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By A. P. Young, F. C. Holden

IN many of the early investigations of the aging of titanium alloys, it was observed that the retained beta phase could be aged to a high hardness without an apparent change in microstructure; moreover, in some alloys it was found that even a very rapid quench produced beta having an anomalously high hardness. What appeared to be hard beta phase was designated as "beta prime," to distinguish it from the softer normal retained beta phase. Frost and co-workers'.' demonstrated that the increased hardness is caused by the formation of a metastable phase, omega, which can be identified from its X-ray diffraction pattern. This reaction also has been studied by Brotzen, et a1.,3 and the structure of the omega phase has been published by Silcock, et al.," and by Austin and Doig.3 The X-ray diffraction analysis by Austin and Doig indicated that omega is low in alloy content, approaching the composition of the stable alpha. The aging process proceeds by the path Thus, overaging results in the formation of alpha and the ultimate disappearance of omega. This is accompanied by a decrease in hardness and a recovery of ductility. The present work represents a study of the aging process in a high-purity Ti-6.4 Mn alloy, in which optical and electron microscopy, X-ray diffraction, and hardness data were correlated. The specimens, 1/8-in. sections cut from 1/4-in.-diam rod, were heated in argon to a temperature just into the beta field (810°C), held for 1 hr, and water-quenched. The aging samples were heated in air for 1 hr at 300 to 500°C. The specimens were ground and mechanically polished, finishing with alternate etching and polishing on a slow-speed wheel to remove disturbed metal from the surface. The surface finally was swab-etched, rinsed with tap water, and air-dried. The etchant used was an aqueous solution containing 1 1/2 pct HF and 3 1/2 pct HNO3. This procedure is identical with that used for preparing metallographic specimens, except that the optimum degree of final etching required for electron microscopy may vary from that used for light microscopy. Metal replicas, backed with silica, were prepared from plastic strip-pings for examination with the electron microscope. Complete details of the procedures used for preparing these replicas are available elsewhere." Fig. 1 shows a comparison between the structures revealed by the electron microscope and those obtained by conventional light microscopy. Hardness and X-ray diffraction data are listed in Table I. No evidence of a second phase is apparent in the optical micrographs for the as-quenched sample, or for the specimens aged at 300 or 400°C; fine alpha particles can be seen in the specimen aged at 500°C. In the electron micrographs, visible equiaxed pocks appear in the as-quenched specimen, and, with increasing density, in the samples aged at 300 and 400°C. After aging at 500°C, the structure becomes markedly coarser and develops an oriented appearance of a second phase, coincident with the appearance of visible alpha particles in the light micrographs. These observations are confirmed by X-ray diffraction and hardness data. In these specimens, aged for 1 hr, the intensity of the omega pattern increases to a maximum at the 400°C aging temperature, corresponding to maximum hardness and the maximum density of pocks in the electron micrographs. At this point the beta phase shows a definite enrichment in Mn. At 500°C, little or no omega re-

Jan 1, 1959

By W. A. Backofen, M. L. Ebner, D. H. Avery

Slip on (11221 planes, which do not contain a close-packed direction, has been identified on crystals strained at —190°C. Evidence is also presented to show that (1071) pyramidal slip is a mode of deformation at 25ºC in a limited orientation range. PYRAMIDAL slip (l011) <1120> in magnesium at elevated temperatures has been established for some time.1-2 More recently, Burke and Hibbard reported this mode of deformation at 25"C, but their results were questioned by Craig.&apos; Also, as far as the present authors know, slip in magnesium on pyramidal planes of the second order, Type II (1122)) has not previously been reported, although Gilman"," identified (1122) slip in another hexagonal metal

Jan 1, 1959

By D. E. Swets, P. C. Frank, D. L. Fry

THE hydrogen-steel system has received considerable attention in the past several years. Recent experiments in this laboratory have shown that another group of operations, namely, sanding and grinding,&apos; must be added to the already well-estab- lished list of phenomena that cause hydrogen to permeate steel. Qualitative tests which have been made with a very sensitive ion-gage type of detector similar to that used by Norton,&apos; showed that sanding a steel surface caused hydrogen permeation. A metal type 6~6 vacuum tube was used as the ion gage with an electrometer to measure the ion c~rrent.~ The protective enamel coating was removed from the tube.

Jan 1, 1959

By W. J. Feuerstein

STRUCTURAL imperfections in crystals, in recent years, have received the attention of metallurgists concerned with single-crystal properties. Certain etching solutions cause local variations in the rate of etching at dislocations.&apos; Dislocations are expected to be preferred segregation sites for impurities&apos; and Dash" has shown that precipitation of copper occurs preferentially at edge-type dislocations in silicon. Diffusion rates along dislocations are probably higher than rates through more perfect crystalline regions.&apos; These and other phenomena which may be influenced by dislocations may be studied by revealing dislocations as etch pits. This study was undertaken to develop an etch solution for silicon crystals to reveal low-angle grain boundaries associated with dislocation arrays and also etch pits with good contour definition. Additional etches have been reported by Dash," Vogel and Lovell.T rystals were sectioned on crystal lographic planes at various angles from the growth direction for microscopic examination of the etch pits. Experimental Crystals of 75-100 g mass and 3 cm diam were oriented and cut on crystallographic planes. Crystals were grown by Czochralski technique from a quartz crucible. Lapped wafers were chemically polished in a modified CP-4 etch for 1 min (bromine was deleted from the etch solution). The polished wafer was then etched* at room temperature for 5 min. The use of larger amounts of H2O and/or CH2COOH slows the reaction. The Cu(NO3),.3H2O was added to increase the reaction rate and give more definition to the etch pits developed. Some pertinent features concerning the attack of this etch on silicon can be summarized as follows:

Jan 1, 1959

By H. W. Weart, D. J. Mack

SINCE the first recognition of eutectic solidification as the simultaneous formation of 2 solids from one liquid,1 many complex structures thus produced have been observed. Despite many attempts to ascertain it, however, the detailed mechanism by which these structures develop remains obscure. In an attempt to clarify the origin of eutectic structures, the study described in this paper was undertaken. The broad objective of the study was to investigate the effect of different solidification conditions on the structure of single eutectic grains. It was expected that the eutectic structure would thereby be isolated for study, but preliminary experiments revealed another structure within each grain. When a literature surveye revealed only 1 previous mention of the subgrain structure,&apos; it was decided to determine the origin of the structure and to study its response to various growth conditions. Before presenting the details of the investigation, it is apropos to distinguish carefully among the 3 structures produced during eutectic solidification. This is especially desirable in view of the apparent confusion existing in the literature concerning the structures in a eutectic specimen. The confusion is due largely to the existence of the subgrain structure noted above which is easily mistaken for the grain structure. In the following paragraphs, an effort is made to distinguish the 3 structures by briefly describing their origin during solidification, even though the results are thereby anticipated slightly. The 3 structures produced during eutectic solidification are the grain structure, the colony structure, and the eutectic structure, each one contained by the one preceding it. The grain structure is analogous to the polycrystalline structure of a metal or inle-phase alloy in that each grain grows from a single nucleus. A somewhat broader interpretation is needed here, however, because of the presence of 2 phases in each grain. Since in most eutec-tics particles of 1 phase are dispersed in a matrix of the other, a grain may be taken to be the region in which the matrix phase is monocrystalline.*

Jan 1, 1959

By M. E. Steidlitz

FABRICATION of semiconductor devices.requires the formation of electrical contacts, often involving the use of lead, tin, and lead-tin solders. In several recent instances when germanium had been bonded to nickel with lead, a new metallic phase appeared, unattached to either germanium or nickel, but rather suspended as a continuous foil wholly contained within the lead matrix. Since such metallic structures do not form in nickel-lead or lead-germanium systems under identical conditions, a brief study was made of the characteristics and possible formation mechanisms of this material. EXPERIMENTAL Specimens were prepared by placing single crystal, transistor-grade germanium discs, 30-mils in diam by 6-mils thick, into a depression in a graphite boat. =A* nickel washers, 125 mils in diam by 5-mils thick, containing a center hole of 10-mils diam, and coined into a soup bowl shape, were loaded on top of the discs, so that the bottom of the

Jan 1, 1959

By F. J. Darnell

SINGLE crystals of a-titanium are difficult to obtain because of the phase transformation at 882 °C in which the structure changes from the high-temperature body-centered structure ß to the low-temperature hexagonal close-packed structure a. Preparation of single crystals of a-titanium has previously been accomplished by a strain-anneal method which takes several days and is not generally reproducible. Single crystals as large as 5 mm

Jan 1, 1959

By D. C. Buffum, L. D. Jaffee

THAT the path of brittle fracture in temper brit-tle steel differs from the path in unembrittled steel has been known for some years1 More recently it has been recognized that alloy steels, at least, display two forms of temper brittleness1 which may be called upper-nose and lower-nose. No experimental evidence has been published, however, to show that the two forms can be distinguished by simple inspection of brittle fractures. Fig. 1 is a photograph of the fractured surface of a V-notch Charpy specimen of SAE 3140 steel in which temper brittleness developed during heat treatment in the lower-nose temperature range.&apos; The steel was austenitized 1 hr at 900°, tempered 1 hr at 675", and embrittled 48 hr at 500°C, with water-quenching after each treatment. The temperature of transition from ductile to brittle fracture in the Charpy test was + 2°C (criterion of transition: 100 pet fibrous). The hardness was Re 25; the austenite grain size ASTM 8. This specimen was broken at 150°C in the standard Charpy test and absorbed 6.6 ft-lb; its fracture was rated 0 pet fibrous. Fig. 1 shows that the fracture appeared rough to the naked eye. At the side of the specimen away from the notch (the side which was in compression during the test), the surface as a whole deviated sharply from a plane. At higher magnification, X60, the surface gave

Jan 1, 1959

By J. J. Gilman

WHEN a zinc crystal is immersed in liquid N, and then plastically bent, it is observed that it bends a certain amount and then fractures. This fracture occurs at a fairly reproducible radius of curvature, and the radius varies systematically with the crystal orientation. Because this phenomenon provides an unusually well-defined set of fracture conditions, it is possible to examine various possible fracture criteria and select the best one. Zinc crystals of various orientations were bent about axes lying in their glide planes (0001) and perpendicular to their glide directions [1210]. The crystals were 0.150 in. diam and had chemically polished surfaces. They were immersed in liquid N, while they were being loaded on an Instron testing machine. A three-point bending jig was used to load them. Table I lists the beam radius at which fracture occurred for each orientation. Each value is the average of at least six measurements. The normal stresses on the cleavage planes, at the time of fracture, were calculated as follows: First, from the fracture radius, R,, the orientation angle, x,, of the (0001) plane at the outer fiber was calculated. If x, = angle between the (0001) and the rod axis prior to bending, d = rod diameter, and X, = angle between the (0001) plane and a tangent to the bent crystal, then1 sinX, _ 2R, sin xn 2R, + d Since the glide shear strain is given by&apos; y &apos; cot X, — Xo the glide strain at fracture, x,, can be calculated. Next, it is necessary to know the tangential stress in the outer fiber at fracture. This depends on the plastic flow curve, so a series of tensile tests were performed at —196°C on another set of zinc crystals. A curve of the average resolved shear stress versus glide strain was prepared. From this curve, shear stresses, r,, corresponding to the glide strains in the bend tests were taken. These values are listed in Table I. Then the normal stresses in the outer fibres, u,, were calculated from the relation U, = r, tan X, This analysis assumes that the fiber stresses in a bent beam are the same as the stresses in simple

Jan 1, 1959

By J. J. Gilman

THEORIES of fracture in metals have been based primarily on the dependence of fracture strength on grain size. This dependence was first observed by Masing and Polanyi.l These workers, as well as 0rowan12 and Davidenkov and Wittman,&apos; thought the grain size could be equated with the crack size in the Griffith condition for crack propagation.4 In this way they could account for the observed proportionality between fracture stress and the reciprocal square root of the grain size. The chief results of recent investigations of the fracture strength of polycrystals are summarized in Fig. 1. Low for iron,5 Petch for iron,&apos; Greenwood and Quarrel for zinc, Petch and Zein for zinc,&apos; and Hauser, Landon, and Dorn for magnesium&apos;&apos; have all found that the fracture stress of a polycrystal is proportional to the reciprocal square root of the grain size. Also, they find that, for large grain sizes, the fracture stress lies near the extrapolation of the yield-stress curve determined at small grain sizes; where measurable plastic strain precedes fracture. Low9 has gone a step further and shown that, in Range 1, the fracture stress in tension is almost identical to the yield stress in compression for a given grain size. Homes and Gouzou, 10 and Low also have shown that, in Range 2 of Fig. 1, microcracks form when the yield stress is reached, but they do not cause failure until the stress reaches a much higher value. The fact that yielding always precedes fracture suggests that fracture is induced by plastic flow. A mechanism by which this might occur was first proposed by Zener.11 He pointed out that blocked dislocations would cause stress concentrations which might cause fracture. The blocking would be caused by grain boundaries. This idea has been developed by Eshelby, Frank, and Nabarro,12 by Koehler,13 and by Stroh.14 By means of etchpits blocked dislocations have been observed in brass by Jacquet,15 in zinc by Gilman,16 and in lithium fluoride by Gilman and johnston. The blocked-dislocation theory and the experiments that bear on it have been summarized by Mott.18 One of the strong points of the blocked-dislocation theory is that it can explain the grain-size dependence of the fracture stress in a natural way. However, Low9 has pointed out that the theory is not satisfying when it is considered quantitatively. Gilman10 also has criticized the theory and presented an alternative interpretation of the relation between fracture stress and grain size. Since the blocked-dislocation theory is not the only theory that can explain in the grain-size effect, it is necessary to find other evidence either to support or deny it. In this investigation the following points were studied: 1) If blocked dislocations normally cause fracture, then the fracture stress ought to be very high in the absence of plastic glide. Is it? 2) If dislocations must be blocked in order for metals to fracture, then certain types of grain

Jan 1, 1959

By N. E. Rogen, N. J. Grant

AGE-hardening in nickel-chromium-titanium-aluminum alloys in the composition range characterized by the Nimonic alloys, is dependent upon the precipitation of the Ni3(AI,Ti) (y&apos;) phase.1 This phase is found to precipitate at about 700°C.1 Ni,Al is a face-centered-cubic ordered structure with a lattice parameter of 3.560 A. The atomic configuration is such that the aluminum atoms occupy the cube corners and nickel occupies the cube faces with titanium substituting for the aluminum preferentially. A maximum of three out of every five aluminum atoms can be replaced by titanium atoms in the -y &apos;phase.&apos; When the solubility of the y&apos;-phase is exceeded for titanium the T(Ni,Ti) phase is formed. The T-phase is an ordered close-packed hexagonal intermetallic structure which shows little solubility for the other elements in these alloys. The lattice parameter for T-phase is a - 2.5454 kr and c = 8.2900 kx units."

Jan 1, 1959

By R. O. Williams

A study utilizing neutron diffraction, electrical resistance, and chemical extraction has been carried out to clarify the chromium-rich part of the miscibility gap previously proposed for this system. This gap appears to be more symmetrical than previously supposed and is very flat. This flatness is believed to be the result of a cooperative phenomenon perhaps of magnetic origin; in particular, it is believed that the chromium-rich phase is antiferromagnetic. The presence of antiferromagnetism is believed to result in an additional strengthening mechanism very sinzilar to that for ordered precipitates. Likewise, it is proposed that the quenched alloys which are known to be ferromagnetic have antiferromagnetic regions which accounts for the maximum in hardness around 75 pct Cr. It is also believed that the existence of antiferromagnetism in other alloys of transition elements might explain some of their anomolous characteristics. THE iron-chromium phase diagram at low temper atures has recently been constructed as shown in Fig. 1(a)1 based upon that work and the work of Fisher, et al2 This system exhibits a number of extraordinary features but perhaps the most striking is the odd shape of the miscibility gap (in the

Jan 1, 1959

By E. I. Alessandrini, A. U. Seybolt

Selective oxidation (preferential oxidation of a single alloy component) can be an important oxidation mode in the case of alloys such as Si-Fe, Al-Fe, and so on, when heat-treating in atmospheres of low oxygen potential. It was the purpose of this work to identify the oxide films formed on iron containing 31/4 pct Si at low oxygen potentials, and to make some observations on the H2O/H2 ratios required to form them. EXPERIMENTAL METHOD HYDROGEN of varying water content offers a very convenient method of establishing a wide range of oxidizing power, or oxygen potential. The method chosen to provide a desired H2O/H2 ratio was to nearly saturate hydrogen at room temperature by bubbling the gas through water at room temperature, and then to wring out the excess water by passing the wet gas through a long glass spiral immersed in a refrigerated liquid at a desired dew point. After oxidation, the samples were examined by electron diffraction in order to identify any film which might be present. PREPARATION AND TREATMENT OF SAMPLES The samples were prepared from a Si-Fe alloy made in the laboratory using electrolytic iron and 96 pct ferro-silico11, vacuum-induction melted, hot-rolled then cold-rolled to 0.015-in. thickness. After the cold-rolling the samples were annealed at 1000°C. Transmission Laue X-ray patterns showed very little evidence of preferred orientation, but it is likely that a weak texture could have escaped notice. The chemical analysis was as follows: 3.34 Si, 0.01 Al, 0.001 0, 0.006 C, 0.002 N. Before each sample was oxidized, it was polished through 4/0 emery paper, cleaned and dried using benzene, carbon tetrachloride, acetone, absolute alcohol, and ether. The method of oxidizing the samples was as follows: Samples about 1/2 by 1 by 0.015 in. were held in the upper, cool part of a vertical vacuum furnace until the desired temperature was established in the furnace, Fig. 1. The sample was then lowered into the uniform hot zone where it was heated quickly to the desired temperature, and subjected to hydrogen of controlled moisture content (dew point) as mentioned previously. Temperatures from 400" to 1000°C were used, and dew points from room temperature to —78°C were explored. In general, the runs lasted about 6 hr, but a few runs were of different duration, as will be indicated later. After the predetermined period of time had elapsed, the specimens were withdrawn by the hollow stainless steel-quartz holder to a cool zone at the top of the furnace. EXPERIMENTAL RESULTS The results are summarized in Table I. Most of the electron-diffraction patterns were obtained by glancing-angle reflection, but in a few cases it was possible to strip the oxide films from the surface by a total immersion in a 3 pct bromine-methanol solution. Some of the thinner stripped films yielded quite satisfactory transmission patterns. However, both diffraction techniques yielded the same results; the patterns while most frequently showing a (low) tridymite, also showed a (low) cristobalite in some cases. A few patterns showed evidence of both phases. Because of the known nonequilibrium behavior of silica, it is not surprising that both forms of silica could coexist on one sample. It was noted that at fairly low dew points such as —40°C, there was a tendency for diffuse diffraction rings, suggesting a very small crystallite size. Also some subsequent experiments to be reported later show that thin silica films can be amorphous.

Jan 1, 1959

By M. J. Sinnott, J. C. Tobin

A technique for collecting and analyzing the small quantities of gases desorbed on heating metal powders has been developed. The gases collected from copper powders of various types of manufacture have been analyzed and found to contain H2O, CO2, and SO2. The quantities of gas evolved vary from 0.05 to 1.0 milliliter, under standard con-ditons of temperature and pressure, per gram of powder. The source of the gas has been traced to surface films on the powder. IN the powder-metal industry it is recognized that the presence of gases in the powders affects the processing and resulting properties of the sintered product. Several sources for these gases are possible; namely, chemical reactions with the sintering atmospheres, inclusion. of gases during compacting, dissolved gases in the metal powder, and sorbed gases on the metal-,powder surfaces.1, 2, 3 Since the quantities of the gases are relatively small, no direct attempt has been made to collect and analyze them nor to isolate their source. The work that has been done has been concerned with the dimensional stability of the product., its porosity, density, or mechanical and physical properties. Conclusions are drawn on the basis of studies of this type as to the probable effect of gases on these properties. 4, 5, 0 In the present investigation an attempt has been made to obtain quantitative data on the analysis and quantities of gases associated with copper powders manufactured in different ways. Apparatus was constructed to collect gases evolved from loosely packed powder samples on heating to various temperatures under a reduced pressure of helium. The powders were outgassed thoroughly in a vacuum prior to starting the collection operation in order to eliminate mechanically entrapped gases and to prevent the possible reaction of the powder with these gases. Experimental Procedures Materials—The major portion of the investigation was carried out on a single heat of water-atomized, oxygen-free, high-conductivity copper. A relatively coarse fraction of this powder, 20 to 100 mesh, was tested as a mixed powder and as various graded size fractions. The microstructure of this powder showed equiaxed grains and was clean and free from oxides and oxide eutectic. The particle size, as determined

Jan 1, 1959

By F. Weinberg

It has been suggested, for some time, that the behavior of metals under high-temperature creep conditions is strongly influenced by the behavior of the grain boundaries present in the material. This has led to a number of investigations1-6 in which attempts were made to measure the properties of the grain boundaries themselves, both to determine their contribution to the creep process and to obtain further knowledge of their intrinsic structure and properties. The present investigation was undertaken to extend the observed data under more closely controlled experimental conditions and to attempt to clarify sonie of the discrepancies in the literature. A brief report of the preliminary results has been published.7 EXPERIMENTAL PROCEDURE For the present investigation, controlled orientation specimens were used having plane boundaries which terminated at the specimen surface. Trierys-tal specimens were used instead of the bicrystal type of Puttick and King,&apos; in order to simplify the gripping technique required to apply a shear stress along the boundary. The aluminum used was obtained from the Aluminum Francaise and was of 99.994 pct purity, containing 0.002 pct each of Cu, Fe, and Si. The aluminum tricrystals were grown horizontally from the melt, in graphite boats, and in an argon atmosphere, following the technique developed by Chalmers.&apos; The orientation of the tricrystals was controlled by using seed crystals separated by mica inserts. Thermal conditions were adjusted to produce tricrystals having boundaries which were macroscopically straight, parallel to the growth direction, and perpendicular to the top surface. Following growth, the tricrystals were carefully cut with a jeweller&apos;s saw to the test specimen shapes shown in Fig. l(a) and l(b). In Fig. 1(a), the boundaries subjected to shear stresses are indicated AB and CD: Points A and C are the ends of the mica inserts separating the seed crystals; points B and D are at lateral cuts made in the specimen. It was found that the slots left by the mica inserts at A and C had to be enlarged slightly, by cutting, to ensure that the boundary terminated freely at these points. A shear stress was applied to the boundarks by gripping the upper two legs in file-faced Inconel grips and suspending a weight from the lower part of the specimen held in similar grips, as shown schematically in the diagram. The test specimens shown in Fig. 1(b) were cut from long

Jan 1, 1959

By N. J. Grant, F. B. Cuff

DURING a recent investigation of high-purity aluminum,&apos; evidence was found for grain-boundary sliding and migration at room temperature. The investigation was concerned primarily with an electron-microscopic study of a fine (approximately 1 micron) network that appeared on the surface of chemically polished aluminum. The nature of the network will not be discussed since it serves only as markers showing relative displacements in this instance. The pertinent observation, shown in Fig. 1, was made on a flat tensile specimen which had been strained about 8 pct at room temperature. As can be seen in this electron micrograph, a grain boundary has undergone a certain amount of sliding. The reason for the existence of the two differently appearing networks is given in the foregoing reference. For present purposes, however, the transition region can be disregarded. A close study of this picture reveals that in addi- tion to sliding there is a narrow band adjacent to and above the grain boundary. Fig. 2 is a schematic drawing showing the essential features of this band.

Jan 1, 1959

By J. E. Hove

The high temperature physical properties of graphite are reviewed and interpreted in the light of present day knowledge of the mechanisms affecting these properties. The thermal and mechanical behaviors only are discussed and, whenever possible, comparisons are made with other refractory materials. Possible further studies are indicated, including some carbide work. AS long as the term high temperature implied only temperatures up to about 1000°C, the materials problems which arose could usually be handled by fairly conventional metal alloy types, such as the Co-Cr-Ni superalloys, for which there exists a great deal of technology. Perhaps this temperature can still be considered an upper limit for normal applications, but it is certainly true that the number of abnormal applications is increasing rapidly. The advent, in recent years, of ram-jet and rocket missiles and of high power nuclear reactor heat sources has raised a host of questions concerning the basic problem of what material to use in the temperature range up to 2000°C and higher. While there are, of course, many metals, in the second and third transition series, which melt at considerably higher temperatures than this, these metals are, at present, pretty well excluded from practical use by other considerations, such as re-crystallization, chemical activity, or excessive plastic deformation. The behavior of metals, from the standpoint of dislocation theory, is just beginning to be understood and thus there is some hope for the future development of very high temperature metals, but the immediate problems would most logically appear to have solutions involving the nonmetals, such as the refractory ceramics and graphite. For this reason, there is presently a great deal of engineering and experimental research being performed on the latter materials, much of this research being exploratory in the sense of gathering new property data. The situation, so far as graphite is concerned, is somewhat more fortunate than with the other refractory solids in the sense that a great deal is already known about its basic properties. This stems both from the fact that the carbon-carbon bond has been of interest to chemists for a long time (and graphite can be considered as a very large aromatic molecule, if desired) and the fact that its properties, both as a function of temperature and of radiation damage, are of critical importance to nuclear reactor designers. It is still true, of course, that such fundamental questions as why graphite remains solid to such a high temperature and why it has such a high thermal conductivity cannot entirely be answered at the present time. It is, nonetheless, meaningful and instructive to consider such problems in the light of existing knowledge. This is what the present paper will attempt to do. Before going on, it may be appropriate to classify graphite and justify its discussion before readers primarily interested in metals. Graphite is comparatively unique among materials in that there is always a property or group of properties which precludes calling it either a metal, a semiconductor, or a ceramic. It has the high electrical and thermal conductivities of a metal, but the artificial, poly-crystalline types show a negative thermal coefficient of electrical resistivity, generally characteristic of semiconductors. On the other hand, semiconductors, by definition, show an ever increasing resistivity as the temperature is lowered, whereas graphite approaches a finite, and, indeed, a rather low resistivity in the region of 10°K and, furthermore, a good single crystal of graphite has a positive temperature coefficient, as for a metal.&apos; On still another hand, its porosity and brittleness at lower temperatures would put graphite in the ceramic class although, unlike most ceramics, it is readily machinable and has a high resistance to thermal shock. All things considered, it is probably more nearly appropriate to call graphite a metal than anything else. Although certainly outstanding in some ways, graphite has its peculiarities and, especially if the reader is unfamiliar with the data, it is probably valuable to review some representative property variations at high temperature. This review is meant to be mainly illustrative and no attempt has been made to be exhaustive. Review of High Temperature Properties At ordinary pressures, graphite does not melt, but sublimes directly into the gaseous phase at about 3700°C. Although the phase equilibrium diagram is still in some doubt, graphite will melt, at a slightly higher temperature, at pressures in excess of about 100 atm. The chief difficulty of using graphite in an oxidizing atmosphere is that the reaction rate becomes quite high at fairly low temperatures. If a threshold oxidation temperature2 is defined as the temperature at which graphite loses 1 pet of its weight in 24 hr, the value in air is 450°, the value in steam is 700°, and the value in carbon dioxide is 900°C. Efforts are presently being made to raise this threshold temperature either by impregnation with a retardant of some type (sodium tung-state and phosphoric acid, for example) or by a suitable metal or oxide coating. It is probably fair to say that, to date, these attempts have not shown any outstanding success in all respects. Most commercial graphites are fabricated by impregnating a carbon flour (say of coke or lampblack particles) with some type of hydrocarbon

Jan 1, 1959

By J. W. Evans

ONE possible method for the separation of hafnium from zirconium is by the vapor phase oxidation of the purified mixed chlorides. An important factor for success depends on the difference in free energies of the following reactions: ZrCl(g) + O2 = ZrO2(s) + 2C12 [1] HfCl4(g) + O2 = HfO2(s) + 2C12. [2] Prakash and Sundaram1 have calculated these values from certain data and carried out experimental work with partial success. However, their free energy values for reaction 2 diverge rather widely from other data which are presented here. The possibility of successful separation is briefly considered in the light of these alternative values. Table I shows the free energy values (?F°) for reactions 1 and 2 derived from the recent compilation of data on oxides and chlorides by Glassner.2 The corresponding values of Prakash and Sundaram for the same temperatures are also shown. The alternative values to those of Prakash and Sundaram for reaction 2 are substantially confirmed by the data of Kubaschewski and Evans" and Cough-lin&apos; on heats of formation, entropies, and heats of vaporization. In spite of the considerable uncertainty in the hafnium data it is reasonable to suppose that the oxide-chloride free energy difference for two such chemically similar elements would not be very far apart. These figures indicate clearly that the oxidation of hafnium chloride is appreciable at all temperatures. Nevertheless, there is a significant difference between the values for reactions 1 and 2

Jan 1, 1959

By E. Goldin, H. J. Juretschke

THE electrical properties of porous conductors have been studied ever since such materials, usually prepared by pressing and sintering of metallic powders, obtained practical importance. In most cases, however, the results have not been easy to interpret because of the great number of independent variables such as powder size, pressure, pressing rate and temperature, subsequent thermal history, porosity, and so on, which enter into the fabrication of industrially desirable materials. In recent years an interpretation of these properties has become of interest because the electrical properties of many new substances, prepared as fine powders by crystallization from the vapor phase or by reactions in the solid state at high temperatures, have to be determined on porous samples formed from the powder. Since theoretically significant information must refer to the fully dense state of the material, it is important to relate the electrical properties measured on such porous samples to those expected for a sample of 100 per cent density. A simple theory for the porosity dependence of the electrical conductivity and the Hall effect was recently proposed by Juretschke, Landauer, and Swanson.&apos; Their model, assuming a distribution of pores of equal size and shape, sufficiently far apart so that each pore is embedded in a region of uni-

Jan 1, 1959

By J. L. Gissy, G. H. Schippereit, J. B. Schroder

THE existence of mechanical anisotropy in hexagonal metals is well known.&apos; One possible method of measuring the degree of anisotropy in a metal is to make hardness measurements on the different planes of a single crystal. So far as is known, no single-crystal hardness data have been reported in the literature for alpha titanium. To obtain an estimate of the variation in hardness with crystallographic direction in titanium, Knoop hardness measurements were made on single crystals of electrorefined alpha titanium obtained from the Bureau of Mines, Boulder City, Nev. The hardness of arc-melted ingots prepared from this material was 64 Bhn (1000-kg load, 30 sec). Crystals of sufficient size were mounted in epoxy resin with the basal plane exposed. They were then polished and etched to remove excess plastic resin and work-hardened surface material. The crystals were then examined by x-ray diffraction, using the back-reflection Laue technique to insure that the crystals were properly oriented and strain free. After Knoop hardness readings (100 or 200-g load) were taken on the exposed surface, the crystals were remounted to expose a prismatic face and the foregoing procedure was repeated. Hardness values for various orientations are listed in Table I. The data show a difference between the hardness of the basal planes and that of the prismatic planes, as anticipated. More striking is the hardness difference observed between different orientations of the indentor on the prismatic planes as shown in Table I and Fig. 1. All of the indentations shown in Fig. 1 were made with a 100-g load. The long axis of indentation A is parallel to the unique (c) axis of the crystal. The deformation marks visible around indentations A and C of Fig. 1 are probably slip traces of the (10.0) system reported by Rosi, et al. The greater hardness observed when the long axis of the indentor was perpendicular to the unique axis of the crystal (indention B) is probably caused by the opera-

Jan 1, 1959