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By Frank C. Kelton

It is perhaps not widely realized that extraction and saturation processes carried out on oil well core samples alter the properties of these samples to varying degrees. On the other hand it is felt by some that quick-freezing of core samples increases their permeability and porosity significantly. Accordingly, laboratory tests were carried out on 49 pairs of horizontally adjacent samples in order to differentiate between the effect of quick-freezing per se on permeability and porosity of the samples, as distinguished from the effect of the identical saturation treatment on permeability and porosity of the companion samples. Also, additional field data were obtained on comparison of frozen vs unfrozen companion samples. LABORATORY INVESTIGATION OF FREEZING us SATURATION EFFECTS Procedure The samples used in these tests were two-cm cubes cut in horizontally adjacent pairs from cores from eight Gulf Coast and Mid-Continent wells, which cores had not previously been frozen. These samples were extracted with carbon tetrachloride, dried, and air permeabilities run in the conventional manner. They were then evacuated and saturated with brine of 25,000 ppm sodium chloride content, and porosities determined by gain in weight. The samples were partially desaturated by evaporation down to an average brine saturation of 68 per cent. One sample from each pair was quick-frozen by covering with dry ice after wrapping in a single layer of paper, and allowed to remain frozen for about two hours; the companion sample from each pair was not frozen. After thawing the frozen sample, all samples were immersed in tap water overnight in order to leach out most of the brine. Air permeabilities were re-run, and the samples were again saturated with brine to determine a second porosity value. For purposes of averaging of data, the samples were grouped according to four permeability ranges, from 0 to 10, 10 to 100, 100 to 1,000, and 1,000 to 3,840 md. Average permeability and porosity changes for the frozen vs the unfrozen adjacent samples are shown in Table 1. Discussion As may be seen from Table 1, the averages of the per cent permeability increases for the quick-frozen samples ranged from 3.8 to 12.9 per cent among the four permeability groups. The average changes among the four groups of unfrozen companion samples ranged from a decrease of 0.2 per cent to an increase of 9.3 per cent. There was no particular correlation of these changes with magnitude of permeability; however, the increase for each group of frozen samples paralleled the increase for the corresponding unfrozen samples. The differences between the two sets of values are believed to be a valid indication of the effect of the quick-freezing in itself, since the treatment of the two samples in each pair was identical except for freezing. The permeability changes which are strictly the result of the quick-freezing are shown in the sixth column of Table 1. These range from a decrease of 0.9 per cent to an increase of 4.0 per cent; the overall weighted average is 1.2 per cent, as compared to an average increase of 6.8 per cent caused by the saturation treatment of the samples not frozen. The average porosity changes are in general smaller than the changes in permeability, and range from a decrease of 2.3 per cent to an increase of 3.3 per cent. The overall weighted average change ascribed to the quick-freezing is 1.0 per cent of porosity. Many factors can contribute to the changes in permeability and porosity observed when subjecting cores to the simple processes used in these tests. Such are: hydration and swelling of clay, adsorption of ions, changes in surface structure and wettability, expansion and compression effects due to ice formation, shrinking and cracking, leaching of salts and colloids, displacement of particles resulting in either blocking or enlarging of pore openings. Whatever particular mechanisms are involved. however, it is apparent not only from this study but also from other investigations in the literature' not directly concerned with quick-freezing, that the effects produced by commonly used extraction, saturation and drying techniques may be of considerable magnitude The results of this study indicate that for the particular samples and techniques used, such effects are of the order of five to six times the effect of quick-freezing. insofar as changes in permeability are concerned. It may be argued that these samples might not include extremely shaly material where the effect of freezing upon permeability may be much greater. However, had such material been available for these tests, it would undoubtedly have been very susceptible also to alteration by the extraction and saturation treatment used. To investigate this point further, the individual sample data were re-grouped according to the magnitude of the average per cent permeability increases for the pairs of samples, irrespective of permeability. The results

Jan 1, 1953

By D. Harker, E. Epremian

The constitution of the nickel-tungsten system has been studied by a number of investigators, the most recent of which are Ellinger and Sykes.1 On the basis of metallography, electrical resistivity and hardness measurements and some X ray diffraction work, they constructed a consitution diagram (Fig 1). Ellinger and Sykes report a0 for the alpha face-centered cubic phase (nickel saturated with tungsten) as 3.66?.‡ The gamma phase (tungsten saturated with nickel) is given as body-centered cubic, but no value of a. is recorded. At 43 pct tungsten by weight (approximately 20 at. pct) a beta phase is reported, but structure and lattice parameter are not known. Primarily on the basis of metallographic evidence, they conclude that beta is a new phase and rule out the possibility of formation of a superlattice since they- observe a new phase in the microstructure. Grube and Schlecht2 report an ordering reaction in a system similar to the Ni-W (Ni-Mo) and Harker3 has shown that NirMo forms a superlattice upon the pre-existing face-centered cubic lattice of the alpha phase. Elliniger arid Sykes observed a change in the diffraction pattern of a 4.3 pct W alloy upon aging and attributed it to the formation of a new phase. Thus, there appears to be some question as to the character of the beta phase reported by Ellinger and Sykes. The purpose of the present work is to determine precisely the structure and nature of the beta phase in the Ni-w system of X ray diffraction methods. Experimental Procedure One inch diameter ingots of 43.66 pct, 40.15 pct and 34.06 pct .tungsten by weight (19.8, 17.6 and 14.1 at. pct respectively) were melted and cast under vacuum from electrolytic nickel and pure wire filament grade tungsten. The ingots were soaked for 16 hr at 1300°C in a hydrogen atmosphere furnace, then swaged at this temperature and finally drawn to 0.020 in. diam wire. Some of the bar stock was retained in 1/2 in. diam rod for hardness and metallographic studies. Wire samples for diffraction work were sealed in evacuated quartz tubes and subjected to the various solution and aging heat treatments in a hydrogen furnace. X ray photograms were taken in a Debeye-Slierrer camera using copper K radiation filtered by nickel foil. Results X RAY DIFFRACTION The lattice parameters of the alpha and gamma phases at saturation were determined by an X ray photogram of the 19.8 at. pct W alloy after solution heat treatment at a temperature well within the alpha-gamma range above the peritectoid temperature (17 hr at 1150°C—oil quenched). The alpha phase was found to be face-centeted cubic with a. of 3.594 ± 0.001?, while the gamma phase was determined as body-centered cubic with a. of 3.158 ± 0.001?. Exactly the same phases and parameters were ohtained with the 17.6 at. pct W alloy for the same solution heat treatment. Thorough analysis was made of an X ray photogram of a 19.8 at. pct alloy sample in the quenched and aged condition (solution treated 1T hr at

Jan 1, 1950

By R. P. Agarwala, M. S. Anand

Using residual activity technique, the diffusion of palladium, silver, cadmium, indium, and tin in alunzinum has been studied in the temperature range of 400" to 630°C. The diffusivities (in units of square centimeters per second) have been expressed as: IMPURITY diffusion in aluminum,1-9 silverand lead5 for cases of low solid solubility of the impurity in the host metal has yielded frequency factors in the range of l0-6 to l0-9 sq cm per sec whereas the activation energy is practically half the self-diffusion activation energy value. From the observed values of frequency factor, activation energy, and entropy of activation, it has been suggested' that these solutes are not diffusing by vacancy or interstitial mechanisms but by a mechanism which should be consistent with such low values of the diffusion parameters (Do and Q). However in spite of extensive work on these types of systems, the mechanism of diffusion is still not well understood. The present investigation on the diffusion of palladium, silver, cadmium, indium, and tin in aluminum has been carried out to throw further light on the diffusion mechanism in systems, where the solid solubility is very low (except for the case of silver). The results are discussed on the basis of solid solubility and the structural changes involved owing to the presence of the solutes in aluminum solid solution. An attempt has also been made to apply the existing theories of charge5-8 and size8 difference between the solute and the solvent. EXPERIMENTAL PROCEDURE Specimens (1/2 in. diam by 3/8 in. high) were machined out of pure aluminum (99.995 wt pct) rod obtained from Johnson Mattheys. They were sealed under vacuum in quartz tubes and annealed at 620° C for several hours; the grains thus developed were sufficiently large to eliminate the possibility of diffusion along the grain boundaries. The flat ends were prepared carefully after polishing as described previously,10 Radioactive nitrates of cadmium, indium, and tin and chloride of palladium containing, respectively, cd115, 1n114, sn113, and pd103 were dissolved in distilled water and mixed with 30 pct acetone. By means of a micropipet a drop of this solution was placed on a smoothly polished and lightly etched surface of the specimen. Due care was taken to see that the solution spread uniformly on the surface of specimen without trickling down its sides. Radioactive silver was elec-trodeposited using a AgCN-KCN bath. The amount of sample deposited in all the cases was not more than 0.1 µ thick. The samples were then sealed in quartz tubes in vacuum. The cadmium samples were sealed in a purified argon atmosphere to avoid evaporation. The samples were then diffusion-annealed. The temperature of annealing varied between 400° and 630°C and was controlled to ±5°C. On heating to -400°C,the deposits of cadmium, indium, and tin, which were of the order of 0.1 p in thickness, were converted to their respective oxides. The contribution of oxygen present in the lattice of aluminum due to these oxides has been calculated and found to be less than 10 ppm in all cases. Oxide method has already been used by other workers11'12 in diffusion studies without any controversy on the issue. However, in some of these investigations, metallic deposition was also tried. The diffusivities calculated from these measurements were found to agree very well with the diffusivities obtained by using the oxide method. Thus it is assumed that the measured diffusivities represent true diffusion coefficients. Since palladous chloride decomposes at about 500°C, the deposited samples which were to be diffusion-annealed below 500°C were heated in vacuum for a very short time at 500°C to allow the decomposition of palladous chloride to palladium metal. Time taken in decomposition of nitrates to oxides and chloride to metal was negligibly small as compared to the period of the diffusion anneals. The residual activity technique13 was used to study the diffusion profiles where thin layers from the specimen surface were removed by grinding it parallel to a flat surface on a 600-grade carborundum paper. The specimen was washed, dried, and weighed, the differ -ence of the weight being the measure of the thickness of the layer removed. After each such abrasion and weighing, the total residual activity on the surface of the specimens was measured by counting 0.656, 0.94,

Jan 1, 1968

By Charles S. Barrett, Robert F. Mehl

A study of the structures arising from the decomposition of the ß solid solutions in the Cu-Zn and the Cu-A1 systems is of peculiar interest in the study of the mechanism of precipitation from solid solutions. It is possible in these alloys, by varying the composition within the ß field, to precipitate from the same basic lattice phases having different lattice structures. It has been rather generally assumed heretofore that the type of structure originating from the decomposition of a solid solution— the Widmanstätten structure—-is determined by the lattice structure of the parent solid solution alone. In the first paper of this series1 evidence was advanced which threw doubt upon this hypothesis; but definite proof of the correctness or incorrectness of the hypothesis should be forthcoming from a study of these alloys, for if the structures originating, for example, from the precipitation of the a phase from ß and of the ? phase from the ß in the Cu-Zn system should be shown to be crystallographically different, the hypothesis would necessarily be held as definitely disproved. Furthermore, the newer proposal enunciated in the first paper of this series—that the origin of the Widmanstätten structure lies in a cooperation between the lattice of the parent solid solution and that of the precipitate-—which at present seems the only a1ternative hypothesis, would in such an event be given added weight. The work here reported shows that the Widmanstätten structures originating in the precipitation of the a phase from the ß and in the precipit'at.ion of the y phase from the y in the Cu-Zn system are indeed crystallographically different; and incidentally presents some new descriptive information on the Widmanstätten structure from a more general point, of view.

Jan 1, 1931

By J. H. Palm

Experiments have shown12 that the formula -S = St - (St - Se')e ne [1] expresses very well the relation between the true stress S and the true form monotonic deformation of plastic metals in single tension and compression (Fig 1). St (asymptotic or final stress), Se' (threshold stress) and ne (characteristic or specific strain) are constants, the significance of which immediately follows from Fig 1. Only at the initial stages of deformation the experimental curve runs below the theoretical curve. This may be due, at least partly, to the uneven stress distribution caused by the anisotropy of the crystals. It is now obvious that a similar formula H = Ht- (Ht- He)e ne' [2] might be valid for the relation between the hardness and the strain. The hardness can be only a function of the strain and must therefore be independent on the hydrostatic tension present in the neck during straining. Hence, if Eq 2 is correct it will also represent the relation between hardness and strain in the middle section of the neck, the strain of which is generally assumed to be almost or fully homogeneous. For aluminum, copper, and several copper alloys it is indeed observed that Eq 2 agrees very well with the experimental Vickers hardness-strain relation from zero strain on up to the strain at fracture. Moreover nc' is, for several metals fairly equal to nc- Hence H = C1 + C2S [3] C1 and C2 are constants during uniform uniaxial tension. Eq 3 was also obtained for copper by Voce.3 According to Hencky's theory4 the ratio of the Brinell or Vickers hardness and the yieldstress (expressed in kg per mm2) of a non-strainhardening metal is approximately equal to 2.8. If Eq 1, 2 and 3 are correct the ratio Ht/S, must accordingly be the same for all plastic metals and be equal to 2.8. It is now experimentally established that this ratio varies between 2.7 and 3.1 with a mean value of 2.9. Taking into account that Ht and especially St can be obtained only by extrapolation from the range of rather low strains, the agreement with the theory is very satisfactory. References 1. J. H. Palm: Appl. Sci. Res. Al (1948) 198. 2. E. Voce: Jnl. Inst. of Met. (1948) 74, II, 537. 3. E. Voce: Metal Treatment (1948) 15, 53. 4. H. Hencky: Zlsch. f. angew. Math. and Mech. (1923) 3, 241.

Jan 1, 1950

By D. G. Russell

Two techniques have been developed with which the applicability of pressure build-up analyses can be extended to include pressure data which previously have been considered virtually unusable. One of the interpretation methods makes possible the analysis of pressure build-up performance during the wellbore fill-up or after production period which occurs soon after a well is closed in. The other technique is an extension of a method for analyzing pressure build-up performance during the late-time portion of the pressure build-up which occurs after boundary eflects first begin to alter the shape of a conventional pressure build-up curve. With both of these methods it is possible to obtain estitnates of the kh product, the skin factor and the reservoir pressure. In addition, with the late-time analysis technique it is possible to obtain an estimate of the contributory drainage volume of the well being tested. This means that in some cases a check on reservoir limit test and (or) material-balance calculations can now be obtained from pressure build-ups. Both methods are slightly more time-consuming than conventional pressure build-up analysis methods because trial-and-error plots of pressure data must be made. The late-time method for analysis of pressure build-ups is in principle applicable to the late-time portion of a two-rate flow test or a pressure drawdown test. The interpretation formulas and procedures for these types of tests are also outlined. In these cases, as with pressure build-ups, it is significant that an estimate of the contributory pore volume is also obtained. On the basis of limited experience with the new techniques, it appears that satisfactory estimates of the kh product, skin factor, reservoir pressure and, for late-time analysis, contributory drainage volume can be obtained. INTRODUCTION The analysis of bottom-hole pressure build-up behavior in closed-in wells has been a a subject of interest in petroleum engineering circles for many years. In fact, few other subjects have received as much attention as pressure buildup analysis methods have. The cause for this interest is essentially twofold in nature. First, the pressure behavior of a well can normally be measured with a reasonably high degree of accuracy so that good data for analysis can be obtained. Secondly, over a fairly wide range of operating conditions, valuable information as to the quality of the reservoir rock and completion efficiency of the well can be obtained at a nominal cost. In recent years, numerous papers have been prepared on the effects of various operating conditions and reservoir heterogeneities on pressure buildup behavior. Very little work has been done, however, on extension of pressure build-up analysis methods to those pressure data which are not amenable to analysis by the present methods. The theory upon which the analysis of shut-in bottom-hole pressure build-up data is based is derived from the solution of the radial flow equation for a slightly compressible fluid for constant-rate conditions. It requires that the well be closed in for a sufficient period of time to obtain a clearly defined linear portion on the plot of observed bottom-hole pressure vs log (t + ?t)/ ?t (where At is shut-in time, and t is producing time to the instant of shut-in). From the slope of the plot and other normally obtainable data, the formation permeability, the well damage or skin factor, and the reservoir pressure at infinite shut-in time (if the reservoir were infinite) can be estimated. The successful application of this procedure depends on being able to recognize the straight-line section on the basic pressure build-up plot. The presently used pressure build-up interpretation theory also assumes that a well is closed in at the sand face and that no production into the well occurs after shut-in. In practice, of course, the well is closed in at the surface, and inflow into the well continues until the well fills sufficiently to transmit the effect of closing-in to the formation. This adjustment period is commonly referred to as the "afterproduction" or "fill-up" portion of the pressure build-up. During the period that the well fill-up effect is most pronounced, the basic pressure build-up plot is nonlinear. At later shut-in times after the effects of a drainage boundary have been felt at the well, deviation from the straight-line behavior of the pressure build-up plot also results. In many cases either of these effects or a combination of both can make the straight-line portion on the pressure build-up plot difficult to recognize. Obviously, an extension of pressure build-up analysis methods to include the afterproduction period and the period in which boundary effects are being felt would be desirable and might render valuable pressure data which for years have been considered virtually unusable. The principal reference of note concerning pressure buildup analysis during the afterproduction period is a paper by Gladfelter, Tracy and Wilsey.1 In the approach of these authors it is necessary to measure the rate of influx into the well during the afterproduction period. This is done through sonic measurements or through measurement of tubing-head and casing-head pressures simultaneously with

Jan 1, 1967

By Frank Tolonen

IRON-ORE beneficiation is becoming of vital importance to the Lake Superior region, since only a few decades will be needed to exhaust the direct shipping ores even if generous allowances are made for discoveries of new orebodies. If the region is to continue as the chief source of ore to the steel industry of the United States, it must concentrate or otherwise beneficiate the enormous tonnages of available iron formations. A beginning along this line has already been made. About one-fifth of the total shipment, even now, is the product of some sort of concentrating process, while another fifth is improved in physical structure by crushing, screening, drying, sintering, or some other similar method. The amount of material amenable to these processes is limited, however; and, in order to be able to utilize the billions of tons of iron formations that cannot be so beneficiated, an intensive study of their nature and concentration possi-bilities is necessary. As long as high-grade ores are available in ample quantities, ore dressing is economically impossible, but as soon as the supply nears exhaustion it becomes necessary. At first the products of ore dressing must meet as nearly as possible the same specifications as the rich ores in order to be adapted to the reduction processes already established. The future requirements as to quality will depend on the one hand upon the ability of the ore dresser to eliminate undesired constituents and charac-teristics in the material, and on the other upon the adaptation of the reduction processes to the product that can be economically supplied. A better understanding of the blast furnace and other reduction reactions, and improvements in the preparation of raw materials, should enable the Michigan districts to continue to supply ore for a long time to come. The study and testing of the iron formations necessary to bring a utilization of the low-grade ores, though requiring time and effort on the part of all concerned, should not be postponed, especially since the economic situation forces the use of the highest grade ore in preference to the leaner material.

Jan 1, 1933

HenRy Clay FRick, a pioneer in modern coke and steel industry and, in more recent years, one of the outstanding financiers of America, died on Dec. 2, 1919, at his home on Fifth Avenue, New York. Although he had been ill since election day, when he had an attack of indigestion which developed into ptomaine poisoning, his death was a surprise to all. He is survived by his widow, Adelaide Howard Childs Frick; a son, Childs Frick, one unmarried daughter, Helen C. Frick, and four grandchildren. A quiet funeral service was conducted at the Frick home on Wednesday afternoon by the Rev. Dr. Leighton Parks, pastor of St. Bartholomew's Episcopal Church. This was attended by members and very intimate friends of the family, Trustees of the Metropolitan Museum of Art, Officers of the Links Club, and Directors of the United States Steel Corpn. Henry Clay Frick was born Dec. 19, 1849, in West Overland, Westmoreland Co., Pa. His father, John W. Frick, was a farmer of Swiss ancestG, and his mother, Elizabeth Overholt Frick, a member of an old Mennonite family. Until he was sixteen, he spent his time at school, on his father's farm, and in his grandfather's distillery, where he kept books. He attended Otterbein University, in Ohio, for a year. Frick, learning that coke was an essential of the already rapidly developing steel business, early invested every cent he could get in coking-coal lands. With the help of an associate of his grandfather, he formed the corporation of Frick & Co., coke dealers, and acquired fifty-one ovens in the Connellsville region and 300 acres of soft-coal lands. During the panic of '73, with the help of Judge Thomas Mellon, a Pittsburgh 'banker, he bought out his partners and while coke was selling at 90 cents a ton enlarged his purchases of suitable lands. Later, the price of coke increased until it was selling at $5 a ton. Before he was 30 years old, Frick was rated a millionaire. Association with CaRnegie In 1878, he sold an interest in the business Do E. M. Ferguson of New York, and later a share to Mr. Ferguson's brother. Four years later the business was reorganized into the H. C. Frick Coke Co. with vast coal lands, great banks of ovens, and a capitalization of $2,000,000. Then began his association with Andrew Carnegie. He exchanged a part of his coke interest for shares in Carnegie Brothers of Pittsburgh,

Jan 1, 1920

By J. E. Hanafee, G. J. London

MANY studies of the deformation behavior of materials under a superimposed hydrostatic pressure have shown that materials brittle at ambient pressure behave in a ductile manner under pressure. Thus, with a metal such as beryllium which possesses relatively low ductility, but otherwise exhibits quite useful physical and mechanical properties, hydrostatic pressure may be particularly useful for both forming beryllium shapes and studying its deformation behavior. In fact, it has been found1"4 that polycrystal-line beryllium in both ingot and powder form appears to behave in a more ductile manner on a macroscopic level in a hydrostatic pressure environment, and it has been suggested2 that this is due to activating a new slip mode. Furthermore, Andrews and Radcliffe5 have found pressure induced nonbasal dislocation activity in hot pressed beryllium. Recently6 it has been shown in "c-axis" compression tests under hydro-static pressures up to 28 kbars that the shear stress needed to cause slip with a Burgers vector out of the basal plane (pyramidal slip) does not change with increasing pressure in beryllium with a purity of some 99.5 pct. This material is equivalent or more pure than the beryllium used in the previous pressure studies. Thus, it appears, as suggested by Inoue et al.,3 that the hydrostatic pressure affects the fracture stress rather than the stress necessary to activate pyramidal slip in beryllium. However, in "c-axis" pressure tests on high purity 12 zone pass beryllium (˜50 ppm total impurities) the macroscopic compression stress needed to cause pyramidal slip was considerably lower than that at ambient pressure.6 It has further been shown that alloying beryllium with nickel and copper in the range 2-5 wt pct also favors the occurrence of pyramidal slip in "c-axis" compression tests,7'8 while lower amounts of nickel and copper do not have significant effects. In the present study the combined effect of hydrostatic pressure and alloying high purity beryllium on the shear stress needed to cause pyramidal slip has been ascertained. A 2.5 wt pct Cu alloy was selected as the first alloy to study as this level of copper did favor pyramidal slip at room pressure. A high purity (12 zone pass) single crystal of beryllium 0.3 by 0.1 by 0.1 in. was cut and polished by an orientation and lapping technique8 so that the top and bottom compression surfaces were parallel and within 3 min of arc to the (0001) plane and the sides parallel to the {l010} and {ll20} planes. In these compression specimens, therefore, the resolved shear stress was nearly zero on both the basal and prism planes, and slip was restricted to pyramidal systems. Analysis of slip traces on the two lateral surfaces served to accurately identify the active slip planes.6'9 The pressure unit was a modified piston-cylinder device fitted with a manganin transducer coil arrangement which continuously monitored and recorded the hydro-static pressure. The load on the specimen was measured by a strain gage load cell which operated entirely within the pressure chamber. This load cell was calibrated before and after each pressure cycle at room pressure in situ and the calibration did not vary more than ±1 pet. These techniques and devices have been previously described in more detail.6 Successively higher compressive stresses were applied to the single crystal under a superimposed hydrostatic pressure until fracture occurred. The strain rate was (4.5 ± 2.0) x 10-6 sec-1 and the average rate of pressure application and release was approximately 0.3 kbars per min. As the load on the specimen was applied by the piston which was used to increase the hydrostatic pressure, the pressure increased during the compression test. This increase ranged from 0.0 to 0.8 kbars, and the maximum hydrostatic pressures are quoted in Fig. 1. The lateral surfaces of the specimen were examined in a light microscope after each pressurization/stress cycle so that the stress at the onset of {1122} pyramidal slip could be ascertained. Post compression height measurements allowed the plastic strain in the specimen to be evaluated to within 0.03 pct. The resulting compression stress-plastic strain curve is shown in Fig. 1 with results of a "c-axis" test on a similar Be-2.5 wt pct Cu single

Jan 1, 1970

By T. S. Lundy, R. E. Pawel, F. R. Winslow, C. J. McHargue

The volume-diffusion coefficients of Nb-95 and Ta-182 in niobium have been measured over the temperature range 878° to 2400°C. High-temperature specimens (T 21500°C) were sectioned by conventional lathe and grinding techniques while low-temperature specimens were sectioned by an anodizing and stripping technique previously described. Over a range of more than ten orders of magnitude in diffusivity, the self-diffusion data is well-represented by the equation: D = (1.1+ 0.2) exp [- (96,000 * 900)/RT ] cm2/sec Upward deviations from this line are explained by short-circuit diffusion in combination with the loss of resolution of the sectioning techniques. Oxygen contamination within the range 100 to 450 ppm was found to have little or no effect on the self-diffusivitj 01 niobium. In a limited number of measurements, the diffusivity of Ta-182 in niobium was found to be about half that of Nb-95 in niobium. These data may be described by D = (I.0+0.90.5)exp [- (99,300 * 2400) /RT ] PREVIOUSLY reported studies of self-diffusion in niobium were by Resnick and Castleman1 and Peart, Graham, and Tomlin.2 The data of Resnick and Castleman over a temperature range 1585" to 2120°C show considerable scatter but. can be represented by D = (12.4 *0.8) exp [-(105,000+3OOO)/RT] cm2/sec The limited data of Peart, Graham, and Tomlin follow the equation D = 1.3 exp (-95,00O/RT) cm2/sec In view of recent diffusion data for ß zirconium3 and 0 titanium4 in which Arrhenius-type plots do not yield straight lines and of the potential high- temperature uses of niobium and its alloys, we thought it important to investigate the diffusion behavior of this bcc refractory metal over a large temperature range. This has recently been done for molybdenum,5 chromium,6 vanadium,"" and tantalum.9 Niobium was a particularly suitable material for this study largely because of the development of an anodizing and stripping sectioning technique which enables diffusivities as small as l0-19' sq cm per sec to be determined accurately.10 In addition to self-diffusion data, data were obtained for the diffusion of Ta-182 in niobium and for the effect of varying oxygen content on self-diffusion. I) EXPERIMENTAL PROCEDURE Materials. The niobium was obtained from three sources: 1) melting stock from Shieldalloy Corp., 2) electron-beam melted single crystals, 1/2 in. diameter, from Semi-Elements, Inc., and 3) oxygen-doped, electron-beam melted single crystals from Materials Research Corp. The melting stock was electron-beam melted and specimens 5/8 in. diameter by 1/2 in. long were machined from the center of the bar. The analysis supplied by the vendor and an interstitial analysis of one of the specimens are given in Table I. The single crystals (type 2 above) were cut with a jeweler's saw to specimens about 3/8 in. long. The interstitial analysis of this material is given in Table 11. This material was used

Jan 1, 1965

By E. S. Machlin, Morris Cohen

The isothermal formation of martensite in a 71 pct Fe, 29 pct Ni alloy is found to take place mainly by the nucleation of new plates rather than by the growth of existing ones, and is dependent on the temperature and time of the isothermal hold, the amount of atherrnal martensite present, the state of internal strain, and the application of tensile stress. The conclusion is reached that the isothermal nucleation is activated by thermal fluctuations superimposed on localized regions of very high strain. The reaction-path concept, involving martensite nucleation at strain embryos by athermal activation on cooling, or by thermal fluctuations on holding, reconciles the athermal and isothermal modes of transformation, and offers a unified picture of the kinetics of martensite formation. THE phenomenon of isothermal martensite formation has been the subject of considerable study lately,'-' and its reality has been thoroughly established. The existence of such isothermal transformation has been claimed to confirm the nucleation and growth hypothesis as applied to the martensitic reaction.1 " However, the isothermal formation of a martensitic phase is also compatible with the reaction-pathw escription of the transformation. The purpose of this investigation is to examine the isothermal mode of the martensitic reaction, both theoretically and experimentally, in order to test the likelihood of the above two mechanisms. Material and Methods The alloy used in this study contained about 71 pct Fe and 29 pct Ni. The complete analysis was 29.5 ± 0.2 pct Ni, 0.036 pct C, 0.02 pct N, 0.19 pct Mn, 0.09 pct Si, 0.008 pct P, and 0.006 pct S. The amount of isothermal transformation was determined by electrical resistance measurements using a Kelvin double bridge. The specimens were in the form of rods 1/16 in. diam x 4 in. long. Aus-tenitizing was performed at 1095°C for % hr either in a pre-purified nitrogen atmosphere or in an evacuated Vycor tube, followed by oil quenching to room temperature. After this treatment, the samples were completely austenitic, and the transformation studies were conducted at subatmospheric tempera- tures. Lineal analysis and X-ray determinations of retained austenite were used to calibrate the electrical resistance changes in terms of the percentage of martensite. The martensite range curve for this alloy is given in Fig. 9 of ref. 6. The M, temperature occurs at about -18°C (255°K). Reaction-Path Theory The nucleation and growth description of the martensitic transformation that occurs isothermally has already been given by Kurdjumow.','2 However, no comparable treatment exists using the reaction-path concept." Consequently this section presents the reaction-path description of the isothermal mode of the martensitic transformation. According to the latter theory, there is a reaction path (a strain) which describes the relative motion of the atoms during their transference from the austenitic to the martensitic state. The possibility exists, therefore, that the nucleation of a martensitic plate is accomplished by an embryo that comprises a state intermediate between that of austenite and martensite. This state may be attained within a finite volume by means of a homogeneous strain of the austenitic matrix. There is a free-energy barrier (F.) along the reaction path, corresponding to a critical strain in a critical volume. When activation is achieved, a spontaneous increase in strain and volume of the embryo takes place to generate the full-size martensitic plate. This mode of nucleation is to be distinguished from the more conventional mode involving embryos of the martensitic phase that merely grow in size by an interface motion. There are two ways whereby embryos in the austenitic phase can achieve the critical state of strain for isothermal nucleation: 1—through the formation of internal strains by plastic deformation or by stress; and 2—through thermal fluctuations, preferably superimposed upon existing internal strains. The activation free energy F,, must be supplied either mechanically by internal strains or thermally by energy fluctuations, or by a combination of both.

Jan 1, 1953

By D. M. Collingwood

The year 1928 has been marked by scattered but considerable wildcat drilling considering the overproduction prevalent in the oil industry. This wildcatting has resulted in the discovery of oil in one interior salt dome, several small fields producing from serpentine, Taylor marl, and Austin chalk, and two new fields producing from the Edwards limestone, one of which in the Bruner or Salt Flat area of Caldwell County promises to be of major importance if not quite as large as the Luling field. Fig. 1 shows production by fields for 1928, compared in the case of the older fields, with that for 1927. New Production From Salt Domes Boggy Creek Dome This salt dome was discovered in 1927 by core drilling and deep tests recommended on the evidence of surface geology. A large block of acreage controlled by the Humble Oil & Refining Co. covers the dome. Since the discovery of the deep-lying salt plug and oil on the flank of this dome in 1927, development has proceeded slowly. Since most of the acreage is controlled in a large block by one company there has been no intensive competitive drilling. Approximately 8 producers and 16 dry holes have been drilled. Production is from the eastern equivalent of the Woodbine and Eagleford horizons of the Mexia district, and has been found extending in a northeast-southnlest direction across the River Neches on the southeast flank of the dome. One of the best wells, the T. H. Jones No. 2, had an initial production of 90 bbl. per hr., and others have made from 100 bbl. per day to 50 bbl. per hr., but production of the better wells has either been shut in or choked down to less than 2000 bbl. per day. At the end of the year the daily average, production was about 2500 bbl. from 8 wells. Considerable gas pressure is present and one well, the Elliot Clark No. 2, drilled in 1927, was shut in with an estimated wet gas production of over 5,000,000 cubic feet.

Jan 1, 1929

By Henry Fay

The erosion of guns is a complex problem which can be solved only by a detailed study of all the factors involved. In the present paper it is proposed to submit the results of observations and experiments which have extended over several years. Certain facts have been established and preliminary reports have been published.' Conclusions have been reached since then which, it is believed, are worth presenting for the purpose of promoting discussion, and with the hope that ultimately the whole truth will be known. The particular phase of the problem which will be presented here is the hardening of the inner surface of the gun tube. It has been known for a long time that after firing a large caliber gun for some time, the surface of the metal becomes hard and brittle, cracks, and wears away.2 The character of this phenomenon may be learned by an examination of some results which are typical, obtained from a detailed study of a 12-in. gun. The gun was trepanned and rings representing various sections of the gun were cut out for examination. The first ring represented the metal at the muzzle end, and the other rings were removed at approximately 120, 240, 325 and 337 in., respectively, toward the breech end. The appearance of the surface of the metal in the different sections is shown in Fig. 1. The greatest amount of wear and maximum amount of hard surface layer was found in section marked E, which is nearest the powder chamber and at the beginning of the rifing; the maximum amount of cracking in D; the maximum amount of copper deposition in C. Section B showed heat cracks on both lands and grooves, and section A showed heat cracks on the driving edge of the lands only. There was much flow of metal in both sections E and D, and there was progressively less effect noticeable toward the muzzle end. The polished cross-section of these same pieces is shown in Fig. 2. The extent of the erosion is greatest nearest the powder chamber, and diminishes toward the muzzle end.

Jan 1, 1917

By C. S. Barrett, O. R. Trautz

PREVIOUS investigations have shown that lithium is body-centered cubic from near its melting point to the temperature of liquid air1,2,3 Nevertheless there was an incentive to search again for a transformation at low temperatures: C. Zener4,5 has remarked that the low value of the elastic modulus (11 - C12)/2 of body-centered cubic metals and alloys implies a rapid increase in free energy with decreasing temperature, which would favor the occurrence of a transformation; also the shear movement to which this modulus applies is one that can transform a body-centered cubic structure into a slightly distorted face-centered cubic structure. It was confirmed that lithium does not transform spontaneously at liquid nitrogen temperature, but it was found that a transformation can be induced at this temperature by cold working the metal,6 and at slightly lower temperatures a different transformation occurs spontaneously. Transformations were also found in solid solutions of magnesium in lithium. THE NATURE OF THE TRANSFORMATIONS Before presenting a detailed treatment of the technique used and the specific data obtained, it would be well to give a brief survey of the general characteristics of the transformations. A new crystal structure begins to form spontaneously in a specimen when it is cooled below a critical temperature. In lithium the structure appears to be close-packed hexagonal; in the Li-Mg alloys most of the diffraction lines (but not all) are near those for hexagonal lithium. The transformation proceeds as the temperature is lowered and soon stops if the temperature is held constant or raised. The course of the transformation on cooling is indicated by the curve at the lower left of Fig I. This curve is drawn in stepped fashion because the transformation goes by discrete steps with audible clicks. Because of the similarity to the decomposition of austenite into martensite in steels, it seems appropriate to designate the beginning of the spontaneous transformation on cooling as the Ms point. On heating to room temperature after the low temperature treatment only the usual BCC structure is found. On heating, the low temperature modification begins to disappear at a temperature which will be designated as Ac. The transformation on heating is halted when the rise in temperature is halted. The temperature must be raised 40 to 50°C above Ac before the reversion is complete-an interval which is sensitive to the rate of heating and to the time of holding at temperatures above Ac. This behavior on heating thus has similarities also to the formation of martensite in the cooling of steels.

Jan 1, 1948

On August 22, 1966, just 20 months after writing sales agreements with Japanese steel mills, Hamersley Iron Pty Ltd made its first contractual shipment of 52,000 tons to Yawata Iron and Steel Co Ltd. As world news, the event was not spectacular. But to mining men it represented an astonishing achievement. Twenty months of feverish effort and the expenditure of more than $100 million had created a large open-pit iron ore mine, a crushing plant, storage and load-out facilities for sized ore, and a supporting township at Mt Tom Price; a 182-mile standard gauge railway; and port and loading facilities for a nominal loading capacity of 6000 tph as well as a port town, Dampier, on King Bay. The project was executed in the sparsely inhabited, arid, dissected-plateau country of the Hamersley Range in the northwest sector of Western Australia. Hamersley, 60% owned by Conzinc RioTinto of Australia Pty Ltd and 40% by Kaiser Steel Corp, holds sale contracts for ore and pellets valued at $770 million. The main contracts are with the Japanese, who are to be supplied with 65.5 million tons of ore during the 16 years from August 1966 to August 1982 and 18 million tons of pellets over a 10-yr period beginning in early 1968. Of lesser importance at this time, but presaging trends, are contracts for a total of nearly 1.5 million tons to be shipped to England and Western Europe during the years 1966-68. A pellet plant with a capacity of 2 million tpy is now under construction at Dampier and is to be on stream early next year. Another interesting project for which the plans are set, is a desalination plant for sea water to supply industrial requirements and potable water for the town of Dampier.

Jan 5, 1967

By E. R. Gilbert

Klopp et al. have reported data on tensile and creep properties of are-melted and electron-beam-melted tungsten. We would like to point out some similarities between their creep results and ours on are-melted and powder-metallurgy tungstenand offer some alternate interpretations of the data which lead to different rate-controlling mechanisms. Klopp et al. observed extensive transient creep in the 2250" to 3500°F temperature range where the acthetivation energy is 106,400 ± 3300 cal mole-1. They suggest cross slip as the possible rate-controlling mechanism since "cross slip may occur only by the movement of screw dislocations, and not edge dislocations, thus allowing only partial recovery of the strain hardening and resulting in a continuously decreasing creep rate, characteristic of transient creep.'' Our tests on are-melted tungsten and powder-metallurgy tungsten displayed an extensive transient creep stage in the same temperature range but they also displayed a steady-state creep stage. The stress levels and strain rates of the two investigations, 1.4 to 7 kg mm-2 and 10-7 to 10-4 sec-1 for that of Klopp et al. and 4.5 to 10 kg mm-2 and 10-8 to 10-4 sec-1 for those of Gilbert et al., are overlapping. The activation energy for creep of our powder-metallurgy tungsten is 105,000 cal mole-1. This close agreement with the value found by Klopp et al. suggests that the same mechanism may be common to both types of materials. It may be possible to gain some additional insight into the rate-controlling mechanism by applying the analysis made by conrad22 in terms of the activation volume. The value for the activation volume depends upon the dislocation mechanism and may be used to help differentiate between various dislocation mechanisms. The activation volume for cross slip is 10 to 100 b3. Values determined from our data were 130 to 220 b3 which are in better agreement with the range of 102 to 104 b3 for nonconservative motion of jogs. Values determined from data of Klopp et al. are 240 to 1040 b3. We suggest that the rate-controlling mechanism for creep in this temperature range may be nonconservative motion of jogs assisted by enhanced diffusion along grain boundaries or dislocations. This model requires an activation energy lower than that for self-diffusion as is observed. If this is the controlling mechanism, a higher activation energy which is closer to that for self-diffusion might be expected for tests conducted on specimens with reduced grain boundary area, i.e., larger grain size. In support of this premise we have conducted tests on speci- mens having one to three grains per cross section for which a higher activation energy of 130,000 cal mole-1 was observed. Also, in agreement with results of Klopp et al., a faster creep rate resulted from the increased grain size and probable decreased dislocation density. Further argument against the cross slip mechanism is based on the constant activation energy we obsserved over the range of 5.0 to 10 kg mm-2,21 whereas the activation energy for cross slip displays a significant stress dependence.23 In addition to this, the powe on the stress term for the strain-rate dependence appears too high for cross slip. The values for tungsten fall in the range of 4.5 to 7 (Klopp et al. and Ref. 21) whereas values of approximately 2 are observed in creep by cross slip for hep metals.24 At temperatures from 3300° to 4000°F, Klopp et al. observed an activation energy of 141,000 ± 4000 cal mole-' which is less than that for self-diffusion, 153,000 cal mole-1, 25 and less than that reported for high-temperature creep of powder-metallurgy tungsten by Green:' 160,000 cal mole-1. Klopp et al. suggested that recovery of strain hardening by dislocation climb is the rate-controlling mechanism in this temperature region. Actually an activation energy larger than that for self-diffusion is predicted by this model if the temperature dependence of the elastic modulus27 is taken into consideration. Klopp et al. attributed this lack of correlation to the compositional differences between powder-metallurgy tungsten used by Green and their own arc- and EB-melted tungsten. It is our position that the value of 141,000 ± 4000 cal mole-' represents a transitional value between the higher and lower activation energies and does not represent a discrete mechanism of deformation. If the tests had been extended to higher temperatures the higher value of activation energy would have been observed as has been demonstrated by both our analy sis'' of isothermal test data by GE-NMPO28 and differential temperature tests28 on are-melted tungsten. The activation energies determined (170,000 and 160,000 cal mole-1) were greater than that for self-diffusion and are in better agreement with the value found by Green. The model of dislocation climb is associated with ; activation volume of 1 b3 according to Conrad.22 The activation volume calculated for the creep data of fou independent investigators in the temperature range where the activation energy for creep is near that for self-diffusion (Klopp et al. and Refs. 26, 28, and 30) is 110 to 2800 b3 and is in the range suggested for nonconservative motion of jogs. In the case of tungsten this mechanism may be preferred over that of dislocation climb. Sincere appreciation is expressed to Messrs. J. E. Flinn and F. L. Yaggee for the helpful discussions and suggestions in the preparation of this communication.

Jan 1, 1967

By R. J. Brison, W. C. Bleimeister

The International Salt Company has long been interested in finding an efficient process for the removal of impurities from rock salt, and particularly from the rock salt produced at the Detroit mine. Methods which have been used previously at the Detroit mine included both hand sorting and differential crushing in a Bradford breaker with rejection of the oversize. Neither of these methods was adequately effective. The development of a new process for removal of impurities from rock salt has been described in the preceding article." Basically, the process consists of exposure of crude rock salt to radiant heating which selectively heats the impure particles. Subsequently the heated particles are removed by adhesion to a heat sensitive-coating on a conveyor belt. This method was found to be effective in improving both the quality and appearance of rock salt by the rejection of the dark, impure particles.

Jan 8, 1960

By Benjamin Tillson

To MOST mine operators, it seems evident that there is a drill-steel problem, although under certain conditions the amount f drill-steel breakage does not appear serious. What is at fault? It may be one or a combination of circumstances. The development of rock drills to the hammer-drill type in place of the old reciprocating piston drill, probably is one important cause for the greater steel breakage. Perhaps the manufacturers of drill steel have failed to realize what alloys are needed for this new service or have overlooked the changes in the art of rock drilling. Although suitable alloys are provided, they may not be so handled in the manufacturing process as to be in the best condition to withstand the demands of rock drilling. The mine operator may be at fault in desiring to use the smallest possible drill-steel section in order that the gage of the drill bits may be correspondingly small and the amount of footage drilled per unit of labor may be greater; or the black-smithshop practice at the mines may need improvement. The miner who uses the drill steel may also require more intensive supervision and education. Again, the manufacturer of rock drills may have failed to study the types of blows the steel alloys will withstand satisfactorily and which their tools are delivering and so he may not know whether or not a different design of rock drill would equal or excel their present drilling speeds without treating the drill steel so severely. All of these hypotheses probably have their supporters. On the other hand, perhaps we are seeking too much service from drill steel and we need a fuller realization of the fatigue strains developed and should prepare to relieve these by a periodic heat treatment of the steels. However, we do not find any proof that any of these suppositions may be held responsible for the drill-steel breakage attendant to mining. The proposed investigation of this matter by the U. S. Bureau f Mines is fully warranted in the promotion of the conservation of labor and material, and the safety of the workmen. The great range of field conditions, as well as the scope of the research, requires the energetic cooperative support of many interests. Various manufacturers of steel,

Jan 5, 1921

By M. R. Geer, H. F. Yancey, P. S. Jacobsen

Two growing problems confront the preparation engineer—still further restrictions on stream pollution and a greater proportion of fine coal as more and more continuous miners come into use. The de-watering screens, centrifuges, and settling ponds that sufficed a few years ago must often be supplemented by more effective equipment, and in some instances the finest solids are now being recovered by vacuum filtration, once considered too costly in many coal washeries. Unfortunately not all slurries can be readily filtered. Some are virtually unfilter-able, and others do not permit the high cake rates needed to hold costs within reason. Often, however, these difficult slurries can be rendered filterable by flocculation. Starch and lime have long increased settling rate in thickeners and can also aid filtration. More recently there have been other flocculants, primarily synthetic polymers or gums. Reports of their great effectiveness prompted the U. S. Bureau of Mines to test their use in filtration as part of its program on recovering and cleaning fine coal. The object of the present work was to compare some of the newer flocculants with starch and lime and to test the reaction of different types of slurries to flocculation. Three natural slurries, three synthetic coal-clay mixtures, and five flocculants were tested, all with a laboratory filter leaf. Slurries and Flocculants Tested: First of the slurries was the thickener feed from the washed coal section of the Michel colliery of Crow's Nest Pass Co. Ltd., Michel, B.C. As shown by the screen and ash analyses in Table I, practically all this slurry was finer than 28 mesh, and about half the solids were finer than 200. Substantially all material coarser than 200 mesh was clean coal, but the ash content of 20.7 pct in the finest size indicates a moderate amount of impurity. This coal is medium-volatile bituminous in rank. Slurry was also obtained from the Black Diamond washery of Palmer Coking Coal Co., King County, Wash. This underflow of a 1/2-mm vibrating slurry screen is a waste product discharged to a settling pond. As shown by the data in Table I it contained about 20 pct material finer than 200 mesh, which analyzed 77.5 pct ash, indicating a high proportion of clay. The coarser sizes were also high in ash content because of the presence of bone and shale. Black Diamond coal is on the dividing line between bituminous and subbituminous rank. The third slurry, from the washery of Roslyn-Cascade Coal Co. in Kittitas County, Wash., was the underflow of a battery of 8-in. cyclones that is wasted with the coarser refuse from the plant. In both size and ash content this slurry was similar to the Black Diamond, but since it was a cyclone-underflow product, the fraction finer than 200 mesh undoubtedly contained a much smaller proportion of finely divided clay than was present in the Black Diamond slurry. The Roslyn-Cascade coal is high-volatile A bituminous in rank. Initially the slurry samples were stored in the laboratory at a concentration of 40 pct solids. A progressive change in filtration and flocculation characteristics during storage was noted, however, presumably because of progressive disintegration of the clay. After careful testing demonstrated that the slurry solids could be dried and then repulped without changing the original filtration characteristics, this procedure was adopted. The flocculants used were Separan 2610, Kylo 27, Jaguar MD-A, Idaho potato starch, and lime. The lime was used as a slurry, the Separan 2610 as a 0.10 pct aqueous solution, and the other flocculants as 0.50 pct aqueous solutions. The starch was causti-cized with sodium hydroxide before use. As some of the flocculants are reported to deteriorate on storage, new batches of flocculants were prepared every two days. Test Procedure: Fig. 1 diagrams the filter leaf apparatus employed. This consists of a test leaf having an effective area of 0.10 sq ft., a 60°, 8 1/2-in. diam cone to hold the slurry sample, a variable-speed electric stirrer, vacuum gage, bleed line for controlling vacuum, filtrate flask, moisture trap, and vacuum pump. Two filter cloths were selected—Saran SA 603, a monofilament, coarse-weave cloth, and polyethylene PO-801 HF, a monofilament, fine-weave cloth. The Saran was selected to provide maximum cake rate without regard to clarity of filtrate and the polyethylene to provide the best quality of filtrate without regard to cake rate. The test charge for the filtering vessel consisted of 800 g of solids (previously soaked for 24 hr) and

Jan 1, 1960

By A. B. Kinzel

IT is with sincere appreciation and a deep sense of responsibility that I accept the honor of delivering the Howe Memorial Lecture. In our time metallurgical research has delved into phenomena ever more complex in character, so that much of this work must be done by organized teams. We, of the Union Carbide and Carbon Research Laboratories and the Electro Metallurgical Co., are truly a team and the lecture which is to follow is due largely to a few of the outstanding members of that team here cited and well known to you in their own right: John Lamont, whose work deserves special recognition, Walter Crafts, William Forgeng, William Binder, Robert Fowler, and David Swan. Thus it is a fitting pleasure to accept this honor for the research institution which I represent. Henry Marion Howe was truly inspirational and his breadth of view was such that he could use either the method of rigid proof or implied mechanism to further the broad understanding. But perhaps his outstanding characteristic was his sound judgment and his ability to weigh observations. The following is replete with observations and judgment as to their implication. We would like to feel that it is the type of dissertation which Howe, himself, would have enjoyed. Chromium carbides in stainless steel have, in a sense, been a key to the progress of the stainless steel industry. Brearley's first stainless steel, containing approximately 0.30 pct C, was successful because it differed from previous Fe-Cr alloys in that the carbides were present in such amount and form as to provide the stainless and cutting properties. Again, the 12 pct Cr die steels with a carbon content of 1.5 pct or more achieved their properties by virtue of the form and distribution of the chromium carbides. In the straight chromium steels with from 12 to 25 pct Cr, chromium carbide amount and distribution are once more the determinants with respect to physical properties and corrosion resistance. In each of the cases just cited the role of the carbide is in essence similar to that of iron carbide in carbon steel and this story is so well known that we need not dwell on it. In the austenitic stainless steels, the chromium carbide plays only a minor positive role with respect to mechanical properties, and for corrosion resistance the objective is to achieve maximum homogeneity. Thus the tendency has been to ever lower the carbon and correlatively the chromium carbide content of the austenitic steels. Today stainless steels with carbon less than 0.03 pct are beginning to find their true place, and this is being aided in major degree by our Company's introduction of a new type of ferrochromium having negligible carbon. The role of carbon in these austenitic steels has long been recognized, and a truly enormous amount of research has been devoted to the subject. Much has been learned about corrosion resistance, but the attack on the problem from the point of view of carbide precipitation and the nature thereof has left much to be understood. Because of its fundamental importance to the future of stainless steels, further study of the mechanism of carbide formation in austenitic steels is of interest. In the 18-8 Cr-Ni steels, which are typical of the austenitic class of steel, carbon is soluble in the aus-tenite to the extent of approximately 0.15 pct at 1832°F (1000°C). The solubility becomes less with decreasing temperature, so that at 1652°F (900°C) the solubility limit is approximately 0.06 pct, at 1472°F (800°C) it is about 0.03 pct, and at 932°F (500°C) it is approximately 0.01 pct. Obviously this is the ideal situation for the classical precipitation phenomenon. It is true that, in a sense, 18-8 austenite resulting from annealing is in a metastable state, and the phenomenon of ferritic precipitation from this austenite should be considered. However, in a so-called fully austenitic 18-8 steel this is secondary with respect to the carbides and can be neglected. A wide variety of chromium carbides is known. Only one chromium carbide, Cr,,C,,, frequently referred to as Cr,C, has been found in unmodified low carbon 18-8 steel. This, fortunately, simplifies the study of the precipitation phenomena. For this study an A.I.S.I. type 304 stainless steel of commercial quality was selected. Analysis showed C 0.07 pct, Mn 0.49 pct, Si 0.33 pct, Ni 9.34 pct, Cr 18.91 pct, and N 0.033 pct. As expected with this composition the steel was homogeneous after annealing except for a few pools of ferrite, which were disregarded in that their location was not coincident with subsequent carbide precipitation areas. On corrosion testing the steel was normal and typical of good commercial quality. Specimens 1/2xlx4 in. were annealed and subjected to precipitation heat treatment for 100 hr at 100" temperature intervals from 1000° to 1500°F and for ten selected intervals ranging from 5 min to 64 hr at 1300°F. These blocks were then cut into small pieces for metallographic studies. The first problem

Jan 1, 1953