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By John H. Schemel

Large Zircaloy-2 hammer or press forged bars did not exhibit the uniform excellent corrosion resistance to steam normally expected of the alloy in wrought form. Weight gains of coupons cut from forged bars were 40 to 60 mg per sq dm compared to the 28 mg per sq dm obtained on hot rolled sheet and strip. The mechanism of this corrosion and its relation to microstructure is discussed. After these observations, a solution heat treatment followed by rapid cooling was tried on coupons from eight heats of forged bars. The basket-weave structure resulting from the quench did not show agglomeration of intermetallic compounds. These coupons showed uniformly good corrosion resistance with weight gains very close to the 28 mg per sq dm expected of Zircaloy-2 after the 14-day steam test. A large forging was solution heat-treated and tested for structure, mechanical properties, and corrosion resistance. Corrosion resistance was excellent and uniform throughout the section. The normally anisotropic mechanical properties were changed to completely iso-tropic. Strength levels were raised while there was a small loss in ductility at the service temperature. Tempering of the quenched structure below the ß transus did not improve the ductility at the test temperature of 600°F. ZIRCALOY-2 end caps for nuclear fuel elements Zircaloy-2 is an alloy of 1.5 pct Sn, 0.1 pct Fe, 0.1 pct Cr, and 0.05 pct Ni, balance hafnium-free zirconium patented by Westing-house Electric Co. used in pressurized water reactors are machined from large hammer-finished forged bars. The bars forged for this application range from 3 1/2 in. to 7 1/2 in. squares. One of the tests used to evaluate Zircaloy-2 for this service is a 14-day exposure to very high-purity steam at 75oF and 1500 psi pressure. Normally wrought products will exhibit a black lustrous film of zirconium oxide after the test and show a weight gain of approximately 28 mg per sq dm. Coupons representing the forged bars exhibited a wide variety of corrosion results ranging from acceptable black coupons to some covered with white crystals of zirconium oxide that had weight gains of nearly 100 mg per sq dm. One of the most common effects was a white corrosion product that looked like a stain to some observers and the outline of massive metal grains to others. Very careful specimen preparation prior to the corrosion test did not affect the result and metallographic examination did not reveal a structural feature of a size and shape that would correlate with the corrosion product. In addition to the relatively poor corrosion resistance, the forged product was not as ductile as desired. The problem of obtaining uniform and more acceptable properties in these heavy-forged sections seemed to be a function of the microstructure of the metal and, in particular, the distribution of the intermetallic compounds. None of the four alloying elements however are appreciably soluble in a zirconium and are present as compounds which usually appear scattered randomly throughout the structure. Moudryl showed that the distribution of these intermetallic compounds in Zircaloy could be related to a ostringero corrosion failure. Grozier et al. discussed another structural defect that causes a similar corrosion effect. In this case, an elongated gas-void was determined to be the cause. M. L. Picklesimer held in his discussion of the Grozier paper that at least a part of the observed =stringersn were caused by the distribution of the intermetallic compounds. He proposed a

Jan 1, 1962

By Carmine D. &apos, Lemuel Tarshis, Joel Hirschhorn, Antonio

Vapor-deposited nickel films in the thickness range 700 to 4360A were tested in uniaxial tension utilizing a microtester designed specifically for this study. Contrary to the findings of some investigators, a definite thickness-strength relationship was observed below 3000 A with a four-to sevenfold increase in strength over that of bulk nickel. The films were characterized by high elastic strains and little plasticity. On the basis of these and other reported data, it is suggested that the high strength level in metal films is related to the manner in which they are produced. Vapor deposition, owing to its severe quenching effect. is believed to promote the formation of point defects which inhibit dislocation movement. IN recent years it has been reported that metals, when in the form of thin films, exhibit extraordinarily high strengths. The data published to date have been primarily concerned with silver and gold because of the ease with which these metals can be vapor-deposited and their high resistance to surface oxidation. Previous investigations into the mechanical strength of thin films has uncovered an apparent dichotomy of view on film behavior. Beams and his co-workers have reported a definite dependence of strength on film thickness. Other workers,"-'6 however, in separate studies on the strength of poly crystalline and single-crystal films have found no such thickness-strength relationship. The study of thin-film strength is extremely difficult because of the many variables associated with film preparation, handling, and testing. Moreover, the manner of test employed by different investigators has varied quite radically, ranging from simple uniaxial-tension to biaxial-bulge testing. The work reported herein was conducted in order to determine the mechanical behavior of a structural metal when in the form of a vacuum-deposited thin film and to gain some insight into the reasons for the high strengths exhibited by metals having such a con- figuration. In this study a method of test was chosen which would yield results which are easily interpreted and lend themselves to comparison with properties of the same material in bulk form. Moreover, specimen-preparation parameters and film-handling techniques are set forth so that other investigators can properly compare their findings with ours. EXPERIMENTAL PROCEDURES A) Film Preparation. Vacuum deposition was performed in an 18 by 30 in. bell jar using a standard New York Air Brake Co. vacuum station with a 6-in. oil-diffusion pump. Before evaporation the system was pumped to a pressure of less than 2 x 10"5 torr. A shield was employed to protect the substrates from the emission of contaminants during the critical melting and outgassing of the evaporant. The source consisted of from one to six filaments (0.020 in. diam). The length of each filament was about 5 in. and was placed 6 to 8 1/2 in. from the substrate in such a manner that the substrate face was at 90 deg incidence with respect to the evaporant beam. The temperature of the source was 2000"~ during evaporation. The films were deposited onto a substrate arrangement which was composed of four basic components, that is: 1) a 3 by 1 in. glass slide; 2) a 22-mm sq micro cover glass on 1); 3) a 22 by 50 mm micro cover glass coated with collodion on 1); and 4) a stainless-steel sheet mask containing twelve rectangular openings of 1-mm and 2-mm widths and lengths of 5 mm laid over 3). Thus, test specimens of 1-mm and 2-mm widths are deposited onto a collodion substrate which precludes epitaxial effects in the specirnens. 17-le The rectangular cover glass and square cover glass were positioned in such a manner that a strip of film would be evaporated along the length and across the width of the large glass slide. This boundary of evaporated film was used to determine the film thickness by multiple-beam interferometry. The square cover glass was used for X-ray and chemical analyses. Thickness of the deposits was varied by changing the number of filaments used (one to six). The duration of evaporation was 30 sec for each filamgnt which resulted in a deposition rate of 8 to 20A per sec. Evaporations were all performed at room temperature; however, radiant heat from the source raised the substrate temperature to 40" to 60°C. All substrates were cleaned by ultrasonic agitation in a solution of spectranalyzed isopropyl alcohol. Collodion was deposited on the micro cover glass by immersion in solution of collodion in amyl acetate. B) Thickness Control and Measurement. Film

Jan 1, 1963

By Joel Bradley, Kot Unrug, Kyle Perry, Mark Klimek

"A West Kentucky mine operation in No. 11 seam, encountered floor heave, due to the localized increase of the thickness in the fireclay mine floor. Floor heave has overridden seals installed in two mined out panels. The third seal’s location was planned for isolating that area from the Mains. A plan of support has been developed to prevent repetition of the floor heave and related problems outby the seals. The applied ground control measures were successful. An attempt of a 3D numerical modeling was made, that it would match the observed behavior of the mine floor and could be used as a design tool in similar conditions. In the paper described are sequence of events, a mitigation ground control system applied, and the first stage of numerical modeling.IntroductionWest Kentucky coal mine operation in the No. 11 seam has approached an area where the fireclay floor thickness increases to about 5 ft. Delayed floor heave has been observed. In the first sealed area (Panel KK), floor heave has overridden the seals and has begun to progress outby. This could not be allowed at the new seal location PP described in this paper because such override would affect the mains. Previously numbers of standing support systems were used, but their effectiveness was not satisfactory as is illustrated in Figure 1, Figure 2, and Figure 3. Therefore, with sealing of the next area approaching, mine management was looking for a better solution that would eliminate the problem. In that stage, our team looked closely at the geology and strength of roof and floor, as well as the characteristics of ground failure around the seal at the first site. The technical literature has been searched as well Haramy et al. 1986 Hsiung and Peng 1987, Peng 2008, Santos & Bieniawski 1987. This led to the identification of factors responsible for the ground failure. It has been determined that increased fireclay thickness from typically less than 2 ft to about 5 ft, when compared to the previously mined parts of reserve, and the strong limestone strata just above the immediate roof composed of 2 ft of black shale, created conditions of failure. It appeared that, when pillars were punching into the mine floor, the strong roof did not break but was bending, exerting high pressure on outby pillars, causing rib failures and closure. This is a similar mechanism to a strong roof affecting longwall shields. The remedial action planned was additional support."

Jan 1, 2015

By W. J. Garmoe

Distinct areas of mineralized and non-mineralized brecciated rock are found in the Ontario Unit of the United Park City Mines. These breccias contain an appreciable fraction of the present ore reserves and probably will become even more important as deeper levels are developed. Three ages of breccia have been recognized in the Ontario Mine to date: ( 1) Pre-mineral breccia that may or may not be mineralized, (2) early post-mineral breccia with mineralized fragments, and ( 3) late post-mineral breccia that is younger than and cuts through ( 1) and (2) above. These age ,relations are known from detailed underground mapping. Locally, both pre-mineral and post-mineral breccias have sufficient lead-zinc-silver content, either as matrix or fragmental material, to constitute ore. Ore-grade breccias occur throughout a vertical range of over 1000 feet and are found in both the Humbug and Deseret formations from above the Ontario 1300 level to below the Ontario 2200 level. The mineralized breccias, which are difficult to distinguish on mine maps, are not well delineated by underground workings, and their horizontal outline and vertical extent are not known beyond the immediate mining areas. The greatest amount of truly mineralized breccia is exposed on the 2200 level.

Jan 1, 1968

By J. P. Timmons, J. H. Moran, G. K. Miller, M. A. Coufleau

A prototype equipment has been designed and built for the digital recording of well logs on magnetic tape at the same time that the regular film recording is made. The format of the digital tape produced is such that it can be used directly at the input of the ZBM 704, 7090 or other models of ZBM computers which accept digital magnetic tape. This apparatus has been used for the experimental field recording of dipmeter tape logs which were subsequently computed by means of an ZBM 704 or 7090. In this paper the equipment and the digital tape are described briefly, and their application to the computer-interpretation of dipmeter data is discussed. A principal element in the interpretation of the dipmeter log is the correlation of the three microresirtivity dipmeter curves to determine the depth displacements between them. Several correlation methods for computer use are considered, with particular attention to their sensitivity to error and their consumption of computer time. The tape data were used to compute information content of the dipmeter microresistivity curves in terms of their frequency spectra. The results show that the sampling rate used in recording the digital information is quite adequate and illustrate a use of the digital tape in evaluating the characteristics of new tools. Some examples of field results are shown. It can be foreseen that, when digital tape recording becomes available for general field use, a whole new realm of possibilities will be opened up for the processing of other well logs through computations, which hitherto were not feasible because they were too laborious and time-con.sunzing. INTRODUCTION The last few years have seen a revolution in the design and production of data-processing equipment. Stored-pro-gram digital computers have progressed from a research curiosity to the basis of a major industry. There are now hundreds of such machines in daily use in the United States. With the acceptance of a technique that was, in fact, already clearly described by John von Neumann in 1945, the last decade has seen great strides in the development'of components, reliability, programming systems and, most spectacularly, in the sheer number of machines built and in use. In 1957 there were enough digital computers available to the oil industry to justify the suggestion that it would be worthwhile to investigate the possibility of using these machines in processing well log data.' The first result of this investigation was the appearance of what may be referred to as the input-output bottleneck. Well logs are customarily recorded on film. To get these data into a machine required then (and still does): a time-consuming semi-automatic reading of the film; conversion of the log data to digital form; and recording these digital data in some medium acceptable for computer input, such as cards, magnetic tape, or punched paper tape. However, the recording, reading, and re-recording could only result in deterioration of the data. Therefore, it was concluded that the fist step should be the development of a new, more direct recording technique supplemental to the film recording, which would provide easy access to the digital computer. There are many solutions to the problem of recording log data in an easily recoverable form. After careful consideration it was decided to adopt the boldest solution which, it was felt, was also the most elegant. It was decided to record well logs directly, in the field, on magnetic tape in such a way that this tape could be used without further modification as an input to the IBM 704 or 7090 computer. To realize practical field recording of magnetic tape logs, it became necessary to develop in a rather small package, an analog-to-digital converter, a tape recorder, and the necessary multiplexing and control circuits to allow the simultaneous recording of a multiplicity of logging signals. The magnetic tape recording was to be made simultaneously with the conventional logging operation in such a way as not to interfere with it. Along with the development of hardware, it was necessary to begin development of interpretation techniques and machine programs that would exploit the power of the digital computer. Here, again, there is a long list of possible applications. After much consideration it was decided to concentrate on the interpretation of the dipmeter log as a first application. It is the object of this paper to describe in some detail the developments sketched in the last three paragraphs.

By M. I. Brittan

Cyanide leaching of two West African gold ores was studied in multiple laboratories to provide process-engineering data for plant design purposes. Unusually slow leach extraction kinetics were observed using routine cyanide leach test procedures with pregnant solution monitoring. Carbon-in-leach testing, however, showed normal rapid kinetic behavior, with the same ultimate gold extractions being reached in much shorter elapsed times. Kinetic curves based on washed solids assays also exhibited more typical leaching characteristics. The results bore all the hallmarks of some equilibrium inhibition of the leaching. It was surmised that ore gangue components were responsible for weakly adsorbing gold-cyanide complex reaction product, thereby holding up soluble gold in the solid phase. The term ?preg-borrowing? was coined to describe the reversible phenomenon and distinguish it from irreversible preg-robbing. The two ores investigated may represent an extreme form of a more general effect that in most instances might pass unnoticed. The effect can lead to misinterpretation of kinetic leach curve data and underestimation of calculated ore head grades.

Jan 1, 2008

By R. O. Williams

It is shown that Ni3Al precipitates homogeneously from nickel-rich alwminum alloys as plates on the (100) planes. Prior to actual precipitation a process occurs which is believed to be one of increasing short-range order. After precipitating the Ni3Al plates enlarge through competitive growth. Discontinuous precipitation can occur simultaneously with the above processes. Recent ideas of the origin of precipitation strengthening appear adequate to explain the hardness changes. REMARKABLY little appears to be known about the precipitation process in Ni-Al alloys in spite of their technical importance. This investigation originated to supply additional information about precipitation in general, this system in particular. Information on the structures and kinetics have been obtained through the use of hardness, X-rays, microscopy, calorimetry, and resistivity on high-purity alloys. PROCEDURES Six alloys, Table I, were prepared by melting carbonyl nickel and high-purity aluminum in alumina crucibles in vacuum and casting into 1-in. graphite molds. All rods were homogenized at least once at 1300°C for 24 hr prior to swaging and this was repeated on the first three alloys after 75 pct reduction. Alloy 4 could be reduced only 10 pct at 1000°C (probably in two-phase field) prior to fracture but 1/4-in. samples quenched from 1100°C were readily reduced cold. Alloy 5 was reduced 15 pct cold but failed on the next pass while alloy 6 of essentially the same aluminum content failed inter-granularly without apparent flow up to 1000°C. The alloys were heated in hydrogen at the elevated temperatures and formed thin, coherent aluminum oxide coatings which provided excellent oxidation resistance at lower temperatures. However, freshly prepared surfaces showed considerably less resistance at 500"to 700°C in air and apparently resulted in internal oxidation. As a consequence, low-temperature agings were carried out in evacuated tubes. RESULTS The isothermal hardening behavior of these alloys at 500"and 565C is given in Figs. 1 and 2. These results were obtained from samples cold worked 75 pct, recrystallized at 1000°C (1100°C for the 7.8 pct Al) and quenched in water. This recrystallization was used to give smaller grain sizes so as to obtain more uniform hardness values and the points represent an average of five readings. The electrical resistivity was measured on 1/16-in. wires quenched from 1000°C during aging at 495°C to give Fig. 3. The energy release and its rate are given in Fig. 4 for the 6.9 pct Al alloy during aging around 500°C. Inasmuch as this was a single run, its accuracy is not known but certainly the general shape and magnitudes are correct. The method used to obtain these results is described elsewhere.' Data for the aging at 600°, 700°, and 800°C of these alloys cold worked 50 pct are given in Fig. 5. Supplementary information from microscopy and X-ray diffraction have been included to indicate recrystallization, discontinuous precipitation and the appearance of superlattice lines from the Ni3Al. The hardness of these alloys as annealed, aged, cold worked, and cold worked and aged is given vs composition in Fig. 6. Those samples which were isothermally aged, Figs. 1 and 2, were reaged at 532°C and at successively higher temperatures for the indicated times to give the data of Fig. 7. These results as well as certain others, support the idea that the level of hardness reached for temperatures above 600°C are equilibrium values more or less independent of path. This being the case, the breaks in the curves would be the complete solution of the Ni,Al. The electrical resistivity versus temperatures for some of these alloys, both aged and unaged, is given in Fig. 8 along with those data from heating slowly (10 deg per day) to high temperatures. Interesting points include the lowering of the Curie temperature (the change in slope), the lack of any indications of a solubility limit and the large temperature coefficient for the Ni3Al. A slight break for Ni3Al around 100C shows up but this is not a Curie temperature as Ni3Al is not ferromagnetic down to -190°C. Metallographically both the nickel-rich solid solution and the Ni3Al appear very much like pure nickel. Profuse twin boundaries are present both

Jan 1, 1960

By Derek. C. Arnold

In order to meet increasing demands for water supply in the next millennium, the Government of the Hong Kong Special Administrative Region have identified the need for a major new water treatment and transfer scheme. The project now under construction will comprise a 1.2 km long raw water aqueduct, a major new 0.25million cubic metres per day water treatment works (with space for upgrading to 1.2Mm³/d), a 12.1 km long treated water aqueduct and a primary service reservoir. In order to meet increasing demands for water supply in the next millennium, the Government of the Hong Kong Special Administrative Region have identified the need for a major new water treatment and transfer scheme. The project now under construction will comprise a 1.2 km long raw water aqueduct, a major new 0.25million cubic metres per day water treatment works (with space for upgrading to 1.2Mm³/d), a 12.1 km long treated water aqueduct and a primary service reservoir.

Jan 1, 1999

By Junfeng Shen

Geosynthetic clay liner (GCL) is a kind of waterproofing material used widely in engineering. The waterproof mechanism is understood in terms of bentonite particles becoming a water-obstructcolloid layer after they absorb water and swell. The swell stress, however, has not been determined directly till now. In our experiment, the swell stresses of the GCL under saturated and water absorbing conditions are measured directly using a custom-made instrument. The results show that (1) the instrument designed by the authors performed satisfactorily and the test results are reproducible; and, (2) The trend line of swell stress variation with time can be divided into three segments. The first segment is characterized by a quick increase of the swell force in the first 0?50 hours. The swell stress increases by 7.00×10-4?1.00×10-3 Mpa / hour. The second segment shows a slow increase of the swell stress during the 50?1730 hours. The swell force increases by 7.54×10-6?2.02×10-5Mpa / hour. The third segment is characterized by little variation in the swell stress after 1730 hours. In this segment, the average value of the swell stress measurements is 0.0719Mpa and the maximum value is 0.0729Mpa. It is suggested that the swell stress is mainly raised by water entering the pores among montmorillonite particles and the interstitial layers in individual montmorillonite crystals, leading to an increase in volume and manifested as a stress on the surrounding materials.

Jan 1, 2006

By H. H. Podgurski, H. E. Knechtel

Annealed Fe-Al alloys do not react readily to form AlN when held at 500ºC in NH3-H2 gas mixtures, but do so upon the introduction of dislocatims. Nuclea-tion of the nitride phase occurs on dislocation sites. In turn, the growth of the aluminum nitride particles causes the ferrite phase to yield plastically, generating more dislocations for the nucleation process. The nitride phase extracted from an Fe-2 pct A1 alloy nitrogenated at 500°C was identified as stoichio-metric aluminum nitride with a hexagonal crystal lattice. THIS investigation reveals the role that dislocations play in initiating and sustaining the nitriding reaction in Fe-A1 alloys. As early as 1931 the work of Meyer and Hobrock1 suggested that the initiation of the nitriding reaction could involve a nucleation controlled process. Recently Bohnenka2 depicted the gas-phase nitriding process below 600°C as one of mixed control limited by nitrogen penetration through the surface, by nitrogen diffusion, by aluminum diffusion, and by nucleation of the nitride phase, Fig. l(a). In our research in a comparable alloy (0.57 pct Al) at 575ºC, we have observed a nitrogenation which we feel is better described by Fig. l(b). In the case of a 2 pct-A1 alloy partially nitrided at 500°C we propose the profiles shown in Fig. l(c). For a complete and accurate description of the process, a concentration profile of the dislocation density in the test specimen would be needed. EXPERIMENTAL Nitrogenization was conducted between 500" and 575°C in a variety of NH3-H2 gas mixtures on three Fe-A1 alloys: 1) zone-refined iron + 0.16 i 0.2 pct Al—levita-tion melt, 2) zone-refined iron + 0.57 0.02 pct Al— levitation melt, 3) plastiron + 2 pct Al—melted by induction heating. To demonstrate the effect of dislocations on reactivity, both cold-worked and annealed samples were investigated. All nitrogenation rate studies were conducted gravimetrically with a gold-plated invar balance4 contained in a gas-flow system. To avoid contamination of the specimens in the reaction zone of the system, the reaction chamber was constructed of high-purity dense alumina. The activity of nitrogen was fixed by specific NH3-H2 gas mixtures whose compositions were continually monitored by calibrated thermal conductivity gages and checked by chemical analysis. Variations of ± 0.1 pct NH3 could easily be detected by both methods. Throughout this paper the activity of nitrogen is defined as PN3 /PH23/2 , where PNH3, and Ph2 are partial pressures in atmospheres. Electron transmission, density measurements, and chemical analyses were made on specimens before and after nitrogenating in order to reveal structural and chemical changes. Similar studies as well as X-ray diffraction studies were conducted on nitride extractions from the nitrogenated 2 pct-A1 alloy. These extractions were obtained by the use of an anhydrous bromine-methyl acetate solution which dissolves the iron and leaves the insoluble nitrides as a residue. Hardness profiles were obtained on cross-sections of partially nitrided specimens to demonstrate the extent of nitriding through the thickness of the specimens. RESULTS AND DISCUSSION The nitrogen activity in the NH3-H2, atmospheres was never allowed to reach a level capable of producing iron nitride (Fe4N). Hence, the term nitriding as used in this paper refers only to the formation of aluminum nitride whereas nitrogenation refers to the total uptake of nitrogen regardless of how it is distributed throughout the alloy. The weight increases observed during the initial stage of a nitrogenating treatment are due primarily to the solution of nitrogen in the ferrite phase, particularly when starting with annealed specimens. Because this initial nitrogenation rate in the case of the 0.57 pct A1 alloy, see Figs. 2 and 3(a), was most rapid the weight change that occurred thereafter might be attributed to the nitriding reaction with the exception of a small weight increment due to the irreversible pickup of oxygen by aluminum. The oxygen (<70 ppm) came from traces of H2O and 0, in the hydrogen and ammonia gases. On the basis of discrepancies between total weight increase and the increase in the nitrogen content of the sample as determined by chemical analysis, it was estimated and later established by activation analysis, that as much as 200 ppm of oxygen were taken up by a fully nitrided Fe-0.57 pct A1 specimen at 575°C. Most of the oxygen could have been picked up from the nitriding atmosphere on the surface of the samples during cooling to room temperature. Even 50 ppm of water in the gas phase will become oxidizing to iron before the sample has cooled to room temperature. The lack of reactivity* of these alloys in the annealed

Jan 1, 1970

By Stuart Brashear, Steven Bell, Donald Smith

This Limestone site required the removal of a bench containing high quality rock located in close proximity to sump and pit pump. The sump and associated electrical box were below the bench that needed to be blasted and located only 70 feet (21.33 meters) away. It was proposed to blast material needed without moving or disturbing the pump or electrical junction box. The area to be blasted was flown with a drone and stitched into a 3D orthographic model. Proprietary blast design software was used to create a blast design solution that would allow for a series of blasts that would provide required fragmentation and muckpile development and not disturb the sump or electrical box. A drill patten was created for each blast using the burden heat map and 3D face profiles created from the 3D orthographic model. Timing contour lines and muckpile displacement predictions were used to create timing sequences for proper face movement away from the sump and electrical box. The shot designs were laid out with a GPS system to insure hole location accuracy. All holes were remeasured for proper face burdens using the 3D model and Boretrak®® data collected after drilling operations were completed to insure hole locations were within face burden tolerances as derived by the burden heat map. After each drill log was reviewed, a charge loading scenario was created for each hole. The timing design created, for each shot and was reviewed and adjusted as needed with regard to desired muckpile displacement away from the sump and electrical box. Blasting results were satisfactory, with fragmentation and muckpile development consistent with the calculated throw from the blast model, with no rock hitting the pump or box and no rocks striking the sump water and disturbing the pump motor.

Jan 21, 2025

By Toshiyuki Hiranami, Keizou Uematsu

"AbstractionSpray freeze drying method was proposed to fabricate ·ceramic: powder granules for high performance ceramics. Obtained granules and their green compact were characterized by SEM and immersion liquid technique. Granules have a spherical shape of about l 00 µ m in diameter, ·and no binder segregation was found with the immersion liquid technique. The granule had sufficient strength for usual handling, while the strength was much lower than that of granules prepared by conventional spray-drying. Unlike the conventional green compacts, granule interfaces were not observed in the green body made with the spray-freeze dried granules.IntroductionIt was well kn9wn that spray-dried granules have binder segregation layer on their surfaces. The trace of this segregation layer survives in the later processes i.e. compaction and firing, and behaves as fracture origins in the sintered bodies. The origin of this problem is solvent drying process from slurry droplet surface in spray drying system. As alternative method we propose a freeze drying process for granule preparation, in which solvent is remove by sublimation though the process, binder segregation can be prevent.In this study, we applied spray freeze drying method to fabricate alumina powder granule, and characterized the obtained granules, and their green compacts and sintered bodies.Experimental procedure.Alumina powders and additives aqueous solution were mixed for 24 hours using a ball mill to· prepare the slurry for spray freeze drying. .The solid loading of the slurry was 30vo1%. Deflocculant (PAA-NH4 type) was added 0.5mass%, and the binder (PVA). 1.0, 2.0 or 3.0mass% of alumina. Obtained slurries were sprayed into liquid nitrogen to make frozen granules. The granules were freeze dried.For the characterization of the granules, immersion, liquid technique observation and single granule strength test were carried out. Strength of individual granule wa5 measured using micro compression testing machine (MCTM, SHIMADZU). Granule strength was calculated from following equation."

Jan 1, 2000

By C. J. Lewis

Commercial lime plants in the United States and Canada are depicted in Figure 1. The sources of lime available to Nevada and bordering states are rather few. Chemstar Lime has plants at Apex, Nevada and Nelsen, Arizona. Continental Lime, Inc. has plants at Tacoma, Washington, Townsend, Montana and Delta, Utah. However, Continental Lime has a new rotary kiln installation underway near Wendover, Nevada. This plant is scheduled to come on stream in the fall of 1989.

Jan 1, 1989

By C. E. Corson

Discussion of the Paper by C. E. Corson, which was presented at the London Meeting, July, 1906. (See Bi-Monthly Bulletin, No. 11, September, pp. 725 to 742.) ALBERT SAUVEUR, Cambridge, Mass. (communication to the Secretary*) :-On close examination I think it will be found that the evidence by which Mr. Corson claims to have shown the inaccuracy of a statement I made a few years ago, is far from convincing. This statement is:- "Hot work, as such, has no influence upon the structure of the metal. Indirectly, however, by retarding crystallization until a lower temperature is reached, it may influence its structure most decidedly ; but the same results could be accomplished by heat-treatment alone, i.e., by reheating the unworked metal to the temperature from which the unworked piece was allowed to cool undisturbedly." The Metallographist, vol. ii., p. 267 (under the head of "Changes of Structure Brought About by Work"). Of the two pieces of steel tested by Mr. Corson, one was worked and finished at a " cherry red " and &apos; the other reheated to a cherry red, and the assumption was made that, in both cases the temperature was about 715° C., although no pyrometric device was used to record it. The possibility of considerable difference in temperature alone is so great as to invalidate Mr. Corson&apos;s inferences. The apparent temperature, moreover, is so close to the critical point of the steel as to render any conclusion very hazardous. -It is not at all evident that the annealed piece was reheated past its critical point, in which case reheating would have had practically no effect. Nor is it evident that the piece forged and finished at a cherry red was not actually below the critical point, which, by definition, I have called cold worked. Mr. Corson&apos;s explanation that " Under proper reheating, on the other hand, the steel becomes a solid solution from which crystals of approximate homogeneity and uniform size may separate" will not be readily understood.

Jan 1, 1907

By C. Atkinson

A method is described for the numerical solution of the diffusion equation with a composition-dependent diffusion coefficient and applied to the radial growth of a cylinder; the radial growth of a sphere, and the symmetric growth of an ellipsoid. Sample applications of the method are made to the growth of particles of proeutectoid ferrite into austenite. RECENTLY&apos; we described a method for numerical solution of the diffusion equation with a composition-dependent diffusion coefficient for the case of the growth of a planar interface. In this paper we extend this method to describe the radial growth of a cylinder, the radial growth of a sphere, and the symmetric growth of an ellipsoid. In the latter case, limiting values of the axial ratios of the ellipsoid reduces the problem to one of a cylinder, a sphere, or a plane depending on the axial ratio. A check on these limiting values is made in the results section. In all of these cases we consider growth from zero size. A natural consequence of this assumption as applied to the sphere, for example, is that the radius of the sphere is proportional to the square root of the time. This is consistent with the condition that the radius is zero initially, i.e., grows from zero size. It may be argued that it is more realistic to consider particles which grow from a nucleus of finite initial size; even in this case the analysis of this paper is likely to be applicable. This can be seen if a comparison is made of the work of Cable and Evans,2 who consider a sphere of initially finite size growing by diffusion in a matrix with a constant diffusion coefficient, with the results of Scriven3 for growth from zero size. This comparison shows that the rates of growth in each case differ trivially by the time the particle has grown to about five times its initial size." This investigation is a generalization of those of Zener,4 Ham,5 and Horvay and cahn6 to the situation often encountered experimentally, in which the diffusion coefficient varies with concentration. First let us consider each of the cases separately. I) GROWTH OF SPHERICAL PARTICLES FROM ZERO SIZE In this case the differential equation in the matrix depends only on R, the radius in spherical coordinates, and can be written: ? 1 <^\ ^13D . , dt U\dRz + R 3Rj + dR dR [ J where C is the composition, t is the time, and D is the diffusion coefficient which depends on c. The boundary conditions will be: c = c, at the moving interface in the matrix, c = c, at infinity in the matrix (and at t = 0, everywhere in the matrix), c = X, is the composition in the spherical particle. Each of the above compositions is assumed constant. In addition there is the flu condition at the moving interface which can be written: , dR0 ~/3c dt \dR/H =Ra where R,, which is a function of t, is the position of the moving interface. We make the substitution q = RI~ in [I] reducing this equation to: & - m - *ws) »i where we have written D = D,F(c) or simply D,F, and Do = D(c,). Thus F[c(q0)] = 1 where q, = ~,/a is the value of the dimensionless parameter q evaluated at the interface. Multiplying Eq. [2] by dq/dc and integrating, we find: where the lower limit of the integral has been chosen so that dc/dq — 0 as c — c,, thereby satisfying the boundary condition at infinity. We require, then, to solve Eq. [3] subject to the condition c = c, when q = q, (this follows from putting R = R, at the interface) together with the flux condition which can be rewritten in terms of q as: Eqs. [3] and [4] together with the condition c = c, at q = q0 enable us to find 77, and the concentration profile c = c(q). Numerical Method. We treat Eq. [3] in the same way as we did the corresponding equation for the planar interface problem&apos; i.e., by dividing the interval c, to c, into n equal steps so that: cr = ca -rbc [5] where r takes the values 0, 1, ... n and we call no,, q1, ... nn the values of n corresponding to the compositions c,, c,, ... c,.

Jan 1, 1970

By A. Szirmae, Hsun Hu

The early stages of recrystallization in a 70 pct cold-rolled Si-Fe crystal of the (110) (0011) orientation were studied with a Siemens electron microscope. Orientation studies based on electron-diffractzotz. patterns confirm the results of previous texture analysis. The driving energy for recrystallizatior and the critical radius for growth were calculated from the dislocation energy and the energy of the subgrain bourzdaries, and it was found consistent with the observed size of the recrystallized grains. The recrystallization characteristics of crystals with different initial orientations are discussed. The recrystallization of cold-rolled (110)[001] crystals of Si-Fe has been widely studied by various investigators.1-4 Their results on both deformation and annealing textures are in good agreement. The rolling texture after 70 pct reduction consists mainly of two crystallographically equivalent (111) [112] type textures and a minor component of the (100) [011] type. The latter is derived from the deformation twins, or Neumann bands, which are formed during the early stages of deformation and later rotate to the (100) [011] orientation upon further rolling reduction. Between the two main (111) [112] type textures, there is some orientation spread, because of which very low intensity areas appear in the pole figure. If these very low intensity areas are considered to be a very weak component in the texture, then a (110) [ 001 ] orientation may be assigned to them. When this rolled crystal is annealed at a sufficiently high temperature for recrystallization, the texture returns to a simple (110) [001]. The purpose of the present investigation was primarily to seek a better understanding of the recrystallization process by using the electron transmission technique. The (110) [0011 type of crystal was selected because orientation data for it are well known from previous studies with conventional techniques. Direct observations on the recrystallization of such a crystal have also been made by using a hot-stage inside the electron microscope, and the results will be reported in another paper. MATERIAL AND METHOD A single-crystal strip of the (110) [001] orientation was prepared from a commercial grade 3 pct Si-Fe alloy by the strain-anneal technique.= The strip was approximately 0.014 in. thick, and was rolled 70 pct at room temperature to a thickness of 0.004 in. Specimens were cut from the rolled strip and were annealed in a purified hydrogen or argon atmosphere. They were then electrolytically polished in a chromic-acetic acid solution to very thin foils. Best results were found by polishing first between two narrowly spaced flat cathodes with the specimen edges coated with acid-resisting paint, followed by polishing between two pointed electrodes until a hole appeared in the center as described by Bollmann.6 It was found that a thin transparent film always formed along the thin edges of the polished specimen. This film was then removed by rinsing the specimen very briefly in a solution of alcohol with a few drops of HF or HCl. RESULTS AND DISCUSSION 1) The Deformed Crystal. From the electron-diffraction patterns taken at various areas of an as-rolled specimen, the texture components as deduced - from ordinary pole-figure analysis were confirmed. Over most of the areas where orientation was examined, a (111) pattern with a [112] direction parallel to the rolling direction was obtained. This corresponds to the main deformation texture of the (111) [112] type. In a few areas the diffraction pattern was (100) [Oil], corresponding to the minor-texture component derived from the Neumann bands. The (110) [001] orientation, which corresponds to the very weak intensity area in the pole figure, was found infrequently. A typical example of the deformed matrix having the (111) type main texture is shown in Fig. 1, where (a) is the microstructure and (b) is the diffraction pattern taken from that area. It was also frequently observed that in other areas more or less continuous rings of weaker intensity were superimposed on the simple (111) diffraction pattern, suggesting the presence of a wide range of additional orientations. Other evidence indicated that the recrystallization characteristics are different in these two different types of areas. The hot-stage observations which provide this evidence will be discussed in another paper. AS shown in Fig. l(a), numerous dislocation-free areas of very small size are embedded in the "clouds" of high-dislocation density. This indicates that the deformation of a single crystal, even after a rolling reduction of 70 pct, is far from uniform on a micro-

Jan 1, 1962

By Clinton C. Knox

To demonstrate the feasibility of unwatering ore reserves previously indicated in the virgin Melrose zinc-lead field in Kansas and Oklahoma, the Bureau of Mines .carried through a pumping project on the Park Walton farm in 1943. Due to success of the undertaking, attention was turned to exploration of other indicated ore deposits and promising areas within that part of the field benefited by .the drainage project, By the middle of 1946, four private mining companies were active in the field; and a considerable output of lead and zinc concentrate had been obtained from the new Park Walton mine and mill. The Park Walton Ore deposit was unwatered by installing deep-well pumps in three holes previously drilled by the lessee. Pumping was continued by the Bureau of Mines at a rate of 2,500 to 1,000 g.p.m. for about 6 months; during this time the local static artesian water-level, originally averaging 108 feet below the surface in the old drill holes was lowered to the 316-foot level, which justified beginning the initial shaft sinking and mine developments. In all, about 750,000,000 gallons of water was pumped by the Bureau of Mines at an operating cost averaging $0.024 per 1,000 gallons.

Jan 1, 1948

By Jei Y. Choi, P. G. Shewmon

The surfrrce-diffusion coefficients (DJ for Aulg8 on (100) and (111) surfaces of copper have been determined between 1050" and 780°C using a new avuzlysis imd experimental procedure. The results are: D, has also been determined fm cua4 at 870°C, and the values found are 4.5 times larger than those measured by the grain boundary grooving technique for the same surface orientations. This difference is felt to result from the approximate nature of the mathematical solution used in the present work. Attempts to measure D, for silver on copper and silver surfaces indicated a means of matter transport different from surface diffision was dominant in moving tracer from the source out over the surface. Cnlculations and experiment both indicate that this is the flow of silver through the vapor phase which completely masks the much smaller flow due to surface diffusion. The previous self-difhsion studies of D, for silver and copper are discussed in terms of our own analysis and found to yield values of D, factors of lo5 or more greater than those found by the grain boundary grooving tech -nique. UNTIL about 5 years ago it was widely believed that the activation energy for surface diffusion, AH, , was less than that for grain boundary diffusion, AHb,, which in turn was less than that for diffusion through the lattice, AHz.&apos; This was concluded from various evidence that D,> Db>Dl, and one tracer study of D, for silver on silver from which AH, was inferred.2 In 1959 Mullins and Shewmon demonstrated that D, could be determined from the kinetics of the growth of grain-boundary grooves.3 Using this procedure, Gjostein measured D, on copper between 800" and 1050°C and found that the activation energy was roughly equal to AHl .4 Subsequent work on copper,5" silver,&apos;,&apos; and goldg between the melting temperature T, and 0.87 T, confirmed that AH, as determined using the grain boundary grooving or scratch-relaxation technique was equal to or greater than AHz. During the same period, Drew and Pye again determined AH, for silver on silver using a tracer techniquelo and a mathematical solution similar to that of Nicker son and arker.&apos; Though the values of D, Drew and Pye measured at any given temperature were about 200 times smaller than those reported by Nickerson and Parker, they again found a low activation energy of about 10 kcal, or about one fifth that found at the higher temperatures with the mass transport technique. A distinguishing characteristic of these two previous tracer studies is that they have worked at low temperatures (-1/2 T,) where they felt volume diffusion was negligible and then analyzed these data as if all tracer atoms leaving the source flowed out into and remained in a homogeneous high-diffusivity surface layer of undefined thickness. This is totally different from the model used in the mass-transport studies or the studies of grain boundary diffusion, which assume the high-diffusivity surface layer to be only a few angstroms thick. If this latter model is applied to the earlier tracer studies, it is shown that the tracer has really pe!etrated into the lattice a mean distance of 1000A. Thus the tracer distribution observed after an anneal is thought to be due to the combined effects of surface and volume diffusion. Independent of the relative validity of the two models, it seems evident to us that any comparison of the values of D, as determined in these two ways is meaningless and misleading, since the values of D, and AH, obtained in these two ways would be totally different for the same physical distributions of tracer. Once the fundamental difference in the approaches of the two techniques is established, we are faced with the question of which model better approximates physical reality. Here all the evidence seems to be on the side of the &apos;&apos;thin surface layer" analysis. In fact, the authors of Refs. 2 and 9 do not argue for the "thick-layer model" we have described; they simply invoke it through the equation they use to calculate D, . The primary evidence for the thin-film approach is: a) grain boundary grooves and scratches widen in proportion to tU4 and Mullins&apos; rigorous analysis shows that this is only valid for a surface layer which is quite thin relative to the width of the groove;11 b) all accepted or seriously discussed models of solid-vapor interfaces and high-angle grain boundaries assume that the disturbed region of the interface is at most a few a0 thick. With the above in mind, it was desirable to determine D, using a radioactive tracer and a "thin-

Jan 1, 1964

By D. V. Edmonds, C. J. Beevers

Tensile tests were carried out on a! titanium containing 850, 1250, and 2700 ppm 0, and up to -500 ppm H. The tests were performed at -196", -78", 20°, 150°, and 300°C at a strain rate of -1.0 x 10??3 sec-1. Increasing oxygen content, increasing grain size, and decreasing test temperature resulted in enhanced embrittlement of the a titanium by the hydrogen additions. Metallographic observations showed that this can be correlated with the influence of these parameters on the introduction of cracks into the a! titanium by fracture of titanium hydride precipitates. CRAIGHEAD et al.1 reported that the hydrogen content normally found in commercial-purity a! titanium (60 to 100 ppm) was sufficient to cause a substantial lowering of the impact strength, and they attributed this embrittling effect of hydrogen to the precipitation of titanium hydride. Lenning et al.&apos; found that in commercial-purity a titanium there is an almost complete loss of impact strength at about 200 pprn H, which is approximately half the value needed to eliminate the impact strength of high-purity a titanium. They also showed that the presence of 3000 ppm hydrogen reduces the room-temperature tensile ductility of commercial-purity material to a value of the order of 10 pct; the corresponding hydrogen concentration for high-purity titanium is over 9000 ppm. It thus appears that the detrimental effect of hydrogen on the mechanical properties of commercial-purity titanium becomes evident at much lower hydrogen contents than for high-purity titanium. The main difference between the two types of a titanium might be expected to be the higher level of interstitial impurity in the commercial-purity grade. Jaffee et a1.3 studied the influence of temperature and strain rate on the hydrogen embrittlement of high-purity and commercial-purity ! titanium. In general, the behavior was the same for both materials; embrittlement was enhanced by decreasing temperature and increasing strain rate. Recent results from tests on commercial-purity a titanium containing 850 ppm O and varying amounts of hydrogen up to -500 ppm showed that the degree of embrittlement by hydrogen is intimately related to the fracture characteristics of titanium hydride precipitates.4 The present paper considers the interrelationship between the mechanical properties and micro-structural features of commercial-purity a! titanium containing 850, 1250, and 2700 ppm 0 and varying amounts of hydrogen up to -500 ppm. 1. EXPERIMENTAL PROCEDURE Three types of commercial-purity titanium supplied by IMI* were used in the investigation, and for the *Address: Witton, Birmingham 6, United Kingdom. purpose of this paper are designated Ti 115, Ti 130, and Ti 160. The principal impurity elements are given in Table I. The material was received in the form of 12.7 mm diam bars having a fully recrystallized structure. Tensile specimens with a round cross-section of 4.5 mm diam and a gage length of 15.2 mm were machined from the bars. In order to develop the same grain size (mean linear intercept of grain boundaries) in each of the three types the specimens were annealed under a dynamic vacuum of <10?5 mm Hg, Table 11. Specimen hydriding was carried out in a modified Sieverts apparatus;&apos; hydrogen was taken into solution at 450°C and after holding the specimens at this temperature for 24 hr they were furnace-cooled to room temperature at an average rate of -100 C deg per hr. By this method nominal hydrogen contents of 0, 50, 100, 250, and 500 ppm were introduced into specimens of Ti 115, Ti 130, and Ti 160 (100 ppm (wt) -0.5 at. pct). The actual hydrogen contents were calculated from the weight differences obtained by weighing the specimens before and after the hydriding treatment. Tensile tests were carried out at temperatures of -196", -78", 20°, 150°, and 300°C on a 10,000 kg In-stron machine at a nominal strain rate of -1.0 x 10-3 sec-1. Fractured specimens were sectioned in planes parallel to the tensile axis, mechanically polished to 0.25 µm grade of diamond paste, and then attack polished using a solution containing by volume 99 parts H2O, 1 part HF, and 1 part HNO3. Although the latter treatment unavoidably opened out cracks and voids visible after mechanical polishing, it did reveal the grain structure, titanium hydride morphology, and deformation twinning structure.

Jan 1, 1970

By J. R. Cahoon, H. W. Paxton

The mechanical properties of unidirectionally solidified A1(rich)-Mg and A1(rich)-Cu castings containing up to 15 wt pct solute have been determined with re -spect to the volume fraction of interdendritic eutectic. Pioperties were determined in the directions pumllel and Perpendicular to that of solidification; the volume fraction of eutectic was varied between the "as-cast" and equilibrizcm amounts by approperiate heat treatment following solidification. The principles of fiber strengthened composites and dispersion strengthened materials are adapted to explain the mechanical properties of these castings. It is generally accepted that castings often have inferior mechanical properties when con~pared to wrought products. However, there is little quantitative data available concerning the factors which make apparently sound castings weak and/or brittle. The relative ease and inexpensiveness of the casting process have always been attractive and, therefore, an understanding of the factors which contribute to the mechanical properties of castings would seem desirable. Such an understanding may lead to an improvement in the mechanical properties to an extent where castings would become competitive in applications where presently only wrought products are considered to have the requisite properties. Such an understanding could also improve the reliability of present cast products. Much of the recent research on castings has centered about determining the extent of segregation in cast alloys. Macrosegregation, particularly inverse segregation, has been studied in some detail 1-8 and a considerable understanding of microsegregation has been obtained.9&apos;10 The effect of solidification rate on dendrite spacing and on the amount of interdendritic eutectic in binary alloys has been established, particularly for Al(rich)-Cu alloys.""0 However, the extension of these ideas to relate the amount of interdendritic eutectic, concentration gradients, micro-segregation, dendrite spacings, and so forth, to the rnechanical properties has been limited. Dean and spear" have related the mechanical properties of an Al-Si-Mg alloy, A356-T62, to the dendrite spacing and have shown that the mechanical properties improve with decreasing dendrite spacing. Passmore et al.12 have shown that annealing at high temperature improves the mechanical properties of Al(rich)-Cu al- loys and Archer and Kempf 13 have shown that an Al-1 pct Mg-1.75 pct Si alloy behaves in a similar manner. Ahearn and Quigley 14 have shown that high temperature homogenization also enhances the mechanical properties of an SAE 4330 steel. However, in the above investigations, no underlying reasons were suggested for the improvement in mechanical properties. The purpose of the present investigation is to relate the mechanical properties of castings to some of the solichfication variables and to derive some equations by which calculations of the mechanical properties may be attempted. In particular, the effect of the amount of interdendritic eutectic and the effect of stress direction with respect to that of solidification on the mechanical properties will be considered. The Al(rich)-Mg and Al(rich)-Cu binary alloy systems were chosen for study. The A1-Mg system was chosen because its constitutional relationships are such that large volunles of eutectic (up to 24 vol pct) may be obtained in the as-cast condition and then be completely dissolved by subsequent heat treatment at about 440°C. This allows a comprehensive study relating the mechanical properties of castings to the amount of interdendritic eutectic. Also the Al(rich)-Mg eutectic is almost a single phase 15 which should make the experimental results more amenable to theoretical interpretation and calculation. The A1-Cu system was chosen for study because of the large amount of related information available concerning segregation, dendrite spacing, and so forth. Unidirectionally solidified castings were used throughout the investigation so that the effect of solidification direction with respect to the direction of applied stress could be determined. THEORETICAL It is well known that upon solidification of binary alloy castings, the nonequilibrium amount of eutectic which forms is given by 10 where fe o is the weight fraction of eutectic, Cs is the solid solubility of solute at the eutectic temperature, k is the equilibrium partition coefficient, and C, is the average composition of the alloy. In the development of Eq. [I], it is assumed that the effects of inverse segregation and diffusion in the solid are negligible, and that no porosity is present. If the casting is homogenized at a high temperature for a long period of time, some (or all) of the eutectic is dissolved and the amount of eutectic for this "equilibrium" condition may be calculated directly from the constitutional diagram. By appropriate intermediate annealing, the

Jan 1, 1970