the grain boundary peak in the high-purity metal, since it has an activation energy of only 22 kcal per mole, but that the hypothesis may still be correct for the grain boundary peak in various silver solid solutions, for which activation energies in the range 40 to 50 kcal per mole are observed. If the elementary step in the grain boundary relaxation process were the same as that for self-diffusion, it would be expected that the relaxation process could be hastened by quenching, 2.c. by introducing a non-equilibrium excess of lattice vacancies. Such a quenching effect has already been demonstrated in the case of another anelastic relaxation process, viz., the Zener relaxation effect. The Zener effect, which occurs in essentially all solid solutions, may be attributed to the reorientation of pairs of solute atoms in the presence of an applied shear stress,&apos; and therefore must take place by means of a volume diffusion mechanism. The hastening of this process through quenching9 has been one way of demonstrating that atom movements in the lattice take place through a defect mechanism, presumably single vacancies. In order to see if the grain boundary relaxation is affected by quenching, it is particularly convenient to compare the grain boundary relaxation with the Zener effect, by choosing a specimen for which both relaxation effects appear. Specifically, a fine-grained sample of a solid solution shows in the curve of internal friction vs temperature, first a peak due to the Zener effect, then a second rise (and eventually a peak at substantially higher temperatures) due to the grain boundary relaxation. The same phenomena are also observable in static anelastic measurements, such as creep at very low stress levels. Thus, for the same fine-grained solid solution, the creepstrain, when plotted against log time, falls on a sigmodial curve with a sharp inflection point, due to the Zener effect, which is followed by a second rise and inflection resulting from the grain boundary relaxation. To look for a quenching effect, static measurements are preferable to the dynamic internal friction measurements, due to the fact that quenching effects tend to anneal out too rapidly at the temperatures at which the internal friction is measured.9 RESULTS AND DISCUSSION Creep experiments in torsion were carried out in an apparatus similar to that described by Ke1, whereby a wire is held under constant torque and its angular displacement is observed as a function of time. The alloy Ag-30 at. pct Znwas selected because of the large Zener relaxation that it displays. The two samples used were a "coarse grained" wire with a mean grain size about twice the diameter of the wire (diam = 0.032 in.), and a "fine-grained" wire which had several grains across the diameter. In Fig. 1 a comparison is made of the creep curves at 160°C of these two samples after they had been cooled slowly from 400°C. Curve A, which represents the coarsegrained sample, shows a unique relaxation process due to the Zener relaxation, with a relaxation time, T , in the vicinity of 100 sec. Curve B, which represents the behavior of the fine-grained sample, on the other hand, shows first the same relaxation process as that in A, followed by a turning up of the curve which corresponds to the onset of a second overlap-

Jan 1, 1962

By R. E. Smallman, I. Milne

The deformation behavior of single crystals of Nb-Mo alloys has been investigated with particular reference to the influence of composition, orientation, and temperature. Strong solid-solution hardening was observed reaching a maximum at the equiatomic cotrlposition and can be attributed to the difference in atomic size between niobium and molybdenutrz. Changes in the form of stress-strain curve, as shown by a high work-hardening rate and restricted elongation to fracture, were observed at a composition of Nb-85 pct Mo and are attributed to the presence of MozC DreciDitate. Conjugate slip was only extensive in dilute alloy samples; at the 50/50 composition deformation rnainly occurred by primary slip, and the onset of conjugate slip gave rise to failure by cleavage on (100). The variation of yield stress of Nb-50 pet Mo with orientation was consistent with slip on (011)(111) slip systems. The temperature deperndence of the yield stress between -196" and 250°C was similar to that of pure bcc metals, but at a much higher stress level; no evidence for twinning %as found. IN recent years the deformation behavior of various pure metals in groups VA and VIA has received considerable attention, but surprisingly little work has been carried out on binary alloys made by mixing metals from the two groups. Such an investigation would be of interest since single crystals of metals of group VA have been shown to deform characteristically with a multistage deformation curve1"3 while a parabolic type of deformation curve has been reported for most of the group VIA metals.4&apos;5 It has been suggested by Law ley and Gaigher~ that the difficulty encountered in obtaining multistage deformation curves for molybdenum in group VIA was possibly because of the presence of a microprecipitate of MozC which they observed even at carbon contents as low as 11 ppm. Recently a multistage deformation curve has been reported for molybdenum ," although the stages are not so definitive as those for group VA metals. The binary alloys of the particular refractory metals which have been investigated in single-crystal form include Ta-w,&apos; Ta- Mo,&apos; and Nb- Na." While a large amount of hardening was observed for alloys of the Ta-W and Ta-Mo systems, associated with room-temperature brittleness for alloys approaching the equiatomic composition, Ta-Nb remained ductile over the complete composition range with little or no solution hardening. Other systems have been investigated by hardness measurements on polycrystalline material and a discussion of the hardening of these alloys has been presented by ~udman." The purpose of the present investigation was to examine the deformation behavior of Nb-Mo alloys in detail, with particular reference to alloy composition and single-crystal orientation. In this way it was hoped to shed some light upon the restricted ductility of these alloy specimens. 1) EXPERIMENTAL PROCEDURE The starting materials were obtained in the form of beam-melted niobium rod and sintered molybdenum rod of suitable dimensions. Since niobium and molybdenum form a complete solid-solution series at all temperatures, alloy single crystals were produced by melting the two constituents together in an electron bombardment furnace (EBM). To produce specimens free from segregation a molten zone was passed over the length of each rod six times in alternate directions at a speed of 10 in. per hr. Typical specimens were analyzed for interstitial impurities by gas analysis and for metallic impurities by spectrographic analysis. The results of this analysis are shown in Table I. Many of the tensile specimens were also analyzed (after testing) by scanning the gage length in an electron beam microanalyzer, from which it was found possible to predict the approximate composition of a specimen from the original proportions of each element in the EBM. The tensile specimens were made with a gage length of 0.5 in. and diameter of 0.075 in., using a Servomet Spark machine. By careful machining on the finest range for the final i hr of this technique, surface cracks could be reduced to the level where they were easily removed by electropolishing in a solution of nitric and hydrofluoric acids. The specimens were strained at a rate of 10 4 sec-&apos; using friction grips designed to prevent accidental straining and maintain a good alignment before straining. The orientations of the individual specimens tested are shown in Fig. 1 and the corresponding compositions listed in Table I1 together with collated experimental data. 2)RESULTS a) General Deformation Behavior. The effect of composition on the room-temperature deformation curves of similarly oriented specimens is shown in Fig. 2. The yield stresses of the pure constituents, while not the lowest reported to date, were at least comparable with existing data. Although the solution hardening was large for alloys at either end of the phase diagram, and comparable with the Ta-W solution-hardening data of Ferris et a1.,8 the low work-hardening rate characteristic of niobium was sustained until a composition of Nb-85 pct MO had been reached. Associated with the peak yield stress ob-

Jan 1, 1969

By D. H. Thompson, A. W. Tracy

Season-cracking is a type of failure of brass that results from the simultaneous effect of stress and certain corrodants. The object of this paper is to present data that will aid in a more complete understanding of the mechanism of season-cracking and related phenomena. Results presented show that certain high copper alloys are susceptible to season-cracking or stress-corrosion cracking, and possible explanations are discussed. Starting at least as far back as 1906, many papers have been devoted to this subject but the symposium1 held in Philadelphia in 1944 is the richest source of information. In order to study season-cracking, several of the many variables were held constant so as to learn the effects of others. Season-cracking is generally understood to refer to the corrosion cracking of brass having internal stresses;²,³ it is a special case of the general stress-corrosion cracking. Inasmuch as applied stresses are more readily produced and controlled, they were used exclusively in this research and the resulting phenomenon must he called stress-corrosion cracking.²,³ Only constant tensile stresses were used. The agents believed to be most frequently responsible for season-cracking are ammonia. amines and compounds containing then]. Both moisture and oxygen also appear to he necessary. Therefore, an atmosphere containing ammonia, water-vapor and air was selected for these tests. Briefly, the work consisted of exposing sheet metal specimens, having a reduced section ¼ by 0.050 in., of copper-base alloys to the effect of static tensile stresses between 5,000 and 20,000 psi and simultaneous contact with a. continuously renewed atmosphere containing 80 pct air, 16 pct ammonia and 4 pct water vapor at 35°C. The gas mixture and the speci- mens were maintained above the dew-point. The time-to-failure in minutes was the primary measure of results. In order to limit the experiment to finite time, it was considered that a specimen which had neither failed nor undergone microscopically detectable cracking in 40,000 min. (4 weeks) while under a stress of 10,000 psi or more could be considered immune to cracking. This is merely a convenient limit and is not to be considered proof of immunity. Supplementary tests in the absence of stress using weight loss or microscopical appearance as measures of attack were made. Apparatus The apparatus used in this research is shown in Fig 1. To facilitate the description it may conveniently be divided into six parts: stress-producing units, test chamber, gas train, electrical controls, timers and gas analysis device. A stress-producing unit is shown in an exploded view at the left in Fig 2. At the right is an assembled unit with a specimen in place in the lower portion; it is this part that remains in the ammonia atmosphere during a test. The upper part contains a spring, a central threaded rod, a large nut and necessary washers, pins, and so forth. Stress is produced in the specimen by screwing down the top nut against the spring, thus putting a tensile load on the central rod and so on the specimen. The wrench that turns the nut by extending through the upper cap, is seen at the upper right of the figure. The magnitude of the load is gauged by measuring from the pin that extends through the side of the tube, to a fixed point on the large flange. Measurement is made with a vernier beam caliper, shown at the right of the figure. The necessary spring compression to give a desired stress is calculated from the calibration curve of the spring and the dimensions of the specimen. The test chamber, center Fig 1, consists of a thermally insulated steel box 32 in. long by 10 in. high by 7 in. wide. A horizontal baffle reaching nearly to each end divides the chamber equally. Below this baffle are inlets for air and ammonia, a heating coil and a fan. Thus the gases are warmed and mixed in the lower level and flow past the specimens in the upper level. A thermo-regulator and thermometer project into the upper space. The top is pierced by 12 ports flanked by 3/8 in. threaded studs. A test starts when a port is opened and a unit containing a stressed specimen is thrust through it and bolted down against a neoprene gasket. The test chamber is held at 35°C. The gas train, right rear Fig 1, carries ammonia and air continuously to the test chamber. Tank ammonia passes through two reducing valves, a needle valve, a flow meter and into the test chamber. The air from either the plant compressor or a small laboratory compressor passes through wool towers and flow controls to the flow-meter. It then bubbles through water at 34°C and through a heated line to the test chamber. Electrical controls, left rear, Fig 1, provide rectifiers and mercury relays for the test-chamber and humidifier-heating-control circuits and outlets for

Jan 1, 1950

By Eugene McAuliffe

MANAGEMENT and control of any body of workmen can be effected through various - well-known methods &apos; though many managers hold certain personal theories of control that range from an absolute dictatorship to what might be called the principle of laissez faire, or the "let alone" method. Managements of many mines, factories, smelters and refineries have been compelled in recent years to accept certain restrictions in hours of labor as well as in methods employed, imposed either by federal or state laws, or by labor agreements. Broadly speaking, regulations of the above character were brought on the industry by managers who failed to recognize fully the rights of their employees to proper sanitary, safe, and compensatory working conditions, as well as working hours that were proper and just. As an example of poor management and worse psychology, we have only to hark back to the days when many bituminous coal operators maintained an inadequate number of pit cars for the number of miners or loaders employed, or if the pit cars were available, the haulage facilities

Jan 1, 1934

By G. D. Hobson, I. Mrosovsky

Material balance relationships for expansion-type reservoirs above the bubble point have been discussed in recent years by H. N. Hall and M. F. Hawkins, Jr. The former drew attention to the effects of rock compressibility, and the latter to the effects of the compressibility of any interstitial water. Hawkins derived a general relationship which took note of oil, water and rock compressibilities, and then by assuming first, the absence of interstitial water and secondly, no effects due to interstitial water or rock compressibility, obtained relationships symbolically of the same form as those previously derived directly by Hall. Hawkins defined his terms clearly, and noted an approximation, due to the omission of a term involving the product of two compressibilities, made in reaching his final equation. There is, however, possible ambiguity with regard to Hall&apos;s written definition of the oil production term. One interpretation would mean that the produced oil and initial oil-in-place, if both required as stock tank barrels, would be obtained by using different formation volume factors in the conversion in order to satisfy the conditions implicit in Hall&apos;s formulae. The other would involve an approximation which, if used, was not stated. In both Hall&apos;s and Hawkins&apos; studies the effect of the approximation or inconsistency is very slight. Nevertheless, it may be justifiable to give a somewhat more consistent derivation, since the complexity of the final relationships is not thereby greatly increased. Let the initial pore volume be Vi and the volume after a pressure drop of Ap, V. Then V = V, (1 - ?p . cr), where c, is the rock compressibility expressed as change in pore volume/unit pore volume/unit pressure change, and is defined with the initial pressure as base. Suppose that the interstitial water saturation under the initial conditions is Sw. The total volume of interstitial water under those conditions will be ViSw. If the volume of the water after dropping the pressure by ?p is V., this will be given by Vw = ViSw (1 + ?p.c,w), the compressibility of the water being c, which is ex- pressed using the volume at the initial pressure as base. If Boi is the reservoir volume factor of the crude oil for the initial reservoir conditions, and B. the equivalent factor when the pressure has dropped by ?p, with N the initial amount of oil expressed as stock tank barrels. N.Boi = Vi - V,Sw If n is the amount of oil produced (in terms of stock tank barrels) as a result of the pressure drop ?p, the remaining subsurface volume of oil at the lower pressure will be Bw, and Bc can be related to the compressibi:ity of oil and the pressure drop, and if c, is the compressibility factor based on the initial conditions, The relationship derived by Hawkins was case to a standard at the lower pressure, while a term (Ap) had been dropped, because of its smallness arising from its dependence on the products of two compressibility factors. When there is no interstitial water and rock compressibility can be ignored, Eq. 1 reduces to:

Jan 1, 1957