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By M. Nakai, G. Itoh, R. Yamada, A. Kurumada

In order elucidate the mechanism for poorer fatigue properties of 7000 series aluminum alloys than those of 2000 series alloys, the effect of the test environment on the fatigue crack growth of the two series alloys was investigated, considering the relationship to hydrogen embrittlement. Two 7075-type and four 2024-type alloys were T6-tempered and T4-tempered respectively, and subjected to fatigue crack growth tests in moist and dry environments. When tested in moist air, fraction of intergranular crack path was higher in the 7075-type alloy with coarse equi-axed grains than in the 2024-type alloy with similar microstructure. This was in accord with the larger crack growth rate in the 7075-type alloy. Both the intergranular crack path fraction and the crack growth rate were decreased by changing the environment to the flow of nitrogen gas. Thus, the large fatigue crack growth rate of the 7075-type alloy is attributable to hydrogen embrittlement.

Jan 1, 2015

By Enever JR

In-situ rock stress measurements conducted in the Phillipines and on Bougainville have revealed stress fields significantly different from expectations based on intraplate experience. The results have potential significance in regard to both engineering practice in the Pacific Rim region as well as in helping to understand the geological processes occuring near subduction Zones.

Jan 1, 1987

Computerisation of mine planning activities, specifically in the areas of geology, design and surveying, has been in progress for a number of years in underground metalliferous mines. The combined capabilities of hardware and software have only recently reached the levels necessary to fully support integrated mine planning computer systems in a graphical three dimensional environment. Despite the apparent capabilities of systems now on the market, many sites retain significant parts of their older paper based systems, and employ a mix of two or more software packages. This paper examines what underground metalliferous sites are doing in this area in Australia, and discusses the factors which are inhibiting full implementation of fully integrated computerised systems. Future directions for further developments in the field are also briefly discussed.

Jan 1, 1995

By H. R. Page, A. E. Jr Lee

From the inception of zinc retorting on a commercial scale in the United States in 1890,' the retort employed has undergone wide variations in its composition and manufacture, facilitating in part equally remarkable improvements in furnace capacities. The early day hand made clay retort was charged with carbonates or silicates or with coarse dead roasted concentrates mixed with a large proportion of charge fuel resulting in a relatively low zinc burden and fired 24 hr in direct coal fired furnaces. Its modern counterpart is fabricated in hydraulic presses from clay mixtures containing sizeable amounts of either silicon carbide or silica flour, charged with sintered flotation concentrates to more than three times the early day zinc burden and fired 24 to 48 hr in gas fired furnaces. This paper does not attempt to describe in detail the early day clay retort practice as it is well outlined in treatises by Ingalls,2 Lodin,3 Liehig,4 Hofman5 and others. A brief review of clay retort practice is presented together with a description of the major developments since 1912. Clay Retorts The Belgian type retort, both in the circular and elliptical forms, has been used almost exclusively. Typical dimensions of press made clay retorts around 1910 are shown in Table 1. Variations in these dimensions were used at different plants according to local conditions to a maximum inside diameter of 9 in. and inside length of 54 in. However, the effective heat penetration in a 24 hr firing cycle and the tendency of the retort to bend limited the retort size. Use of the elliptical vessel was an attempt to present a stronger cross-section resisting the tendency to bend and to increase the burden without increasing the depth of heat penetration. One exception to the 48-54 in. length was the 60 in. retort used as early as 1905 at Palmerton by means of supporting the last 12 in. at the butt end with a specially designed furnace back-wall. This backwall construction with the 60 in. retort had been developed and used at Bethlehem by G. G. Con-vers and A. B. DeSaulles. An attempt was made at Blende, Colo. to use even larger retorts of the Rhenish type based on European practice and requiring much higher furnace temperatures. Satisfactory plastic clays capable of withstanding these temperatures were not available, and the plant never operated successfully. PREPARATION OF BATCH Composition of the clay retort by weight was 40 to 50 pct raw clay and the balance "grog." Generally speaking the mix consisted of 7 parts plastic clay to 9 parts grog by volume. Principal source of the clay used was the Cheltenham vein—sometimes referred to as "St. Louis city clay." A typical analysis was A12O3-31.0 pct, SiO2-50.0 pct, Fe2O3-2.5 pct, MgO-0.3 pct, CaO-1.5 pct and loss on ignition 14.0 pct. At the smelter the clay was weathered whenever possible and then crushed to 0.08 in. or finer. Grog consisted of calcined adobies, cleaned retort scrap and cleaned refuse fire brick such as old furnace brick, blast furnace linings, and others. Saggers from ceramic plants and calcined flint clay were later used. The grog materials were ground to 0.12 m. or finer. Occasionally coke dust up to 10 pct of the mix was substituted for a part of the grog following European practice.² Particle size of the grog has a major influence on the retort properties—the larger the grain, the better can the retort withstand thermal shocks, resist bending at furnace temperatures and resist corrosion from slag; the smaller the grain, the lower the loss of zinc vapor through the retort walls. Grog forms the skeleton of the retort, and the clay shrinks around its grains to act as a binder. In the drying process, the grog has a stabilizing effect on the drying rate, decreasing shrinkage and giving up previously absorbed water to the surrounding clay to minimize the danger of cracking or checking.² Grog and clay were mixed through a horizontal pug mill with 10 to 20 pct water added, depending on whether the retort was to be formed by hand or mechanically, more water being required for the hand process. The batch or "mud" extruded from the pug mill was cut in convenient lengths for handling, stacked in piles or in special rooms, covered with wet burlap and allowed to "rot" or age from 1 to 8 weeks to increase plasticity. HAND MOLDING If the retort was to be molded by hand, the mud was repugged after the rotting period and given to the molders. Their molds consisted of 3 sheet iron or wood cylinders, each one third the retort length and defining the outer shape of the retort. Beginning with the bottom section, mud was placed in the form and tamped with a rammer

Jan 1, 1950

By Klaus-Peter M. Hanke

INTRODUCTION In early 1979 Jim Walter Resources, Inc. (JWR) of Brookwood, Alabama approached TMCI Construction, Inc (TMCI) to make a proposal on a program that involved the sinking of up to 10 ventilation shafts of approximately 6.7 m (22 ft) diameter and ranging in depth from 500 to 700 m (1650 to 2300 ft) for the JWR coal mines in Alabama. At this time TMCI was already constructing the first spiral underground bunker (capacity 2000 tons) in North America for the JWR organization at their No. 4 mine in Alabama. The TMCI proposal was based on the use of the mos modern large diameter shaft boring machine rather than sinking the shafts using the conventional drill blast-muck technique. The proposal was made based o: the experiences by the parent company, Thyssen Schachtbau, which has been using this type of machin in Germany for shaft boring since 1971. As a result of the TMCI proposal JWR issued a purchase order to TMCI for the construction of four 6.7 m (22 ft) diameter, concrete lined, unfurnished ventilation shafts ranging in depth from 500 to 700 (1650 to 2300 ft). An order was thus placed with WIRTH Machinen- and Bohrgeraete Fabrik GmbH, in Germany for the manufacture of a model 650/850 E/Sch "Schachtbohrmaschine" (Vertical Shaft Borer = V-mole which arrived on site in Alabama in early 1981. The first V-mole GSB 450/500 was introduced in Germany in 1971 and was capable of enlarging in one step a pre-drilled 1.2 m (4 ft) pilot hole to 4.5 - 5.0 m (14.7 - 16.4 ft). This machine has sunk 9 staple shafts and deepened one surface shaft for a total of 2360 m (7740 ft) of shafts. On the last shaft boring operation in 1978 the machine was converted as an experiment to drill without a pilot hol using a hydraulic pumping system to remove the cutting debris. A second generation machine, the SB VI 500/650, was introduced in 1977 for enlarging the pilot hole to a range of 5.0 - 6.5 m (16.4 - 21.3 ft) diameter. This machine is still in operation and has already drilled well over 2000 m (6500 ft) of shaft. The third generation of V-mole, the SB VII 650/85( for diameters from 6.5 to 8.5 m (21.3 to 27.9 ft) was: commissioned in May 1980 and has been used for two surface shaft deepenings totalling 606 m (1990 ft) with another scheduled for 1982. The main advantages favouring the use of such V-moles were identified as: 1) A reduction in manpower to the crew required in a conventional shaft sinking operation. 2) A considerable reduction in time to complete a shaft compared to conventional techniques. 3) The use of the V-mole eliminates many of the hazards encountered in conventional sinking. Based on the successful performance of the first three V-moles in Germany, Thyssen Schachtbau decided to employ this principle abroad. In 1980 a second machine of the third generation was built and is now operated by TMCI Construction, Inc. in Alabama. The first shaft was completed at the end of 1981 and this paper describes the method of operation including some unique aspects not attempted on prior V-mole operations and some of the statistics arising out of the experiences during the first shaft boring operation. THE NO. 7 MINE FAN SHAFT SITE Jim Walter Resources, Inc. was formed in 1970 to exploit the coal field in Alabama on the southern tip of the Appalachian coal field. The coal reserves amount to around 650 million tons of mainly good quality coking coal of which about 350 million tons are to be extracted over the next 30 years. Shaft sinking and preparatory work began in 1972, and at present 6 mines are producing around 5.4 million t.p.a. Annual production is to expand to 10 million t.p.a. as soon as possible, and the ventilation shafts to be sunk by TMCI play a vital role towards attaining this goal. The first shaft site is located at the No.7 mine, near Brookwood, Alabama. The actual location of the shaft relative to the production shafts is shown on the mine plan (Fig. 1), which also shows the room-and-pillar extraction system used at present. The mine plan further shows the conveyor route used for the muck removal. The geological survey showed that the strata consisted of horizontal layers of mainly sandstone, sandy shale and shale interspersed with several coal seams. The seam being extracted at the No. 7 mine is a combined seam made up of the Blue Creek and Mary Lee seams at a depth of 513 m (1682 ft) and having an average seam thickness of about 2 m (6 ft). At the beginning of September 1980 the surface site preparation and pre-grouting work was completed by JWR, and TMCI was able to commence with Stage I of the shaft sinking program - the drilling of the pilot hole.

Jan 1, 1982

By C. A. Stickels

A heat of electrolytic iron, to whzch alunzinutn and nitrogen had been added, was hot-rolled, cold-rolled 90 pct, and recrystallized at temperatures from 500" to 700°C. Primary recrystallization textures appear to arise from competitive growth of two types of nuclei: 1) those having orientations belonging to the "usual" primary recrystallization texture found in riming steel, and 2) those with the {111} (110) ovientation. Development of a (111}(1 10) component in the primary recrystallization texture occurs only over a certain interval of isothermal recrystallizatzon temperatures when the material is supersaturated with respect to the precipitation of AlN. Lowering the degree of supersaturation depresses the temperature interval in which a (111)(110) component occurs. An elongated, 'pancake-shaped" recrystallized pain structure and a marked delay in the start of recrystallization were found in all specimens which were supersaturated with respect to A1N precipitation after cold work, regardless of their recrystallization texture. ONE of the consequences of killing low-carbon steel with aluminum is a significant change in recrystallization behavior. About 15 years ago, Solter and eatttiel showed that this behavior was largely controlled by aluminum and nitrogen in the steel. If complete precipitation of A1N was prevented before cold rolling, an increased "recrystallization temperature" was observed in subsequent. annealing, and the recrystal-lized grains were not equiaxed. Leslie et a1.2 studied this phenomenon in some detail and clearly demonstrated the relationship between A1N precipitation, recrystallization kinetics, and the development of "pancake-shaped" grains. It has also been known for some time that aluminum-killed steels, processed to produce elongated "pancake" grains, develop a (11 I}( 110) primary recrystallization texture. This texture has not been found in iron or low-carbon rimming steel as a primary texture4j5 but has been observed following grain growth in electrolytic iron.5 The present work was undertaken to study in more detail the effect of A1N supersaturation on recrystallization textures in iron. LITERATURE REVIEW The deformation texture in heavily rolled iron has been studied in detail by Bennewitz.~ The texture consists primarily of a partial fiber texture about a (110) axis in the rolling direction, designated here as fiber texture A. It includes the range of orienta- tions (111)[110] - (001)[ 110] - (11l)[110]. A weak secondary texture also is present.6 This is a duplex partial fiber texture about two (110) fiber axes located 60 deg from the rolling direction and 30 deg from the sheet normal. The range of this texture, designated here as fiber texture B, about the [101} fiber axis is (112)[110] - near (545)[252] - (211:1[011] *The range given here follows Bennewit~.~ A few pole figures from re-crystallized material indicate a broader range than this.' However, the components which are strongest in the recrystallization texture are in this range.'________________________________________________________ Primary recrystallization textures in unkilled steels can be accounted for by growth of members of fiber texture B present in the deformed metal.5 However, while members of fiber texture B dominate the primary texture, other orientations survive primary recrystallization as well. In particular, some {111}(110) members of fiber texture A must also grow during primary recrystallization, because a well-defined {1ll)( 110) texture develops during subsequent grain growth at 700°C.5 The unusual recrystallization behavior of deformed supersaturated solid solutions has been attributed to: 1) retention of the solute in solution,' 2) formation of coherent, preprecipitation solute clusters prior to and during re~r~stallization,~ and 3) formation of a precipitate prior to and concurrent with recrystallization.'~-'~ When aluminum is supersaturated with iron, the difference in grain boundary mobility between general high-angle boundaries and certain special coincidence site boundaries is apparently eliminated.' In aluminum-killed steels, precipitation of A1N can take place at ordinary subcritical recrystallization temperatures. The rate of precipitation increases with increasing aluminum or nitrogen contents.2'13 There is some doubt, however, as to whether true precipitates form during the time at temperature needed to complete recrystallization. Leslie ef a1.2 found that precipitation in one steel was complete after about 100 min at 700GC, or after about 1000 min at 650GC, as measured by chemical analysis for AlN. Aoki et a1.,13 using internal friction for dissolved nitrogen, showed that a large fraction of the dissolved nitrogen was removed from solution within a few minutes annealing time at temperatures from 400" to 800°C. However , the rate of formation of AlN, as detected bv chemical analvsis. was much slower than the apparent rate of nitrogen removal. Hasebe,'~~ using carbon extraction replicas, has identified A1N precipitates by electron diffraction in a 0.2 C steel, solution-treated at 1300°C and annealed 2 hr at 700°C. Borchers and kim,I6 also using a replication technique, observed precipitates after annealing treatments as short as 2 min at 640°C. However, Leslie et a1.' state that no A1N precipitate can be seen while recrystallization is being inhibited in aluminum-killed steel.

Jan 1, 1967

By J. E. Johnson

Discussion of the paper of Mark A. Amnion, presented at the Cleveland meeting October, 1913, and printed in Bulletin No. 70, October, 1913, pp. 1167 to 1 79. J. E. JOHNSON, JR., Ashland, Wis.:-Mr. Amnion's paper is of great value and it is not with any desire to reflect. on it that I wish to make some comments on the use of the scleroscope. We obtained one on trial some months ago and made very extensive tests with it, but without getting very satisfactory results. The readings on good car-wheels and bad ones were just about the same, but the factor in its use most likely to produce erroneous results is a variation in the thickness of the specimen being tested. In order to find out something on this subject we cut a thin disk, about -1,: in. thick, off a brass rod; then one 0.25 in. thick, and another about 1 in. thick, put the surfaces of all in about the same condition and tested them with the scleroscope. The readings varied widely; the resilience of the anvil evidently entered to a very high degree on thin pieces. By putting a soft cushion between such a thin specimen and the anvil, and forcing it up against the tube, the softness of the cushion cut off the resilience of the anvil and the results were more like those of the thick specimens. After a thickness of 1 in. was passed, the specimen seemed to have mass enough so that the resilience of the anvil no longer played an important part, and above that thickness results on different materials would be comparable, even though the thicknesses differed. But in order to get comparable results on thin specimens the thickness of those being tested should be the same. I should like to ask how the variations introduced in this way were overcome, so as to get strictly comparable results. Mr. AMMON :-.I have not had much experience in that, but the scleroscope is simply an indication of the surface harden-

Dec 1, 1912

By R. M. Campoy

In this review of the copper, lead, zinc, and molybdenum industries in Mexico, the author describes how the country's position in the world base metals market and the climate for foreign investment in Mexico have been affected by the nationalization of the banking industry in that country and by recent mining legislation and the current national debt crisis.

Jan 1, 1988

By R. Bhaskar

The results of an experimental and theoretical research study into the spatial and temporal behavior of dust in mine airways is summarized in the paper. Experiments were performed in mine airways under both controlled and normal operating conditions. The controlled experiments included the documentation of the concentration and deposition patterns for two types of dust under three different velocities. Real-time dust concentration data were obtained using Real-time Aerosol Monitors (RAM-I) to determine the dispersion coefficient in a mine airway. On the theoretical front, a convection-diffusion model of the dust flow phenomenon in mine airways was developed considering dominant mechanisms affecting dust transport and deposition in mine airways. Mechanisms modeled include turbulent and gravitational deposition, coagulation, and dispersion.

Jan 1, 1988

By D. Cronan

Both the S.W. Pacific and Caribbean island arcs are host to submarine hydrothermal mineral deposits, the former more so than the latter. In the S.W. Pacific, the whole gamut of such deposits occurs on the sea floor, sulphides, silicates, Fe and Mn oxides, whereas during the limited studies to date in the Caribbean only hydrothermal Mn oxides have been found to occur on the sea floor there in any abundance. Comparisons between the respective Mn oxides are thus the most fruitful at the present limited stage of our knowledge of Caribbean submarine hydrothermalism. Southwest Pacific hydrothermal ferromanganese oxides and their possible relationships to underlying sulphide mineralisation or bedrock anomalies (the geochemical window concept) has been a fruitful field of research in recent years. Certain element enrichments in crusts may be indicative of underlying mineralisation (Rogers et al. 2001). However, more recent studies on buried hydrothermal crusts in the S.W. Pacific (Cronan et al. 2002) have shown there to have been postdepositional diagenetic uptake of elements in them, and thus these elements cannot be used with any confidence in geochemical exploration for buried sulphide deposits unless sediment involvement can be eliminated.

Jan 1, 2002

By Alan Bauer

In recent years there has been increasing interest in employing the technique of explosive casting in strip mines. Explosive casting, applied to dragline mining, describes the procedure whereby the overburden is horizontally displaced by blasting causing a portion of the material to be thrown into the previously mined cut. Therefore, it is not necessary to subsequently excavate this waste with the stripping equipment. In addition, allowing the overburden unrestricted movement during the blast most often results in improved fragmentation and enhanced machine productivity. The method is receiving attention because the capital and maintenance costs of stripping equipment has risen rapidly in relation to explosives costs. Improved productivity is necessary to maintain unit costs at an acceptable level.

Jan 1, 1983

By Antônio Peres, Rogério Guimaraes, Lauro Takata, Harold Wyslouzil

"Serrana’s concentrator, in Araxii, has capacity of producing approximately 800,000 tomes of phosphate rock per year, with a P205 grade of 36%, from ores grading between 13% and 14% P205. The annual mining rate is around 3.5 Mt and the apatite is concentrated according to the following sequence: crushing, grinding, desliming, flotation and low and high intensity magnetic separations.Two industrial flotation columns (3.0 m x 4.5 m x 14.5 m) were commissioned in 1993 as a replacement for circuits of barite and apatite fines flotation. In 1994, four additional flotation columns of the same dimensions were put into operation, replacing 42 Wemco mechanical cells (300 ft3 each), in the barite pre-flotation, coarse apatite (two columns) and tailings regrind flotation Circuits. Overall, 66 mechanical cells (300 ft3 each) were replaced by only six flotation columns, designed by Cominco Engineering Services Ltd - CESL.The present paper describes the stages of process development, pilot plant testwork and implementation of the industrial project, highlighting the most relevant aspects of each of them. The results achieved are related to improvements in flotation selectivity, gains in P205 recovery and better process control, resulting in better quality products and lower production costs. Other benefits included a decrease in flotation reagents, electrical energy and maintenance materials consumption. These factors contributed to a significant decrease in the cost of phosphate rock production at Serrana, in Araxa - M G."

Jan 1, 1999

By W. A. Tiller

The conventional techniques1 for determining the liquidus and solidus surfaces of an alloy system containing more than two components are extremely tedious to use and do not provide a complete picture of the equilibrium relations between solid and liquid alloys. These techniques are unable to yield "tie-line" information concerning solid and liquid phase equilibrium, a very important parameter in the solidification description of the liquid alloy and a very necessary parameter in the preparation of "zone-levelled" alloy crystals. The tie-line in a polycomponent system is analogous to the partition coefficient, k, in a binary system, it gives the composition of the solid, Cs, in equilibrium with a liquid of composition CL. That is, CS = kCL, where CL = [CO, c1,...Cn] denotes the concentrations of the n + 1 constituents, and k = [ko, kb ,...kon] denotes the gross partition coefficient for the elements between the two phases; thus, k is the tie-line in this system. To provide complete information concerning the two-phase equilibrium in a polycomponent system it is necessary to know both the liquidus surface and the gross partition coefficient. From these two the solidus surface is obtainable. The conventional techniques are unable to provide this information in other than a binary system so we must look elsewhere. In recent years considerable insight has been gained into the correct description of the liquid-solid transformation, 2,3 and controlled solidification experiments may now be designed to both facilitate and enhance equilibrium diagram studies. In the present paper consideration is given to two methods for obtaining the liquidus surface and the tie-lines in polycomponent systems. The methods to be described below deal with the solidification of an n + 1 constituent liquid alloy of initial composition [co, Co ,...Con], where the superscripts refer to the elements present and the Oth element is considered as the solvent in which the n solutes are dissolved. The general assumptions made in the treatment are the following: (i) the solid and the liquid at their interface are in equilibrium during the growth of the solid phase, (ii) there is no diffusion in the solid phase, and (iii) the liquid phase is completely mixed and is therefore homogeneous in concentration. METHOD I In this section a general experimental method will be described which, in principle, is capable of giving an exact description of the liquidus surface and tie-lines. Consider the unidirectional solidification of the liquid alloy specimens L cm in length (freezing either horizontally or vertically). Allow the sample to be frozen very slowly from one end, as indicated in Fig. 1, with complete mixing in the liquid, and then analyze the solid bar to determine its chemical constitution as a function of position, x, along the bar. Let Fig. 2 represent a possible distribution of the ith constituent along the bar. At the point x' the concentration of the ith constituent is C1/5(x'), and the average concentration of the rest of the bar between x' and L is given by C:(X' - L) where [ A(x) c1s(x)dx c1/s(x'-L) = f A(x) dx x * and A(X) is the cross-sectional area at x. In a similar manner. all the Cf(x' - L) may be determined. Thus, the gross artition coefficient k for liquid of composition [Cs(x - L),...Cs(X' - L)] is given by -ko= CUs')/Cos(x' - L)" k0 = csn{x')/cs(x'- L) During the freezing of a charge of length L, the liquid composition may vary over a wide segment of the phase diagram and the gross partition coefficient over this segment of the phase diagram may be determined from one, bar. Fig. 3 illustrates the magnitude of Cs(g)/C,' as a function of the fraction g of the bar which has Solidified.2 We can see that the concentration in the bar will vary over a range of about 15 wt pct for C,' = 10 wt pct and k0 = 0.5. It appears that 4 or 5 specimens would be adequate to study a simple binary eutectic phase diagram.

Jan 1, 1960

By J. L. Brimhall, R. A. Huggins, M. J. Klein

IN a recent paper1 it was shown that small additions of magnesium, copper, and oxygen decrease the stacking fault probability in plastically deformed silver. Correlation of :X-ray data with measurements of the twinning frequency and metallographic observation of slip lines indicated that this effect may be due to an increase in the relative stacking fault energy. It is the purpose of the present note to report experiments on these alloys of known stacking fault probability that relate to the stacking fault contribution to electron scattering in silver. There has been a considerable amount of disagreement in the literature concerning the contribution of dislocations to electrical resistivity in metals. Calculations relating to the scattering of conduction electrons by the dislocation core, its strain fields, and related effects have been made by several authors.2-5 In addition, experimental work on Cu,3,6-10 Ni,11 and Ni-Co alloys1:! has shown that dislocations in those materials exert an anomalously large effect on resistivity, this effect being related to the presence of extended dislocations. The problem of the scattering due to the stacking fault region between partial dislocations has been treated theoretically using several different procedures.13-'0 The results of these calculations show a wide range of predicted values, from rather large to completely negligible, due to the models and the approximations involved. Almost all of this work has been applied to the specific problem of scattering in copper. The more complete theoretical treatments depend rather critically. on the shape of the Fermi surfaces and the effective mass parameters. The scattering from stacking faults might therefore be expected to vary significantly from one metal to another. Unfortunately, sufficiently accurate information is not available .which would allow one to have any confidence in a theoretical prediction of the scattering in the case of silver. Resistivity measurements were made on small wire specimens of vacuum melted high purity silver and several dilute alloys. These measurements were made at approximately 20°C, with great care being given to temperature measurement and corrections. At such temperatures phonon scattering is significantly greater than imperfection scattering. However, the experimental technique was such that measurements on similarly treated samples were reproducible to better than 2.4 X 10-9 ohm-cm, which was satisfactory for this work. Values of the stacking fault probability and the increase in resistivity after severe plastic deformation are given in Table I. The increases in resistivity resulting from plastic deformation agree quite well with previous work on silver.14,21-26 In addition, it has been previously found that interetitials and vacancies in silver anneal out well below room temperature, so that the residual imperfection scattering in these experiments could be expected to be primarily associated with dislocations;. It is seen from the data in Table I that large changes in stacking fault probability are accompanied by differences in resistivity due to imperfection scattering of only a few percent. If the scattering of conduction electrons by extended dislocations is primarily due to the area of the stacking fault region independent of the stacking fault energy, and not the core and strain fields associated with partial dislocations, the resistivity increase should be proportional to the product of the dislocation density and the stacking fault width, as is the stacking fault probability. Therefore, the imperfection scattering caused by plastic deformation of the alloys should be considerably smaller than that found in the pure silver. Since this was not observed, these results sugegest the possibility that stacking faults do not themselves contribute appreciably to the electrical resistivity of silver deformed at room temperature. An alternative explanation of these observations is that the scattering by the stacking fault is roughly proportional to the stacking fault energy, rather than being determined by the total area of such faults. Sufficient evidence is not yet available to separate these two possibilities. This work was supported by the Office of Naval Research.

Jan 1, 1962

By J. J. Gilman

RECENT experiments have shown that the mar-tensite reaction in a standard tool steel is influenced by the history of the reacting austenite. The martensite reaction proceeds to a given extent at a higher temperature in austenite formed from an annealed spheroidal structure than it does in austenite formed from tempered martensite: that is to say, a prequench lowers the temperature range of the martensite reaction in this steel. Furthermore, a prequench alters the appearance of the martensite as compared to martensite formed from "virgin" austenite. The steel used for this work was a manganese nondeforming tool steel of the following composition: C, 0.85 pct; Mn, 1.43; Si, 0.23; Ni, 0.12; Cr, 0.59; and W, 0.59. A piece from a rolled bar was forged to Vz in. sq and annealed to a spheroidal structure. Disks 1/8 in. thick were used in the experiments. Several sets of the specimens were prepared and austenitized in lead at 1475°F. Three sets were held at 1475°F for 15 min, quenched in water, and immediately reheated to 1475°F. After this quench the three sets were held for 5 min, 30 min, and 3 hr, respectively, at 1475°F and then the individual specimens in each set were quenched in molten baths to 325", 300°, and 275ºF, respectively, and held for 15 sec. Immediately after the quenching treatments? the specimens were tempered at 600°F for 30 sec to darken the martensite and then were quenched in water. The holding times were short enough to avoid isothermal transformation. Three other sets of specimens were given the same final quenching treatments as the first three, but they were austenitized at 1475°F for 20 min, 35 min, and 3 hr and 15 min, respectively, without an intermediate water quench. The two individual specimens from each set were quenched together in order to eliminate quenching errors. Several experiments were run in this way but only the six sets referred to are described in detail because they all yielded essentially the same results. By means of the stated heat treatments the martensite reaction was investigated for: 1—austenite formed from an annealed structure, and 2— austenite formed from tempered martensite. The martensite reactions of the two types of austenite differed significantly. This is shown by Figs. 1 and 2. Comparison of the photographs on the left in Figs. 1 and 2 with those on the right shows that less martensite was formed by a quench to a given temperature from prequenched austenite than from austenite formed from an annealed structure, that is, the prequench suppressed the martensite reaction. The effect was observed for specimens austenitized 20 min, 35 min, and 3 hr and 15 min. A comparison of Figs. 1 and 2 shows that the martensite reactions in both sets of specimens occurred at lower temperatures with increased austenitization time. This was caused by the increased dissolution of residual carbides with time at 1475°F. However, the difference in amount of martensite formed in prequenched austenite as compared to virgin austenite was approximately the same for all three austenitization times. One further observation should be noted. The needles of the martensite formed from virgin austenite are somewhat longer on the average and show less of a tendency to cluster than those of martensite formed from prequenched austenite. This causes the martensite on the right in Figs. 1 and 2 to have a more acicular appearance than the martensite on the left. Since the sets of specimens used in this work were austenitized for identical total lengths of time and the effect of a prequench on the martensite reaction existed for specimens held at the austenitizing tem-

Jan 1, 1953

By H. Bartholomew, L. Bernstein

Experirmental bonds of Ag-In SLID to gold, copper, nickel, Kovar, Dumet, nickel-plated molybdenum, ALL-Pt vrzetallizing on ceramic, and nickel-plated Mo-Ti metallizing on Al2O3 ceramic have been made and exalnined. Metallographic taper sections were made of all bonds for Phase identification. In addition, failuve temperatltres were determined for bonds heated undev shear load. Sowe bonds were subjected to electvon-probe microanalysis. The ternary SLID system Ag-In-Sn was similarly evaluated. Plwtonticrogaphs and electron-probe data are discussed and related to over-all phase diagrams , diffusion effects, and wetting properties of the liquid. The rates of diffusion andor compound formation of gold, silver, and copper in indium and In-Sn liquids are found to descend in that order, as measured by the relatzve thickness of phase layers in the interfaces of bonds made at 200to 400°C. It is concluded that the SLID process enables one to produce bonds which are stable at high temperatures, after having been fabricated at low temperatures, and meet most requirements for application in interconnection, silicon bonding, or hermetic sealing of electronic devices. In a previous publication by one of the authors, the theory and principle of SLID bonding in simple binary systems was discussed. SLID bonding can be defined as a joining process using a preform with a laminated structure: a low-melting material clad on a high-melting core. During the bonding cycle a metallurgical transformation occurs such that the resultant structure can withstand, without reliquefying, temperatures higher than those at which the bond was made. The phases present in the joint interface region are high-melting compounds or, if equilibrium is reached, solid solution. Properties will depend on the phases present, ultimately approaching those of the solid substrate core of the SLID preform. The high-temperature strength of bonds made in these systems was correlated to phase identification in the interfacial bond regions from metallographic cross sections and electron-probe X-ray microanalysis data. When complex electronic components are assembled, multiphase systems result at the interfaces of bond regions. In these ternary and higher-order systems very little diffusion and phase-diagram data is available with which to predict the nature of the interfacia1 region that will result from conventional fabrication processes and subsequent field use of the component. Clarke and Rhines did early work in the field using the Al-Mg-Zn system as a vehicle. Guy and co-workers studied a portion of the Cu-Zn-Ni system and some mathematical4 and general5 aspects of multiphase diffusion. irkald' and coworkerse recently studied diffusion in the Cu-Zn-Sn system. No diffusion studies have been made on ternary systems directly applicable to this investigation but some phase diagrams are available: g-u-n, u-i-n, and CU-i-n. I) DESCRIPTION AND EXPERIMENTAL PROCEDURE The elements Au, Ag, Cu, Ni, In, and Sn are found in alloy systems commonly used in fabricating and assembling electronic components. Reference to the appropriate phase diagrams9 and the previous publication will indicate that the binary systems of any of the two low-melting elements, indium and tin, with any one of the other four higher-melting elements are capable of producing SLID bonds. It is not unreasonable to assume that any ternary or higher-order alloy system containing at least one of the elements gold, silver, copper, or nickel and one or both of the elements indium and tin could be used to produce bonds which are stable at high temperatures, after having been fabricated at low temperatures. Indeed, the liquidus contours and isothermal sections of those phase diagrams available7 do so indicate. The parameters affecting the bond and the equipment used in evaluating the bond were described in an earlier paper.' The current investigation used the

Jan 1, 1967

By W. C. Winegard, J. E. Gruzleski

The stages in the development of cells in the Sn-Cd eutectic have been studied by unidirectionally solidifying specimens under known conditions of growth rate, temperature gradient, and impurity concentration. By means of a quenching technique, the shape of the solid-liquid interface during solidification could be observed. Cells are shown to develop from defects or depressions which exist on the solid-liquid interface of even a pure eutectic. Th,e container wall, grain boundaries , and fault lines in the micro structure play an important role in cell development. A two-phase eutectic dendrite found at high impurity concentrations is described. L HE cellular or colony structure formed during eutectic solidification caused much confusion among early workers in their attempts to classify eutectic structure, and a major advance in our understanding of eutectic solidification was made when the colony structure was identified with a cellular interface caused by the presence of impurity elements.1'2 To date, however, no detailed investigation of the stages in cell development in eutectics has appeared, and little work has been done on the cell morphology. The present paper presents the results of such an investigation. EXPERIMENTAL The Sn-Cd eutectic was used in this investigation with lead as the impurity element to promote cell formation. Materials of 99.999 pct purity were zone-refined to an approximate purity of 99.9999 pct. Alloys as close as possible to the eutectic composition, 67.75 wt pct Sn, were prepared using zone-refined tin and cadmium. The alloys were then zone-melted to the exact eutectic composition, as described by Yue and Clark,3 using a total of twenty to twenty-five passes on each eutectic bar. Only those portions of the bars which contained no primary phases were used in subsequent experiments. Master ternary alloys were prepared by melting together the required amounts of lead and the Sn-Cd eutectic in a sealed Pyrex tube under argon, and these alloys were then diluted with pure Sn-Cd eutectic as required. Unidirectional solidification was conducted vertically upward using the two-element resistance furnace shown schematically in Fig. 1. With this apparatus, it was possible to obtain a range of growth rates and temperature gradients so that the cell structure could be studied under different growth conditions. Growth rates in the range of 0.4 to 9.1 cm per hr and temperature gradients from 2.25° to 1475°C per cm were used. To ensure steady-state solidification, power to the furnace was stabilized, and the entire apparatus was enclosed in a large box in which the temperature was controlled to within ± 0.2°C. The apparatus was arranged so that the specimen could be quenched in situ during solidification by pouring a large quantity of water down the central vycor tube. Excellent delineation of the solid-liquid interface at the time of the quench was obtained because of the structural changes accompanying the rapid solidification of the remaining liquid. The quenching rate was always extremely rapid as no evidence of a transition in the lamellar spacing at the interface was ever observed. The alloys to be solidified were contained in degassed, high-purity graphite tubes, 1 cm OD and 0.6 cm ID. The specimens were about 18 cm in length. After the alloy had been cast into the tubes, three 0.8-mm holes were drilled in the tube to the center of the specimen. One hole was placed 1 cm from the bottom of the ingot, and the other holes were placed at 8 and 9 cm, respectively, from the base of the ingot. Thirty-four-gage chromel-alumel thermocouples were sealed in the holes. The lower thermocouple was used to position the solid-liquid interface at the start of each experiment,

Jan 1, 1969

By G. Derge, Ling Yang, John Henderson

Self-diffusion coefficients of aluminum have been measured by the capillary reservoir technique in liquids of the CaO-SiO2-Al2O3 system containirg equimolar portions of CaO and SiO2, in the temperature range 1400o to 1520oC. For the melt containing 6.0 mole pct Al2O3 the results can be expressed by the equation D = [(4.3 * 0.3) x 104] exp [(-85,000 * 17,500)/RT]cm~sec-'; for the melt containing 12.5 mole pct Al2O3, by the equation D = (5.4 i 0.2) exp [(-60,000 * 10,00O)/RT] cm2sec-: A concept of the constitution of these melts has been developed which proposes that the ratios O/(Si +Al) and Al2 O3/CaO determine the dominant species present. In recent years considerable work has been done on the system CaO-SiO2-A12O3 with a view to elucidating the structure of liquids of this system. Conductivity,' electromotive force, 3 density,4 and activity5 measurements have all contributed to an understanding of the problem, but as yet there is no entirely satisfactory theory that can account for all the observed experimental results. In principle, measurements of self-diffusion coefficients should give further information which is not obtainable in other ways. Self-diffusion coefficients for calcium, silicon,6 and oxygen7 have been measured. The present investigation extends the series of self-diffusion coefficient measurements to include aluminum and develops a concept of constitution for liquids of this system which is compatible with all the observed experimental results. EXPERIMENTAL The technique used in the measurement of the self-diffusion coefficient of aluminum was the capillary-reservoir method,8 A126 being incorporated in the melt contained in the capillary and inactive aluminum in the reservoir melt. This method was chosen because it minimizes convection currents' in the diffusion zone. It has the further advantage that the boundary conditions are such that a slight modifica- tion of Anderson and saddington's8 solution of Fick's Law yields an equation which requires measurements only of the length and the initial and final average specific activities of the solidified melt within the capillary. where t is the diffusion time, Co is the average specific activity of the melt contained in the capillary at t = O, C is the average specific activity of the melt contained in the capillary at t = t, 1is the length of the capillary, and D is the self-diffusion coefficient. The equation is of somewhat different form than those previously applied to high temperature diffusion studies because the specific activity of the melt in the reservoir was negligible both before and after diffusion. The diffusion apparatus is shown schematically in Fig. 1. It consisted of a silicon carbide resistance furnace containing the inactive melt in a graphite crucible, internal dimensions 7/8 by 4 in., enclosed in a 1 1/4- by 1 1/2- by 30-in. McDanel tube. A 3/8-in. McDanel thermocouple sheath to which graphite radiation shields and a graphite protection tube were affixed, was used to hold the capillaries. By moving this sheath through a rubber seal at the top of the McDanel tube the capillaries could be lowered into or raised out of the melt. Temperature was measured with a Pt-Pt 13 pct Rh thermocouple which was calibrated periodically against a standard couple. The furnace temperature was controlled to ± 5°C through the thermocouple immersed in the melt. Temperature variation over the depth of the melt was less than ± 3°C. The inactive melts were prepared from pure native quartz, assaying better than 99.9 pct SiO2, A.R. grade A12O3 and A.R. grade CaCO,. The same quartz and CaCO3 were used in preparing the active melts, the

Jan 1, 1962

By B. S. Gottfried

A generalized mathematical model is presented, which describes the thermal recovery of oil in linear systems with convective external heat loss. Three-phase fluid flow, conduction-convection heat transfer, chemical reaction between oxygen and oil and aqueous phase change are included, making the model applicable to a variety of thermal recovery processes. The complexity of the model requires that numerical solutions be obtained with a digital computer. The numerical approximations of the partial differential equations are presented and the algorithm is discussed in detail. The analysis is applied to the detailed simulation of a forward-combustion laboratory tube experiment. The results of the simulation exhibit with reasonable accuracy all of the principal thermal and bydrodynamic characteristics observed in the laboratory (i.e., the propagation of a combustion zone, a steam plateau, and oil and water banks). Numerical variations in the parameters appearing in the external heat-loss term, the aqueous phase-change term, and the Arrhenius combustion term are examined, and their effects on the predicted temperature profiles, oil and water production histories, and the air-oil ratio are presented. In spite of an oversimplification in the chemical reaction mechanism and the considerable computer time required to obtain solutions, the model represents a major advance in the ability to simulate the phenomena observed in thermal recovery experiments. INTRODUCTION Thermal oil recovery refers to a class of recovery processes where heat is supplied to a reservoir to provide the necessary expulsive energy. This thermal energy can be supplied externally as steam or hot water, or it can be generated in situ by forward or reverse combustion. In either case, however, thermal recovery processes are charac- terized by the simultaneous flow of two or three fluid phases in a variable - temperature field, accompanied by possible chemical reaction or phase-change effects.2-5, 7- 10, 14,16-18,20-22 Although a physical understanding of the thermal recovery processes is far from complete, it is possible to construct mathematical models which describe approximately all of the principal physical and chemical phenomena. However, attempts to solve such models, even with high-speed computers, involve formidable mathematical difficulties. Consequently, theoretical solutions have been obtained only for idealized cases in which important physical phenomena are neglected. For example, consider the process of forward in situ combustion. All such theories which have been developed consider only certain aspects of the process, such as heat transfer, 2-4,8,9,l6,l8 heat transfer with phase change,10 heat transfer with chemical reaction,7 or the hydrodynamics of three-phase flow.21 A general theory including all of the above phenomena has not been developed to date. This paper presents a unified theory of thermal recovery processes in linear systems. A mathematical model is developed which explicitly includes conduction-convection heat transfer with convective external heat loss, chemical reaction between air and oil, aqueous phase change, and the hydrodynamics of three-phase flow. A system of equations is developed which can be solved numerically on a high-speed digital computer, resulting in predicted temperature, pressure, and saturation histories in space and time. The model allows a more detailed simulation of thermal recovery tube experiments than had previously been possible. THEORETICAL DEVELOPMENT Consider the linear flow of gas, water and oil in a homogeneous porous medium. Assume that the oil will react with gaseous oxygen, and that mass is transferred between the water and gas phase by evaporation or condensation. In the absence of hydrocarbon phase change, intraphase diffusion,

Jan 1, 1966

By D. W. Frommer, M. M. Fine

Mineral dressing research showed that iron concentrates of commercial quality could be produced from the Pea Ridge deposit near Sullivan, Mo. Magnetic separation and flotation, on a laboratory scale, yielded concentrates ranging from 60.0 to 71.4 pet Fe at recoveries of 89.9 to 97.6 pct. Early in 1957 St. Joseph Lead Co. announced discovery of three new centers of iron ore deposition in east central Missouri.' The discovery resulted from exploratory drilling in the vicinity of a magnetic anomaly or high. A deep hole, drilled into the Pre-Cambrian porphyry to determine the cause of anomaly, penetrated a magnetite-rich deposit.' The magnetic surveys that revealed the anomalies were begun by the Missouri Geological Survey in 1929 and were continued into the 1930's.3 At that time an area in Crawford County known as the Bourbon magnetic high was located. A second anomaly, similar in size and intensity, was discovered about six miles to the northeast in Franklin County in the vicinity of Sullivan. This area was not studied in detail until the early 1940's. To determine the cause of the anomalies, the USBM drilled four holes at the Bourbon site in 1943-1944.4 The most productive hole penetrated four mineralized zones having a total thickness of 127.5 ft. These zones contained magnetite in rhyo-lite porphyry at depths between 1600 and 2000 ft. Additional iron ore below 2000 ft was considered a possibility at the time. Some years later, the USGS, Missouri Geological Survey, St. Joseph Lead Co., and others cooperated in sponsoring an aeromagnetic survey of the Bourbon-Sullivan area which resulted in discovery of a third high, about eight miles distant at Pea Ridge. This deposit, with an estimated reserve of 50 to 100 million tons of iron ore, is considered the main find, although other orebodies in the same general area are of definite interest.' When this discovery was announced, St. Joseph Lead also revealed plans to exploit the deposit jointly with Bethlehem Steel Co. A new corporation, Meramec Mining Co., was established, and 1962 was set as the target date for completing physical plant.2 Annual production goals of 2 million tons of beneficiated and pelletized iron ore were proposed. Preparation of the site for shaft sinking was begun in June 1957. Orebody and Character of Samples: Drilling to date has established the Pea Ridge orebody as a crescent-shaped pipe about 300 ft wide. Some drillholes were bottomed at 2800 ft in very good iron ore.6 he orebody is predominantly magnetite, although hematite and a mixture of hematite and magnetite occur in the upper areas and peripheral walls. Drillhole cores representing ore from different areas and depths of mineralization were grouped into composites for testing. These composites consisted of: 1) medium grade mixed hematite and magnetite, 2) high grade mixed hematite and magnetite, and 3) high grade magnetite. The medium grade ore came from three holes with depths of i350 to 1900 ft. The high grade mixed hematite and magnetite cores came from the 1790 to 2812-ft interval of one hole. Partial chemical analyses of the three samples, given in Table 1, reveal that all are excessively high in phosphorus and sulfur, and sample 1 is

Jan 1, 1960