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By J. L. Kenty, J. P. Hirth

The concept of a critical super saturation, below which the nucleation rate is essentially zero and above which it is essentially infinite, is discussed with reference to vapor-solid nucleation. The necessary and sufficient conditions deduced for observations of this type of behavior are: 1) the nucleation rate must exhibit a sharp dependence on super saturation, 2) the growth rate must be sufficiently large that nuclei become observable in the time period of the experiment, and 3) the number of highly preferred nucleation sites must be small. Experiments reveal that the nucleation of silver on sodium chloride is visually detectable at all experimentally accessible super saturations and does not exhibit critical nucleation behavior. Failure to observe a critical super saturation is attributed to the insensitivity of nucleation rate to supersaturation as a consequence of the particular values of the contact angle and the surface free energy for this system. THE concept of a critical supersaturation, below which the nucleation rate is essentially zero and above which it is essentially infinite, arises naturally in homogeneous nucleation theory. Experimentally this type of behavior has been found by Volmer1 and others for water and other low surface tension liquids, as reviewed by several authors.2'3 The same type of behavior has been predicted and observed for heterogeneous nucleation of solids by Yang et al.4 and others,596 as also recently reviewed.2,7,8 In the work reported here on the heterogeneous nucleation of silver on NaC1, however, no critical super-saturation was found. Similar observations have been made recently for other systems.9-11 These results led to a reexamination of nucleation theory which revealed that there are conditions for which critical behavior is not predicted, either for homogeneous or heterogeneous nucleation. Although heterogeneous nucleation is of primary importance in this paper, some insight into critical behavior for such a case can be gained by considering homogeneous nucleation. Accordingly both types of nucleation theory are reviewed briefly. The requisite conditions for critical supersaturation behavior are then considered. The experimental results for the nucleation of silver on NaCl are presented and interpreted in terms of the theoretical presentation. REVIEW OF NUCLEATION THEORY There are essentially two approaches to nucleation theory, the so-called classical theory involving the concepts of bulk thermodynamics, and the statistical mechanical theory in which nuclei are regarded as macromolecules. The classical theory is based on the work of Volmer and Weber12,13 and Becker and. Doring14 and has been extended by Pound et al.15 The crucial assumption in the classical theory is that the small clusters or nuclei can be characterized by the same thermodynamic properties as those of the stable bulk phase. Thus, the nuclei are assumed to have a surface free energy, y, and a volume free energy of formation (relative to the vapor phase), ,, identical to that of the bulk. For deposition under low super-saturation conditions, the nuclei are large and this assumption is satisfactory. However, in many cases of interest, the nuclei contain only a few atoms and this assumption is highly questionable. The statistical mechanical models originated, for the specific case of a dimer as the critical nucleus, with the work of Frenkel16 and were extended later to larger sizes by Walton,17,18 Hirth19 and, more recently, Ht Zinsmeister. These models describe the nucleus in terms of a partition function, the estimation of which is tractable for clusters of 2 to 10 atoms, but extremely difficult for clusters larger than 10 atoms. Although the classical and statistical mechanical models are expected to apply for the limiting cases of large and small nuclei, both are uncertain for intermediate sizes. In this paper we shall treat only the classical model, recognizing that it is exact only for large nucleus sizes and regarding it as a phenom-enological description for small nucleus sizes. When analyses of experimental data using bulk properties show the nucleus size to be small, the resulting parameters should be regarded as largely empirical parameters describing the relative nucleation potency of the system. Considerable justification for the continued use of classical theory is provided by its general success in predicting nucleation behavior as a function of supersaturation and temperature. We emphasize that the qualitative features of the statistical mechanical models, particularly the critical super-saturation behavior that is central to the present work, are the same as those of the classical model. Of course, potential energy terms and surface partition functions replace the volume and surface energy terms of the latter model. The most recent versions of classical nucleation theory have been extensively reviewed.2,3,7 so that only the results are presented here. For homogeneous nucleation of a condensed phase from the vapor phase, the volume free energy change is ?Gv=vrT = =^ln£ [1] where v is the molecular volume of the condensing species. The supersaturation ratio,

Jan 1, 1970

By P. B. Bucky

IN previous publications1 it was shown that a scalar model of any weighty structure, where the stresses produced are mainly due to gravita-tional forces, will behave similarly to its prototype if the model and proto-type material are the same arid the model is placed in a centrifugal field of force (substituted for the gravitational field) which has been increased in the same proportion as the linear model scale is decreased. This principle is now being applied to a study of the behavior of mine structures in an attempt to determine the laws that control them. The importance of using models as a means of solving mining problems has been discussed previously, yet it may be well to reiterate here that this method of research and type of solution have been productive of far-reaching results in other engineering fields, notably those of hydraulics and aerodynamics. Model research is a comparatively cheap and time-saving method of checking practical experience. Up to the present time the best guide to what will happen underground has been the judgment of the man who has observed most and been longest acquainted with the local conditions. His experimental work underground is limited by the elements of cost, time and danger, while failure is seldom permissible. In a model, as in a test tube, all elements that affect the solution of a problem are under definite control. Each element may be varied independently at will and its effects noted. Experimental work on a small-scale model may there¬fore be performed at a comparatively small expenditure of time and money, instead of on a full scale in the mine, where the cost, time, and danger elements are necessarily great. The experiments described in this paper are to be considered as part of a progress report and a continuation of Technical Paper 425. One object of the research program is to make each experiment point the way for further research on that particular phase of the general problem. The results and conclusions presented here are therefore not to be considered as final, but are to be interpreted as

Jan 1, 1933

By D. C. Jillson

WORK in recent years1-' has indicated a complexity of the processes of deformation of metal crystals not previously appreciated and not fully accounted for by any hypothesis so far advanced. Furthermore, the nature and mechanism of formation of nuclei of recrystallization have not been determined precisely. The deformation of single crystals of zinc has been studied frequently, but the purity of the zinc and the perfection of the specimens sometimes have been given little consideration. Methods have been developed recently that readily yield zinc single-crystal specimens of high quality.5 In the present work, such specimens were deformed in various ways under various conditions, and deformed specimens were annealed to obtain information regarding recovery, recrystallization, and grain growth. The paper attempts to correlate and evaluate previous data, as well as to present new data, in order to determine areas in which more detailed work might be done most profitably. Tests at temperatures from the freezing point (419.46°C) to room temperature revealed glide only on basal planes in a close-packed direction (100)*, as reported by Mark, Polanyi, and Schmid6 and others. Markings probably similar to those observed by Kolesnikov7 and Boas and Schmid8 were noted in specimens stretched at elevated temperatures, but it seemed clear that these were not caused by prismatic or pyramidal slip (see second paragraph of section on Phenomena Involving Bending of the Basal Plane). Twinning Twinning on the octahedral plane of a face-centered cubic metal has been pictured as a process of simple homogeneous shearing along that plane in a [112] direction. It was recognized by Mathewson and Phillipsv and others""" that the (102) twinning of zinc required a somewhat more complex mechanism and might be considered as a homogeneous shearing of (102) planes in a [211] direction plus slight adjustments of atoms to positions of greater stability or lower energy, or as a single movement of each of the atoms in the same sense into the final positions. Gough and Cox" modified Mathewson's mechanism to obtain a more stable lattice configuration, but it is not clear that they succeeded, and their mechanism requires movement of some of the atoms in a sense opposite to that of the overall twinning movement. They also suggested that twinning may occur as a result of previous basal slip. This conclusion was based on the observation that twinning caused by alternating torsion was clearest and most profuse at positions for maximum basal slip rather than for maximum stress on (102) planes in the close-packed direction (not the twinning direction), and no mechanism was described. It might be wondered whether resolution of stresses on the twinning plane in the twinning direction would have afforded a simpler explanation. If twinning is essentially a simple homogeneous shearing along the twinning plane, it would seem that twins should grow by a smooth, continuous mechanism, and, indeed, that a simple reversal of stresses should reverse the shearing and de-twin the crystal. Cylindrical tablets 1/8 to 1/4 in. thick were cleaved from singlle-crystal specimens and were squeezed, perpendicular to a second-order prism plane, to give a tensile stress perpendicular to a first-order prism plane. A "click" was heard and a thin, needle-like twin appeared on the basal cleavage face. If squeezing was continued smoothly, the twin, viewed at magnifications up to X500, grew smoothly and quietly (fig. 1). X-ray examination verified that the twinning was of the (102) type. Rotating the compression axis 90" to reverse the stress then caused a smooth, continuous shrinkage and ultimate disappearance of the twin (fig. 2). The squeezing also caused a rumpling of the basal

Jan 1, 1951

By F. N. Rhines, A. H. Grobe

At elevated temperatures the oxide of silver is unstable in the air at atmospheric pressure, consequently no external oxide scale forms upon pure silver under conditions of high-temperature annealing. When small quantities of certain alloying elements are present in the silver, the formation of a thin external scale is possible1 and in addition there may form a subscale composed of the oxide of the solute element precipitated within the body of the silver. Norbury2 and Leroux and Raub3 have reported internal oxidation (subscaleiormation) in alloys of silver with 2, 7.5, and 30 per cent of copper. The presence of the subscale is believed to be responsible, at least in part, for the objectionable "fire mark" in Sterling silver.4 Several other alloys of silver, after oxidizing heat-treatments, are known to exhibit undesirable polishing characteristics that may be the result of internal oxidation. Except for the absence of an external scale of silver oxide, it is to be anticipated that silver alloys will prove to be very similar in their oxidation behavior to the alloys of copper, the oxidation characteristics of which have been studied in some detai1.5,6 The present research confirms this anticipation. The oxidation of a series of 20 dilute alloys of silver has been studied metallographically; some types of subscale not encountered among the copper alloys have been found. Instances of internal oxidation in alloys of most of the metals of the 1-6 and VIII groups of the periodic system are on record, but evidence of this type of oxidation in alloys of the metals of the intermediate groups is lacking. A number of the metals of the intermediate group, among them cadmium, lead, tin, and zinc, appear to provide the conditions essential to subscale formation; i.e., they form oxides with a relatively low negative free energy of formation, they dissolve other metals that form more stable oxides, and, presumably, oxygen will diffuse through them. In a study of 40 alloys of these white metals only a few cases of internal oxidation have been found. The probable reasons for this difference in behavior will be discussed presently. Experimental Procedure Silver.—The silver alloys employed in the oxidation studies were prepared in heats of 30 grams each from high-purity silver (99.993 per cent Ag)* and the purest avail-

Jan 1, 1942

By J. W. Freeman, E. E. Reynolds, A. E. White

Fram a study of 63 systematic alloy modifications it was found that molybdenum, tungsten, and columbium, added individually or simultaneously, and increases in chromium cause major improvements in 1200°F rupture strengths of Cr-Ni-Co-Fe base alloys. Rupture strengths were a function of the effect of composition modifications on both the inherent creep resistance and the amount of deformation the alloy would tolerate before fracture. THIS paper describes the results of an investigation of a series of alloys with systematic variations of the chemical composition of the following basic alloy: C, 0.15: Mn, 1.7; Si, 0.5; Cr, 20.0; NI, 20.0: Co, 20.0: Mo, 3.0; W, 2.0; Cb, 1.0; N, 0.12; Fe, 32.0 pct. The 62 modifications of this alloy were produced under conditions which minimized all factors influencing properties at high temperatures except composition. Melting, fabrication, and heat treatment were carefully maintained constant. Stress-rupture properties at 1200°F were used as the primary criteria of evaluation of the alloy. The objective of the study was to obtain data for determining the fundamental role of the influence of alloying elements on properties of heat-resistant alloys at high temperatures. In addition the results should be useful in determining optimum chemical compositions, the sensitivity of properties to variations in composition, and the degree to which alloy content could be reduced while retaining worthwhile properties. It is difficult or impossible to develop correlations between properties at high temperatures and systematic variations in chemical composition from published data for wrought heat-resistant alloys developed for gas turbines.' ' The main reason for this is the extreme dependence of the properties on conditions of treatment of the alloys." In most cases variation in final treatments between alloys so influences the properties that the influence of chemical composition is obscured. In addition it is recognized Table I. Basic Alloy and Some Modifications Used Basic AllOy, Variations in Element Pct Composition, Pct C 0.15 0.08. 0.40. 0.60 Mn 1.1 0.03,0.25.0.50,1.0,2.5 S1 0.50 1.2, 1.6 Cr 20.0 10, 30 Ni 20.0 0, 10,30 Co 20.0 0. 10, 30 MO 3.0 0. 1.2.3, 5, 7 W 2.0 0, 1, 5, 1 Cb 1.0 0.2,4,6 N 0.12 0.004, 0.08, 0.18 Fe 32.0 that variations exist between heats of the same alloy which are related to melting practice and that there is a strong possibility that conditions of hot working influence response to final treatments. The development of heat-resistant alloys has been based on the gradual accumulation of data roughly related to composition from extensive testing programs. There is every reason to believe that in most cases the optimum compositions have been achieved by this procedure in the alloys commercially available. There are, however, very little data showing the influence of systematic variations of composition free from the influence of other factors, particularly for alloys of the type investigated. Several investigators of cast alloys have demonstrated compositional effects, notably Grant,1-6 Epremian,t Guy,8 and Harder and Gow.9 Sykes10 eviewed the work on the wrought alloy Rex 78 and the systematic variations of carbon, copper, molybdenum, and cobalt leading to the development of the stronger Rex 337A alloy. From the papers by Wilson11 and Henry12 it is possible to deduce the beneficial effect of substituting cobalt for iron in 0.45 pct C-20 pct Cr-20 pct Ni-4 pct Mo-4 pct W- 4 pct Cb alloys. Wilson mentioned but did not present the extensive compositional studies involved in developing these alloys. Binder" showed optimum properties for 3, 2, and I pct, respectively, for molybdenum, tungsten, and columbium in 20 pct Cr-20 pct Ni-20 pct Co-30 pct Fe alloys for limited systematic variations of these

Jan 1, 1953

By M. Metzger, L. E. Hendrickson

Grain boundary attack in 16 pct HCl was found to be substantially the same at low penetrations in zone-refined aluminum (individual impurities 0.1 at. ppm), superior electrolytically refined aluminum (51 at. ppm), and aluminum with various impurities at much higher levels. It was concluded that impurity atom segregation affecting corrosion would have been detected and that the corrosion susceptibility did not originate in this segregation but in the structure of the boundary. It was pointed out that the significance of most previous studies in this system had been obscured by an unrecognized autocatalytic copper reaction. Although general corrosion rate was also impurity-insensitive , there was shallow pitting attributed to iron seg-regation and a hillocked surface texture associated with copper; these were interpreted as due to cathodic damage affecting- cathode distribution. THE grain boundary corrosion suffered by high-purity aluminum in hydrochloric acid has been the object of some interest (the earlier work is summarized in Ref. 1). The central metallurgical question here is whether the corrosion susceptibility of a boundary originates in its structure or in impurity atom segregation. An attempt to study this question revealed large catalytic effects associated with small quantities of the copper impurity in the aluminum, 220 ppm, or in the corrodent, and these magnified the preferential boundary attack and obscured the intrinsic susceptibility question.' After the catalytic effects had been examined,"' test conditions could be designed to avoid them and methods developed for studying the shallow boundary penetrations prevailing when they were absent.= It then became possible to determine whether intrinsic boundary corrosion in aluminum involves impurity segregation. The older work provided no firm information on grain boundary segregation of specific solutes influencing corrosion although several studies suggested iron segregates.4,5 perryman4 found in 10 pct HC1 that the slowly developing (microns per month) grain boundary grooves were deeper in material of higher iron content (range 10 to 550 ppm, with 5 to 80 pprn Cu) but he did not measure the depths of general corrosion, which were probably several times greater, and his reference surfaces may have varied more than did the groove depths. Metzger and Intrater's results for 20 pct HC~,' which yielded higher time-average rates (mm per month) and deeper penetrations, suggested that boundary segregation of iron (range 4 to 230 ppm, with 22 pprn Cu) decreased the penetration rate. However. in the stronger acid the autocatalytic l.E. HENDRICKSON, Student Member AIME ,and M. METZGER, Member AIME, are Research Assistant and Professor of Physical Metallurgy, respectively, Department of Mining, Metallurgy and Petroleum Engineering, University of Illinois, Urbana, Ill. Manuscript submitted March 11, 1968. IMD effect of the copper impurity is greater1,2 and it is now evident that their corrosion rates had been much magnified by this effect and did not provide a proper basis for the analysis of segregation. In exploratory studies of other solutes (made under the same conditions), 1000 ppm additions of Mg, Mn, or Si or 100 ppm Ca were without effect.1 Montariol6 noted that boundary attack in 22 pct HC1 persisted after zone refining although with fewer deep fissures (ranges 0.06 to 4 ppm Cu, 4 to 23 ppm Fe). Autocatalytic effects may have influenced these results also and those of Perryman.4 The present objective was, as a first step, to see whether quantitative tests designed to exclude autocatalytic influences would indicate the existence of low-level impurity effects on intrinsic boundary corrosion. Comparison of electrolytically refined with zone-refined aluminum of lower copper and iron contents revealed no differences in boundary corrosion, but certain impurity-sensitive differences in general corrosion morphology were noted and investigated further at higher impurity levels. I) EXPERIMENTAL PROCEDURE A) Material. A selection from commercially available material was made with the cooperation of several producers. An electrolytically refined lot (III-A, 1 ppm Cu, 2.4 ppm Fe) studied previously3 provided a starting point. Since material of substantially higher copper content could not be used if the autocatalytic corrosion reaction were to be avoided,2,3 a lot (III-B) of about the same purity was added as a check, two zone-refined lots (I-A and I-B) with copper and iron an order of magnitude lower were selected for comparison, and an intermediate lot (11) was included. Analytical data are given in Table I. For copper in I-B and iron in I-A and I-B, the actual concentration is thought to be near the limit given, i.e., about 0.1 ppm. For titanium, vanadium, and chromium in III-B, the actual amounts are thought to be, like those in III-A, substantially lower than in the zone-refined lots (these elements concentrate in the solid on freezing). Data are given later for some additional lots surveved. B) Specimen Preparation and Procedure. one by 3 cm blanks with a notched stem were cut from 1-mm cold-rolled sheet, annealed 24 hr in air at 650°C, and quenched in an air stream (42°C per sec initial cooling rate, 100 sec to cool to 100°C). The high annealing temperature maximized diffusivities and approach to equilibrium impurity distributions. A water quench, in principle more efficient in preserving the distribution established, was undesirable because the boundaries were almost plane and they tended to shear and migrate during the quench and thus to be separated from any existing impurity atmosphere. Test procedures, previously described,3 involved electropolish-ing, etching 2 min in 10 pct HF at 24.0°C, and expos-

Jan 1, 1969

By V. A. Phillips

Transmission-electron micrographs of electro-thinned samples of bulk-aged Cu-3.1 pet Co alloy show an aging sequence, supersaturated solid solution — coherent particles — quasi -coherent particles — noncoherent particles. Hardening is due to precipitation of coherent spherical fee coball-rich particles showing coherency strain fields, which are resolved at between 15 and 30A diameter. Loss of- full coherency did not occur until well into the overaged region, even with the assistance of deformation after aging. Different average particle diameters of 123, 92, and 149 ± 10Å were observed in samples aged to peak yield strength at 600°, 650°, and 700°C, respectively, indicating that there is no critical size for peak hardening. Noncoherent particles tended to develop (111) faces and became octahedral in shape. Dislocations tended to nucleate spherical coherent particles which eventually grew together forming large elongated particles. The surface energy of a noncoherent (low-angle) inter-phase boundary is estimated to he about 50 ergs per sq cm. A number of particle lining-up phenomena were observed. Overaging is principally attributed to increase in particle spacing, progressive loss of coherency, and increase in amount of discontinzdous precipitation. COPPER dissolves about 5.6 at. pet (5.2 wt pet) of cobalt at 1110oC1 and the solubility decreases to 0.75 at. petl (0.54 at. pet)2 at 650°C and to 0.1 at. pet or less at lower temperature.' It has been known for many years3-5 that Cu-Co alloys are capable of age hardening. Since cobalt is fee above 417°C and its atom size is only about 2 pet smaller than that of copper, precipitation of coherent particles would be expected. The equilibrium phase precipitated at 700°C and below contains about 10 pet Cu in solution which tends to stabilize the fee structure, lowering the transformation temperature to 340oc.l The alloy is known to undergo discontinuous precipitation in addition to general precipitation; while the former can be seen with an optical microscope, the latter precipitates are not visible except in the grosly overaged condition.5, 6 Extensive use has therefore been made of the ferromagnetic properties of the precipitate in order to follow the course of aging, and it has proved possible to measure the average particle size, spacing, approximate shape, and volume fraction and to determine that the particles are coherent without ever seeing a particle (see for example Refs. 2, 7, and 8). The magnetic measurements of particle size are limited to diameters below about 120Å.7 The present study was undertaken using the techniques of transmission-electron microscopy in order to check the above conclusions, to extend the previous magnetic work to larger particle sizes, and to attempt a more detailed correlation of properties and structure. A portion of this work has already been published.9-11 The present paper is concerned with the metallographic features of precipitation in relation to aging curves. Bonar and Kelly12'13 have published preliminary results of a similar study on single crystals of Cu-2 at. pet Co. EXPERIMENTAL Preparation of Alloy. A Cu-Co alloy, containing 3.12 wt pet (3.36 at. pet) Co by analysis, was prepared from 99.999 pet purity oxygen-free copper and electrolytic-grade cobalt. The alloy was melted and cast in vacuo in a high-frequency furnace using a graphite crucible and mold: Analysis showed chat 0.004 pet C was picked up during melting. The 1-1/2-lb ingot was homogenized in hydrogen for 24 hr at 1000°C. Slices were cold-rolled to 0.005 or 0.003 in. thickness, with an intermediate 650°C anneal in hydrogen at 0.080 in. thickness. Batches of six to ten strips were solution-treated in sealed-off quartz tubes in high vacuum in a vertical furnace and quenched by dropping into iced brine containing a device which snapped off the nose of the tube. Solution treatment consisted of 1 hr at 990°C or 2 hr at 965°C. The latter was employed for all mechanical-property studies, since a tendency was noted for the higher temperature to give porous material. Strips were usually aged individually in a horizontal vacuum furnace, inserting into the hot zone and withdrawing into a cold zone without breaking the vacuum. This method gave a rapid heating rate, permitting the use of short aging times. In some cases, particularly for the longer aging times at the higher temperatures, samples were sealed individually in quartz tubes in high

Jan 1, 1964

By A. H. Grobe, F. N. Rhines

At elevated temperatures the oxide of silver is unstable in the air at atmospheric pressure, consequently no external oxide scale forms upon pure silver under conditions of high-temperature annealing. When small quantities of certain alloying elements are present in the silver, the formation of a thin external scale is possible1 and in addition there may form a subscale composed of the oxide of the solute element precipitated within the body of the silver. Norbury2 and Leroux and Raub3 have reported internal oxidation (subscaleiormation) in alloys of silver with 2, 7.5, and 30 per cent of copper. The presence of the subscale is believed to be responsible, at least in part, for the objectionable "fire mark" in Sterling silver.4 Several other alloys of silver, after oxidizing heat-treatments, are known to exhibit undesirable polishing characteristics that may be the result of internal oxidation. Except for the absence of an external scale of silver oxide, it is to be anticipated that silver alloys will prove to be very similar in their oxidation behavior to the alloys of copper, the oxidation characteristics of which have been studied in some detai1.5,6 The present research confirms this anticipation. The oxidation of a series of 20 dilute alloys of silver has been studied metallographically; some types of subscale not encountered among the copper alloys have been found. Instances of internal oxidation in alloys of most of the metals of the 1-6 and VIII groups of the periodic system are on record, but evidence of this type of oxidation in alloys of the metals of the intermediate groups is lacking. A number of the metals of the intermediate group, among them cadmium, lead, tin, and zinc, appear to provide the conditions essential to subscale formation; i.e., they form oxides with a relatively low negative free energy of formation, they dissolve other metals that form more stable oxides, and, presumably, oxygen will diffuse through them. In a study of 40 alloys of these white metals only a few cases of internal oxidation have been found. The probable reasons for this difference in behavior will be discussed presently. Experimental Procedure Silver.—The silver alloys employed in the oxidation studies were prepared in heats of 30 grams each from high-purity silver (99.993 per cent Ag)* and the purest avail-

Jan 1, 1942

By M. R. Geer Olds, H. F. Yancey

THE increase in the number of coal-cleaning plants employing dense-medium processes occurring since 1946 is especially interesting when viewed historically. Both sand and magnetite were introduced as material for heavy mediums at about the same time, sand in the Chance process in 1921 and magnetite in the Conklin process in 1922, but from that point on their records diverge. The Chance process enjoyed a steady growth from its inception, whereas no additional magnetite plants were built in the United States for over 20 years. Then, following the close of the World War 11, magnetite was again introduced, this time with marked success. During the following years some 47 plants employing magnetite medium were built. This rapid growth of dense-medium cleaning has been concurrent with widespread adoption of full-seam mining on one hand and a return to a more competitive market on the other. At a time when changing mining practice has provided cleaning plants with dirtier coal and changing market conditions have simultaneously demanded a cleaner product, the industry has through necessity turned to improved preparation. The inherently greater sharpness with which dense-medium processes can separate coal from impurity is thus helping to hold the line against ever-increasing mining costs and at the same time assisting materially in retaining badly needed markets. Although dense-medium cleaning unquestionably offers a distinct advantage when the washing problem is difficult, other methods can provide almost equally high efficiency when the coal is easy to wash. Moreover, fine coal cannot yet be treated by heavy medium in a proved process, although the Driessen cyclone is in the pilot-plant stage. Most of the present types of dense-medium equipment have been in use only a few years, and the dearth of information in the literature concerning their performance characteristics is entirely understandable, Nevertheless this information is necessary if the process is to be intelligently applied to individual cleaning problems. Without data on the efficiency of a process in a particular type of separation, it is difficult to assess the advantage to be expected from it. Similarly, the role of particle size in heavy-medium separation is important in some cases, yet there is little published information on this aspect. To mention only one more of the numerous points on which essential information is lacking, the bearing of medium characteristics on performance has been discussed only in qualitative terms. It was with the hope of providing information on some of these points that the Bureau of Mines built a dense-medium pilot plant for cleaning coal at its Northwest Experiment Station in Seattle in 1950. The plant has been operated continuously since that time, and over 50 runs have been made on 7 coals exhibiting a wide range of washability characteristics. An idea of the magnitude of this work will be gained from the fact that examination of the plant products has involved some 600 float-and-sink separations and about 2500 ash determinations. A laboratory pilot plant is especially well adapted to investigate many aspects of performance because close control over test conditions can be exercised and because a large number of tests can be made rapidly. On the other hand, factors such as consumption of medium and other cost items can be investigated satisfactorily only in a commercial plant. Actually, the two forms of investigation should be complementary, with the laboratory work pointing the way for confirming tests in commercial units. Pilot Plant The dense-medium pilot plant employed for this work comprises a 24x30-in. drum-type separating vessel, a 12-in. densifier, a 12-in. magnetic separator, a 26-in. x 9-ft vibrating screen, and the necessary pumps and conveyors for handling materials. Arrangement of these units corresponds with the flowsheet used in most commercial plants, except that a thickener is not provided for the feed to the magnetic separator. Coal and refuse discharge from the separating drum to the vibrator, which is divided longitudinally down the center. Medium draining through the first 3 ft of the screen is recirculated directly to the drum. Sprays on the middle 3 ft of the screen rinse medium from the products, and the last 3-ft section is for dewatering. Dilute medium from the rinsing and dewatering sections is pumped to the magnetic separator, where magnetic solids are recovered. These are pumped to the densifier, from which they return to the medium-drainage sump by gravity through a demagnetizing coil. The drum-type separating vessel is a scale model of a commercial unit. Feed enters axially at one end of the drum just below the surface of the bath, and float material overflows through a circular opening at the other end. Particles sinking to the bottom of the bath are picked up by lifting flights bolted to the inner wall of the drum, elevated out of the bath, and sluiced to the vibrating screen. Baffles suspended in the bath prevent float material from entering the sink-lifting flights. About 8 gpm of medium is used to sluice the feed into the drum. An additional 15 to 24 gpm, depending upon operating conditions, is added through two pipes dipping into the bath behind the baffle on the side where the sink-lifting flights enter the bath. The bath available for separation is 2 ft long and 13 in. wide, giving an area of 2.08 sq ft. Depth of bath from the surface to the top of the sink-lifting flights, measured vertically below the axis, is 6 in.

Jan 1, 1954

By M. Khorouzan, L. Thomas

Designers of semiconductor devices have been strivi,ng to resolve problems associated with Au-A1 alloys in bonded in.tercomzeclions. One approach now being- used is that of waintaining a physical seyav-atioz between the two metals in bond areas. This is accolrzplished by alunzincnz-plating a bonding area on the tips oJ the kovar leads and using alcminurn wires to join the senzicondictor device to the leads. The portion of the kovar lead which is on the externul side of the sealed package is gold-plated to provide an oxide-free surface for soldering or welding. A discoloration condition originally thought to be sinilar to purple plague, occuving in the yluled uluninur bonding area after package sealing, has been investigated to determine its efiects ipm bond integrity. Electron-micro-probe analysis determined that no1 only gold, but lead, zinc, and silicon were also present in the discolored area. A series of samples conlaining' conkrolled umonts of these inzpitrities weve prepared and subjected to a sil.zuluted sealing process. The investigations swcued that, of the contawiinants, only zinc toas detrinenlul to Lhe bond integily. The discoloration condition itself was found not to be detrimental to the bond integrity. DESIGNERS of semiconductor devices have been striving to resolve problems associated with Au-A1 alloys in bonded interconnections. One approach now being used is that of maintaining a physical separation between the two metals in bond areas. This is accomplished by aluminum plating a bonding area on the tips of the kovar leads and using aluminum wires to join the semiconductor device to the kovar leads. The portion of the kovar lead which is on the external side of the sealed package is gold-plated to provide an oxide-free surface for soldering or welding. Contamination as evidenced by discoloration of the aluminum-plated area was observed in a number of integrated circuits undergoing examination for defect characteristics which cause electrical failures.' This paper contains the results of an investigation to determine the nature of this discoloration, its cause, and its effect upon the integrity of the interconnection bond. I) THE NATURE AND EXTENT OF ALUMINUM-BOND CONTAMINATION The initial hypothesis in the investigation was that the discoloration was caused by reaction of the aluminum film with some unknown contaminants during the sealing of the hermetically sealed integrated-circuit flat package. The package is a rectangular ceramic container sealed with glass which surrounds the kovar leads as well as joining the top to the bottom. The seal is made hermetic by heating and cooling the package to devitrify the glass. In the case of the packages under investigation, the hermetic sealing had been accomplished with dry air as internal atmosphere. The apparent effect of contaminations as observed by microscopic examination was the formation of surface oxides having variations in color encompassing the whole spectrum of visible light. The contamination appeared to be related to one of the more notorious examples of these colorations, the so called purple plague.' In addition to purple plague, Fig. 1 shows the tarnish in the luster of the aluminized surface in the bond area which had been observed in many of the integrated circuits. To identify the contaminant in the bond area electron-probe microanalysis techniques were used.3 Fig. 2 shows the result of this analysis. The contaminants identified were gold, aluminum, zinc, lead, silicon, and cobalt. Fig. 2(a) is a back-scatter display of the area under study. The back-scattered electrons provide a general indication of the distribution of elements in the specimen surface. Elements with higher atomic number scatter more electrons back from the surface and are seen as light areas in the picture. The sample current, Fig. 2(b), is the amount of current conducted by the specimen as a result of electron-beam striking it and is an indication of element distribution. The Sample current is the reverse of back-scatter and complements it. Other pictures in Fig. 2 are produced by characteristic X-rays generated by the elements, allowing the isolation of the element of interest. The isolated element appears white and all other elements are dark. In this manner a comparative study provides a correlation between different surface areas and the elements which are in these areas. The area covered by the gold film, Fig. 2(c), shows that the boundary between the gold film and the kovar is not sharp as expected and that some sort of diffusion has taken place. Fig. 2(c) shows that some gold particles have been carried to the bond area and are in the proximity of the bonded wire in spite of the presence of a physical barrier in the form of the un-

Jan 1, 1967

By P. Chiotti, J. T. Mason

Ther?nal, X-ray, metallographic, and vapor pressure data were obtained to establish the phase diagram and standard free energy, enthalpy, and entropy of formation for the compounds in the Sw-Zn system. Four compounds, SmZn, SmZn2 , SmZn4.s, and SmZn8.5, melt congruently at 960°, 94Z°, 908°, and 940°C, respectively. The cornpounds SlnZns, Sm3Znll, and SnzZn7.3 undergo peritectic decomposition at 855", 870°, and 890C, respectively. Another compound of uncertain stoichiometry, SmZn11, undergoes peritectic decomposition at 760°C. Four entectics were observed with the following compositions in weight percent zinc and eutectic tenzperatures in degrees Centigrade: 12 pct, 680°C; 36 pct, 890°C; 58 pct, 850°C; and 72 pct, 900°C. An allotropic transformation and a composition range were observed for the SmZnz compound. The transfor)nation varies from 905" to 865°C as the zinc content increases from 16.0 to 48.5 wt pct, respectively. The free energy of formation of the compounds at 50PC varies between -15.9 kcal per mole for SmZn to -51.1 kcal per mole for SmZn,.,. Corresponding enthalpies vary between -19.2 to -78.3 kcal per mole. The ther-modynamic properties for the liquid alloys are described by the relations: A search of the literature revealed very little information on the Sm-Zn system. Chao et al.&apos; as well as Iandelli and palenzonai have reported the structure of SmZn to be cubic B2 type and Kuz&apos;ma et al3. have reported the structure of -sm2zn17 to be of the Th2Ni17 type. The purpose of this work was to establish the phase diagram of this system, to determine the zinc vapor pressure over the solid two-phase regions of the SYstem, and to calculate the thermodynamic properties of the compounds. MATERIALS AND EXPERIMENTAL PROCEDURES The metals used in this investigation were Bunker Hill slab zinc 99.99 wt pct pure and Ames Laboratory samarium. Analysis of the samarium by chemical, spectrographic, and vacuum-fusion methods gave the following average impurities in ppm: Nd, <200; Eu, <100; Gd, <100; Y, <50;Ca, 225; Ta, 400; Mg, 10; Cu, ~50; 0, 175; H, 20; and N, 15. The elements Fe, Si, Cr, Ni, Al, and W were not detected. The samarium was received as sponge metal and was kept under argon except when being cut with shears and when being weighed. Tantalum was found to be a suitable container for alloys with zinc contents up to the Sm2Znl, stoichio-metry. At higher zinc contents the grain boundaries of the tantalum containers were penetrated by the alloy and the containers failed during prolonged annealing. About 25 g of massive zinc and samarium sponge were sealed in tantalum crucibles equipped with thermocouple wells. These crucibles were in turn sealed in stainless-steel jackets. All closures were made by arc welding under an argon atmosphere. The samples were equilibrated in an oscillating furnace and in some cases were given various heat treatments in a soaking furnace. After appropriate heat treatment the steel jackets were removed and the alloy subjected to differential thermal analysis. The apparatus was calibrated against pure zinc and pure copper and found to reproduce the accepted melting points within 1°C. Alloys were subsequently subjected to metallographic examination and those of appropriate compositions were used for X-ray diffraction analysis and for zinc vapor pressure determinations. The vapor pressures were determined by the dewpoint method. Both the differential analysis and dewpoint measuring apparatuses have been described in earlier papers.4, 5 All alloy samples were etched with Nital (0.5 to 3 pct nitric acid in alcohol) except the samarium-rich alloys. These more reactive alloys were electro-polished in a 1 to 6 pct HClO4 in methanol solution at -700c at a potential of 50 v. EXPERIMENTAL RESULTS Phase Diagram. The results of thermal analysis are indicated by the points on the phase diagram, Fig. 1. Eight compounds and four eutectics were observed. The composition of the compounds and their melting or peritectic temperatures are given on the phase diagram. The four eutectic compositions in wt pct zinc and eutectic temperatures in % are: 12 pct,- 680°C; 36 pct, 890°C; 58 pct, 850°C; and 72 pct, 900°C. The stoichiometry of the most zinc-rich compound is still uncertain, but is very likely either SmZnll or SmZnlz. However, to simplify the presentation which follows it will be referred to as SmZnll. As shown on the phase diagram the phase regions for some of the samarium-rich alloys have not been unambiguously established. A sample of pure samarium was observed to transform at 924°C and to melt at 1074"C, in good agreement with corresponding val-

Jan 1, 1968

By M. E. Wadsworth, K. C. Bowles, H. E. Flanders, R. M. Nadkarni, C. E. Jelden

The kinetics of precipitation of copper on iron of various purity were carried out under controlled conditions. The rate of reduction has been correlated with such parameters as copper and hydrogen ion concentration, geometric factors, flow rate, and temperature. The character of the precipitated copper as a function of flow conditions and rate of PreciPitation has been observed under a variety of conditions. ThE precipitation of copper in solution by cementation on a more electropositive metal has been known for many years. Basile valentine&apos; who wrote Currus Triumphalis Antimonii about 1500, refers to this method for extraction of copper. Paracelsus the Great2 who was born about 1493 cites the use of iron to prepare Venus (copper) by the "rustics of Hungary" in the "Book Concerning the Tincture of the Philosophers". Agricola3 in his work on minerals (1546) tells of a peculiar water which is drawn from a shaft near Schmölnitz in Hungary, that erodes iron and turns it into copper. In 1670, a concession is recorded4 as having been granted for the recovery of copper from the mine waters at Rio Tinto in Spain, presumably by precipitation with iron. Much has been published in recent literature on the recovery of copper by cementation, the majority of the articles being on plant practice.5-24 The rest include articles on investigation of the variables involved25-28 and a review of hydrometallurgical copper extraction methods." This literature has established: a) The three principal reactions in the cementation of copper are Cu + Fe — Fe+4 +Cu [ 11 One pound of copper is precipitated by 0.88 lb of iron stoichiometrically. In actual practice about 1.5 to 2.5 lb of iron are consumed. 2Fe+3 + Fe — 3Fe+2 [21 Fe +2H&apos;-Fe+2 + H2 [3] Reactions [2] and [3] are responsible for the consumption of excess iron. Wartman and Roberson&apos;28 have established that Reactions [ I] and [2] are concurrent and much faster than Reaction [3]. b) Acidity control is important in the control of hydrolysis and the excessive consumption of iron. he commercial workable range is approximately from pH = 1.8 to 3." c) Iron consumption is closely related to the amount of ferric iron in solution. Jacobi" reports that, by leaving the pregnant mine waters in contact wi th lump pyrrhotite (Fe7S8) for 3 hr, all the iron was reduced to the bivalent condition and scrap iron consumption was cut to 1.25 lb scrap per pound of copper precipitated. He also reported that SO2 has been used successfully to reduce ferric iron to the ferrous state. d) The ideal precipitant is one that offers a large exposed area and is relatively free of rust. e) High velocities and agitation show a beneficial effect upon the rate of precipitation, as it tends to displace the layer of barren solution adjacent to the iron and also dislodges hydrogen bubbles and precipitated copper to expose new surfaces. Little work, however, has been published on the reaction kinetics of copper precipitation on iron. Cent-nerszwer and Heller20 investigated the precipitation of metallic cations in solutions on zinc plates. They found the cementation reaction to be a first-order reaction. The rate constant was independent of stirring for high stirring rates and they concluded that the rate is governed by a diffusional process at low stirring speeds and by a "chemical" process at higher stirring speeds where the rate reaches a constant value. This conclusion has been challenged by King and Burger30 who could not find any region where the rate was independent of the stirring speed, although the rate constant they had obtained for high stirring speed was greater than the maximum value of the rate constant reported by Centnerszwer and Heller (by a factor of six). King and Burger, therefore, concluded that the rate of displacement of copper was controlled only by diffusion. Cementation of various cations on zinc has been summarized by Engfelder.31 APPARATUS A three-necked distillation flask of 2 000-mm capacity was used as a reaction vessel. A pipet of 10-mm capacity was introduced through one of- the side necks, the sample of sheet iron, mounted in a rigid sample holder, through the other, the stirrer being in the middle as shown in Fig. 1. The whole assembly was immersed in a constant-temperature bath. The stirrer was always placed at the same depth in the solution. EXPERIMENTAL PROCEDURE Reagent-grade cupric sulfate (J. T. Baker Chemical Co., N.J.) was used to make up a stock solution containing 10 g of copper per liter which was then diluted to various concentrations as required. Experimental data were obtained by measuring the amount of copper and iron ions in solution at successive time intervals. The initial volume of the solution was always 2000 ml, 10-ml aliquots being removed each time for chemical analysis. Because the total volume change of the solution was less than 10 pct, no correction was used for solution volume change. Nitrogen was bubbled through the solution before and

Jan 1, 1968

By A. E. Nehrenberg

THE mechanism by which austenite forms in steels has received a great deal of attention in the literature in past years.&apos;-&apos;* Our present knowledge concerning this mechanism has been recently summarized quite concisely by Bain and Vilella,1 while a few years ago the literature was carefully reviewed by Roberts and Mehl.² The consensus is that any ferrite-carbide interface is a potential site for the nucleation of austenite during heating above the Acl temperature, and that the new austenite generally grows freely to produce approximately equiaxed grains, whether the carbides are initially present in the lamellar or the spheroidal form. In the case of eutectoid steels, growth of the new grains of austenite continues until contact is established with other grains. Then growth stops and an initial austenite grain size is established which does not change until the heating is continued to some high temperature at which grain coarsening begins. In the case of pearlitic steels which are not of eutectoid composition, the proeutectoid ferrite or carbide may interfere with the growth of the austenite if the temperature is not above that designated the A63 or the Acm, respectively. Although a large amount of work has been done to establish the mechanism of austenite formation in steels, it became clear to the present author while he was studying the transformation characteristics of a new 0.25 C Mn-Si-Ni-Mo hypoeutectoid steel" that the manner in which austenite grows in steels depends upon some factor, or factors, not previously considered. This was indicated by the fact that when this steel in the spheroidized condition was heated above the Ae1 temperature the new austenite which was formed did not envelop the carbides and grow in an equiaxed manner as described by Bain³ or spheroidized steels. Instead, in this steel, the austenite was observed to grow much more readily in certain directions than in others with the result that at temperatures within the Ac1-Ac³ ransformation range the austenite grains were acicular in shape. The excess ferrite was also found to be acicular with the distribution of these phases being such that a lamellar pattern was developed. This unusual directional growth of austenite in this new steel initially in the spheroidized condition is illustrated by fig. 1. A search of the literature revealed that this type of growth was not necessarily peculiar to this steel for similar microstructures had been observed by other investigators.4-8 However, the full significance of these microstructures does not appear to have been appreciated, and no work has been done to determine the conditions responsible for this directional growth of austenite or to arrive at an understanding of it. It was for this purpose that the work described in the present paper was carried out. Material: During the course of this investigation a total of 15 steels was studied. They consisted of hypoeutectoid, eutectoid and hypereutectoid carbon steels, and hypoeutectoid and hypereutectoid alloy steels, all of which were obtained in the annealed condition from commercial warehouse stock. As received, the carbon and alloy hypereutectoid steels had microstructures which consisted of spheroidal carbides in ferrite, whereas the eutectoid steel and the hypoeutectoid steels were pearlitic. The grades of steel represented were 1050, 1080, 10110, 3310, 4140, 4340, 4615, 6145, 8620, 9260, 9442,

Jan 1, 1951

By P. K. Strangway, H. U. Ross

Briquets consisting of pure artificial magnetite, pure artificial hematite, and mixtures of the two were reduced by hydrogen in a loss-in-weight furnace at temperatures in the range 500° to 1000° . The rate of reduction of the pure hematite briquets increased continuously with increased temperature. In contrast, the pure nmgnetite briquets exhibited a pronounced rate ninimutn at about 700°C. Metallographic studies of partially reduced briquets rerlealed that, at this temperature, the he.matite samples reduced in a topo-chemical manner while the magnetite ones reduced uniformly throughout, and after partial reduction their cross sections contained a mixture of iron and unreacted wustite grains. No iron shells could be detected on the surfices of any of these uwstite grains. X-ray diffraction investigations indicated that these grains had a rzinimum lattice parameter when they had been formed at the rate rninimum temperature. Also, it was found that an activation energy of 41,000 cal per mole zoas required for reduction when only these wustite grains were present. Thus, it is suggested that the overall reduction rate of the rnagnetile su?nples at temperatures in the range influenced by the rate nzinirnum phenomenon was limited by the rate qf iron ion diffusion in the unreacted wustite grains. THE rate minimum phenomenon, which has often been observed when reducing iron oxides at a temperature of about 700°C, is one of the most interesting, yet unresolved, problems in the field of reduction kinetics. Basic principles of chemical kinetics and &apos;In some instance, a second rate minimum has been observed at about 900°C. Since most investigators are in agreement that this minimum is directly related to the transformation from a to y iron (which takes place at 911°C) and since it was not encountered during the present reduction tests, it will not be referred to in this vaver. fundamental laws of diffusion all agree that, as the temperature is increased, the rate of reduction should also increase. However, with certain ores, it has been found that their reduction rate actually decreases with an increase in temperature up to some value X where a minimum reduction rate is reached. With further temperature increases beyond X the rate becomes more rapid again. Temperature X is usually referred to as the "rate minimum temperature", while the overall type of behavior constitutes the "rate minimum phenomenon". This phenomenon has been reported by numerous investigators. They have found rate minima during the reduction of both artifiial&apos; and natural374 magnetites and artificia15j6 and natural5" hematites. Rate minima have been observed when reducing high-purity material2 or low-grade ores,3&apos;4 when studying particles in the micronsize range5 or relatively large agglomerates,g10 and during reduction with either hydrogen7 or carbon monoxide.11"2 Previously, this phenomenon has been attributed to many factors; these include sintering and recrystallization of the iron formed during reduction374 changes in microporosity of the ore upon redction,"" formation of dense iron shells around retained wustite grains,11716 and chem-isorption,17 to name only a few. However, most investigators who have reported a rate minimum merely speculated as to what seemed to influence it and they did not examine the fundamental causes. Consequently, the present experimental study was initiated in order to evaluate the basic factors which could be associated with this phenomenon. MATERIALS AND METHODS The experimental techniques, followed during this investigation, are similar to those which have been described previously.18 The chemically pure magnetic powder was prepared by partially reducing Fisher reagent-grade hematite with a gaseous mixture of carbon monoxide and carbon dioxide in a rotating-drum furnace. Three-quarter-inch diam cylindrical briquets which weighed about 12 g were formed from this magnetite powder and pure hematite powder. All of the briquets were sintered while they were slowly raised through the 1200°C hot zone of a vertical tube furnace. An argon stream was continually flushed through this furnace in order to prevent oxidation of the magnetite briquets, while in the case of the pure hematite briquets sintering was carried out in air. The sintered hematite briquets had a density of 5.06 g per cu cm while the density of the sintered magnetite briquets was 4.27 g per cu cm. The sintered briquets were reduced by purified hydrogen in a loss-in-weight furnace at temperatures in the range 500" to 1000°C. In all instances, the critical reducing gas velocity was exceeded and, in order to ensure that the results were reproducible, duplicate briquets of each type were reduced under each set of experimental conditions. A continuous record of the weight loss during reduction was obtained with the aid of a Statham transducer. The present experimental setup was capable of detecting a change in weight as small as 10 mg. Since a weight loss of over 2 g usually occurred during each reduction test, an accuracy of better than 0.5 pct of the total weight loss could be achieved. RESULTS AND DISCUSSION Reducibility Tests. In the first set of experiments, pure hematite and pure magnetite briquets were used.

Jan 1, 1969

By Klaus Schwerdtfeger

The rules of solution of oxygen from H2O-H2-He gas and of nitrogen from N2-H2 gas in shallow melts of liquid iron were measured at 1610o and 1600o C, respectiuely. Concentration profiles were detemined in the liquid iron. Tire rate data indicate that the solution process is controlled by diffusion in the iron melt. The diffusivities for oxygen and nitrogen in liquid iron, as calculated from the present data, are DFe-o = (12 ± 3) < 10-5 sq cm per sec and DFe-N = 11 ± 2) X 10-5 sq cm per sec at the temperatures employed. AN attempt was made by Shurygin and Kryukl to measure the diffusivity of oxygen in liquid iron. In their experiments a silica disc was rotated in liquid iron containing oxygen, and the rate of formation of liquid iron silicate was measured. Assuming that the rate of dissolution of silica is controlled by diffusion of oxygen in the iron, the oxygen diffusivity was computed from the rate data giving Dfe-0 = 6.1 X 5 sq cm per sec at 1600°C. Although this value seems to be of the right order of magnitude, there is no proof of the correctness of the assumptions involved in the interpretation of these rate data. The oxygen concentration in the iron at the iron-iron silicate interface was taken to be that in equilibrium with the silica-saturated silicate melt. That is, it was assumed that no concentration gradient existed in the liquid silicate. This is a questionable assumption, unless it is proved that the thickness of the silicate layer is very much smaller than that of the diffusion boundary layer in the iron. Furthermore, Shurygin et al.1 used the Levich equation2 to interpret their rate data. This equation was derived for mass transfer between a solid disc and a single-phase liquid. The hydrodynamic and diffusion boundary layers in the iron stirred by a disc, via coupling of the silicate melt, may be appreciably different from those predicted by Levich&apos;s derivations. In the present work the diffusivities of oxygen and nitrogen in liquid iron were measured at 1610" and 1600oC, respectively. EXPERIMENTAL METHOD Iron melts contained in high-purity gas-tight alumina crucibles were reacted with H2O-H2-He gas for the determination of the oxygen diffusivity and with N2-H2 gas for the determination of nitrogen diffusivity. At the end of the reaction period, the samples were quenched in a cold H2-He gas stream at the top of the furnace. Oxygen or nitrogen contents in the iron were determined by chemical analysis. Two different types of diffusion experiments were perforxed. To determine concentration profiles, a few rate measurements were made using 4-cm-deep melts. The solidified samples were sliced into discs and each disc was analyzed for oxygen or nitrogen. In another series of experiments, oxygen or nitrogen was diffused into shallow melts (about 0.5 to 1 cm in depth) and the total sample was analyzed to obtain an average concentration of the diffusate. In most experiments, 4- to 5-mm-ID alumina crucibles were used. Some experiments were also made in smaller (3 mm) and larger (7 mm) diam crucibles. This variation in diameter caused no difference in the reaction rate, within the limits of experimental uncertainty. To promote the establishment of a stable density profile in the melt, all the samples were suspended in the lower end of the hot zone so that the top of the melt was hotter by a few degrees. Molybdenum wire resistance heating was used. The reaction tube of the furnace was a gas-tight recrystal-lized alumina tube. In most experiments the furnace was heated by an ac power supply. To check the possibility of inductive stirring, some experiments were carried out in a dc operated furnace, with essentially the same results. The temperature of the furnace was controlled automatically in the usual manner. The temperature was measured with a Pt/Pt-10 pet Rh thermocouple and is estimated to be accurate within ±5°C. The iron used was prepared by melting and vacuum-carbon deoxidizing electrolytic "Plastiron" in a zir-conia crucible. The main impurities are: Si 0.004 pct P, S <0.002 pct Cr 0.005 pct N 0.001 pct Zr 0.002 pct O 0.003 pct Mn 0.004 pct C 0.002 pct The gas composition was controlled by constant pressure head capillary flowmeters. Oxygen was removed from the gas mixture by passing it through columns of platinized asbestos (450°C) and anhydrone. Selected H2O contents were obtained by passing the purified gas through oxalic acid dihydrate-anhydrous oxalic acid mixtures held at constant temperature in a water bath. Water vapor pressure data for the oxalic acid dihydrate-anhydrous oxalic acid equilibrium were taken from the 1iterature.3 The flow rate used was about 1.5 liters per min. The whole system was checked for tightness at regular intervals.

Jan 1, 1968

By A. W. McReynolds

One characteristic of plastic deformation which distinguishes it from elastic strain is the essential inhomo-geneity of plastic strains. Elastic strain varies continuously through a material, and average relative displacements of initially adjacent atoms are only small fractions of their initial spacing, (strains of the order of 0.01 or less). On the other hand, plastic flow corresponds to the appearance of discontinuities in strain of the lattice, such as dislocations or slip bands, where local strain, on an atomic scale, is several orders of magnitude higher. These discontinuities are visible on a microscopic scale as the familiar slip lines (Fig 1). In spite of this obvious microscopic inhomogeneity, however, macroscopic measurements almost invariably show a smooth curve of stress vs. strain (Fig 2b) even if measurements of linear strains be made to an accuracy of one part in 107. This macroscopic homogeneity of strain indicates that the discontinuities in strain on slip planes occur in increments too small or too slow to be recorded individually, and further that they occur sufficiently independent of one another so that the small increments add at random to a smooth stress-strain curve. The present paper describes observations of plastic strain in aluminum of commercial purity and in high purity Al-Cu alloys, where there exists a strong coupling between slip in various regions of the specimen such that once initiated it spreads rapidly through a large volume. The total effect is that of relatively large, rapid, and regularly spaced steps of strain followed by periods of only elastic strain. Fig 2a illustrates the type of "stair-step" stress-strain curve which results. The properties of this cooperative slip phenomenon will be described further in the section on results: in par- ticular it will be shown that each step corresponds to the propagation of a wave of plastic deformation through the specimen. Some interpretations of the mechanism by which it occurs will be made in the following section. Although the type of plastic wave phenomena to be described has not previously been reported, there are numerous cases of related effects in the plastic yielding of metals: YIELD POINT PHENOMENA The most familiar of such effects is the "yield point" observed in low carbon steels, brass, duralurninum, and the like. It consists in the sudden termination of the elastic portion of the stress-strain curve by a large plastic strain. Since the usual tensile machine is such that yielding of the specimen relieves the load, the resulting curve is as shown in Fig 3. As the strain continues, deformation occurs at a lower stress for some time, then follows a rising curve, but with no further sudden yielding. This effect has been observed in brass by Sachs and Shojil and later by many others. Edwards, Phillips and Jones2 made extensive studies of the effect in steel, and of the role of various alloying elements. Although there seems to be fairly clear evidence that the yield point is caused by a hardening of the material by precipitation of impurities, no satisfactory explanation for the sudden yielding has been given. Winlock and Leiter3 have shown that the strain. level of the upper yield point depends strongly on the rate of loading, the yield point increasing by almost a fac- tor of two as the strain rate goes from 0.002 in. per in. per min. to 4.4 in. per in. per min. This effect would seem to imply an incubation period before yielding is initiated at a certain stress. On the other hand, by going to very slow loading rates, Edwards, Phillips and Jones2 showed that the yield point does not become lower and eventually disappear as might be expected, but, on the contrary, begins to rise at loading rates below about 25 Ib per in. per min. becoming much higher than at rapid loading rates. STRAIN AGING If, instead of continuing straining of a specimen after occurrence of a yield point, the load is removed and the specimen aged, resumption of the test results in occurrence of another yield point as shown by the dotted curve of Fig 3. The new yield stress is generally higher than the previous maximum applied stress. This hardening of the material by straining and subsequent aging is undoubtedly related to quench age-hardening resulting from the aging of a specimen quenched from high temperature. Since neither effect is observed in pure metals, it is generally accepted that quench-aging in all cases is the result of hardening by precipitation of a supersaturated alloying element, and that strain-aging is probably a similar precipitation, accelerated by disruptions of the lattice by previous strain. Pfeil4 has shown that strain-aging does not occur in iron from which all of the carbon has been removed, but that only a very small carbon content, around 0.003 pct, is necessary to cause strain-aging. In accord with this observation is recent work by Dijkstra5 in this laboratory showing that the solubility limit of carbon in iron is extremely low, less than 0.001 pct at 400°C. Edwards, Phillips and Jones2 have shown that the strain-aging effect is also removed by the addition of small quantities of elements such as Mo, Mn, Ti, and the like, which readily form carbides. Their results demonstrate the

Jan 1, 1950

By Donald E. Meyer

High current-low voltage EB-gun evaporation in an oil-free ultra-high vacuum system was found to be necessary, though not sufficient, for stability (300°C, 106 v per on) of aluminium gate MOSFET&apos;s and MOS capacitors not stabilized by a phosphorous glaze. five characteristics of the equipment used: 1) Vacuum purification of the aluminum charge, 2) Ionization of the evaporant by the electron beam, 3) X-ray formation, 4) Residual gases during evaporation, and 5) Metal film structure were studied as Possibly significant in MOS fabrication. EVAPORATION of contact metals common to the semiconductor industry historically has been accomplished with oil diffusion pump systems and various resistance heated evaporant sources as dictated by the type of metal evaporated. To meet a need for greater reliability of semiconductor devices, other metallization methods were developed. A good example would be application of the moly-gold contact system to integrated circuits with deposition by RF or triode sputtering.&apos; More recently, fabrication of stable metal-oxide-silicon devices and circuits has put new demands on metallization. The purity of the thin metal films composing MOS structures is critical, particularly at the metal-oxide interface, and ultra-high vacuum metallization using sputter-ion pumping and electron beam gun (EB-gun) evaporation are well suited for the task. At this laboratory aluminum has been the most common contact-gate metal for both MOS capacitors and MOSFET&apos;s. In the earliest work with MOS capacitors, aluminum was evaporated from wetted tungsten filaments using both diffusion pump and ion pump vacuum systems. In spite of clean oxide techniques these capacitors were unstable under bias-tempera-ture stressing. Only after a switch to EB evaporation of aluminum were stable capacitors produced. Using the same techniques it was possible to make MOSFET&apos;s with equivalent stability. Stability data for a discrete MOSFET is shown in Fig. 1. This is a "clean" oxide gate (no phosphorus stabilization or no etch back of a thicker gate) having a thickness of lOOO? thermally grown on the (111) plane. Gate length after diffusion was 0.24 mils, and the devices were hermetically sealed. Stressing conditions were 300°C and 106 v per cm applied alternately as a positive and negative field for 10 min, 50 min, and 4 hr for a total stress time of 10 hr. An initial shift in turn-on voltage of 0.1 v was detected for 10 min of positive bias. All evidence at this laboratory indicated that while EB-gun evaporation of ultra-high purity aluminum was not sufficient for 300°C stability, it did seem to be necessary. There may well then be something inherent in the EB-gun deposition used which enhanced stability, and probably no single factor existed but rather a series of factors. It is the purpose of this paper to report on some of the investigations carried out to learn more about EB-gun evaporation in ultra-high vacuum systems. EXPERIMENTAL DESCRIPTION The EB-gun was self accelerated, had a maximum power rating of 10 kw, and used a water-cooled copper crucible able to hold a 20-g aluminum charge. The electron beam was bent 180 deg and focused by an electromagnet which also provided movement of the beam across the crucible. Normal power conditions in this work were 9 kv and 300 to 600 mamp. The gun can be described as high-cur rent/low-voltage and was quite different in its mechanism of operation from EB-guns with much higher acceleration potentials. An oil-free vacuum system capable of 5 x 10- l0torr, a quartz crystal rate and thickness monitor and a quadruple mass spectrometer completed the evaporation system, Fig. 2. A typical evaporation cycle consisted of a 3 to 4 hr pumpdown to the upper l0-9 range and evaporation at l0? per sec with the pressure in the bell jar not rising above 1 x 10"7 torr. Thickness control was 5 pct or less and could be automatically monitored and controlled. Five phenomena associated with the EB evaporation and considered as possible contributors to Ma performance included a purification effect, ionization of evaporating aluminum, X-rays, constitution of vacuum ambient during evaporation, and film structure dependence upon evaporation rate. These phenomena are now discussed. Vacuum Purification. The design of the EB-gun permitted purification of the aluminum charge by vacuum outgassing. Particular features included an efficiently water-cooled copper hearth with a capacity of over 20 g of aluminum and the capability for sweeping the beam across the charge. Such capacity meant that aluminum had to be added only after about every fifth evaporation. A new charge was not required each evaporation as is necessary with filament evaporation. An oxide "scum" which appeared on the charge could be completely cleared from the top hemisphere of the charge by sweeping with the beam prior to opening the shutter. An indication of the purifying effect was obtained by a series of analytical measurements on incoming aluminum, after melting but with little vacuum out-gassing, after 30 min outgassing, and the evaporated film itself. Either a solids (spark source) mass spectrometer or an emission spectrometer were used for analyzing the aluminum charge. Analysis of the evapo-

Jan 1, 1970

By Mark J. Klein

The creep of 043 was studied over the temperature range 1650" to 3200°F and over the stress range 3000 to 44,000 psi. The steady-state creep rate over this range of stress and temperature can be expressed by the equation where A is a constant, is the stress, and is -0.8 x 103 psi-&apos;. Over a narrow range of stress variations c0 a and for this proportionality n varies from 3 to 30 in accordance with the relation n = aB. Above about 2400° F, H, the apparent activation energy for creep, is 110,000 cal per mole, a value about equal to that estimated for self-diffusion in this alloy. Below 2400°F, H increases with decreasing temperature reaching a value of -125,000 cal per mole at 1700° F. In this temperature region, H appears to be a function of the interstitial concentration of the alloy. MOST of the detailed creep studies of dispersion strengthened metals have been concerned with metals having fcc structures. However, there are a number of important refractory alloys with bcc structures that derive part of their high temperature strength from an interstitial phase and whose creep behavior has not been well defined. This paper describes the creep behavior of the bcc alloy, D43, over the temperature range 1650" to 3200°F (0.4 to 0.7 Thm) and over the stress range 3000 to 44,000 psi. In addition to colum-bium, this alloy contains 10 pct W. 1 pct Zr, and sufficient carbon (-0.1 pct) to form a carbide dispersion throughout the matrix of the alloy. The effects of variations in temperature and stress on the steady-state creep rate of this alloy are presented in this paper. EXPERIMENTAL PROCEDURES Creep tests were made in a vacuum of 106 torr under constant tensile stress conditions using a Full-man-type lever arm.&apos; Creep specimens were machined from 0.020-in. D43 sheet (grain size -5 x l0-4 in.) processed in a duplex condition (solution annealed -2900°F, 40 pct reduction in area, aged 2600°F). The specimens were tested in this condition without further heat treatment. Specimen extensions over 1-in. gage lengths were continuously recorded using a high temperature strain gage extensometer. Differential temperature and stress measurements were used to determine temperature and stress dependencies of the creep rate. Activation energies were calculated from the changes in strain rate induced by abrupt shifts in the temperature during constant stress creep tests. The 100°F temperature shifts used in most of the activation energy determinations required 15 to 90 sec depending upon the temperature at which the shift was made. The dependence of strain rate on stress was determined by measuring the change in strain rate for incremental stress reductions during constant temperature tests. It has been shown that columbium-base alloys such as D43 are susceptible to contamination by gaseous interstitial elements during vacuum heat treatments.&apos; In this regard, it is unlikely that these alloys can be heat treated without some loss or gain of interstitial elements despite the precautions taken to control the heat treating environment. However, several factors suggest that changes in interstitial concentrations of the specimens during testing did not affect the results presented in this paper. First, the dependence of the creep rate on the stress or temperature determined during the course of a single creep test showed no variations with the duration of the test. A variation would be expected if a loss or gain in interstitial concentration during the course of the test affected results. In addition, precautions taken during this investigation to minimize interstitial contamination by wrapping the gage lengths of the specimens with various foils2 (Mo, Ta, W) did not produce a detectable change in the stress and temperature dependencies relative to the unwrapped specimens. The averages of duplicate analyses for carbon and oxygen in several specimens determined before and after creep testing are listed in Table I. The combined nitrogen and hydrogen concentrations which were ordinarily less than 50 ppm did not change in a detectable way with creep testing. The analyses show that only minor changes in carbon concentration occurred during creep testing except for specimen 4. This specimen which was tested at 3100°F lost a significant amount of its carbon concentration to the vacuum environment. Specimen 1 gained 100 ppm of O, while specimens 2, 3, and 4, which were tested at progressively higher temperatures, lost increasing portions of their initial oxygen concentrations during testing. RESULTS AND DISCUSSION The Temperature Dependence of the Creep Rate. The apparent activation energy for creep, H, was de-rived from creep curves similar to that shown in Fig. 1. Steady-state creep was rapidly attained at the beginning of the test and with each change in temperature. This behavior suggests that the alloy rapidly attains a stable structure with each shift in temperature or that the structure is constant throughout the test. Since the dispersion will tend to stabilize the structure, the latter is probably the case. The activation energy was found to be independent of the direction of the temperature shift and the magnitude of the shift (50" or 100°F). Although H was approximately independent of the strain, there was a tendency for it

Jan 1, 1970

By S. R. Goodman, Hsun Hu

The formation of secondary - recrystallization tex-tlires in cube-textured 18-8 stain less steel (Type 304) Ilas been studied at three temperatures. Prolonged annealing at 100°&apos;C protluces a PredoninanGly (520) [OOZJ-type texture, which is related to the cube te.ture of the primary lnatrix by a rotation of approxivzately 22 deg around the [001] axis in the rolling direction. Annealing at 1200 or 1300°C facers the formation of the (123)[272/-type texture, which is related to the matrix texture by a [111] rotation of app.voxiniately 40 deg. These observations suggest that in the secondary recrystallization of cube-texlut-ed stainless steel an apparent actilation energy for growth is higher for grains related to the tncrtuix Og [111] rotations thun those reloted by [100] rotations. THE formation of secondary-recrystallization textures in cube-textured primary matrices of fcc metals has been studied widely by various investigators. For Fe-40 pct Ni alloys, Pawlek&apos; and wassermann2 reported that the orientations of secondary grains were related to the cube texture by rotations of 30 and 38 deg around [001] in the rolling direction. However, Rathenau and custers3 found that, while in one Fe-48 pct Ni alloy, most of the secondary grains were oriented with respect to the cube-textured matrix by rotations around [001] of 26.5 deg, in another alloy of a different origin, the orientations of secondary grains were related to the cube texture by rotations of approximately 35 deg around a [lll] axis. Similar orientation relationships were also observed between the secondary grains and the cube-textured primary matrices of copper.4"a No attempt was made to differentiate these two types of orientation relationships; reorientation by either a [111] or a [100] rotation was considered to be equally favored. The present investigation consisted of a study of the secondary recrystallization textures in cube-textured stainless steel. It was noted that the secondary grains formed in stainless steel were considerably smaller than those of Fe-Ni alloys or copper. This offered the advantage that the secondary recrystallization texture could be determined by the texture-goniometer technique, and a more detailed study of the textural development during the course of secondary recrystallization could be made. The effect of annealing temperature on the formation of secondary-recrystallization textures was also investigated. MATERLAL AND METHOD It was shown earlier"-" that a strong cube texture can be obtained in 18-8 stainless steels by rolling at 800°C to produce the copper-type deformation texture, followed by annealing at 800" to 1000°C for recrystallization. To improve the cube texture for the present study, a commercial-grade 18-8 stainless steel (Type 304) was rolled at 800°C first to 5 mm (0.2 in.) thick plates. Three of these plates were then stacked and welded together along the edges into a sandwich assembly. After annealing at 900°C for 20 min: the assembly was finally rolled at 800&apos;C to 90 pct reduction in thickness with reheats and end-for-end reversals after each pass. Only the central strip, which was reduced from 5.0 to 0.50 mm (0.7 in. to 0.020 in.) thick, was used. The chemical composition of the steel in weight percent was as follows: C, 0.06; Mn, 0.38: Cr, 18.71; Ni, 9.56: P, 0.011; S, 0.009; and Si, 0.39. The purpose of rolling the strip in a sandwich assembly was to prevent direct contact between the central strip and the rolls. It was observed earlier" that, when the strip was rolled at 800°C without being enclosed in a sandwich assembly, the cube texture obtained by subsequent annealing at 900" or 1000° C for recrystallization was largely confined to the central section of the strip, while most of the recrystallized grains formed in the surface section of the strip were not cube-textured. This was obviously due to the fact that the actual temperature at the strip surface during rolling, as a result of direct contact between the strip and the cold and massive rolls, was considerably lower than 800°C. By using a sandwich assembly for hot rolling, the cube texture obtained upon subsequent annealing for recrystallization was found to extend through the entire thickness of the strip. After rolling, the central strip was taken from the sandwich assembly. and cut into specimens. Prior to annealing. the specimens were etched to 0.25 mm (0.010 in.) thick. A tube furnace provided with a purified, dry argon atmosphere was used for annealing. Textures were determined by the reflection technique. using a Siemens automatic texture-goniometer and ZrOz-filtered MoKa radiation. With a time constant of 4 sec. the preferred orientation of the secondary grains could be measured satisfactorily by the integrated intensities. Both (111) and (200) reflections were measured, and corresponding pole figures were constructed according to the techniques described previously.10 The agreement between results deduced from these two reflections was excellent. RESULTS AND DISCUSSION Secondary-Recrystallization Texture due to Prolonged Annealing at 1000°C. Fig. 1 shows the primary-recrystallization texture of a specimen annealed at 1000°C for 30 min. A substantial improvement in both sharpness and intensity of the cube texture, owing to the present processing method, can be noted readily by comparing Fig. 1 with similar pole figures shown earlier in Refs. 9 and 11. Secondary recrystallization

Jan 1, 1967

By E. S. Bumps, W. F. Craig, M. Baeyertz

Metallurgists have looked at fractures macroscopically for many years and have evolved a vocabulary in which such words as "cleavage," "brittle," "shear," "ductile," "granular," "fibrous," and "silky" are used to describe the appearance of the fractured surface; but the meaning of these words in terms of metal structure is not well established. Observations of the structural meaning of "brittle" and "ductile" fractures in plate steels have been made, notably, by Kramer and coworkers1 and by Tipper.2 Grossman3 has studied the fracture of tempered martensite and combinations of ferrite and martensite. Notwithstanding these and other less concerted attacks on the problem, present understanding of fracture rests more on assumptions and logic than on experiment. It is the purpose of this paper to add a little to the growing fund of experimental observations of the nature of fractures in steel. The particular fractures to be described were obtained in conventional impact testing of an ordinary structural steel shape. In impact tests of the Charpy type, the specimens fail in a characteristic manner that depends on the steel and the temperature of testing. With ordinary structural and many other steel products, an appropriate range of testing temperature will cause a considerable change in the energy absorbed before fracture. This change is known as the energy transition and is accompanied, more or less closely, by alteration in the appearance of the fracture. At testing temperatures below the transition, the terms "brittle," "cleavage," or "granular" are used to describe the fracture; above it, "ductile" or "shear" are often used. Within the transition, the fracture changes from "brittle" to "ductile" by a progression in appearance, wherein the "brittle" portion of the fracture becomes restricted to a smaller and smaller central area of the fracture, as the testing temperature is raised and the "ductile" type of fracture is approached. All this is well known to metallurgists. The specimens under consideration are of the conventional V-notch Charpy type. They were taken from a structural steel shape of the following analysis: C Mn P S Si 0.17 0.46 0.009 0.029 0.03 Both longitudinal and transverse specimens were tested and examined. Fig 1 shows the energy values plotted against testing temperatures in the usual way. The curves for both longitudinal and transverse specimens show an energy transition, although the maximum energy absorbed at testing temperatures above the transition is greater in the case of the longitudinal specimens. Fig 2 illustrates the changes in the macroscopic appearance of the fractures that are associated with the energy transitions shown in Fig 1. Macrographs of the fractured surfaces of the specimens have been identified with the fracture ratings at intervals on each fracture curve. The numerical fracture rating on the ordinate of Fig 2 indicates the percentage of the area of each fracture that was considered to be "brittle" on macroscopic observation. Such rating of the fracture type in impact tests is now customary in many laboratories. Hereafter the specimens will be identified by reference to Fig 2 as 90 pct "brittle," 80 pct "brittle," and so on, in accordance with this macroscopic evaluation of the appearance of the fractured surface. The discussion of the fractures is divided into three parts. The first two are concerned with the mode of fracture and its relation to general structural features. As the same metal-lographic features occur in both longi-

Jan 1, 1950