Search Documents

By H. I. Aaronson, C. Laird

Although the past results of X-ray experiments indicate that the broad faces of 0' plates are coherent with their matrix, dislocations lying in arrays have frequently been observed at these boundaries by transmission electron microscopy. Critical experiments employing the latter technique have been carried out in order to determine the origin of these dislocations. It is concluded that 8' plates are essentially coherent with the matrix at their broad faces throughout the aging temperature/time envelope studied. Virtually all of the dislocation arrays observed are deduced to have been formed by plastic deformation accompanying transformation. The proportion of dislocations arising from convexity of the plates is shown to be negligible by comparison with that from plastic deformation. At the higher aging temperatures, a[001] dislocations appeared in moderate numbers. These dislocations were traced directly, however, to the ledgewise dissolution of 0' occasioned by the formation nearby of 0 crystals. On the other hand, since there is a parametric difference normal to the broad faces of the ?' plates, mismatch dislocations do form at their edges. A previous conclusion that these dislocations have Burgers vectors of type a[001] was confirmed directly. The edges of 0' plates were observed to develop octagonal shapes when growing, but circular shapes during dissolution. 1 HIS paper presents the results of an investigation of the interfacial structures of plates of the transitional phase, 8', formed in an A1-4 pct Cu alloy. In a companion paper, Part 11, the effects of these structures upon the migration kinetics of a:?f boundaries are reported. This work is pa.rt of a general program designed to establish the basis of precipitate morphology. The present authors in Al-Ag,1 and whitton2 previously in U-C alloys, have used transmission electron microscopy to examine directly the vander Merwe3-6 networks of dislocations anticipated7 to compensate the small amount of lattice misfit normally founda at the broad faces of Widmanstatten plates. Since the broad faces of 0' plates are considered to be perfectly coherent with the corresponding habit planes in the a matrix,' no dislocations should be present at these faces. Many reports have been published, however, giving evidence to the contrary.10-18 The primary objective of this investigation was therefore to ascertain the nature of these dislocation structures. An attempt to do this is described in the first three sections of this paper. Inspection of the matching of the a and 8 ' lattices at the orientations of the 0:0' boundary corresponding to the edges of 0' plates raises the possibility that these edges may be made up of rather closely spaced edge- type misfit dislocations oriented so as to be sessile with respect to the lengthening or shortening of the plates. Since this structure should severely inhibit migration of the plate edges (Ref. 7, Part II), a situation not originally anticipated,' an experimental determination of the interfacial structure of the edges of 8' plates was clearly in order, and is reported in Section III. Those aspects of the experimental procedure applicable to both Parts I and I1 are presented in the next section. Specific procedures applicable to individual aspects of each investigation, and also the relevant surveys of the literature, are then individually reported in the appropriate sections. I) GENERAL EXPERIMENTAL PROCEDURE The material used in both parts of these studies was the same as that of a previous investigation:" strips of A1-3.93 pct Cu, 0.009 in. thick, prepared as before, solution-annealed at 548°C for 6 hr, and quenched. Details of subsequent aging, and in some cases deformation treatments, are given in the Experimental Procedure sections of the individual parts of both papers. Specimens of the heat-treated strips were electro-thinned as beforeLg and examined in a Philips EM 200 microscope equipped with a goniometer stage. A commercial hot stage, of the grid-heater type and capable of * 30-deg tilt about one axis in the plane of the specimen, was also used for kinetic studies. The usual precaution of calibrating for the additional heat supplied by the electron beam was taken.19 A 16-mm cine cam-I era mounted outside the viewing window was frequently used to record the transformations. Conventional selected-area diffraction and dark-field viewing techniques were used to identify the precipitates in the foils. Normal bright-field images corresponding to two-beam diffracting conditions or dark-field images were employed to characterize the dislocations observed at the interfaces of the precipitates. The application of these techniques to the study of an interphase boundary, and the interpretation of the images,20'21 has been fully described in a previous paper.'

Jan 1, 1969

By Frank T. Sisco

DURING the past year the iron and steel industry of the world as a whole operated on a satisfactory basis. No discoveries nor new processes of outstanding importance were announced either here or abroad, although rumors of such were persistent. In general, the efforts of operating and research men in the iron and steel industry seemed to be confined to improving quality and reducing cost, not a new aim, for the same efforts have been noted in nearly all of the annual reviews of the industry for at least fifty years. Most published papers on blast furnace or steel-making practice were of minor importance, at least when judged by the criterion of radical change in process or product. If the number of reports in the technical journals, or articles in the trade press, is used as a basis for comparison, continuous-mill sheet is still the most important single product of the industry; papers were published detailing the results of the investigation of various phases of the

Jan 1, 1938

By Charles T. Holland

In connection with the problems of bumps in coal mines, much has been written concerning the manner in which roof action and methods of mining enter into the pressure effects observed but little has been written upon the relationship between the physical properties of coal and associated rocks and the cause of bumps. The question of the strength of bituminous coal, of its elasticity, and the energy it can absorb and store has been neglected to a considerable extent in connection with this problem. It is generally agreed that one of the necessary conditions for bumps is a strong coal.' Therefore, since the strength of coal is an important factor, a study of this property should be made. several investigators have conducted laboratory experiments upon the strength of coal, using small cubical or rectangular test pieces. An attempt will not be made to go into all of these investigations, but they are briefly summarized in Table I. Experiments of the U. S. Bureau of Mines and of Lawall and Holland (pp. 27 and 34 of ref. 9), however, indicate that regardless of whether the coal is a soft and weak semihituminous (Pocahontas No 3) or a hard and strong bituminous (Stockton) coal, the maximum ultimate unit stress developed in laboratory tests is higher in small cubes than in larger cubes of the Same coal.

Jan 1, 1942

By Charles T. Holland

In connection with the problems of bumps in coal mines, much has been written concerning the manner in which roof action and methods of mining enter into the pressure effects observed but little has been written upon the relationship between the physical properties of coal and associated rocks and the cause of bumps. The question of the strength of bituminous coal, of its elasticity, and the energy it can absorb and store has been neglected to a considerable extent in connection with this problem. It is generally agreed that one of the necessary conditions for bumps is a strong coal.' Therefore, since the strength of coal is an important factor, a study of this property should be made. several investigators have conducted laboratory experiments upon the strength of coal, using small cubical or rectangular test pieces. An attempt will not be made to go into all of these investigations, but they are briefly summarized in Table I. Experiments of the U. S. Bureau of Mines and of Lawall and Holland (pp. 27 and 34 of ref. 9), however, indicate that regardless of whether the coal is a soft and weak semihituminous (Pocahontas No 3) or a hard and strong bituminous (Stockton) coal, the maximum ultimate unit stress developed in laboratory tests is higher in small cubes than in larger cubes of the Same coal.

Jan 1, 1942

Adolph Thies, Haile Gold-Mine, S. C.: I have little to add to what my friends Messrs. Nitze and Wilkens have said on this subject. What I have done in mining during the last fortyone years, partly in this country and partly in South Africa (though not, as to Africa, in the diamond- or the gold-fields,

Jan 1, 1896

By A. Bar-Or, G. Hirsch, P. Wynblatt

The rates of nucleation and growth at various temperatures in U-0.45, 1.1, and 1.35 at. pet Cr alloys were determined, providing a qliantitative proo-f that the transformation under study is controlled by nucleation and diffusional growth. The linear relationship, log NV US 1/T2, suggests that the nucleation of the alloys is probably controlled by the classical law of nuclei formation. The established linear relationship, log G us 1/T, in the three alloys suggests that growth in each alloy is controlled by bidimensional nuclei formation. It was established that both the rate of growth and the rale of nucleation are retarded by chromium. Using classical formulae, sApN and sApG (the ß/a apparent interphase energies) were calculated from the nucleation and growth results, respectively. It was found that their values are of the same order of magnitude. However, why sApN is unaffected by the addition of chromium, while sApG increases with it, needs further clarification. In the last few years, an appreciable amount of research has been conducted into the mechanism of the ß —a transformation in uranium and uranium alloys. Recently, Burke1 reviewed the published work concerned with this transformation. He concluded that in dilute uranium alloys it is firmly established that, at high temperatures in the a range, ß decomposes by a diffusional process. However, he based his conclusion either on experiments dealing with the over-all transformation,2 or on qualitative results.3 It was felt that a better insight into the mechanism of transformation could be gained from the quantitative measurements of growth and nucleation rates, which we aimed to carry out in the present work. In a previous work,4 we have shown that a rim of a grains is formed during isothermal transformation at the high a range. The thickening of this rim with time was used for measuring the rate of growth in U-0.45 at. pet Cr.5 This method was also used in the present work for two more alloys. For meas- uring the rate of nucleation, a method of quenching and macroetching was developed to allow differentiation between a grains formed during isothermal transformation at the upper a range and those formed after quenching, Fig. 1. EXPERIMENTAL WORK AND RESULTS The Rate of Nucleation. The measurement of the rate of nucleation was effected by counting the num-

Jan 1, 1965

By P. R. Sperry, M. L. Holzworth, P. A. Beck

The basic work of Z. Jeffries 1,2,3 has long ago established the main features of grain growth in the presence of a dispersed second phase. Working with sintered specimens of initially fine grained tungsten, to which various amounts of thoria had been added, Jeffries found that grain growth was inhibited, that is, practically prevented, up to a certain annealing temperature. When this temperature, which increased with the thoria content, was exceeded, extremely large grains developed abruptly from the fine grained matrix. such a a temperature" was later found by Gross-mann~3,14 in certain types of steel, while ~~i~ showed15 that in other steels grain growth was gradual, without inhibition and abrupt coarsening. Later work demonstrated the connection between coarsening and aluminum additions. The viewpoint that the direct cause of inhibition and coarsening in aluminum killed steels is a fine dispersion of aluminum oxide21 appears to be held quite generally, although some doubts have been voiced even very recently, 20 In steels with Ti additions the titanium carbide phase is considered responsible for the inhibition effects found.20 The coarsening temperature increases with the titanium content. That coarsening may occur, as a result of, certain heat treatments, even in low carboil rimmed sheet steel, was shown by Samuels. He also detected coarsening in a steel ingot of similar composition, after giving it the same heat treatment as that used by him for the sheet material. In this instance the identity of the inhibiting phase has not been determined. However, Tangerding found23 that even the few small carbide particles, which occur in carbonyl iron, have a very marked inhibitive effect. As a result of an oxidizing anneal at 850°C, large grains begin to form at the surface of the carbonyl iron specimen, and gradually grow inward as the subsurface oxidation of the carbide particles progresses and eliminates the obstruction. The carbide particles Can be removed also by annealing in hydrogen, with similar results. But if the oxidizing anneal of 10 hr at 850°C in air is followed between in hydrogen for 75 hr at the Same temperature, grain growth is more hibition usually the whole specimen is transformed into a single crystal (approx. 1.5 cm X 5 cm). The reversed procedure, that is, hydrogen annealing followed by oxidizing annealing leads to smaller grains. Tangerding arrived at the logical conclusion that there must have been in his specimens some additional obstruction to grain growth hibition from a minor impurity, other than carbon. It seems likely that this obstruction was caused by an oxide formed during the oxidizing anneal and gradually reduced rtain the subsequent hydrogen-annealing. Other instances of grain growth inhibition of varying severity, associated with a dispersed second phase, have been described by several authors. Notable examples: lead with 0.06 pet Cu or Ni,24 cartridge brass with 0.12 pct Cr,25 or with up to 0.15

Jan 1, 1949

By Roy T. Southin*

The finding of Wojciechowski and Chalmers that columnar crystal growth can be stopped by an already existing network of equiaxed crystals is supported by other recently reported work on both static and mechanically stirred ingots.27,28 Support for the dendrite arm fragmentation mechanism, too, can be obtained by comparing the effects of spinning on ingots cast in a material which has a dendritic solid/liquid interface during solidification with one which has a cellular interface, since no evidence of change at the interface, other than a change in growth direction, can be found in the cellular material. But there are alternative ways of producing an equiaxed zone in materials which characteristically solidify with a cellular interface. For example, interrupted spinning of a normally cast ingot can do it. In this case it has been shown28 that the equiaxed zone is caused by the showing of dendritic particles from the top surface of the ingot to nucleate a network of equiaxed crystals which forms a barrier to the continued growth of the columnar grains. This showering has also been shown to be a major effect in materials solidifying with dendritic interfaces. It has also been proposed28,29 that, for an equiaxed zone to form, not only must nuclei be present ahead of the interface but also these nuclei must grow to sufficient size to be a barrier to the further growth of the columnar grains. A further requirement is that a network of these grains continually be produced ahead of the growing interface; otherwise there is a reversion to columnar growth as found by Wojciechowski and Chalmers. In normally cast ingots it is possible for nuclei ahead of the interface to grow since during a large portion of the solidification time there is no temperature gradient ahead of the interface,30 and, in fact, the liquid may be slightly undercooled.29 This then is the major difference between the work under discussion and normal industrial solidification conditions, since Wojciechowski and Chalmers imposed a considerable temperature gradient in the liquid. The statement in the Wojcieshowski and Chalmers paper that one needs to lower the temperature gradient

Jan 1, 1969

By H. B. Aaron

DIFFRACTION micrography provides a useful tool for studying complex strain fields. Newkirk1 observed an X-ray diffraction effect due to strain interactions at the intersection of slip lines in LiF and attributed this effect to a reduction of primary extinction due to strain-gradients. It is the object of the current study to correlate the observations and rationalize the appearance of the strain fields around intersecting slip planes in x-ray extinction contrast studies by a semiempirical method. Fig. 1 shows typical strain patterns observed by Newkirk on (200) reflections in LiF using chromium radiation and a variation of the Berg-Barrett (B-B) technique. In other reflections (e.g., {220}) the strain patterns are absent. If one defines a vector 3 along the striations from the more to the less dense portion, and a vector v in the direction of the X-ray beam parallel to the reflecting plane, such that both 3 and v lie in the reflecting plane of the crystal, Fig. 1, then the positions of the striations on (200) reflections appear to be confined such that where is the angle between v and 3, Fig. 1. It will be seen later that 3 is the projection of Sb on the reflecting planes. In some instances the strain effects are observed only above the slip plane, Fig. l(C), in others only below, Fig. l(D). Where the strains appear to be both above and below the slip line, Fig. l(E), one sees that there are actually two slip lines present, Fig. 1(F) and (G). Furthermore, the strain patterns are observed in the vicinity of the intersection of emergent edge dislocations (diagonal lines on the micrograph) with emergent screw dislocations (horizontal lines on the micrograph), e.g., Fig. l(HJ) and (MN), but not in the vicinity of the intersection of two emergent edge dislocation arrays, Fig. l(HJ) and (KL). However, in the latter case the contrast of the dislocation array itself is reduced along (KL) and increased along (HJ). Line (AB) in Fig. 1 is the edge of the crystal, and the material to the left of (AB) is the side of the sample. Any empirical correlation will have to include most, if not all of the aforementioned observations. Lang2 and Newkirk1 determined that the condition for contrast is given by g-b#0 [2] where b is the Burgers' vector and g the diffraction vector (normal to the reflecting planes). If one assumes that the direction of the observed strain fields is given by the vector sum of the slip vectors (Sb1) of the intersecting slip planes, then Eq. [2] may be generalized to g-S#0 [3]

Jan 1, 1963

By J. H. Richards, G. Derge, W. Peifer

A WIDE Variety of metallurgical defects in steel have commonly been attributed to the presence of excessive amounts of hydrogen. These defects include flakes in rails and forgings, cracks in welds, and blisters in enameled ware, They have been related to hydrogen largely by circumstantiat evidence because the proper sampling and analysis of steel for hydrogen has been known to be difficult and because no general agreement has existed with regard to the relative suitability and accuracy of available methods. The need for a direct approach to these problems has been widely recognized. This general situation has been well described in the publications of the British committee on the Heterogeneity of Steel Ingots,' and current approaches to the problem were reviewed in a recent symposium of the AIME.I2 The plans described by the present authors at that meeting have now matured satisfactorily and will form the basis of this paper. The principal object of this paper will be to describe the methods of sampling and analysis which have been developed and to illustrate the usefulness and limitations of these methods by citing their application to a specific problem, The first problem selected for investigation was the behavior of hydrogen in rail steel. This was chosen because the occurrence of shatter cracks in rails has received a great deal of study over a long period of time. The defect can be controlled by the expensive procedure of slow in However, it may be hoped that a better understanding of the factors influencing the hydrogen content of rails will lead to a more efficient control of this problem. Rails provide a good starting point for a more general study of hydrogen in steel because they provide a cross-section large enough to develop cracks, Yet small enough for convenient handling and sampling. Rails were also selected because they arc made from plain carbon steel and it was felt that the initial work should not be complicated by the effects introduced by alloying elements. The considerations which determined the original planning of this study should be reviewed. The inherently complex nature of most hydrogen problems requires that a large amount of data must be collected for any particular case. This in turn means that any analytical procedure should be as rapid and simple as possible. The same objective of speed is indicated by the fact that it may finally become desirable to use hydrogen analysis in process control. It is also known that hydrogen is highly fugitive and that any procedure which facilitat~s early analysis of the sample without prolonged or irregular storage has distinct advantages. The most important previously reported methods of analysis may be classified into three general groups: (I) Vacuum fusion, (2) Vacuum extraction (without melting), (3) Total combustion.

Jan 1, 1949

By G. Derge, W. Peifer, J. H. Richards

A WIDE Variety of metallurgical defects in steel have commonly been attributed to the presence of excessive amounts of hydrogen. These defects include flakes in rails and forgings, cracks in welds, and blisters in enameled ware, They have been related to hydrogen largely by circumstantiat evidence because the proper sampling and analysis of steel for hydrogen has been known to be difficult and because no general agreement has existed with regard to the relative suitability and accuracy of available methods. The need for a direct approach to these problems has been widely recognized. This general situation has been well described in the publications of the British committee on the Heterogeneity of Steel Ingots,' and current approaches to the problem were reviewed in a recent symposium of the AIME.I2 The plans described by the present authors at that meeting have now matured satisfactorily and will form the basis of this paper. The principal object of this paper will be to describe the methods of sampling and analysis which have been developed and to illustrate the usefulness and limitations of these methods by citing their application to a specific problem, The first problem selected for investigation was the behavior of hydrogen in rail steel. This was chosen because the occurrence of shatter cracks in rails has received a great deal of study over a long period of time. The defect can be controlled by the expensive procedure of slow in However, it may be hoped that a better understanding of the factors influencing the hydrogen content of rails will lead to a more efficient control of this problem. Rails provide a good starting point for a more general study of hydrogen in steel because they provide a cross-section large enough to develop cracks, Yet small enough for convenient handling and sampling. Rails were also selected because they arc made from plain carbon steel and it was felt that the initial work should not be complicated by the effects introduced by alloying elements. The considerations which determined the original planning of this study should be reviewed. The inherently complex nature of most hydrogen problems requires that a large amount of data must be collected for any particular case. This in turn means that any analytical procedure should be as rapid and simple as possible. The same objective of speed is indicated by the fact that it may finally become desirable to use hydrogen analysis in process control. It is also known that hydrogen is highly fugitive and that any procedure which facilitat~s early analysis of the sample without prolonged or irregular storage has distinct advantages. The most important previously reported methods of analysis may be classified into three general groups: (I) Vacuum fusion, (2) Vacuum extraction (without melting), (3) Total combustion.

Jan 1, 1949

By M. C. Miller, M. R. Tek, D. L. Katz

Previous studies have shown the influence of geometric variations on the drawdown and buildup behavior of a well completed in a homogeneous, porous medium containing a single-phase fluid. Other papers have shown the influence of leakage across caprocks on drawdown tests on water wells and aquifers. This paper presents two additional cases: the influence of an adjacent, expansible source of fluid, and the effect of caprock leakage on the measured in situ permeability and compressibility. Methods are developed for evaluating the influence of these two deviations from a uniform aquifer upon the in situ permeability and compre.ssibility indicated by pump tests. INTRODUCTION In relatively uniform porous media, saturated with a single-phase fluid, the effective reservoir permeability and compressibility of the water porous bed system may be obtained from drawdown and build-up data on wells, such as the initial wells drilled in aquifer storage reservoirs (Homers and Katz et al.7). In considering the drawdown and build-up behavior of wells adjacent to storage reservoirs, or of wells in an aquifer adjacent to an unknown gas field, questions arise as to the quantitative effect of the presence of compressible fluid on the measured permeability and compressibility of the formation. The solution to this problem is presented, using a constant pressure line source at some distance from two wells, one of which is pumped and the other used as the pressure observation well. The solution is presented in terms of the geometry of the wells and their relationship to the constant pressure line. It is known that caprock leakage can interfere with the determination of an aquifer in situ permeability and compressibility (Mantush4, Hantush and Jacob3 and Witherspoon and Muellerll). This paper presents an added case of the effect of leakage through a caprock of known thickness and permeability on well tests starting with a reservoir of constant pressure above the caprock and contributing to the aquifer as a result of the pump test. Again, both a pumping well and an observation well are postulated. These two cases are presented together since both a compressible fluid nearby and. caprock leakage show similar deviations from expected in situ permeability and compressibility results, and the solutions are similar in form. EFFECT OF A NEIGHBORING CONSTANT PRESSURE SOURCE ON IN SITU MEASUREMENTS Bixel, Larkin and van Poollenl presented a paper showing the influence of a discontinuity of the properties of a fluid formation system on drawdown and build-up performance. Their case is related to the present study but may be distinguished from it in that their numerical work and resulting graphs relate only to the performance of the pumping well itself. Incorrect values for in situ reservoir properties may be obtained from pump test data if the skin effect of the well and after-flow are not considered. Analysis of this problem by the authors indicated that the more reliable method of evaluating the influence of an adjacent gas field or expansible source was to use one well for pumping and another well at some distance from it for observing the aquifer pressure performance. An expansible source of hydrocarbon near a pumping well may be approximated by a line of constant pressure for a pump test. Fig. 1 illustrates a pumping well and an observation well near a line of constant pressure. Pressure behavior of an observation well during a drawdown pump test near a line of constant pressure is given by the diffusivity equation:

Jan 1, 1967

By C. M. Schwartz, A. E. Austin, W. B. Wilson

High-temperature X-ray diffraction studies were conducted with Th-U alloys with up to 10 wt pet U. The solid solubility of uranium in thorium as a function of temperature was determined by the method of lattice parameters. Thorium will dissolve up to 2.5 wt pet U at 950°, 4.5 wt pet U at 1150°, and 7.5 wt pet U at 1250°C. Determinations were made of the temperature of the transition of thorium and of the effect of uranium upon the transition. The a to ß transition for thorium was observed to occur at 1330' ±20°C. Mean coefficients of expansion were calculated for thorium and two alloys, and for ThO2 in contact with the thorium, using X-ray lattice-parameter data. Values obtained at 950° C for thorium and Tho2 were 12.1 and 9.40 X 10-6 per OC, respectively. Impurities obtained during the X-ray exposure were identified by diffraction and were essentially Tho2 and ThC, with two additional unknown phases being detected. The effect of the impurities upon the results is discussed. DIRECT investigation (i.e., high-temperature X-ray diffraction studies) of the phase diagram of thorium-rich uranium alloys has been shown to be necessary since recent work1 - has disclosed the presence of an allotropic transformation near 1400°C in pure thorium with the room temperature face-centered-cubic phase transforming to a body-cen-tered-cubic structure at the elevated temperature. The effect upon the transition of the addition of uranium to thorium and of the solubility of uranium in thorium at high temperatures remained unknown, yet was of interest in understanding fabrication procedures and elevated temperature use. The present work was undertaken to provide information in this area by determining the transition temperature, the effect of uranium on the transition temperature, and the solubility of uranium in thorium as a function of temperature. Experimental Work The high-temperature diffraction data were obtained using a camera especially designed for the purpose3 and capable of reaching temperatures in excess of 2000°C at pressures as low as 1 x 10." mm Hg. Temperature regulation was provided by regulating the power input to ±0.1 pet variation, and by regulating the water flow through the camera jacket to provide a constant thermal load. The X-ray sample was a rod nominally 80 mil diam, which was further turned down to 20 mil and then etched to 18 mil over 1/2 in. of one end. This was placed in the sample holder and mounted on the camera so that the smaller part was surrounded by a cylindrical tantalum-sheet radiation-type heating element. Diffraction from the sample was recorded on film after passing through a slot in the heating element and radiation-baffle shield and through beryllium vacuum windows. The X-ray film mounting was of the Straumanis type4 with a camera diameter of 114.59 mm. Since previous work of Chiotti' indicated that impurities considerably alter the transition temperature, chemical analysis of the arc melted iodide crystal-bar thorium samples was obtained prior to testing. The analysis disclosed the material to contain as low as 0.001 ±0,0002 wt pet H and 0.007 ±0.001 wt pet 0. Carbon was 0.003 wt pet and nitrogen less than 0.002 wt pet. This material was sealed in mild steel in an inert atmosphere and subsequently hot rolled to 3/8-in. diam rods at a temperature of 732°C. Following removal of the jacket, the material was pickled and cold swaged to 1/8-in. rods, from which the diffraction samples were prepared. The alloys were similarly prepared, with the uranium being added during arc melting. The uranium analyses of the alloys prepared appear in Table I. The experimental procedure for diffraction examination of the three samples of high-purity thorium differed from those of the Th-U alloys. The original practice, later modified, consisted of pumping down the camera with the diffusion pump on and then admitting liquid nitrogen to the cold trap of the system. This was modified for the Th-U alloys by maintaining liquid nitrogen in the cold trap at all times before and while the diffusion pump was heated. This minor change produced a reduction in the amount of carbon pickup by the sample during exposure to the diffusion-pump vapors. The sample was brought to the desired test temperature and exposed for 21/2 hr at pressures which were usually 2 x 10-6 mm Hg, or lower. Exposures were made at

Jan 1, 1959

By A. D. Hoagland

The Flat Gap mine is located on Copper Ridge in the Appalachian Valley of East Tennessee. Large zinc orebodies occupy zones of solution and collapse breccias in Lower Ordovician Kingsport limestone and dolomite. The principal ore mineral is sphalerite and the principal gangue is dolomite. The ore-bearing horizons are considerably more than 95% covered by thick residual clay. Geochemical prospecting under these conditions is an effective exploration technique which played an important role in the discovery and exploration of the Flat Gap mine. The Flat Gap mine, an important zinc deposit of the Appalachian type, is located in the Copper Ridge Belt in the Valley of Tennessee approximately 30 miles northeast of the Mascot-Jefferson City zinc district (Fig. 1). The ore is found in large stratigraphically linked deposits in solution and collapse breccias of the Kingsport, Longview and, to a minor extent, Mascot formations of Lower Ordovician age. The principal ore mineral is sphalerite and the principal gangue minerals are dolcrmite and accessory pyrite. There is local minor galena and very rare barite and fluorite. The ore-bearing horizons are covered by about 20 to 100 ft or more of residual clay except where active streams have exposed bedrock. Considerably less than 5% of the favorable host rocks are exposed in outcrop, and this nearly complete soil cover creates a serious obstacle to exploration. Comprehensive surveys of the base metal content of soils furnish a useful substitute for outcrop geology, and under favorable conditions, this technique is a very valuable prospecting tool. The analysis is carried out by a titrametric dithizone method modified after the procedure described by E. B. Sandell.' This method detects

Jan 1, 1962

By W. A. Koehler, M. E. Hinkle

THE oxidation of pyrites and marcasite in coal-mine strata to produce discolored acid mine drainage has long been explained by chemical reactions occurring in three stages: 1. The iron sulphide minerals are oxidized by air in the presence of water to produce a nearly colorless solution of ferrous (iron) sulphate and some sulphuric acid. 2, The ferrous sulphate and part of the sulphuric acid combine with additional oxygen to form ferric (iron) sulphate. 3. The ferric sulphate hydrolyzes to form a reddish-brown precipitate and additional sulphuric acid. Nevertheless, chemical data alone are not sufficient to explain all observed facts. The first stage seems to proceed more slowly in the chemical laboratory than in nature. The second and third stages normally occur to a marked extent within a few days with raw mine water. However, the visible discoloration that characterizes these latter stages will not develop for months if the mine water is sterilized. By means of bacteriological techniques, two organisms have been isolated from acid coal-mine drainage. These are believed to have a part in promoting the chemical reactions summarized above. One organism is considered to be Thio-6aciZZus thiooxidans, which is known to be able to convert elementary sulphur or sodium thiosulphate to sulphuric acid. Its relation to the formation of acid mine drainage has not Yet been established. Apparently it does not Promote the oxidation of ferrous iron compounds in mine drainage to the ferric state. The second organism has not been identified by name but is described as " Gram negative, non-spore-forming rods approximately 0.4 micron wide by 0.8 to 1.0 micron long." A pure culture of this when added to sterile mine water will cause the latter to become amber and develop the typical rusty precipitate. However, it has not been found to convert sulphur compounds to sulphuric acid. Further research is planned to determine whether or not these microorganisms play a significant part in the first stage of oxidation whereby pyrites and marcasite minerals go into solution. Some of the conditions found in abandoned sections of coal mines, such as a reduced Percentage of Oxygen and increased percentage of carboll 'dioxide in the air, as well as the existing temperature and acidity ranges, will be tolerated by and may even be favorable to these organisms. Acid Coal Mine Drainage Acid coal mine drainage has been the concern of coal mine operators, industrialists, sportsmen and public officials for many years. The amount of acid flowing into our streams in the coal-producing areas varies considerably from mine to mine and from

Jan 1, 1949

By S. S. Lisnyak, B. D. Lichter

An important question in the study of the "hot corrosion" phenomenon (i.e., accelerated oxidation of Ni-Cr-base and Co-Cr-base "superalloys" in the presence of sulfur) is the condition of formation and physico-chemical properties of different sulfides in the system Ni-Co-Cr-S. seyboltl has observed the presence of ((chromium-rich" sulfides of compositions M3S4, and M2S3 in the scales of superalloys after hot corrosion; however, as he points out, much confusion exists as to the identity of sulfides found on Ni-Cr and Co-Cr alloys. According to the literature,&apos; ternary sulfides, e.g., NiCr2S4, (monoclinic structure) and CoCr2S4 (spinel structure) exist in these systems. Therefore, it is likely that these compounds are found in the scale of alloys undergoing hot corrosion. At present this possibility has not been discussed in the literature. Accordingly, we have begun investigation on the properties of ternary sulfides and their mutual interaction. The magnetic and electrical properties of pure NiCr2S4 and pure CoCr2S2 have been previously determined."-&apos; In the present work we have studied the interaction between these sulfides at high temperatures by means of X-ray diffraction and magnetic susceptibility. The measurements were carried out at room temperature after appropriate thermai treatment of samples at high temperatures. For alloys containing the three elements nickel, chromium, and cobalt this interaction may be of importance in the study of the hot corrosion phenomenon. The compounds NiCr2S4 and CoCr2S4 were prepared from stoichiometric amounts of the elements and CoS in the case of CoCr2S4. Samples were mixed and sealed in evacuated silica tubes, slowly heated to 900&apos; to 100O°C, and held at temperature for a total of 120 hr (NiCr2S4) or 80 hr (CoCr2S4) with intermediate re-grinding~. X-ray diffraction showed a pure spinel structure for CoCr2S4 and a monoclinic structure for NiCr2S4. Samples of mixtures of the two ternary sulfides were prepared with different compositions, see Table I. Samples were heated for a total of 80 hr at 900&apos;C in evacuated silica tubes, with one intermediate regrind-ing. The capsules were slow-cooled to room temperature in the furnace. Samples for X-ray diffraction were prepared by placing powder on a glass slide with vaseline. Dif- fractometer traces were obtained on a G.E. XRD-5 diffractometer using a chromium tube and vanadium filter (40 kv and 31 ma). The size of the spinel unit cell was determined from the eight highest angle peaks (80 deg < 2b < 160 deg). Calculated lattice parameters were plotted vs cos B . ctn 0 and extrapolated to 6 = 90 deg. Measurements of the magnetic susceptibility were carried out at room temperature using the Faraday method in a magnetic field of 3000 G and using a Cahn RG automatic microbalance. The system was calibrated from the well-established susceptibilities of platinum and palladium: (Xpt = 0.968 x 10-6 g-1 and Xpd 7 5.04 x 10~6 g-1). The weight of samples were approximately 0.200 g. X-ray diffraction patterns for unsintered mixtures showed Bragg peaks from both the spinel (CoCr2S4) and monoclinic (NiCr2S4) structures. After sintering no monoclinic peaks were observed in samples containing up to 40 mol pct NiCr2S4. However, the spinel peaks were displaced to higher angles with increasing

Jan 1, 1970

By R. E. Reed-Hill, W. H. Hartt

The "(3034) " twin lamellae in magnesium have been reexamined using electron microscope replicas. The double twinning sequence (1011)- (1012) for these lamellae has been reconfirmed. Evidence is presented showing that the irrational habit of these twins is probably the direct result of a need to accommodate the second-order twinning shear. This accommodation vesults in both an external shear in the matrix and an internal shear in the lamella. A model is presented which accounts for the experimental obseruations. HCP metals are generally characterized by strong textures in which the basal planes of the grains tend to be aligned parallel to the rolling direction in sheet or plate and to the axis of wires and rods. Since these directions are normally those in which finished products may be subjected to uniaxial tensile stress, these textures effectively place the crystals of magnesium objects in orientations unfavorable to both basal slip and {1012) twinning, the easily activated deformation systems of this metal. However, when magnesium crystals are deformed near room temperature by a tensile stress nearly parallel to the basal plane, another mode of deformation can occur.&apos;,2 This is _a double twin which forms first by twinning on {loll), followed by retwinning on (1012). The resulting twin lamellae are of considerable interest. First, they can be responsible for "brittle fracture&apos;&apos;; when a tensile stress axis lies in the basal plane, because the reoriented crystal structure of a lamella is well-aligned for slip while the parent crystal is not. Extensive internal deformation may thus occur while the matrix is still almost undeformed. This deformation can produce a ductile rupture running the length of the twin.3 Complete failure of the specimen may then oc-cure by the spreading of cracks from one twin lamella to another, or by the movement of a crack into a grain boundary and then along grain boundaries. Since the the volume fraction of these doubly twinned elements is normally very small, the fracture may be classed as brittle, although the fracture mechanism is, of itself, essentially ductile. It is quite probable that the ductility transition, normally observed4 in magnesium at or slightly Obove room temperature, is closely associated with these second-order twins. The detailed relationship of these twins to fracture in single crystals will be the topic of another paper. The present paper is concerned with a second important aspect: the irrational habit of these double twins that falls close to {3034), or approximately 7 deg away from the primary (10il) twin habit. There is good reason to believe that this habit results directly

Jan 1, 1968

By R. Ruka, E. A. Gulbransen

ONE of the interesting questions in the understanding of the reaction of iron with oxygen is the kinetics and the mechanism of the crystal structure changes occurring in the formation and breakdown of the oxide film. In a recent work&apos; the electron diffraction method was applied to the study of the solid phase reactions occurring internally in the oxide film on iron above 570°C and under vacuum conditions. The higher oxides which form under oxidizing conditions are reduced internally by the diffusion of iron into the oxide leaving the bulk of the oxide FeO (wüstite). This paper will ask and consider three questions. These are: (1) What is the nature of the depressed transformation temperature for the Fe3O4 + Fe ? 4Fe0 reaction in very thin oxide films?; (2) What is the mechanism of the forward transformation of Fe3O4 to FeO at or near 570°C?; and (3) What is the nature and mechanism of the slow reverse transformation of FeO to Fe3O4 and Fe below 570°C?. The composition of the oxide as well as its electrical, magnetic, mechanical, and chemical properties depends upon the kinetics of these reactions. Literature Survey The composition of the oxide film on pure iron at temperatures up to 570°C has been studied extensively by the electron diffraction method. Nelson,&apos; Jackson and Quarrell,3 and Gulbransen and Hick-man4 have shown that Fe3O4 and a-Fe2O8 are formed in the temperature range of 225" to 570°C, while y-Fe,O, is most likely formed below 225°C for long oxidation times. In very thin films FeO (wüstite) is found to form at temperatures as low as 400°C, and its existence depends essentially upon the extent of initial oxidation. FeO is not found in the formation of thick films below 570°C. On the other hand X-ray5 and micrographic8 studies show that Fe3O4 may exist as the main com-ponent in thick films up to 650°C, well above the equilibrium temperature for the reaction Fe3O4 + Fe ? 4FeO. The kinetics of the forward and reverse trans-formations of the reaction Fe3O4 + Fe ? 4FeO have not been studied in oxide films. However, analysis of the data of Gulbransen and Hickman4 on the heat-ing and cooling of iron oxide films of known thick-nesses in high vacua shows that the forward reac-tion proceeds rapidly even for a 4000 A film at temperatures of 575" to 600°C. The reverse reaction or decomposition of wüstite (FeO) occurs at tem-peratures below 450°C. Thus, there appears to be a slow process in the decomposition which is in marked contrast to the rapid forward reaction. The decomposition of a bulk sample of FeO free from the metal has been studied extensively by Chaudron, Benard and coworkers using the micro-graphic, X-ray diffraction, and thermomagnetic methods. Chaudron7,8 first observed that the de-composition of FeO (wüstite) below 570°C occurs by the reaction 4FeO ? Fe3O4 + Fe. Later work by Chaudron and Forestier5 showed that the rate of de-

Jan 1, 1951

By M. C. Udy, P. C. Rosenthal

The use of minute boron additions to steel has been given considerable attention in recent years. Comparisons made between boron-free and boron-containing heats of otherwise identical analysis have indicated, beyond a doubt, the potent effect of boron for increasing hardenability. The actual hardenability increase has been expressed as a multiplying factor, or in terms of the percentage of other alloying elements boron might replace. The present study presents the results of a systematic survey of the use of boron as an addition to two types of alloyed steels and as a replacement for various percentages of molybdenum in these steels. Both cast and wrought steels of the two basic compositions were prepared. Molybdenum, because of its scarcity at the time the investigation was started, was selected as the element to be replaced by boron. Attention was also given to the effects of boron on other properties of the steels, particularly their notched-bar toughness, to investigate other advantages or limitations of this element as an alloy addition. Experimental Work Steels Investigated The two basic compositions used in this survey of boron as a replacement for molybdenum are given in Table I. The cast steels were higher than the wrought steels in silicon content to improve the casting quality. Table i.—Nominal Composition of the Two Base Alloy Steels composition. Per Cent. steel Mn Si Ni Cr Xi-Cr-Mo Wrought 0.28-0.70-0.15-o.45—o.45-0.32 0.90 0.25 0.55 0.55 Cast 0.28-0.70-0.35-0.45-o.45- 0.32 0.90 0.45 0.55 0.55 hln-&lo Wrought 0.28-I .go-o.0.15-0.32 1.70 0.15 Cast o. 28- I.40- 0.35- 0.31 1.70 0.45 a Sulphur was 0.040 per cent maximum and phosphorus 0.030 per cent maximum. Each of these base steels, both cast and wrought, was made with molybdenum contents of o, 0.10, 0.20, 0.30, or 0.40 per cent, the total comprising 20 steels. In addition, each of these 20 compositions was prepared with boron additions in amounts of o.oo15, 0.003, and 0.006 per cent as ferroboron; and in amounts of 0.002 and 0.005 per cent as two commercial intensifier alloys, designated alloys A and B. The boron-treated steels totaled 100, and the entire series 120. Test Procedure Each steel was melted as an individual heat in a 50-lb., magnesia-lined induction furnace. Ingot iron scrap or mixtures of ingot iron and steel. scrap were used for melting stock. Deoxidation was carried out in the furnace with 0.10 per cent aluminum. The aluminum addition was

Jan 1, 1948

By M. C. Udy, P. C. Rosenthal

The use of minute boron additions to steel has been given considerable attention in recent years. Comparisons made between boron-free and boron-containing heats of otherwise identical analysis have indicated, beyond a doubt, the potent effect of boron for increasing hardenability. The actual hardenability increase has been expressed as a multiplying factor, or in terms of the percentage of other alloying elements boron might replace. The present study presents the results of a systematic survey of the use of boron as an addition to two types of alloyed steels and as a replacement for various percentages of molybdenum in these steels. Both cast and wrought steels of the two basic compositions were prepared. Molybdenum, because of its scarcity at the time the investigation was started, was selected as the element to be replaced by boron. Attention was also given to the effects of boron on other properties of the steels, particularly their notched-bar toughness, to investigate other advantages or limitations of this element as an alloy addition. Experimental Work Steels Investigated The two basic compositions used in this survey of boron as a replacement for molybdenum are given in Table I. The cast steels were higher than the wrought steels in silicon content to improve the casting quality. Table i.—Nominal Composition of the Two Base Alloy Steels composition. Per Cent. steel Mn Si Ni Cr Xi-Cr-Mo Wrought 0.28-0.70-0.15-o.45—o.45-0.32 0.90 0.25 0.55 0.55 Cast 0.28-0.70-0.35-0.45-o.45- 0.32 0.90 0.45 0.55 0.55 hln-&lo Wrought 0.28-I .go-o.0.15-0.32 1.70 0.15 Cast o. 28- I.40- 0.35- 0.31 1.70 0.45 a Sulphur was 0.040 per cent maximum and phosphorus 0.030 per cent maximum. Each of these base steels, both cast and wrought, was made with molybdenum contents of o, 0.10, 0.20, 0.30, or 0.40 per cent, the total comprising 20 steels. In addition, each of these 20 compositions was prepared with boron additions in amounts of o.oo15, 0.003, and 0.006 per cent as ferroboron; and in amounts of 0.002 and 0.005 per cent as two commercial intensifier alloys, designated alloys A and B. The boron-treated steels totaled 100, and the entire series 120. Test Procedure Each steel was melted as an individual heat in a 50-lb., magnesia-lined induction furnace. Ingot iron scrap or mixtures of ingot iron and steel. scrap were used for melting stock. Deoxidation was carried out in the furnace with 0.10 per cent aluminum. The aluminum addition was

Jan 1, 1948