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By G. L. Fisher, R. Maddin

The rate of self-diffusion in silver single crystals during torsional strain was measured over a wide range of tenperatures and strain rates. The torsional strain was applied in a cyclic manner by reversing the divection of twist aftev a certain fraction of the total strain was reached. Surface roughening was substantially reduced in cyclically strained specimens compared with those strained entirely in one direction. Diffusion coefficients were determined using a special sectioning technique whereby slices could be removed parallel to the surface of a strained torsion specimen. The accuracy of the technique was first determined using roughened specinzens statically diffused to various penetration depths. Simultaneous strain-diffusion experiments were then performed over the temperature range 550' to 800 using strain rates of lo-' to 2.2 x 10' sec-'. The increase in silver self-diffusivity during torsional plastic strain was small. In only one case was the enhancement greater than twice the static value of the diffusion coefficzent. The enhancements increased slightly with decreasing tenzperature and increased with increasing strain rate. A dislocation pipe mechanism of enhancement was found to explain the present results satisfactorily. The magnitude of the enhancements and the tenzperatures at which they were observed suggests that a steady-state dislocation density of 10' to 109 dislocations cm-2 existed during the high-tempevature deformation. The diffusion coefficient in an fcc metal can be increased by 1) an excess concentration of point defects and 2) the rapid movement of diffusing atoms along grain boundaries and dislocations. Plastic deformation may be responsible for enhancements by both of these mechanisms. The results of simultaneous strain-diffusion experiments have caused a controversy over the magnitude of the effect of strain upon diffusion. Experiments reporting large enhancements have been criticized for the manner in which the diffusion coefficient was measured. Several of these experiments have been contradicted by other results obtained in a similar manner. No enhancement was reported by two investigators''' who performed their experiments at high temperatures where strain is expected to have little effect upon diffusion. A recent theoretical analysis3 indicates that strain enhancement by a vacancy mechanism will be small and exist only at low temperatures and rapid strain rates. It is doubtful whether any experiment has yet proven that strain enhancements can occur by an excess vacancy mechanism. In the present research, experiments were performed at extremes of temperature and strain rate in an attempt to measure enhancement of diffusion by strain. Self-diffusion of silver during torsional deformation was chosen since two separate investigations4j5 indicate that at least a small enhancement exists under these conditions. An accurate method of determining the diffusion coefficient, mechanical sectioning, was employed after a technique was devised to apply it to the roughened surface of a strained torsion specimen. First, the accuracy of the technique was determined for different penetrations (2fi). Then diffusion experiments were performed at temperatures between 550" and 800°C at strain rates of 10"" to 10"" per sec. I) EXPERIMENTAL PROCEDURE Preparation of Specimens. The experimental procedure detailing the growth of crystals, their treatment, and the annealing techniques during straining are contained in earlier publications.496 The dimensions and orientations of the torsion specimens are shown in Figs. 1 and 2. licatiin of Torsional Strain. The torsion procedure was changed from that used previously. The specimens were strained in cycles to reduce the gross distortion and surface roughening inherent in twisting to high strains in one direction. For all strains above 0.30, the total surface strain was obtained as follows: the specimen was twisted to 0.10 strain in a clockwise direction; the direction was reversed; the backlash in the torsion apparatus was taken up; and the specimen was twisted in a counterclockwise direction through the initial starting point to 0.20 surface shear strain in the opposite direction. The strain reversal took no more than 1 min to accomplish. The total accumulated surface shear strain was arbitrarily taken as the sum of the 0.10 strain in the clockwise direction plus 0.20 in the counterclockwise direction. Although cyclic torsional strain was beneficial, surface roughening and specimen distortion were not entirely eliminated. The variation in the diameter at various points on the specimen was often much greater than the depth of the diffusion zone. Since serious errors might result if specimens were sectioned in the usual manner, a technique of lathe sectioning was devised to allow sections to be taken parallel to planes of equal concentration.

Jan 1, 1967

By C. G. Dunn

This note describes a method for revealing micro-structures in tantalum through an {hkl} effect which was found during a study involving amorphous and crystalline tantalum oxides. Three steps are involved. First, an amorphous film of Ta2O5 is formed anodically1 on the surfaces of the tantalum. The amount needed depents on the thickness of the substrate tantalum; 1600Å of oxide, however, is generally satisfactory in the 1/2 -mil (12 µ) to 3-mil thickness range. The film at this point is uniform in color. Second, the anodized specimen is heated to 800°C in. vacuo during a period of a few minutes and held at temperature for 5 min. During such a treatment some oxygen leaves the oxide through dissolution into the substrate tantalum (e.g., at 10-8 Torr oxide films thinner than about 800 completely disappear on 5-mil strip). Considerable crystallization of the amorphous oxide also occurs adjacent to the metal/oxide interface but not throughout the oxide thickness. The thickness of oxide needed as a residue is not too critical; values between 500 and 1000Å (optical estimate') may be most useful. (The presence of a residual oxide layer after the heat treatment usually is not apparent; in fact, in most instances the surfaces have a metallic appearance. This can be the case also for thicker residual films needed in the detection of a crystalline oxide phase by surface X-ray diffraction—e.g., if the initial oxide film is 3000Å thick, enough crystalline oxide is formed for X-ray detection.) The third step eliminates the metallic appearance at the surface and reveals (or produces) a varied colored film of tantalum oxide. Two methods can be used and both will be illustrated. If the specimen is dipped in 48 pct HF for a few seconds, the tantalum is found to be covered with a multicolored oxide. Moreover the colors correlate with the grain structure of the substrate tantalum through an orientation dependence of the type illustrated in Fig. 1, where results of an orientation analysis by the Laue X-ray method have been plotted. Other film conditions would have provided a different set of colors for the same orientations. Any such correlation, however, provides a basis for a qualitative assessment of the distribution of orientations in a large sample. The cold-rolled structure of tantalum foil produced by a recrystallization anneal and a cold-rolling reduction of 95 pct is not easily obtained by usual metal-lographic methods.' The structure, however, is easily revealed by the present technique, Fig. 2. The alternative third step is to reanodize the specimen. The optical nature of the final anodic film depends on the presence of the crystalline oxide phase; color variations indicate optical thickness variations in the film. As the reanodization voltage is increased from zero, straw and light brown colors appear (as usual), followed now by variations in color from grain to grain depending on {hkl} orientation. Suitable end points, which are determined by voltage and time, are easily established. Fig. 3, for example, shows the grain structure developed in 5-mil tantalum strip during a 30-min re-crystallization and grap-growth anneal at 1800°C. An anodic film of 1600Å thickness was added prior to the special 800°C anneal. Reanodization was done at constant current density (near 0.2 ma per sq cm) in an aqueous solution of H3PO4 (0.01 pct) and was terminated after the initial brown colors were replaced by a deep blue at approximately 30 v. Although good grain contrast can be obtained with bright field illumination, even greater contrast for black and white micrographs may be obtained with polarized light, as in Fig. 3.

Jan 1, 1967

By S. J. Rothman, J. J. Hines, D. Rokop, R. Bastar

Self-diffusion has been measured along each principal axis of mosaic-structured and relatively perfect a uvaniuln single crystals. The anisotropy reported before,' D[loo] = D[001] » D[010 , has been confirmed. and is shown to represent volume diffusion. Diffusion along the dislocations in the imperfect crystals increases the measured value of D significantly without changingr either the shapes of the penetvation plots or the anisotropy. SINCE a uranium is orthorhombic, the diffusion coefficients along the principal crystallographic axes need not be equal. A previous measurement of self-diffusion in three mosaic-structured a uranium single crystals1 showed that Dil00| =D[001 >>D[010 This result was deduced from penetration plots similar to those in Fig. 1. That is, the plots for diffusion along [loo] and [001] were the straight lines expected from the solution of the diffusion equation for the boundary and initial conditions of the experiment, while the plot for diffusion along [010] was curved so that the part near X = 0 had a much steeper slope than the straight lines did. The observed anisotropy could have been due either to volume diffusion or to diffusion along the very large number of dislocations in such crystals, according to the model proposed by aart.' The authors of Ref. 1 argued that the latter possibility was unlikely, but did not exclude it. To settle the question, a simultaneous measurement of self-diffusion in mosaic-structured and mosaic-free crystals was carried out. This paper reports the results, as well as a measurement of the temperature dependence of self-diffusion in a uranium. EXPERIMENTAL The single crystals of a uranium used in this experiment are listed in Table I, together with other experimental details. '(Mosaic structured" here means crystals grown by solid-state phase change.3 The Laue spots from such crystals spread over about 6 deg of arc, which indicates that the density of dislocations is very large. "Perfect" crystals are ones grown by Fisher's grain-coarsening tehnique; the Laue spots from these are quite sharp,4 which shows that the dislocation density in these crystals is much lower. All of the crystals were made from high-purity uranium; a typical analysis is given in Ref. 1. The actual diffusion directions were within a few degrees of the nominal. The thin layer-sectioning method was used. The experimental techniques were similar to those described in Ref. 1, except the sectioning. Crystals Za, 2b, 2c, 4a, 4b, and 4c were sectioned on a lathe. Crystals 3a, 3b, and 3c were sectioned by grinding on glass plates, with 1 p diamond powder in butyl alcohol as the abrasive. The use of glass plates eliminated the crowning noted in Ref. 1, and the use of butyl alcohol as a lubricant simplified the chemical processing of the sections. Material from the sides of these three crystals was removed by spark machining before sectioning. The sectioning of crystals 4a and 4c was botched. A small cylinder of metal was extruded from the center of crystal 4a while turning off material from the side of the sample. Crystal 4c was misaligned by 2.1 x 10"3 cm because the dial gage jammed during the alignment. Notwithstanding these errors, useful data were still obtained from these crystals. The points

Jan 1, 1967

By O. J. Kensok

THE literature contains three previous studies of the chemical activities of Cd-Mg alloys: Trumbore, Wallace, and craigl obtained chemical activities of magnesium at 543" by performing electromotive-force measurements; Borg and irchenall calculated activities at 590°K based on vapor-pressure measurements of cadmium over Mg-Cd alloys by the Knud-sen effusion method; and Terpilouski and arcka measured the chemical activities of magnesium in the liquid state by means of electromotive-force measurements. In this study the vapor pressure of cadmium was determined over nine solid Cd-Mg alloys in the temperature range 571" to 697°K by the Knudsen effusion method. This study was prompted by the lack of precision in the data reported by Borg and Birchenall. Knudsen cells used during this study were fabricated from 0.020-in.-thick tantalum tubing. The tops and bottoms of the cells were cut from 0.010-in.-thick tantalum sheet, press-fitted, and welded in place. The orifice was drilled in the cell top and subsequently reamed to a knife-edge thickness. The diameter of the orifice was measured with a traveling microscope capable of reading 0.0001 cm. The sample was heated by means of a resistance

Jan 1, 1967

By Charles E. Lundin, Albert S. Yamamoto

The phase eqrtilibria of the binavy system, Nb-Al, weye established by metallographic, X-ray diffraction, hardness, and thermal-analysis techniques. Three intermediate phases were identijied, Nb3.41, Nb2Al, and NbAl,. The N6,Al and NbAl phases have extensiue ranges of solubility. The maximum solubility of aluminum in niobiuitz is 23.0 i0.5 at. pct (8.0 wt pct) A1 clt 1960" i 10°C. The maximumt solinbility oj niobiunz in alunzinutz is less than 0.3 i0.1 at. pct (1.0 tut pct) fib. The A'h,Al and Nbgl interinediate phases forr peritectically at 1960° 1UcC and 1870" i 1U9C, respectively, zuhereas the NbAl, phase fornts congvuently fio11t the melt at 1605" 10°C. A eutectic reaction oc- The intent of an over-all research program was to develop niobium-aluminide-base coatings for refractory metals since these coatings exhibit useful high- curs at 57.8 i 0.5 at. pct (28.5 wt pet) A1 and 1500" 10°C between NbAl and NbAl3. The aluminum-rich region of the system has a peritectic reaction where liquid reacts with NbAL, to form aluminum solid solution. Structural and lattice -pavarneter data /or the intermediate phases were determined by X-ray diffraction analyses. An unusual transfov?nation having charactevistics of nucleation-growth and martensitic nleclanisils occurs on cooling the niobium solid solution in the region from 18 at. pct (6 wt pct) Al to 23 at. pct (8 wt pct) Al. The transfornzed product on rapid quenchirzg is an acicular NhAl. temperature resistance to oxidation. As a prerequisite, the phase equilibria of the system, Nb-A1, were necessary. A literature survey revealed that the phase-equilibria data reported by various investigators were rather incomplete and conflicting. Therefore a complete determination of the system was undertaken by the authors from 0 to 100 pct of solute addition.' In 1959 Savitskiyz presented a tentative diagram which

Jan 1, 1967

By J. E. Pavlick, W. W. Mullins, H. W. Paxton

The growth of nitrogen-austenite during nitriding of large-gvained ferrite between 650" and 800°C has been studied as a functimz oJ time and nitrogen potential of the atmosphere for a variety of alloying elements dissolved in the ferrite. Whereas no preferential grain boundary penetration by aistenite takes place in zone-refined iron, the addition of aboit I zot pct oj' V, Ti, Cr. Si, or A1 produces significant protrusion of austenite down former 0-0 boundaries. This group of elernents also produces ragged y-0 interfaces in the a grains which advance more slozuly than those in zone-refined iron. 1 zut pct of Ni, Co, Mn, or Mo dissolved in ferrite leads to essentially smooth 7-a interfaces, which advance at rates very sinzilar to those in zone-refined iron. The reasons for this effect are not clear. Two possible limiting hypotheses, not necessarily exclusive, are discussed. One involves p.referentia1 grain bozindary diffusion of nitrogen. the other the production of a zone of ferrite, stpersatcirated with respect to nitrogen, from which austenite precipitates helerogeneozisly on the a-a gyain boundaries. In many phase transformations which occur by nu-cleation and diffusion-controlled growth, information on the parameters which control the rates of nuclea-tion and of growth is frequently not quantititative. The present work describes a continuation of efforts to obtain quantitative data on growth rates in systems where thermodynamic and kinetic factors are known independently, and thus where comparison with models of the growth process is possible. The work has been conducted on iron-base systems for two reasons. Firstly, there are several interesting transformations in this system, many of which are of practical importance. Secondly, the information on phase diagrams, activities, and diffusion coefficients is more complete in this system than in any other. These data are still inadequate to check all the effects of alloying elements, but at worst it is possible to make estimates and thus examine order-of-magni-tude agreement with proposed models. It will be necessary to do this in the present work. The problem studied was the growth of nitrogen-austenite in from the surface of large-grained columnar crystals of alloyed ferrite exposed to an NHs/Hz atmosphere. Earlier work by Grozier et al.' had suggested that growth of austenite may occur preferentially along grain boundaries of alloyed ferrite, whereas zone-refined (ZR) iron shows no such effect, Fig. 1. Growth of nitrogen-austenite into single crystals of zone-refined ferrite is planar provided sufficient nu-cleation sites for austenite are present, and is controlled by volume diffusion of nitrogen through the

Jan 1, 1967

By P. M. Robinson, J. S. LI. Leach

The heats of solution oj' indiurrr, tin, lend, nrzd tellurium have been calculated from the measured heat effects when mechanical mixtres of indium and telLuium tin and tellurium, and lead and tellurium were added to liquid bismuth. The results are in good agreement xith publislzed values.s for the separate sollction of each eleltzent in bismuth. The heats oj solution of cadmium and tellurium calculated from the rneasuved heat effects on adding trechanical mixtures of these elements do not ugree zc,itl the published values jbv the separate solution of each element. It is shown that at 623°K Ile interaction between cadmium and tellurium dissolved in liquid bismuth is strong enough to led lo preciPitation of solid CdTc. The heats oj- jor-mation of CdTe at 273" nd 623°K (1)-c crilculated fi-or the measured heat ejlfecls. The calcnlaled az'erage deviation from the Kopp-l\'ez?,zunrz rule fov solid CdTe is less than 0.06 cat per g-atom- C over this lertzperalure range. Tlze importance 0.f these oDserl.ations to the determination of heals of formation hy metal solution calorimetry is considered. LIQUID metal solution calorimetry is a convenient method for determining the heats of formation of solid compounds. In this technique the heat of formation is the difference between the measured heat effects on dissolution of the compounds and of mechanical mixtures of the components in the liquid metal.' The heat of solution of the mechanical mixture may be calculated from the measured heat effect. At infinite dilution of the solutes, this heat of solution is equal to the sum of the heats of solution of the separate components. If the heat of solution of one of the components is known, the value for the other can be derived; if both are known, they may be used to check the accuracy of the calorimetric technique. The heats of formation of the tellurides of cadmium, indium, tin, and lead have recently been measured by metal solution alorimetr. The heats of solution of indium, tin, lead, and tellurium at infinite dilution in liquid bismuth at 623"K, calculated from the measured heat effects on solution of the mechanical mixtures, are in good agreement with the published values. The heats of solution of cadmium and of tellurium calculated from the measured heat effect on solution in bismuth at 623'K of mechanical mixtures of cadmium and tellurium, however, do not agree with values estimated from the literature. 1) EXPERIMENTAL PROCEDURE AND RESULTS The Heats of Solution of Indium, Tin, Lead, and Tellurium in Bismuth. The heat effects were measured when mechanical mixtures corresponding to the compounds In,Te, InTe, In2Te3, In2Te5, SnTe, and PbTe were dissolved in bismuth. The calorimetric procedure and the method of calculation have been described elsewhere.' The heats of solution of the mechanical mixtures were obtained by subtracting the change in heat content per gram-atom of the sample between the addition temperature (273°K) and the bath temperature (623"K), (H623°K - H273°K)S, from the measured heat effects. The calorimeter was calibrated with pure bismuth. The reported values of the measured heat effects are based on (HGoK - ^273oK)Bi = 4.96 kcal per g-atom.3 The measured heat effects are found to be linear functions of the solute concentrations of the bath in the dilute solution range. The values, extrapolated to infinite dilution, are listed in Table I, together with the heats of solution of the mechanical mixtures calculated using the published values of (H 623°K - H273°k)s for indium, tin, lead,3 and tellrium. All the error limits quoted in this work represent the spread of values obtained. The heats of solution in liquid bismuth at 623°K of mechanical mixtures of indium and tellurium in four different proportions were determined. Values of the heats of solution of the two components were then calculated from the resulting four simultaneous equations: The heats of solution at infinite dilution of tin and lead in liquid bismuth at 623°K were calculated from the heats of solution of the mechanical mixtures of tin and tellurium and of lead and tellurium using the heat of solution of tellurium calculated above. These values of the heats of solution are listed in Table I1 together with some published values for comparison.

Jan 1, 1967

By J. W. Spretnak, V. Kandarpa, G. W. Powell, R. A. Erickson

Neutron-diffraction pattens were obtained at room temperature from two austenitic manganese steels, oxc wth n carbon content of 1.23 zct PC/ and the olher 0.63 wt pct. Analysis of the data showed that the curOan atols (ions) occupy tile octaheadral site. any of the cavbon is in the tetahedral site, tile concentration is too strzall lo be detecled by tzeulr-on d$facton. AS part of a long-range investigation of eutectoid decomposition, a study of the structure of austenite has been undertaken in order to characterize as best as possible the parent phase in the Fe-C system. Austenite is a solid solution of carbon atoms in fcc iron, the carbon atoms being located in interstitial positions. Of the two kinds of interstitial sites, octahedral (0, +, +) and tetrahedral (4, a, $), in this lattice, the larger octahedral sites are usually considered as the probable position of carbon. The strain energy involved in pytting a neutral carbon atom (covalent di; ameter 1.54)' into a tetrahedral site (diameter 0.56A) is considerably greater than that associated with p2tting carbon into the octahedral site (diameter 1.08A). Dayal and arkeen' observed the migration of carbon in iron under the influence of an electric field indicating tha tye carbon is ionized. pe ionic radius of C'4 is 0.15A and that of C" is 0.29. Consequently, it is conceivable that carbon ions could occupy the tetrahedral sites if one considers only the strain energy. Size considerations alone do not provide a complete explanation for the formation of the interstitial solid solutions, however. because metals such as copper, zinc, and cadmium do not appear to form interstitial solid solutions in spite of their large interstitial sites. Evidence which suggests partial ioni-zation of the carbon in austenite has also been obtained by Speiser, Spretnak, and alor who derived formu- lae relating the observed change in lattice parameter to the effective diameter of the solute atom. Assuming the carbon occupies the octahedral site. they computed the effective diameter of carbon to be 1.33A whereas the covalent diameter is 1.54A. petch5 concluded from powder X-ray diffraction experiments that carbon atoms occupy the octahedral sites. His conclusion was based on the general trend observed in the intensity of the diffraction lines as the carbon content of the austenite was increased. However, the major difficulty associated with application of X-ray diffraction to this particular problem is the relatively poor scattering power of carbon. Williamson and smallman' utilized X-ray diffraction to study the line broadening caused by the strain field about carbon atoms inferrite: these investigators concluded that at least 85 pct of the carbon in ferrite occupies the octahedral (0: 3, B) site rather than the larger

Jan 1, 1967

By Donald B. Hunter

The Ti-Be system up lo 10 wt pct Be with cortlmercialll' pure titanium has been determined using metal-lographic techniques. Beryllium forms the p eutec-toid type of dinqarz with titanium; the eutectoid occurs at 0.5 xt pct Be azd 1500 F, while an eutectic is foritled at 1887°F containing between 5 and 7.5 rt pct Be. The solubility of beryllium in a titanium does not exceed 0.05 zvt pct and the maximum solubility of Beryllim in 0 titanium is 1 wt pet. PREVIOUS work on the Ti-Be phase diagram indicated that additions of beryllium to titanium resulted in a marked lowering of the melting point,''' and it was also observed that the solubility of beryllium in both a and p titanium was retricted.- With the limited data available, McQuillan and McQuillan constructed a partial hypothetical diagram of the Ti-Be system,= assuming that beryllium formed the 0 eutectoid type of system with titanium. Subsequently Bedford published a tentative diagram of the entire Ti-Be system: although he did not investigate alloys containing less than 10 wt pct Be. This diagram included several compounds described by other authorities.'-l3 The following account describes work directed toward determining the form of the titanium-rich portion of the phase diagram up to 10 wt pct Be. MATERIALS AND METHODS Titanium sponge and a master alloy containing 25 wt pct Be were used in the formulation of the alloys; the chemistry of these materials is detailed in Table I. The alloys were prepared as 250-g compacts which were melted into buttons by arc melting under purified argon on a water-cooled copper hearth. Results of analyses for beryllium are shown in Table I but no determination of other elements was made at this stage. The alloys were hot-rolled at 1650°F to produce 0.1-in.-thick sheet. Alloys containing up to 1.5 pct Be were fabricated without trouble; samples containing 2 pct cracked initially upon hot rolling, but thereafter rolled satisfactorily to sheet. The 5, 7.5, and 10 pct were hot short and could not be rolled. Surface contamination was removed from samples by sandblasting and pickling in a 5 pct HF-35 pct HNOs solution. For metal log raphic examination, heat-treated specimens were mounted transverse to the rolling direction, ground on silicon carbide papers of increasing fineness to 600 grit, and electropolished at 20 v using a solution containing 600 ml of methanol, 60 ml of perchloric acid, 360 ml of butyl cellosolve, and 2 ml of Solvent "X". Difficulty was experienced in obtaining good electropolished surfaces with samples containing more than 1 pct Be, the degree of difficulty in-

Jan 1, 1967

By J. D. Wood, J. T. Smith

ALLOYS with a dispersed second phase in a metallic matrix are generally much stronger than the matrix itself. Plastic deformation in dispersion-strengthened alloys is usually confined to the matrix phase when recovery processes are active, while in the absence of recovery both phases may yield.' The alloy system studied in the present research was WC-12 wt pct Co and consisted of noncoherent WC particles dispersed in the cobalt matrix. Some particle-to-particle contact existed but not enough to produce a continuous WC skeleton. The microstruc-ture of the WC particles was characterized by very straight edges, forming a trapezoidal shape in any plane of polish. Previous investigations with WC-Co alloys at room temperature have shown that fracture of the WC particles occurs in transverse rupture testing.' Room-temperature slip was reported for WC particles after indentation for hardness measurements.3 Elevated-temperature deformation of WC particles in a WC-12 pct Co alloy was suggested by recent electron microscope studies of specimens deformed at 900' to 1000°C.4 In highly deformed alloys, the WC edges were serrated in contrast to the usual straight or smooth appearance. WC-12 pct Co and WC-15 pct Co alloys have been previously studied under elevated-temperature com-pressive-creep conditions by the present authors. Electron microscope studies of two-stage replicas from deformed specimens showed no evidence of slip or fracture of the WC particles. These specimens were brought to temperature and allowed to equilibrate prior to the application of the creep load. It was believed that the load-application rate, a crosshead speed of 0.005 in. per min on an Instron universal testing machine, was sufficiently low that recovery within the cobalt matrix was sufficient to limit the deformation to this matrix. A series of experiments was performed to evaluate the influence of loading rate on the deformation of WC-Co alloys. A WC-12 pct Co alloy was selected for these determinations. The average WC particle size was 4.45 p with an average linear separation between particles of 0.59 p. The selected temperature was 800°C and was monitored with a Chromel-Alumel thermocouple attached to the specimen. Testing was conducted in an argon-atmosphere chamber to prevent oxidation of the WC-Co specimens. This chamber was mounted on an Instron universal testing machine equipped to apply the load at a fixed rate. Each specimen was loaded to 110,000 psi compression stress at 0.05 and 0.5 in. per min. The loading rate was monitored prior to insertion of the test chamber and was found to be almost precisely the nominal rate selected. The specimens were raised to temperature and held to equilibrate with the surroundings, and then the load was applied and held for 4 hr to duplicate the exposure time utilized for the creep specimens. The time to reach full load at a crosshead speed of 0.005 in. per min was some 500 sec and was reduced to 50 and 5 sec as the loading rate was increased to 0.05 and 0.5 in, per min, respectively. The model developed by Ansell,' when recovery processes do not occur, considers that fracture or deformation of the dispersed particles is necessary to relieve back stresses on dislocation sources and allow dislocations piled up against particles to sweep out in the matrix to cause plastic deformation; he further states that, even at elevated temperatures, the dispersed-particle deformation is necessary for yielding in the absence of recovery. For the case of straight dislocation segments piled up against a straight barrier, such as the straight-sided WC particles, the shear stress, 7 exerted on a particle is: where h is the spacing between particles (0.59 p), a is the applied stress (110,000 psi), p, is the shear modulus of the matrix (6.7 X lo6 psi at 80O°C), and b is the Burgers vector of the matrix dislocation. From Eq. [I], the shear stress, 7, exerted on the WC particles when no recovery occurs is of the order of 6 X lo6 psi at 800°C. The limiting stress, F, that will

Jan 1, 1969

By O. R. Morris, G. L. Hawkes

A galvanic cell, using as electrolyte a fused salt solution of calcium carbide and as electrodes carbon and a Fe-C alloy of known composition, has been set up to study the thermodynamics of Fe-C alloys in the temperature rmzge 800" to 1000°C. Time independence and reproducibility of the cell electromotive force were taken as evidence of the reversible behavior of the cell. Carbon was believed to be present in the electrolyte as the so-called acetylide ion, C;-. The plots of the cell electromotive force us temperature for a specific alloy composition were straight lines within the limits of experimental error. The average Partial molar enthalpy of carbon in iron relative to pure carbon was found to be +10,610 i 93 cal per g-atom C. Thermodynamic analysis of the data has led to the following equation for the carbon activity, ac, based upon pure carbon as the standard state: In ac = In Zc + 10,560/RT + (10.02 + 77O/T)ZC - 2.350 where ZC is the lattice ratio [nC/(nF, - nc )] and T is the absolute temperature. This equalion gives carbon activity values generally slightly lower than those from gas equilibration studies reported in the literature. METAL LOGRAPHIC examination of a polished cross section of the steel anode used in the electrolysis studies of fused salt solutions of calcium carbide by Morris and Harry revealed extensive carburization of the steel by the electrodeposited carbon. This carburization was reflected in the variability, with time, of the applied potential to the electrolysis cell, necessary to maintain a constant current density at the electrodes. This observation suggested the setting up of a galvanic cell of the "alloy concentration" type to study the thermodynamics of some metal-carbon alloys. Cells of this general type have been widely used for the study of alloy systems.2 In view of the availability of published data in respect of the austenite phase of the Fe-C system, it was decided to carry out measurements upon these alloys before proceeding to studies of less well documented systems. The galvanic cell may be written: where [C] is carbon dissolved in iron. The electrolyte was a fused salt solution of calcium carbide, containing some 5 to 10 mol pct of carbide. The cell reaction is believed to be: C(s)-[CI [I1 Carbon forms an interstitial solid solution in iron, with the atoms located in the octahedral interstices. In the fcc crystal structure of austenite there is one octahedral interstice per iron atom. Thus, the lattice ratio, ZC, shown by Gurney3 to be the fundamental concentration parameter in the context of interstitial solutions, is given by: where nc and nFe are the number of carbon and iron atoms, respectively. chipman4 has recently shown empirically the advantages of using this concentration parameter instead of the more usual atom ratio or atom fraction. The cell electromotive force, E, assuming reversible behavior, is related to the carbon potential or the partial molar free energy of carbon in the solid solution relative to pure carbon at the same temperature and pressure, GP at the composition ZC, by the equation: where z is the carbide ion valency and F is the Faraday constant. An activity of carbon, ac, in the solution relative to the value of unity assigned to pure carbon, and an activity coefficient, qC , are defined such that: where R is the gas constant and T the absolute temperature. GF is further related to the relative partial molar enthalpy Hm, and the temperature coefficient of the cell electromotive force, (aE/aT)Zc, by the equations: Measurement of the cell electromotive force thus enables calculation of the relative partial molar thermodynamic properties of carbon in iron, if z is known. At E = 0, the solid solution is in equilibrium with pure carbon. More convenient for many purposes is the standard state based upon the infinitely dilute solution, Henry's law. The relationship between the activity coefficient of carbon based upon this standard state, and that based upon the pure carbon standard state, qC , may be obtained by considering the free energy of transfer of carbon from the latter standard state to the former. The relationship is: where +:H is the activity coefficient of carbon in the hypothetical standard state based on a reference of

Jan 1, 1969

By Lucien F. Trueb

The microstructure of explosion-bonded pairs of similar and dissimilar metals has been investigated by electron microscopy. A review of the specific problems encountered and the methods used for obtaining surface replicas and thin-film transmission specimens of the bond interface is given. The bond area is mainly characterized by continuous and practically diffusionless metallurgical bonding. The very large shear stresses induced along the collision front of the plates being joined causes extreme grain elongation and a symmetrical pattern of subgrains in the bonding direction. The bond zone is also characterized by a very high density of dislocations and pressure-induced twins. Localized heating occurring during the cladding process can result in partial re-crystallization or the formation of thin layers of molten material. The force of precisely controlled explosions causing a high-velocity impact between metal plates has been used for several years to achieve metallurgical bonding between an extremely wide variety of metals. This method essentially consists of accelerating a plate to high velocity toward a stationary plate by a detonating explosive. Since the restrictions to bonding are not those encountered with conventional nethods, it becomes possible to bond pairs of metals having widely different mechanical properties that are immiscible or form brittle intermetallic compounds. Many applications of such composite metals are found in the field of corrosion protection as well as numerous other fields; for example, explosion bonding is being applied for fabricating the materials used by the United States Mint in the new sandwich-type coins. The primary condition for establishing a metallurgical bond is that absolutely clean metal surfaces be brought together. Any metal exposed to the atmosphere is covered with oxides, adsorbed gases, and other contaminants; even a very forceful impact of two such surfaces is not sufficient for bonding. Cowan and Holtz-man,"' who reviewed the dynamics of colliding plates in detail, showed that in order to achieve a good bond the explosion conditions must be chosen in such a way that the plate collision velocity is less than the sonic velocity, in which case no oblique shock waves are attached to the collision front. A pressure wave is then generated ahead of the collision line, and the material forming the colliding surface of each of the plates flows forward and is ejected in the form of a spray, the so-called jet. The dynamic elastic limit of the metals must be exceeded so that there is sufficient plastic deformation. At the point where the jet formed by the junction of the inner surface layers of both plates separates from the combined plates, the material experi- ences a very high shearing strain and the pressure can reach several hundred kilobars. This process strongly influences the microstructure of the bond zone as will be seen later. Behind the collision front, uncontami-nated layers of internal material are brought together under high pressure and are thus metallurgically bonded. I) STRUCTURE OF EXPLOSION BONDS The different types of explosion bonds that can be obtained depend on the explosion conditions, and have been investigated by Cowan and Holtzman,1'2 Holtzman,3 Klein; Bahrani and crossland,' and Buck and Horn-bogen. The preferred kind for practical applications is the so-called wavy bond, typical examples of which are given in Fig. 1 showing light micrographs of various metal-to-metal interfaces. In forming this type-- of bond the collision energy is mainly expended in jetting, the formation of waves, and localized melting. Beyond the crest of the waves, eddy-shaped areas are observed in which the two metals are mixed in a complex pattern of streaks. Cowan and Holtzmanl first proposed that this wavy pattern is analogous to periodic eddy shedding in the flow of a viscous fluid around an obstacle (Von Karman's eddy street). The mass of metal ahead of the stagnation point, which is associated with the jet and has forward momentum, plays the role of an obstacle and the eddies created in the flow of solid metal around the stagnation point are preserved in the final clad specimen. This idea has been reviewed more recently by Klein4 and the variables involved in the wave formation have been discussed in some detail by Bahrani and crosslands and Buck and Hornbogen.6 Several studies of the structure of explosion bonds by light metallography have already been published.1-6 Aside from the waviness and the eddies which were mentioned above, the most striking characteristic of the area in the vicinity of the bond interface is a very considerable longitudinal grain deformation which appears to be strongest at the metal-to-metal boundary and dies out as one moves away from it. Large twins are often observed within the deformed grains, and molten areas are found in the center of the eddy-shaped structures situated beyond the crest of the waves. The large hydrostatic pressures and shear stresses occurring at the interface modify the mechanical and chemical properties of the bond zone. Increases in hardness in this area have been reported by various authors.396 The defects along the interface can also cause a local increase of the chemical reactivity and thus might be expected to boost the etching rate. However, the effects of this preferential etching cannot be observed by light microscopy due to its inherently limited resolution power. The same limitation precludes the observation of morphological features directly along the bond interface as well as the interface itself. Furthermore, no information can be gained by light-

Jan 1, 1969

By Robert E. Green, Edmund G. Henneke

The velocities of the three possible modes of elastic wave propagation have been calculated for single-crystal copper at 1-deg intervals throughout the standard stereographic triangle. The results are presented as isospeed contours. The deviations between the direction of maximum energy flux and specimen axes are presented as arrows drawn on stereographic projections. The orientation factors associated with dislocation motion on the primary slip system only and also on all twelve slip systems have been calculated for all three elastic waves. The orientation factors associated with a longitudinal standing wave resonant specimen have also been calculated. The results of these calculatians are compared with earlier calculations on aluminum crystals and the relevance of these results to ultrasonic attenuation measurements is discussed. GREEN and inton' have recently outlined a procedure for determining the orientation factors for dislocation damping in single crystals. The stress field associated with a given mode of elastic wave propagation is resolved on the slip systems in which it is assumed the dislocations are vibrating due to this stress field. Since, at room temperature, slip has been observed on a limited number of systems in fcc and hcp metal crystals, it is assumed that these systems are the appropriate ones to use for the resolving process. This procedure has been applied to aluminum crystals by Hinton and hreen' and to zinc crystals by Henneke and Green.3 In the present paper the modes of elastic wave propagation have been calculated for copper crystals possessing orientations at 1-deg intervals throughout the standard stereographic triangle. The deviation of the energy flux vectors from the specimen axis have been calculated for each mode. The stress field associated with each mode has been resolved on the (111)(110) slip systems in order to determine the orientation factors for dislocation damping. MODES OF ELASTIC WAVE PROPAGATION For any direction in a linear homogeneous aniso-tropic material three elastic modes can be propagated. In general, each of these modes has both longitudinal and transverse particle displacements associated with it and thus are not "pure" modes. However, there is always one component which is greatest in either a longitudinal or transverse direction and hence the mode may be referred to as quasi-longitudinal or quasi-transverse, accordingly. For certain axes of high symmetry pure modes may be propagated. The phenomenon of refraction of the energy flux vector is also exhibited by elastic wave propagation in ani-

Jan 1, 1969

By A. C. D. Chaklader, Ashok K. Kakar

The basic density equation originally dericed ' to predict the increase in density of a compact of spherical particles with the progressive deformation at the points of contact has been further modified to include the yield strength of the material. This has been done by assuming that the contact areas grow to stable sizes under a fixed stress which is equal to three times the yield strength. The final equation has the form: where Do and D me the initial and final bulk densities of the compact, u is the applied pressure, and Y is the yield strength of the material. This equation was tested with the data obtained on spheres of lead, K-Monel, and sapphire. The calculated yield strength t~alues for lead and sapphire are within the range of values reported in the literature. A few of the earliest hot pressing models proposed to explain the mechanism by Murray, Livey, and williams2 and then by McClelland3 are based on a plastic flow mechanism. However, more recent investigations suggest that the overall densification process is a combination of several mechanisms, such as particle rearrangement, fragmentation, plastic flow, and stress-enhanced diffusional creep. While fragmentation and particle rearrangement are considered to be responsible for the densification in the early stages,"475 it has been concluded that the final stages of hot pressing are controlled by stress-enhanced diffusional creep.516 The manner in which the densification takes place, i.e., by fragmentation, particle rearrangement, plastic flow, or stress-enhanced diffusional creep, would depend upon the type of material, the temperature, and the stress level used during the hot-pressing experiments. Metal compacts can be expected to have a much greater contribution from plastic flow than ceramic oxides. Also, plastic flow would be a significant contributing factor to densification at high temperatures and high stresses. Most of these works, directed towards elucidation of densification mechanism, have dealt with kinetics of the process. The results of most of the authors vary from one another and they have proposed either new empirical or semiempirical equations to fit their data. The densification rate was found to vary with the type of the powder, shape and size of the powder, initial packing density of the compact, and a few other factors such as rate of heating, pressure, and so forth. Beyond the initial stages, the densification process has been considered to be as time-dependent flow, controlled by a diffusional process, e.g., Nabarro-Herring creep. Palm our, Bradley, and johnson' have attempted to use modified creep rate equations to interpret the data of densification under hot-pressing conditions. Beyond the initial stages, however, the densification would be controlled by a process depending upon the temperature, pressure, and size of the powders. It is the authors' belief that such densification cannot be exclusively controlled by a single process and so attempts should be made to study some observable phenomenon like microstructure, yield strength, and so forth. The emphasis of this work has been toward studying the densification problem from a more fundamental point of view. Some of the principal variables, like initial packing density, mode of packing, and size of the powders, have been controlled to a great extent. The total strain produced on pressure application (instantaneous) in such a case can be considered to be due to plastic and elastic deformation. The elastic component of the strain can be determined by decreasing the load to the initial value. The strain remaining then can be correlated with the contact areas produced by deformation and the corresponding applied load. In a previous paper,' the possible deformation behavior of spheres in a compact has been theoretically analyzed and experimentally tested. The change in contact area radius a relative to the particle radius R was related to the bulk density and the bulk strain for simple and systematic modes of packing. Tt was found that a density equation relating the above parameters can be represented by: where D and Do are the bulk densities of the compact at any value of a/R and a/R = 0, respectively. This basic equation should hold for any material as it was derived from geometrical considerations alone. An attempt has been made in this work to include the yield strength in the above density equation, so that a knowledge of the properties of any material can be used in predicting the densification behavior during the hot-pressing process. THEORETICAL CONSIDERATTONS The deformation of two spheres in contact under a static load can be compared to the deformation occurring between a hard spherical indentor and the flat face of a softer metal. Tt has been shown theoretically by both ~encky~ and lshlinskyg and experimentally by ~abor" that, for a material incapable of appreciable work hardening, the mean pressure required to produce plastic yielding (for deformation occurring between flat face and a hemispherical indentor) is approximately equal to three times the elastic limit, Y, of the material (in tension or compression experiments). Tabor has further observed that the same relationship is valid in the case of work-hardening materials, if the elastic limit at the edge of the indenta-

Jan 1, 1969

By Jaakko Kajamaa

Neutron diffraction was applied to determine sheet textures by the transmission method. Cold-rolled and recrystallized copper sheets were investigated. The amount of cube texture was determined for three compositions, in which the phosphorus content was, respectively, 0, 0.005, and 0.03 wt pct. The heat treatment was in every case 8 sec at 650°C. In the two latter cases the cube texture was prevented. In addition a comparison with the X-ray diffraction transmission method was made with the 96 pct cold-rolled copper sheet. Outer parts of both (111) pole figures can be considered to be rather identical. This is seen from the fact that the intensity ratio ITD/120" was 0.45 for neutron diffraction and 0.40 for X-ray diffraction. Differences between the methods were discussed in detail. Features peculiar to neutron and X-ray diffraction in texture studies were listed and compared. In this work neutron diffraction was applied to determine sheet textures. Specifically, it was desired to ascertain whether this method can be used to reveal differences when compared to other methods. In addition, the amount of the cube texture in copper sheets was determined as a function of phosphorus content. Previous applications of neutron diffraction to texture problems include the following: nickel wires,' wire of some bcc metals,' and uranium bars.3 In the neutron diffraction technique the greatest difference is in the sample—its method of production and its volume. A sample needs no treatment and its volume is roughly 105 times larger than the volume of an X-ray diffraction sample. The cold-rolled sheet was investigated both by neutron diffraction and by X-ray diffraction, because it is expected that, due to large number of defects, possible differences in the results of the two methods would be revealed. It is a well-known fact that X-ray lines show broadening when cold-worked. Analysis has shown that this is based chiefly on small crystalline size, micro-stresses, and/or faults.4'5 Neutrons are sensitive to the above-mentioned disturbing factors as well, but circumstances in diffraction are different from the X-ray case. Because the sample represents a larger volume, the result is an average over that volume. In addition, it can be assumed that the sample has preserved its original structure, because it needs no special preparation. The particular limitation of neutrons is the relatively low neutron intensity available from nuclear reactors. This decreases the resolution as compared to the X-ray diffraction methods. Furthermore, absorption mainly reduces diffracted X-ray intensity, while multiple scattering effects, i.e., secondary extinction, disturb neutron diffraction. SO neutrons and X-rays behave in a different way when interacting with matter. As in other structural investigations, one can utilize this difference in texture studies as well. One cold-rolled and three recrystallization textures in copper sheets were investigated by neutron diffraction. The samples were produced at the Outokumpu copper factory to the specifications shown in Table I. The paper is divided into five parts. The first deals with the theory of the measurement. In the second, experimental procedures are described. Results are presented in the third part. Both cold-rolled and re-crystallized samples are studied. Discussion is in the fourth part, and finally in the fifth part some conclusions are drawn. 1) THEORETICAL CONSIDERATIONS Properties peculiar to neutron diffraction are the following: a) the scattering length varies greatly between one element and another; b) many of the elements do not absorb neutrons appreciably. In this connection it is of primary interest to know the interaction of neutrons with lattice imperfections. As with X-rays this problem leads to diffraction analysis of deformed and recrystallized metals. From the physical point of view the main difference is that neutrons are scattered by nuclei (magnetic scattering is not considered here), whereas X-rays are scattered by electrons. The features peculiar to neutron and X-ray diffraction methods in texture studies are listed in Table 11. Pole figures are an important tool in performing structural analysis of deformed or recrystallized metal. Present texture research technology requires pole figures which are as precise as possible. The choice between these two methods depends on the technical information which is required. The X-ray diffraction transmission technique may give results which are not necessarily representative of the average physical state of the sample. Although foil samples normally contain enough crystallites for diffraction, they may not necessarily represent the whole structure. An example of this problem is the frequently observed difference between the "surface" and the "inside" texture of a sample. The production of foil samples may disturb the original structure of the parent material. The selection and orientation of the foil from the sample is quite arbitrary. Normally, a highly deformed piece of metal has several texture components. Different components are deformed in a slightly different manner. This is a re-

Jan 1, 1969

By Craig S. Hartley

Expressions are obtained for the energy changes associated with the reaction of (a& (111) slip dislocations on intersecting (110)planes in anisotropic bcc metals. An energy criterion for assessing the likelihood of dissociation of the products of such reactions is also presented. It is found that the "burrier reactions" which form a(100) dislocations at the intersection of two active {110) slip planes are more energetically favorable in metals which exhibit a high value of Zener's anisotropy factor, A, than those which have a low value. The results are presented in a form which permits the stacking fault energy to be obtained from a measurement of the separation between par-tials in a dissociated configuration. However, until accurate calculations or measurements of the stacking fault energies involved are available, it is not possible to assess the physical importance of dissociated dislocations. In a recent paper,' the energy changes associated with several types of reactions between two slip dislocations, (a/2)(111){110), in bcc structures were calculated.* Isotropic elasticity and the approxima- tion v = -3- were employed. The purpose of this work is to present calculations of the energy changes for many of the same reactions using anisotropic elasticity. The problem of dissociation of a(100) and a(110) dislocations is also considered, and maximum fault energies for which dissociation will be energetically favorable are calculated for several bcc metals. Two general types of reactions are considered; those for which the reactant (a/2)(111) dislocations have long-range attractive forces and those for which the reverse is true. An example of the former is: (a/2)[lll] + (a/2)[lll]-a[l00] while the latter are typified by: (a/2)[lll] + (a/2)[111] -a[011] Only reactants lying in different slip planes are considered; therefore, the products must lie along (111) or (100) directions, which are the intersection of two {llO} planes. It will be assumed that the reactants and products are infinitely long parallel dislocations, since in this case the energy change associated with the reactions is a maximum.' THEORY The self-energy per unit length of a straight mixed dislocation in an anisotropic medium can be written? where b is the Burgers vector, K is an appropriate combination of the single-crystal elastic constants, and R and ro are, respectively, outer and inner cut-off radii of the elastic solution. The energy given by Eq. [I] does not account for any variation of the core energy with orientation. This could be manifested by an orientation dependence of the core radius or, equivalently, the Peierls width, of the dislocation. However, the energy contribution due to this source is expected to be small, and current models of the dislocation core are not sufficiently accurate to justify such a refinement. It has already been shown that for the isotropic case the energy contributions due to nonzero tractions across the cores of the reactants and products exactly cancel one another in the reaction.' Accordingly, it will be assumed that this contribution to the total energy change in the anisotropic case is small. In the subsequent discussion it is also assumed that the core radii of the reactant and product dislocation are the same and that, where stacking faults are formed, the faulted region is bounded by the centers of the partials. Consequently only changes in elastic energy due to the reactions will be considered. When the dislocation is parallel to either the (111) or the (100) directions, K may be written:375 K = (Ke sin2 a + Ks cos2 a) [2] where K, and Ks are the combination of elastic constants corresponding to an edge and screw dislocation lying along the same direction as the mixed dislocation, and a is the angle between the direction tangent to the dislocation line and the Burgers vector. Eq. [2] should not be confused with the isotropic approximation to the variation in energy with line Orientation.6 It should be noted that the essentially isotropic expression for K is a result of the characteristic symmetry of the (111) and (100) directions and is not, in general, valid for other dislocation directions in anisotropic cubic metals. The energy* change for a reaction in which the re- actant and product dislocations are parallel perfect dislocations can be written: where Ep and E, refer to the self-energies of the products and reactants, respectively. For dislocations parallel to (100) and (111) directions, Eq. [3] becomes:

Jan 1, 1969

By O. P. Quist, S. H. Carpenter

The internal friction of single-crystal LiF has been investigated as a function of crystal orientation, while simultaneously applying a static compressive load. Three different crystal orientations were used, these being samples with the (100), (110), and (111) directions along the sample length. Results show the internal friction measured to be similar in appearance but different in magnitude for the (100) and (110) samples. With both orientations, application of a static stress caused an increase in the amplitude-independent damping, which recovered with time. The amplitude-dependent damping was found to first decrease with load up to the yield stress and then increase rapidly. Samples of the (111) orientation were found to have no measurable internal friction even when loaded up to 3.0 kg per sq mm. These results are discussed in terms of current theories of internal friction in ionic crystals. This paper presents the results of an investigation of the dislocation mobility and dislocation motion in high-purity single-crystal LiF at applied stress levels up to the yield stress. Internal friction measurements of the dislocation damping while simultaneously applying a static compressive load was the experimental technique used. LiF crystals much prefer to slip on the (110) set of slip planes. However, with a suitable stress distribution they can be forced to slip on the {100} set of slip planes.' For both sets of slip planes the slip direction is the same, namely (110). The object of this investigation was to try and learn more about the dynamical properties of dislocations on the different slip systems in LiF crystals. EXPERIMENTAL PROCEDURE Dislocation damping was measured, in terms of the log decrement, using a five-component marx2 composite piezoelectric oscillator. The apparatus used allows one to measure the dislocation damping while simultaneously applying a static compressive load on the sample. A schematic of the apparatus used is shown in Fig. 1 and a complete discussion of its construction and operation is given elsewhere.3 All measurements were made at room temperature with a longitudinal resonant frequency of approximately 50 kc per sec. Damping measurements were made using longitudinal vibrations with strain amplitudes from up to 10. Static loads up to 3.0 kg per sq mm were applied to the sample while the damping was being measured. Samples used in this investigation were high-purity single crystals of LiF purchased from the Harshaw Chemical Co. All samples were approximately of an in. square and were hand-ground in length to one-half wavelength to match the resonant frequency of the driver and gage quartz crystals. Samples of three different crystallographic orientations were used, namely, crystals with the (loo), (110), and (111) direction parallel to their length. The magnitude of the damping is a direct measure of the area on the slip planes swept out by dislocations during a complete oscillatory cycle. With a sample of the (100) orientation the maximum applied shear stress, due to longitudinal vibrations, is on the (110) slip planes with zero shear stress on the (100) slip planes. Thus, damping measured from a (100) sample is due entirely to dislocation motion taking place on the (110) set of slip planes. Similarly, for a sample of the (111) orientation there is zero shear stress on the (110) set of slip planes and the damping measured from a sample of this orientation is due entirely to dislocation motion on the (100) set of slip planes. Samples of the (110) orientation will have a component of shear stress on both sets of slip planes and

Jan 1, 1969

By K. E. Manchester, A. H. Clark

Hall measurements have been made on three groups of silicon samples, which were implanted with boron, aluininunz, and phosphorus ions. Boron and phosphorus implants show essentially bulk properties when annealed under previously determined conditions, i.e., 15 min at 80O°C for boron and 10 mirz at 600°C for phosphorus. Bulk properties were also observed in alu?rrilzum-implanted samples when ion concentrations were below 1015 cm-'; however, strange low- terrzperalure behavior and a falloff in cavrier concenlralion was observed for sarnples with ion concenlration abozle loL5 ern-'. This can be attributed to solid-solubility effects since the average bulk concentration f0.r the sample exhibiting bulk properties was below the reported solubility limit for aluminum, while the other samples were above the limit. A detailed study of the annealirzg process indicates that mobilities reach bulk c~alues in phosphorus implants at 250DC, in alutninurn ivnplanls a1 500°C, and in boron implants at 700°C. A simple rnodel has been proposed to fit the annealing data. THE technique of semiconductor doping by direct injection of energetic ions is being actively investigated at this Center. Previously published results from this laboratory have been concerned with the investigation of implanted profiles in silicon and the electrical properties of the resultant p-n structures.1'2 Diode properties have been reported for junctions produced by implantation of boron ions and aluminum ions in n-type silicon and phosphorus ions in p-type silicon. Reverse current characteristics, as well as electron diffraction data, have indicated that the damage to the structure produced by the energetic ions during the stopping process can be minimized by a gentle anneal. The time-temperature conditions for this process have been determined by sheet resistivity studies.3 Electrical profiles of implanted phosphorus ions which have been determined by a differential sheet conductivity technique are based on the assumption that the carrier mobilities and activation energies for ionization are the same as in bulk silicon. These assumptions have been substantiated experimentally by measurement of Hall constants and resistivities of these implanted structures and the results are presented herein. After establishing the bulk property equivalence of the annealed layers, a more detailed examination of the annealing process was undertaken. The sheet resistivity technique initially used to determine anneal conditions has been broadened to include Hall constants and mobility behavior with anneal conditions. PROCEDURE The three groups of samples prepared for Hall measurements were implanted with mass-separated beams of boron, aluminum, and phosphorus ions. Ion concentrations are equivalent surface concentration and they have been determined from beam current measurements during implantation. Implant conditions for the samples are shown in Table I. In all cases, the samples had previously been angle-lapped on one edge and the junction delineated with an HF-copper sulfate solution. The surfaces of the implanted regions of the samples were exposed to the delineating solution and in the case of the "n"-type phosphorus-implanted areas it was later discovered that a finite amount of surface could be etched by this solution. This is discussed in the section on results. The implanted samples were ultrasonically cut into the conventional six-probe configuration shown in Fig. l.* In the annealing studies, for which the meas- urements were taken at room temperature, the sample holder was a Bakelite jig incorporating tungsten mechanical probes. The other measurements were taken using a conventional liquid-nitrogen cryostat, in which the temperature could be varied from 77" to about 400°K. For these measurements, 10-mil aluminum dots were evaporated onto the samples and alloyed with the silicon at 530°C for 15 min. The samples were then mounted in flat-packs, and gold wires were bonded to the aluminum dots. The electrical measurements were made with either a Leeds & Northrup K-5 Potentiometer, a Keithley 147 Nanovolt Null Detector, or a Keithley 601 Electrometer. No external bias was applied to the substrate. There is an inherent self-biasing arrangement at one or the other of the current contacts (depending upon the polarity), and this suffices to limit the current flow to the implanted layer. The magnetic field was generated by an A. D. Little electromagnet. Except for occasional checks of linearity of the Hall voltage with

Jan 1, 1969

By Kiyoshi Azuma, Hiroshi Kametani

The variation of the dissolution behavior of a ferric oxide with calcining temperature has been investigated. Samples were prepared by thermal decomposition of ferric hydroxide, nitrate, oxalate, and sulfate at low temperature, followed by the calcination in the temperature range between 600" and 1200°C. The samples of eight series and a fine crystalline sample of hematite were dissolved in 1 N hydrochloric acid at 55.2°C and the results are represented on double-log graphs for convenience. It is confirmed that all dissolution courses follouj either the accelerated process or the parabolic process except in the special case of the crystalline hematite which dissolced in accordance with the uniform dissolution of a particle. Examinations of the physical properties of the oxide powders revealed that the surface area measured by the permeability method is strikingly relevant to the dissolution behavior of the oxide. In the previous paper,' detailed data were presented on the effect of the kind of acid, the solution temperature, and the concentration of acid on the dissolution of two ferric oxides. It was also shown that these sam ples dissolved in strikingly different ways. The present investigation was carried out on the dissolution of various calcined samples prepared from various ferri salts by various methods to ascertain the course of dissolution. Pryor and Evans2 pointed out a change of the dissolution rate at around 700°C for a series of calcined ferric oxides prepared from the hydroxide. Several papers374 reported also the dissolution of ferric oxide samples. It seems, however, that a systematic account of the relationship between the dissolution behavior and physical properties of the oxide has not yet been given. This paper presents the variation of the dissolution of the oxide in relation to the calcining temperature and the change of physical properties of the calcines. EXPERIMENTAL Raw materials were prepared by precalcination of ferric hydroxide, thermal decomposition of ferric nitrate, oxalate, and sulfate, and aerial oxidation of ferric chloride vapor, at as low a temperature as possible. The products were crushed, ground, if necessary, and sieved with a 100-mesh Tylor screen prior to calcination, after which the specimens were dissolved in acid solution. The following is a detailed description of the preparation of the samples. Sample H. About 500 g of ferric chloride (guaranteed reagent) were dissolved in 5 liters of deionized water and filtered. Ferric hydroxide was precipitated by addition of the minimum amount of ammonium hydroxide solution, and the precipitate was washed continuously till chloride ion was not detected by silver nitrate solution, and then filtered. The filter cake was dried at 120°C for a week and ground, and the -100 mesh portion was used. Sample S. Ferric sulfate (guaranteed reagent) was pyrolytically decomposed in a crucible at 700°C for 24 hr and the product was sieved. In this case the following calcination was carried out at temperatures over 700°C. Sample B. Commercial ferric oxide (guaranteed reagent). About 15 kg of ferric nitrate were decomposed in a furnace maintained at 800°C for 2 hr. The actual temperature of the decomposition was not measured. The product was crushed and sieved, and the -100 mesh portion was used. Sample N. About 50 g of ferric nitrate (guaranteed reagent) were decomposed in a beaker in a sand bath until a red-brown dense solid was produced. This product was crushed and sieved, and subjected to complete decomposition at 500°C. The precalcined product was again sieved and used. Sample N2.5. Since the decomposition temperature was not controlled for sample AT, a different sample was prepared in a temperature-controlled furnace. The subscript represents the decomposition at 250°C. The product was treated in the same manner as sample N. Sample Nc. Under atmospheric pressure it is prac-tically inevitable that ferric nitrate hydrate melts to form a brown liquid at about 50°C before pyrolysis. For this reason, the salt was first slowly heated under reduced pressure (about 10-3 mm Hg measured in a trap refrigerated by dry ice-alcohol) to achieve dehydration without melting. About 5 hr were required for the dehydration and the partial decomposition. Then the temperature was elevated to 500° C in air for complete decomposition. The relatively porous product was sieved and used. Sample Ov. About 200 g of ferric oxalate hydrate (extra pure) were dehydrated under reduced pressure (as described above) followed by thermal decomposition at 500°C for 6 hr in air. The decomposition of this salt was accompanied by liberation of carbon monoxide, by which the ferric salt was initially reduced to a black powder. The powder changed in turn into brown ferric oxide as the gas liberation decreased and reoxidation predominated. The product consisted of sparkling fine particles passing through a 100-mesh screen. However it was ground and sieved as for the other samples. Sample D. Commercial fine powder for magnetic tape purposes. The preparation was as follows.5 Ferric chloride vapor and preheated excess air were mixed and passed into a reaction tube where oxidation took place at 450°C. The fine powder formed was collected in a cottrell chamber. The product was vacuum-degassed at 450°C for 1 hr and sieved.

Jan 1, 1969

By R. E. Reed-Hill, W. H. Hartt

High-purity magnesium single crystals, oriented with basal plane parallel to stress axis, were deformed in tension at room temperature so as to form second-order (1011)- (1012) twins. Investigation by optical and replica electron microscopy revealed that these twins can support very large internal strains approaching 1000 pct and that fracture follows as a direct consequence of these strains. At low strains the basal plane within (10i1)- (10i2) twins rotates with increasing deformation until it becomes approximately parallel to the twin boundaries. This rotation is not consistent with deformation by basal slip in the double twin, but may be rationalized in terms of nonbasal slip. Observations also indicate that basal slip in the primary (1011) twin, prior to second-order twinning, may contribute to this rotation. Deformation markings that are difficult to interpret in terms of slip on previously observed systems were noted in the twin bands. A possible means of rationalizing these markings is to relate them to grain boundary shear. The fracture associated with these twins is thought to initiate by formation of voids at the twin boundaries. These grow into microcracks, with ultimate fracture occurring by tearing of the interconnecting regions. THE significance of deformation and fracture along second-order (10i1)-(10i2) twins in polycrystalline magnesium1-' and many magnesium alloys4 was discussed in a previous paper.5 It was noted that this double twinning mode effectively controls the room-temperature deformation and fracture of magnesium single crystals, oriented with the basal plane parallel to the tensile axis. Reed-Hill and Robertson6 observed that, although these crystals may fracture at a macroscopic strain of less than 1 pct, shear strains up to 1000 pct could occur within (1011)- (10i2) twin bands. They concluded that the fracture was "ductile" and resulted from concentrating the deformation into a very small volume. The present research was undertaken to obtain a better understanding of the deformation and fracture associated with these twins. EXPERIMENTAL TECHNIQUES The experimental procedure involved straining high-purity magnesium single crystals in tension at room temperature. Specimens were oriented so that the applied stress direction was within 2 deg of the [1010] crystal axis. The specimen cross section was rectangular with faces closely parallel to (1210) and (0002). The specimen preparation and testing procedures have been described in detail elsewhere.5 Second-order (1011)- (1072) twins nucleated during straining were studied using optical and replica electron microscopy techniques. For the latter, cellulose acetate-carbon double replicas were employed. Observations were made on the (1210) crystal surface, which is the plane of shear for twins of this type. In addition to the single crystals, some investigations were carried out on longitudinal polycrystalline specimens. These were obtained from high-purity magnesium plate with a texture in which the basal planes of the grains tended to be nearly parallel to the specimen axis. All. polycrystalline specimens were annealed prior to testing in order to produce a coarse grain structure, permitting Laue back-reflection X-ray photographs to be obtained from grains of interest. EXPERIMENTAL RESULTS a) Deformation Within ( 101 l )-( 1072) Twins. Fig. 1 is an optical micrograph demonstrating the extent of the plastic deformation that can occur in (1011)-(10i2) twin bands. Notice at the lower left corner of the photograph the large displacement of the dark band (a polishing step) where it crosses the twins. The shear strain in the twin band at this position was computed to be 700 pct, using the ratio of step width to twin thickness. It was previously shown5 that (1011)- (1012) twin bands are often composed of small, separate twins aligned close to the macroscopic habit plane. Fig. 2 is an electron micrograph from near the upper end of the twin band in Fig. 1. This illustrates that the twin habit often does not coincide exactly with the macroscopic band habit, but may be inclined at a slightly smaller angle to the matrix basal plane. Figs. 3 through 6 are electron micrographs of the twin band in Fig. 2 illustrating successive steps in the development of deformation inside the band. A twin band such

Jan 1, 1969