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THE 143d* meeting of the American Institute of Mining and Metallurgical Engineers was held in New York, Feb. 19-22, 1934. It consisted of the annual business meeting, 45 technical sessions at which 217 papers were presented (more than ever before), two sessions of the Board of Directors, two meetings of Section delegates, three formal lectures, luncheon Monday, Tuesday and Wednesday for members and guests, three group dinners and five group luncheons, eight committee meetings, the annual reception and dinner dance, an informal dance, a smoke-dinner and the annual meeting of the Woman's Auxiliary. Attendance showed an improvement over the previous year, as follows: Total, 1534 in 1934, and 1475 in 1933; dinner dance, 543 and 426; smoker-dinner, 425 and 394; informal dance, 295 and 250; women's registration, 160 and 133; luncheons, 1034 and 734. The Coal Division held a general session and three sessions on Coal classification. It also met jointly with the Committee on Ground Movement and Subsidence. At the general session, John B. Dilworth presented a report of the Committtee on Methods of Valuing Coal Proper- ties and R. H. Sweetser a report of the Committee on Evaluation of Coal for Blast-furnace Coke. C. B Huntress, Executive Secretary of the National Coal Association, was the principal speaker at a luncheon on Monday, at which Eli T. Conner, Chairman of the Division, presided H. N. Eavenson, newly elected President of the Institute, and Edward W Parker, Secretary-Treasurer, of the Anthracite Institute, also spoke. The Executive Committee of the Division met Tuesday afternoon. The Sectional Committee on Classification of Coal met on Thursday morning and the Technical Committees on Scientific and Use Classification of Coal on Thursday afternoon. The Committee on Mining Methods held two general sessions, one for presentation of a half-dozen interesting papers on strictly mining subjects, the other on Aerial Transportation. It also participated jointly in a session on ground movement and subsidence with the committee of that name and the Coal Division The Institute's Standing Committee on Aviation had its inception at the session on Aerial Transportation

Jan 1, 1935

By Robert C. Mackenzie

EXAMINATION by the differential thermal analysis technique of a large number of samples of soil and other clays from various parts of the world has demonstrated that the occurrence of cold-precipitated hydrated ferric oxide (the identification of which in soil clays has previously been reported5) is not confined to any limited locality; e.g., it has been found in varying amounts in clays from Britain, France, South Africa, and Mexico. The conditions necessary for its formation have not yet been elucidated in spite of the number of samples examined, and, since this mineral is so easily prepared in the laboratory, it has been considered that a study of its constitu¬tion and of the conditions under which it could be prepared might. lead to a better understanding of the factors involved. These experiments form the subject of this short paper, and while they do not supply a definite answer to the mode of formation of this mineral in nature, they may, it is hoped, add something to our knowledge in this direction. EXPERIMENTAL WORK Materials Employed The a-Fe2O3.H2O (goethite) and B-Fe2O3.H2O (Weiser and Milligan") were obtained from the Department of Geology, University of Aberdeen, and Rothamsted Experimental Station, respectively. The beta material had been prepared by the method described by Fernelius.1 y-Fe2O3.H2O (lepidocrocite) was prepared at the Macaulay Institute for Soil Research by the method of van Schuylenborgh.6 The method employed for preparation of the cold-precipitated hydrated ferric oxide* was as follows: 1: 1 NH4OH was added (moderately fast with rapid stirring) to a solution of 5 grams FeC13.6H2O in 250 ml distilled water at room temperature (ca. 15°C) until red litmus turned blue, and then 2 to 3 drops in excess. The resultant precipitate was filtered off on a Buchner funnel and washed free from Cl- with distilled water. Previous experiments having shown that drying over P2O5 at room temperature or drying at 105°C in the oven did not significantly affect the thermogram,5 samples were nor-

Jan 1, 1952

By W. A. Backofen, N. E. Paton

DEFORMATION twinning contributes to the plasticity of titanium and has been reported on {10i2}, {1121},{11221}, {1123}, and {1124} planes.' Twinning on {1011} has also been reported in commercially pure titanium as a result of quenching from 1000', but no reports of such twins occurring as a result of deformation alone are known to the authors. In this study single crystals grown from iodide titanium were compressed parallel to the c axis over the temperature range 25 to 800°C. From it has come evidence that (10il) deformation twins are the most common type between 400" and 800°C. Crystals about 6 mm high with a rectangular cross section measuring about 3 by 4 mm were used in the tests. Chemical analyses showed them to contain 220 to 250 ppm 0 arid 10 to 15 ppm N. The long axes of the crystals were made parallel to the c axis while the sides were parallel to (i100) and (1120), as shown in Fig. 1. The crystals were compressed between polished alumina platens with no lubrication, in a vacuum of 2 x lom5 torr. Between 25' and 300°C the predominant twinning mode was (11221, which is to be expected for this stress state. There was an abrupt change at 400°C, however, and from that temperature up to 800°C twinning took place almost entirely on (10i1 ). Twin planes were determined by the two-surface rnethd. Measurement of the crystal orientation within several (10i1) twins by means of polarized-light4 and transmission electron microscopy showed that the lattice reorientations were the same as those previously reported for the (10i l ) twinning in magnesium.' This reorientation is consistent with The calculated (10il ) twinning shear for titanium is 0.10, while that measured from surface offsets on a (1150) face, which is parallel to the shear direction, was about 0.15. Slip within the twins, or in the matrix adjacent to the twins, probably accounts for the larger measurement. Slip in both locations was judged to have occurred from observations on foils in the electron microscope.

Jan 1, 1970

By J. H. Jackson, R. B. Fischer

VERY little is known about the properties of relatively pure refractory metals in the cast state since these metals are customarily made by powder-metallurgy methods. Recently, the development of the vacuum arc-melting furnace has made possible the study of large, sound ingots of cast molybdenum. The ductility of molybdenum, at room temperature, may vary considerably. Often the brittle state depreciates the utility of molybdenum, even for the material prepared by powder-metallurgy techniques. At the outset of this program, it was felt that if the causes of brittleness were to be determined it might be possible to reduce the difficulties met in the working of molybdenum. Also, the problem of producing strong, fusion weldrnents of molybdenum might be solved. The conditions required to produce ductile molybdenum by the powder-metallurgy technique are reasonably well known. A fibrous or elongated-grain structure has good ductility although this property is probably directional. The equiaxed structure obtained in fully recrystallized molybdenum has been considered to be brittle. The reasons for this behavior of molybdenum have not been established. The preparation of cast molybdenum ingots by the vacuum-arc method was described by Parke and Ham.' By this method, ingots weighing up to 250 Ib were produced." Parke and Ham1 found that molybdenum ingots containing more than approximately 0.0025 pct oxygen could not be hot forged. On the other hand, they found that fully deoxidized cast molybdenum, containing up to 0.06 pct carbon, could be hot forged. However, the as-cast ingots containing either carbon or oxygen were brittle at room temperature. Parke and Ham,1 Woodside,3 and Zapffe, Landgraf, and Worden4 noted the inter-granular precipitation of oxide or carbide in cast molybdenum. This study—part of an extensive program on mo-lybdenum—had the following objectives: (1) To investigate the causes for the brittleness of cast molybdenum at room temperature, (2) to investigate the means for improving the ductility of cast molybdenum at room temperature, and (3) to add to the knowledge of cast molybdenum and of cast metals in general.

Jan 1, 1951

By Willard Ide Pierce

CONSIDERING the extent of the gold-fields of Nova Scotia, which occupy an area of 6000 to 7000 square miles, a few words as to the cost of extracting and reducing the ores may prove of interest. The veins are for the most part narrow, 4 to 12 inches being the usual width, while a 2-foot vein is considered a large one; but they offer many advantages, inasmuch as the ore is easily treated, wages are low, and machinery and supplies can be transported by water to within from 1 to 15 miles of the different mining districts. The Salmon River vein is a notable exception to the above rule as to width. It ranges from 4 to 20 feet in width. The Hall-Anderson of Fifteen-Mile Stream (where there are several veins) is another, which has proved to be from 3 to 4 feet in width. The ownership of all mining claims is retained by the Crown, but persons wishing to work them can lease as many unoccupied areas as they please for twenty-one years, with option of renewal, on the payment of two dollars for each area, and a royalty of 2 per cent. on all gold extracted. Each area consists of 150 Get along the lode and 250 feet across, upon which a certain amount of work must be done, which by recent law is reduced from one hundred to forty days' work for each area every year. Any person erecting in a district the first mill, having at least 8 stamps, and being at least 10 miles from any other mill, is entitled to a lease of 10 areas, free from advance-payment and from royalty, for twenty-one years. A record must be kept, and filed every quarter at the office of the Commissioner of Mines, of the total number of days' work done on the areas, and a monthly statement of the number of tons crushed, and the actual yield of gold. Gold-mining in Nova Scotia may be said to have commenced with the finding of a nugget of gold in a brook at Tangier in 1858. The following table, taken from the official report," gives the total yield

Jan 1, 1885

By K. P. Fournier

This report describes work on the problem of predicting oil recovery from a reservoir into which water is injected at a temperature higher than the reservoir temperature, taking into account effects of vis cosity-ratio reduction, heat loss and thermal expansion. It includes the derivation of the equations involved, the finite difference equations used to solve the partial differential equation which models the system, and the results obtained using the IBM 1620 and 7090-1401 computers. Figures and tables show present results of this study of recovery as a function of reservoir thickness and injection rate. For a possible reservoir hot water flood in which 1,000 BWPD at 250F are injected, an additional 5 per cent recovery of oil in place in a swept 1,000-ft-radius reservoir is predicted after injection of one pore volume of water. INTRODUCTION The problem of predicting oil recovery from the injection of hot water has been discussed by several researchers. l-6,19 In no case has the problem of predicting heat losses been rigorously incorporated into the recovery and displacement calculation problem. Willman et al. describe an approximate method of such treatment. 1 The calculation of heat losses in a reservoir and the corresponding temperature distribution while injecting a hot fluid has been attempted by several authors.7'8 In this report a method is presented to numerically predict the oil displacement by hot water in a radial system, taking into account the heat losses to adjacent strata, changes in viscosity ratio with temperature and the thermal-expansion effect for both oil and water. DERIVATION OF BASIC EQUATIONS We start with the familiar Buckley-Leverettg equation for a radial system:* This is sometimes referred to as the Lagrangian form of the displacement equation. If the injection water is at some temperature Ti which is above reservoir temperature TR, then lw, the flowing fraction which is water, is a function of both water saturation and fluid temperature at any given time and position in the reservoir. The temperature dependence of lw results from the temperature dependence of the viscosity ratio µo/µm of which lw is a function. Thus:

Jan 1, 1966

By L. W. Bahney

MANY methods for the determination of gold or silver, or both, in cyanide solutions have been published, which with care in manipulation, and modification in some cases, will give results that are satisfactory. It is possible to classify or group these methods as follows: 1. Evaporating the solution in a porcelain or agate-ware dish containing litharge,1 or in a "boat" made of lead foil. 2. Forming a lead sponge containing the precious metals by means of zinc shavings,2 zinc dust,3 or a piece of aluminum.4 3. Decomposing the cyanide solution with an acid and precipitating the precious metals by the use of one or a combination of some of the following: Copper sulphate,5 sodium sulphite,6 hydrogen sulphide,7 cement copper.8 4. Precipitating the silver by zinc dust held in a Hooch crucible and determination with a standard solution of sulphocyanate.9 5. Electrolysis.10 6. Colorimetry.11

Jan 2, 1915

By R. R. Joseph, K. A. Gschneidner

The solid solubility of magnesium in the close-packed modifications of lanthanum, cerium, praseodymium, neodymium, gadolinium, dysprosium, and lutetium was determined from approximately 250°C to the eutectoid temperature by the X-ray parametric technique. The maximum solubility at the eutectoid temperature was found to be 9.4 at. pet in La, 8.2 in Ce, 9.2 in Pr, 8.2 in Nd, 14.1 in Gd, 13.5 in Dy, and 11.5 in Lu. The solubility when compared at a homologous temperature was found to increase in a smooth manner with decreasing size factor from lanthanum to lutetium. It was found that the apparent atomic radius of magnesium when dissolved in the lanthanide solvent was constant and could be accounted for by consideration of several models dealing with departures from Vegard's law. The eutectoid temperatures for the lanthanide-rich La-Mg, Ce-Mg, and Nd-Mg systems were found to be 544", 505°, and 551°c, respectively. A study of the solid solubility of magnesium in the lanthanide metals was initiated because the valency and electron egativity difference are essentially constant and favorable for solid-solution formation, and thus one could examine the effect of size on terminal solid solutions. The size factors based on metal radii range from 14.6 pet for lanthanum to 7.5 pet for lutetium. Since the electronegativity difference is small, 0.06 unit or less, application of the Hume-~othery' and Darken and Gurry2 criteria for solid-solution formation indicates that the solubility of magnesium would be expected to be small for lanthanum (between 5 and 10 at. pet) and to increase as we proceed along the lanthanide series. A search of the literature revealed no data on the solubility of magnesium in the lanthanide metals except for the solubility of magnesium in cerium at 450oCA typical phase diagram for the lanthanide-rich end of the lanthanide-magnesium phase diagram is shown in Fig. 1. EXPERIMENTAL PROCEDURES Preparation of Alloys. The lanthanide metals used in the preparation of alloys for this investigation were prepared from the corresponding fluoride by metallothermic reduction techniques previously described by Daane.4 The chemical analyses of the metals used are listed in Table I.

Jan 1, 1965

By L. R. Raymond, J. L. Hudson

This paper proposes a comparatively short-term (8 to 10 hours) well test for detecting and characterizing well-bore damage and for measuring mean formation permeability. The proposed test is made by injecting fluid at constant pressure, recording injection rate as a function of injection time. After one to four hours of injection, the well is shut in and fall-off of bottom-hole pressure is obtained as a function of shut-in time. Formation permeability is estimated by an iterative technique. First, a value of formation permeability is assumed. Then, a plot of the recorded injection rate as a function of dimensionless time is made, using the assumed pertneability value. From the slope of the injection-rate curve. a new value of formation permeability is calculated. If the new value agrees with the original assumed value, the assumption was the correct formation permeability. If the values do not agree, the process is repeated using the new permeability value in the calculation. Convergence is rapid, and a reliable permeability value results. Pressure fall-off data are used to check the result. Graphs of pressure and injection rate us functions of time given in the paper show that changes in permeability of the formation in the neighborhood of the wellbore are disclosed by this technique. Thus, the short-term test can he used to detect formation damage. Also, a rough measure of the radial extent of damage can be inferred, which is helpful in designing stimulation treatments. The mathematical model used for this work was a single-zone, horizontal reservoir with a damaged zone in which permeability decreased continuously as radial distance to the wellbore decreased. This model is more realistic than the usual two-zone, discontinuous permeability model used in published works; calculations indicate the realistic model is valid. Vertical variations in horizontal permeability were studied with this model, and results indicate that the permeability measured by the short-term test is the mean horizontal permeability for the vertical interval tested. The proposed short-term test thus should be useful in detecting and characterizing formation damage and in measuring formation permeability needed in calculating reservoir transmissibility. INTRODUCTION To plan the most efficient production or injection schedule for a well and to design or evaluate the optimal stimu- lation treatment, it is necessary to know the properties of the reservoir adjacent to the well, particularly the reservoir transmissibility and characteristics of a damaged zone, if one exists. Several techniques for determining reservoir transmissi-bility from well tests have been presented in the literature. 1,2,3,4 All these techniques rely on conducting constant-rate well tests that often are difficult to execute. A constant-pressure well test is generally easier to carry Out. and this paper contains the first available method for the analysis of constant-pressure well tests. Determination of wellbore damage from transient well tests has been the subject of several papers."" From these studies it is apparent that information necessary for determination of the characteristics of a damaged zone is available shortly after the transient well test is initiated. Consequently, it may not be necessary to carry out an extensive well test (for example, a pressure build-up test) if the primary purpose of the test is to detect the existence of wellbore damage. All previous studies of well testing to determine wellbore damage have been based on the two-zone perrneability model. In this model the damaged zone has a permeability k,, extending to a radius r,,, and the formation permeability k obtains from r, to the drainage radius r,.. Consequently, there is a discontinuity in permeability at r = r,,. This discontinuity can be eliminated by assuming a continuous variation in permeability through the damaged zone. The effect of this assumption on transient well tests is discussed in following sections of this paper. In addition, all formations have within them vertical permeability variations associated with lithology changes throughout the zone of interest. This paper also considers the effect of these variations on transient well tests. ANALYSIS OF CONSTANT-PRESSURE WELL TESTS The mathematical analysis associated with the injection of fluid at constant wellbore pressure into a single-zone, horizontal reservoir completely filled with a fluid of small and constant compressibility and constant viscosity is given in Appendix A. In this analysis it is assumed that the well is located at the center of an undamaged, circular drainage area. From this analysis, the formation permeability can be obtained as follows. 1, Estimate a value for the formation permeability k. 2. Prepare a plot of injection rate q vs

Jan 1, 1967

By E. R. Gilbert, D. E. Munson

Creep of copper under 75 to 1.50 kg per sq cm stresses at temperatures near the melting point was found to he a complex reaction controlled by three mechanisms acting in parallel. In order of appearance with decreasing temperature, the auerage activation energies, Qc , are 168, 79, and 24 kcal per mole. Stress dependence of the minimum creep rate was found to he an exponential for the two high-Qc processes and a power law for the low-Qc, process. Transition of control occurs from one mechanism to another. The relative transition temperature depends upon the applied stress, and the range oiler which the transition occurs depends upon the difference in the activation energies of the mechanisms. The creep behavior at high temperature is explained by the climb of dislocations through thermally or mechanically formed jogs, CREEP in pure fcc metals at temperatures in excess of one half the absolute melting point is normally controlled by dislocation c1imb.1,2 A climb model which seems applicable, according to an extensive analysis of data,3 was derived by Weert-man.4 This model assumes jog-saturated dislocations and predicts an activation energy, Qc, nearly equal to that for self-diffusion, USd. Although the requirement of jog saturation is restrictive, agreement between theory and experiment seems adequate. Many other theoretical treatments, including an early model by Mott,5 include detailed consideration of jog formation as an initial requirement for climb. These models predict activation energies for creep which differ from those of self-diffusion. Seeger6,7 postulates an observed activation energy related to the stacking-fault energy. Thus, Usd <Qc<5Usd + Uj where Uj is the jog-formation energy. Seeger incorporated qualitatively the influence of the relative numbers of thermal and athermal jogs. Expanding this concept, Shoeck8 explicitly states a function based on formation mode and relative numbers of vacancies: e r ci exp {-uf/k T} exp {-Um/k T} [ 1 ] where Uf and Um are energies of vacancy formation and migration, respectively. The concentration of jogs, Cj, depends upon the manner of jog production. For intersection jogs, Cj is not sensitive E. R. GILBERT, Junior Member AIME, formerly with De- to temperature; for thermal jogs, Cj is proportional to exp {—Uj /k T}. Schoeck regards each mode as a distinct mechanism; therefore, the mechanisms may act together.299 The diversity predicted by theory, surprisingly, has not been substantiated by experimental results. A significant investigation must include the extremes in stacking-fault energy. Extensive creep studies of aluminum10 and nickel,11,12 high stacking-fault energy metals, have been made. Comparable studies on a low stacking-fault energy metal, such as copper, have not. It is the purpose of this paper to report the results of an investigation of the creep of copper under conditions which favor thermal jogs. EXPERIMENTAL Cylindrical compression creep specimens (0.240 in. in diameter by 0.400 in. long) were machined from cold-drawn rods of electrolytic tough-pitch copper containing 0.0007 Mg, 0.002 Fe, 0.001 Ni, 0.0005 Ag, Cd < 0.005, and Pb < 0.005 wt pct impurities. Undetected spectrographically was a nominal 0.04 wt pct 0, which occurs as a Cu2O constituent distributed discontinuously at grain boundaries. Vacuum annealing at 900°C for 15 min produced a stable 0.03-mm average grain diameter. Testing was carried out using apparatus similar to that described by sherby,13 modified by enclosing platens and a push rod in a vacuum cylinder. Normally this arrangement resulted in pressures less than 10 Only a slight surface tarnish, less than 0.0005 in. in thickness, occurred during the test. The applied stress, corrected for atmospheric pressure, was maintained within 2 pct of the desired true stress by the addition of lead shot at fixed strain increments. Creep strain was measured with dial gages as a relative displacement of the upper and lower platens; accuracy of measurement was 0.0001 in. Two creep-test methods were used, the differential or cyclic temperature14 and the isothermal, to obtain creep data at stress levels of 150, 100, and 75 kg per sq cm over the temperature range of 620° to 1032°C. Minimum creep rates were used from both test methods; this was considered proper because comparable temperature tests or cycling to the same temperature gave the same creep rate, within experimental error. The cold vacuum test chamber, with the unstressed specimen in place, was heated to temperature by placing a preheated furnace over the chamber. Temperature equilibrium was attained within 30 min. For the cyclic tests, the stress was removed during the 5 to 10 min necessary to effect the temperature change and re-

Jan 1, 1965

By Hollis L. Caswell

The present status of integrated circuits utilizing. superconductive switching. elements is reviewed with special attention given to fabrication techniques, methods for interconnecting completed circuits, and refrigeration requirements. Cryoelectronics has been largely an "inte- grated-circuit" technology since its conception because the switching speed of superconductive devices is attractive only when these devices are fabricated with thin-film techniques. It is true that cryotron circuits can be constructed from wires of appropriate materials (as indeed was done by Dudley Buck 1 in his early investigations) but these circuits will switch in times characteristic of milliseconds whereas similar circuits fabricated by thin-film methods have potential switching times of nanoseconds. Furthermore, cryo-electronic devices such as the cryotron lend themselves readily to fabrication by thin-film techniques since these components may be made from polycrys-talline thin films and are relatively insensitive to the presence of impurities (as measured by semiconductor standards). Therefore, during the past decade considerable effort has been devoted to developing techniques for batch fabricating circuit arrays containing superconductive switching elements. Technology had developed to the point several years ago that fabrication of cryoelectronic arrays containing up to one hundred devices was rather straightforward. However, larger arrays containing between lo4 and 106 components which are required for commercial development of cryoelectronics still pose very severe yield problems. Thus in a sense cryoelectronics found itself in 1962 at the point semiconductor technology finds itself today; namely, individual devices and small groups of integrated devices could be fabricated with acceptable yield and the outlook for building larger integrated-circuit arrays was bright. Unfortunately, problems associated largely with yield have made fabrication of these larger arrays difficult. Unlike semiconductor technology, cryoelectronics had to solve the problems of large-scale integration before it could become economically attractive. This has proven to be a sizable burden to bear. Since several reviews exist on superconductivity,2 superconductive devices,3 and cryoelectronic technology, no attempt will be made in this paper to summarize these areas. Instead a few specific topics will be dealt with in more detail. First, a brief description is given of selected superconducting switching and storage devices with special attention to several metallurgical techniques which improve the performance of these devices. Second, techniques used to fabricate cryoelectronic devices are described with emphasis on problems affecting yield. Third, techniques for interconnecting a number of cryoelectronic planes are described. And last, refrigeration of cryoelectronic components is discussed briefly since the low operating temperature of superconductive devices is an important consideration in this technology. SUPERCONDUCTING STORAGE AND SWITCHING DEVICES The basic superconductive switching device is the thin-film cryotron. The geometry of this device is attractively simple, since it involves only the intersection of two lines that are electrically insulated from each other. The switching element (gate) and control element (control) of a crossed-film cryotron are arranged as illustrated in Fig. 1. The material for the gate is selected to permit the gate to be switched from the superconducting to the normal (resistive) state by the application of a control current. Tin, which has a critical temperature (T,) of 3.7°K, is commonly used for the gate and the cryotron is operated at a temperature just below T, (for example, 3.5°K). The control material (normally lead, with T, = 7.2°K) is chosen so that the control is never driven normal during circuit operation. To improve cryotron operation, a ground plane, also of lead, is placed under all of the circuitry to act as a diamagnetic shield and improve the current-density uniformity across the width of various thin-film elements. Normally, line widths vary from 0.005 to ^ 0.020 in. and film thicknesses from 300 to 10,000A, although new fabrication techniques make narrower lines feasible. In fabricating cryotrons it is important that the edges of the gate elements be geometrically sharp to avoid undesirable switching characteristics associated with a thinner edge region, Fig. 2. One technique which has been used extensively to form patterns consists of placing a physical mask containing the film pattern between the evaporation source and the substrate and depositing through the mask. Film strips formed in this manner possess a penumbra at the film edges due to shadowing of the evapor-ant under the mask. Several techniques have been proposed for minimizing effects due to this penumbra. One of the more promising metallurgical techniques

Jan 1, 1967

By D. Turnbull, J. H. Hollomon, J. C. Fisher

Application of the concepts of nu-cleation and growth to the analysis of experimental transformation data has led to valuable descriptions of phase transformations, an outstanding example being the transformation austenite —* pearlite which has been examined with particular care by Mehl and co-workers.&apos;-5 In addition to the pearlite transformation, the proeutectoid fer-rite and proeutectoid carbide transformations are known to proceed by nucleation and growth. Martensite, on the contrary, until recently was thought to form by a mechanism involving neither nucleation nor growth; however, extension of standard nucleation theory6 suggests that martensite, bain-ite, and the other products of austenite decomposition all grow from nuclei in the parent phase. The theory that martensite forms by nucleation and growth is strongly supported by recent experimental work of Kurdjumov and Maksimova.7 The concepts of nucleatioli and growth have been fruitful also in providing a sound basis for quantitative theoretical treatments of the kinetics of phase transformations. For example, Volmer and Weber8 and Becker and Döring9 developed the theory of nucleation from fundamental considerations to a point where excellent agreement was obtained with the results of experiments on the condensation of supercooled vapors. As a result of their analysis, the kinetics of vapor-liquid transformations now can be predicted. It seems probable that application of the theories of nucleation and growth to a quantitative study of austenite decomposition similarly will clarify the nature of the individual transfor: mations involved, and will enable the calculation of austenite transformation kinetics. In the present paper, the theories of nucleation and growth are applied to the austenite ? martensite transformation in steels. The analysis begins with a discussion of nucleation in single component systems. Martensite appears to be coherent with the parent austenite, hence the nucleation theory is modified to include the effects of elastic distortion. Nucleation in the two component iron-carbon system then is discussed, for most steels are primarily alloys of these two elements. Finally, M. temperatures and martensite transformation curves are calculated for each of several alloy steels of varying carbon and chromium content, and are compared with those determined experimentally by Lyman and Troiano10 and Harris and Cohen.11 Nucleation Theory NUCLEATION IN SINGLE COMPONENT SYSTEMS6,12-14 The work required for reversible formation of a region of phase within the parent a phase is expressed conveniently as the sum of two terms: W1 = Aa, the product of the area of the interface and the interfacial free energy, and W2 = VAf, the product of the volume of the region and the free energy increase per unit volume associated with the transformation. The total work is therefore W = Aa + VAf. When a is more stable than ß, Af is positive and W increases without limit as the volume increases. The transformation a ?ß does not occur. It is nevertheless true that small regions of phase ß enjoy temporary existence here and there in the a. The equilibrium number of ß regions of given size is proportional to exp(— W/kT) per unit volume of a, assuring that larger (ß regions occur with diminishing probability. When a is less stable than ß, Af is negative and W passes through a maximum as V increases. Assuming for simplicity that regions of ß are spherical, as is true when the interfacial tension is isotropic and there are no elastic strains, W = 4r2a + (4/3)*r3Af The maximum value of W is W* = 16iro3/3Af2 [1] for regions with radius r* = -2o/Af. [2] For single component condensed systems it has been shown14 that the steady rate of nucleation of 0 per unit volume of untransformed a is nearly proportional to exp[- (W* + Q)/kT] where Q is the activation energy for the unit processes of adding or removing one atom from an embryo or nucleus. If To is the temperature at which a and ß are in equilibrium, the rate of nucleation is a maximum at a temperature 0 < T < To where (W* + Q)/kT is a minimum. P regions smaller than critical size are called embryos; they tend to grow smaller and disappear, only exceptionally growing larger. Regions equal to or larger than critical size are called nuclei. A critical size nucleus may grow indefinitely large or may shrink back to a, either process decreasing the free energy of the region.

Jan 1, 1950

By Y. Nakada, U. F. Kocks, B. Chalrners

THE orientation dependence of work hardening has previously been studied over the entire range, i.e., including special orientations of high symmetry, in aluminum1-3 and silver.* The differences between various orientations are substantial, and the trend is the same in all fcc crystals. However, there are quantitative differences in behavior between aluminum and silver at room temperature, particularly in the (100) orientation. While many experiments on gold have been reported,5&apos;9 none included the special orientations. We therefore undertook tension tests on gold crystals of the axial orientations (100),(110), (Ill), (211), and, as a representative of single slip, one whose Schmid factor was 0.5 (hereafter referred to as m = 0.5). Single crystals of dimensions 1/4 by .1/4 by 6 in. were grown from the melt1&apos; at a rate of 4 in. per hr, using gold of 99.99 pct purity obtained from Handy and Harman. A growth rate of 1/8 in. per hr, or a purity of 99.999 pct,ll produced no difference in results. The crystals were annealed at 1000°C in air for 24 hr and furnace-cooled down to room temperature, after which they were electro-polished in a solution of potassium cyanide (40 g), potassium ferrocyanide (10 g), soda (20 g), and enough distilled water to make 1 liter of solution, at a current density of 0.02 amp per sq mm. After 2 or 3 hr, a very smooth surface was obtained by this method. Nine m = 0.5, two (loo), one (110), three (Ill), and one (211; crystals were tested at room temperature in a floor-model Instron machine at a tensile strain rate of 0.5 pct per min. The accuracy of the stress measurement was k10 g per sq mm up to 500 g per sq mm, k2 pct for higher stresses. The strain measurement was accurate to t2 pct. The scatter of stress at a given strain among the crystals of the same orientation was small, *5 pct being the largest. The representative stress-strain curves for various orientations are shown in Fig. 1. Table I summarizes the work-hardening parameters as used by seeger.12 Results of other investigators are also included in this table for comparison. There are no previous data for the corner orientations. Values for m = 0.5 crystals agree fairly well with those of Berner. Tm is defined as the stress at which the stress-strain curve begins to deviate from linearity of Stage 11. However, in practice this is a very difficult value to estimate because each investigator has a different idea of linearity. Therefore, the comparison with the values of other investigators may not be valid. In the present experiment, (100) crystals had the highest 111, followed by (111) crystals. The work-hardening rate in Stage I1 was highest for the (111) crystal followed by the (100) crystals. Our value for 0x1 of m = 0.5 orientation agrees very well with those of other investigators. 1) Single-Slip orientation (m = 0.5). These crystals were oriented so that the primary-slip vector was contained in one side face. The dimension perpendicular to this side should then not change if single slip indeed takes place. Within the accuracy of measurement, this dimension did not change during the deformation. Since the accuracy is k0.2 pct, the amount of secondary slip is less than 0.7 pct of the slip on the primary-slip system at 30 pct tensile strain. This is in conformity with the results obtained by Kocks" for aluminum crystals. The tensile axis moved toward (211) during the deformation. Slip bands, Fig. 2(a), were very fine and closely spaced. Some deformation bands were observed. There were no clear-cut cross-slip traces such as the ones observed on aluminum m = 0.5 crystals. 2) (111) Crystals. The orientation of the tensile axis was stable during the deformation. From this observation, one can deduce that at least three slip systems must have operated, and probably all six because the remaining three all have cross-slip relation to one of the first three.&apos; It was very difficult to observe the slip markings. Consequently, we could not confirm by this method that six systems were operative during the deformation. At high strains (above 50 pct shear strain), this orientation had the highest stress-strain curve. At 80

Jan 1, 1964

By O. D. Sherby, J. E. Dorn

SEVERAL years ago Zener and Hollomon1 suggested that the flow stress of metals might be related to the temperature and strain rate in accord with the functional equation: s=s(eeh/rt) [1] at the same state. The Zener-Hollomon relation also contains the significant tacit implication that the energy for activation, AH, is substantially independent of the state of the material. The utility of this method for correlating creep data with tensile data was illustrated in a preliminary reportL on the effect of alloying additions on the secondary creep rates of binary a solid solutions in aluminum as shown for Al-Mg alloys in Fig. 1. Not only do the secondary creep and ultimate tensile properties for each alloy correlate well on this plot but in addition the activation energy, AH = 17,900 R cal per mol. is identical for the various compositions. Similar types of curves correlating the secondary creep rates at 422°K (300°F) with ultimate tensile data were also obtained for a series of dilute a solid solutions of copper, germanium, zinc, and silver in aluminum. In all cases the activation energy for creep, AH, was found to be equal to about 17,900 R cal per mol independent of the type of alloying element or its concentration. The coincidence of the activation energies for these alloys is probably due to the fact that the activation energies for creep are rather insensitive to small composition changes. Alloying additions, as shown in Fig. 1, however, increased the stress necessary to obtain equivalent values of ee?H/RT. Thus the stress level in curves of the type represented by Fig. 1 gives the relative creep strengths of the various alloys. Inasmuch as these correlations between creep and tensile data were obtained for creep tests at 422°K, it appeared advisable to ascertain the range of secondary creep rates and temperatures over which correlations between creep and tensile data could be made on the basis of the Zener-Hollomon relationship. If, for example, the activation energy changes with different ranges of creep temperatures, the utility of the proposed analysis would be severely weakened. But if AH is constant not only for all dilute alloys, but also over wide ranges of temperature and secondary creep rates, confidence will be developed not only in the broad utility of the Zener- Table I. Chemical Analysis&apos;:&apos; and Grain Size of Alloys Mean Grain Residual Impurities. Wt Pct Diam- Alloying Atomic-eter, Element Pct Si Fe Cu Mg Mn mm Aluminum 99.987 0.003 0.003 0.006 0.001 0.25 Magnesium 1.617 0.003 0.004 0.006 0.26 Copper 0.101 0.003 0.003 0.0006 0.001 0.29 Zinc 1.616 0.003 0.005 0.007 0.001 &apos;0.26 * The authors express their appreciation to the Aluminum Company of America for the preparation and chemical analyses of these alloys. Hollomon relationship, but also in its theoretical justification. The following investigation was instituted in order to ascertain the range of validity of the Zener-Hollomon relationship when it is applied to creep and tension data. Materials, Equipment, Technique In view of rather extensive correlatable data now available on the plastic properties of a series of binary a solid solutions of magnesium, copper, germanium, zinc, and silver in aluminum, these alloys were again selected for the present investigation. Previous reports have already covered the tensile properties of these alloys at subatmospheric temperatures," and at elevated temperature," as well as their fatigue properties at atmospheric temperatures, nd their creep properties at 422°K.&apos; Resistivity measurementsa as well as precision lattice constant determinations? evealed that these alloys are a solid solutions. After homogenization. 0.100 in. thick sheets were rolled to 0.070 in. in thickness and recrystallized to about the same grain size. Since the details of these treatments were reported previously," they will not be repeated here. In order to appropriately reduce the scope of creep tests in this investigation only the typical alloys shown in Table I were investigated. All creep specimens were machined such that their tensile axes were in the rolling direction of the 0.070 in. sheet stock. The gage section was 2 in. long and 0.500 in. wide with a reduced section of 3.50 in. Creep strains were measured by means of rack and pinion type extensometers with a sensitivity of 0.005 in. per division. Readings were estimated to 1/10 of

Jan 1, 1953

By J. P. Morris

PREVIOUS investigations1,2 have shown that alloying elements in liquid iron influence the thermodynamic activity of sulphur and thereby affect the partition of sulphur between metal and slag in the desulphurization process. For example, the greater efficiency of desulphurization in the blast furnace as compared to the open hearth can be attributed in part to a higher level of sulphur activity in blast-furnace metal due to the higher concentration of carbon and silicon. In the present investigation, a short study was made of the influence of manganese on the activity of sulphur in liquid iron and iron-carbon alloys. In contrast to carbon and silicon, manganese was found to decrease the activity coefficient of sulphur; and in iron-carbon alloys it counteracts to some extent the influence of carbon. However, at manganese concentrations normally present in the blast furnace or open hearth, the effect of manganese is small. Since manganese sulphide has a limited solubility in iron, manganese can act, under certain conditions as a desulphurizing agent. Considerable data on the manganese-sulphur product in carbon-saturated melts were obtained in the investigation and have been included in this report. The experimental procedure was the same as that used in the earlier investigations on the effect of silicon&apos; and carbon&apos; on sulphur activity. Briefly, the method was as follows: The molten alloy, contained in a graphite or sintered alumina crucible, was brought to equilibrium at a constant temperature with a mixture of hydrogen and hydrogen sulphide of constant composition by bubbling the gas through the metal. Samples of&apos; the melt were taken for analysis at regular intervals by suction through a 2 to 3 mm bore silica tube dipped into the metal. The experiments were run in a graphite spiral resistance furnace with melts weighing 50 to 60 g. The gas bubbling tubes were made of sintered alumina and were 5/16 in. OD, 1/16 in. ID, and 24 in. long. Equilibrium was assumed to have been attained when the sulphur content of the liquid metal reached a constant value. During an experiment there was a rapid loss of manganese from the melt by volatilization. To offset this loss, small additions of manganese were made periodically. The rate of manganese addition needed to maintain a constant manganese concentration was determined in preliminary tests. In all of the experiments, deposits of manganese sulphide formed above the melts in a cooler region of the furnace. Apparently, these deposits resulted from a reaction between manganese vapor and hydrogen sulphide in the gas. To prove that manganese sulphide did not volatilize from the melts to a measurable extent, an experiment was run in which helium was bubbled through liquid iron containing both manganese and sulphur. Although manganese volatilized rapidly in this test, there was no appreciable loss of sulphur. Volatilization of manganese sulphide from a melt would have led to an apparent equilibrium condition in which the sulphur content of the metal was lower than the true equilibrium value. The experimental results are shown in the first seven columns of Table I. The data in the last two columns were obtained from the previous work on the effect of carbon&apos; and show what the results would have been in the absence of manganese but with temperature, gas composition, and carbon content of the metal remaining the same. Comparison of the last four columns show that, in the presence of manganese, the sulphur content of the metal increased at equilibrium and the activity coefficient of sulphur decreased. However, the results show that, for manganese concentrations below 3 pct, the effect of manganese is small. The values for activity coefficient of sulphur given in Table I were calculated from the following relations: S (in liquid metal) + H2 (gas) = H2S (gas) [l] K ph2s/?s X %S X phg = 0.00251 [2] where K is the equilibrium constant for the reaction, PH2S and ph2 are the partial pressures of hydrogen sulphide and hydrogen, respectively, and ?s is the activity coefficient of sulphur. The standard state for sulphur was taken to be a 1 pct solution of sulphur in pure iron. The numerical value for K at 1600°C was determined in the earlier work. For the purpose of showing graphically the results of the tests run at 1600°C, the activity coefficients of sulphur were recalculated so as to correspond to a manganese concentration in the metal of 2 pct in each case. In the calculation it was assumed that the increase in sulphur content of the metal at equilibrium caused by the presence of manganese

Jan 1, 1953

By M. G. Benz

The superconducting performmce of diffusion-processed Nb3Sn is influenced by its micro structure. High isotropic transverse current density may be achieved in this material by a process which forms a precipitate of ZrO, within the Nb3Sn. FOR an ideal type-I1 superconductor, little or no transport current can be carried in the mixed state; i.e., little or no transport current can be carried above the lower critical field H,,, where the field penetrates abruptly in the form of current vortices or fluxoids, even though full transition to the normal state does not occur until the upper critical field H,,.&apos; Fortunately, nonideal type-I1 superconductors can be readily obtained and these carry large transport currents up to the upper critical field H. Both theoretical and experimental investigations have attributed this current-carrying capability for nonideal type-I1 superconductors to pinning of the fluxoid lattice by heterogeneities in the microstructure of the superconducting material. These heterogeneities may take the form of dislocations or dislocation clusters,2"5 grain boundaries: structural imperfections introduced by phase transformations; radiation damage,8"10 or precipitates.11"15 Nb3Sn formed by diffusion processing is a type-I1 superconductor. Heterogeneities are needed for high superconducting critical currents above H,,. This paper will cover: a) what the microstructure of diffusion-processed NbSn looks like; b) what changes in the microstructure take place when the system is doped with precipitates, and c) how these changes in microstructure influence the superconducting critical currents. EXPERIMENTAL Preparation of Samples. Diffusion processing was used to form the Nb3Sn. The procedure used was as follows: a) coat the surface of a niobium tape with tin; b) heat-treat this tape at a temperature above 930°C to form a layer of Nb3Sn at the Sn-Nb interface. Such a layer of NbsSn is shown in Fig. 1 The thickness of the NbsSn layer formed was controlled by the time and temperature of the heat treatment. The same general procedure was used for preparation of both undoped samples and samples doped with a precipitate. An additional step was included in the preparation of the doped samples which consisted of internal oxidation of zirconium to form ZrOn. The details of the doping process will be reported in a later paper. Sample Testing. The Nb3Sn tape samples were soldered to a copper or brass shunt. Current and voltage leads were then attached to the sample in the usual four-probe resistance measurement configuration. The sample was cooled to 42°K. In some cases it was cooled in the presence of a high magnetic field and in other cases with the field turned off. The results were the same for both cases. The samples were oriented in a configuration with field transverse to current but could be rotated such that the angle between the field vector and the wide side of the tape sample could be changed. Measurements up to 100 kG were done in a superconducting solenoid and measurements above 100 kG in a water-cooled copper magnet at the MIT National Magnet Laboratory. Once the test field was reached, the current in the sample was increased until voltage was detected across the sample. The critical current was taken as the current at which voltage was first detected in excess of background noise. In most cases this was 1 to 2 x 10~6 v for a— in.-wide sample carrying several hundred amperes with a in. separation between voltage leads and with a 10 "-ohm shunt resistance. RESULTS AND DISCUSSION Microstructure. Examination of the microstructure of the undoped Nb3Sn shows rather large-diameter (1 to 2 columnar grains growing outward from the niobium surface toward the tin surface. As the layer is made thicker by longer diffusion times, these grains grow longer. Few new grains are started. Transmission electron microscopy shows little or no second-phase material within the bulk of the Nb3Sn layer. The microstructure of a diffusion-processed NbsSn layer changes quite drastically when the system is doped so as to form a precipitate within the NbsSn layer. Instead of large-diameter columnar grains of NbaSn forming, smaller-diameter (0.5 to 1 ) equiaxed grains of Nb3Sn decorated with the precipitate form. Fig. 2 shows a transmission electron micrograph of a Nb3Sn layer doped with zirconium oxide. This layer has been etched so that one may look between the grains

Jan 1, 1969

By K. A. Slagle

Hydraulic analysis of the wellbore has become increasingly inzportant for designing cementing operations and selecting equipment, materials and techniques to complenzent modern well-c-ompletion practices. Non-Newtonian fluid technology has advanced beyond the point where former empirical methods of analysis adequately define the hydraulic system and fluid properties. In view of these factors, this paper describes a series of rheological calculations which have been found practical, through field usage, for assistance in selecting a cementing program. A relatively simple laboratory method using standard viscometric equipment is suggested for determination of the rheological properties of slurries, and clrrta are presented on some of the more common cementitrg conzposition.A. A criterion for divergence from laminar-flow characteristics has been proposed. Usefulness of the calculations is indicated by examples of cementing operations where they have been used. INTRODUCTION With the changing aspects of well-completion practices during the past few years, it has been increasingly important to have a relatively simple method of analyzing the flow conditions existing in the well during cementing operations. This is particularly true in view of the improved economics toward which most of the changes have been directed. Rheological characteristics of slurries used for cementing should be a major consideration in the trend toward smaller casing sizes, either single or multiple strings. Receiving increased attention is the practice advocated in 1948 by Howard and Clark&apos; of attaining turbulent flow with the fluids circulated during a primary cementing operation. While there may still be a difference of opinion concerning this technique, most available information indicates that superior primary-cementing results are generally obtained when high displacement rates are employed. Fluid properties of the slurry to be used must be available, as well as calculation methods, to determine what flow rates should be attained and the probable consequences in terms of frictional pressure and horsepower utilization. It would certainly be inappropriate to attempt high displacement velocities if sufficient pressure might be developed to create lost circulation. Since cementing slurries are non-Newtonian fluids, it is not possible to define their rheological or fluid properties by the single factor of viscosity and then make calculations for the quantities just described. Because the shear stress-shear rate ratio is not constant: it becomes necessary to establish at least two parameters for adequate fluid-flow calculations. It is not the purpose of this paper to delve into the mathematical development of non-Newtonian technology, nor to discuss the arbitrary classification system under which a single fluid may resemble two or three different classes depending upon experimental conditions. Rather, the intention is to present a useful series of calculations based on a concept applicable to both Newtonian fluids and to the preponderance of non-Newtonian fluids encountered in the oil-producing industry. Development of this approach was begun some 32 years ago,&apos; and has most recently been brought to fruition by Metzner and his co-workers at the U. of Deleware. Some non-Newtonian fluids encountered in the petroleum industry, other than cementing slurries, have also had the benefit of this method of analysis."&apos; The two parameters required to define the fluid are usually denoted by the symbols n&apos; and K&apos; and, for the purposes of this discussion, are called "flow behavior index" and "consistency index", respectively. These two slurry properties permit calculation of the Reynolds&apos; number and the "critical" velocity, or the velocity at which departure from laminar flow begins. EXPERIMENTAL DETERMINATIONS The two principal instruments used for rheological studies are the pipeline (capillary-tube) viscometer and the rotational viscometer. When conveniently possible, a capillary-tube viscometer (where the pressure drop and flow rate of the material can be measured) is the better method for rigorous determination of the flow behavior index and consistency index for non-Newtonian fluids. With pressure-drop data at various flow rates, it is then possible to prepare a logarithmic plot of shear rate as the abscissa-shear stress as the ordinate. For fluids which do not exhibit time-dependency, these data will usually produce a straight line. The flow behavior index n&apos; represents the slope of this line, while the consistency index K&apos; becomes the intercept of this line at unity shear rate in accordance with the mathematical derivations associated with this concept of rheology. Due to the difficulties anticipated in maintaining a uniform, pumpable cement slurry for the time interval required to obtain measurements from the pipe viscometer, the n&apos; and K&apos; data reported herein were obtained using a direct-indicating rotational viscometer (Fig. 2). The

By Henry Howe

THIS paper is a discussion of the extremely valuable one of Mathewson and Phillips, The Recrystallization of Cold-Worked Alpha Brass on Annealing,1 which not only gives us a wealth of important data reached with great intelligence, but also shows both uncommon powers of imagination and a perfectly fair spirit. Let me try to analyze some of their results, and to enunciate some of the laws to which these results point. For brevity I refer to them as "The Authors." The page numbers in parentheses refer to their paper. 1. The Visible Aspects of Recrystallization.-Plastic deformation in a single direction, as in cold-rolling, draws the grains out in that direction, that is, inequiaxes them, without destroying their apparent individuality. Each grain seems to endure the drawing out without changing either its volume or its orientation materially, for the etching tint of each grain seems to be as uniform after as before the drawing out, save for the "etch bands,"2 shown in Figs. A to D of the Authors&apos; Plate I. As the drawn-out metal is heated progressively, on reaching a certain "disintegration temperature" these old grains seem to break up, probably into submicroscopic fragments, and this apparent disintegration is followed by a coalescence of these fragments, so that their number decreases and their size increases, till they may become very much coarser than before the drawing out. Ruder3 coarsened the grains of silicon steel in this way till single grains reached an area of 50 cm. (7.75 sq. in.). The deformation itself probably breaks each grain up into many very small fragments, but the fact that each grain continues to have a uniform etching tint different from that of its neighbors indicates that these fragments retain the initial common orientation of the grain. The apparent disintegration which we observe, then, is rather an abandonment of community of orientation for individuality of orientation, by the fragments formed earlier during the deformation. Hence we

Jan 10, 1916

NEW MEMBERS The following list comprises the names of those persons who became members during the period Oct. 10, 1916 to Nov. 10, 1916. ANCHOR, HANS C., Supt., Dome Extension Mines Co., Ltd., South Porcupine, Ont., Canada. BABCOCK, EARLE J., Dean, College Of Min. Engrg., State Univ. Of N. Dak., Grand Forks, N. Dak. BIDDIS0N, P. McDONALD, Construction Engr., The Ohio Fuel Supply Co., 52 W. Gay St., Columbus, 0. BOISE, PAUL T., Asst. Mgr., Stimpson Equipment Co., Felt Bldg., Salt Lake City, Utah. BOURNE, RALPH H., Vice-Pres Whiting Foundry Equipment Co., Harvey, Ill. BROWN, EDWARD OTTO FORSTER, Min. Engr., 706-707 Salisbury House, London, E. C., England.

Jan 12, 1916

Born, March 28, 1853 Died July 23, 1916 On July 25, 1916, the Board of Directors adopted the following minute and directed that it be entered upon the records of the Board & published in the Bulletin of the Institute, and thereof be sent to the family of Mr. Kirchhoff: Charles Kirchhoff was one of the oldest, most widely known and most highly esteerned members of the American Institute of Mining Engineers. On his return in 1875 from a European course of study, to begin a highly successful professional career, he joined the Institute, then just commencing to acquire both national and international recognition. He became at once, and-for more than 40 years continued to be, its earnest , and active supporter, serving it as Manager, Councilor or Director for terms aggregating 9 years,-as Vice-President for 2 years, and twice as President for a year. During the critical period of the erection of the United Engineering Society&apos;s building, and the transfer of the Institute to permanent head-quarters therein, Mr. Kirchhoff was its efficient representative in the United Engineering Society, and on the Building Committee. In 1910, amid the congratulations of colleagues, employees, and innumerable friends, Mr. Kirchhoff retired from active business as Vice- President and General Manager of a great technical publishing company. Yet he sacrificed of his well-earned leisure yet one more year, to serve the Institute as President through 1911.

Jan 9, 1916