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By Luc Delaey, Horace Pops

The structure and morphology of thermoelastic and burst type martensitic phases that form upon cooling in Cu-Zn-Si p phase alloys have been studied by transmission electron microscopy. The martensitic phases are composed of a lamellar mixture of two close-packed structures with different stacking sequence, namely ABCBCACAB (orthorhombic) and ABC (fcc). Striations within thermoelastic martensite are most likely produced during interaction with impinging burst-type martensite and not as a consequence of secondary shears. In a study of the martensitic transformation in ternary Cu-Zn based 0 phase alloys1 the dependence of the martensitic transformation temperature, M,, with composition shows variations for elements within a constant valence subgroup and between different subgroups. Such variations are not reflected in a change in habit plane, which is approximately the same for each ternary alloy, namely in the vicinity of (2, 11, 12 Ip. The fact that the habit plane remained constant, despite large differences in M, temperature and electron concentration, suggested2 that the crystal structures of the martensitic phases could be nearly the same. Crystal structures of ternary Cu-Zn based martensites have been determined recently for alloys containing the three-valent elements gallium3, 4 and aluminm. The present studies have been made to examine the structures and morphology of the martensitic phase in ternary Cu-Zn based alloys containing a four-valent element, silicon. I) PROCEDURE Two alloys were prepared by melting and casting weighed quantities of the component high-purity metals in sealed quartz tubes under half an atmosphere of argon. They were subsequently remelted by levitation under a protective atmosphere of argon. After allowing for losses of zinc as determined by the difference in weight before and after casting, the compositions in atomic percent of both alloys were established to be Cu-33.5 Zn-1.8 Si and Cu-27 Zn-5.0 Si. These alloys were homogenized in the P-phase field for 2 days at 800" C. Bulk samples consisted of a martensite phase at room temperature, the M, temperature being approximately 30' and 200" for the 1.8 and the 5 pct Si alloys, respectively. Thin disks were cut from the ingots using a spark machine, and they were heated for 5 min at 800' and quenched into water in order to obtain martensite. These slices were thinned chemically at room temperature in a solution consisting of 40 parts HN03, 50 part H3PO4, and 10 parts HC1 and thinned further electrolytically by the Window technique, using a voltage of 15 to 25 v and a mixture of 1 part HN03 and 2 parts methanol, which was kept at a temperature near -30° c. Foils were examined by transmission electron microscopy using a Philips EM 200 electron microscope. 11) RESULTS AND DISCUSSION 1) Structure and Morphology. Fig. 1 shows the martensitic phase in the alloy containing 1.8 at. pct Si. This phase is composed of contiguous platelets, each containing striations. The direction of the striations changes at the boundary between individual platelets. These internal markings resemble the striations that are usually identified as stacking faults, as for example in Cu-A1 martensites6-a or the lamellar mixture of two close-packed phases in Cu-Zn-Ga marten-sites.3p '9 lo In the present alloys, selected-area diffraction experiments have been obtained in order to determine the nature of the striations. Figs. 2(a), (61, and (c) are electron diffraction patterns of an area inside a single martensite plate. Fig. 2(a) contains diffraction spots which correspond to two close-packed structures with different stacking sequences, namely ABCBCACAB (orthorhombic) and ABC (fcc). Spots belonging only to the fcc structure are indicated by arrows. By tilting the foil either the orthorhombic structure, Fig. 2(b), or the cubic structure shown in Fig. 2(c) may be obtained. When the foil is oriented so that only the diffraction spots of the orthorhornbic structure are present, bright-field illumination shows small lamellae, as seen in Fig. 3. In this figure the lamellae that belong to the fcc structure are bright bands inside the dark extinction contours of the orthorhombic structure. The boundaries of the lamellae are parallel to the basal planes of the orthorhombic structure and to the {Ill} planes of the cubic structure, the close-packed directions of both structures being parallel. The 5 pct Si alloy shows similar features as those described for the 1.8 at. pct Si alloy.

Jan 1, 1969

By W. D. Robertson

A considerable quantity of experimental data is available on the effect of time, temperature, work-hardening and composition on the mechanical properties and corrosion resistance of aluminum alloys. There is, however, little unity of purpose or generality in most of the studies and, because the number of variables is large, the results often are not applicable to the problems that arise in practice. Furthermore, the customary expression of the effect of different variables as separate curves or as three-dimensional diagrams with time, temperature and the dependent variables as coordinates, renders the task of selecting the optimum time and temperature of precipitation treatment, with respect to their simultaneous but unequal effect on the various properties, exceedingly difficult. The introduction of precipitation treatments to raise the yield strength of 24S-type alloys has created a need for a method of correlating and summarizing the effect of time and temperature on the mechanical properties and susceptibility to intercrystalline corrosion. Correlation is particularly desirable because the choice of time and temperature is of necessity a compromise to achieve the best possible combination of properties, and the optimum conditions cannot be accurately specified unless the data are in a form in which the alternatives can be readily compared. Therefore, the work described here was undertaken to find a common relationship that would be capable of expressing the effect of precipitation time and temperature on the various properties and by means of which the relative effects could be correlated and their interpretation facilitated. REVIEW OF PREVIOUS WORK In connection with an investigation of the interrelation of age-hardening and creep performance of nickel-silicon-copper alloys, Jenkins and Bucknall1 demonstrated that the time required to attain maximum hardness and tensile strength could be expressed by a relationship of the form: t = K. 10m/T (I) where t is the time in hours and T the temperature of aging, expressed in degrees absolute, and m and K are constants. They also pointed out that this expression is analogous to that which describes the change of the diffusion coefficient D with temperature; namely, D = Ae-Q/RT (2) Later, Cohen used this relationship in studies of the mechanism of age-hardening of a silver- rich copper alloy2 and of duralumin. 3

Jan 1, 1945

By E. D. Levine

A study was made to delineate the dislocation mechanistns controlling prismatic and basal slip at low temperatu.ves in titanium containing approximately 100 ppm interstitial impurities. Mechanical tests were performed on single crystals prepared from electron-beam zone-refined crystal bar by a thermal cycling technique. Measurements were made of critical shear stress and activation volume as functions of temperature. The study revealed that, for prismatic slip, two overlapping thermally activated processes aye rate-controlling below 300°K. From absolute zero to 220°K, reasonable agreement is obtained between experiment and the mechanism which involves overcoming 01. the Peierl's stress by the formation of double kinks in dislocation lines lying in close-packed directions. Between 220°and 300°K prismatic slip is controlled by a difierent thermally activated process which is characterized by a stronger temperature dependence of critical shear stress. Above 300OK, prismatic slip is controlled by an athermal process. Basal slip is controlled by a thermally activated process between 195' and 400°K. An approximate analysis suggests that it is the Peierl's mechanism that is rate-controlling in this range. Below 195°K, the occurrence of mechanical twinning obscures the features of basal slip. Above '400 X, basal slip is controlled by an athermal process. TITANIUM, an hcp metal with a less-than-ideal axial ratio of 1.587, deforms primarily by {1010}( 1210) slip at low temperatures.'-3 The critical shear stress is strongly affected by the presence of interstitial impurities. For example, churchman3 reported that a decrease in the combined nitrogen plus oxygen content from 0.1 to 0.01 pct decreased the critical shear stress for prismatic slip from 9.19 to 1.4 kg per sq mm at room temperature. Basal slip, however, is less strongly affected by interstitials; Churchman's data for {0001}( 1210) slip show a smaller decrease in critical shear stress from 10.90 to 6.3 kg per sq mm at room temperature, for the same decrease in interstitial content. In the present study, an attempt has been made to delineate the dislocation mechanisms controlling prismatic and basal slip in titanium. In view of the strong effects that impurities may have on at least some of these mechanisms, experiments were performed on material of reasonably high purity, so that comparisons could be made with investigations on material of both lower and higher impurity contents. EXPERIMENTAL PROCEDURE Single crystals were prepared from 1/4-in.-diam electron-beam zone-refined crystal bar, furnished by Materials Research Corp. The analysis of the zone-refined material is given in Table I, indicating that the interstitial content is of the same order as the purer of the two materials investigated by Churchman. The Vickers diamond pyramid hardness number of this material was 85.6. This value agrees quite well with the correlation between hardness and interstitial content proposed by Ogden and Jaffee.4 On the basis of hardness, the material investigated in this study is, therefore, intermediate in purity between that studied by Churchman (Vpn = 97) and that studied by Spangler and Herman (Vpn = 62.4).5 Single crystals were grown by thermal cycling through the a-ß transformation. Two-in.-long by 1/2-in.-diam bars were heat-treated in vacuum for 4 hr at 1160°C. The temperature of the furnace was then lowered to 865°C, and the specimens were maintained at this temperature for 7 days. This procedure was repeated several times, resulting in the growth of individual crystals up to 14 in. in length. Specimens for testing in direct shear along either {l010}(1210) or {0001}(1210) were obtained by electrical-discharge machining. The specimens were in the form of cylinders 1/4 in. in diam by approximately 1/2 in. in length, with a reduced section 1/8 in. in diam and a in. in length. All specimens were electropol-ished subsequent to electrical-discharge machining. The ends of the test specimens were cemented in stainless-steel collars. Shear tests were conducted in an Instron Tensile Tester at various temperatures from 4.2° to 508°K. Measurements were made of critical shear stress and activation volume. The latter measurements were performed by measuring the change in flow stress upon cycling the strain rate between ?1 = 2.62 x 10-3 per sec and ?2 = 2.62 x 10-4 per sec. Activation volume, I)*, is defined as u* =ßkT [11

Jan 1, 1967

By Walter R. Hibbard

THE term lineage was first introduced by Buerger1 to denote dendritic branches, grown from a crystal nucleus during solidification from the liquid, with imperfections in alignment of the order of 10-1 [ ] to 10-4 cm. These imperfections are associated with distortions of the dendrites caused by the subsequent solidification of the inter-dendritic filling and are illustrated in Fig 11. Buerger etched zinc single crystals made by the Bridgman2 method to develop this dendritic pattern and show lineage microscopically. Since maximum crystal growth in face-centered cubic metals occurs along the [100] directions,2 lineage would appear in copper as small deviations of dendritic branches from these directions. Elam4 and others5-8 reported etching effects resulting from coring along the cube face traces of various copper alloy crystals. These traces were used by Samans8 to orient single crystals. Williamson9 noted a "lined structure (in alpha brass single crystals) consisting of an ordered arrangement of etch pits" which were probably traces of {100} planes but somewhat uncertain because they were never straight (i.e., apparently lineage). Greninger10 described the distortion and multiplicity of Laue spots in back reflection X ray photograms of copper single crystals showing lineage. The present investigation includes metallographic, X ray and property studies of lineage in large grain and single crystal specimens of copper. PREPARATION OF SPECIMENS Specimens 10 in. long and ½ in. diam of oxygen-free copper (O.F.H.C.) (99.985 pct) were melted vertically in graphite crucibles under an atmosphere of commercial nitrogen purified through a gas train including sulfuric acid, potassium chloride, magnesium chlorate, hot copper chips and phosphorus pentoxide. After a temperature of 1200°C had been maintained for ½ hr, the furnace was turned off and the specimens cooled at a rate of about 150°C an hour. This method of direct solidification was used by Greninger10 to produce single crystals with marked lineage and in this investigation produced a fair yield of single crystal specimens in

Jan 1, 1947

By John Wulff, J. H. Hollomon

SERVICE failures in brass condenser tubes are often due to corrosion. One of the commonest types of corrosion reveals a surface structure of redeposited copper.1 The study of the effect of alloy addition agents in preventing this kind of corrosion has received impetus from the extensive publication of Bengough and May.2 The mechanism of such inhibition gauged from immersion corrosion,' cell corrosion,' and electrolytic potential' measurements is controversial. This controversy may be attributed in part to a lack of definite knowledge5 of the actual nature of the initial films formed in corrosion experiments. In our own corrosion work on copper and 70-30 brass this lack became apparent and necessitated an examination of surface films by X-ray and electron diffraction methods. The present paper is but an introduction to the subject of inhibition and reports the structure of films formed in active corrodents. EXPERIMENTAL PROCEDURE The arsenical coppers and brasses used in this investigation were made by one of the writers* in the plant of the Revere Copper and Brass Co. at Rome, N. Y. Standard commercial melting practice was followed. O.F.H.C. copper and electrolytic zinc were used and all arsenic additions were made by additions of a master copper alloy. The final analyses are given in Table I. Likewise, the extrusion, cold-rolling, and intermediate anneals employed followed standard usage. Final heat-treatments were given to all materials to produce a similar grain size (0.035 to 0.045 mm. for brass and 0.015 to 0.025 for [ ] copper) for the alloys containing various percentages of arsenic. Samples ¼ by ½ by ¼ in. were cut from sheet material resurfaced and finished with 0000 metallographic papers. The surfaces were all degreased in alcohol and benzine before immersion. After immersion in the corrodent, they were washed gently in tap water, dipped first in alcohol and then in pure benzine and allowed to dry in the diffraction apparatus. The reflection electron diffraction apparatus used has been described by Burwell

Jan 1, 1941

By W. C. Leslie

The textures of cold-rolled and of annealed iron are compared with those of an iron-0.8 pct copper alloy in which the amount of precipitation after cold rolling was controlled. Previously published pole figures -for cold-rolled and for annealed iron are confirmed. The effects of precipztatiotz after cold rolling are to retain the cold-rolled texture after annealing, to inhibit the formation of the usual allnealing texture, and to produce elongated recrys-tallized ferrite grains. It is suggested that the inhibition of new textures by precipitation after cold rolling is a general phenomenon. A great deal of attention has been paid to the development of texture during the secondary or tertiary recrystallization of ferritic alloys, but very little work seems to have been done on the control of texture during primary recrystallization. If such control were attained, it might be possible to simplify the processing of oriented materials or to change the characteristics of current cold-rolled and an-nealed products. From previous experience, it seemed likely that texture could be controlled by recrystallizing a supersaturated solid solution. Green, Liebmann, and Yoshidal found that the formation of preferred orientation in aluminum (40 deg rotation about <111> relative to the deformed matrix) was inhibited when iron was retained in supersaturated solid solution in the strained aluminum. The authors attributed this inhibition to iron atoms in solid solution. There is, however, an alternative explanation. Green et al, took a highly supersaturated solution of iron in strained aluminum and heated it to an unspecified temperature for recrystallization. It is probable that precipitation occurred prior to and during recrystallization, and it is proposed that the inhibiting agent is this precipitate, rather than the iron atoms in solid solution. It is important to note that precipitation before cold work is ineffective; the effective precipitate is that formed after cold working and either before or during recrystallization. The location and distribution of the precipitate are critical. Precipitation in such a manner has been found to have profound effects upon kinetics of recrystallization and the microstruc-ture of the recrystallized alloys.2-4 It would be surprising, indeed, if this were accomplished with no change in texture. Because of the relative simplicity of the system, and because of previous experience,4-7 it was decided to determine the effect of precipitation on texture in an alloy of iron and copper. Bush and Lindsay5 found an unspecified change in texture in cold-rolled and annealed low-carbon rimmed steel sheets when the copper content exceeded 0.1 pct. MATERIALS In earlier work, the rate of recrystallization of a low-carbon steel was greatly decreased by 0.80 pct copper, and, after the proper treatment, the recrystallized ferrite grains were greatly elongated.4 Accordingly, it was decided to investigate the effect of precipitation on texture at this level of copper content. The iron and the iron-copper alloy were made from high-quality electrolytic iron and OFHC copper, vacuum-melted in MgO crucibles, cast, hot-rolled to 0.2 in., then machined to 0.150 in. The compositions are given in Table I. The plates were heated to 925°C and brine quenched, twice. This produced a ferrite grain size of ASTM 0 in the iron and ASTM 1 in the Fe-Cu alloy. Disk specimens were cut from the heat-treated plates, repeatedly polished and etched, then used to determine (110) and (200) pole figures by reflection. Despite the complication of large grain size, these pole figures strongly indicated a random texture. PROCEDURES The copper content in solid solution in ferrite before cold rolling and recrystallization, and hence, the amount that could precipitate during the recrys-tallization anneal, was controlled at three levels by heat treatment. The specimens as quenched from 925° C were presumed to have all the copper, 0.80 pct, in solid solution. Other samples of the quenched alloy were aged 5 hr at 700°C to retain about 0.5 pct Cu in solid solution.6 A third set of quenched specimens was reheated to 700°C, then slowly cooled in steps, to reduce the amount of copper in solid solution to a very low level. All specimens were cold-rolled 90 pct, from 0.150 to 0.015 in. thick. The rolling was done in one direction only, i.e., the strip was not reversed between passes, with a jig on the table of the mill to keep the short specimens at 90 deg to the rolls. The rolls were 5 in. in diameter and speed was 35 ft. per min. Machine oil was used as a lubricant. In a supersaturated alloy, the maximum effect of the copper precipitate on microstructure and on recrystallization can be developed by a treatment at 500°C, after cold rolling and before recrystallization.&apos;

Jan 1, 1962

By AIME AIME

CHARACTERIZING the annual meeting as the biggest and best one yet has happened so repeatedly that some may suspect it has become a conventional phrase, like "good morning," and yet, what else can be said when not only one but all the many who express an, opinion use exactly those words, while the first of the two adjectives refers to a physical fact that is established by count. Perhaps there were hose who thought it not the best, but if so they were too tactful to say it out loud. So without further ado . let us admit that the one hundred and thirty-ninth general meeting of the A. I. M. E. held; as our By-laws provide, in New York during the third week of February, was the biggest and best of them all. Perhaps it is because we appreciate friendliness more as we grow older, but somehow the friendly atmosphere seems to rise slowly in temperature with the passing of the years. The actual atmosphere on the opening morning was quite the contrary, Feb. 17 being the coldest clay, so far, of this winter, and the suburbanites claimed that their thermometers had registered 5" below zero when they got up, though the official records of the weather man did not go that low. From that time the temperature steadily rose until Thursday afternoon, when the official record of 66" constituted a record high for Feb. 20 at New York. The weather man explained it by learned discussion of highs and lows, but we suspect the A.I. M. E. members must have had something to do with it.

Jan 1, 1930

Minnesota School of Mines Experiment Station, University of Minnesota, Minneapolis, Minn For copies of Bulletins or a list of publications, address The Director Bulletins available are Bulletin 1, Iron mining in Minnesota, by C. E Van Barneveld (1912), $1, 2, Concentration tests on Mesabi ores, by W R Appleby and E Newton (1913), 3, Preliminary concentration tests on Cuyuna ores, by W. R Appleby and E Newton (1915), 4, Bibliography of Minnesota mining and geology, by W. Gregory (1915); 5, Manganiferous iron ores of the Cuyuna district, by E Newton (1918), 7, The future of the Lake Superior district as iron-ore producer, by E W. Davis (1920), 9, Magnetic concentra¬tion of iron ore, by E. W Davis (1921); 10, Ball mill crushing iii closed circuit with screens, by E W Davis (1925); 12, Minnesota manganiferous iron ore in relation to the iron and steel industry, by T L Joseph and S P Kinney (1927). A mining directory of Minnesota, by J J Craig, is published annually.

Jan 1, 1933

By Cyril Wells, Robert F. Mehl

THE practical importance of a knowledge of the rates of diffusion of carbon and of alloying elements in steel has often been pointed out 1,13 This importance ties chiefly, though by no means only, in the study of carburizing and in the study of reactions involved in the eutectoid inversion in steel The present paper is one of a series in which the rates of diffusion of carbon and of alloying elements in y-Fe and their mutual effects are measured with precision The present paper reports the results of meas¬urements of diffusion coefficients D of manganese in -y-Fe over wide ranges of temperature, and concentration, and of the effect of carbon and impurities and of grain size upon these diffusion coefficients [$] Data on the rate of diffusion of manganese in y-Fe have been reported by Fry2 and by Paschke and Hauttmann 3 Claussen4 calculated the diffusion coefficient D from Fry&apos;s data as 3 X 10-10 sq cm per sec at 960°, which will be shown to be incorrect by many orders of magnitude Paschke and Hauttmann gave a rough value of 9 6 X 10-8 sq cm per sec at 1400° Neither of these workers considered the variation of D with concentration and with temperature, and neither studied the effect of carbon METHOD AND MATERIALS Method -The method employed is identical with that used earlier 5 Owing to the volatility of manganese and the consequent loss of manganese from the outer surfaces of the welded diffusion bars, these surfaces were discarded The diffusion anneals were performed in argon, except one (test 18) in which hydrogen was employed for purposes noted below The temperatures reported are accurate to about 3° and constant within I°C , the diffusion times are reported to the nearest half hour The specimens were furnace-cooled after the diffusion anneal at about 2° per min Layers, usually 0 004 in thick, parallel to the welded interface, were machined from the specimens after the diffusion anneal and analyzed for manganese to 0 01 per cent Mn, the distances of these layers from the interface was measured to ±0 0002 in The concentration-penetration curves obtained were analyzed by both the Grube and the Matano solutions e The D values calculated by these two methods do not differ appreciably 5 The Grube solution was used except in tests in which the effect of concentration upon D was greater than 50 per cent, in such cases the Matano method, which explicitly provides for the variation of D with concentration,1 is the more reliable In general, the number of layers machined from the zones of variable composition were large enough to permit the use of the average layer composition as that of the center of the layer In 3 of the 30 tests, however, this is not justified, and the D values

Jan 1, 1941

By E. S. Shaw

Utah continues as a very unimportant factor in oil and gas production. During the year 1933 there were no developments of noteworthy significance. There were no completions of either producers or dry holes during the year. The only commercial production in Utah has been from the five fields listed in Table 1 and mentioned briefly in the following paragraphs. Virgin Field.—Oil was first discovered in 1907 in a well drilled in sec. 13, T.41S., R.12W. Of the first 15 wells drilled, all but three were failures. In 1920 production from four wells from 550 to 600 ft. deep was about 20 bbl. per day. San Juan Field.—The majority of drilling in this area was done during the years 1908 to 1911. Probably the first commercial well in the field was that drilled in NW.$i sec. 6, T.42S1 R.l9E., in 1909. Farnham Dome.—The discovery well was drilled by the Utah Oil Refining Co. in April, 1923. A number of years later a second well was drilled by The Carbon Dioxice a mica1 Co., the present operator, which markets the gas in the manufacture of "dry ice" and allied products. The eslimated production, as indicated in Table 1, includes 500 million cubic feet estimated to have been lost while drilling the discovery well. Cisco Dome.—The discovery well was drilled by The Utah Oil Refining Co. in November, 1924. The Crystal Carbon Co. markets the gas in the manufacture of carbon black. Ashley Valley Dome.—The discovery well was drilled by the Utah Oil Refining Co. in April, 1925. The Uintah Gas Co. markets the gas in Vernal, Utah, for fuel. Woodside Structure.—In February, 1924, gas was discovered in a well drilled by Utah Oil Refining Co. in sec. 12, T.l9S., R.l3N., on the Woodside structure. This gas is reported to contain an appreciable amount of helium. It is being held as a helium reserve by the United States Government. Salt Valley Structure.—The test of the Utah Southern Oil Co. in sec. 31, T.23S., R.21E., on the Salt Valley structure has been shut down for some time at 6120 ft. This test has passed through more than 5000 ft.

Jan 1, 1934

By -1ng. B. Stampa

Introduction The author investigates the technology available for mine flow reversibility by reversing the operation of the main fans. Air bypassing, reversed sense of impeller rotation, and blade angle overturning are indentified as the most important means of achieving the flow reversibility. The axial flow fan is studied in some detail and the important variables of fan capacity, pressure, and efficiency are compared. Although some references to rotor velocity vectors are made (Fig. 1) the subject is not presented in detail in order to stay within the paper&apos;s intent. Background Where applicable, the requirements of mine air reversibility are an important part of the overall mining operation and should be given proper attention during mine ventilation design. The need for preventing mine fires in intake shafts is the reason that, for a long time, it has been a requirement in the Eastern European countries to have the feature of flow reversibility built into all main mine ventilators. In West German coal mines, this requirement does not apply since higher priority is given to safety precautions. Some metal and mineral mines in the United States are required to have ventilation currents readily reversible in order to fight underground fires as stipulated by the Title 30, Mineral Resources, and by the mining laws of some states. From an equipment performance standpoint, assuming the validity of other flow technical data, the question of the ventilator flow reversibility is of ten related to the actual scenario of the begining of an emergency. Most of the time, the system operates within half of the design capacity only and assuming that the equivalent mine orifice is stable only a quarter of the design pressure must be produced by the fan in any event. For such an operating point, the effective gradient and energy losses can be disregarded since fire emergencies usually last for only a short period of time. Equipment Comparison Fig. 2 depicts a ventilation arrangement wherein the flow can be reversed through changing the functional position of the dampers (slideplate or blade type). Both axial or radial fans can be applied in this case. For the mine exhaust operation, slide plates (guillotines) I and I11 are opened while plates I1 and IV are closed. The flow thus encounters two reroutings, which due to the reverse ventilation flow cannot be equipped with proper sheet metal deflectors, and therefore large pressure losses result during the air reroute operation. When the intake operation is considered, the dampers I and I11 are closed and the pair I1 and IV is opened. For this arrangement, almost the same flow conditions exist as for the foregoing exhaust mode; however, one additional corner must be applied. Even if a highly efficient ventilator is used in one of the previously described systems, the additional pressure losses resulting from air rerouting plus those encountered between the planes "0" and "00" significantly affect the overall energy requirements. Illustrated in Table 1 are the relative component losses of even flows for evaluation of various detour arrangements. From the aerodynamics standpoint, when these components are being designed, serious consideration must be given to the individual levels of the pressure loss coefficient, 5. Should there be an increase in the sum of the loss coefficients due to the particular component selection such that, for example, the net change is 1.2; the resulting pressure to be produced by the fan can be calculated from:

Jan 1, 1981

By Atmaram H. Soni

A statistical study of machinability led the writer to examine existing data in regard to a thermal conductivity-mechanical properties relationship. Various functional relationships were proposed and the multiple correlation determined by statistical methods discussed in detail in the 1iterature.4 The relationship having the greatest engineering and statistical significance is discussed below. The data for cast and wrought aluminum alloys&apos;&apos;2 have been compiled and analyzed separately. The results of these analyses are presented in Table I. The empirical equations for estimating the thermal conductivity (cal sec-&apos; cm-I °C) are: for cast aluminum alloys K = 0.584 -0.00984/ -0.00263 Hb [1] and for wrought aluminum alloys K = 0.604 -0.00424l- 0.00170 Hb [2] where 1 and Hb are the percentage elongation and Brinell hardness number. Results of a similar analysis pertaining to stainless steel3 are also given in Table I. This empirical equation for estimating thermal conductivity (Btu per sq ft per hr per °F per in.) is K = 216.19- 1.716l - 0.0997 Hb [3] The relationship between electrical resistivity and mechanical properties was also examined for aluminum alloys using the same statistical method. These results are listed in Table II. While the situation of exact multiple correlation

Jan 1, 1965

By J. L. Walter

Tungsten ingots with and without small amounts of aluminum, silicon, and potassium were prepared by conventional powder-metallurgy techniques and hot-rolled to 0.001-in.-thick sheet. The ingots were rolled three different ways: I) initially cross-rolled, then straight-rolled; 2) straight-rolled, and 3) straig-ht-rolled, then cross-rolled. Total reductions were 99.8 pct. Samples of the sheet were annealed, in vacuo, at temperatures from 1200" to 2200°C. Substructures and textures were determined by transmission electron microscopy and X-ray diffraction. The doping agents were found to form particles of mullite (Al,Si,O,,) which, on rolling., appear to form strings of particles aligned parallel to the rolling direction. These strings of purticles are associated with an increase in the recrystal-lizution temperature from 1400°C for undoped sheet ThE addition of small amounts of aluminum, silicon, and potassium to tungsten has a profound effect on the grain structure of drawn and recrystallized wire.&apos;" The grains generally do not occupy the entire cross section of the wire and have lengths many times the wire diameter. The grain boundaries are, in general, nearly parallel to the wire axis and have long &apos;&apos;fingers" projecting one into the other as may be seen in Fig. 1, a longitudinal section of recrystallized doped wire. When unsuitable additions or no additions are made or the wire is not sufficiently reduced during draw- to 2000°C for doped sheet. The particles also limit lateral grain growth while allowing growth parallel to the rolling direction. This results in the formation by primary re crystallization of large highly elongated grains of specific preferred orientation. The undoped sheet recrystallized to small equiaxed gains. The preferred orientations of the large grains were determined from axis density plots to be (113) (691). (0i3) (362), (116) (304) types for the initially cross-rolled then straight rolled material. These orientations are similar to the orientations of the large elongated grains in doped and heavily drawn tungsten wire. The nucleation of the large grains appears to occur by recovery of certain of the polygonized regions of specific orientation, followed by strain-induced boundary migration.

Jan 1, 1968

By G. H. Billings

The following data were collected some years since in the course of a series of experiments with oil as a substitute for coal in puddling, the earlier form of the Archer apparatus being employed. While I am not at liberty to give the name of the works at which the tests were made, I can personally warrant the accuracy of the record, and feel justified, therefore, in making this contribution to a discussion in which all trustworthy figures are likely to be helpful. Oil in Puddling. Starting with everything cold, we used 3518 pounds of oil to get the furnace in condition to charge a heat. After charging, we ran until we took out 13,340 pounds of puddled-blooms, using 8437 pounds, or 1333 gal1ons of oil. At 49 cents per gallon, the cost for oil per gross ton of blooms was $10.91; with coal at $4.871/2, and a consumption of 1.5 tons per ton of blooms (the usual quantity is ton for ton), the cost for coal is $7.31. The quantity of oil consumed in heating the furnace to the proper temperature and to make the bottom, is not included in the oil consumption stated. The quantity used was 3518 pounds of oil. Usually, the furnace need only be lighted, but in this case it had been repaired, and it was necessary to make a new .bottom. This required the quantity of oil necessary to make three heats of blooms, or 2109 pounds. Deducting this from the 3518 pounds leaves 1409 pounds of oil as the quantity consumed every week in lighting up.

Jan 1, 1889

By David J. Mack, Dennis R. O’Boyle

The mechanical properties of three polycrystalline intermediale Phases that have the y bvass structure were measured in compression between 400° and 900°K. At the lower testing temperatures— termed Region I— no plaslic deformalion occurred prior to brittle fracture. A1 higher temperatuves— lermed Region III—the three phases deformted plaslically for all strain rules from 1.5 to 36.0 x 10-4 sec-1. The homologous temperature for the inilialion of plaslic flow increased in the same order as tile atomic size difference between atoms in each plmse (0.50Tmp for y AgZn, 0.62Tmp for y CuZn, and 0.76 Tmp for y CuCd). The increased flow resistance oj. y CuCd above 0.50Tmp is explaitzed in terms of the more random arrangement of copper and cadmiwn atoms on the laltice sites of&apos; the y bvass slrilctlire. Plaslic deformation of the three intevmiediate phases in the smooth plaslic flow region appears to be controlled by oacancy-induced dislocation climb. RESULTS are reported in this paper on the mechanical properties of three binary intermediate phases having the complex cubic y brass structure (D82, 143m) containing fifty-two atoms in the unit cell. The structure of the prototype y brass phase, y CuZn, was first analyzed by Bradley and Thewlis&apos; based on the X-ray work of Westgren and Phragmen.2 The y brass structure can be visualized as an arrangement of twenty-seven body centered cubes (3 by 3 by 3) into a unit cell containing fifty-four atoms. From this cell the center atom and the four corner atoms are removed and the remaining fifty-two atoms are slightly displaced, resulting in the y brass structure. Approximately thirty-five binary intermediate phases that have an electron to atom ratio of 21:13 have been reported to form the y brass structure. The three y brass intermediate phases selected for this study are formed from elements in groups IB and IIB of the first and second long period. Each compound exists over a composition range that includes the stoi-chiometric composition A5BB (Cu5Zn8, Ag5Zn8, Cu5Cd8). y CuZn (a0 = 9.944A) has solubility limits extending from 57 to 67 at. pct Zn at 500°C apd melts peritecti-cally at 834°C; y AgZn (a0 = 9.326A) has a solubility range from 58 to 63 at. pct Za and melts peritectically at 661°C. y CuCd (a, = 9.596A) melts congruently at 563°C and has a maximum range of solubility extending from 58 to 64 at. pct Cd. In order to minimize coring and microporosity during solidification, the composi- tion of the compounds that form peritectically (y CuZn and y AgZn) was chosen so that the temperature difference between the liquidus and the solidus is only a few degrees centigrade. In selecting y brass compounds for this study, phases were avoided that contained a transition element or that had an electron to atom ratio greater than 1.70. Hume-Rothery et a1.3 has observed that, when the electron to atom ratio of y brass intermediate phases exceeds 1.70, atoms begin to drop out of the unit cell to maintain a constant electron to unit cell ratio. EXPERIMENTAL PROCEDURE Each of the three intermediate phases was prepared from elements having a purity >99.999 pct. Only trace quantities of impurities (<0.0001 pct) were detected by spectrographic analysis of the four elements. A master alloy of each compound weighing approximately 130 g was prepared by melting the elements in evacuated sealed quartz tubes heated in an electric furnace. Specimens having the desired diameter for the mechanical property measurements were obtained by remelting the master alloy in a Vycor tube under argon and drawing specimens from the melt into 4-mm-diam quartz tubes. To obtain a smooth surface on specimens drawn from the melt, the inside of the quartz tube was coated with a thin layer of graphite formed by thermally decompositing acetone. All specimens were imbedded in sealing wax and cut to the proper length with a diamond-impregnated cut-off wheel. The individual specimens (4 mm diam by 8 mm long) were then sealed in evacuated Pyrex tubes and homogenized at 500°C for 72 hr. Chemical composition, variations over the length of the rod (10 cm long) were generally less than 0.2 pct and the analyzed composition differed from the desired composition by less than 0.5 pct.4 Metallographic evaluation of the specimens after homogenization showed no evidence of a second phase in any of the compounds. The average grain size of the specimens after homogenization was 1 to 2 mm. All mechanical property measurements were carried out in an argon atmosphere using a compression apparatus described previously.5 Prior to applying the load to the specimen, the test apparatus was evacuated, back-filled with argon, and heated to the test temperature for 20 min to establish thermal equilibrium. EXPERIMENTAL RESULTS Mechanical properties measured as a function of temperature and strain rate for each of the three y brass phases were similar. From room temperature to the ductile-to-brittle transition temperature, only elastic deformation was observed prior to fracture.

Jan 1, 1968

By B. Russell, J. R. Tuffley

The stability regions of the normal sulfate and the various basic sulfates of lead in 02-SO2 and PhS-SO2 gas atmospheres were calculated from available thermodynamic data over the temperature range 600o to 1200oC. Roasting experiments were carried out on samples of synthetic lead sulfide at various temperatures and for various roasting times. Compounds in the oxidized surface layers were identified by X-my techniques. The order in which the various reactions occurred was determined by considering the experimental results in conjunction with the thermodynamic calculations. The results of these investigations were applied to the sintering of lead sulfide concentrates, and the extent to which each of the various reactions occurred in a sinter bed was determined. PRODUCTS which may be formed during the roasting of PbS are PbS04, PbSO4. PbO, PbSO4 . 2PbO, PbSO4 . 4PbO, PbO, and Pb.1, 2 The formation of some sulfates has been studied by a number of investigators3-&apos; whose experiments have mostly been confined to the temperature range between 500" and 850°C. Culver, Gray, and spooner3 were able to isolate the reactions which formed PbSO4 and PbSO4. PbO by carefully controlling the compositions of the roasting gases. The kinetics of these reactions were studied within the temperature range 673" to 807°C. Gardner4 and Merrick5 also studied the kinetics of oxidation of PbS. Gardner disregarded the existence of basic sulfates and, as a result, misinterpreted some of his experimental data. Merrick recognized that most of the products of oxidation were mixtures of basic sulfates, but was unable to identify the individual components. Hoschek6 oxidized PbS in air at temperatures up to 800°C. The products were analyzed after roasting had proceeded for some minutes, but no attempt was made to follow the reactions in the initial stages. Consequently, no roasting mechanism was derived. He also showed that PbSO4 and PbSO4 - 4PbO decomposed slowly in air at temperatures in the region of 800°C. The stability regions of sulfates in the temperature range 500" to 850°C have been only partially determined, while practically no information is available for higher temperatures. In the lead sintering process, PbS concentrates are roasted at temperatures within the range 1000o to 1200°C. In order to investigate sulfate formation, it is necessary to determine the effects of temperature and gas composition upon the various oxidation reactions and upon the regions of stability of the various lead sulfates. Kellogg and Basu2 introduced an equilibrium diagram similar to those in Figs. 1 to 3 which provides a convenient medium for represent-

Jan 1, 1964

By G. Sachs, G. Espey

It is well known that high residual stresses are created in tubing by the sinking process, in which no internal tool or mandrel is used.l-4 In this process, the wall thickness is usually slightly increased.4,5 In tube drawing with a plug, the residual stresses are gradually reduced with increasing reduction of wall, and, by the conventional approximation methods of stress measurement, no stress is detected in tubing in which the wall reduction is over 50 per cent of the total reduction in cross-sectional area.l,6 It has also been observed that higher stresses are set up by the use of wide-angle dies than by acute dies.? &apos;he actual stress distribution in drawn tubing has been investigated by a few workers,7"9 but has not been related to the conditions of the drawing process. Some attempts have been made to correlate accelerated cracking tests for stresses in brass products with the magnitude of the circumferential stresses and the condition of the metal. l-10 Materials and Preparation The material used in this experiment was commercial cartridge-brass (70 per cent Cu, 30 per cent Zn) tubing in the straightened and annealed condition.† This tubing was of various outside diameters and had a wall thickness of 0.032 in. It was sunk to 1/2-in. outside diameter by means of an experimental 10,000-lb. hydraulic draw-bench, with a maximum drawing speed of 8 ft. per minute. All dies were of 1/2-in. diameter. Various die materials and contours were used in order to study the effect of die design on residual stresses in the tubing; the specifications for the dies are given in Fig. I. Castor oil was used throughout as lubricant. Chattering occurred in sinking some of the tubing, particularly when dies having the more acute angles were used, but the tubes showing chatter marks did not differ markedly from those with a smooth surface regarding the magnitude of residual stress. The source of the chattering has not yet been revealed. The wall of the tubing becomes thicker during s nking, as previously observed.4,5 The increases in wall thickness were 3, 7 and 9.5 per cent on the average for 6, 20 and 33 per cent reductions in diameter, no effect of the contour being noticeable. The reductions mentioned throughout this investigation do not consider this fact, but are reductions of the outside diameter. For the evaluation of the residual stress two fundamentally different methods were used. The approximation methodl,ll of splitting a certain length of tubing was employed to obtain a bird&apos;s-eye view of the effect of various drawing conditions on the circumferential stress measured by the diameter change occurring On splitting. For

Jan 1, 1942

By G. Sachs, G. Espey

It is well known that high residual stresses are created in tubing by the sinking process, in which no internal tool or mandrel is used.l-4 In this process, the wall thickness is usually slightly increased.4,5 In tube drawing with a plug, the residual stresses are gradually reduced with increasing reduction of wall, and, by the conventional approximation methods of stress measurement, no stress is detected in tubing in which the wall reduction is over 50 per cent of the total reduction in cross-sectional area.l,6 It has also been observed that higher stresses are set up by the use of wide-angle dies than by acute dies.? &apos;he actual stress distribution in drawn tubing has been investigated by a few workers,7"9 but has not been related to the conditions of the drawing process. Some attempts have been made to correlate accelerated cracking tests for stresses in brass products with the magnitude of the circumferential stresses and the condition of the metal. l-10 Materials and Preparation The material used in this experiment was commercial cartridge-brass (70 per cent Cu, 30 per cent Zn) tubing in the straightened and annealed condition.† This tubing was of various outside diameters and had a wall thickness of 0.032 in. It was sunk to 1/2-in. outside diameter by means of an experimental 10,000-lb. hydraulic draw-bench, with a maximum drawing speed of 8 ft. per minute. All dies were of 1/2-in. diameter. Various die materials and contours were used in order to study the effect of die design on residual stresses in the tubing; the specifications for the dies are given in Fig. I. Castor oil was used throughout as lubricant. Chattering occurred in sinking some of the tubing, particularly when dies having the more acute angles were used, but the tubes showing chatter marks did not differ markedly from those with a smooth surface regarding the magnitude of residual stress. The source of the chattering has not yet been revealed. The wall of the tubing becomes thicker during s nking, as previously observed.4,5 The increases in wall thickness were 3, 7 and 9.5 per cent on the average for 6, 20 and 33 per cent reductions in diameter, no effect of the contour being noticeable. The reductions mentioned throughout this investigation do not consider this fact, but are reductions of the outside diameter. For the evaluation of the residual stress two fundamentally different methods were used. The approximation methodl,ll of splitting a certain length of tubing was employed to obtain a bird&apos;s-eye view of the effect of various drawing conditions on the circumferential stress measured by the diameter change occurring On splitting. For

Jan 1, 1942

By J. T. Eash, N. B. Pilling

As a group, the alloys of iron, nickel and carbon are, in application, one of the most versatile of the ferrous alloy family, and while many investigations have been made of their properties and structure only a few attempts have been made to show the microstructural relation over a wide range of composition. It is neccssary at once to be clear about the kind of structural relation with which we are concerned. The equilibrium diagram provides quantitative data regarding structural make-up and its variation with temperature, but equilibrium structures often are far different from the technologically significant ones. This is particularly true of steels and cast irons, and some of the most important cornponents of their structure may have no place in the diagram of equilibrium. The fugitive phases that result from transFormation are dealt with in a highly detailed and descriptive manner in. the S-curves originated by Bain. The present effort lies in a still different field, which is to portray the relation between sweeping changes in composition and in resultant structure, under controlled but nonequilib-rium conditions of thermal treatment. The Guilletl diagram presented in 1903 was the initial attempt at such a correlation and frequent references to this diagram can be found in the literature even though later work has shown it to require considerable modification. Kase2 published in 1925 a diagram that was based upon observations of furnace-cooled melts: At low carbon contents this diagram agreed fairly well with experience in showing pearlite to exist in alloys containing up to 6 per cent nickel, and completely austenitic steels at 24 to 28 per cent nickel. At high carbon contents, however, irons were indicated as being martensitic at around 2 per cent nickel and austenitic at 12 per cent nickel; these interpretations are open to serious question. Based upon experience with steels and cast iron, Wickenden3 suggested a diagram that agreed fairly well with Kase&apos;s in the steel range but which presented a wide departure from it in the higher-carbon, high-nickel regions. In its broad picture, the Wickenden diagram is an approach to "practical" structures. Reference may be made at this point to the efforts of Marsh5 to construct an iron-nickel-carbon equilibrium diagram by the coordination of scattered data, which are described at length in the monograph cited. None of these preceding attempts has appeared to meet the need for an exposition of the structural sequence that is encountered in traversing the nickel-steel nickel-cast-iron range under what might be called ordinary departures from equilibrium. ExperimentaT. Conditions Since the intent has been to establish structures developed under arbitrary conditions, experimental details have perhaps more than usual importance. Three of these conditions should particularly be noticed: (I) the material was cast and not subjected

Jan 1, 1942

By J. T. Eash, N. B. Pilling

As a group, the alloys of iron, nickel and carbon are, in application, one of the most versatile of the ferrous alloy family, and while many investigations have been made of their properties and structure only a few attempts have been made to show the microstructural relation over a wide range of composition. It is neccssary at once to be clear about the kind of structural relation with which we are concerned. The equilibrium diagram provides quantitative data regarding structural make-up and its variation with temperature, but equilibrium structures often are far different from the technologically significant ones. This is particularly true of steels and cast irons, and some of the most important cornponents of their structure may have no place in the diagram of equilibrium. The fugitive phases that result from transFormation are dealt with in a highly detailed and descriptive manner in. the S-curves originated by Bain. The present effort lies in a still different field, which is to portray the relation between sweeping changes in composition and in resultant structure, under controlled but nonequilib-rium conditions of thermal treatment. The Guilletl diagram presented in 1903 was the initial attempt at such a correlation and frequent references to this diagram can be found in the literature even though later work has shown it to require considerable modification. Kase2 published in 1925 a diagram that was based upon observations of furnace-cooled melts: At low carbon contents this diagram agreed fairly well with experience in showing pearlite to exist in alloys containing up to 6 per cent nickel, and completely austenitic steels at 24 to 28 per cent nickel. At high carbon contents, however, irons were indicated as being martensitic at around 2 per cent nickel and austenitic at 12 per cent nickel; these interpretations are open to serious question. Based upon experience with steels and cast iron, Wickenden3 suggested a diagram that agreed fairly well with Kase&apos;s in the steel range but which presented a wide departure from it in the higher-carbon, high-nickel regions. In its broad picture, the Wickenden diagram is an approach to "practical" structures. Reference may be made at this point to the efforts of Marsh5 to construct an iron-nickel-carbon equilibrium diagram by the coordination of scattered data, which are described at length in the monograph cited. None of these preceding attempts has appeared to meet the need for an exposition of the structural sequence that is encountered in traversing the nickel-steel nickel-cast-iron range under what might be called ordinary departures from equilibrium. ExperimentaT. Conditions Since the intent has been to establish structures developed under arbitrary conditions, experimental details have perhaps more than usual importance. Three of these conditions should particularly be noticed: (I) the material was cast and not subjected

Jan 1, 1942