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By John Wulff, J. H. Hollomon

Service failures in brass condenser tubes are often due to corrosion. One of the commonest types of corrosion reveals a surface structure of redeposited copper.' The study of the effect of alloy addition agents in preventing this kind of corrosion has received impetus from the extensive publication of Bengough and May.2 The mechanism of such inhibition gauged from immersion corrosion,2 cell corrosion,3 and electrolytic potential4 measurements is controversial. This controversy may be attributed in part to a lack of definite knowledge5 of the actual nature of the initial films formed in corrosion experiments. In our own corrosion work on copper and 70-30 brass this lack became apparent and necessitated an examination of surface films by X-ray and electron diffraction methods. The present paper is but an introduction to the subject of inhibition and reports the structure of films formed in active corrodents. Experimental Procedure The arsenical coppers and brasses used in this investigation were made by one of the writers* in the plant of the Revere Copper and Brass Co. at Rome, N. Y. Standard commercial melting practice was followed. O.F.H.C. copper and electrolytic zinc were used and all arsenic additions were made by additions of a master copper alloy. The final analyses are given in Table I. Likewise, the extrusion, cold-rolling, and intermediate anneals employed followed standard usage. Final heat-treatments were given to all materials to produce a similar grain size (0.035 to 0.045 mm. for brass and 0.015 to 0.025 for copper) for the alloys containing various percentages of arsenic. Samples 1/2 by M by 1/4 in. were cut from sheet material resurfaced and finished with 0000 metal-lographic papers. The surfaces were all degreased in alcohol and benzine before immersion. After immersion in the corrodent, they were washed gently in tap water, dipped first in alcohol and then in pure benzine and allowed to dry in the diffraction apparatus. The reflection electron diffraction apparatus used has been described by Burwell and Wu1ff.6 Fig. I is typical of the photo-grams obtained. The X-ray apparatus was

Jan 1, 1941

By John Wulff, J. H. Hollomon

Service failures in brass condenser tubes are often due to corrosion. One of the commonest types of corrosion reveals a surface structure of redeposited copper.' The study of the effect of alloy addition agents in preventing this kind of corrosion has received impetus from the extensive publication of Bengough and May.2 The mechanism of such inhibition gauged from immersion corrosion,2 cell corrosion,3 and electrolytic potential4 measurements is controversial. This controversy may be attributed in part to a lack of definite knowledge5 of the actual nature of the initial films formed in corrosion experiments. In our own corrosion work on copper and 70-30 brass this lack became apparent and necessitated an examination of surface films by X-ray and electron diffraction methods. The present paper is but an introduction to the subject of inhibition and reports the structure of films formed in active corrodents. Experimental Procedure The arsenical coppers and brasses used in this investigation were made by one of the writers* in the plant of the Revere Copper and Brass Co. at Rome, N. Y. Standard commercial melting practice was followed. O.F.H.C. copper and electrolytic zinc were used and all arsenic additions were made by additions of a master copper alloy. The final analyses are given in Table I. Likewise, the extrusion, cold-rolling, and intermediate anneals employed followed standard usage. Final heat-treatments were given to all materials to produce a similar grain size (0.035 to 0.045 mm. for brass and 0.015 to 0.025 for copper) for the alloys containing various percentages of arsenic. Samples 1/2 by M by 1/4 in. were cut from sheet material resurfaced and finished with 0000 metal-lographic papers. The surfaces were all degreased in alcohol and benzine before immersion. After immersion in the corrodent, they were washed gently in tap water, dipped first in alcohol and then in pure benzine and allowed to dry in the diffraction apparatus. The reflection electron diffraction apparatus used has been described by Burwell and Wu1ff.6 Fig. I is typical of the photo-grams obtained. The X-ray apparatus was

Jan 1, 1941

By C. E. Birchenall

DaVIES and Richardson1 have measured composition changes for Mn1-Owith variation in the equilibrium partial pressure of oxygen at 1500°, 1575°, and 1650°C, where 6 is the deviation from the simple stoichiometric ratio. No trend in this relationship with temperature was observed, so these remarks apply to the whole temperature range. The experimental values and their spread are shown by the bounded straight-line segments in the figure. Davies and Richardson assume that the stoichiometric oxide shows Frenkel disorder, given by the equation Mn2+Mn2+ + Mn" [1] where Mn2+ is a normal lattice ion which moves into an interstitial site (Mni2+) leaving a vacancy (Mn,"). The superscript dashes indicate that there is an excess charge on the anions surrounding the vacancy. Schottky disorder, in which a perfect crystal acquires equal numbers of cation and anion (Ov") vacancies according to leads to equations of exactly the same form in the analysis below. In view of the neutron diffraction results of Roth,2 which shows extensive Frenkel disorder in Fe1-o, Eq. [I] seems more appropriate. The deviations from stoichiometric composition, MnO, result from the addition of oxygen according to In effect Davies and Richardson ignore the variation of Mn3+in constructing equilibrium constants. This approximation might not appear to be very serious if the excess positive charges are really free posi- tive holes (P+) and if the intrinsic conductivity is high as a result of the reaction where e- is a conducting electron. However, the calculations of Buessem and Butler3 indicate that the partition functions of the free charges make a large contribution to the equilibrium constant, and this approximation would not be a desirable one. Furthermore, Morin4 shows that the positive holes are probably localized in MnO and therefore are described better by the mass action principle than by a degenerate gas model. Davies and Richardson's considerations require that activity corrections be applied beyond a 6 value of about 0.005. It is proposed to show here that the dissociation pressure curve can be calculated reasonably well over the whole composition range studied without activity corrections by choosing appropriate equilibrium constants for reactions 1 and 2. As stoichiometric composition The mass action constant, Ks, for reaction [I.] and the equilibrium constant, K,, for reaction [2] are given by where the concentrations which do not change significantly have been absorbed into the constants. The brackets denote concentrations. The deviation from stoichometric composition is given by gardless of [Mni2 ]. Substituting Eq. [7] in [6]

Jan 1, 1961

By R. G. Hudson, H. W. Paxton

The activities of iron and chromium in their binary system have been measured by the Knudsen orifice technique. The system is close to ideal at 1200- 1250 C with slight positive deviations. The effect of this on the liquidus and solidus is discussed and these are calculated for two limiting assumptions. The vapor pressure of pure iron, pure chromium, and one alloy is compared with that obtained by the Langmuir free-evaporation technique. The results agree, indicating a vaporization coefficient close to 1. HERE are no reported data on the thermody-A namic activity of iron and chromium in the solid portions of the phase diagram. It is generally assumed'.' that the liquid solutions of iron and chromium obey Raoult's law. Chen and Chipman" have inferred that the liquid solution obeys Raoult's law. The minimum' in the solidus-liquidus curves for the iron-chromium system shows that there are deviations from ideality in either the solid or liquid or both. Because of the importance of developing reliable techniques for obtaining activity data in alloy systems at high temperatures and of the data so obtained, an investigation of the activity of iron and chromium in the binary at 1200-1250 C was undertaken using the Knudsen cell. Experimental The apparatus and experimental procedures used in this investigation were described in a previous paper." The various materials used for the construction of the Knudsen cell are listed in Table I. The metal used for the determination of the vapor pressure of chromium was 99.9 pet pure, obtained from A. D. Mackay, Inc. The iron was obtained as electrolytic grade from Fisher Scientific Company. The Fe-Cr alloys used are shown in Table 11. The largest error and controlling factor in the accuracy in the experiments reported here are introduced because of temperature uncertainty, which is estimated to be ±3 deg C. This introduces an uncertainty of 26 pet in the measured pressure. Experimental Results and Calculations The vapor pressures of pure iron and pure chromium were calculated from the measurements by the Knudsen equation3 on the assumption that only Cr or Fe atoms were present in the gas phase. For metal vapors, at the high temperatures and low pressures involved in these measurements this should be a reasonable assumption. Whitman" corrections for the cells were made. The data and results of pressure calculations are presented in Table I, and are plotted in Fig. 1. In addition the heat of vaporization at the absolute zero, ?E°°, has been calculated for chromium. These calculations are summarized in Table 111. The data used in these calculations are as follows kelley13 (F° - E°°/T ) Speiser, et al H (F° - E°°/T ) 010 where F" — standard free energy at temperature T It was necessary to calculate the total pressure above the alloys from the total weight loss W of the cell during a run, because of the fact that all attempts to collect and analyze the chromium and iron vapor effusing from the cell in the apparatus available failed. The failure resulted from the re-

Jan 1, 1959

By J. M. Silliman

Several years ago, having occasion to determine the amount of some very slight atmospheric depressions, I devised and had constructed by a skilful tinsmith the manometer shown in the accompanying drawing. It consists of a cylindrical vessel of tin, 6.6 inches in diameter, through the sides of which, at a distance of about six inches from the top, passes an air-tube D, its upper extremity terminating 3.5 inches from the top. Within the vessel is a combined cylipdrical

Jan 1, 1889

By H. R. Deeth, C. J. Koenig

ABOUT ten years ago nepheline syenite was introduced to the ceramic industry and the material has now found application in the various branches of the industry, namely, as a vitrifying agent in white-wares and as a source of alumina and alkalies in glasses, glazes, and porcelain enamels. Spence1 described the early mining and processing developments regarding this material. Since the time of his article many research papers have appeared relating to the usage of nepheline syenite but little has been published relative to subsequent mining and processing methods. A complete indexed review of the literature was issued recently by Ohio State University.2 Since "Lakefield" nepheline syenite from Blue Mountain, Ontario (Fig I); is the only commercial ceramic-grade nepheline syenite now being marketed on this continent, discussion is limited to this operation carried on by American Nepheline Limited, Lakefield, Ont., and its subsidiary in the United States, American Nepheline Corp. at Rochester, N. Y. THE DEPOSIT Blue Mountain is a long ridge rising to an elevation of about 350 ft at its highest point above the surrounding country and extending for about six miles in a northeasterly direction across Methuen township, Peterborough County. The rock of the Blue Mountain central alkaline intrusive is fine grained, light colored and of even granitic texture. The solidified mass consists predominately of the minerals albite, nepheline and microcline in order of abundance. The average mineralogical composition for the deposit obtained by a study of 112 samples taken regularly over the whole formation is as follows:3 [VOLUMETRIC MINERAL PER CENT Albite54 Microcline20 Nepheline22 Muscovite2 Mafics2 (Biotite, hastingsite, magnetite) Total100] The mineral nepheline (or nephelite) belongs to the feldspathoid group but is lower in silica content than are the feldspars. It is a silicate of soda, potash, and alumina (NaKAl2O32SiO2). Its presence with albite (Na2O A1203 6SiO2) and microcline (K20 A12O3 6SiO2) gives the Blue Mountain deposit an alkali content of 15 pct and an alumina content of approximately 24 pct. While this huge alkaline mass is somewhat variable mineralogically, it has a uniform chemical composition. When the iron, mostly in the form of magnetite, is removed during processing, a remarkably uniform finished product is obtained. Geological surveying has disclosed many million tons of nepheline syenite in the Blue Mountain deposit. Diamond drilling in the area where mining is now being carried out indicates a block of about one million tons.

Jan 1, 1947

By Nicholas J. Grant, John Chipman

A full understanding of the behavior of sulphur in the basic open-hearth process has been delayed by lack of dependable data covering a wide range of slag conditions in the absence of other complicating factors that are likely to be present under operating conditions. It has been generally accepted that strongly basic slags will dissolve more sulphur than will those in which the acidic and basic constituents are more nearly balanced. It has been known that both slag and metal would absorb sulphur from the flame when the sulphur content of the fuel is high. It has been generally believed that manganese assists in the transfer of sulphur from metal to slag. Almost nothing has been known of the effect of temperature upon desulphurization;the data have been conflicting and there has been a tendency to depend rather heavily upon analogy with the better understood desulphurizing reaction of the blast furnace. Many attempts have been made to build up from open-hearth data a satisfactory theory of the chemical reactions involved in the process. That this has been an extremely difficult task is owing to two reasons that have been very generally recognized. The first is the parallel and simultaneous reaction of sulphur transfer between flame and bath. The second is the difficulty of distinguishing between the factors that ultimately limit the transfer from metal to slag—i.e., equilibrium conditions—and those whose influence is only upon the rate of transfer. It has been the object of the work here recounted to study the equilibrium distribution of sulphur between slag and metal in the absence of these complicating factors, with the hope that a better understanding of the equilibria involved would lead to a better control of the over-all process in the open-hearth furnace. A review of the literature on steelmaking reveals almost complete agreement with respect to the paramount importance of lime in promoting desulphurization. In expressing this effect in quantitative terms, however, and in formulating the relationships between the concentrations of other slag components and the desulphurizing power, there is little agreement. For instance, the "basicity" or basic strength of the slag is commonly measured by some kind of ratio of concentration of basic to acidic oxides. Of these, the following have been reviewed by Schenck:20 V radio = % CaO/%SiO, Körber and Oelsen's _ % CaO'. 100/% CaO + % SiO2 where (%CaO') = % CaO- 1.18% P2O5. Herty's B ratio = % CaO - (0.93 . % SO2) - (1.18% P205) Among others there appear the following: 20 References are at the end of the paper.

Jan 1, 1947

By Nicholas J. Grant, John Chipman

A full understanding of the behavior of sulphur in the basic open-hearth process has been delayed by lack of dependable data covering a wide range of slag conditions in the absence of other complicating factors that are likely to be present under operating conditions. It has been generally accepted that strongly basic slags will dissolve more sulphur than will those in which the acidic and basic constituents are more nearly balanced. It has been known that both slag and metal would absorb sulphur from the flame when the sulphur content of the fuel is high. It has been generally believed that manganese assists in the transfer of sulphur from metal to slag. Almost nothing has been known of the effect of temperature upon desulphurization;the data have been conflicting and there has been a tendency to depend rather heavily upon analogy with the better understood desulphurizing reaction of the blast furnace. Many attempts have been made to build up from open-hearth data a satisfactory theory of the chemical reactions involved in the process. That this has been an extremely difficult task is owing to two reasons that have been very generally recognized. The first is the parallel and simultaneous reaction of sulphur transfer between flame and bath. The second is the difficulty of distinguishing between the factors that ultimately limit the transfer from metal to slag—i.e., equilibrium conditions—and those whose influence is only upon the rate of transfer. It has been the object of the work here recounted to study the equilibrium distribution of sulphur between slag and metal in the absence of these complicating factors, with the hope that a better understanding of the equilibria involved would lead to a better control of the over-all process in the open-hearth furnace. A review of the literature on steelmaking reveals almost complete agreement with respect to the paramount importance of lime in promoting desulphurization. In expressing this effect in quantitative terms, however, and in formulating the relationships between the concentrations of other slag components and the desulphurizing power, there is little agreement. For instance, the "basicity" or basic strength of the slag is commonly measured by some kind of ratio of concentration of basic to acidic oxides. Of these, the following have been reviewed by Schenck:20 V radio = % CaO/%SiO, Körber and Oelsen's _ % CaO'. 100/% CaO + % SiO2 where (%CaO') = % CaO- 1.18% P2O5. Herty's B ratio = % CaO - (0.93 . % SO2) - (1.18% P205) Among others there appear the following: 20 References are at the end of the paper.

Jan 1, 1947

By S. A. Mersol, C. T. Lynch, F. W. Vahldiek

The room-temperature microhardness of as-grown and annealed MoaC single crystals was measured on the (0001), {2110), and1012) planes using Knoop and Vickevs indenters at loads ranging front 25 to 1000 g. The orientatimz dependence of hardness with respect to crystal axes was also studied. The average random hardness of as-grown crystals was determined to be 1520 kg per sq nm. Annealing to 2000°C decreased. the average hardness by 150 units. An increase in hardness after annealing at 2200 ;C was noted. Optical and electron microscopy revealed slip and twin traces on all planes studied, as produced by mi-cvohavdness indentations. Basal (0001)(2i10) slip was determined to be the primary slip system and was substantiated by electron transmission microscopy. A secondary {1010)(2110) slip was produced by mi-crohardness indentations. The lattev also produced twinning- of the {10i2)[0001] type, as proven by electron diffraction. Electrical resistivity and elastic-modulus anisotropy were found and correlated with hardness anisotropy and Mo2C crystal structure. Elastic-modulus values were obtained by microhard-uess and ullrasonic methods. Bonding mechanism of Mo2C is discussed. ROOM-TEMPERATURE microhardness indentations are useful for studying hardness anisotropy, slip, and twinning in brittle materials. Slip has previously been produced in this manner in and WS~~.~ Recently, the authors5 reported slip of the {10i0)(11~0) type produced by high pressure and microhardness indentations on hexagonal TiBz (c/a = 1.066) single and polycrystals. This slip system was also reported by French and ~homas' and Taka-hashi and ~reise~ for hexagonal WC (c/a = 0.976) crystals. These results suggest that prismatic rather than basal slip is favored in hexagonal nonmetallic materials having a c/a ratio considerably less than the ideal (1.633). Buerger precession and cylindrical X-ray rotation patterns were previously' taken on cleaved sections of the Mo2C single crystals studied in this work. Th~y were found to be hexagonal MoaC with a. = 3.0233A, co = 4.7344A, and C/O = 1.5660. The latter ratio is close to that of the beryllium metal (c/a = 1.57), which slips primarily on the (0001) plane: but also slips on the (1010) planes.'0 ~irconium" (c/u = 1.59) and titanium12 (c/o = 1.59) deform mainly by slip on nonbasal planes which contain a close-packed direction. This is due to the fact that for these two metals the initial resolved shear stress for slip on the (10i0) prismatic planes is lower than that on the (0001) plane. The prominence of basal rather than prismatic slip in metals of high c/o ratios is shown by cadmium (c/a = 1.89), zinc (c/a = 1.86), and magnesium (c/a = 1.62) which deform mainly by basal slip. However, in case of the latter, by stressing magnesium crystals in tension or compression parallel to the basal plane, slip on (10i0) planes can also be produced.13 Several hardness values for polycrystalline Mo2C are reported in the literature: Biickle ' and Samsonov'~ give a value of 1800 and 1479 kg per sq mm, respectively, at a 100-g load; and Kieffer and Benesovsk~'~ report a value of 1950 kg per sq mm at a 50-g load. ~ott'~ reports a Vickers hardness value of 2000 kg per sq mm, with the load unspecified. A Rockwell A hardness value of RA = 88 has also been reported.'' In the present work, for comparison with single-crystal Mo2C hardness values, a Khnloo value of 1600 It 150 kg per sq mm was found on 99.6 pct pure and 99 pct dense hot-pressed Mo2C. This work was undertaken partly to explain the considerable differences in hardness values reported for polycrystalline Mo2C. EXPERIMENTAL The Mo2C single crystals investigated were prepared by a Verneuil-type process using an electric arc by the Linde Co. of the Union Carbide Corp.lg The largest specimens grown were boules 7 mm in diam by 40 mm in length. The crystals had an average density of 9.04 g per cu cm, with a Mo + C content of 99.8 wt pct. The major impurities were: 100 ppm each Na, Zr, and Ca; 85 ppm 0, 55 ppm Fe, and 10 ppm each Cr and Ta. The crystals were found to be carbon-poor, the average carbon content being 5.73 wt pct (stoichiometric value is 5.89 pct). The molybdenum content was found to be 94.08 pct, which is nearly stoichiometric. Electron-microprobe traverses of selected specimens were done with a Phillips-AMR microanalyzer. Thin-film and carbon replicas were used to prepare electron micrographs. This work was done with a JEM-6A electron microscope. Prior to optical and electron-optical studies, specimens were mounted in Lucite and polished on a vibratory polisher using diamond-paste grades ranging from 9-1 p and Linde A powder for up to 48 hr. Dilute nitric acid was used for thin-section polishing and chemical etching for 1-15 min. Electrical-resistivity measurements at room temperature were taken with a Rubicon bridge, using gold contacts. For hardness measurements, a Tukon Microhardness Tester Type FB with Knoop and Vickers indenters was used. Measurements were taken at loads ranging from 25 to 1000 g; however, the 100-g load was chosen as the standard load. All measurements were taken at room temperature. Indentations of cracking classes 1 and 2 only were considered for hardness determinations.20)21 (There are six cracking classes, ranging from "class 1" for a perfect inden-

Jan 1, 1967

By K. E. Gray, M. S. Seth

In Part I of this work,1 equations of elasticity were formulated for transversely isotropic, axisymmetric, homogeneous, porous media exhibiting pore fluid pressure. Equations of elasticity and the thermal analogy method were used to determine transient horizontal, tangential, and vertical stresses and radial displacement in a semi-infinite cylindrical region when either a constant rate of pressure or a constant rate of flow is maintained at the wellbore. In this paper, the approach presented earlier is extended to finite reservoirs for the cases of (1) steady-state flow, (2) constant pressures at the well bore and outer boundary and (3) constant pressure at the wellbore and no flow at the outer boundary. Results of this work show that radial and tangential stress gradients are high near the wellbore but diminish rapidt; away from the well; the vertical stress gradient behaves in the same way but is less severe. Radial stresses are compressive or neutral, whereas tangential and vertical stresses may be tensile, neutral or com-pressive, depending upon the boundary conditions, the physical Properties of the system and the radial distance involved (vertical stresses are always compressive in an unbounded system1. For constant boundary pressures, both radial and tangential stresses increase with time whereas they both decrease for a closed outer boundary and constant pressure at the wellbore. The vertical stress decreases with time for both systems. For steady-state systems, radial displacement may be positive or negative, depending upon the dimensions of the system, the pressure differential and the porosity. Radial displacement may be positive or negative for a closed outer boundary but is positive for constant pressures at both boundaries. INTRODUCTION The importance, utility and complexity of a realistic appraisal of the stress state at and local to a wellbore were indicated in Part I. In this paper the analytical approach presented earlier is extended to finite, cylindrical reservoir geometry for the cases of (1) steady-state flow, (2) constant pressures at wellbore and outer boundary and (3) constant pressure at the wellbore and no flow at the outer boundary. Other than the outer boundary of the reservoir being finite, the physical system and assumptions pertinent thereto are the same as before. The reader may wish to review the mathematical development through Eq. 49 of Part I before proceeding here. SOLUTION FOR GENERAL BOUNDARY CONDITIONS THERMAL STRESSES AND DISPLACEMENTS Since normal stresses do not exist at the free boundaries of a system subjected to thermal stresses, the boundary conditions are

Jan 1, 1969

By Maurice L. Pinel, Thomas Wright, Thomas T. Read

The erroneous, but until recently widely prevalent, belief that iron castings were first made in Europe in the fourteenth century has been adequately refuted in a number of earlier papers;1, ll, l² but except for an unpublished metallographic study by the late William Campbell, and a phosphorus analysis by T. A. Wright of the metal of the Han dynasty cast-iron stove described by Laufer,2 nothing was known of either the chemical composition or the metallographic structure of the metal of any ancient iron casting, although data on castings of recent manufacture are abundantly available. The metal of the stove, which had been buried in a grave for at least 15 centuries, was so corroded that it seemed inadvisable to publish the results on it until morc and better specimens from other ancient castings could be obtained. During a journey through China, extending from Peking in the north to Canton in the south and Cheng-tu, Ssu-chuan, in the west, in the late summer and autumn of 1936, one of us (T. T. Read) was fortunate enough to obtain, and bring home for study, nine castings, all more than 1000 years old, of which the date of manufacture is precisely known because of inscriptions cast in them, another that lacks a dated inscription, but of which the date, as explained below, can be otherwise established, and a specimen from the largest iron casting ever made," which is also dated. The period covered by this suite of specimens extends from 502 to 1093 A.D. In addition, a sample was obtained from a casting of date unknown, but which was certainly cast before 1000 A.D. and possibly as early as the third century A.D. A detailed description of these specimens is as follows: I. 502 A.D. A pair [right and left) of recumbent lions (Fig. 1). Base 7¼ by 113/8 in., height over-all 5½in, weight 26 Ib. each. The tranblation of the inscription on the base is: "Made on the twenty-fourth day of the seventh month of the third pear of Ching Ming of Great Wei" (Sept. 11, 502 A.D.).

Jan 1, 1938

By Maurice L. Pinel, Thomas Wright, Thomas T. Read

The erroneous, but until recently widely prevalent, belief that iron castings were first made in Europe in the fourteenth century has been adequately refuted in a number of earlier papers;1, ll, l² but except for an unpublished metallographic study by the late William Campbell, and a phosphorus analysis by T. A. Wright of the metal of the Han dynasty cast-iron stove described by Laufer,2 nothing was known of either the chemical composition or the metallographic structure of the metal of any ancient iron casting, although data on castings of recent manufacture are abundantly available. The metal of the stove, which had been buried in a grave for at least 15 centuries, was so corroded that it seemed inadvisable to publish the results on it until morc and better specimens from other ancient castings could be obtained. During a journey through China, extending from Peking in the north to Canton in the south and Cheng-tu, Ssu-chuan, in the west, in the late summer and autumn of 1936, one of us (T. T. Read) was fortunate enough to obtain, and bring home for study, nine castings, all more than 1000 years old, of which the date of manufacture is precisely known because of inscriptions cast in them, another that lacks a dated inscription, but of which the date, as explained below, can be otherwise established, and a specimen from the largest iron casting ever made," which is also dated. The period covered by this suite of specimens extends from 502 to 1093 A.D. In addition, a sample was obtained from a casting of date unknown, but which was certainly cast before 1000 A.D. and possibly as early as the third century A.D. A detailed description of these specimens is as follows: I. 502 A.D. A pair [right and left) of recumbent lions (Fig. 1). Base 7¼ by 113/8 in., height over-all 5½in, weight 26 Ib. each. The tranblation of the inscription on the base is: "Made on the twenty-fourth day of the seventh month of the third pear of Ching Ming of Great Wei" (Sept. 11, 502 A.D.).

Jan 1, 1938

By Walter S. Weeks

THE process of cooling air by allowing it to expand and do work in an engine is well known, but the theory of obtaining cold air by free expansion without the aid of an engine operating with cutoff has been little discussed. Case 1 If compressed air is allowed to expand to the atmospheric pressure through an orifice in an sir line that is supplied by a running compressor, no cooling results except the negligible one due to the Joule-Thompson effect, which will not be considered in this discussion. The effect of moisture will also be neglected. Symbols P1 Absolute pressure in pipe line, lb. per sq. ft. P3 Atmospheric pressure, lb. per sq. ft. V1 Volume occupied by 1 lb. air in pipe. V3 Volume occupied by 1 lb. air at P3. Cv Specific heat of air at constant volume, B.t.u. per lb. per deg. F. Cp Specific heat at constant pressure, B.t.u. per lb. per deg. F. T1 Absolute temperature of air in pipe, deg. F. T3 Absolute temperature of air outside after expansion and mixing and coming to rest. J Mechanical equivalent of heat (778 ft. lb.) R Gas constant for air, ft-lb. per deg. F. = 53.3. Formulas Consider 1 lb. of air moving out of a pipe. The air behind it does work on it to the extent of P1V1 ft-lb. This work is really done by the thrust of the piston of the compressor and does not come from any energy that exists in the air in the pipe. The pound of air contains

Jan 1, 1937

By A. A. Smith, J. S. Smart

THe elements silver, cadmium, tin, antimony and tellurium either are found as impurities in commercial coppers or are intentionally added to produce coppers for special uses. When present in small quantities, each of these metals exerts a significant effect on the properties of copper, but despite considerable experience and investigation much remains to be learned about their specific behavior. The presence of extraneous elements, particularly oxygen, complicates the study of individual additions of impurities in the range o to 0.05 per cent, and the use of high-purity copper has been found extremely helpful in circumventing these difficulties. It has also been found that individual behavior can be satisfactorilv determined from measurements of two commercially important properties, conductivity and softening temperature, according to methods that have been described previously.1,2 Data of this type, obtained in the course of a general investigation of the effects of impurities on the properties of both oxygen-bearing 2nd oxygen-free copper are presented herewith. All samples except those containing cadmium were prepared as 3/8-in. continuously cast oxygen-free rod from pure copper and weighed amounts of a similarly synthesized master alloy. Composition was checked by chemical methods, which in every case verified the calculated analyses. Test wires were cold-drawn one B, and S. number per pass on a string-up machine using four 30-min. 600°C. anneals at 0.3125, 0.257, 0.204 and 0.162 in. and a final cold reduction of 75 per cent to 0.081 in. Some oxygen-bearing samples were produced directly from the 3/8-in. oxygen-free rods by diffusion from a surface scale at 850°C. and drawn to size as listed above, but most of the alloys were remelted in pure graphite and cast through air into a I-in. round mold. The slugs were then hot-rolled to 5/16 in., drawn to 0.162 in., annealed as listed above, and drawn to 0.081 inch. Conductivity wires were annealed for I hr. at 100°C. intervals between 300' and 800°F., using a hydrogen atmosphere for oxygen-free samples and purified nitrogen for the oxygen-bearing type. All samples were quenched from the annealing temperature in 10 per cent H²S04 and conductivity was measured on a Hoopes bridge that had been checked against a Kelvin bridge. The softening temperatures represent the half-hard stage of the annealing curve as determined by a standard procedure.' Effect of Silver Numerous excellent studies of the argentiferous coppers are found in the literature, and these have been adequately

Jan 1, 1943

By A. A. Smith, J. S. Smart

THe elements silver, cadmium, tin, antimony and tellurium either are found as impurities in commercial coppers or are intentionally added to produce coppers for special uses. When present in small quantities, each of these metals exerts a significant effect on the properties of copper, but despite considerable experience and investigation much remains to be learned about their specific behavior. The presence of extraneous elements, particularly oxygen, complicates the study of individual additions of impurities in the range o to 0.05 per cent, and the use of high-purity copper has been found extremely helpful in circumventing these difficulties. It has also been found that individual behavior can be satisfactorilv determined from measurements of two commercially important properties, conductivity and softening temperature, according to methods that have been described previously.1,2 Data of this type, obtained in the course of a general investigation of the effects of impurities on the properties of both oxygen-bearing 2nd oxygen-free copper are presented herewith. All samples except those containing cadmium were prepared as 3/8-in. continuously cast oxygen-free rod from pure copper and weighed amounts of a similarly synthesized master alloy. Composition was checked by chemical methods, which in every case verified the calculated analyses. Test wires were cold-drawn one B, and S. number per pass on a string-up machine using four 30-min. 600°C. anneals at 0.3125, 0.257, 0.204 and 0.162 in. and a final cold reduction of 75 per cent to 0.081 in. Some oxygen-bearing samples were produced directly from the 3/8-in. oxygen-free rods by diffusion from a surface scale at 850°C. and drawn to size as listed above, but most of the alloys were remelted in pure graphite and cast through air into a I-in. round mold. The slugs were then hot-rolled to 5/16 in., drawn to 0.162 in., annealed as listed above, and drawn to 0.081 inch. Conductivity wires were annealed for I hr. at 100°C. intervals between 300' and 800°F., using a hydrogen atmosphere for oxygen-free samples and purified nitrogen for the oxygen-bearing type. All samples were quenched from the annealing temperature in 10 per cent H²S04 and conductivity was measured on a Hoopes bridge that had been checked against a Kelvin bridge. The softening temperatures represent the half-hard stage of the annealing curve as determined by a standard procedure.' Effect of Silver Numerous excellent studies of the argentiferous coppers are found in the literature, and these have been adequately

Jan 1, 1943

By J. E. Andersen, J. Halpern, C. S. Samis

In the presence of oxygen, galena is oxidized in an aqueous medium containing sodium hydroxide, in accordance with the following reaction: PbS + 2O2 + 3OH ? HPbO2 + SO4 = + H2O A novel method was devised for following this reaction which takes place in an autoclave under oxygen pressure, by measuring the concentration of HPbO2- in the solution with a cathode ray polarograph employing stationary platinum electrodes. Using this procedure a study was made of the kinetics of the reaction, in which the effect of a number of variables, including temperature, oxygen pressure, NaOH concentration, and agitation, on the rate, were determined. The results of this study are described and discussed, in terms of possible mechanisms for the reaction. IT is known that sulphide minerals can be oxidized in aqueous media in the presence of oxygen under pressure. Reactions of this type can be classified in two categories depending upon whether the products of oxidation are insoluble or soluble in the aqueous medium. An example of the first type of reaction is the oxidation of pyrite in alkaline solutions, giving rise to an insoluble iron oxide. The kinetics of this reaction have been investigated and its mechanism established.' It has been applied in the oxidation of iron sulphides present in refractory gold ores to improve the subsequent recovery of gold by cyanida-tion.' The second type of reaction finds application where it is desired to recover the metal by leaching during oxidation. In this case a medium is selected in which the oxidized mineral is soluble. For example, nickel, copper, and cobalt sulphide ores can be treated by oxidation in the presence of a solution of ammonia dissolving the nickel, copper, and cobalt as the metal ammine sulphate salts.:' A similar treatment has also been reported to apply to zinc sulphide ores.' It is known that lead sulphate is soluble in solutions of sodium hydroxide or ammonium acetate. It should therefore be possible to dissolve the lead from galena ores by aqueous oxidation with oxygen under pressure using either of these solutions. The applicability of the ammonium acetate treatment has been investigated and confirmed in a recent study.' The present paper describes the results of a similar study in which the kinetics of the oxidation of galena in sodium hydroxide solutions have been investigated. The object of this investigation was pri- marily to obtain information of a fundamental nature relating to the kinetics and mechanism of this reaction. The use of pulps of comminuted ore in a study of this type is undesirable because of the difficulty in controlling and measuring the surface area. The measurements were therefore made using galena crystals of measured surface area and in this way absolute reaction rates were obtained. The reaction was found to proceed as follows: PbS + 2O2 + 3OH ? HPbO2 + SO4 + H2O [I] The products are sodium plumbite and sodium sulphate salts which dissolve in the aqueous solution. The reaction studies were carried out in an autoclave in which a desired pressure of oxygen was maintained. The reaction was followed by measuring the concentration of lead in the solution with a cathode ray polarograph. The results obtained in this investigation are presented and discussed in the present paper. Experimental Materials: The crystals of galena were obtained from Violamac Mines, Sandon, B. C. The following impurities were indicated by spectrographic analysis carried out by the Provincial Assay Office, Victoria, B. C.: Sn, 0.02 to 0.2 pct; Sb, 0.07 to 0.7; Zn, 0.01 to 0.1; Si, Fe, Mg, As, less than 0.05 pct; Ag, 126 oz per ton, the latter determined by standard fire-assaying procedure. After cutting a specimen measuring approximately 5x7x12 mm, the crystal surfaces were ground parallel to the 100 axes using No. 2 emery, washed with water, and the dimensions were measured with a micrometer. Oxygen gas was of commercial grade and supplied in cylinders by Canadian Liquid Air. Solutions were made up with chemicals of chemically pure grade and distilled water. Equipment: Autoclave: The autoclave used in these studies was constructed of stainless steel and

Jan 1, 1954

By G. R. Speich, R. A. Oriani

The rate of coarsening (Ostwald ripening) of copper precipitate Particles in an a-iron matrix has been studied in three Fe-Cu alloys containing 2.3, 4.0, and 5.4 wt pct Cu in the temperature range 730° to 830°C. The kinetics of Ostwald ripening ohevs the law as deduced by Lifshitz and Wagner so that matrix diffusion controls the process. The equilibrium shape of the particles is that of rods with hemispherical caps, and is due to the BOTH Lifshitz and slyozov and wagner3 have recently solved the complex problem of determining the rate of coarsening of an assembly of precipitate particles. Both workers, although using somewhat different mathematical approaches, arrive at the same equation for the relation between the mean radius 7 of a distribution of spherical particles and the bulk solubility co, diffusivity D in the matrix, large anisotropy in the copper/a iron interfacial free energy (y,/y, = 4). A Lifshitz -Wagner equation, modified for both the nonsphericity of the diffusion field and the reformulation of the Gibbs-Thomson equation for the case of nonsphencal particles, is derived. Using the modified equation, the calculated values of the mean radius are smaller than those observed by a factor of about one -half. absolute temperature T, gas constant R, surface free energy y of the precipitate/matrix interface, and the time t of aging: where V, is the molar volume of the precipitate and v is a stoichiometric factor. Wagner's treatment is restricted to the case of spherical particles growing in a liquid matrix where strain energy is not a factor. Lifshitz and slyozov have discussed the effect of anisotropic particle shapes and internal strains on the rate of coarsening. Oriani4

Jan 1, 1965

By B. M. Loring, I. R. Kramer, F. M. Walters

No observations on the mechanical properties of iron-manganese alloys have been published since pure manganese became readily available, either distilled manganese or electrolytic manganese. The purpose of this investigation was to determine whether iron-manganese alloys are essentially brittle, as reported by Hadfield,1 or whether the lack of ductility observed was due to impurities in the alloys that he studied. EXPERIMENTAL METHODS The alloys were melted in a 100-lb. high-frequency induction furnace with a magnesia crucible, using ingot iron and electrolytic manganese. The heats were split, the first giving six 18-lb. ingots and the second three. Each successive heat was higher in manganese and about 0.3 per cent silicon [ ] was added as ferrosilicon. The alloys forged without difficulty-and were given a 90 percent reduction to 3/4-in. rounds. The chemical analysis of the alloys is given in Table I. The forged rods were homogenized for [ ] 24 hr. at 1150°C. (2100°F.) and air-cooled. The alloys were tested in two conditions of heat-treatment: (I) as normalized from 870°C. (1600°F.) and (2) normalized from 870°C. and tempered for z hr. at 540°C. (1000°F.) and air-cooled. The tensile specimens were 0.505 in. in diameter with a 2-in. gauge length, and the impact specimens were V-notch Charpy. The proportional limit and yield strength were taken from the stress-strain diagrams drawn by a Peters recorder. The data obtained from the tensile test and the hardness measurements are given in Table 2. Specimens that were normalized are designated by N; those that were normalized and tempered by T. Alloy G was tempered after forging (D).

Jan 1, 1941

By H. I. Aaronson, S. M. Kaufman, G. M. Pound

WHETHER the carbides associated with bainite precipitate from the ferritic component of bainite, as originally suggested by Davenport and Bain, or from the adjacent austenite at aus-tenite:ferrite boundaries, as proposed by Hultgren, 2 remains in dispute.3-5 A recent metallographic investigation in this laboratory has shown that the bainitic carbides formed in a plain carbon hypo-eutectoid steel nucleate entirely within ferrite at higher reaction temperatures3 and in both ferrite and austenite at lower temperatures." The present investigation was undertaken to ascertain the effects of alloying elements on the nucleation sites selected by bainitic carbides. The compositions of the steels studied were as shown in Table I. Specimens of these steels were

Jan 1, 1958

By V. B. Kurfman

The temperature, strain rate, and orientation deDendence of defbrnzation of single-crystal MgAg has been examined. The crystals exhibit a tendency to single glide and little or no hardening at 25°C for many orientations. A much higher hardening rate is observed when multiple glide occurs, such as can be initiated by surface defects. The tendency for easy glide becomes less dependent on surface preparation and orientation as T — 100°C and bars so tested often fail after one-dimensional necking-. At T > 200°C (transition temperature for single-crystal notch sensitivity and poly crystalline ductility) single glide diminishes and two-dirnensionul necking begins. The crystals do not strictly obey a critical resolved shear stress law, but show the influence of {loo) cracks in determining the slip mode. The results are correlated with the difficulty of sciperdzslocation intersection and semibrittle behavior of this compound in single-crystal and poly crystalline form. Comparisons are made with the slip selection mode observed in tungsten, with the reported observations of easy glide in bee metals. and with the mechanical behavior of poly crystalline MgAg. PREVIOUS work on tensile deformation of polycrys-talline MgAgl and bending deformation of single-crystal MgAg2 has shown that the compound is semi-brittle (i.e., notch and grain boundary brittle). If this semibrittleness is supposed to result from the difficulty of multiple glide (associated with the problems of superdislocation intersection) one might expect single crystals deformed in tension to show pronounced single glide and strong orientation dependence of hardening rate. These experiments were done to examine this supposition and to study the tensile deformation of a highly ordered system which may be considered bcc if the difference between the two kinds of atoms is ignored (actual structure: CsC1). EXPERIMENTAL Single-crystal ingots were grown by directional freezing as previously described.' These ingots were sliced into a by a by 2 in, rectangular bars by electric discharge machining, then round tensile bars were conventionally machined to 1/8-in.-diam by 1-in.-long reduced section. The bars were typically tested without an anneal because of the problem of magnesium vapor loss and they were typically tested as mechanically polished. The analyses are within the same limits as those reported earlier; i.e., the average composition for each specimen is within 0.5 at. pct of stoichiometry, while the total range from end to end in a given specimen varies from 0.7 to 1.4 at, pct. There has been no indication in the results of any variation in slip or fracture mode attributable to the composition fluctuations. The slip systems were determined by two-surface analysis of the bars after testing to failure at room temperature. Single glide was so dominant that there was little difficulty in identification of the dominant slip system even though the tensile elongation to failure often approached 7 to 8 pct in room-tempera- ture tests. Elevated-temperature testing was done in a silicone oil bath and low-temperature testing was done in liquid Np or a dry-ice bath. All stress measurements are reported as engineering stress unless otherwise specified, and crosshead travel is used as the strain measurement. RESULTS The tendency toward single glide is best seen in the pictures, Figs. 1, 2, and 3, which depict deformation at fracture as a function of test temperature. While it is possible to find regions of secondary slip by careful microscopy, such regions are very small. The development of a ribbon-shaped configuration from an initially round section bar pulled at 100°C is typical, occurred by single glide, and illustrates the degree to which such glide continues. At temperatures =100°C the bars typically show elongation of 20 to 50 pct by predominently single glide. Despite the large elongation, fracture even at 150°C occurs in a brittle mode, Fig. 2, in the sense that it is an abrupt failure which shows no discernible necking in the second dimension of the bar's cross section (i.e., there is no appreciable action of any slip modes which would decrease the broad dimension of the cross section). Near 200°C the fracture mode changes slightly. Although most of the sample extension is by single glide, after the bar develops the characteristic ribbon shape it begins to neck in the second (i.e., broad) cross-sectional dimension. The bar becomes very thin in the "necked down" region, Fig. 3, and the reduction in area approaches 100 pct. Often there oc-

Jan 1, 1967