Second Israel Materials Engineering Conference

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INIS-mf—10066

SECOND ISRAEL MATERIALS ENGINEERING CONFERENCE

February 21-23,1984

BEN-GURION UNIVERSITY OF THE NEGEV

BEER-SHEVA, ISRAEL

CONFERENCE PROCEEDINGS

Edited by

A.GRILL & S.I.ROKKLIN SECOND ISRAEL MATERIALS ENGINEERING CONFERENCE

February 21-23.1984

BEN-GURION UNIVERSITY OF THE NEGEV

BEER SHEVA, ISRAEL

CONFERENCE PROCEEDINGS

Edited by

A.GRILL & S.I.ROKH LIN ORGANIZING COMMITTEE Chairman - Prof. S. Rokhlin, Ben-Gurion University of the Negev Dr. U. Arnon, Israel Aircraft Industry Prof. D. Brandon, Technion -Israel Institute of Technology and Israel Institute of Metals Dr. B. Cina, Israel Metallurgical Society and Israel Aircraft Industry Prof. M. Dariel, Ben-Gurion University of the Negev and Nuclear Research Centre - Negev Dr. S. Kenig, Rafael - Haifa Dr. G. Metzger, National Council for Research Development Dr. H. Paruz, Ministry of Defense Prof. J. Pel leg, Ben-Gurion University of the Negev Dr. M. Polak, Ben-Gurion University of the Negev Prof. M. Ron, Technion - Israel Institute of Technology Dr. A. Stern, Nuclear Research Centre - Negev Prof. B.Z. Weiss, Technion - Israel Institute of Techology

EDITORIAL COMMITTEE

Chairman - Prof. A. Grill, Ben-Gurion University of the Negev Prof. L. Kornblit, Ben-Gurion University of the Negev Prof. I. Minkoff, Technion - Israel Institute of Technology Mr. B. Rnbin, Israel Aircraft Industry Prof. M. Rosen, Ben-Gurion University of the Negev Prof. M. Schieber, The Hebrew University of Jerusalem Dr. J. Zahavi, Israel Institute of Metals Mr. Z. Wagner, Israel Military Industries

THE CONFERENCE WAS SPONSORED BY:

jpartment of Materials Engineering, Ben-Gurion University of the Negev Department of Materials Engineering, Technion - Israel Institute of Technology Faculty of Engineering, Tel Aviv University School of Applied Science and Technology, The Hebrew University of Jerusalem Ministry of Defense National Council for Research and Development Israel Atomic Energy Commission, Nuclear Research Centre - Negev Israel Military Industries RAFAEL - Israel Armament Development Authority INTEL Electronics Ltd. ASHOT - ASHKELON Ltd. ISCAR BLADES Ltd. ISKOOR Ltd. URDAN, Associated Steel Foundries Ltd. CONTENTS

I. TRANSFORMATIONS

Kinetics and Metastable Structure Formation in Rapid Solidification Processing. (Invited). D. Turnbull 1 Some Techniques for the Study of Atomic Motions with Applications to Ceramic Materials. (Invited). A.S. Nowick 11 An Atomic Resolution Study of Radiation-Induced Precipitation in Fast Neutron Irradiated Tungsten (Rhenium) Alloys. (Invited). R. Herschitz, D.N. Seidman 21 Some Aspects of the Crystallization of Rare Earth-Noble Metal Amorphous Thin Film. L. Shikhmanter, M. Talianker, M.P. Dariel 30 Dendritic Growth and Dendrite Arm Spacing in the Solidification of Steel. M. Bamberger, I. Minkoff 35 Transition from Fibrous to Lamellar Morphology in Undierctional Solidified Ni-W Eutectic. S.F. Dirnfeld, Y. Zuta 39

The Microstructure of Rapidly Solidified Cu-Fe Alloys. A. Munitz, Z. Livne 46 Phase Stability 5 Massive Transformation in Y203 - Completely Stabilized Zirconia (Y-CSZ). A.H. Heuer, R. Chaim, M. Ruhle 51

Microstructure Evolution S Ordering in Commercial MGO-Partially Stabilized Zirconia (MG-PSZj. R. Chaim, D.G. Brandon 55

Deformation Induced Decomposition of Uranium-Titanium Martensite. G. Kimrnel, J. Sariei, A. Landau, M. Talianker 59

II. ALLOY PHASES AND STRUCTURE

Applications of Analytical Electron Microscopy to Materials. J.I. Goldstein, O.B. Williams, M.R. Notis 63 Field Electron and Ion Emission from Zirconiated and Zr Free W Cathode. J. Pelleg, J.L. Fink 67 The Transmission Electron Microscopy of NBi-5 Brazed Joint of Inconel 718. B. Grushko, 0. Botstein, B.Z. Weiss 71 II

Size Effect in Radiation Induced Segregation (RIS). L. Kornblit, A. Ignatiev 74 Applications of Advanced Color Computer Graphic Display Techniques to Materials Problems .Phase Equilibria § Diffusional Growth. M.R. Notis, S.K. Tarby, J.I. Goldstein 77

Coercivity and Squareness Ratio in Co-W Thin Films. U. Admon, G. Kimmel, M.P. Dariel, E. Grunbaum, J.C. Lodder 82 Structural Analysis of High-Vacuum, High-Temperature BNi-5 Brazed Joint of Inconel 718 Superalloy. B. Grushko, B.Z. Weiss 88 Phase Relations in the Cu-Nd System on the Cu Rich Side. C. Laks, J. Pelleg, L. Zevin 92 Effect of Small Additions on Grain Refinement of a 14 Carat Au-Ag-Eu-Zn Alloy. M. Fishman, L. Gal-Or, A. Iram . 95

Reinvestigation of the Pr-Ga System in the 66-100 at% Range. J. Pelleg, D. Dayan, G. Kimmel 100 Texture in Low Alloyed Uranium Alloys. J. Sariel, G. Kimrael, J. Pelleg 104 III. MECHANICAL BEHAVIOR Microstructure and Properties of Tungsten-Based Heavy Alloys. (Invited). D.V. Edmonds 108

Investigation of the Creep Failure Mechanism in the Mo-5%W Alloy. A. Freund, D. Agronov, A. Rosen 119 Isochronous Creep of Copper-Beryllium-Nickel Alloy (Cu-0.4 Be - 2.0 Ni) Solution Treated and Aged . N. Nir 122

Hot Tearing of Lead Alloys. P. Arigur, F. Weinberg 128 The Embrittlement of Steels by Low Melting Point Metals. N. Breyer 133

The Effect of Substructure on Creep Properties of the TZM Alloy. D. Agronov, E. Freund , A. Rosen 137

Fracture Toughness Evaluation of Brittle Materials Using Indentation Method. Z. Nissenholz 142 Ill

IV. ENVIRONMENTAL BEHAVIOR

Failure of Welded Inconel-600 Pipe in the Cooling Systems of a Nuclear Reactor. G. Kohn, B. Herrman, E. Rabinovitz, A. Stern, S. Addess 146

Effects of Metallurgical Variables on Hydrogen Embrittlement in Types 316, 321 and 347 Stainless Steels. P. Rozenak, D. Eliezer 152

Martensitic Transformations in 304L and 316L Types Stainless Steels Cathodically Hydrogen Charged. E. Minkovitz, D. Eliezer 158

Mechanical Properties Degradation of Hydrogenated Austenitic SS. I. Gilad, Y. Katz, H. Mathias 163

Tensile Flow and Fracture Behaviour of Austenitic Stainless Steel After Thermal Aging in Hydrogen Atmosphere. Y. Rosenthal, M. Mark-Markowitch , A. Stern, D. Eliezer 167

Corrosion Behaviour of Al-Cu Alloy Thin Films in Microelectronics. J. Zahavi, M. Rotel, H.C.W. Huang. P.A. Totta. 175

V. PROCESSING AND TESTING

Quantitative Nondestructive Evaluation Using Ultrasonic Waves. (Invited) L. Adler 185

On the Homogenization Problem in Sintered Alloys. L. Levin, A. Stern 201

Heat Transfer to Water and its Importance for Metal Casting and Heat Treatment. M. Bamberger, B. Prinz 209

Fast, Non-Destructive Elechtrochemical Detection of Surface Inclusions in Metallic Substrates. I. Rubinstein 214

Telephone Tokens Produced by Powder Metallurgy. A. Sharon 216

Automatic Detection of Fluorescent Indications. K.M. Jacobsen 221

Cryoforming of 301 and 302 Stainless Steel. T. Livni, S. Bar-Ziv, A. Rotem, A. Rosen 225

Ausforming of H-ll Mod Steel G. Elkabir, A. Rosen 232 IV

VI. SURFACE PHENOMENA

Thermodynamic and Kinetic Phenomena in Adsorbed Layers. (Invited). M. Grunze 235

Supported Silver Catalysts Have Some Important Properties in Common with Rough Silver Films: EHipsometry § Raman Data. P.H. McBreen, D. Hall, J. Lalman, M. Moskovits 245

XPS Studies of Si Films Deposited from SiCl4 by an RF Cold Plasma Technique. E. Grossman, M. Polak, A. Grill 249

Quantitative XPS of NaS-Alumina: Evidence for Sodium Anisotropic Segregation. Y. Grinbaum, M. Polak 253

VII. SURFACE TREATMENT AND COATINGS

Structure of Protective Diffusion Coating for Niobium Alloys. M. Kazinets, 0. Gafri, L. Zevin, B. Rabin 258

The Effect of B and P Doping on the CI Concentration and the Deposition Rate of Si from SiCl4 in RF Plasma. R. Manory, E. Grossman, R. Avni, A. Grill 262

Silicon Nitride Coatings by the Low Pressure RF Plasma Technique. U. Carmi, A. Raveh, A. Inspektor, R. Avni 267

Boridation of Steels in Low Pressure R.F. Plasma. A. Raveh, A. Inspektor, U. Carmi, E. Rabinovitz, R. Avni 272

Laser Induced Copper Electroless Plating. S. Tamir, J. Zahavi 277

Surface Hardening of Steel by Boriding in a Cold RF Plasma. I. Finberg, R. Avni, A. Grill, T. Spalvins, D. Buckley 283

VIII. MATERIALS AND PROCESSES FOR ELECTRONICS

Current Metallization Issues in Microelectronic Devices. (Invited). K.N. Tu 288

Schottky Barrier Height for Ti-W on Silicon. M.O. Aboelfotoh, K.N. Tu 291

Interfacial Reactions in Laser Annealed Ni/GaAs Contacts. A. Lahav, T. Brat, C. Cyterman, M. Eizenberg 299

Enhanced Opto-Electronic Activity of Semiconductors by Means of (Photo) Electrochemical Etching. R. Tenne, V. Marcu 304 Laser Induced Metal Deposition on GaAs substrates. J. Zahavi 308

Vapor Phase Soldering of Surface Mounted Electronic Assemblies. E. Falkenstein, I. Fainaro 315

Lattice, Grain Boundary and Short Circuits Diffusion of Phosphorus in TaSi Thin Films. J. Pelleg 320

The LIMM Technique for Determination of the Spatial Distribution of Polarization of Space Charge in Polymer Electrets. S.B. Lang 324

Chalcogenide Infrared Glass Fibers A. Bornstein, N. Croitoru 328

Programming of Crystal Diameter in Czochralski Growth by a Cooling Plot: Application to InSb. M. Azoulay, Z. Burstein 332

IX. MATERIALS FOR ENERGY CONVERSION

The Dependence of the High Temperature, High Flux Stability Materials on Surface Structure and Composition. (Invited) . A. Ignatiev 336

Transparent Conductor Films as a Material for Photovoltaic Junctions with Polycrystalline Silicon. Z. Harzion, M. Zafir, J. Rishpon, S. Gottesfeld, N. Croitoru 343

Photoelectrochemical Characterization of ^, G. Dagan, G. Hodes, S. Endo, D. Cahen 347

MN. 15Pe~ oc - Hydride Compacts for Hydrogen Heat Pumps. Y. Josephy, Y. Eisenberg, M. Ron 350

X. POLYMERS AND COMPOSITES

Crosslink Density of Polymers - Can it be Determined by Solvent Swelling? M. Gottlieb 356

Phase Separation in Rubber Modified Flame Retardant Epoxy Systems. Hemi Nae 361

Recent Advance in the Strength and Time Dependent Failure Process of Kevlar Monofilaments and Composites. H.D. Wagner, S.L. Phoneix, P. Schwartz 365

Deformation Processes in Impact Modified PVC. A. Hadas, A. Siegmann 369

Fatigue Crack Propagation Mechanisms in Polymers. A. Bussiba, Y. Katz, H. Mathias 373

AUTHORS' INDEX 378 KINETICS AND METASTABLE STRUCTURE FORMATION IN RAPID SOLIDIFICATION PROCESSING

D. Turnbull

Division of Applied Sciences, Harvard University, Cambridge, MA 02138 USA

ABSTRACT

The general procedures in metastable structure synthesis are surveyed and the reasons for kinetic preferences for metastable structures are discussed. In melt solidification the motivation for forming metastable states arises from the interfacial undercooling which results when high thermal gradients are imposed at the solidification front. The conditions for formation of glasses, supersaturated solutions and microcrystalline solids in rapid solidification processing are discussed.

INTRODUCTION

The general procedure in low pressure metastable structure synthesis is to energize a material--by melting, dissolution, irradiation or cold working--and then deenergize it by quenching or some condensation process [1,2]. To trap a metastable structure, the end stage of deenergization must be one in which the kinetic processes within the structure are too sluggish to alter its configurational state during the period of its intended study or use. Thus, the structure would be in a "configura- tionally frozen" rather than metastable equilibrium state. Nevertheless, it is common practice to apply to such states the label of the one from which they were frozen.

Three types of metastability might be distinguished:

(11 Compositional--e.g. supersaturated solutions. (2) Topological--e.g. glasses and other amorphous solids and certain crystalline phases. (3) Morphological--e.g. solids with a considerable density of such extended imperfections as: dislocations, interphase and intergrain boundaries.

Numerous examples of all three types of metastability have long been recognized. This paper surveys the formation of novel alloy structures in the above three categories by the melt quenching technique. First, however, a digression on the problem of kinetic preference in metastable structure synthesis is presented. KINETIC PREFERENCE IN METASTABLE STRUCTURE FORMATION

Deenergization may open up possibilities foT evolution of the system into some metastable, or oven unstable, state. Its outcome will be determined by: (a) what thermodynamic options develop, (b) which of these options is kinetically preferred, and (c'J whether the last stage of deenergization is rapid enough to kinetically trap a meta- or unstable state, if formed.

Actually, much experience indicates that metastable states often are kinetically preferred over the most stable ones. The early experience of this nature so impiessed Ostwald that he formulated his well-known "step rule", that, in deenergization a system evolves through the succession of available metastable states of decreasing free energy [3]. Such a rule has not been justified theoretically and it must break down when the driving free energy, AG, of forming the metastable structure is sufficiently small. If force coupling contributions are neglected the net flux of matter, Jj, in a particular transition is the product of a kinetic factor, p^, and a thermodynamic factor, f(AG^"), which reduces to AG:. in the linear kinetic regime, RT » |AGJ|, thus:

p. = J./f(AG.). (1) Then the interesting question is how do the p- for the various possible transitions correlate with the corresponding AG^ or other thermodynamic constants.

While there are, indeed, many examples where kinetic preferences, as indicated by the p^, in evolving systems are different from those most favored thermodynamically, there are clear exceptions to this behavior. A more general pattern appears to be kinetic preference for the state with entropy nearest to that of the initial one. For example, in the down- quenching of vapor or liquid, amorphous over crystal state formation seems always to be preferred kinetically, while in crystal upquenching liquid always forms faster than dilute vapor. Also, in phase separation, the most favored kinetic path is generally through a succession of states of decreasing dispersion. This pattern toward step entropy evolution in structural change may reflect that the processes most favored kinetically are likely to be those requiring the least changes in the correlations of the positions and motions of atoms in ';he evolving group. For example, in condensation the atomic positional correlations which must develop are less extended for amorphous than for crystalline structures. However, the favored sequence in the structural evolution of superheated quartz (quartz -*• liquid -*• Cristobalite) apparently departs from a step entropy pattern. This path probably is determined by some poorly understood features of the interface structures (quartz - Cristobalite and quartz - liquid). It seems evident, also, that the step sequence in structural evolution may be altered by heterogenous nucleation, e.g., it is likely that the first appearing phase in undercooled Fe alloy melts [4] should revert from b.c.c. to f.c.c. by addition of appropriately chosen heterogeneous nuclei. In these instances a step entropy preference, while not apparent in the bulk behavior, might still be exhibited by the interface processes.

From these considerations we expect that in rapid solidification processing formation of supersaturated solutions and microcrystall ine and amorphous structures would be kinetically favored. Whether or not these structures form will be determined by the magnitude of the thermodynamic driving free energy which is developed and their entrapment would be determined by the rate of quenching, T, through the configurational freezing temperature.

KINETICS OF SOLIDIFICATION

The optimum melt disposition for rapid quenching is as a thin layer in good thermal contact with a massive heat sink of high thermal conductivity. The sink is usually composed of a crystalline metal or alloy with a composition often very different from, but in some experiments identical with, that of the melt. It is the undercooling, AT^, developed in the melt at the interface during rapid quenching which may drive the formation of non-equilibrium structures. AT- = Tβ - Tj where T^ is the liquidus and Tj the interface temperature.

To sustain substantial undercooling a melt must be highly resistant to homogeneous crystal nucleation. Simple ["classical") nucleation theory relates the steady state nucleation frequency, I, to the scaled

undercooling, - T)/T£, as follov.s [5]:

~ nk. exp (2) 3(AT )2T

where n = number density of melt atoms or molecules kj = frequency of crystal-melt interfacial rearrangement

3 = ASm/R AS,,, = molar entropy of melting, here assumed temperature independent

Tr = T = absolute temperature of undercooled melt 1 1/7 assumed T independent and isotropic AH m a = crystal-me±t interfacial tension AH,,, = molar enthalpy of melting N = Avogadro's number V = average molar volume At a given value of the scaled crystal-melt interfacial tension, ex, and

with k^ in the range M/picosec., the calculated I (ATr) rises sharply with

ATr from immeasurably low to measurable levels at some onset undercooling AT^(T] which depends on the quench rate, f, but only weakly in the high atomic mobility ("labile") range. Reflecting this weak dependence, at the lowest realizable cooling rate AT'(T') approaches some minimum Threshold value AT^Q which increases sharply with increase in a. At a > 1, 1 does not reach a measurable level at any undercooling.

When the interfacial rearrangement is thermally activated, its fre-

quency decreases with falling T so that I(ATr) would be depressed to levels well below those calculated with k; - 1/pic.osec. This effect has been roughly evaluated for the case where k; scales as the shear relaxation frequency. It is found [5] that with increasing scaled glass temperature. Trg = Tg/T^ where T is the actual glass temperature, I(AT ) is depressed, its maximum sharpens, and AT^.Q is increased. Also, the value of a at which the maximum in I(AT ) falls below the measurable level decreases.

Experience indicates that AT' for pure metal and alloy melts is quite high, typically ^0.2 to 0.3 corresponding toa^ 0.6 when k^ ^ 1/picosec. This AT^0 would be the upper limiting interfacial undercooling in the rapid quenching of metal melts. Its magnitude is sufficient to permit formation of structures which depart widely from that at equilibrium.

The movement of a planar crystal-melt front occurs in a sequence of at least two steps: (1) interfacial rearrangement and (2) transport of the heat of crystallization away from the interface. In alloys the movement may, in addition, be limited by transport of impurity from the interface. In pure melts the speed, u, of the front is related to the interfacial rearrangement frequency, kj, as follows [6,7]:

AG/RT. u = fk.XCl-e *) (31 where f = fraction of surface sites at which growth can occur A = interatomic spacing AG = molar free energy of crystallization, equal to -AS AT. at small departures from equilibrium

In the linear kinetic regime (RT » |AGJ1 equation (3) becomes

AT. u = fkjXB ~ (4) i At steady state u must also be proportional to the rate of transport of the crystallization heat from the front and neglecting temperature non- uniformity in the liquid, it is related to (grad T)., the thermal gradient in the solid at the interface, as follows: ~l

K(grad T).V u ~- sp-i- (51 m where < is the thermal conductivity in the solid. Then from eqs. (4) and

K(grad T). VT. fk.A BAH" (61 3 m

SOLIDIFICATION OF PURE METAL MELTS

Arguments have been presented elsewhere that in the crystallization of pure metal melts (al f is near unity, (bl the interfacial rearrangement process is limited only by the collision frequency of atoms from the melt on the interface and it needs no thermal activation f8]. Thus, fk-A should approach as an upper limiting value the sound speed, us, in the liquid metal [1]. Also, it follows that a melt in contact with a crystallization front would not form a glass having an appreciable lifetime at any rate of quenching. Assigning typical values to the constants which appear in equation (6) 2 (e.g^, fkjX = u5 = 4000 meters/sec, K - 0.2 cal/cm sec.deg, 3=1, and 3 AHm/V = 300 cals/cm ) the maximum scaled interfacial undercooling at the crystallization front in a pure metal melt is estimated to be

ATi -q -=-=- = AT . ^ 1 x 10 " (grad T) . (7) ij, ri l

Q With the highest thermal gradients now realizable, M0 deg/cm in quenches (-T ^ 10*~deg/sec) following high energy pulses of a few pico- seconds duration [9], ATrj would only reach levels of order 0.10 which are well below AT^O, the scaled undercooling at measurable nucleation onset. Supporting this estimate, experiments of Lin and Spaepen [10] indicated that the crystallization front in pure liquid I-"e reached velocities as high as 500 meters/sec, in quenches following SO picosecond laser pulses, corresponding to ATp^ 'v- 0.10 [eq. (4) with fk^A *\J 4000 meters/sec]. Thus, even at such quench rates there should be no measurable homogeneous crystal nucleation in advance of the crystallization front in pure metal melts. Regrowth into pure molten layers produced by such processing should be epitaxial as was observed by Lin and Spaep^n [101. However, the interfacial undercooling should be sufficient for the generation of considerable densities of stacking raults and other extended imperfections.

When liquid layers are quenched on substrates with compositions very different from their own the crystallization front is formed by growth of n"clei originating in the near vicinity of the melt-substrate interface, m-s. If the substrate structure is very different from that, of the crystallized mets 1 the m-s temperature might fall to a level where- homogeneous nucleation occurs rapidly. Then a fine grained crystallization front would develop and evolve by further growth into a fine grained columnar structure. A similar course of structural evolution would be followed if nuclei formed heterogeneous ly a*: isolated sites on the substrate surface. Ther nucleation would begin at lesser undercooling and the ultimate grain density/area would be proportional to the number density of the sites activated.

In the foregoing it was assumed that the transient period for dei'el- opment of the steady r.ucleation frequency is negligible. A recent analysis of the transient by Kelton et a 1. [11] indicates that this assumption, while quite unsatisfactory in the glass transition regime, should be approximately valid for the highly labile regime which appears to extend to Tr well below 0.5 for pure liquid metals.

SOLIDIFICATION OF ALLOY MELTS

Crystal Growth In near equilibrium crystallization homophase impurities in the melt must, in general, be redistributed, but at high interfacial undercooling ''partitionless" crystallization, i.e. that unattended by any long range impurity transport, may become possible [12]. The thermodynamic conditions for non-equilibrium inclusion of impurities in growing crystals as well as partitionless crystallization have been defined by Baker and Cahn [13,14] and recently reviewed by Boettinger et at. [15]. Partitionless crystallization may occur when the interfacial undercooling exceeds that at temperature T0(x) where the free energies of crystal and melt alloy phases of identical composition, x, are the same. In this process, impurity may actually be "trapped" in the crystal at a higher chemical potential than it had in the melt provided the total free energy of crystallization is negative. The course of T0(xJ for a binary alloy with a falling liquidus, T^(x) , is shown schematically in Figure 1. Often T0(x") is, for one or another alloy structure, well above 0°K at every composition, meaning that at any x the alloy will have a lower energy in some "homogeneous crystallized chan in an amorphous solid form. However, there are some alloy systems having solution thermodynamic parameters such that the To(x) for both terminal structures plunge to zero, say at x0 and xo, on non-intersecting courses so that at any x between x0 and xo the alloy at 0°K will be more stable as an amorphous solid than in any homogeneous crystalline form. We have to consider two regimes, one with and the other without solute partitioning, and the transition from one regime to the other, of crystal growth in alloy melts.

Liquidus, Tj (x)

Limiting Temperature T0(x) of t Partitionless Crystallization T Nucleation onset Temperature To'(x)

Glass Temperature /TgU)

Figure 1. Schematic of composition dependences of homogeneous nucleation, onset temperature, To(x), and maximum temperature To(x) for partitionless crystallization. Consider a planar front between a crystalline alloy and a thin layer of its melt where heat is extracted through the crystal. The front moves and the impurity redistributes itself in response to an imposed thermal gradient and impurity concentration in the body of the melt. Given the atomic and thermal transport coefficients and the thermodynamic parameters the front velocity and impurity redistribution, including the conditions at transition to partitionless growth, can, in principle, be calculated [16].

When impurity is rejected into the melt the velocity of the front should at steady state, and if the liquid solution is dilute, satisfy the relation: D. u = - ^i- (grad c) . i where c. = the impurity concentration in the melt at the interface i. (grad c)j = the concentration gradient in the melt at i. Dj = the impurity diffusivity in the melt at i.

The upper limiting speed of the front imposed by this relation is D.

With values D^ ^ 2 * 10~5cm2sec~ ' and X^ 2 x lutein, typical of liquid melts, u £ 10 meters/sec, which is more than 2 orders of magnitude below the maximum velocity possible f^4000 meters/sec) in collision limited growth. The impurity transport in melts is thermally activated and its rate generally falls to negligibly low values as the temperature is lowered through the glass transition range.

?n the partitionless regime the formal analysis of the growth rate is similar to that for pure metals, but with TQ(x~) replacing the crystal-melt equilibrium temperature. However, two limiting microscopic growth mechanisms may be distinguished. In one, crystallization proceeds with no change in impurity near-neighbor coordination, i.e. without reordering of the compositional short range order (C-SRO) of the melt. Then the growth could be collision limited and k-jX might, as in pure melts, approach the sound speed. In the other limit where some essential reordering of the C-SRO is necessary for crystallization, k^X would be approximately equal to Dj/X. As when partition occurs, this growth should be thermally activated and become negligible at temperatures well below T

Kinetic analyse; [16,17] indicate that the transition from growth with to that without partition should occur within about one order of magnitude of growth velocity centering at u^Dj/X. In this transition, the actual solute distribution coefficient goes from near equilibrium to a level near unity corresponding to complete solute trapping by the crystal.

Nucleation

Experience indicates that in the labile range the temperature To(x) at measurable nucleation onset in alloy melts roughly parallels the liquidus Tp (x). This behavior was accounted for by Thompson and Spaepen [18] by assuming that (a) The scaled crystal-melt tension, a, as in Spaepen's model, is mainly topological in origin and so only weakly dependent on x. (b) The crystal nucleus is always in interaction with melt of average composition.

If Tβ(xj falls sharply with x and if, as is often observed, Tg(x) varies little with x, TQ(x) on a course parallel with Tg(x) would cross T (x) at some composition. However, provided k^ always scales with D j, a! x increases, T'(x) will fall below its parallel course and drop to zero before reaching its projected intersection with Tg(x).

Glass Formation

From the foregoing it follows that a necessary condition for metal glass formation in melt quenching and, indeed, in condensation on substrates containing crystal nucleants, is That some redistribution of impurity, either by partition or short range reordering, be thermodynamically demanded for crystallization [20,5,8,7]. Then the interfacial rearrangement frequency has to scale with Dj and so would be thermally activated and suppressed to negligible levels at temperatures below To.

In addition to the impurity redistribution requirement the melt quench rate between Tofx) and Tp(x) must exceed the critical value, -f_r, for bypassing homogeneous crystal nucleaticn. Applying simple nucleation theoi/ with the scaled tension a^0.6, as indicated by the nucleation resistance of pure metals, and assuming that kj scales with reciprocal shear viscosity as described by the Vogel-Fulcher equation. Spaepen and the writer [21] related -T to the scaled glass temperature, Tra- The relation indicates that -fcr should fall sharply with increasing Trg to 6 10 deg/sec at Trg^l/2 and 1 deg/sec at Trg^2/3. Since Tg(x) is relatively insensitive to composition Trg(x) should increase and -Tcr decrease with a falling liquidus T#(x) so that -fcr should be near a minimum at minimum T^(x).

These considerations indicate that those alloy melts with Trg >_ 0.4 and in which crystal growth is reconstructive and thermally activated should be capable of forming glass when quenched at the highest rates now realizable.

Boettinger [22] has demonstrated that in certain alloy systems where partitioning in crystallization is required, the compositions most favorable for glass formation are, indeed, those falling in the estimated x0 to x0 range over which the alloy is more stable in amorphous solid than in homogeneous crystalline form. In contrast, recent, experiments of Lin and Spaepen [23] clearly show that some alloys with T (xl well above Tg(x) can be quenched to glasses. Apparently, in these alloys the kinetic resistance to short range reordering, without partition, retarded crystal growth sufficiently for T to be reached without appreciable crystallization. Some observations of Bagley and the writer [24] indicated an interesting transition, with a composition change of only 0.1 at.%, from thermally activated partitionless to partitioning growth of Ni^P crystals in amorphous films. The partitionless growth resulted in a monophasic polyhedral crystal morphology while growth with segregation required temperatures more than 75° higher foT an equivalent growth rate and led to a dendritic morphology. Grain Structure of Rapidly Solidified Alloys

When partitionless crystal growth occurs without thermal activition, it appears that homogeneous nucleation in advance of the crystallization front should be as unlikely as in pure metals and so the grain morphology of the alloy should resemble that of pure metals crystallized under similar conditions.

When growth is thermally activated, there are conditions where fine equiaxed grain structures should develop. In the crystallization of thin molten alloy layers, attended by impurity redistribution, the constitutional undercooling can become very large, even with high thermal gradients into the solid at the interface. Thus, interface shape instabilities might develop so that crystallites could be injected into the melt by the "meltback" mechanism [25]. However, the minimum times for onset of shape instability or the liquid thickness needed for initiation of the meltback mechanism have not been evaluated. Such evaluations might indicate a minimum thickness of molten alloy film required for forming an equiaxed grain structure.

Under the conditions that growth is thermally activated and Trg(x) is relatively high, in sufficiently rapid quenches the interface temperature may be lowered to levels below To(x) so that the frequency of homogeneous nucleation in advance of the interface becomes appreciable. The alloy melt might then crystallize to an equiaxed grain structure or, at a somewhat higher quench rate, undercool to a glass in which a high number density of microcrystallites are embedded. The latter structure would evolve to an equiaxed microcrystalline one upon reheating. The individual crystallites might be mono- or bi-phasic depending on the alloy composition. Thompson et at. [26] have suggested that the number density of crystallites formed in treatments of this type can be limited considerably by the transient times needed to reach steady state nucleation. Experience, indeed, indicates [27] that certain glass forming alloys crystallize to equiaxed fine grained structures when quenched at intermediate rates. Also there is considerable experience which shows that glasses of such alloys form similar structures, often composed by biphasic grains, upon reheating.

ACKNOWLEDGMENTS

Much of this paper was written during the author's stay at the National Bureau of Standards [November 1983) as a visiting scientist in the Center for Materials Science. The Harvard phase of the research was supported by ONR Contract N00014-77-C-0002 and MRL NSF Contract DMR 80-20247. The author is indebted to W.J. Boettinger and S.C. Coriell (NBS) and M.J. Aziz, A.L. Greer and F. Spaepen (Harvard) for useful discussions.

REFERENCES

1. D. Turnbull, Met. Tra->s. 12A, 695 (1981). 2. D. Turnbull, Ann. Rev. Mat. Sci. (ed. R.A. Huggins, R.H. Bube and D.A. Vermilyea) 13_, pp. 1-7 (1983). 3. See discussion in M. Volmer, "Kinetik der Phasenbildung," pp. 200-05, Steinkopff, Dresden (1939). 10

4. R.E. Cech, Trans. A.I.M.E. 206, 585 f!956). 5. D. Turnbull, Contemp. Phys. 10, 473 (1969). 6. D. Turnbull and M.H. Cohen, "Modem Aspects of the Vitreous State" (J.D. Mackenzie, ed.) !_, p. 38, Butterworth's, London (1960). 7. F. Spaepen and D. Turnbull in "Laser Annealing of Semiconductors" (J.M. Poate and J.W. Mayer, eds.), pp. 15-4], Academic Press, New York (1982). 8. D. Turnbull and B.G. Bagley in "Treatise on Solid State Chemistry" (ed. N. B. Hannay) 5_, p. 513, Plenum, New York (1975). 9. See N. Bloembergen, "Laser-Solid Interactions and Laser Processing" (S.D. Ferris, H.J. Leamy and J.M. Poate, eds.), A.I.P. Conference Proc. 50_, pp. 1-10 (1979). 10a. C.J. Lin and F. Spaepen in "Chemistry and Physics of Rapidly Solidi- fied Materials" (B.J. Berkowitz and R.O. Scattergood, eds.), pp. 273-280, A.I.M.E. Conference Proc. (1983). 10b. See also S.C. Coriell and D. Turnbull, Acta Met. 30^, 2135 (1982). 11. K.F. Kelton, A.L. Greer and C.V. Thompson, J. Chem. Phys. 79_, 6261 (1984). 12. H. Biloni and B. Chalmers, Trans. Met. Soc. A.I.M.E- 2S3, 375 (1965). 13. J.C. Baker and J.W. Cahn, Acta Met. 17_, 575 (1969). 14. J.C. Baker and J.W. Cahn, "Solidification", pp. 23-58, Am. Soc. Metals, Metals Park, Ohio (1970). 15. W.J. Boettinger, S.R. Coricll and R.F. Sekerka, in pres.-;, Mats. Sci . and Eng. (1984). 16. M.J. Aziz, J. Appl. Phys. 53_, 1158 (1982); see also for references to earlier literature. 17. M.J. Aziz, Appl. Phys. Lett. 43_, 552 (1983). 18. C.V. Thompson and F. Spaepen, MRS Symposia Proc. (ed. G.E. Rindone) 9_, p. 603, Amsterdam, North Holland (1982); see also, Acta Met., 51_, 2021 (1983). 19a. F. Spaepen, Acta Met. 23, 729 (1975). 19b. F. Spaepen and R.B. Meyer, Scripta Met. H), 257 (1976). 20. D. Turnbull, "Solidification", pp. 1-22, Am. Soc. Metals, Metals Park, Ohio (1970). 21. F. Spaepen and D. Turnbull, "Rapidly Quenched Metals" (N.J. Grant and B.G. Giessen, eds.) pp. 205-229, MIT Press, Cambridge. Mass. 22. W.J. Boettinger, "Proc. of 4th Intl. Conf. on Rapidly Quenched Metals" (T. Masumoto and K. Suzuki, eds.) J_, pp. 99-102, Japan Inst, of Metals, Tokyo (1982). 23. C.J. Lin and F. Spaepen, Appl. Phys. Lett. 4J_, 716 (1982). 24. B.G. Bagley and D. Turnbull, Acta Met. 18, 857 (1970); see Figs. 5 and 7. 25. K.A. Jackson, J.D. Hunt, D.R. Uhlmann and T.P. Seward III, Trans. Met. Soc. A.I.M.E. Z56_, 149 (1966). 26. C.V. Thompson, A.L. Greer and A.J. Drehman, "Proc. of the 4th Intl. Conf. on Rapidly Quenched Metals" (T. Masumoto and K. Suzuki, eds.) J_, pp. 743-746, Japan Inst. Metals, Tokyo (1982). 27. A.J. Drehman, Ph.D. Thesis, pp. 100-102, Division of Applied Sciences, Harvard University, Cambridge, Mass. (1983). 11

SOME TECHNIQUES FOR THE STUDY OF ATOMIC MOTIONS WITH APPLICATIONS TO CERAMIC MATERIALS

A.S. Nowick

Henry Krumb School of Mines Columbia University, New York, N.Y. 10027, USA

ABSTRACT

A review is presented of the use of the techniques of a.c. impedance measurements, dielectric relaxation aril anelastic relaxation to study technologically interesting ceramic materials. Illustrations are given of the application of these methods to the study of solid electrolytes as well as quartz crystals used for frequency control.

INTRODUCTION

Knowledge of rates of atom movements plays an important role in the understanding of most materials. In the case of ceramic (ionic) materials, the migrating species are often charged ions, and then their migration rates can be studied by electrical means. Such studies are of particular interest in the case of materials that are relatively good ionic conductors, the so-called solid electrolytes or "superionic conductors" which are of interest as electrolytes for solid-state batteries and fuel cells. Such applications are of sufficient interest that in recent years there has been an outpouring of review books [1- 3] and conference proceedings [4-7] in this area. Materials studied include those that ccd-jct by migration of alkali ions, silver ions, protons, oxygen io" arn fk j.ine ions, among others.

The present paper will review three techniques that have been ac- tively used in recent years in the author's laboratory, with illustrations of the kind of information, both basic and applied, derivable from them. The techniques are: (1) a.c. impedance measurement, (2) dielectric relaxation, and (3) anelastic relaxation. We will see that valuable information about technologically important systems has been obtained using these techniques.

A.C. IMPEDANCE MEASUREMENT AND ANALYSIS

An ideal ionically conducting material should be represented elec- trically as a capacitance C in parallel with a resistance R, the former representing the dielectric properties of the medium and the latter its conduction. Real materials are not so simple, however, due to a number of complications. First, there is the problem of introducing and discharging the conducting ionic species at the respective electrodes. Second, one finds another effect, loosely called the "grain-boundary" effect, which is due to blocking of carriers at internal interfaces within the electrolyte. As a result of these factors the actual equivalent circuit of a sample can be extremely complex. A relatively simple representation is shown in 12

Fig. la, consisting of three R-C units in series with each other, one representing the bulk or lattice behavior, another the grain-boundary effect and the final one, the electrode effect. In order to study such an equivalent circuit, a.c. measurements are made over as wide a frequency range ("frequency window") as possible (usually 1 Hz to 106 Hz). The sample can then be represented by its complex admittance, Y* containing both real and imaginary parts (representing the current in-phase with and 90° out-of- phase with the applied voltage). One may write Y* = G(w) + iuC(u) (1) where G(ui) is the effective conductance and C(CD) the capacitance both, in general, functions of the frequency. We may also introduce the reciprocal quantity, Z* = 1/Y*. called the complex impedance, which also has real and imaginary parts Z* = Z'(u>) - iZ"(u) (2) A convenient and widely used analysis is the complex impedance plot in which Z" is plotted as a function of Z1 with to as parameter. For the case of the equivalent circuit shown in Fig. la, such a plot yields three semi- circular arcs, one for each R-C unit, as shown in Fig. lb. In this plot the frequency increases as we go from right to left (arc l to arc 3). In this way, it becomes possible to separate out the true lattice conduction from the grain-boundary and electrode phenomena, and to study each separately. (a) Rμ, R. —wv 0 50 Lac to.c.

5067 •c 0 : 6 V. 1^0} c. 2

500 • o 0.92 »10"Zalm © • t.O » 10"*a!m (W / V 1

Fig. la. Equivalent cir-uit representing lattice [I), grain boundary (gb) and electrode (e) effects. Fig. lb. Schematic diagram of com- Fig. 2. Examples of complex plex impedance plot corresponding impedance plots for CeO2:6% to circuit of part (a). Arrow Y2O3 at three different temper- indicates direction of increasing atures. Arcs are labeled to frequency. match Fig. lb. From Ref. [8]. 13

Actually, the experimental "frequency window" is usually not wide enough to display all three effects at any one temperature, but by covering a range of temperatures, we can see all of them. Figure 2 shows an 3+ example for the case of Y doped CeO2- This material is an oxygen-ion conductor through the migration of oxygen-ion vacancies which are introduced into the lattice as charge compensation for the lower valerit cation dopant. Such ionic conductors may be used as the solid electrolyte in high-temperature fuel cells or in oxygen sensors. From Fig. 2, we see that at Tow temperatures (e.g. 178 °C) arcs 2 and 3 appear while at high temperatures arcs 1 and 2 are obtained. Further we see that only arc 1, the electrode arc, has a strong dependence on partial pressure of oxygen of the ambient gas. These results show, however, that the behavior is more complex than that predicted from the schematic diagram of Fig. lb. Firstly, the arcs are not full semi-circles but are somewhat depressed (as shown by the sloping lines drawn in Fig. 2). This result can be interpreted as meaning that each arc corresponds to a narrow distribution of R-C circuit elements rather than to single values. For simplicity we ignore this fact for the present. Secondly, it is found that the electrode contribution can involve more than a single arc. This will be discussed below. Nevertheless, the model of Fig. lb can serve as the basis ;T extracting appropriate parameters. The two intersection points R,~ and R23 are clearly related to the equivalent-circuit parameters by

R23 (3) and (4) R,12

Fig. 3. Conductivity plots for a CeO2:l% Y2O3 sample. Upper curve, obtained from the intersection R23 represents the lattice conductivity. Lower curve, consisting of a.c. measurements from R12 intersection, as well as d.c. 4-probe data, is dominated by the grain-boundary effect. From Ref. [9].

2.4 ZJo 14 electrolyte. Both curves can be fitted to the conventional Arrhenius-type relation al = B exp (-H/kT) (6) in which H is the activation enthalpy. We have obtained useful information from each of the three arcs. From arc 3 we have obtained the lattice conductivity as a function of temperature, composition and dopant [9,10]. Considering that oxygen-ion vacancies are introduced as charge compensation for the dopant ions, one might expect a monotonically increasing conductivity as the dopant concentration increases. It is found instead that a goes through a sharp maximum as a function of dopant concentration as shown for Y2O3 dopant in Fig. 4. At the same time, the activation enthalpy goes through a minimum. These results show that strong defect interactions suppress the conductivity beyond a concentration of 4-6 mole % dopant. The study of defect interactions in systems of high concentration is a subject of great interest and one which requires further study. Similarly, it was found that for different trivalent cation dopants of the same composition, the highest conductivity and lowest activation enthalpy H occur for a dopant whose size is closest to that of the host cation [10]. Such a result means that strain energy terms play an important role in determining H, a result that has recently been verified by computer simulation calculations [11]. From a practical viewpoint, these studies show how best to optimize the lattice conductivity with respect to dopant species and dopant concentration.

In most applications of solid electrolytes, one wishes to achieve the maximum possible overall d.c. conductivity. Accordingly the grain-boundary effect cannot be ignored. In fact, Fig. 3 shows that the overall conductivity can be as much as lOOx lower than the lattice conductivity due to the grain-boundary effect. Similar effects are prevalent for other solid electrolytes, e.g. in β-alumina which is the electrolyte in the much pub- licized sodium-sulfur battery [12]. In the present case, particularly because Hgb » r'-i (see Fig. 3), we can only interpret the electrical

(il-cm),-'1

Fig. 4. Variation of conductivity (ev) . a (et 182 °C) and of activation enthalpy with composition, for CeOo - YoOq solid solutions. From Reff [9]fJ

O-β -

0.6 05 .1 15 effects in terms of the existence of a continuous (very poorly conducting) blocking layer, presumably associated with grain boundaries. In order to verify the presence of such a layer and to determine its origin, we examined thinned (ion-milled) samples by STEM (scanning transmission electron microscopy) in combination with microanalysis by EDAX (Energy Dispersive X-ray Microanalysis) and EELS (Electron Energy Loss Spectroscopy) [13]. A number of microstructural features were observed in this study, but the one most relevant seemed to be the presence of "thick boundaries", as shown in Fig. 5, i.e. layers of ^SOOA thickness which seem to be continuous. These were generally found to be an amorphous phase which had Si as its major cationic constituent. This silica-type phase did not surround each grain, however, but appeared around a small agglomerate of grains. In order to determine whether the presence of Si as an impurity is indeed responsible for the grain-boundary effect, we have very recently succeeded in preparing doped ceria samples from starting materials that were essentially silicon free. Our measurements show that in such a material the grain boundary effect is virtually eliminated.

The remaining arc (arc 1 in Figs, lb ar. 2) is that due to electrode effects. Actually electrode processes are very complex and rarely give rise to just a single arc [14]. For the case of oxygen-ion conductors, we have carried out extensive studies with porous Pt-paste electrodes as well as a variety of others [15]. The simplest behavior (i.e., a single depressed arc) was observed for freshly prepared Pt-paste electrodes at relatively high oxygen partial pressures. With the aid of an auxiliary technique, called the current-interruption method [15], it was possible to show that, in this simplest case, the electrode process is controlled by a charge-transfer mechanism, in which an adsorbed oxygen adatom 0a(j under- goes the reaction

2e + V °ad " + vo = ad where VQ and Vacj are lattice oxygen vacancy and vacant adsorbed site, respectively. The cathodic reaction goes from left to right and the anodic

Fig. 5. STEM microstructure of a sintered CeO2:6% Gd2(h sample showing thick boundaries" of an amorphous silica phase. From Ref. [13]. 16 in reverse. At low oxygen pressures or when the electrode has aged under high temperature or high current, an increase in Z" on the low-frequency side of the arc takes place. This is related to a transport limited process, or "concentration polarization." DIELECTRIC AND ANELASTIC RELAXATION Because of important similarities we will treat these two techniques together. In both cases an alternating field is used, electric and stress fields, respectively. We measure the fractional energy loss per cycle in the form of the quantity tan 6, where 6 is the angle by which the response (polarization or strain, respectively) lags behind the applied field. In both cases the simplest behavior takes the form of the well-known Debye peak as a function of frequency:

tan 6 = A • U)T/(1 + U2T2) (7) where A is a measure of the strength of the relaxation process and T is the relaxation time. In the case of dielectric or anelastic relaxations due to defect pairs (or higher clusters), A is proportional to the defect concentration and to the square of the electric or elastic dipcle strength of the defect, respectively. On the other hand, the reciprocal x"1 is related to the frequency of the controlling ionic migration process and is therefore thermally activated: 1 T- = v0 exp (-Hr/kT) (8) where H is the activation enthalpy for relaxation. Because of Eqs. (7) and (8)r, it is possible to observe tan 6 either as a peak in frequency at constant temperature or a peak in temperature at constant frequency. For more complex processes, tan 6 must be written as a summation over expressions of the type Eq.(7),and the corresponding peak is then broader than a simple Debye peak [16]. For dielectric relaxation an alternative to the a.c. tan 6 measurement is available. It is called the TSDC (thermally stimulated depolarization current) method (also known as ITC) [17]. This method offers an extremely high sensitivity and therefore can detect the reorientation, under electric field, of electric dipoles at concentrations as low as 1 ppm. It has been applied to ceramics of the type already discussed [18], but for lack of space we will not be able to cover this work here. To illustrate the use of these relaxation techniques we turn instead to a different problem, that of the frequency instability of quartz crystals. It is well known that α-quartz, due to its piezoelectric property, can be fabricated into resonators that vibrate at a fixed frequency. A common application is to control time pieces (i.e. watches). But applications of such resonators to controlling satellites and guidance systems demand far greater frequency stability, often to as much as 1 part in 109 or 1010, even in radiation environments. Frequency instabilities under irradiation are found to be closely linked to impurities present in these crystals, one of the most important of which are alkali ions, e.g. Na+ or Li+ [19]. These alkalis are present in the crystal as compensation for the impurity Al3+, which sits on Si"+ site and is deficient one positive charge. In recent years, techniques have been developed for replacing alkalis by H+ or one alkali with another by a process of electrodiffusion or 17

50 - QUARTZ OSCILLATORS ACCUMULATED OFFSET — Fig. 6. tffect of * 30 10 MeV ELECTRONS Z-GROWTH electron irradiation 10 DtffP SYNTHETIC on the room-temperature frequency change, 1 \- Q BBS El A AAAAAAtt 6f, of 5 MHz quartz- <-> o _ ° O O O SWEPT Z-GROWTH crystal resonators. O SYNTHETIC From Ref. [21]. o -10 o o s °o 2 -30 -50 - ^NATURAL 1 1 I 104 105 ' 10w 10' RADS (Si)

"sweeping" [20]. In this process the crystal is treated at elevated temperature under a strong d.c. electric field in an appropriate environment. Also, methods of hydrothermal growth have made possible the production of synthetic quartz crystals that are often far lower in impurity content than natural crystals. Figure 6 shows the effect of electron irradiation at room temperature on three quartz crystals, a synthetic and a natural, both as grown, and on the same synthetic crystal after sweeping to replace alkali ions with H+. The scale of frequency changes in this graph is very large compared to the best present-day requirements. Nevertheless, the figure shows that H+ sweeping greatly reduces the frequency change that takes place under irradiation. Thus we see that alkalis contribute to 6f. One reason for a frequency offset at room temperature can be the occurrence of an anelastic relaxation peak at lower temperatures. The

Fig. 7. Anelastic loss (internal friction, Q"1) of a Na- doped quartz crystal as a function of temperature. From Ref. [20].

* .f0-100 '» '«> INIMIWJJO 18

formal theory of anelasticity [16] shows that the presence of in anelastic peak given by Eq. (7) produces a frequency depression

Sf/f = - A/2 (9) at all temperatures above the peak, where 5f is the change in frequency relative to a defect-free crystal. In other words, to understand a frequency offset at room temperature requires that we know about relaxation peaks that occur below room temperature. Such peaks are indeed prevalent. For quartz crystals containing Na+, a prominent pair of anelastic peaks is found at low temperatures, as shown in Fig. 7. A comparable pair of peaks is observed in dielectric relaxation [22]. These have low activation energies (Hr = 0.05 and 0.14 eV, respectively) and are attributed to the Na+ being bound to a substitutional Al3+, with the Na+ residing in the interstitial channel. No such peaks due to Al-Li pairs have been found in spite of a careful search for them [23], presumably because the Li+ sits on a 2-fold axis of symmetry. Irradiation produces important changes in these defects. Electron-hole pairs generated serve to free alkali ions from the Al3+ and to create Al-h (aluminum-hole) pairs instead. The migration of the alkali can be observed as conductivity enhancement immediately after irradiation [24]. It is not yet clear which defects serve as the traps at which the alkali ions terminate, but such trapping sites probably also capture one or more electrons. The Al-Na peaks are eliminated or reduced by the irradiation and, instead, new large peaks can be observed both by anelastic and dielectric loss measurements. Figure 8 shows such dielectric loss peaks for both Na+- and Li+- containing crystals. Similar anelastic peaks have been observed [25]. It is not yet clear as to whether these peaks are due to Al-h centers or to alkali centers which are created by the trapping of Li+ or Na+ freed by the irradiation. Nevertheless,

120 .TO5 T0Y0 SUPREME tIRRAD.) •

100 jj^Li SWEPT \. n \ 80 - i 1 Fig. 8. Dielectric loss of Li- ' '" No SWEPT / swept, Na-swept and H-swept synthetic quartz following \\ room-temperature X-irradiation. 60 1 From Ref. [24].

40 IT _^ H SWEPT 20 - // // A. i • t J 1 1 L 0 4 8 12 16 20 24 28 32

100/T 19

Eq. (9) shows that the elimination of a low-temperature anelastic relaxation peak by irradiation results in an increase in frequency, while the creation of a new peak results in a frequency decrease. Thus, it is clear that the solution to questions concerning frequency instabilities at room temperature lies in a better understanding of these low-temperature relaxation phenomena. In conclusion, it was intended to show how the methods described herein give insight into ionic migration in ceramic materials. These methods can be usefully complemented by other widely used techniques such as diffusion and NMR measurements. Together, they give the type of insight into defect migration mechanisms that make it possible to eliminate practical problems or to determine conditions of maximum performance.

ACKNOWLEDGMENTS

The author wishes to acknowledge the contributions of his present and former associates: Drs. Da Yu Weng, H. Jain, J. Toulouse and R. Gerhardt- Anderson, and the support received from both the U.S. Department of Energy and the U.S. Air Force, RADC.

REFERENCES 1. S. Geller, ed., "Solid Electrolytes", Topics in Applied Physics vol. 21, Springer-Verlag, Berlin, 1977. 2. P. Hagenmuller and W. van Gool, eds., "Solid Electrolytes", Academic Press, New York, 1978. 3. E.C. Subbarao, ed., "Solid Electrolytes and Their Applications", Plenum Press, New York, 1980. 4. W. van Gool, ed., "Fast Ion Transport in Solids", North-Holland, Amsterdam, 1973. 5. G.D. Mahan and W.F. Roth, eds., "Superionic Conductors", Plenum Press, New York, 1976. 6. P. Vashishta, J.N. Mundy and G.K. Shenoy, eds., "Fast Ion Transport in Solids", North-Holland, Amsterdam, 1979. 7. J.B. Bates and G.C. Farrington, eds., "Fast Ionic Transport in Solids", North-Holland, Amsterdam, 1981. 8. Da Yu Wang and A.S. Nowick, J. Solid State Chem. 35 (1980) 325. 9. Da Yu Wang, D.S. Park, J. Griffith and A.S. Nowick, Solid State Ionics 2 (1981) 95. 10. R. Gerhardt-Anderson and A.S. Nowick, Solid State Ionics 5 (1981) 547. 11. V. Butler, C.R.A. Catlow, B.E.F. Fender and J.H. Harding, Solid State Ionics 8 (1983) 109. 12. R.W. Powers and S.P. Mitoff, J. Electrochem. Soc. 122 (1975) 226. 20

13. R. Gerhardt-Anderson, A.S. Nowick, M.E. Mochel and I. Dumler, Proc. Conf. High Temperature Solid Oxide Electrolytes, ed. F.J. Salzano, Brookhaven National Laboratory 1983, vol. 1, p. 225. 14. J.R. Macdonald, in "Electrode Processes in Solid State Ionics", M. Kleitz and J. Dupuy, eds., p. 149, Reidel Publ. Co., Dortrecht- Holland, 1976. 15. Da Yu Wang and A.S. Nowick, J. Electrochem. Soc. 126 (1979) 1155, 1166; 127 (1980) 113; 128 (1981) 55. 16. A.S. Nowick and B.S. Berry, "Anelastic Relaxation in Crystalline Solids", Academic Press, New York, 1972. 17. C. Bucci, R. Fieschi and G. Guidi, Phys. Rev. 148 (1966) 816. 18. Da Yu Wang and A.S. Nowick, J. Phys. Chem. Solids 44 (1983) 639. 19. J.C. King and H.H. Sander, Rad. Effects 26 (1975) 203. 20. D.B. Fraser, in "Physical Acoustics" vol. 5, W.P. Mason, ed., p. 59, Academic Press, New York, 1968. 21. B.R. Capone, A. Kahan, R.N. Brown and J.R. Bucfrnelter, IEEE Trans, on Nucl. Sci. NS-13 (1966) 130. 22. D.S. Park and A.S. Nowick, Phys, Stat. Sol. (a) 26 (1974) 617. 23. J. Toulouse, E.R. Green and A.S. Nowick, in "Proc. 37th Ann. Symp. on Frequency Control", 1983, p. 125, U.S. Army EP.ADCOM. 24. H. Jain and A.S. Nowick, J. Appl. Phys. 53 (1982) 485. 25. J.J. Martin, L.E. Halliburton and R.B. Bossoli, in "Proc. 35th Ann. Symp. on Frequency Control", 1981, p. 317. 21

THE CHEMISTRY ON A SUBNANOMETER SCALE OF RADIATION-INDUCED PRECIPITATION AND SEGREGATION IN FAST-NEUTRON IRRADIATED TUNGSTEN-RHENIUM ALLOYS

Roman Herschitz and David N. Seidman

Cornell Un.versity, Department of Materials Science and Engineering and the Materials Science Center, Ithaca, New York 14853-0121

ABSTRACT

The phenomena of radiation-induced precipitation and segregation have been investigated in W-10 at.% Re and W-25 at.% Re alloys, employing the atom-probe field-ion-microscope technique. The W-10 at.% Re alloy is subsaturated with respect to the solvus line of the primary solid solution (3 phase), while the W-25 at.% Re alloy is supersaturated with respect to the same solvus line. The specimens had been irradiated in the Experimental Breeder Reactor II to a fast-neutron fluence of 1.4x10 neutrons cm"2 (E>0.1 MeV) at 575, 625 and 675 C. This corresponds to 8.6 dpa and an average displacement rate, for the iiwo year irradiation time, of 1.4x10" dpa s~. The results of the present investigation show a very significant alteration of the microstructure of both alloys as a result of the fast-neutron irradiation. In the case of the W-10 at.% Rβ alloy coherent, semicoherent and possibly incoherent precipitates with the composition ^WRe and a disc-shaped morphology — one or two atomic planes thick — were detected at a number density of n<10^ cm , and a mean diameter of o>57 A. For the W-25 at.% Re alloy coherent, semicoherent and incoherent precipitates with the composition ^WRe3 were detected; the precipitate's number density is vLoJ-7 cm"3 with a mean diameter of 40 X. None of the ^WRe precipitates or the ^WRe3 coherent precipitates were associated with either line or planar defects or with any impurity atoms. Therefore, a true homogeneous radiation-induced precipitation occurs in these alloys. The semicoherent WRe3 precipitates were associated with He atoms; that is, these precipitates may have been heterogeneously nucleated. In the W-25 at.% Re alloy a two dimensional WRe3 phase has been observed at a grain boundary. A physical argument is pr:^rnted for the nucleation of WRe or WRe3 precipitates in the vicinity of displacement cascades produced by primary knock-on atoms. It is suggested that in both cases the first step in the nucleation of a precipitate is due to the formation of tightly-bound mobile mixed dumbbells which react to form an immobile di-rhenium cluster. Possible sequences of point-defect reactions which can lead to either WRe or WRe3 cluster are detailed. The further growth of a cluster (WRe or WRe3> into a precipitate is most likely driven by the irreversible vacancy: self-interstitial atom annihilation reaction, as suggested recently by Cauvin and Martin.24 Point-defect mechanisms for all the other observations are also discussed. INTRODUCTION

Over the last few years there has been a rapid growth of interest in the phenomena of radiation-induced (as opposed to accelerated) segregation and precipitation.1-4 Different types of irradiation — electrons, ions or neutrons — can induce significant segregation of alloying elements either toward or away from grain boundaries, voids or free surfaces. Radiation can also cause the heterogenous or homogeneous precipitation of a phase in subsaturated solid solutions and it can alter the phase stability of alloys. Radiation-induced segregation and precipitation are of paramount technological importance, since they play a crucial role in the nucleation and growth of voids and have a strong effect on the physical properties of metals alloys used in the fuel cladding and core structure of the fast-breeder reactor, as well as in the materials used for the first wall of fusion reactors. These phenomena are also of considerable fundamental interest.

The study of W(Re) alloys is of technological importance, as they are used in thermocouples for the measurement of temperature in nuclear reactors. As a result of an exposure to a fast-neutron flux the decalibration of W(Re) thermocouples occurs. ' The alloys W-10 at.% Re and W-25 at.% Re are of particular interest in understanding the phenomenon of radiation-induced precipitation, as the former alloy is subsaturated with respect to the solvus line of the primary solid solution (6 phase), while the latter alloy is supersaturated with respect to this solvus line — it is in the 3 plus a phase field.7' Sikka and Moteff^ and Williams e_t al_- have identified the crystal structure of radiation-induced precipitates in fast-neutron irradiated W-25 at.% Re alloys using transmission electron microscopy — for specimens which had been irradiated at 1100 C and higher — and it corresponds to the y.-phase which has the composition WRe3- Williams et^ al^. 10 also investigated fast-neutron irradiated W-5 at.% Re and W-ll at.% Re alloys. For specimens which had been irradiated at 1100 C and above all the precipitates analyzed by electron diffraction were consistent with the x-phase crystal structure. Whereas Williams e_t al_- were unable to obtain interpretable electron diffraction patterns from any of the specimens which had been fast-neutron irradiated at 900 C or lower.

In this short summary paper we present the results of an extensive atom-probe field-ion microscope (FIM) study of radiation-induced precipitation and segregation in fast-neutron irradiated W-10 at.% Re and W-25 at.% Re alloys. ' Our atom probe FIM allows us to determine the chemical identity of all the elements in the periodic tabled3"16 In addition, the atom-probe FIM has a lateral spatial resolution, for chemistry, of a few-tenths of a nanometer and a depth resolution which is determined by the interplanar spacing of the region being analyzed. We found very significant alterations of the microstructure between 575 to 675 C. In the case of the W-10 at.% Re alloy precipitates with the composition ^WRe (o phase) were detected at a number density of VL0l6cm~3. They were not associated with linear or planar defects or with any impurity atoms; i.e., a true homogeneous radiation-induced precipitation occurs in this alloy. Coherent, semicoherent and incoherent precipitates were detected. For the W-25 at.% Re alloy coherent, semicoherent and incoherent precipitates with the composition 23

(X phase) were detected at a number density ^1017cm~3. The coherent precipitates were not associated with either line or planar defects, or with any impurity atoms. This strongly suggests that the coherent WRe3 precipitates were a result of a homogeneous radiation- induced process. The semicoherent and incoherent MJRej precipitates were found to be associated with He atoms; i.e., they may have been heterogeneously nucleated. In addition we found evidence for a two- dimensional t»WRe3 phase at a grain boundary; this phase is the result of a radiation-induced segregation process.

EXPERIMENTAL DETAILS

Wire specimens of W(Re) alloys were irradiated to a fast neutron- fluence of i-4xlO22 neutrons cm"2 (E>0.1 MeV) at elevated temperatures (575, 625 and 675°C) in Experimental Breeder Reactor II (EBR-II) at Richland, Washington. This corresponds to 8.6 dpa for row 7 of EBR-II. Hence, the average displacement rate for the two year irradiation time is 1.4xlO~7 dpa s"1. The wir3 specimens were electroetched into sharply-pointed FIM specimens. Next the specimens were analyzed chemically by the atom-probe technique at an ambient pressure of ^4x10"^ Torr, with the specimens maintained at 45 K. A pulse fraction (f) of 0.15 was used for all the analyses. The quantity f is the ratio of the pulse voltage to the steady-state dc voltage. A constant pulse frequency of 60 Hz was employed. The average field-evaporation rate — average number of ions evaporated per field-evaporation pulse — was equal to 0.02 ions pulse . Using these experimental conditions we were able to obtain good agreement between the nominal Re concentration, and the Re concentration as determined by the atom-probe technique for unirradiated alloys. These experimental conditions were used in all of our chemical analyses. The specimens were imaged employing 3He as an imaging gas. The reason for.using -%e, rather than He gas, was to minimize the concentration of He present in the atom probe and therefore, to make it possible to identify *He atoms which have had their origin in the neutron- irradiated specimens. The basic mode of displaying the data in the present experiment is in the form of an integral profile. A Re integral profile is obtained by plotting the cumulative number of Re events versus the cumulative number of W plus Re events. The average slope of such a plot corresponds to the average Re concentration of the volume analyzed, since the cumulative number of all events detected is proportional to depth. In analyzing a particular precipitate the slope of an integral profile lower limit to the actual Re concentration in pitate (*), as in most cases the dimensions of the analyzed cylinder are greater than the size of a precipitate. The superscript ppt stands for precipitate, the subscript u on the bracket means an uncorrected value and the superscript * implies a corrected value. The relationship between and * for different possible precipitate morphologies is presented in Appendix A of Herschitz and Seidman. 24

EXPERIMENTAL RESULTS: W-10 AT.% Re

Four radiation-induced precipitates were detected and analyzed, whereas no voids were found in the W-10 at-% Re alloy. The density of the radiation-induced precipitates is equai to ilO an" ; it was determined following the procedure used by Brenner and Seidman. ' The following summarizes the main experimental results: (1) Coherent, semicoherent and possibly incoherent precipitates have been observed. The number density of precipitates is * ^ 52 at.% Re. This result indicates that the WRe precipitates in the W-10 at.% Re alloy are radiation resistant in the temperature range 575 to 675 C — in the presence of a fast-neutron flux. (4) The precipitates were not associated with either linear or planar defects, or with any impurity atoms; i.e. a true homogeneous radiation-induced precipitation occurs in this alloy. (5) No voids were detected in this alloy. This indicates that the addition of 10 at.% Re to W suppresses void formation, as voids have been detected in pure tungsten — which had been subjected to fast-neutron irradiation.

EXPERIMENTAL RESULTS: W-25 AT.% Re

Six precipitates, three voids, a grain boundary and a region immediately adjacent to a grain boundary were chemically analyzed by the atom probe technique. The following summarizes the main experimental results for this alloy: (1) Coherent, semicoherent and incoherent precipitates have been observed. Their n\number density is M.0cm and the mean diameter is ^40 A. (2) The precipitates observed have either a disc shaped or spherical morphology. (3) The composition of the radiation-induced precipitates corresponds to "WRe^; i.e., * is approximately equal to 75 at.% Re. (4) The coherent precipitates (^WRe3) were not associated with either linear or planar defects or with any impurity atoms; i.e., a true homogeneous radiation-induced precipitation occurs in this alloy. (5) The semicoherent and incoherent precipitates were associated with ^He atoms; i.e., heterogeneous precipitation may have occured in this alloy. 25

(6) Voids at a number density of M.0 cm and a mean diameter of 90 A have been detected in this alloy. No significant Re enrichment or depletion at these voids had occured. (7) Formation of a two-dimensional ^WRe3 phase has been observed at a grain boundary.

DISCUSSION

W-10 at.% Re alloy

The fact that the precipitates in the subsaturated alloy are not associated with either structural defects or with any impurity atoms indicates that a true homogeneous radiation-induced precipitation occurs in this alloy. Experimental evidence for homogeneous radiation-induced precipitation has been presented recently by Cauvin and Martin in the case of Al (Zn) alloys, 8'19 by Brager et_ al^.20 in the case of a 316 stainless steel, by Mukai and Mitchell for a Ni (Be) alloy, ^ and Kinoshita and Mitchell and Wahi arid Wollenberger23 for Cu (Be) alloys. Theoretical treatments of this physical phenomenon have been considered by Cauvin and Martin and Maydet and Russell; ° the latter authors only considered the possibility of the nucleation of incoherent precipitates, whereas Cauvin and Martin also considered coherent precipitates . We now describe a possible sequence of plausible events which can result in the homogeneous micleation of WRe (the o phase) precipitates, in a subsaturated alloy which is subject zo irradiation with fast neutrons. ' The primary source of radiation damage, in the case of fast neutrons, is the displacement cascade. Each displacement cascade is created by a primary knock-on atom (PKA) with a mean recoil energy of 4 keV. In the case of pure tungsten it is known from FIM experiments that a displacement cascade, created at 15 K, consists of a vacancy-rich core (^2 to 30 at.%) surrounded by a distribution of self-interstitial atoms (SIAs), which is created by the replacement collision sequence mechanism. The concentration of SIAs on the periphery of a displacement cascade can be as high as 'VL to 3 at.%. Since the radiation damage is highly localized in the displacemenr cascade — the point defect supersaturation in between the displacment cascades is initially negligible — it is probable that the nucleation of a WRe precipitate occurs in its vicinitv The absolute efficiency of this nucleation process is as low as the final density of radiation - induced WRe precipitates is ^lO-Lfc>cm~J> — which is significantly less than the number density of PKAs that produce displacement cascades.^ Employing the known properties of point defects in W and W(Re) alloys it can be demonstrated that plausible first steps in the nucleation of a WRe (a phase) precipitate involve the migration of tungsten SIAs to Re atoms to form mobile mixed dumbbells — in the immediate vicinity of a displacement cascade — which in turn react to form an immobile di-Re cluster.11 The di-Re cluster can then grow by the accretion of mixed dumbbells, and pure tungsten or rhenium SIAs. Specifically, the formation of a WRe cluster is envisaged to occur via the following possible reactions: (a) two mixed dumbbells react to form an immobile di-Re cluster; (b) the di-Re cluster reacts with a pure tungsten SIA to 26

form a WRe2 cluster; and (c) the WRe2 cluster reacts with a second tungsten SIA to form W2Re2 (or 2WRe) cluster. During the course of the two year irradiation the displacement cascades dissolve slowly (see Appendix B of Herschitz and Seidman-'--'-) and they provide the vacancies which can result in the shrinkage of clusters. Recent experiments by Averback and Ehrhart^8 on Ni-1 at.% Si also suggest strongly that point defect clustering and trapping reactions occur in the vicinity of displacement cascades. Further specific details of the growth or shrinkage of a cluster are difficult to state, but they can be rationalized in terms of the Cauvin-Martin model for radiation-induced metastability. The physical basis of the Cauvin-Martin model is that the irreversible vacancy- SIA annihilation reaction drives solute clusters towards a larger solute content and hence to precipitation. A plausible mechanism for the suppression of void swelling, in this alloy, involves the dominance of vacancy: SIA recombination over the destruction of these point defects at a biased sink — the dislocation. This is possible, in particular, by the recombination of vacancies wieh SIAs which are trapped in immobile clusters involving SIAs and rhenium atoms. This strong recombination process prevents the accumulation of a sufficient number of vacancies for the nucleation and growth of voids. This mechanism for the suppression of voids is consistent with the mechanism suggested for the homogeneous nucleation of WRe (o phase) precipitates.

W-25 at.% Re alloy

A very striking observation is the detection of coherent precipitates with the composition WRe3, in the W-25 at.% Re alloy, which are not associated with either structural defects or impurity atoms. The latter observation strongly suggests that they were homogeneously nucleated. The basic problem is to explain the radiation—induced precipitation of WRe3 if the x phase is not in thermal equilibrium with the primary Cβ) solid solution between 575 to 675 C. The x phase may also be nucleated by the homogeneous nucleation mechanism suggested in the previous subsection for the a phase in the vicinity of displacement cascades. The formation of a WRe^ cluster is envisaged to occur via the following reactions: (a) two mixed dumbbells react to form an immobile di-Re cluster; (b) the di-Re cluster reacts

with another mixed dumbbell to form a Re3 cluster; and (c) the Re3 cluster reacts with a pure tungsten SIA to form a WRe3 cluster. Once again the further growth (or shrinkage) of this elementary cluster can be understood to occur via the Cauvin-Martin model**4 for radiation- induced metastability. A necessary condition for the formation of a large number density of WRe3 precipitates is that the nucleation current of these immobile WRe3 clusters be considerably greater than the nucleation current of immobile WRe clusters — in the W-25 at.% Re alloy. In the previous subsection we suggested three point defect reactions which can lead

to the formation of a W2Re2 (or 2WRe! cluster. If it is assumed that these nucleation reactions, as well as the three reactions postulated

above for the nucleation of a WRe3 cluster, are in detailed balance 27

then it is readily shown that t 3 c <* c c WRe3 W-W Ii and 2 2 CW Re " CW-wcii ' where cw_w is the concentration of tungsten SIAs and the constants of proportionality are the products of the rate constants of the individual point defect reactions that lead to WRe3 of fr?2Re2 (2WRe) . The above equations show that c^g should be greater than cw Re . This is because the value of c^ grows more rapidly — in the direct vicinity of a displacement cascade — in the supersaturated W-25 at.% Re alloy than in the W-10 at.% Re alloy; note that cw_w is initially the same in both alloys and is in the range 10~2 to 3*10~2 at. fr. This model for the homogeneous nucleation of WRe3 precipitates, from a supersaturated W-25 at.% Re alloy, shows qualitatively that the nucleation current of immobile WRe3 clusters is greater that of W2Re2 (2WRe) clusters. However, the model does not explain why in the subsaturated W-10 at.% Re alloy the Re concentration of the precipitates stops at ^50 at.% Re. This question stands as an unsolved problem at present. 4 Another very interesting observation is the detection of He atoms inside semicoherent or coherent participates. The detection of 4He atoms in this alloy is at first glance somewhat surprising, as the cross section for 4He production in pure tungsten is quite small. The cross section for the production of Tie atoms on rhenium atoms is not available, however L.R. Greenwood (Argonne National Laboratory, private communication) estimates that the 4He production rates on Re and W should be quite similar. Thus, for a fluence of 8.6 dpa the estimated 4He concentration is "^4x10"3 appm in W-25 at.% Re; this assumes the displacement threshhold energy of a Re atom is identical to that of a W atom — 52 eV.23 This suggests that the 4He atoms were most likely produced on the impurity atoms present in this alloy. The elements B, C, N, O and S all have rather large cross sections for the production of He atoms (L.R. Greenwood, private communication). The absence of He atoms in precipitates in the W-10 at.% Re alloy ^ may simply be a result of a lower level of impurity atoms in this particular alloy.

Herschitz and Seidman 2 discuss four possible mechanisms which can result in the detection of 4He atoms in the WRe3 precipitates. Two of the mechanisms imply that the WRe3 precipitates were heterogeneous ly nucleated and two that they were homogeneously nucleated. There is no obvious way to distinguish among these four mechanisms. Hence, we are left with the distinct possibility that the semi or incoherent WRe2 precipitates were heterogeneously nucleated. In the voids analyzed by the atom probe technique we were unable to detect 4He atoms. The specimen temperature was 45 K during the chemical analyses/ hence, we can rule out the possibility of 4He atoms diffusing out of the voids as they were dissected, since the measured mobility of He atoms at 45 K is extremely small.30"32 The most likely reason we were unable to detect 4He atoms in the voids was simply that the volume fraction of the void analyzed was small — typically much less 28

than 0.1, and since the absolute number of 4He atoms per void is not expected to be very large, the probability of detecting one ^He event is very small. The corrected Re concentration in a grain boundary was determined to be "W5 at.% Re. This value corresponds to WRe3 (x phase). The grain boundary Re concentration was found to fall to the bulk value (25 at.% Re) in *v>4 A. Thus, the x phase forms along a grain boundary and we have an example of a two-dimensional phase, which is the result of a non-equilibrium radiation-induced segregation process. The atomic mechanism for the formation of this phase can be explained in terms of the migration of mixed dumbbells to the grain boundary. This mechanism is consistent with our suggestion that the WRe3 precipitates form in a W-25 at.% Re alloy, subject to a fast-neutron radiation field, as the result of a homogeneous nucleation process which involves mixed dumbbells reacting to form an immobile di-Re cluster.

ACKNOWLEDGEMENTS

This research was supported by the U.S. Department of Energy. Additional support was received from the National Science Foundation through the use of the technical facilities of the Materials Science Center at Cornell University. We wish to thank Mr Robert Whitmarsh for enthusias- tic technical assistance, Dr Alfred Wagner (now at Bell Laboratories) for preparing the specimens for irradiation, Dr Martin Grossbeck (Oak Ridge National Laboratory) for arranging for an irradiation in EBR-II, Dr Robert S Averback (Argonne National Laboratory) and Dr Avner Brokman (Hebrew University) for useful discussions, Dr L R Greenwood (Argonne National Laboratory) for kindly performing calculations for us employing the ENBF/B-V code, and Dr Georges Martin (Centre d1Etudes Nucleaires de Saclay) for useful questions and comments on the more extended version of this manuscript.

a) Presently at R.C.A. Corp., Astroelectronics Division, Princeton, New Jersey 08540. b) Presently on a leave of absence at the Hebrew University of Jerusalem, School of Applied Science and Technology, Bergmann Building, 91904 Jerusalem, Israel.

REFERENCES

1. J. Nucl. Mater. £3_, 1 (1979). 2. J.R. Holland, L.K. Mansur and D.I. Potter, editors, Phase Stability During Irradiation (Metallurgical Society of AIME, Warrendale, Pennysylvania, 1981). 3. F.V. Nolfi, Jr., editor. Phase Transformations During Irradiation (Applied Science, London, 1983). 4. K.C. Russell, Phase sLability Under Irradiation, to appear in Progress in Materials Science (Pergamon Press, Oxford). 5. W.E. Browning, Jr. and C.E. Miller, Jr., in Fourth Symposium on Temperature: Its Measurement and Control in_ Science and Industry (Reinhold Publishing Company, 196 2), Vol. 3, part 2, pp. 271-276. 6. R.L. Shepard and B.H. Montgomery, Oak Ridge National Laboratory Report No. 5108, Oak Ridge Tennessee, November 1976, pp. 418-427. 29

7. J.M. Dickinson and L.S- Richardson, Trans, of the ASM j>^, 758 (1958). 8. M. Hansen and R. Anderko, Constitution of Binary Alloys (McGraw-Hill, New York, 1958), pp. 1153-1154. 9. V.'K. Sikka and J. Moteff, Metall. Trans. 5_, 1514 (1974). 10. R.K. Williams, F.W. Wiffen, J. Bentley, and J.O. Stiegler, Metall. Trans. A 14A, 655 (1983). 11. R. Herschitz and D.N. Seidman, Cornell Materials Science Center Report No. 5014 (1983). To appear in Acta Metallurgica (1984) . 12. R. Herschitz and D.N. Seidman, Cornell Materials Science Center Report No. 5015 (1983). To appear in Acta Metallurgica (1984) . 13. A'. Wagner, T.M. Hall, and D.N. Seidman, Rev. Sci. Instum. 46, 1032 (1975) . 14. T.M. Hall, A. Wagner, A.S. Berger and D.N. Seidman, Scripta Metall. 1(3, 485 (1976) . 15. T.M. Hall, A. Wagner and D.N. Seidman, J. Phys. E: Scient. Instrum. 10_, 884 (1977) . 16. A. Wagner, T.M. Hall and D.N. Seidman, J. Nucl. Mater. 69_ and 70, 413 (1978) . 17. S.S. Brenner and D.N. Seidman, Radiat. Effects 2A_, 73 (1975). 18. R. Cauvin and G. Martin, J. Nucl. Mater. J33_, 67 (1979) . 19. R. Cauvin and G. Martin, Phys. Rev. B 23_, 3333 (1981). 20. H.R. Brager and F.A. Garner, J. Nucl. Mater. T3/ 9 dq78) . 21. T. Mukai and T.E. Mitchell, J. Nucl. Mater. 105, 149 (1982). 22. C. Kinoshita and T.E. Mitchell, Electron. Micros. £, 236 (1980) . 23. R.P. Wahi and H. Wollenberger, J. Nucl. Mater. 113, 207 (1983) . 24. R. Cauvin and G. Martin, Phys. Rev. B 23_, 3322 (1981). 25. S.I. Maydet and K.C. Russell, J. Nucl. Mater. 64_, 101 (1977). 26. L.A. Beavan, R.M. Scanlan and D.N. Seidman, Acta Metall. 19_, 1339 (1971) . 27. C.Y. Wei and D.N. Seidman, Philos. Mag A 43_, 1319 (1981). 28. R.S. Averback and P. Ehrhart, to appear in J. Phys. F: Metal Phys. (1984) . 29. M.I. Current, C.-Y. Wei and D.N. Seidman, Philos. Mag. A 47_, 407 (1983) . 30. A. Wagner and D.N. Seidman, Phys. Rev. Lett. £2_, 515 (1979) . 31. J. Amano, A. Wagner and D.N. Seidman, Philos. Mag. A 44_, 177 and 199 (1981) . 32. J. Amano and D.N. Seidman, Cornell Materials Science Center Report No. 4153 (1982). To appear in J. Appl. Phys. (1984). 30

SOME ASPECTS OF THE CRYSTALLIZATION OF AMORPHOUS COPPER-RICH COPPER-HOLMIUM THIN FILMS

L. Shikhmanter^, M. Talianker^'-1 and M,P. Dariel^b>C')

a) Materials and Process Engineering, Engineering Division, Israel Aircraft Industries Ltd., Lod, Israel. b) Department of Materials Engineering, Ben-Gurion University of the Negev, P.O. Box 653, Beer-Sheva, Israel. c) Nuclear Research Centre-Negev, P.O. Box 9001, Beer-Sheva, Israel.

ABSTRACT Transmission electron microscopy was used to study the composition dependence of the crystallization of amorphous copper-rich Cu-Ho thin films. The phases formed during the crystallization of these films correspond to those predicted by the equilibrium phase diagram. In addition, the crystallization of a hexagonal, phase which so far has not been reported in rare earth-copper systems, takes place. Some aspects of the morphology of the precipitates crystallizing within an amorphous matrix are discussed.

1. INTRODUCTION In recent years amorphous alloys have attracted considerable attention. Amorphous metals not only exhibit technologically promising properties, they also constitute convenient materials for the study of the physical aspects of the amorphous-crystalline phase transformation in solids. In our previous studies we have observed that the crystallization of rare earth (R)-Cu and R-Ag amorphous films of near-equiatomic composition results in the formation of an equilibrium intermetallic compound [1-3]. In order to investigate the composition dependence of the crystallization process we have also studied the crystallization of amorphous copper-rich Cu-Ho thin films.

2. EXPERIMENTAL METHODS

Amorphous thin films, approximately HOoR thick, were prepared by vacuum evaporation at 10"G Torr of an appropriate crystalline alloy onto a carbon film substrate held at room temnerature. The initial Cu-Ho crystalline alloys were prepared by arc-melting mixtures of the holmium and copper that was followed by homogenization anneal in evacuated quartz capsules. The crystallization process was observed in situ during the various annealing treatments by using a JEOL 200B transmission electron microscope equipped with a hot stage attachment. 31

In the course of this study thin amorphous films having the compositions 68 at % Cu-Ho, 74 at % Cu-Ho and 78 at % Cu-Ho were investigated.

3. RESULTS AND DISCUSSION

The crystallization of the amorphous 68 Cu-Ho film begins at 190°C and results in the formation of lense-shaped crystals (Fig.l) of the orthorhombic C£Cu2-type H0Q12 intermetallic compound. As the original crystal grows new, differently oriented HoCu2 crystallites emerge from its external surface, Fig.2a. The new crystallites have a twin-like orientation with respect to the original parent crystal. A schematic representation of this system is shown in Fig.2b. The twin planes are the [103] planes and the [002] planes contain the direction of the fast growth of the HoCu2 crystals in the amorphous matrix.

Fig.l. The structure of a thin 68 at Cu-Ho amorphous film after crystallization at 190°C. The crystallization of amorphous 74Cu-Ho film starts at 260°C and leads to the appearance of the crystalline phase which consists of a large number of parallel thin plates, Fig.3a. The selected area diffrac- "tion (SAD), pattern taken from this phase can be indexed in terms of a hexagonal unit cell with parameters a = 11.45A and c = 16.82A. As far as we know, there exists no previous report on a stable hexagonal phase at room temperature in the 20-25 at % R-Cu concentration range. Franceschi, in his studies of the R-Cu (R = Nd, Gd, Dy) [4r6] systems, suggested on the basis of the observed thermal effects, the existence of a high- temperature R2CU7 phase. According to Franceschi, this phase cannot be retained at room-temperature, its structure therefore had not been determined. Possibly, the hexagonal phase, which we observed in the courss of the crystallization of 74Cu-Ho films corresponds actually to a com- '-ound with the same stoichiometry, namely H02CU7 In contrast to crystallization taking place at 260°C, the isothermal annsal of the amorphous 74Cu-Ho film at 280°C results in the simultaneous crystallization of two different morphologies, denoted A and B, of the crystalline phases as shown in Fig. 3b. The analysis of the SAD-pattern, taken from this sample, revealed that the A-type crystalline phase is identical with the hexagonal phase, obtained during the isothermal anneal at 260°C, and that the B-type phase has a tetragonal structure with parameters identical to the parameters of the Dy2Cu9 compounds [4]. 32

002 103

Fig.2 a - The H0C112 crystal grown in the amorphous matrix with emerging twins; b - The schematic representation of the crystal shown in (a).

Fig.3. The crystalline phases formed in the amorphous 74 at % Cu-Ho film during isothermal crystallization at a) 260°C; b) 280°C. 33

The crystallization of amorphous 78Cu- Ho thin films begins at 240°C and results in the emergence of elongated rhombic-like crystals (Fig.4) having the same tetragonal structure as was determined for crystals formed during the crystallization of 74Cu- Ho film at 280°C. The detailed morphology of the precipitates crystallizing within the amorphous matrix is the result of highly complex interactions involving the volume change associated with the crystallization, the relative values of the elastic constants of the precipitate and the matrix and the values of the various precipitate matrix interfacial energies.

Fig.4. The crystal formed in an amorphous 78 at % Cu-Ho film during isothermal annealing at 240°C. The crystals of RCu and RAg phases with cubic CsC£-type structure, formed during the crystallization of amorphous films, have a circular shape 1-3 . Such a shape was also observed for the crystallization of ot-Fe crystals (bcc type structure) from Feeo(Ci-xBx)2o amorphous alloys [7], Y-FeNi crystals (fcc-type structure) from the Fe^oNiuoBao amorphous alloy [&] and crystals of a Pd-rich fee phase from the Pds3Sii7 amorphous alloy [9]. We may assume that high symmetry cubic crystals possess a relatively low anisotropy of their elastic constant. Neglecting the possible anisotropy of interfacial energy terms in the first approximations, we may define our problem as that of determining the morphology of isotro pic precipitates in an isotropic medium. Elastic theory [1O] predicts that if the precipitate is stiffer than the matrix (higher value of the elastic constants) it will display a spherical shape. Such a spherical shape will transform into a cylindrical one in a thin film, when the dimensions of the precipitate are of the order of the thickness of the film. This indeed was observed for the growth of the equiatomic CsC£- type RCu and RAg compounds. Crystals with a lower than cubic symmetry crystallizing within the amorphous matrix may possess, in principle, a higher degree of anisotropy of their elastic constants. Again neglecting surface energy effects we may expect such crystals to display departures from the spherical morphology exhibited by the cubic-type crystal. This indeed was observed for the crystals appearing in the copper-rich Cu-Ho amorphous matrix. The crystals with orthorhombic, hexagonal and tetragonal structures are formed during crystallization of these amorphous films. All these crystals have an elongate.", rorm (Figs. 1-4). A similar elongated shape of the crystals was also observed for crystallization of orthorhombic Fe3B phase from the Fe7s-B25 amorphous alloy 111], and for the hexagonal Te- Se tshase within amorphous Te-Se films [12], 34

REFERENCES

1. L. Shikhmanter, M. Talianker and M.P. Dariel, Thin Solid Films, 90, 51 (1982). 2. L. Shikhmanter, M. Talianker and M.P. Dariel, J. Phys. Chem. Solids 44_, 745 (1983). 3. L. Shikhmanter, M. Talianker and M.P. Dariel, Bulletin of the IPS, 29_, 80 (1983). 4. E. Franceschi, J. Less Common Met., 87^ 249 (1982). 5. M.M. Carnasciali, G.A. Costa and E.A. Franceschi, J. Less-Common Metals 92^ 97 (1983). 6. M.M. Carnasciali, S. Cirafici and E. Franceschi, J. Less-Common Metals 92^ 143 (1983). 7. K. Shimomura, P.H. Singii and R. Ozaki, J. Mat. Sci. 15_, 1175 (1980). 8. K. Muller, M. Von Heimendahl, J. Mat. Sci. 17_, 2525 (1982). 9. M. Scott, Rapidly Quenched Metals III, (Edited by B. Cantor), Vol.1, 198, Metals Society, London, 1978. 10. D.M. Barnett, J.K. Lee, H.I. Aaronson and K.C. Russel, Scripta Metal1., 8, 1447 (1974). 11. J.S. Vermaak and J. Petruzzello, J. Appl. Phys. 53_, 6809 (1982). 12. M.M. Carnasciali, S. Cirafici and E. Franceschi, J. Less-Common Met. 81, 115 (1981). 35

DENDRITIC GROWTH AND DENDRITE ARM SPACING IN THE SOLIDIFICATION OF STEEL

M. Bamberger & I. Minkoff Department of Materials Engineering Technion, Israel Institute of Technology Haifa, Israel

INTRODUCTION

Dendritic growth is characteristic of materials growing under conditions of thermal or diffusional instability. While non-metallic materials may be observed to grow bounded by low index crystallographic planes, metallic growth is invariably unstable for even small values of undercooling. The crystals grow in a characteristic manner shown in Figure 1. The difference between the case for metals and that for non-metals is in the relatively small values of surface energy for metals between solid and liquid phases. The higher values for non-metals act to stabilise the crystal faces in growth.

Figure 1 shows the characteristic parameters of a metallic dendrite - the tip radius P and the spacing between the dendrite branches X2. This is termed the secondary spacing while the primary spacing denotes the distance between the primary arms of growing dendrites.

A2 is an important parameter in cast metals since it determines the scale of the segregatxon and hence the homogeneity of the material and the degree of separation of the inclusions. The tip radius P is an inverse function of the growth rate R, decreasing with an increase of T?. A number of papers have attempted to correlate X with p. However the most recent analysis (1) tends to separate the two phenomena of radius of the tip and instability of the surface behind the tip. The tip is stable in growth while the branching phenomenon is an instability in growth which operates by its own mechanism in conditions generated along the surface. In addition to the instability mechanism, X2 may be dependent on dissolution mechanisms, subsequent to growth. These are not dealt with here.

A MODEL FOR BRANCHING INSTABILITY

The problem becomes one of determining an appropriate model for the spacing between the branches. One such model proposed by Adams (2) is shown in Fig.2. The arms are taken as cylindrical and grow radially into a liquid characterised by a solute distribution which was determined by experimental methods. Adams determining that there was a constant ratio of solute between that in the liquid at the surface of

the branch Cγ and that in the liquid at the centre between branches Co. He determined X by heat flow calculations.

In the model which we have adopted, the value of CjjCo is also taken as constant but we refer this to a physical requirement of the system which is the optimum spacing of arms for the most rapid dissipation of the 36

undercooling. We suggest that the spacing is such that the initial value of solute concentration at the centre between dendrite arms does not vary by more than 17C from the solute concentration of the bulk liquid, and this occurs at a sharp minimum. If there were no sharp minimum, the arms would be too far spaced for thft maximum dissipation of the undercooling. If the value of solute concentration at the centre varies too much from Co, the arms are too closely spaced and their rate of radial thickening would be reduced. We can then calculate X which is twice the thickness of the boundary solute layer for a ratio of CL/CO taken to be 1.005. We have calculated the values of dendritic growth velocity R for steel cast on copper chills of varying thickness, using a computer programme and then obtained ^2- We have compared these calculated values with values experimentally obtained on cast steel plates. The calculated values were obtained from the following relationship resulting from the model:-

A = . llnn k R l-ko °

The experimentally measured values of X are in good agreement with the predicted values. In the equation, ko is the distribution coefficient (for carbon in y iron) and D is the diffusion coefficient (for carbon in liquid iron).

EVALUATION OF X

The evaluation of X is made as follows:

The distribution of solute ahead of a growing dendrite arm is given by: _ El D CL = Co [1 + £±2 exp ] o

Co is the original concentration of solute in the liquid. We then evaluate the distance L at which C^ = 1.005 Co and suggest that this is the limiting initial concentration between dendrite arms. This was confirmed by experiment.

We then have: D , 0.005 . L = " ¥ ln T^- ' k°

and for ko = 0.45, X2 = 2L = 11 . D/R.

EVALUATION OF R

The solidification rate, or rate of growth, R was determined from the temperature distribution in the cast plates. This was considered as a one-dimensional problem with all the heat being extracted in the direction of the chill. The calculations give at each instant of time, the progress of the solidification front. The influence of the chill was included in the programme by simultaneously calculating the 37

temperature field in the chill itself. Thus the different thicknesses of chill resulted in different values of R.

COMPARISON OF X (CALCULATED) WITH X (EXPERIMENTAL)

Figure: 3 shows the comparison between calculated and experimental values for different plate and chill considerations. This correlation is good for the 75 mm thick cast plate but shows some deviation for the 50 ram plate, it is suggested due to departure from the boundary conditions as the system becomes smaller in dimensions. Figure 4 shows the variation of X with plate and chill thickness.

REFERENCES

1. S.C. Huang, M.E. Glicksman. Acta Met. 29_, 1981, 717. 2. P.K. Rohatgi, CM. Adams. Trans AIME 239, 1967, 1729. A, (mm)

t 50mm thick plate 75mm thick plate 0.35 » 50mm chill 0.30 . x 100mm chill wafer cooled 0.25

0.20

0.15

0.10

0.05 0.02 10 20 30 40 50 10 20 30 40 50 60 70 Fig 4 Distance from chill ( mm ) —» K Spacing vs. Distance from Chill

0.05 J L 0.01 Fig 1 Fig 2 10 20 30 40(mm) 10 20 30 40 50 60(mm) Characteristic Model of Rohatgi and Distance from chill -—» parameters of Adams (2) Fig 3 dendrite p and X2 X Spacing vs. Distance from Chill Calculated Measured 39

TRANSITION FROM FIBROUS TO LAMELLAR MORPHOLOGY IN UNIDIRECTIONAL SOLIDIFIED Ni-W EUTECTIC

S.F Dirnfeld and Y. Zuta Department of Materials Engineering Technion, Israel Institute of Technology Haifa 32000 Israel

INTRODUCTION

Observations of transition from fibrous to lamellar structure in unidirectionally solidified (UDS) eutectics were reported. In a Ni-W eutectic by Kurz and Lux!, in an Au-Cu eutectic by Livingstone2 and in Al-A^Ca eutectic by Street et al.3. This phenomenon was reported as an irregular one2, and partially explained by Kurz and Luxl on the basis of changes in interfacial free energy. Hunt and Chilton^ attributed the transition to certain given percentage of inclusion elements which force the lamellar to change their orientation and thus change the interfacial energy per unit area of interphase boundaries. Their explanation does not show any relationship with growth rate-R. Jackson and Hunt5 gave a mathematical criterion for predicting the preferred structure, the system with a volume fraction of the strengthening phase greater than 32 pet yields a lamellar structure, while at lower volume structure a fibrous structure tends to be generated. The purpose of this paper is to analyze the dependence of the morphology of UDS eutectic alloys as a function of inter-fibrous or inter-lamellar spacing and to explain the transition from fibrous structure to a lamellar one with increasing the solidification rate. Experimental data from the Ni-W eutectic are presented.

EXPERIMENTAL

Specimens of Ni-W eutectic alloy were solidified unidirectionally by a special apparatus" where the rate of solidification - R was controlled by the velocity of the furnace and a temperature gradient - G of 10.5° °C/mm at the solid-liquid interface was created. The as grown composite structures of UDS Ni-W eutectic at relatively slow R (2 to 10 mm/hr) consists of W fibres in a Ni-W solid solution matrix. The strengthening W phase has a volume fraction of about 6 pet. Fig.l shows the transverse section of a UDS specimen growth at 8 mm/hr (G/R = 1.312), the structure consists of parallel fibers having a uniform distribution. By faster R, transition from fibres to lamellar has been observed, the faceted fibres are transformed into lamellae and two kinds of morphology exist simultaneously. At a growth rate R = 16.8 mm/hr (G/R = 0.628). most of the structure are lamellar, Fig.2 shows the transverse microstructure.

ANALYSIS OF MORPHOLOGY

Now the relationship between the volume fraction of both phases and the geometry of either the lamellar or fibrous structures, will be discussed. Let us consider a cube of unit edge dimension (unit volume) within which 40

40 #

Fig.l. Fibrous structure, Lamellar structure transverse section transverse section the second α-phase appears either as fibers or lamellae in the γ-phase matrix, Fig.3. The number of existing fibers in the unit volume is nf = 1/X^f, where Xf is the interfibrous spacing. The number of lamellae should otherwise be n£ = 1/X^, where X& is the interlamellar spacing. The volume of the α-phase in the unit volume cube is equal to the volume frac- 1 tion F of this phase. Dividing it by the specific number of fibers or lamellae, we have the average volume of a fiber or lamella, thus: 2 Vol. of a fiber = F/n£ = FX f (1); Fig.3. Transverse cross-section of Vol. of a lamella = F/n£ = FX& (2). unit volume containing fibers or Therefore: df = 2v^T/TT'Xf (3) lamellae. and d& = FXo (4) Disregarding the fiber end interfaces or the lamella end surfaces, we obtain the specific lamella-matrix boundary area - Sf: = 2/A (5) 2/X - /irF (6) 0 S£ = nf 7 f

It is seen that S« depends only on the interlamellar spacing, and Sf on the interfibrous spacing and the volume fraction F. There exists a certain volume fraction of the α-phase Fe, for which the specific interfacial area for both morphologies is the same. By SJJ, = Sf, we have

(7) e h * and when Xf = X^ the volume fraction of the α-phase will be Fe 0.318 or approximately 32%. 41

Assuming that the crystallographic orientation relationship is the same between the phases in both morphologies, the lamellar and fibrous one, we have the following expression for the energies:

Es = a S^ (8) ; Es = a Sf + A(r) (9) where Es£ and ESf stand for the interfacial energies for the lamellar and fibrous structures respectively, a is the surface tension for a planar interface and A(r) the surface tension due to the curvature of the fiber boundary. The latter would not exist in lamellar interfaces, but in fibrous structure the A(r) term will increase with decrease of the fiber radius. This can be expressed by: OSf 1 E = OS + —i- = SO(1 + ±-) (10)

Sf t r£ 1 rf where r is the fiber radius.

Substituting eq. (5) and (6) for Esj,j Egf respectively, we obtain E. =£L . =2 2a (11) E Af

Substituting the value for r_ from eq. (3) we get:

Eq -£ (13) ; E =2| S£ h Sf Xf \\ The expression for Egf compared with the expression of Eg^ shows that by decreasing the interfibrous spacing Aj (increasing the solidifications

rate-R) increase considerably Egf of the fiber morphology (the second term of eq. (14) is inversely proportional to x£). However, increase of the volume fraction-F of the α-phase will not affect the interfacial energy of the lamellar structure, but will affect it in the fibrous one. During eutectic solidification, separation into two solid phases takes place, each phase with its own concentration of the two components. In order to obtain it, diffusion in the melt must take place to build up the necessary distribution. In this section, the energies needed for it in different morphologies - lamellar or fibrous - will be evaluated. Therefore, we look for some diffusion parameters that will depend on the geometrical features only, such that will show the differences in energy needed in the different structures. We return to the unit cube model of the material, used in the preceding section, and observe the diffusion field around a lamella or a fiber. The diffusion path equals the interlamellar or interfibrous spacing. Let us put the lamells or fiber at the midpoint of such a spacing, which represents the width of the diffusion field. In the case of a fibrous structure the width of the diffusion field is the same in both orthogonal directions around a fiber having a square form. In the lamellar case the diffusion field length in the direction normal to the spacing was taken so that the area of the field equals that of the fibrous case, see Figs. 4a and b. In the lamellar case, since the length normal to the spacing direction was A|/\£, the area of the a phase is {\y\i)A%. In the fibrous case 42

the area of the a- phase is irdf/4. Sub- stituting the values of d£ and df from *1 r - 1 equations (3) and (4), we get the following: fa.r 2 %; -r^=FXf (15)

Since the areas of \\V the diffusion-field configurations are equal in both cases, that of the matrix phase-y is also the same. In the lamellar case the diffusion took place only in Fig.4. The diffusion field a) around a lamella the spacing-width b) around a fiber, c) equivalent around direction (in the a fiber. melt in front of the solidification interface) so that the representative geometrical parameters of the process are the lamellar half thickness d£/2 and the diffusion field half dimension \%/2 (see Fig.4a). The situation around a fiber is different, and the diffusion- field configuration is more conveniently treated in terms of a circle, with the same diffusion distances around the fiber in each direction, provided this circular configuration has the same area as the square. Thus, the diameter of the circular diffusion-field will become 2Xf//if, and the representative geometrical parameters will be df/2 - the fiber half-diameter, and Xf/i/iF - the diffusion field dimension. Since the amount of both phases in the equivalent diffusion-field for both morphologies is the same, it is hypothesized that the total energy needed for the diffusion is proportional to the average diffusion distance for both components A and B for build-up of the concentration distributions in the solid state. A mechanistic approach is suggested, whereby the average diffusion distances for the components are X^ and XJJ. These distances are between the "centers of gravity" for the concentration fonfigurations (within the diffusion-field element) for the melt and for the two-phase solid states. This approach accounts for the concentration differences as well as for the geometrical configurations. In order to calculate the "center of gravity" points, the static rule of mements will be considered. It states that the distance of that point of a given area from another chosen arbitrary point equals the sum of moments of the partial areas about that same chosen point, divided by the sum of the partial areas. Let the following designations be used: I/j. and L2 are the width of the phases a and y respectively; C^ and C2 - weight concentrations of component A in phases a and y respectively and (1-Ci), (I-C2) - the same for component B in both phases (Fig.5). Thus, be the rule of moments, one gets: 43

L +L L C +L C (L +L ) X l 2 l l/2 2 2 l 2/2 Cβ A = L C +L C 1 —— 2 2 1 1 and

1-C, - Ll(1-CJ/2+L2(1-C2)(Ll+L2/2) Ll+L2 "TJ T f "i p^ _LT ("\ C \ O Co 2 2 11 (17) for the lamellar structure Li=d£/2 and T L2=X -d! /2 see Fig.5. If we use the d£ 1-C, T 9 ; value from eq.(4), then Li=FX£/2 and L2=(l-F)X£/2, therefore we get the average diffusion distances: -L,- (a) (y> (a) (y) (18) Fig.5. Concentration profile of components A and B. and

X£ (F -F)(C2-C1) (19)

However, for the fibrous structure L =dJ2 and L =Xf//rT - j— ; using df from eq.(3), one gets:

(20) = Xf/F (!-• = V^TTT • X, therefore:

(21) ; x = (22) XAf = Bf 2/rr 2/F

Now, assuming that the total energy for diffusion within a diffusion field element is proportional to the average diffusion distance, the analysis for diffusion in both morphologies are given by:

(23) EDf -

Where: PA and Pg are some energy functions incorporating the diffusion coefficients and concentration parameters for components A and B, and n is the number of diffusion field elements within the actual cross section of the specimen; E^ and EDj are the total required energies for diffusion for the lamellar and fibrous morphologies respectively. The source of these energies in the system is the amount of supercooling below the solidification temperature^. It should be recalled that the system itself will choose the morphology, lamellar or fibrous, which needs the lower amount of supercooling. Let us denote E£ for the energy needed to obtain lamellar structure and E£ Tor the fibrous one:

= E (25) ; E = + E,. (26) si f c 44

Substituting eqs. (13) and (14), we get: A (F2-F)(C -C ) X (F*-F)(C -C )

X. (F-/F)(C_-C ) A (F-/F)(C -C ) E 20 /iJ + 2TO+ n{ j_E 2_J^ + p r_J L_^ £ A A B ff XX^^ 2/ir v¥(C221-C1)+C2 2 2/ir y^^

(28) CONCLUSIONS AND DISCUSSIONS

From the above expressions (27) and (28) the following conclusions can be drawn i) Increase of the interface spacing in both morphologies reduces the interfacial surface energies. ii) The fibrous interfacial surface energy, compared with the lamellar one, increases considerably as the interfibrous spacing X_ increases. iii) Decrease of the interphase spacing in both morphologies reduces the average diffusion distances, and thereby the energy needed for diffusion.

Following these conclusions, the transition from one morphology to another by changing the rate of solidification can be explained as follows: When an eutectic system solidifies, the rate of solidification dictates the maximum available average diffusion distance. A lower rate will increase the average diffusion distances and vice versa. Eutectic systems will seek to reduce the surface energies in both morphologies. In the fibrous structure the surface energy is reduced also by increasing the curvature (eq.lO), since the number of fibers decreases with increase of the interface spacing and therefore their radii increase for a constant volume fraction. To sum up, increase of the interphase spacing, A. or A-, reduces the surface energy terms in the total expressions; on the other hand, it increases the energy needed for diffusion, and the two increases may cancel each other out, in which case decrease of the solidification rate would not increase the interface spacings. This accoun ~ for the upper bound on these spacings (and thereby on the thickness of the lamellae or fibers) existing in different eutectic systems. With a given rate of solidification, which dictates the maximum average diffusion distances, the structure calling for lower energy expansion will be preferable. Increase or decrease of the solidification rate may bring about a situation favoring the second morphology, and transition to that structure will then take place.

It seems that at relatively low solidification rates the total energy for the fibrous structure is lower than that for the lamellar structure. But when the solidification rate increases the dimension of the fibres and X decrease and the curvature component of the surface tension increases until the total energy needed for the fibrous structure exceeds that for the lamellar, and at that point the latter is favored. 45

This finding is supported by the micrographs (R=14 mm/hr) which show a structure composed of both fibers and lamellar, since at that solidification rate the total energy is the same for both morphologies. The fibers have square or rectangular cross-sections; the lamellae exceed the fibers through joining of several fibers end to end. Certain lamellae are seen to havt an "appendage" at the end, as though a fiber is attached there (Fig.2). The model shown in this section is applicable to other systems, and may account for transition phenomena from a fibrous to a lamellar structure reported as unexplained irregular effects by Livingstone^ in the Au-Cu system and by Street^ in the Al-Al^ Ca system.

REFERENCES

1. W. Kurz and B. Lux, Met. Trans, l^ (1971), p.329. 2. J.D. Livingstone, J. Appl. Phys. 41 (1970), p.192. 3. K.N. Street, F.C. St. John and G. Piatti, J. Inst. Met., f5 (1967) p.326. 4. J.D. Hunt and J.P. Chilton, J. Inst, of Metals, 91^ (1963), p.338. 5. K.A. Jackson and J.D. Hunt, Trans. AIME, 236 (1966), p.1129. 6. S.F. Dirnfeld and D. Schwam, Proc. 2nd Ristf International Conf. on Metallurgy and Materials Science, Roskilde, Denmark, 1981, p.225. 46

THE MICROSTRUCTURE OF RAPIDLY SOLIDIFIED Fe-Cu ALLOYS

A. Munitz and Z. Livne

Nuclear Research Center-Negev P.O.B. 9001, Beer-Sheva, Israel.

ABSTRACT rhe impact of cooling rate on the microstructure of Fe-Cu alloys was investigated. A variety of solidification techniques v;ere employed,in order to achieve cooling rates of several order of magnitude variation. Our experimental results Teveal an extension of copper solid solubility in y iron. The cooling rates have a drastic influence on the microstructuTe. \t cooling rates of about 10* °C/s a dendritic solidification mode could be observed, while at high cooling rates the iron has g. circular appearance embeded in a copper rich matrix. The similarity of the microstructure obtained by electron beam surface melting to that obtained in molten pockets by explosive bonding of Fe-Cu plates, suggests a similarity in the cooling rates (i.e. 105 °C/s) (1).

INTRODUCTION Explosive bonding of metals is a method for producing large area bonding between metal plates, particularly when the metal mechnical properties differ grossly. A well known example is the bonding between large plates of lead and steel, which can not be achieved by any other means. A typical industrial application of explosive bonding is the welding of pipelines for heat exchangers. In such systems the pipeline has to be corrosive resisting inside, and highly thermally conducting for the rest of its thickness. Steel - Cu explosive-bonded materials are used for protection walls of safes. The outer layer is a very hard thin steel layer aimed to prevent mechanical penetration. The inner layer consists of a thick copper plate for heat absorption aimed to prevent penetration with a torch. Explosive bonding involves very fast solidification rates [1). In our present work we report a metallurgical study of fast solidification rates in Fe-Cu alloys involving explosive bonding, as well as electron beam surface melting and melt quenching on a wateT-cooled copper plate. With these three techniques it was possible to scan over a large range of solidification rates.

EXPERIMENTAL In the present work we have studied specimens of Fe-Cu alloys obtained by the following methods: A) Resolidification of melt on a water-cooled copper plate: about 10 gr of Fe 50 w/o Cu bulks were melted in a ceramic crucible in an induction furance. After melting, the material was poured over a water-cooled copper plate. The cooling rates evaluated on the basis of the cast material thickness (2) range between 10 and 100 °C/s. B) Electron beam surface melting and subsequent solidification by self quenching. The melting conditions were as follows: driving voltage 60 kV, current 47

between 5 and 10 mA, electron beam scanning velocity about 500 cm/min. The electron beam focus was about 1 mm beneath the molten surface. Cooling rates as high as 105 't/s were achieved (3,4^. C) Resolidification of molten pockets in an explosive bonded Fe-Cu plates. A 5 mm thick copper plate of dimensions 200 mm x 300 mm was used as the flyer plate, and a 10 mm thick iron plate of the same dimensions was used as a parent plate. The flyer plate was placed 1.5 mm parallel to and above the parent plate, the latter was lying on the ground. The explosive material used T=ras TNT+NH4N03 giving rise to a 2200 m/s detonation velocity. Selected areas from the above mentioned three solidification techniques were cut with a slow diamond wheel. The surface was then mechanically polished with diamond pastes up to 1/4 vm, then etched v-th 2gr FeCl3+10 cc HC1 in 95 cc ethanol (Cu etchant) for 10 s. The specimens were examined with a scanning electrom microscope (Philips SEM 505) with EDAX options.

RESULTS AND DISCUSSION In Fig. 1 we show secondary electron images (SEMs) of Fe - 50 w/o Cu quenched on a water cooled copper plate. The figure reveals a dendritic solidification structure, consisting of primary arms of about 100 pm long and 10 vim wide. EDAX analysis showed, that between 9 to 13 w/o Cu is miscible in the dendrite arm (see Fig. 1). Observation of Fe-Cu equilibrium phase diagram (5) indicates an upper limit of about 8% miscibility of copper in y-Fe phase. In our case first to solidify are primary y-Fe dendrites, which dissolve up to 13 w/o copper. The relatively high cooling rates involved (10-100 °C/s) allow this enhanced miscibility of Cu in y-Fe phase. The y-Fe phase is unstable at room temperature. It transform into a-Fe phase. Indeed no y-Fe phase could be detected by x-ray diffraction. All the sample volume contained a-Fe and e-Cu phases solely. Cu is immisible in a-Fe. It then precipitates in the dendrite arm as the primary y-Fe crystals transform into the a form. Upon etching with nital, a fine microstructure is revealed in the primary y-Fe crystals (see Fig. IB). The dendrite arms are embeded in a copper-rich matrix (68 w/o Cu). It consists of a primary e-Cu, as well as a-Fe secondary phases. Precipitated copper is expected to exist within the a-Fe, but this detail could not be seen in the present resolusion. in Fig. 2 we show a secondary electron image of a resolidified surface melted by an electron beam. The copper etching of the specimens reveals two types of structure: A) Big particles (>10 pm) with flat surfaces, within which fine rounded second phase precipitants may be observed. Elemental x-ray micro-analysis of these particles made by the scanning electron microscope showed, that about 30 to 35 w/o Cu is present. On the other hand, when x-ray diffraction was performed on electron beam treated surfaces, only a-Fe phase could be seen. We atribute the high copper concentration in these particles to the impact of the high cooling rate on the Cu miscibility in y-Fe. In this case, cooling rates as high as 105 °C/s are effected, which in turn enhance the Cu miscibility up to 35 w/o. Presumably, the big particles first solidify as primary y-Fe crystals. They are then vigorously swept by the convections in the molten pool. When solidification continues the y phase transforms into an ct-Fe phase, and this process is accompanied by Cu segregation. Compared to the previously described solidification of Fe-Cu alloy quenched on copper (Fig. 1), only short cooling times are involved. Therefore, a different type of Cu precipitation could be observed; (B) Small tiny spheres with diameters 48 between 0.2 vm and 2 ym embeded in a copper-rich matrix. In Fig. 3 we show a secondary electron image of Fe-Cu alloy resolidified as molten pockets by explosive bonding of Fe-Cu plates. Upon proper Cu etching the sample reveals two types of structure: (I) Big flat particles, with surfaces which appear smooth even at high magnifictions. X-ray microanalysis of this structure showed only 0.9 w/o Cu. These results suggest that the particles did not melt al all. Apparently, the particles were torn away from the iron plate by the large melt convection, a result of the large turbulent flow induced by the high pressure involved in the process. They did not have sufficient time to melt; (II) Tiny round spheres, about 0.1 vm diameter, embeded in a Cu rich phase, which was disolved by the etchant. The presence of tiny spheres attached to the big flat particle indicates, that the particle envelope has melted, and resolidified. The similarity between the microstructure of Figs. 2 and 3 suggests, that cooling rates as high as 10s °C/s could be obtained during resolidification of molten pockets, consistent with theoretical predictions (1).

SUMMARIZING REMARKS The microstructure of Fe-Cu alloys were investigated. It was found that the major impact of high cooling rates is to enhance copper miscibility in Fe. The primary Y-Fe transforms into a-Fe, and induces copper precipitation. The character of the a-Fe, which transforms from y-Fe particles is different in the different solidification processes. Microstructural similarities between electron beam molten surfaces and molten pockets in explosive bonded plates suggest, that cooling rates as high as 105 °C/s are obtained in both cases.

ACKNOWLEDGMENT The authors wish to thanks Mr. C. Cotler and Z. Barkai for their technical aid, and to Dr. Z. Burshtein for his editing remarks. Finally, the authors thanks Prof. B. Weiss for his helpful discussions.

REFERENCES 1) Z. Livne, Structure phenomena in the bond zone of explosively bonded plates, Nuclear Research Center-Negev, NRCN 464, 1979. 2) F.A. Joly and R. Mehrabian, J. Mater. Sci. 9_ , 1446 (1974]. 3) S. Kou, S.C. Hsu, and R. Mehrabian. Metall. Trans. 12B , 33 (1981). 4) R. Mehrabian. Int. Metals Rev. 27^ , 185 (1982). 5) P.M. Hansen, Constitution of Binary Alloys, 2nd Ed. (McGraw Hill, 1958). 49

Secondary electron images (SEMs) of Fe-50 w/o Cu qunenched on water cooled copper plate. (A) General view. A typical x-ray spectrum of a dendrite arm is given in the right side; (B) A cross-section of dendrite arm after etching with nital „ 50

Fig. 2 : Secondary electron image of an electron beam molten surface. An x-ray spectrum of the big particle is given on the right side.

Fig. 3 : Secondary electron image of resolidified molten pocket in explosive bonded Fe-Cu plates. An x-ray spectrum of particle (type I) is given on the right side. 51

PHASE STABILITY AND MASSIVE TRANSFORMATION IN Y 0 -COMPLETELY STABILIZED ZIRCONIA (Y-CSZ)+

A.H. Heuer, R. Chaim* and M. Ruhle** Dept. of Metallurgy and Materials Eng., Case Western Reserve Univ., Cleveland, U.S.A.

INTRODUCTION Zirconia (ZrO2) alloyed with Y2O3 stabilizer has recently become an important material due to the high strength and fracture toughness relative to other stabilized oxide-zirconia systems. The improvement of these mechanical properties is directly connected to the volume fractions of ZrO2 polymorphs, namely cubic (c), tetragonal (t) and monoclinic (m) phases, within the material microstructure. The presence of high concentrations of oxygen vacancies needed for electrical charge neutrality and the sluggish diffusion of the stabilizer cation causes instability of the high-temperature polymorphs (c.,_t) under certain conditions. Thus different phase-diagrams for the _t + c_ two-phase field have been reported (1). EXPERIMENTAL As-received mixed powders of ZrO2 - 12 wt% Y2O3 were pressed and sintered at 1600°C for 2 hr. Small specimens were subsequently heat-treated as follows: 1550°C/l hr; 1400°C/1,5 hr, 1 week; 1250°C/20 hr. These were checked by X-ray diffractometry using Cu-Ko. radiation and later prepared for transmission electron microscopy (TEM), by mechanical and ion thinning followed by carbon coating. The alectron microscope (JEOL 200-CX) was operated at 200 KV. RESULTS AND DISCUSSION TEM observation of the sintered specimens revealed c^ grains with a tweed structure containing small _t precipitates. Careful observation of these grains in very thin areas showed fine modulations in two directions (Fig.l) due to the strain field around the coherent t_ precipitates. Selected area diffraction patterns (SADP) from such £ grains for [lll]c zone-axes showed the 3 variants of [112]t reflections (Fig.2) indicating the presence of conventional primitive t_ structure. A completely different microstructure was observed for 1550°C heat-treatment corresponding to new peaks in the X-r:-iy profile (Fig.3), the c_ grains have converted to twinned and acicular crystals. The internal microstructure of these crystals was characterized by curved antiphase boundaries (APB). However observation of appropriate diffraction patterns (Fig.5) showed the presence of only 2 variants of [112]t reflections indicating a different t_ symmetry, namely body centered tetragonal (BCT), t_'. This phase was reported recently (2) as a product of quenching 8 wt% Y2O3 PSZ in the skull melting process. The same microstructure was found for heat- treatments at 1400°C for 1 and 5 hrs. where several tweed c^ grains were

+ To be published in Advances in Ceramics, 1983 * Dept. of Materials Engineering, Technion, Haifa, Israel. ** Max-Planck Institut fur MetalIforschung, Stuttgart 1, F.R.G. 52

also observed. EXAX microanalysis showed an 13 wt% Y2O3 content in c^ grains, where 12 wt% Y2O3 was found for the _t' phase. The rapid phase transition of c^ to _t' at constant composition suggests a massive transformation. According to the ZrO2~Y2O3 phase diagram (3), at the above temperatures a drastic composition change from 12 wt% to 3-4 wt% Y2O3 should occur to form the conventional _t precipitates. Thus the new _t* phase could form without composition change (a displacive transformation) at short annealing times, but must necessarily change crystal symmetry. Indeed this metastable form, which was stable relative to the t_ -> m transformation, disappeared with further annealing for 120 hrs. In addition, a new microstructure of conventional _t phase then appeared, containing lens-shaped to colonies which were internally twinned (Fig.6). This t_ phase was not stable relative to the t_ •> m transformation. At 1250°C the same tweed c^ structure was found, confirming the sluggish nature of diffusion process in this material. ACKNOWLEDGEMENTS R. Chaim acknowledges the MPI association support in liis stay in Stuttgart. REFERENCES 1. A.H. Heuer, N. Claussen and M. Ruhle, Advances in Ceramics, 1983, to be published. 2. V. Lanteri and A.H. Heuer, Advances in Ceramics, 1983, to be published. 3. H.G. Scott, J. Mater. Sci., 10, 1527-1535 (1975).

Fig. 1: £ grain with tweed structure in sintered material. 53

Fig. 2: [111] zone-axis diffraction from c_ grains.

Fig. 3: {400} peaks for £,_t»£' phases 54

Fig. 4: APB's in the acicular _t* crystals.

Fig. 5: [111] zone-axis diffraction from t1 phase.

Fig. 6: _t lens-shaped colonies, internally twinned. 55

MICROSTRUCTURE EVOLUTION AND ORDERING IN COMMERCIAL MGO-PARTIALLY STABILIZED ZIRCONIA (MG-PSZ)*

R. Chaim and D.G. Brandon Department of Materials Engineering Technion-Israel Institute of Technology, Haifa, Israel

INTRODUCTION The polymorphic nature of ZrO2 on cooling from cubic (c_) to tetragonal (_t) and finally monoclinic (m) structures is now well understood. For technological applications ZrO2 is usually "alloyed" with the cubic oxides, for example MgO, CaO, Y2O3, to stabilize the high temperature polymorphs (£,_t) at ambient temperature. The jt -*• m stress induced martensitic transformation is accompanied by a 3% volume increase, and is responsible for the fracture toughness improvement. Thus the mechanical properties of PSZ are sensitively dependent on the material microstructure, a fact which explains the comprehensive microstructural studies (1-6). EXPERIMENTAL Sintered commercial 9 mole % MgO-ZrO2 (PSZ) specimens were heat-treated under the following conditions: 900°C/20 hr, 1000°C/5 hr, 1100°C/l hr, 1400°C/120 hr. The specimens for transmission electron microscopy (TEM) were prepared by mechanical and ion thinning methods, followed by carbon coating. The electron microscope (JEOL CX-100) was operated at 100 KV.

RESULTS AND DISCUSSION The as-received material was composed of c grains with small ellipsoidal shaped _t precipitates, dispersed homogeneously in the £ matrix (Fig.I), Other £ grains contained larger lenticular shaped _t precipitates (Fig.2), which were transformed to the m symmetry and characterized by internal twinning. Large m twins were observed as a grain boundary phase (Fig.3), with a glassy phase at grain boundary triple points. Selected area diffraction patterns were used to identify the different phases mentioned above. The £ grains containing small t_ precipitates have related diffuse scattering intensity (DSI) in diffraction patterns, which has proved to be associated with the £ matrix (7). The microstructure after 1400°C annealing was changed completely to £ grains containing m particles in an eutectoid microstructure with an MgO-rich phase. No changes in the volume fraction of the different phases was observed after annealing at 900-1100°C, except for appreciable strain-field contrast around the small £ precipitates. Selected area diffraction from such areas revealed an ordering phenomena characterized by strong superlattice reflections (Fig.4). The ordered phase was identified as β-phase (Mg2ZRsOi2) with rhombohedral structure, located in the interspace between the small _t particles (Fig.5).

To be published in the Jour. Mater. Sci. 56

This rhombohedral structure is a distorted version of the c^ fluorite structure, with strings of oxygen vacancies located along a specific

<111>C direction. The β-phase has been reported (8) as a high temperature phase formed by quenching from 1850°C. Thus the formation of this phase at lower temperatures has been considered a result of non-equilibrium.

The DSI in PSZ was related previously (7) to short-range ordering of oxygen vacancies, which produced trigonal relaxation of the c^ fluorite structure. The slow diffusion of stabilizer cation enables only short- range diffusion, which causes local enrichment of c_ regions between small _t particles. The appropriate high saturation concentration enables the β-phase formation (28.6 mole % MgO), which was assisted by the presence of ordered oxygen vacancies. This microstructure is metastable and for long annealing times dissociation should occur (Fig.6), to reach bulk equilibrium. The formation of β-phase enhances the stabilization of _t particles by retarding further t_ growth, but destabilizes it by increasing the coherency strain field at the t-0 boundaries.

ACKNOWLEDGEMENTS

This work was in part performed at MPI fiir Metallforschung, Stuttgart.

REFERENCES 1. G.K. Bansal and A.H. Heuer, J. Amer. Ceram. Foe, j>B (5-6), 235-238 (1975).

2. R.H.J. Hannink and R.C. Garvie, J. Mater. Sci., Y]_, 2637-2643 (1982).

3. D.L. Porter and A.H. Heuer, J. Amer. Ceram. Soc, 62^, (5-6), 298-305, (1979).

4. L.H. Schoelien, Ph.D. Thesis, Case Western Reserve Univ., Cleveland, Ohio, USA, 1981.

5. R.H.J. Hannink, J. Mater. Sci., 18, 547-470 (1983).

6. R.H.J. Hannink, J. Mater. Sci., l^, 2487-2496 (1978).

7. R. Chaim and D.G. Brandon, Advances in Ceramics, 1983, to be published.

8. 0. Yovanovitch and C. Delamarre, Mat. Res. Bull., 11, 1005-1010 (1976). 57

Jig. 1; Ellipsoidal _t precipitates in a £

Fig. 2: Lenticular m phase in a £ matrix.

Fig. 3: Large m twins at grain-boundaries. 58

Fie. 4: [lll]c zone-axis diffraction from ordered β-phase regions.

Fig. 5; D.F. image showing β-phase between £ particles.

Fig. 6; Dissolution of the ordered β-regions. 59

DEFORMATION INDUCED DECOMPOSITION OH URANIUM-TITANIUM MARTENSITE

* * * ** G. Kiiranel , J. Sariel , A. Landau and M. Talianker

* Nuclear Research Center Negev ** Ben-Gurion University of the Negev.

Uranium and titanium form a solid solution at high temperatures at all concentrations. The crystal structure of this solution is A2 tungsten type as that of the high temperature phases of the pure elements (y uramium or 3 titanium). The uranium-rich alloys (below 25 at % titanium) decompose into a uranium and S-^Ti when cooling from the y solid solution, but rapid cooking of y phase by water quenching results in a metastable a' phase, which is a supersaturated solution of titanium in a uranium. The first reaction y -*• a+S is an eutectoidic decomposition and the second is a martensitic transformation γ-^a1.

The a1 martensite may transform into a+S during heat treatment above 300 C.

The S phase has an A1B2 crystal structure, which is similar to the to phase of zirconium or titanium alloys. The formation of u phase from the bcc structure (H-T phases of zirconium and titanium) is v.'ell established and its analogy with the formation of 6 from y uranium has been indicated (1) .

The discovery of the diffusionless transition a1 ->• m in zirconium under high hydrostatic pressure (2) led us to examine the stability of a1 martensite under stresses. We found that plastic deformation of a1 is always accompanied by the a' -> a transition.

Geometric analysis shows that the transition a1 -»• 6 is similar to a1 -*• to in zirconium, so that the diffusionless transformation a1 *- a+S takes probably place during plastic deformation of (U-Ti) martensite.

EXPERIMENTAL Uranium 5% Ti was cast into graphite molds at high vacuum. The cast was annealed for 24 hours at 850°C after hot working in order to get an homogeneous titanium distribution. The martensitic structure was obtained by resolution of the titanium at 800°C and water quenching. Tensile specimens were machined and the tensile test showed elongation of 25 ± 5%. The specimens, which have been broken in the tensile test were cut for XRD studies. For comparison, several specimens were subjected to various heat treatments leading to different structures.

RESULTS AND DISCUSSION The diffraction spectra giving the different lattice unit sizes, which obtained according to the heat treatment given,are: a) Slow furance cooling from y phase, giving a mixture of a uranium and

6-IJ2Ti (Figure 1-1). 60

b) Air cooling from y, giving a mixture of a uranium which dissolves titanium (but unsaturated) with S-U-Ti (Figure l-II). c) Water quenching giving an a1 martensite, which is suppersaturated solid solution of titanium in uranium (Figure l-III).

The lattice parameters calculated by least-squares routine are indicated in the figure. The diffraction spectrum obtained from the deformed martensite (at the neck of the tensile specimen) is similar to that of figure l-II,but without the 6-phase peaks. It indicates that a transition a' > a occurred during the deformation. In spite of the absence of the 6-phase peaks, the excessed titanium is probably precipitated as 6-phase, giving the complete reaction af -+ a+tS, since the a phase dissolves no more than 3 at% titanium. As the reaction occurs rapidly (during the tensile test) and at room temperature, it is assumed that it is diffusionless.

A similar transition was found in zirconium. The diffusionless transition a1 -*• a) in zirconium was explained by a model showing the crystal lographic relationships between the zirconium a' and u phases-

The similarity of a' uranium to a' zirconium and 6-U2Ti to u phase leads to a similar model. Figure 2 shows the crystallographic orientation relationship between a'-uranium and 6-U2Ti:

(loo) [oio]a II(oooi) [iioo]6 The lattice distortions between uranium and lUTi are 0.46%, 4.7% and 2.7% along [l00]a , [0lQ]a and [00l]a respectively. The missfit is in the range which plastic deformation can accomodate. Moreover, a comparison of the atomic arrangements in the twolattices reveals that the atomic shuffles G on the (100)a plane to form (1100)6 plane are small - 0.13 A°and 0.2 A alonl g[OOl [001]]^ and [01l]a directions respectively, but shifts of 0.7 A° must occur in the directions of ±[l00]a to form two layers of uranium and titanium in U»Ti.

REFERENCES 1. R.D. Tomlinson, J.M. Silcock and J. Burke, J. Inst. Metals, 1970, 98, 154.

2. A. Rabinkin, M. Talianker and 0. Botstein, Acta Met. 1981, 29, 691. 61

Figure 1: Diffraction spectra obtained according to the heat treatment given. o

1 2

o o

Figure 2 : Atomic positions on (100)(J Cleft), compared with (OOODy Ti (right).

t i 63

APPLICATIONS OF ANALYTICAL ELECTRON MICROSCOPY TO MATERIALS

J. I. Goldstein, D. B. Williams and M. R. Notis

Department of Metallurgy and Materials Engineering Lehigh University, Bethlehem, PA 18015 USA

The analytical electron microscope (AEM) incorporates an energy dispersive spectrometer (EDS) to detect x-rays from an electron-transparent foil. The combination of a thin foil and a focused high energy electron beam permits a nominal x-ray spatial resolution of ~ 10-50 nm. This resolution is two orders of magnitude better than that obtained with bulk specimens in the scanning electron microscope (SEM) . In addition, quantification of the x-ray data is relatively straightforward using the ratio method, where the concentration ratio of elements A and B in the thin foil is directly proportional to the x-ray intensity ratio: (1) CA/CB " AB

1. Study of Precipitate-Free Zones in Al-Ag (1) The occurrence of precipitate-free zones (PFZ) around grain boundaries in aged Al-based alloys has been attributed to solute depletion due to grain boundary precipitation, or vacancy depletion. In Al-Ag alloys two distinct zones free o£ the metastable Y' precipitate are observed (Fig. 1): a wide 'grey' PFZ (GPFZ), and a narrow ( ~ 500 nm wide) 'white' PFZ (WPFZ) immediately adjacent to the grain boundary. A typical TEM microstructure of an aged alloy is shown in Fig. 1,

DISTANCE FROM GRAIN BOUNDARY x 100 t&HCt PROM C«*!N SOUNQARV «2Q0»* Fig. 1: A typical micro- Fig. 2: (a) AEM concentration profile across structure of a Al-16 wt% a grain boundary WPFZ for a sample aged at Ag sample aged at 433 K 433 K for 50 hrs. (b) AEM concentration for 50 hrs. profile across a GPFZ in the grain boundary region. 64 exhibiting a well-defined WPFZ and GPFZ. Solute concentration profiles were measured using an electron-beam diameter of 5 nm. X-ray data were analyzed using equation (1). The results (Fig. 2) indicate that the WPFZ is caused by marked solute depletion (Fig. 2a) and the GPFZ by vacancy depletion (Fig. 2b). The solute content of the GPFZ is equal to that of the bulk. Therefore both solute depletion and vacancy depletion mechanisms explain the formation of PFZ in this system.

2. Early Stage Growth of the Ni Al Intermediate Phase in Ni-NiAl Dif- fusion Couples (2) The major experimental techniques used during the past decade to study the kinetics of interdiffusion and intermediate phase growth have been the optical microscope and the electron microprobe. Because both of these techniques have spatial resolution limits of~1 \M m, it has not been possible to examine early growth kinetics. For example there is little or no information available on the kinetics and/or morphology of Ni Al layer growth for times < 3 hrs. Figure 3 is a typical microstructure showing the growth of Ni Al in a Ni-NiAl diffusion couple produced at 1100°C. The protrusions in the Ni,Al phase are always associated with grain boundaries in the Ni Al layers since grain boundary diffusion significantly contributes to Ni Al layer growth below 1100°C.

Fig. 3: TEM micrograph of an Ni_Al layer protrusion associated with a single Ni_Al grain boundary.

AEM profiles obtained across the diffusion interfaces are shown in Fig. 4. To obtain the upper data set in Fig. 4 the data were corrected

Fig. 4: AEM profiles for a specimen held at 1100°C for 15 minutes.

DISTANCE (/*«) for x-ray absorption effects (3) . The equilibrium concentrations at the appropriate two-phase boundaries at 1100°C are shown in Fig. 4. The data obtained at the interfaces appears consistent with the interface concentrations predicted from phase equilibria. It thus appears that interfacial equilibrium is established even at very short times.

3. Chemical Identification a: Submicron Particles in Steel Weld Metal(4) The complexity of the variables that affect the welding process make it difficult to determine the relationship between the microstructure and the mechanical properties of welds. Interpretation of the significance of microstructural differences has been limited by the lack of localized chemical information. However AEM, in combination with optical and scanning microscopy can overcome some of these problems. The identification of submicron particles in pressure vessel steel weldments is of particular interest because of their effect on mechanical properties. Fig. 5a shows a TEM image of precipitates in such a weldment. The EDS spectrum of the precipitates, Figure 5b, taken in the AEM shows large amounts of Si and Mn indicating that the precipitates are silicates, responsible for poor impact properties. If such precipitates are analyzed in bulk specimens using the SEM the matrix Fe will so dilute the EDS spectrum that silicate identification is impossible.

(a) (b) Fig. 5:(a) TEM image showing the presence of precipitates characteristic of G80 weldments. (b) EDS spectrum from matrix precipitates in Fig. 5(a).

4. Low Temperature Diffusivity Measurements (3) Y In Fe-Ni alloys, a knowledge of the diffusivity of Ni in Fe-Ni (D ) below 800°C is necessary to model the growth of ferrite in austenite. Experiments that report D values have been carried out only above 1000°C. Simulation of ferrite growth therefore requires a considerable downward extrapolation of high temperature diffusivity data. Diffusion distances obtained for temperatures < 1000°C and for times < one month are smaller than the spatial resolution of the EPMA. Because of the improved resolution of the AEM, diffusion profiles that are 40 times as small as those necessary for the EPMA technique can be measured at temperatures down to 750°C. 66

Fig. 6 shows three profiles in an Fe-10.4 wt% Ni/Fe-15.5 wt% Ni diffusion couple. In Fig. 7 the measured D* values are compared with other high temperature data. The diffusivities agree well with the extrapolated values of the high temperature data.

3 o"i i 96S1 !

a ' "*| - «ENe-ra aiiRrrn rieret

Fig. 6: Measured Ni Fig. 7: Variation of measured Ni diffusiv- concentration profiles ity with temperature. The open circle (©) across the weld interface. is the measured D^ in Fe-12.5% Ni-0.15% P, and the filled circles (•) are those measured in Fe-12.5% Ni.

In summary the AEM can be used for a large number of materials applications. With improved specimen preparation techniques higher beam currents and higher operating voltages, we may eventually have even better x-ray resolution than is available on our present AEM instruments.

Acknowledgment s We wish to acknowledge the support of the National Science Founda- tion through Grants Nos. DMR 79-23278, 80-23955, EAR 82-12531, the Department of Energy through Contract No. EY 76-5-02-2408 and the National Aeronautics and Space Administration through Grant NGR 39-007-043.

References 1. Merchant, S. M., Notis, M. R. and Williams, D. B., 1982, Solid- Solid Phase Transformations, The Metallurgical Soc. A1ME, Warrendale, PA, p. 733. 2. Glitz, R., Notis, M. and Goldstein, J. I., 1982, Solid-Solid Phase Transformations, The Metallurgical Soc. AIME, Warrendale, PA, p. 691.

3. Goldstein, J. I., Costley, J. L., Lorimer, G. W. and Reed, S. J. B.3 1977, SEM/1977 I, 0. Johari, ed. Chicago, p. 315. 4. Sankar, J. and Williams, D. B., 1981, SEM/1981/I, SEM, Inc., AMF 0'Hare (Chicago), IL, p. 159. 5. Narayan, C. and Goldstein, J. I., 1983, Met. Trans., 14A, 2437. 67

FIELD ELECTRON AND ION EMISSION FROM ZIRCONIATED AND Zr FREE W CATHODES

J. Pelleg * and J.L. Fink**

Materials Engineering, Ben-Gurion University of the Negev, Beer Sheva, Israel.

**at Bell Laboratories, Murray Hill, N.J. 07974, U.S.A.

ABSTRACT Electron Beam Exposure Systems (EBES) for patterning future submicron devices have been one of the major fields uf investigation in current years [1]. The successful operation of EBES cathodes is not fully understood. The cathodes are basically single crystal emitters having a <100> orientation, and they are processed before use. The term given to this processing is activation [2], and it basically consists of heating the emitter in an oxygen ambient, after Zr has been deposited on it by some desirable technique. The purpose of activation is the lowering of the work function (WF) of preferential planes of the cathode by the ads orb ant, in this case Zr T3-8], to achieve strong emission from these planes during uninterrupted use for thousands of hours. These are {100} planes. However, reports indicate [6,9] that other planes also are associated with the lowering of the work function and thus with bright emission.

In this work the emission characteristics of zirconiated W tips is reexamined. Field electron emission microscopy (FEM) is supplemented by field ion microscopy (FIM).

EXPERIMENTAL

The system could be operated in FIM or FEM mode due to the reverse polarity of the power supply.Since the fields necessary for ionization of the imaging gas is about an order of magnitude larger than the negative fields required for electron emission, it was essential to work with ion emitters having radii in the 500-1000 A range. W single crystal emitters having <100> orientation were prepared by electropolishing.

The Zr was evaporated in situ onto the W cathodes. Before the evaporation, the Zr source assembly was outgassed for an extended period and in addition it has been kept always at a sufficiently high temperature during the stand-by period.

Evacuation of the microscope was performed by a combination of a diffusion pump, Ti sublimation pump and ion pump. The diffusion pump was cooled by a liquid nitrogen trap, which was kept permanently cold to

*~~ftork performed in Bell Laboratories, Murray Hill, N.J. 07974, U.S.A. while on sabbatical leave. 68

prevent oil contamination of the chamber. The best residual pressure achieved was 1.7x10" torr. Frequent changes in polarity of the power supply were performed on turning to FEM experiments or vice versa, and therefore it was desirable to keep as short as possible the time interval between the two modes of operation, and therefore FEM runs were performed in the 1x10" - 5x10" torr range. For FIM experiments, He gas purified through a heated vycor leak was introduced to the chamber and the ion pump valve turned off just before this operation. Only the diffusion pump was used to maintain a dynamic pressure of 1.9-2.3x10 torr for the imaging. The tip was cooled by liquid nitrogen. Heating and flashing of the tip for very short times, and its field evaporation (FE) was performed during the experiments as needed.

RESULTS AND DISCUSSION

Attention in this work was directed toward obtaining additional information by FIM on the effect of adsorbed Zr on W tips, to determine the conditions favouring this adsorption to be localized on certain planes and to see if it corresponds to the observed FEM emission from {100} crystal planes. As can be seen in Fig. 1 such a confinement of the emission is indeed possible under certain conditions even if a cold cathode is the emitter. In the absence of Zr the pattern of the FEM micrograph shown in Fig. 2 was obtained. Unlike in the presence of Zr, where only {100} planes are emitting, in the absence of Zr a smaller central (100) plane, and planes in the vicinity of {111} faces are emitting. Flashing of the tip at a higher temperature is essential to arrive at the pattern seen in Fig. 1, but it is not yet clear if it is sufficient. The purpose of the flashing is to provide sufficient thermal energy for the adsorbed atoms to rearrange themselves by overcoming barriers in their way to the preferential sites, in our case, to {100} planes. There is a great probability to find deposited atoms at a random distribution in metastable sices without this rearrangement, as can be seen in Fig. 4. Compare this FIM pattern with a Zr free pattern seen in Fig. 3. It is not yet clear that in order to obtain the desired emission shown in Fig. 1, the presence of Zr and flashing of the tip are sufficient. There are indications that this unique emission in <100> oriented emitters might be associated with blunting and faceting of the tip. These processes might occur during the application of a heated cathode, since during the operation of cold cathodes the pattern of Fig. 1 was repeatedly reproduced on tips which has undergone blunting and faceting.

CONCLUSION

The results of this work indicate that the patterns obtained by thermal- field electron cathodes can be reproduced also in cold cathodes. The presence of Zr is essential for achieving it. FIM patterns indicate the presence of Zr atoms -after flashing- on certain planes and not on others. No conclusive evidence emerges from the work if faceting is an essential prerequisite to obtain confinement of the emission to the {100} planes. It can be safely concluded that in the absence of Zr no such confinement of the emission is possible. 69

REFERENCES

1. D.R. Herriott, J. Vac. Sci. Technol., 20 (1982) 781. 2. R. Liu,unpublished work, 1983. 3. M. Good and E.W. Miller, Handbuch der Physik, 21 (1956) 176. 4. Shrednik, Sov. Phys. - Solid State, 1 (1956) 1137. 5. Shrednik, Sov. Phys. - Solid State, 3 (1961) 1268. 6. Fursei and S.A. Shakerova, Sov. Phys. Techn. Phys., 11 (196u, 827. 7. Swanson and L.C. Crouser, J. Appl. Phys. 40 (1969) 4741. 3. Swanson and N.A. Martin, J. Appl. Phys., 46 (1975) 2029. 9.

Fig. 1. Field-emission pattern of zirconiated W [100] tip. 0.65 KV; 9.1xlO"IU torr.

Fig. 2. Field-emissionQfrom Zr-free W [100] tip. 0.84 KV; 1x10 torr. 70

Fig. 3. FIM pattern of the tip before5Zr evaporation, 7.48 KV; He pressure 2.2x10" torr.

Fig. 4. FIM pattern. Distribution of the Zr on the tip before flashing. 7.75 KV; the pressure 2.2x10" torr. 71

THE TRANSMISSION ELECTRON MICROSCOPY OF NBi-5 BRAZED JOINT OF INCONEL 718

B. Grushko, 0. Botstein, B.Z. Weiss

Dept. of Materials Engineering Technion, Haifa 32000, Israel.

The structure of a brazed joint of Inconel 718 with BNi-5 filler metal was investigated. Transmission electron microscopy, X-ray diffractometry and EDSA were used in the study of the diffusion zone and of the overheated joint.

The diffusion zone in the base material (fig.l) can be divided into two sub-regions. In sub-region I* the main precipitate is a Nb-rich G-phase, while in sub-region II, the depth of which can be directly related to the width of the gap, the precipitates are chiefly carbides of Nb and Ti (fig.l, 2).

In an overheated brazed joint the filler metal penetrates into the grain boundaries of the base metal (fig. 3) which results in the formation of new phases. The dominant one was identified as a hexagonal Laves phase (Cr,Ni,Fe,Si)2 (Nb,Ti,Mo). It forms an eutectic with the γ-phase in the grain boundaries (fig. 4a,b).

A fine precipitation of y" and (Nb,Ti)C phases was revealed in the y present in the grain boundaries (fig. 4c), while in the central parts of affected grains only (Nb,Ti)C precipitation was found (fig. 5)•

A "binary" phase diagram for Inconel 718 - BNi-5 system is proposed (fig. 6).

Acknowledgement

The authors would like to express their appreciation to the Wolf Foundation and the Technion Research Fund for their financial support.

*(close to the actual liqud-solid interface) 72

Fig.l. The diffusion zone SEM Fig.2. The carbide precipi- image, deep etching tation dark field XI,500 X66,000

Fig.3. The structure of the overheated braze joint. XI,500 73

a) b)

Fig.4. TE microscopy of the grain boundary region (the overheated joint). a) The y-A, eutectic with y" precipitation c) in γ-phase, bright field, X50,000 b) Diffraction pattern of Ai Z.A.=[2423]. c) Diffraction pattern taken from γ-y". Z.A.=[001]y.

1350

1260

1090

Fig.5. Diffraction pattern of the MC precipitation in the y matrix in the vicinity of the grain boundary. Z.A.[lll]y.

(Ifi-718) %BNi -5 BNi-5

Fig.6. The scheme of the phase diagram. Inconel 718-BNi-5 system. 74

SIZE EFFECT IN RADIATION INDUCED CCGREGATION (RIS)

* ** L. Kornblit and A. Ignatiev * Materials Engineering Department , Ben-Gurion University ** Physics Department, University of Houston

A net correlation has been previously established between the volume misfit factor of a solute atom in a binary metallic alloy and the direction of its radiation induced migration [1]. Following the results obtained by H.W . King [2], which established that the atomic volumes of allotropes change little in the transition points, the volume misfit factor was defined as the mean atomic volume misfit of the solute with respect to the solvent, normalized to the concentration of the solute [3l. In Ref. 3 a brief summary of reported observations of radiation induced segregation (RIS) in binary alloys is given containing 26 alloys. The prevailing majority of these alloys (23) obey the rule stating that a positive volume size misfit results in a negative (away from the sink) solute transport and vice-versa. There are however, three exceptions, (Al-Ge, Ni-Ge and Cu-Fe; the first element in the couple is the solvent). It is difficult to see why the mean atomic volume misfit rule is violated only for these three alloys, hence here we redefine the volume misfit concept in terms more applicable to atomic diffusion.

From pseudopotential theory it is known that (at least for simple metals and alloys) the total energy of the metal (alloy), E, can be presented in the following form E = E (V) + Z if, ° where E is that part of the energy which is sensitive to the magnitude of the volume V, but does not depend implicitly on the structure, i.e. is independent of the coordinates of the "pseudoatoms".' The remaining part of energy is the structure sensitive part and is presented as a double sum over all atom pairs in the crystal, ij>.. is the interaction energy between pseudoatoms i and J. Usually E (V) is the dominant term. From the minimal energy requirement a relation exists between the volume dependent energy E and the structure sensitive energy term. Within this frame-work it is easy to understand why enthalpies of allotropic transformations for an element are much less than the enthalpy of sublimation and why the atomic volume differences during ullotrcpjc transformation are so 75 small. The explanation is that E prevails within the total energy expression therefore the tendency in the crystal is to change structure rather than volume. The close-packed behaviour of metals and alloys is also a result of the^ domi- nation of E ( V). o When one characterizes the atomic behaviour of an alloy from the standpoint of total energy, (e.g. density, stacking fault creation) the appropriate characteristic will be the atomic volumes of the solute and the solvent. But for RIS and solute atom transport what matters is, in fact, the diffusivity of the solutes. Diffusion, however, is a process which leaves the volume of the crystal unchanged but depends strongly on the mutual dimensions of the solute and solvent atoms. There- fore, for diffusion processes one should utilize a volume misfit parameter, which follows from the dimensions of the atoms rather than from the mean volumes of the atoms in the unit cell [h]. This can be done by defining the volume mis- 5 fit parameter as ^ /r ) ~1, where r^and rsoly are the metallic radii of the solute and solvent atoms, respectively. For close-packed, or nearly close-packed structures (fee, bee, hep) these radii can be taken simply as half the nearest- neighbour distance or half the bond length of the corresponding elements [5l •

We list in Table I the misfit parameter based on solute/ solvent metallic radii and compare it to that based on mean atomic volumes. It can be seen in Table I that the discrepancies noted on ref. 1, with respect to volume misfit based on mean volume per atom and direction of solute segregation under RIS are eliminated by using the volume misfit based on atomic size. REFERENCES 1. L.E. Rehn in Metastable Materials Formation by Ion Implantation, '('S'.'T. Picraux and V.J. Choyke , eds. )Elsevier (1982) p. 17. 2. H.W. King in Alloying Behaviour and Effects in Concentra- ted Solid Solutions, (T.B. Massalski, ed.7 Gordon and Breach Tl963) p. 85- 3. L.E. Rehn and P.R. Okamoto, Phase Transformation and Solute Redistribution in Alloys durinr Irradiation, (F.V. Nolfi, Jr., ed.) Elsevier (1983). It. See, for example, W .B . Pearson: The Crystal Chemistry and Physics of Metals and Alloys, '//iley-Ir:terscience (1972), p. liHi.

5. Handbook of Chemistry and Physics, 5^th Edition, CRC cress (1973) p. F197. 76

r or Table 1. Volume misfit parameters based on atomic size (r ., solv^ mean atomic volume (King) determined for a number of solute/solvent systems. Note that the discrepancies in the direction of segregation (under RIS) predicted from the King volume misfit parameter are removed when the atomic size volume misfit parameter is used.

Volume Misfit % Direction of Volume Misfit % Alloy 1 Segregation After King [2] rsolv^ sol •* r /r 1 < sol solv> - (from Ref. 1)

Pd-Cu 1.278 -20 + -19 Pd-Fe 1.2115 -27 + -12 Pd-Mo 1.3775 1.36255 -3 + -4 Pd-Ni 1.2458 -26 + -14 Pd-W 1.37095 -2 + -4

Al-Ge 1.2249 -37 + +13 Al-Si 1.4315 1.17585 -45 + -16 Al-Zn 1.3397 -19 + -6

Cu-Ag 1.4447 +44 +44 Cu-Be 1.278 1.1130 -34 + -26 Cu-Fe 1.24115 -8 + +5 Cu-Ni 1.2458 -7 + -8

Ni-Al 1.4315 +52 + 15 Ni-Au 1.44205 +55 +64 Ni-Be 1.1130 -29 + <0 Ni-Cr 1.2490 +1 +10 Ni-Ge 1.2458 1.2249 -5 + +15 Ni-Mn 1.36555 + 32 +23 Ni-Mo 1.36255 + 31 +22 Ni-Sb 1.450 +58 + 21 Ni-Si 1.7585 -16 + -6 Ni-Ti 1.4478 +57 +29

Ti-Al 1.4315 -3 + -20 Ti-V 1.4478 1.3112 -26 + -15

Fe-Cr 1.24115 1.2490 +2 +4 Mg-Cd 1.59855 1.4894 -19 + -21 77

APPLICATION OF ADVANCED COMPUTER GRAPHIC DISPLAY TECHNIQUES TO MATERIALS PROBLEMS: PHASE EQUILIBRIA AND DIFFUSIONAL GROWTH

M. R. Notis, S. K. Tarby and J. I. Goldstein

Department of Metallurgy and Materials Engineering Lehlgh University, Bethlehem, PA 18015 USA

Phase diagrams and the understanding they provide of equilibrium phase relationships in multicomponent systems are of major importance to both basic education and to advanced research in all areas of materials science and engineering. Binary systems are reasonably small in number and simple enough so that phase relations can be fully visualized in two-dimensional displays. However, most materials of engineering interest are ternary or higher order systems and their phase relations are significantly more difficult to visualize. While major advances have been made in the past decade concerning the calculation of phase diagrams from thermodynamic data, the development of methods for the graphic display of these phase diagrams has limited our ability to utilize the information generated and hence to understand and solve applied problems involving phase equilibria. We have recently used a wide variety of CAD (Computer Aided Design) systems in to display both binary and ternary phase diagrams and have found these systems to be of extreme importance in both educational and research areas.

1. Phase Diagrams—Background

The phase diagram for a binary system, which consists of two components, is easily portrayed in two dimensions (again, assuming pressure fixed). For this situation, Gibbs phase rule (F + P = C + 1) indicates a unique composition for each of the phases in equilibrium within a two phase field at constant temperature. Therefore, knowledge of temperature enables the determination of equilibrium phase compositions, and knowledge of the nominal composition of the system allows the calculation of the mole fraction of each of the phases by application of the reverse lever law to the composition tie-line. Computer approaches to display binary phase diagrams have been developed more than a decade ago, and there is currently an extensive program organized through the American Society for Metals and the American Ceramic Society to tabulate data and generate binary diagrams. The National Bureau of Standards has technical responsibility for this program.

If another variable is added, either as a third component, or by making the pressure variable, the use of a third dimension is required to accurately portray the equilibrium phases present. However, it is common to display these three-dimensional diagrams in two-dimensional sectional views, and hence visual perception of proper phase relations very often limits the utility of these phase diagrams. This problem 7S

has generated a host of approaches to the display of ternary phase relations. These include the use of coordinated sets of two-dimensional ternary diagrams, transparent overlay projections, solid models made of plaster of Paris or soap, and the use of transparent sheets of glass or plastic to represent isothermal planes. Also, a film series on phase diagrams is available from Pennsylvania State University which demonstrates some simple binary and ternary computer generated diagrams, Finally, two dimensional isothermal sections of ternary diagrams are generated by computer in the phase diagram development program of the National Bureau of Standards.

2. CAD Systems and Phase Diagram Display

As will be described in the following section, CAD computer systems have been shown to be ideally suited to phase diagram display. The CAD facilities include the following computer-graphics systems: Digital Equipment Corporation VAX 11/780 Evans U Sutherland PS300 Dynamic Display System (E/S) ApplicoTi DECSYSTEM 11/34 IBM 4341 Model Group 2 (vm) with IBM 3250 terminals. Turnkey software systems include McAuto Unigraphics (McDonnell Douglas Automation Company), SDRC (Structural Dynamics Research Corporation), and Applicon's Solids Modeling package. The Applicon software runs on its own hardware (PDP11/34), and the VAX11/780 is used to support the McAuto and SDRC software. The software on the VAX11/780 is operational in an interactive mode using VAX-11 Fortran. In addition, the software packages are capable of utilizing a finite element generator and MOVIE. BYU graphics. The VAX11/780 processor has six CAD work-stations (each having a Tektronics 4014, alphanumeric screen, function keyboard, and data tablet) and six VS11 color terminals.

Most of our work has been with the Tectronics 4014 using Uni- graphics software; this system has an extremely high resolution display but is currently limited to a 'green screen.' A new Unigraphics system with high resolution color capability is currently being installed. So far, the majority of our color display work has been on the VS11 terminals. Both of the above systems (Tectronics 4014 and VS11) are limited in terms of the speed of dynamic display. We have performed limited work with the high speed E/S system and we are developing procedures to shift data and program files developed on the Unigraphics system over to the E/S system. We are also working with both Applicon and IBM 4341 display systems and are comparing the input and output advantages and problems for all systems. Because some systems have unique advantages for either data input or display output we are trying to develop universal file transfer software so that we can take advantage of the benefits of each particular system.

Data input to CAD systems may be accomplished in a number of ways: by entering individual data points numerically into the system from the keyboard, by using equational forms for multiphase surfaces from programs written for a specific system, or by digitizing devices from hard copy of a phase diagram. For example, diagrams published in Bull, of Alloy Phase Diagrams may be laid upon the digitizer pad and traced over to input to the computer in just a few minutes each. 79

Once a phase diagram is available, there are a number of specific features or uses of the diagram that might be of interest. These include: - the number of phases present at a given alloy composition and temperature - the maximum solubility of the components in a given phase at the temperature of interest - the equilibrium compositions of coexisting phases in multiphase regions (i.e., the "tie line" compositions) - the weight fraction or atom (uole) fraction of phases present (i.e., using the lever rule) - the ability to follow the course of crystallization (solidificatiir.x) and tc understand the nature of the phase transformations which occur (i.e., reaction types)

With this in mind, the following present capabilities of the CAD system have been established: - multimode data input with direct access to interactive computation (Fortran) and simultaneous display; data may be input using any temperature scale and in either weight or atomic percent for composition - "wireframe" of solid model in three dimensions by high resolution (single color—green) display - color graphics display of wireframe model - color graphics display of solid model with color shading to enhance 3-D perspective - "exploded view" capability in wireframe or solid model (i.e., "solid puzzle" view) - any arbitrary surface view or rotation as chosen - any isothermal section or composition cut - split screen, multiview capability (four separate views simultaneously) - selection of specific phase diagram area and blow-up (expanded view) - color coding of individual lines - heightened intensity or variable line width for individual lines (future) - hidden line styling (solid, dash, dot, alternating pattern, etc.) - hidden line removal - output to terminal, multicolor printer/plotter, or to videotape.

These capabilities can be demonstrated with a number of examples. First, and easiest, binary phase diagrams can be displayed and compared with experimental data. Figure 1 shows the most recently published 'provisional' version of the Al-Ag phase diagram (1). Superimposed on this diagram is the low temperature experimental data for the equilibrium solvus line, obtained by analytical electron microscopy (2). The experimental data can be 'curve-fit' by a least squares quadratic spline, with the simple push of a button, and this curve compared to the published diagram. Also shown in Figure 1 is the data for the metastable GP Zone solvus (2). In the color version of this figure, the diagram may be displayed in one color and each of the data sets in a separate color thereby making multiple data sets quite easy to compare.

A similar example, this time involving experimental diffusion so

studies, is shown for the Ni-Al system in Figure 2. The interfacial composition data obtained from a recent study (3) concerned with the growth of the Ni Al phase formed during interdiffusion in a NiAl vs Ni diffusion couple, is shown superimposed on the Ni-Al phase diagram. We are now in the process of generating a ternary phase diagram for the Ni-Al-Cr system (4) in order to display the ternary diffusion data obtained from a research program now in progress.

The final example (Figure 3) shows a hypothetical ternary eutectic system with partial solid solubility of the three components. This type of display can be formed from a two dimensional figure with isothermal contour lines, such as are found in textbooks or journals. The perspective view shown in Figure 3 allows easy visualization of the isothermal intersections with the liquidus and solidus surfaces.

The interactive capability allows complete interfacing of display and calculation features of the system. For example, the system is capable of calculating the weight fraction or atom (mole) fraction of each phase in a two-phase field, given the temperature of interest and the nominal composition of the alloy. This is accomplished by first typing the nominal composition of the alloy and the desired temperature. The computer then locates this composition as a point in the Gibb • triangle base-plane (Figure 3) and projects a vertical line through i' parallel to the temperature axis (>.~axis). Selection of the temperature identifies the isothermal plane of interest. Further, tie-lines may be drawn to determine the composition of phases in equilibrium at the temperature of interest. The tie-lines may be input from experimental data, or, with the simplifying assumption that the tie-line is directed to the ternary end-member corner, it may be created by extending a line from the ternary corner in the isothermal plane through the intersection of the vertical (composition) line and the isothermal plane. In any event the intersections of the tie-line wi h, for example, a liquidus surface on the one end, and the solidus surface on the other end, may easily be determined. Software capability built into the system determines the length of tie-line from the composition vertical to either the liquidus or solidus surface and the ratio of this length to the total tie-line length gives the fraction of each phase present under equilibrium conditions. The composition of each ternary phase is also printed out.

In summary, CAD computer display system- can be used for data storage and display of complex phase equilibria. The interactive mode allows for rapid comparison cf the stored phase diagrams with new input data. The dynamic display capability allows for the visualization of relationships not easily discernible from two-dimensional static figures such as commonly found in textbooks and journals.

References 1. R.P.Elliot and F.A.Shunk, Bull. Alloy Phase Diagrams, _1_, 36 (1980). 2. S.M.Merchant, M.R.Notis and D.B.Williams, Met. Trans., 14A, 1825 (1983). 3. R.Glitz, M.Notis and J.I.Goldstein, p.691 in Solid-Solid Phase Transformations, H.I.Aaronson, et al., eds., Met. Soc, A1ME (1982). 4. S.M.Merchant and M.R.Notis, J. Mat. Sci. & Eng., to be published. 81

Acknowledgment The authors thank Mr. Jeff Roeder and Ms. Cathy Curtin for the actual generation of the computer figures.

M Ni Ag A I

1600

1S00 /TX

1400 • ://: \—- 1300

\ 1200 o LΑ / o>L a 1000 a> 900 •

BOO -• 700

600 00 10 20 30 40 50 60 70 80 90 100 00 10 20 30 40 50 60 70 80 90 100 Weig^t Percent Aluminum A.tom ' c Percent

Figure 1. The Ag-Al system (1) Figure 2. The Al-Ni system with with experimental data (2) super- experimental data (3) superimposed. imposed .

Figure 3. Schematic diagram of ternary eutectic system showing isothermal planes and tie-lines. 82

COERCIVITY, SQUARENESS RATIO AND MICROSTRUCTURE IN Co-W THIN FILUS

U. Admon(a), M.P.Dariel^, E. Grunbaum ^ , G.Kimmel(al and J.C.Lodder^

(a) Dept. of Metallurgy (b) Dept. of Physics, Nuclear Research Center - Negev, Israel. (c) Dept. of Electron Devices and Materials, Tel-Aviv University, Ramat Aviv, Israel, (d} Dept. of Electrical Engineering, Twente University of Technology, Enschede, The Netherlands.

ABSTRACT Thin Co-W magnetic films can be electrodeposited in a variety of non-equilibrium structures. By a proper choice of the deposition parameters it is possible to obtain alloys displaying a wide range of microstructures and, consequently, of magnetic properties. Co-W films, 200-500A thick, were electrodeposited under various plating conditions. The coercivities and squareness ratios of these films were determined by magnetometry, their microstructures by transmission electron microscopy. A close correlation between the magnetic properties and the respective microstructures has been established.

INTRODUCTION Cobalt is readily electrodeposited from aqueous solutions as a pure metal or in the form of various binary, ternary and even quarternary alloys. Over 55 such alloys have been reported in the literature (1). Tungsten, on the other hand, cannot be electrodeposited as a pure metal from aqueous solutions neither can it be co-electrodeposited with other elements, with the exception of iron, cobalt or nickel (2). In the case of Co-W, deposits with up to 65 wt./o (37 at./o) tungsten were obtained. Electrodeposited thin films often exhibit non-equilibrium phase structures which depend strongly on the deposition process parameters. This is particularly true in the case of cobalt based alloys for which the presence of the non-equilibrium fee phase is attributed to the sluggishness of the fee to hep transformation. Omi et al. (3) studied thick (20mn) Co-W films by X-ray diffraction. Their deposits contained varying proportions of a crystalline hep phase and a non-crystalline phase corresponding to the C03W stoichiometry. Rachinskas (4) Polukarov f5) and Armyanov and Vitkova (6)determined the magnetic properties of thick (over lpm) Co-W films at various plating conditions. They reported coercivities in the range of 200-600 0e, and squareness ratios of 0.6-0.8.

EXPERIMENTAL The cobalt-tungsten films were electrodeposited onto coppei-coated microscope slides. The slides had been pre-coated with a thin layer of formvar. The specimens for the magnetic measurements were prepared by cutting the slides into 1 r.m^ squares. The specimens for transmission electron microscopy were prepared by cutting 2 mm- squares on the coated surface, floating off the squares 83

by dissolving the formvar, and finally selectively dissolving the copper as described elsewhere (7). The following solutions were used for electrodeposition:

acidic pH,g/l basic pH,g/l

CoSo4•7H20 60 80

Na2WO4•2H70 0-5 0-80

MgSO4-7H2O 50 - 30 - H3BO3 - 350 Rochelle Salt-4H90 - 50 (NH4)2S04 0.1 - Sodium Lauryl Sulfate H S0 NH OH pH correction 4 4

The bath temperatures used were 22°, 50 , and 85 C, at current density of 10±l mA/cm2. The deposits had a gray metallic appearence with increasing brightness at elevated bath temperatures. At pH=2 the solutions were unstable and tended to precipitate. The deposits were dull, and it was not possible to obtain uniform samples for the magnetic measurements. The magnetic measurements were carried out using a Foner type vibrating sample magnetometer at the Twente University of Technology. The microscopic examinations were done at 100 kV on a JF.OL-JEM 7A transmission electron microscope at the NRCN laboratories.

RESULTS AND DISCUSSION The detailed correlation between the microstructure of the deposits (grain morphology, phase constitution and texture) and the deposition parameters is beyond the scope of this paper, and will be reported elsewhere. However, as the magnetic properties are determined primarily by the microstructure, the main microstructural features that have been observed will be described here. Two types of structure were obtained for different deposition parameters, and are shown in Figure 1. In most cases, the deposits were crystalline (Figure la and lb), consisting of the hep and the (unstable) fee phases in various proportions. The hep crystallites showed a varying degree of texture, with anCoO.l] fiber axis perpendicular to the film plane. However, Co-W films deposited from basic baths at room temperature and salt weight ratios above 80/10 (these numbers, and others when quoted, give the C0SO4 to Na2WC>4 salts weight ratios in (g/1) in the plating baths) were amorphous, or composed of crystallites of a very small grain size, as can be seen fTom their diffuse electron diffraction patterns (Figure lc). Figure 2 shows schematic diagrams in which the crystallography and texture of the deposits, obtained from the electron diffraction patterns, are represented as a function of the bath composition, temperature and acidity. It should be noted that the corresponding tungsten concentrations in the films ranged from 7 to 34 wt/o. Its dependence on the plating conditions will not be described here in detail. 84

The typical grain morphology of the crystalline deposits is shown in the transmission electron micrographs of Figure la and lb. The hep and fee grains are qui-axed and their size is in the range of a few hundred Angstrom. In deposits from acidic baths the mean grain size, as well as its spread, increased with the bath temperature (Figure la). IΓ. deposits from basic baths and above room temperature large single crystals, a few thousand Angstrom in diameter were evenly dispersed in a population of small crystallites, a few hundred Angstrom in diameter (Figure lb). The large crystals often showed planar defects and a subgrain structure, probably due to the grain structure of the copper substrate. A similar effect was observed in the amorphous films (Fig.lc)

The coercivity, He, and the squareness ratio, R=Ir/Is (Ir and Is are the residual and saturation magnetization, respectively), obtained from the hysteresis loops, are summarized in Figure 3 and 4. The following conclusions may be drawn: (a) The Co-W films have a medium to high coercivity. It ranges from 100-600 Oe in the crystalline films and decreases to 20-30 Oe in the amorphous films. (b) He is low for films which contain a high proportion of fee crystallites, and increases with increasing content of hep crystallites. For amorphous films He is even lower. This stems from the differences in magneto-crystalline anisotropy, which is the highest for the hexagonal phase and zero for the amorphous phase (compare Figure 2 with Figure 3). (c) He decreases with increasing perfection of texture ([00.l] perpendicular to the film plane) . This can be seer, at pH=8.5; T=85°C when the bath composition changes from 80/20 to 80/80 (see Figure 3), and at pH=6 when the bath temperature is increased from 50° to 85°C. This is consistent with the fact that the[00.l]axis, which is the easy axis of magnetization of the hep crystallites, approaches the normal to the film plane. (d) He increases when the population of large cr>stals increases. This occurs at pH=8.5 at higher temperatures and tungsten ion concentrations in the bath. However, this effect is small compared to the phase composition effect. This result is in full agreement with those of Armyanov and Vitkova (7) but only in partial agreement with those of Rachinskas (4). (e) The squareness ratio, R, increases in deposits of higher fee phase contents, and reaches values as high as 0.9. This can be attributed to the presence of four independent easy axes of magnetization in the fee phase, in contrast to the single axis in the hep crystallites. Thus, when the external field is being removed there is a greater chance for the magnetization vector to jump into an easy direction close to the field direction. The theory (8) predicts R values of 0.866 for fee and 0.5 for the hep phases. Our results agree, within the experimental error, with these values. (f) The amorphous deposits have a high value of R, as might be expected. However, R being less than unity indicates that there is some degree of short range order in the films. This assumption is supported by Omi et al. (3). (g) R decreases when the degree of texture perfection ([00.l] perpendicular to the film plane) increases. 85

(h) R decreases when the spread in grain sizes increases. This was observed for baths with increasing tungsten ion concentrations, and agrees with the results of Rachinskas (4). (i) There is a crystallographic and magnetic azimuthal symmetry in the plane of the films. This was deduced from the symmetry of the diffraction rings and confirmed by in-plane magnetic measurements at various directions.

CONCLUSIONS There is a strong influence of the deposition parameters (particularly the temperature, pH, and composition of the bath) on the microstructure, and hence on the magnetic properties. By a careful choice of these parameters it is possible to facilitate the production of magnetic alloys with a wide range of properties.

ACKNOWLEDGEfENTS The authors wish to thank Dr. T. Wielinga of the Twente University for several helpful discussions, and Mr. B. Yusov of the NRCN for the preparation of the specimens.

REFERENCES 1. Krohn, A., C.W. Bohn, Plating, March 1971, 237. 2. Sastry, B.S.R., Metal Finishing, Oct. 1965, 86. 3. Omi, T., H. Yamamoto and H.L. Glass, J. Electrochem. Soc, 119 1972, 168. 4. Rachinskas, V.S., in "Electrodeposition of Metals" Proc. 10th Lithuanian Conf. of Electrochemists, Dec. 1968, translated from Russion by the Israel Program for Scientific Translations, Jerusalem,1970, p. 51. 5. Polukarov, Yu. M., in "Electrodeposition of Alloys", V.A. Averkin, ed., Moscow 1961, translated from Russian by the IPST, Jerusalem, 1964, p. 52. 6. Armyanov, S., S. Vitkova, Surface Technology, 1_, 1978, 319. 7. Admon, U., A. Bar-Or and D. Treves, J. Appl. Phys. 44_, 1973, 2300. 8. Chikazumi, S., "Physics of Magnetism", John Wiley, 1964, Ch. 12. 86

Typical microstructures of Co-W thin films (for the notation see text) : (a) crystalline, uniform grain size (pH=4; 85°; 60/5). (b) crystalline, non-uniform grain size (pH=8.5; 85°; 80/40). (c) non-crystalline (pH=8.5; 22°; 80/40).

20 ""'40 T,°C S.W.R. Fig. 2: Dependence of phase composition and texture on the deposition parameters: (a) acidic baths, S.W.R. = 60/5 (b) basic baths, S.W.R. = 80/0 to 80/80 S.W.R. is the Salts Weight Ratio (CoSO4 to Na2W04) in the baths. Dots-fcc, bars-hep, shaded area-non-crystalline. The [00.l] texture is illustrated by the degree of alignment of the bars. 87

T,°C

22° S.W.R. Fig. 3: Dependence of the coercivity, He, on the deposition parameters: (a), (b) same as Figs. 2(a), 2(b), respectively.

80/0 30/ 80/ 80/, R0/ 10 2(} 0 R0 S.W.R.

Ir Fig. 4: Dependence of the squareness ratio, R= /js,on the deposition parameters: (a), (b) same as Figs. 2(a), 2(b), respectively. 88

STRUCTURAL ANALYSIS OF HIGH VACUUM, HIGH TEMPERATURE BNi-5 BRAZED JOINTS OF INCONEL 718 SUPERALLOY

B. Grushko, B.Z. Weiss

Dept. of Materials Engineering Technion, Haifa 32000, Israel.

INTRODUCTION

In the present paper the results of structural investigations of the brazed joint of Inconel 718 by BNi-5 filler metal are reported. Interdiffusion of elements between the base material and the brazing joint may lead to the formation of new phases (alloys) in the intermediate zone, which have a lower melting temperature than the parent material, resulting in mobility of the "real" liquid-solid interface. It is obvious that compositional elements of filler and base metals are present on both sides of the former "real" liquid-solid interface, while "location" of these elements is controlled mainly by thermodynamic and "environmental" factors as well as gap clearance.

It is therefore reasonable to assume that the post-brazing structure of the "brazing-influenced zone", which includes the brazing joint and some adjacent areas of the base material, is influenced mainly by three factors, viz. temperature, time, and gap width (TETIG).(1).

THE EXPERIMENTAL PROCEDURE

Three methods were used for the purpose of identification, namely: optical and SEM microscopy, energy dispersion analysis (EDSA) , and X-ray diffractometry. Specimens (fig.l) were brazed in a vacuum resistance furnace in accordance with thermal regimes described in fig.2. X-ray diffraction measurements were conducted on section α-a of the specimen II. A deep etching technique was applied in order partially to dissolve the surrounding y matrix. The etchant was 5% Nital electrolyte applied for 5-8 minutes with a potential of 6V and a current density of 1.3 mA/mm2. For metallographic investigations the specimens were etched with aqua regia. In addition, a thermal etching technique was used, in which the specimens were heated to 700eC, kept for 10 min at that temperature, and then cooled in air.

RESULTS

The depth penetration of the liquid into the base material, measured in the JS-μ section (see fig.3), was found to average 30 um for II-A ) specimens, 60 ym for II-C specimens, and 440 Um for specimen II-D.

*) II, III are specimen configurations (fig.l); A,B,C,D are the thermal regimes according to fig.2. 89

X-ray diffraction results on deep-etched specimens are shown in fig.4. In the II-A and II-C specimens, lines of phase 6-and G-phase were observed (γ-phase was dissolved). In specimen II-D only G-phase lines were observed in addition to the lines of the γ-matrix, which was, in this case, only partly dissolved. Because of the directional solidification, quantitative analysis was practically impossible. Rotation of the specimens did not change the intensity of the 0 lines, while peaks of the G-phase disappeared at certain orientations of the specimens and reappeared in others. This indicates a uniform distribution of the 6 compound, most probably as a fine constituent of the eutectic. The meta.llographic examinations showed the presence of the dominant γ-solid solution, coarse irregular eutectic of the G-phase and y of a flower-like morphology, irregular eutectic of 0 and γ (the 0 and G-phase could be distinguished with great certainty by the "EDS analysis).

In the II-C specimen two fine eutectics could be observed, being coloured differently by thermal etching. Approximate calculations of the eutectics1 compositions led to the following results: the "bright" eutectic: 40% G and 60% 0-phase; the "dark" eutectic: 60% G, 30% 0, and 10% γ-phases. In specimens II-A and II-C the γ-phase was found to contain small quantities of Fe, Nb, and Mo. In the specimen, II-D, the γ-phase contains significant quantities of the base metal elements, but the concentration of Si was found to have considerably decreased. EDS point analysis showed that the Nb replaces Cr in G and its content in the G-phase increases when the heating regime is changed from II-A to II-D, which also results in an increase of the lattice parameter from 11.12A to 11.22A (see fig.4).

The 9-phase does not show any presence of Nb. The results of metallographic studies and EDS analysis of the microstructure in the wide-gap joints, (Specimen III brazed by thermal regime A, b>0.25 mm) were basically similar to those described previously for specimens IIA and IIC. As the gap narrows (<0.1 mm), a joint with a centerline eutectic is formed, as shown in fig.5. The structure consists of large bright G-phase particles and small gray areas of 0. The composition of the G-phase is very similar to that obtained in II-D specimens despite the lower brazing temperature applied.

Two effects were observed as a result of the gap's narrowing: a) the volume fraction of 6 decreases; b) Nb content in the G-phase increases.

In joints with gaps narrower than "V50 ym only the presence of che G-phase could be observed,in joints ^30 ym G-phase is presented as a discontinuous chain of small particles.

DISCUSSION

On the basis of the results obtained it can be concluded that the phase formation in the BNi-5 brazing of Inconel 718 can be basically described by the Ni-Cr-Nb-Si quasi-quarternary system. Titanium can be considered as an element interchangeable with niobium, while nickel and chromium are sometimes replaced by iron. This approximate phase diagram (fig.6) was constructed by using four ternary diagrams known from the literature (2,3,4,5). A few alterations, based on experimental results, were introduced; e.g. the TT-phase (of the Ni-Cr-Si system2) was not included 90

since it was not observed experimentally. The absence of the TT-phase suggested the formation of an eutectic between the G- and the γ-phases, which was subsequently confirmed experimentally.

Occlusion of isomorphous phases, such as NbgNii6Si7, CrgNi^Siy (G) and Xj_- hexagonal Laves phases, leads to the creation of regions in which only a single one of those phases is present. It should be assumed that the phase distribution and compositions at room temperature are very close to those shown in the diagram (fig.6). The initial location of the BNi-5 alloy is represented by point (1) . This point is approximately situated at the boundary of the region y+Q and the region y+9+G. The influence of the base metal on the structure of the NBi-5 brazing can be illustrated by shifting the figurative point from position (1) to position (2), which represents the approximate location of the base material - Inconel 718. (see fig.6b; the straight line approximation was used), Initially, the Y-6-G structure is present in the molten alloy. Gradual shifting along the line (1-2) (see fig.6b) shows that once the 6 phase disappears, the structure should consist only of G- and γ-phases, and finally only the γ-phase remains. Experimentally this situation is simulated by varying the width of the gap. The wider the brazing gap, the closer the phase constitution is to position (1), and the narrower the gap, the closer the constitution is to position (2). The predicted changes in volume fractions of the G-and 6-phases were in practice shown to occur by changing the width of the gap.

The wide spectrum of brazing experiments, conducted in different thermal conditions and for different brazing gaps, showed that all the phases that appear in the constructed phase diagram were actually observed in the brazing joint. T-iis means that, although the brazing process does not seem to follow conditions of equilibrium, yet the formation and outward appearance of the different phases can be predicted from the phase diagram.

ACKNOWLEDGEMENTS

The authors would like to express their appreciation to the Wolf Foundation and the Technion Research Fund for their financial support.

REFERENCES

1. R. Johnson, Weld. res. suppl. 1981, Vol.60, pp.185-193.

2. E.I. Gladyshevskii and L.K. Bornsevich., Russian J. Inorg. Chem. 1963, Vol.8, pp. 997-1000.

3. E.I. Gladyshevskii et al., Inorg. Mater., 1969, Vol.5, pp. 1882-1883.

4. H.J. Goldschmidt and J.A. Brand, J. Less-Common Metals, 1961, Vol.3

5. L. Kaufman and H. Nesor. Proceeding of the conference on In Situ Composites. Sept. 5.8. 1972 Lakerville. Conn. Vol. Ill, Publ. NMAB-308-III, Washington D.C. 1973. 91 11

Fig.l. The specimens. Fig.2. The thermal cycles of brazing

Fig.3. The B-β section Fig.5. Micrograph of of specimen Il-C, aqua III-A specimen, thermal regia, X100. etch., X600 Fig.4. Diffraction patterns for type II specimens .

Ni-Cr

(a) (b) Fig.6. Schematic diagram of the Ni-G-Si-Nb system for 800-1000°C (a) general view; (b) ternary section. 92

PHASE RELATIONS IN THE Cu-Nd SYSTEM ON THE Cu-RICH SIDE 1 2 2 3 C. Laks , J. Pelleg , and L. Zevin ' Israeli Military Industries, Ramat Hasharon, P.O. Box 1044, 2 3 Materials Engineering Department and Institutes for Applied Research, Ben-Gurion University of the Negev, Beer-Sheva, Israel

Major inconsistencies in RCu systems (R = rare earth element) occur in the 75-88 atom % Cu range (1-7). No data are available for the NdCu system. The objective of this communication is to compare the existing RCu systems and decide which of them behaves similar to the NdCu system.

EXPERIMENTAL

Alloys of Nd-Cu were prepared with 63-93 atom % Cu by arc melting and were annealed for homogenity. The specimens were examined by X-ray diffraction, differential thermal analysis (DTA), and metallography. Standard techniques were used for metallographic examination by optical microscopy and scanning electron microscopy (SEM) with energy-dispersive analysis (EDAX).

RESULTS AND DISCUSSION

Table 1 listt- •-. - '• ntermetallic compounds identified in this work together with informa-":u available frore the literature, and Fig. 1 is a tentative diagram of the phase relations in the system under consideration. Five compounds were identified, NdCug being the richest in Cu. A previously unknown NdCu4 compound was also detected. Most of the peaks could be indexed on the basis of an orthorombic cell, but definite structure determination has to await work with single crystals. The absence of certain systematic extinctions rules out the possibility of the Pnnn space group suggested for RCU4 type couponds (1). It is interesting to note that in the LaCu system LaCu4 was not detected, unlike in other systems (1,5) where R was a light rare-earth element. This difference between the LaCu and NdCu systems merits further investigation.

CONCLUSIONS An orthorombic NdCu4 compound was detected in the NdCu system and in this regard it resembles the CeCu and SmCu (7) systems. No definite structure determination was possible by the technique used, and complete evaluation of its structure will require single-crystal work. 93

Table 1. Crystallographic data for phases in the Cu-Nd System

Phase Crystal system and Lattice parameters (A) Reference structure type a b c

NdCU- Orthrombic; CeCu 8.092 5.062 10.105 8 o bc 7.952 5.044 10.203 This work

NdCu Hexagonal; CaCu,- 5.104 - 4.107 This work

5.097 - 4.112 This work

NdCu Unknown - - - This work

NdCu Orthoronibic ; CeCu2 4.387 7.059 7.420 9

4.384 7.096 7.417 This work

NdCu Orthorombic ; FeB 7.32 4.55 5.59 10

7.302 4.569 5.578 This work

REFERENCES

1. T.B. Rinehairaner, D.E. Etter, J.E. Selle and P.A. Tucker, Trans. Metall. Soc. AIME, 230 (1964) 1193.

2. S. Carifici and A. Palenzona, J. Less-Common Met., 53 (1977) 199.

3. E. Franceschi, J. Less-Common Met., 87 (1982) 249.

4. A. Iandelli and A. Palenzona, J. Less-Common Met., 25 (1971) 333.

5. K.H.J. Buschow, Philips Res. Rep., 25 (1970) 227.

5. L.A. Ofolubkov, N.M. Shibanova, Yu. G. Sakonov and G. Ya. Fedorova, Russ. Metall. N6 (1977) 147.

7. K. Kuhn and A.J. Perry, Met. Sci. J. 9 (1975) 339.

8. K.H.J. Buschow and A.S. Van Der Goot, J. Less-Common Met., 20 (1970) 309.

9. A.R. Storm and K.E. Benson, Acta Crystallogr. 16 (1963) 701.

10. A. Iandelli and A. Palenzona, in K.A. Gschneidner and L. Eyring (eds.), Handbook on the Physics and Chemistry of Rare Earths. North Holland, Amsterdam, 1979. 94

So <5 o o noo \r ta t3 5> 1 \

1000 s> \ 917*

866° 2 900 829" 843* 2 S.S00 ..

700

600

0 0 10 15 20 25 30 35 Atomic percent Nd

Fig. 1. Partial Cu-Wd diagram. 95

EFFECT OF SMALL ADDITIONS ON GRAIN REFINEMENT OF A 14 CARAT Au-Ag-Cu-Zn ALLOY * M. Fishman, L. Gal-Or, G. Iram Center for Noble Metals, Institute of Metals, Technion, Israel * Dumax Corp. INTRODUCTION AND EXPERIMENTAL PROCEDURE The process of jewellery fabrication by cold working inevitably involves intermediate annealings aimed at restoring alloy ductility by recrystallization of the work-hardened structure. On the other hand, the annealing treatment can lead to a coarse grained structure as a result of excessive grain growth following recrystallization. Due to the fact that, plastic flow in crystals is orientation dependent on further deformation,, especially if operations like deep drawing or bending are involved, this structure can manifest itself undesirably by producing a surface effect known as "orange peel". The purpose of this work was to find an addition (or a combination of additions) to a 14 carat Au-Ag-Cu alloy containing 6% Zn that would act as an effective grain refiner of the re crystallized structure in a wide range of annealing temperatures and, preferably, as refiner of the cast structure as well. The following additions and their combinations were investigated (wt %). Ir 0.02, 0.05; Zr 0.04, 0.05, 0.1; Co 0.05, 0.1, 0.15, 0.2, 0.3; B 0.005, 0.01, 0.1; 0.1 Zr + 0.005 B; 0.15 Co + 0.005 B; 0.1 Co + 0.05 Zr +0.005B. Alloys were prepared by induction melting in graphite crucible under protection of argon. The solidified structure was investigated on samples cut from both bottom and top sections of the ingots. The rest of the ingots was cold worked into wire 1 mm in dia with intermediate annealings at 680°C. The final annealing treatment preceding determination of recrystallized grain size was performed at 600, 650, 680, 700 and 750°C at various time intervals from 0.5 min to 2 h.

THE INFLUENCE OF SMALL ADDITIVES ON THE ALLOY STRUCTURE Ingots of the unmodified 14 carat Au-Ag-Cu-Zn alloy showed a well pronounced dendritic segregation [fig- 1 left) with dendritic branches being coarser in the upper part of the ingots. The grain size of the annealed microstructure depends drastically on annealing time and temperature. For example, the rise of temperature from 600 to 750°C results in 20-fold increase of average grain size after 15 min annealing. Boron and cobalt introduced separately had no grain refining effect on either cast or annealed structure of the alloy. Zirconium has a slight grain refining effect on the recrystallized structure that is more marked at higher Zr content. Cobalt and zirconium are much more effective when introduced in combination with boron. Combined alloying with 0.1% Co, 0.05% Zr and 0.005% B has a strong and reproducible grain refining effect on the recrystallized structure at all annealing temperatures. However, alloys with the complex Co-Zr-B addition have a higher hardness and lower plasticity than the unmodified alloy. Iridium produces drastic grain refining of the cast structure (Fig.l) and also substantially reduces the recrystallized grain size (Fig. 2B). Both effects enhanced in the alloy with higher Ir content. A fine grained recrystallized structure in iridium containing alloys is retained after 96

threefold remelting. The grain refining effect of iridium on the cast structure was even enhanced by remelting. Iridium does not reduce alloy ductility. Tensile tests on 1 mm wire showed an increased alongation and unchanged strength. In order to verify whether the .grain refinement achieved by addition of 0.05% Ir is adequate to prevent the formation of "orange peel", the Ericksen test has been carried out on sheets of unmodified and Ir-containing alloy annealed for 15 min at 720°C. As can be seen in Fig. 3 a distinct "orange peel" pattern develops on the surface of the unmodified alloy whereas the Ir-containing alloy gives a satisfactorily smooth surface. For the two alloys that demonstrated a significant grain refining effect, the one with 0.05% Ir and the other with the complex Co-Zr-B addition, attempts have been made to reveal the distribution of additives in the alloy by the use of transmission electron microscopy and electron probe microanalysis. However, none of the techniques has proved to be effective.

GRAIN REFINEMENT DURING SOLIDIFICATION Possible mechanisms of the effect of minor additions on grain refinement during solidification can be summarized as follows: 1. Minor additions of highly active elements which may react with the atmosphere, crucible, other minor components of the alloy or with some un- controlled impurities can form minute solid particles of oxides, carbides etc. On further cooling, these particles provide sites for so called "heterogeneous nucleation" of the alloy crystals, thereby diminishing the extent of supersaturation needed for solidification and increasing the number of growing crystals. This mechanism has been experimentally confirmed by X-ray of TiC particles in an Al alloy with 0.1% Ti (1). 2. If an alloying addition forms with the major component of the alloy a phase diagram with an eutectic (or peritectic) point situated very close to the pure major component, even small amounts of the alloying component give an alloy of hypereutectic (or hyperperitectic) composition. In these alloys a more or less wide temperature range ATp exists between the liquidus and eutectic (peritectic) lines where primary crystals of the minor alloying component (or its compound with the major component) will precipitate. Due to a very low concentration of the minor addition very fine particles are formed and when the eutectic (peritectic) temperature is reached these precipitates will provide centers for heterogeneous nucleation and growth of the crystallites of the (essentially pure) major component. An additional factor that accounts for a more effective grain refinement in alloys in comparison with pure metals (1) is a concentration gradient formed in the melt in front of the growing crystals. Due to a difference in composition between solid and liquid phases at a given crystallization temperature and a limited diffusion rate in the melt, the layer adjacent to the growing surface will become enriched with the component that preferably remains in the liquid phase. This compositional change means a more or less sharp decrease, in accordance with the liquid slope, of the equilibrium freezing etemperature of the melt in the layer or, in other words, a retardation of crystal growth at a given cooling rate. Due to this, a greater number of potential centers of heterogeneous nucleation is activated on further cooling and better grain refinement is achieved. The wider the separation between solidus and liquidus curves and the steeper the liquidus slope, the greater retardation of growth rate occurs (2). The well pronounced dendritic segregation 97 observed in the 14 carat jewellery alloy studied (Fig. 1 left) indicates that a significant compositional gradient develops during its solidification which enhances grain refinement if this effect is produced by minor additions according to one of the two mechanisms mentioned above. The observed behaviour of the alloy with additions of Ir, B, Zr, or Co during solidification is in agreement with the corresponding binary phase diagrams. Compositions of the alloys containing B, ZT and Co fall well into hypoeutectic range so that no precipitation of B, Co or ZrAuj can occur during solidification. Contrary to that, alloys with 0.02 and 0.05% Ir are h>pereutectic and grain refinement induced by Ir precipitation at the first stage of solidification can be expected. The enhanced effect observed at the higher Ir content is likely due to a greater numbers of Ir precipitates formed in this alloy. It can be also assumed that the positive effect of remelting is due to a more uniform Ir distribution achieved by remelting.

RETARDATION OF GRAIN GROWTH DURING ANNEALING The additives studied can be divided into two groups according to their behaviour at the annealing temperatures: i) iridium that is practically insoluble in the alloy and probably is present in the form of fine particles formed during crystallization, ii) the rest of additives which at the concentrations used are in the solid solution. According to theoretical predictions (3), both finely dispersed particles and solute atoms can affect the migration of grain boundaries and thus restrict grain growth though by different mechanisms. The driving force for grain growth in a recrystallized structure is the excess of energy associated with grain boundaries and it is inversely proportional to the radius of the boundary curvature. Fine particles cause a local increase in boundary length and thus create a drag of boundary motion. At a given volume fraction f of the particles this effect is inversely proportional to the average particle radius T_. Moreover, since the average boundary curvature and hence the driving force diminish with the increase of average grain size, at some stage the particles can prevent further growth. The upper limit for the grain size D in the pre- m sence of particles is Dm = 4/3 (r/f) . For the alloy with 0.05 wt.% Ir f is 3 x 10 and the effect of the additive depends on the degree of finess of Ir particles. If we assume their size to be 0.1 ym, the maximum grain size Dm in this case is about 0.2 mm. It means that precipitates of this average size cannot prevent appreciable grain growth during prolonged annealings (as it indeed occurred in the Ir-containing alloy afte 1 h annealing at 750°C), but the retardation of boundary migration they produce could result in the grain refinement that was observed after 15 min annealings . The retarding effect of very small amounts of soluble additives on grain boundary migration in pure metals is a well established fact (4).It is explained as an effect of solute atoms segregation along grain boundaries accompanied by energy gain and described by the following equation:

CΓ, = CQ-exp(-E/RT), where C« and CQ are boundary and bulk concentrations or the solute atoms respectively, E is the interaction potential between the solute atoms and the boundary, T is the annealing temperature. Be- cause the boundary region is characterized by a less regular and less dense atomic structure, size misfit between the solute atoms and the matrix lattice is the most evident reason for the boundary segregation. It is likely, however, that other factors such as difference in valency or type and strength of interatomic bonds between host and foreign atoms 98

can play some role, too, in formation of the foreign atoms atmosphere around grain boundaries. When such a boundary is caused to move, the interaction of foreign atoms with the boundary creates a drag and, unless the driving force for boundary movement is large enough, the boundary is compelled to carry the atmosphere of foreign atoms along with it. In this situation, the boundary velocity is controlled by the rate at which foreign atoms can diffuse behind the boundary (3,4") . Obviously, in spite of retardation of grain growth no upper limit can be imposed by boundary segregation of solute atoms on the grain size at long time annealings. The other difference between effect of finely dispersed particles and that of boundary segregation is that the latter diminishes with the increase of annealing temperature as a result of both increased diffusion mobility of impurities and decreased tendency for grain boundary segregation as follows from Eq. 2. However, the grain refining effect of soluble additives (or the lack of it) as observed in this investigation is not consistent with the theoretical predictions. Of soluble additives introduced in various amounts (0.005-0.1% B, 0.04-0.1% Zr, 0.1-0.3% Co) only Zr gave a moderate grain refinement whereas B,which represents two extremes of the size misfit in respect to the alloy lattice, had no effect at all. At the same time the lowest amount of boron used (0.005%) can have significant effect, while added together with 0.05% Zr and 0.1% Co. (Fig. 2 ). It is evident that boundary segregation as described by Eq. 2 cannot account for the effect because the total bulk atomic concentration of the additives in this alloy is lower than in many alloys where the same additives were introduced separately without appreciable grain refining effect. Another observation inconsistent with the effect of the atmosphere of solute atoms at the boundaries is a low rate of grain growth observed at 750°C. The strong enhancement of grain refining observed when complex alloying by B, Zr and Co is used suggests that there must be an interaction between the additives resulting in a limited mobility of grain boundaries. A.s noted by Losch (5) , chemical interaction between two solutes is most likely to occur at grain boundaries where it is promoted by a higher concentration of solute atoms resulting from their grain boundary segregation. Because boron forms borides with both Zr and Co, we can expect that it is precipitation of boride particles at the grain boundaries that reduces mobility of the boundaries and prevents extensive grain growth in the alloy containing additions of B, Zr and Co.

REFERENCES

1. A. Cibula J. Inst. Metals, 1949, 76_, 321-359. 2. W.A. Tiller J. Metals, 1959, 11_, 512-514. 3. R.W. Cahn "Recovery and Recrystallization", Ch.19 in: "Physical Metallurgy", ed. by R.W. Cahn, Elsevier North-Holland Publ. Co, pp. 1129-1197. 4. P. Gordon, R.A. Vandermeyer in "Recrystallization, Grain Growth and Textures", ed. H. Margolin, ASM, 1963, pp. 205-266. 5. W.H.P. I.osch ScriptaMet., 1977, 11, 889-892. 99

Fie. 1: Effect of additions of 0.05% Ir on the as-cast microstructure.

Fig. 2: Effect of small additives on the recrystallized structure annealed for 15 min at 750°C. A) unmodified alloy; B) addition c 0.05% Ir; C) addition of B + Zr + Co.

Fig. 3; Surface of the alloy (left) and the same alloy with the addition of 0.05% Ir deformed by Ericksen cupping test. 100

REINVESTIGATION OF THE Pr-Ga SYSTEM IN THE 66~TU0 AT % RANGE

J. Pelleg - Ben-Gurion University, Beer-Sheva D. Dayan, G. Kinunel - Nuclear Research Centre-Negev

INTRODUCTION The phase diagram of Pr-Ga system has beer, the subject of several investigations(1,3). The existence of the compounds PrGa and PrGa2 in this system has been reported by all investigators. However the existence of some of the previously compounds, i.e. Pr3Ga(l-2), PrsGa3(2,3), PraGa2(1-2,4) and Pr2Ga(3) and their structure are controversial. Never- theless, the topic of this communication is concerned with the Ga-rich side of the Pr-Ga system in the 67-100 at. % Ga range. In previous work (5,6,12) it has been established that some of the RGa2-type compounds (R is a rare earth) have a wide range of homogeneity wMle other not(7). This can be understood on the basis of a model suggested by Pelleg and Zevin(6). Similarly, the possible existence of an additional intermetallic compound on the gallium rich side has been pointed out(8). This has been confirmed(9) and the structure has been identified as RGa6. In order to clarify the predictions based on the Pelleg and Zevin model(6) regarding solid solubilities of Ga in PrGa2 we undertook a study using x-ray diffraction, differential thermal analysis, metallography, SEM and diffusion couples techniques. Particular emphasis was placed OTi determining the crystal structure of the intermediate phases and the evaluation of the change of the lattice parameters of the e phase with compositions within the homogeneity range. (£=Pri-xGa2n+x)). RESULTS AND DISCUSSION Fig.l shows the Ga-rich side of the Pr-Ga system derived from the results of this investigation. Data from reference 1 and from reference 2 were included for comparison. It can be seen that PrGa2 shows a rather wide range of homogeneity which extends to about 78 at. % Ga. The sym- bol e was assigned to this phase. In addition a gallium-rich intermetallic compound was detected in this system having the composition PrGa6. The phase n was assigned to it. These two phases are now discussed. The e phase PrGa2 is the stoichiometric composition of the e phase which forms congruently at 1470°C(3) and has a hexagonal structure of A£B2 type. Its lattice parameters a-e listed in Table 1. In accordance with other RGa2 compounds, where R is a light rare earth metal, the solubility range of Ga in PrGa2 is broad and extends to about 78 at. % Ga. There is also an increase in the lattice parameters with the addition of Ga. 101

An increase in the unit cell is expected according to the Pelleg - Zevin model(6) which is based on pairwise substitution by 6a atoms of the R constituent in the basal plane. Diffusion-couple results indicate the presence of a layer between Pr and PrGa6 as seen in Fig. 2. In this layer a gradient in concentration of Ga was found from the stoichiometric composition to about 79 at. % Ga as indicated by electron microprobe measurements.

The r| phase

PrGa6 forms peritectically at 466±6°C. The formation of PrGa6 was established also by diffusion - couple experiments. In Fig. 2 such a couple is seen between pure Ga and pure Pr. The predominant phase that forms at a relatively low temperature in a short period (360°C/2h) is PrGae. In agreement with similar RGa6 compounds(5-7, 9,10] PrGa6 has a tetragonal structure of the PuGa6 type(ll). Its lattice parameters are presented in Table 1. One of the authors (J. Pelleg) holds the Samuel Ayrton Chair in Metallurgy.

REFRENCES

1. Tandelli, Gazz. Chim. Ital., 19_, (1949) 70. 2. S.P. Yatsenko, A.A. Semiannikov, B.G. Semenov and K.A. Chuntonov, J. Less-Common Met. 64, (1979) 185. 3. S. Cirafici and Franceschi, J. Less-Common Met., 66, (1979) 137. 4. A.E. Dwight, J.W. Downey and R.A. Conner Jr. Acta Crystallogr. 23_, (1967) 860. 5. G. Kimmel, D. Dayan, A. Grill and J. Pelleg, J. Less-Common Met., 75, (1980) 133. 6. J. Pelleg and L. Zevin, J. Less-Common Met., 77_, (1981) 197. 7. J. Pelleg and G. Kimmel, Materials Science and Eng., 52_ (1982) PI 8. R. Manory, J. Pelleg and A. Grill, J. Less-Common Met., 61_, (1978) 293. 9. J. Pelleg, G. Kimmel and D. Dayan, J. Less-Common Met., 81_, (1981) 33. 10. S.E. Hazsko, Trans. Metal. Soc. AIME 221_, (1961) 201. 11. F.H. Ellinger and W.H. Zachariazen, Acta Crystallogr. 19, (1965) 281. 12. D. Dayan, U. Atzmony, and M.P. Dariel, J. Less-Common Met., 87, (1982) 87-98. 102

• -This work *-Ref. 2 • -Ref. 3

1,0 50 60 70 80 90 100 ATOMIC PER CENT GALLIUM

Fig. l. The ga.1 lium-rich side ot the Pr-Ua System. 103

Table 1 Crystallographic data in the 50-100 at. % range of the Pr-Ga System

Compo- Struc- constants Phase sition ture at.% Ga type a b c

PrGa_ 66.7 4.2817 4.2898 This work 4.272 4.298 10 4.3021 4.2924 This work 75. 4.3182 4.3022 This work 78.7 4.3167 4.3113 This work RGa, 85.7 7 PuGao, 6.014 .654 This work

Fig. 2. Diffusion couple of praseodymium-gallium showing the interface layer of the e phase C?r x Ga ). Heat treated at 360°C for 2h. Pr is the grey and Ga is light grey layer, x540. 104

TEXTURE IN LOW ALLOYED URANIUM ALLOYS

* * ** J. Sariel , G. Kiminel , and J. Pel leg

* Nuclear Research Center Negev P.O.B. 9001 Beer-Sheva, Israel. ** Ben-Gurion University Beer-Sheva, Israel.

INTRODUCTION Orthorombic uranium is known to exhibit strong anisotropy and its consequence might be critical during irradiation growth. The degree of the anisotropy expressed by the dimensional changes in polycrystalline uranium, depends on the grain size and the preferred orientation that is present in the material due to the prior processing. Thus, randomly oriented fine grained uranium might be almost free of irradiation growth on a macroscopic scale, due to mutual canceling of this effect, while coarse grained uranium with a strong preferred orientation may change its dimensions in a very anisotropic manner. In practice it is very important to be able to evaluate quantitatively the degree of preferred orientation, and to decrease as far as posible, if not completely to eliminate this preferrred orientation. In this work two uranium alloys, namely, adjusted uranium and uranium chromium alloys were chosen to investigate anisotropy, by means of x-ray diffraction. Schulz method diffraction (1,3) and the inverse pole figure technique (2,3) were used for texture determination.

EXPERIMENTAL The composition of the adjusted uranium is 0.02-0.05% Fe, 0.05-0.12% Al and 0.08-0.09% C. It was vacuum melted and cast into 3.6 cm diameter rods. The rods were 3 solution heat treated in a salt bath at 740°C, water quenched and then vacuum annealed at 540°C. The uranium 0.1% Cr alloy was prepared in the same way, 3 solution heat treated in a salt bath at 720°C, quenched to a second salt bath at 540°C for isotropic transformation and finally vacuum annealed at 520cC. The specimens were slices from the rod, with the surface of interest perpendicular to the rod axis. Cu Kα radiation was used for the regular diffraction, and a curved graphite crystal monochromator was put in front of the proportional detector. The area illuminated by the x-ray beam was such that hundreds of grains were covered,to achieve good statistical sampling. For the Schulz method diffraction a texture attachment was used in conjunction with the goniometer. The recorded intensities were sent in parallel to a TTY paper tape punch for subsequent computerized data processing. RESULTS In the Schulz technique the diffracted intensities were corrected for background and defocusing, normalized and were devided into six groups. The groups were graded on a logarithmic scale. The relative intensities, P ((j),a), were presented on a stereographic projection, yielding the direct pole figure. Here a is the angle of rotation of the specimen about an axis 105

normal to its surface, and is the tilt angle about an axix located in the diffraction plane perpendicular to the goniometer axis. In figure 1 two representative direct pole figures are shown, one of an as cast and the second of fully heat treated adjusted uranium. At the bottom of each figure the relative intensity scale is presented. While the relative intensity 3cale for the as cast material spreads up to 19.32, for the fully heat treated specimen, it reaches only 2.96. Similar pole figures were obtained for the U-0.1% Cr alloy, indicating a similar reduction in the intensity scale for the fully heat treated specimens. In the inverse pole figure technique, pole densities Pi were calculated from the peak intensities according to the equation: I. / 1°. Pi = /' *

Here n is the number of peaks. I. are the intensities from the specified planes of the specimen and I? are the calculated intensities of a randomly oriented polycrystalline material. These P^ values were used to construct the inverse pole figures. Figure 2 shows a quadrant of the inverse pole figure for an as cast and a fully heat treated adjusted uranium. Here also a reduction was observed of the pole densities from a rang of 0.27-2.59 in the as cast material to a rang of 0.58-1.83 in the heat treated specimen. Similar reduction was obtained for the uranium 0.1% Cr alloy.

DISCUSSION AND CONCLUSIONS B .th, the direct and inverse pole figures characterize the manner of preferred orientation qualitatively rather than quantitatively. Furthermore, they are applicable in materials which were subjected to thermo-mechanical treatments, resulting in systematic and definite texture. However, our results indicate that specimens which were not subjected to such treatment, like as cast, or just heat treated specimens, show non reproducible texture even if they came from the same batch. This observation can be attributed to the fact that as cast material, lacking any thermo-mechanical treatments, is supposed to be free of preferred orientation, and if there exists some degree of preferred orientation, it is not systematic and not reproducible. Therefore a technique is suggested to overcome this difficulty. In essence, two parameters calculated form the results of the Schulz method are suggested to characterize quantitatively the fine non-systematic texture in as cast and annealed materials. For a specimen ideally free of preferred orientation, the normalized relative intensities, P (,cO, should all be equal to one, and thus, the standard deviation of these values is zero. For any other specimen the values of P would spread about the value of one, (although the mean value would still be one) and the standard deviation would be greater than zero. The higher the maximum value of P (

In order to confirm the validity of these two parameters, two extreme cases were examined, a sintered silicon, which is expected to have minimal preferred orientation, and a rolled uranium which has a pronounced texture. Indeed the results for P and SD were confirmed, resulting for the silicon 106

inimininmnniMMiiim.nnmH«i..^^.Mm^L- i ••,•„..'••Li.Miniiiiiiiiin.ii.w.,.•.,•••..„. •»••• :::::::

• •••••••if" >»KiW>> •

sssss:::: • — "~----I«***lXI*""«II.M««K"ltM*Hf«M«Mtt"*"-—-***11"**'"** -"""? ass:.KKS".::. :

wMjnnnnnnnnr mi i mm 11 ii"i*' ' *II tttu.*:V.,',..'.'.',lt MMMM D00000000011»IlIWIIHmi*«t

ess xSSSSSXSS

I 1 1 I i •—*• i I mix I I •MΜ I f MH f I *" t I'..— 1 j •••»• i t inn 1 1 ww i I w-a ! i i i !"-* ; i ***** i I mw I i *M*I i i —•• i ! • j:"" j j "'•" j j!!!?? i J "*"• I"***;

Fig. 1 : Representative direct pole figures of an as cast (left) and a fully heat treated (right) adjusted uranium.

002 023 022 021 041 020 002 023 022 021 041 020 0.92 l.?3 1°83 0?81 0779 0.58

1.25

60 200 0 95

Fig. 2 : Representative inverse pole figures of an as cast (left) and a fully heat treated (right) adjusted uranium. 107

in 1.32 and 0.09 and for rolled uranium 11.6 and 1.2 for PmaY and SD lIictA respectively. P and SD for adjusted uranium observed in this work are 14.3±4.9 and l.?£0.2 respectively for as cast material, and 3.2±0.9 and 0.4±0.1 for the fully heat treated specimen. These values are statistical averages from a large number of specimens examined. Similar trend was observed in the U-0.1% Cr alloy, but due to smaller number of specimens investigated, the statistical sampling is not adequate. Hie conclusions of this work are: 1. The suggested parameters seem to be good quantitative indicators for the level of preferred orientation even in the case of as cast materials. 2. Cast materials have some degree of preferred orientation but this is non-systematic and non-reproducible. 3. The preferred orientation can be reduced by appropriate heat treatments.

* One of the authors (J. Pelleg) holds the Samuel Ayrton chaii in metallurgy.

REFERENCES 1. L.G. Schulz, "A direct method of determining preferred orientation of a flat reflection sample using a Geiger counter x-ray spectrometer", J. Appl. Phys. 20 , 1030-1033 (1949). 2. G.B. Harris, "Quantitative measurment of preferred orientation in rolled uranium bars", Phil. Mag. 43_ , 113-123 (1952). 3. J. Sariel, "Texture in low alloyed uranium alloys", M.Sc. thesis in engineering, Ben-Gurion University of the Negev, Beer-Sheva, July, 1981. 108

MICROSTRUCTURE AND PROPERTIES OF TUNGSTEN-BASED HEAVY ALLOYS

D V Edmonds

Department of Metallurgy & Science of Materials, University of Oxford, Parks Road, Oxford 0X1 3PH, UK

ABSTRACT

The embrittlement of tungsten heavy alloys of typical nominal composition (wt%) W-5Ni-5Fe and W-7Ni-3Cu by the interfacial segregation of trace impurity elements has been examined by Auger electron spectroscopy. The occurrence of precipitation at tungsten-matrix and tungsten-tungsten interfaces, and in both the matrix and the tungsten phases, has been examined using high-resolution microanalytical transmission electron microscopy. The influence of certain processing variables and heat treatment is apparent, and the effect of segregation and precipitation on mechanical properties is discussed.

INTRODUCTION

Tungsten-based heavy alloys are manufactured by liquid-phase sintering techniques1•2, and after cooling from the sintering temperature the microstructure consists essentially of a continuous network of approximately spheroidal tungsten grains embedded in a Ni-Fe(or Cu)-W matrix binder- phase. A useful combination of mechanical properties is exhibited, but there is a marked variation in properties, particularly impact resistance, with composition and processing conditions. A better understanding of the change in microstructure with heat treatment, and of the microstructural dependence of mechanical properties, formed an objective of a research programme of which this paper outlines the results obtained so far and refers generally to alloys of typical nominal composition (wt%)90W-5Ni-5Fe and 90W-7Ni-3Cu. Parts of the work are published in more detail else- where3"7 .

EXPERIMENTAL PROCEDURES

A detailed account of the experimental techniques employed is given else- where3 '5~8.

EXPERIMENTAL RESULTS AKJ DISCUSSION

Alloy Microstructure

The microstructural appearance of the alloys is illustrated by Fig.l. It is conventional commercial practice to furnace cool the alloys from the 109

sintering temperature, and this results in basically similar structures, although differences in specific parameters can be identified by measurement. Table I documents the matrix volume fraction, average tungsten grain size, and contiguity of the tungsten particle network as a function of selected treatments. (The contiguity of the tungsten network, C^, can be defined9 as the average fraction of surface area that a given tungsten particle shares with its !0Qum neighbours, and is given by 2N, WW Fig.l Micrograph (optical) of as- (1) + 2N. sintered furnace-cooled N.WM ' WW W-5.2Ni-4.8Fe alloy. (Muddle) where [%\n| is the average number of tungsten grain boundaries and Nyjyj is the average number of tungsten-matrix interphase boundaries that are intercepted per unit length of randomly positioned line).

Table I Measured microstructural parameters Average tungsten Volume fraction Contiguity of Specimen particle size, μm of matrix, % tungsten particles, cw W-5•2Ni-4.8Fe

As sintered, ]FC 21.5 ± 1.2 30.2 + 1.,8 0.351 ± 0.023 2h @ 1200°C, WQ 0.358 lOh @ 1200°C, WQ 30.0 ± 2..0 0.416 ± 0.027 2h @ 1350°C, WQ 0.377 5h @ 1350°C, WQ 0.415 lOh @ 1350°C, WQ 29.6 ± 3,.2 0.452 ± 0.026

W-7.2Ni-2.4Cu

As sintered 23.1 ± 1.3 25.4 ± 2,.1 0.405 ± 0.021 lOh @ 115O°C, WQ 23.8 ± 2.6 0.404 5h @ 1350°C, WQ 0.431 lOh @ 1350°C, WQ 23.7 ± 2.5 26.0 ± 1.6 0.475 ± 0.023

(FC = furnace cooled, WQ = water quenched)

The W-7.2Ni-2.4Cu alloy is representative of a typical commercial preparation sintered at the lower end of the permissible temperature range. This has resulted in irregularly-shaped tungsten grains with a relatively high level of contact between adjacent grains. By comparison, the W-5.2Ni-4.8Fe alloy, prepared at a higher sintering temperature, contains approximately spherical tungsten grains with a lower level of contiguity. The average particle size is comparable in both alloys, but the volume fraction of matrix is higher in the W-Ni-Fe alloy, and may be attributed to the greater solubility of tungsten in the Ni-Fe matrix than in the Ni-Cu matrix. High- temperature solution treatments do not appear to change significantly the 110 average tungsten particle size or the volume fraction of matrix, but do tend to increase the contiguity of the tungsten particle network in both alloy systems. An increase in Cw is observed with increasing time at a given temperature or with increase of solution treatment temperature for a given time.

Figure 2 shows typical transmission electron micrographs of the two alloys, particularly of the tungsten-matrix and tungsten-tungsten boundaries. These boundaries exhibit structural features characteristic of those in other materials, and at least for specimens furnace cooled from the sintering temperature or water quenched from a high solution treatment temperature i 1200°C), there is no evidence of any major microstructural change. The impact resistance of the alloys in these two states can be markedly different, however, and thereby warrants a more detailed examination.

Interfacial Segregation of Trace Elements

For both W-Ni-Fe and W-Ni-Cu alloys a post-sintering heat treatment produced a significant improvement in toughness; for example, the impact energy is approximately doubled (compared with an as-sintered specimen) by a solution treatment at 1350°C for 1 hour followed by a water quench. This behaviour was reversible; specimens toughened by solution treating and quenching could be re-embrittled by solution treating and furnace cooling.

Figure 3 illustrates a fracture surface typical of both alloys in the as- sintered furnace cooled condition. The fracture is predominantly interfacial in character and salient features are marked A, B and C on the micrograph; A corresponds to the tungsten side of a ti'ngsten-matrix interface, B to the matrix side of the same type of interface, and C to the impingement boundaries between tungsten particles.

Similar freshly exposed fracture surfaces were subjected to analysis by Auger electron spectroscopy using an instrument capable of measuring separately the surface composition of individual fracture facets. Figure 4 shows Auger spectra from the matrix side of the tungsten-matrix interface (i.e. area B in Fig.3), and of significance are. the strong phosphorus, and to a lesser extent sulphur concentrations segregated to this surface. These impurity elements derive from the elemental powders used in the manufacture of the heavy alloys, typical analyses of which are given in Table II, and are expected to lower the interface cohesion and thus

Fig.2 Micrographs (TEM) of W-5.2Ni-4.8Fe alloy," (a) tungsten-matrix interphase boundary, (b) tungsten-tungsten grain boundary. (Muddle) Ill

eruph (f=EMl of frac- Fig.4 Auger electron spectra 'f as—sintered from fracture surface of as- -\: I'.'- .•..J!li-2.4Cu sintered furnace-cooled W-5.2Ni-4.8Fe and W-7.2Ni-2.4Cu alloys. (Lea and Muddle) '-•.::, . Lnt •_-; t s! ingly , on the corresponding tungsten side of •<:-:'.:*. Lri x interface (area A), only a small amount of phosphorus ;, ;iiiii similarly, very little phosphorus was evident on the ••-,,•'-• \ .-p. interfaces (area C). This association of phosphorus K-kel-rr.ai,rix side of the fractured tungsten-matrix interface can 1 •• - i • II- L.. .i->m.-•:is "i-ated by rastering the incident electron beam across the

'irfi.-t ctv i selecting the energy of an Auger electron peak respesentative : /l I. .1'' ' . -ui rrr --lenient to form an image. Figure 5 illustrates the simi— '.'t-n. r.h...» phosphorus and nickel images so formed.

Table II Analysis of elemental powders Composition, wt. i uv,ier Ca Fe O(ppm) -,! i 0 .010 0.004 0.002 0.011 0.03 800 irbur.v L • 0 .096 0.011 0.003 30 - 0.026 0.004 •irbui" v i I 0 .012 0.002 0.002 99.93 400

lion i-CLd-r.cnt followed by quenching was found to reduce significantly

'. I.t p'r.osohorus concentration at the interphase interfaces, but not the -? 'J .. I. hur concentration, and also produced some significant changes in the f rue •'ire mode, as shown by Fig.6. Although interface separation remains inr- ,najor component of failure (measured to be > 50pct of the fracture = ii f if"1] 'in improvement of interfacial cohesion is evident from the areas :.;*ilf- matrix failure (Di and an increase in the fraction of tungsten is railing by cleavage (E). Thus, it seems that phosphorus is mainly aside for the measured reversible embrittlement by trace element

wl ly, Auger electron spectra from the facets resulting from tungsten-tungsten boundary fracture (area C in Fig.3) showed segregation of nic!•-..•= I to levels as high as one monolayer in W-Ni-Cu alloys, but less in ,v'-:!i-FV. In the W-Ni-Cu alloy segregation of copper to the tungsten- matrix ':-,'jundory during heat treatment was also detected, which would be 112

A-'

Fig.5 Augergraphs of the elements P and Ni on a fracture surface of W-7.2Ni-2.4Cu alloy. (Lea and Muddle) expected to lower boundary cohesion. It is known that segregation of nickel will embrittle polycrystalline tungsten10 and so this and the possibility of copper segregation could be expected to contribute to the lower impact toughness levels generally found in the W-Ni-Cu alloy system.

Precipitation Reactions

Interface Precipitation

Figure 7 illustrates the fracture surface appearance of a commercially- supplied W-5Ni-5Fe material (of unspecified processing conditions) with particularly low impact toughness. Fracture has occurred predominantly by failure of the microstructural interfaces; the tungsten-tungsten boundary areas are smooth and featureless, but the tungsten-matrix interface areas show evidence of facetting and secondary cracking, suggesting the existence of a brittle interfacial precipitate. In the present work it proved possible to reproduce this fracture surface appearance by first giving a solution treatment (^ 1350°C) and quench, and then ageing in the temperature range 750-850°C. However, there have been other reports of

Fig.6 Micrograph (SEM) of Fig.7 Micrograph (SEM) of fracture surface of W-7.2Ni-2.4Cu fracture surface of as- alloy solution treated 2h at sintered furnace-cooled 1350°C and water quenched. W-5Ni-5Fe alloy. (Jones) (Muddle) 113 similar behaviour, also credited to precipitation, after furnace cooling from the sintering temperature11"13.

By deep-etching this fracture surface and exposing it to X-ray analysis some evidence of another phase was obtained, which was tentatively thought could be an intermetallic compound of the form (Ni,Fe)W, isomorphous with the orthcrhombic intermetallic compound NiW (whicn occurs at approximately 75wt%W in the Ni-W binary system11*). Not all the X-ray lines expected for this phase were detected, however, although a low success rate might be expected from a rough fracture surface. According to the phase diagram, this phase might be expected to dissolve at temperatures 5 1000°C,and some improvement in impact toughness values and fracture surface appearance was obtained after solution treating and quenching from the range 1000-1200°C. Furthermore, Henig et al15 have claimed identification of a (Ni,Fe)W interfacial precipitate in W-7Ni-3Fe alloys heat treated at < 1000°C. Consequently, there is some evidence to support the possible ^occurrence of an interphase intermetallic compound in W-Ni-Fe alloys.

However, more detailed examination has revealed the existence of an alternative (or additional) precipitation reaction. Firstly, high-resolution Auger electron spectroscopy performed by Muddle16 on the precipitate exposed at the fracture surface shows a significant carbon peak. The shape of the peak, and its apparent stability to argon-ion sputtering, suggests that the carbon is in the form of a carbide rather than a surface deposit. Secondly, thin foil electron microscopy has allowed a closer examination of the precipitate structure (and morphology).

Figure 8 confirms the presence of a precipitate ^ 0.4ym thick at ths tungsten-matrix interface. Similar precipitates have been induced in alloys simulating the composition of The matrix in the W-5Ni-5Fe heavy alloy (38.5Ni-40Fe-21.5W). Convergent beam electron diffraction and X-ray powder analysis of these precipitates have indicated a diamond cubic structure (with lattice parameter, ao = 11A) and space group Fd3m (no.227)8>16, whilst electron microprobe analysis gave the W:Ni:Fe ratios as 50.2:20.6:29.2 (at%). An experiment using a laser induced ion mass analyser also detected an additional small carbon peak in the spectrum,

Fig.8 Micrograph (high-voltage Fig.9 Micrograph (TEM) of TEM) of W-5Ni-5Fe alloy solution W-7.5Ni-2.5Cu alloy solution treated lh at 1350°C, water- treated lh at 1350sC, water- quenched and aged lOOh at 850°C. quenched and aged lOOh at 850°C. (Posthill). (Muddle). 114 thought to be real and estimated to result from a precipitate carbon content of approximately 8at%." Both the ternary systems Ni-W-C and Fe-W-C 1B l9 contain n-oarbide phases of the forms Ni6WeC" and ?e6W6C » f respectively, both with Fd3m structure and ao = 10.9A. Consequently, it is con- eluded that the embrittling tungsten-matrix interfacial precipitate in W-5Ni-5Fe alloys is an n-carbide of composition (Ni,Fe)6W6C, although the (simultaneous) presence of an intermetallic {Ni,Fe)W phase cannot be ruled out.

Although the W-7Ni-3Cu system has been subjected to much less metallographic examination, the observation of precipitation at the tungsten- matrix interfaces has been made by transmission electron microscopy, as shown by Fig.9. The precipitates have been observed after ageing for lOOh at 850°C specimens previously given a high temperature solution treatment followed by viuenching. The precipitates may occur singly, or as a group of apparently equiaxed grains. Insufficient analytical metallography has been carried out to identify this precipitate unambiguously at the present time. However, it is believed to have a similar embrittling effect to that found in the W-Ni-Fe system.

Figure 10 shows evidence of precipitation at a tungsten-tungsten boundary in the W-5Ni-5Fe alloy aged for lh at 1050°C. Convergent beam electron diffraction identified the structrue as fee (ao = 3.60A), and thin film microanalysis also demonstrated qualitatively high nickel and iron concentrations in the precipitate. The precipitate thus appears to be similar to the matrix-phase of the W-5Ni-5Fe alloy which is also fee (ao = 3.60&). Evidence for an irrational orientation relationship with either of the adjacent tungsten grains, lying in the region between Kurdjumov-Sachs20 and Nishiyama-Wassermann21,22, has also been ob'tained8. It is concluded that the precipitate forms to relieve the supersaturations cf nickel and iron in the tungsten grains resulting from liquid-phase sintering, which have been reported23 to be approximately O.lat%Ni and l.lat%Fe, respectively .

The raised flat areas representing cleavage along the tungsten-tungsten grain boundaries are readily evident on the fracture surface and normally

0-3pm

Fig.10 Micrograph (high-voltage Fig.11 Micrograph (SEM) of as- TEM) of as-sintered furnace- sintered furnace-cooled cooled W-5Ni-5Fe alloy aged lh W-3.5Ni-l.5Fe alloy aged lh at at 1050°C. (Posthill) 1050°C. (Hogwood) 115 have a smooth featureless appearance. However, precipitation on these boundaries is clearly apparent from the cleaved fracture surface, as illustrated in Fig.I?, which follows the work of Hogwood2", who has clearly demonstrated this effect in W-Ni-Fe alloys heat treated for various times at 1050°C. Other investigations25'26 have also interpreted similar features on the otherwise smooth fra ;ture facet as probably resulting from intergranular precipitation, although without experimental identification of the precipitate.

The appearance of these precipitate features on the tungsten-tungsten fracture surface facets has also been associated with improved ductility25'27. The tungsten-tungsten boundaries are an intrinsically weak link in the microstructure; it has been shown that the impact properties are generally reduced by an increase in contiguity of the structure (essentially an increase in the tungsten-tungsten grain boundary area) and also that segregation of nickel to these boundaries takes place, which should also lower their cohesion. Consequently, any reduction in the tungsten-tungsten grain boundary area, in this case by intergranular precipitation of a 'ductile' precipitate with potentially stronger interfaces, would be expected to result in increased alloy toughness.

Matrix Phase Precipitation (W-Ni-Fe alloy)

Ageing treatments in the approximate temperature range 750-850°C following a solution treatment and quench result in lamellar and Widmanstatten forms of precipitation in the matrix phase of W-5Ni-5Fe (Fig.12). A 'dendritic' precipitate formed in the matrix of W-Ni-Fe alloys slowly cooled from the sintering temperature has been previously reported12, but not unambiguously identified, whilst slow cooling was also shown to lead to a decrease in lattice parameter of the matrix phase and the appearance of a fine dispersion of tungsten precipitates26.

During the present study both forms of precipitation were reproduced in alloys made up to simulate the matrix composition, from which they could be electrolytically extracted for analysis by X-ray powder diffraction. From the X-ra> results, and from electron diffraction in both the 'matrix' alloys and the' heavy alloy itself, both the lamellar and Widmanstatten precipitates were identified a tungsten. It was also possible to

Fig,12 Micrographs (TEM) of W-5Ni-5Fe alloy solution treated lh at 1350°C, water quenched and aged 25h at 850°C; (a) lamellar precipitation and (b) Widmanstatten precipitation. (Muddle) 116

"i 5

•0.1 pm- Fig.13 Microhardness of the Fig.14 Micrograph (weak beam TEM) tungsten phase in as-sintered of the tungsten phase in as- furnace-cooled W-5Ni-5Fe alloy sintered furnace-cooled W—5Ni—5Fe rolled 10.1%RA and aged for In. alloy rolled 10.1%RA and aged for (Posthill) lh at 600°C. Arrows point to defect-dislocation interactions. (Posthill) associate formation of the lamellar morphology with migration of a high- angle grain boundary, suggesting a typical Type 1 discontinuous precipitation reaction from supersaturated matrix, e.g. Ysupers. + Yequl. + W. An independent observation of discontinuous precipitation of tungsten in the matrix phase of W-7Ni-3Fe heat-treated in the range 800-1000°C has also been reported recently15.

The present investigation did not attempt to examine the influence of matrix-phase decomposition on mechanical properties. Conflicting reports exist in the literature15>2B'z9, although it should be noted that the previously discussed precipitation reaction at the tungsten-matrix interface could occur concurrently, and result in simultaneous embrittlement of the alloy.

Tungsten Phase Precipitation (W-Ni-Fe alloy)

Ageing in the temperature rang,3 500-700°C for lOOh leads to an increase in the yield and ultimate tensile strength of the W-5Ni-5Fe heavy alloy. Microhardness measurements «n the tungsten particles reflect this increase in the tensile properties of'-the alloy, whilst the hardness of the matrix phase remains unaltered. Prior deformation of the alloy accelerates this age hardening reaction as «3hown in Fig.13 (as well as giving a work hardening increment). Examination of aged specimens by transmission electron microscopy, using the weak-beam technique to enhance diffraction contrast, identified a fine dispersion of elastic strain fields •** 20A in diameter (e.g. Fig.14). Analysis of the strain fields suggested that they would be caused by the formation of platelet precipitates on {100} tungsten planes, resulting in a tensile strain field in <100> tungsten directions. Although it has not yet proved possible to obtain analytical data from such a small area, it ±s proposed that the precipitates form to relieve the supersaturation of nickel and iron in the x-ungsten; it has already been shown that this can occur at much higher ageing temperatures by precipitation at the tungsten-tungsten grain boundaries. 117

CONCLUSIONS

The mechanical properties of tungsten heavy alloys are very dependent on the interfaces present in the microstructure. Segregation of phosphorus impurity to the tungsten-matrix interfaces results in embrittlement, but this can be reversed by a high-temperature solution treatment followed by a quench. Sulphur segregation to the same boundaries, with an expected loss in cohesion, was also observed. Segregation of nickel to tungsten- tungsten boundaries and the possibility of copper segregation to tungsten- matrix boundaries was similarly detected, and is expected to contribute towards generally lower toughness levels. The relative proportions of tungsten-matrix and tungsten-tungsten boundaries are also shown to vary according to processing and heat treatment conditions and should influence alloy properties. The formation of an n-carbide (Ni,Fe)eW6C and possibly an intermetallic compound (Ni,Fe)W, are shown to occur at the tungsten- matrix interfaces in W-5Ni-5Fe, depending on processing conditions and heat treatment, and can severely embrittle the alloy. Precipitation of an uni- dentified phase at the same interfaces in W-7Ni-3Cu alloys has also been found. Precipitation can also be induced at the tungsten-tungsten grain boundaries, but this phase is similar to the heavy alloy matrix binder- phase, and consequently is thougtt could enhance cohesion across these otherwise relatively weak boundaries. Matrix-phase decomposition was also shown to be possible, and resulted in the discontinuous precipitation of tungsten with a lamellar morphology, or in tungsten with an intragranular Widmanstatten morphology. Low-temperature ageing resulted in intragranular precipitation in the tungsten phase, and this age-hardening response could be accelerated by prior deformation.

ACKNOWLEDGEMENTS

I am grateful to Professor Sir Peter Hirsch FRS and Professor J W Christian FRS for the provision of laboratory facilities. I acknowledge and thank my colleagues B C Muddle, J B Posthill, P N Jones and M C Hogwood, both for stimulating discussion and permission to use some of their published and unpublished results. This research has been carried out with the support of the Procurement Executive, Ministry of Defence.

REFERENCES

1. R H Krock, Metals for the Space Age, Proc. Plansee Seminar, 1964, p.257. 2. D J Jones and P Munnery, Powder Metall., 10, 1967, 156. 3. D V Edmonds and P N Jones, Met.Trans., lOA, 1979, 289. 4. B C Muddle and D V Edmonds, Residuals, Additives and Materials Properties, Royal Society (London), 1980, p.129. 5. B C Muddle and D V Edmonds, Metal Sci., 17, 1983, 209. 6. C Lea, B C Muddle and D V Edmonds, Met.Trans., 14A, 1983, 667. 7. J B Posthill and D V Edmonds, Intl.Conf. Phase Transformations in Solids, Crete, 1983. 8. J B Posthill, D Pnil Dissertation, University of Oxford, 1983. 9. J Gurland, Trans AIME, 212, 1958, 452. 10. M P Seah, Acta Met., 28, 1980, 955. 11. R V Minakova et al, Second Intl. Powder Metall.Conf. 2, 1966, 91. 12. RVMinakova et al, Poro.Metall., 65, 1968, 73. 118

13. Alfennappann, Report NO.R70/71, Chalmers University of Technology, Gothenburg, 1971. 14. J M Walsh and M J Donachie, Met.Trans., 4, 1973, 2854. 15. E T Henig, H Hofmann and G Petzow, Proc. 10th Plansee Seminar, 2, 1981, 335. 16. B C Muddle, Monash University, Melbourne, unpublished research. 17. M L Fiedler and H H Stadelmaier, Z Metallk., 66, 1975, 402. 18. Ya S Umanskii and N T Chebotarev, Izv.Akad.Nauk. SSSR, ^5, 1951, 24. 19. J Leciejewicz, J Less-Common Metals, 7, 1964, 318. 20. G Kurdjumov and G Sachs, Z.Physik, 64, 1930, 325. 21. Z Nishiyama, Sci.Rep.Tohoku Univ., 23, 1934, 637. 22. G Wassermann, Arch.Eisenhuttenw., 6, 1933, 347. 23. L Ekbom, Scand.J.Metall., 5, 1976, 179. 24. M C Hogwood, R.A.R.D.E., Sevenoaks, unpublished research. 25. D G Brandon, E Ariel and J Barta, Proc. 5th Intl.Symp,Electron Microscopy and Strength of Materials, ed. G Thomas, 1972, p.849. 26. P A Verkhovodov et al, Poro.Metall., 196, 1979, 8. 27. R Gero and D Chaiat, Materials Engineering Conf., Technion, Haifa, Israel, 1981, p.46. 28. S S Kiparisov et al, Nauchn.Tr.Vses. Naucho-Issled.Proektn.Inst.Tugo, Met.Tverd.Splanov., 16, 1976, 280. 29. L AKonyukhova, Yu A Eiduk and L S Vodop' yanova, Tsvetny Metally., No.10, 1974, 57. 119

INVESTIGATION OF THE CREEP FAILURE MECHANISM IN THE MO - 5% W ALLOY

E. Freund, D. Agronov and A. Rosen Department of Material Engineering Technion I.T.T. Haifa

INTRODUCTION

Previous, investigations have established the fact that creep failure for most metals and alloys is a result of grain boundary cavitation (1,2,3). In an earlier paper (4) by one of the authors it was reported that near the fractured surface of a Molybdenum alloy specimen which failed in creep at high temperatures, almost every grain boundary which was normal to the direction of the applied stress contained a crack. Moreover it was shown in the same paper that thermo-mechanical treatments which effect the size and shape of the grains, also influnce creep ductility. For example a slightly deformed specimen by swaging contains large almost equi-axed grains and therefore a high fraction of the boundaries are normal to the applied stress, while a heavily swaged specimen is composed of very long and thin grains where most of the grain boundaries are parallel to the direction of the applied stress. The first specimen failed in creep after a few percent of elongation, while the latter specimen crept to approximately 50 percent strain before fracture.

The aim of the recent investigation was to understand the kinetics of crack nucleation and growth. Since the investigation is not yet completed, we are going to report here only the work done on a slightly swaged specimen.

EXPERIMENTAL

All the experiments were carried out on a Mo-5%W alloy which was manufactured by powder compacting, sintering and finally swaging. The reduction by swaging was 18% and the resulting grain diameter was 77 micron, with an aspect ratio of 3,13. Creep experiments were performed in a specially constructed creep apparatus, in dry Hydrogen, under an applied stress of 50 MPa and at a temperature of 1200 deg. centigrade. Three creep experiments were carried out: (1) up to the end of the primary region, (2) up to the end of the steady state region, (3) up to fracture. The specimens were removed from the creep apparatus and i^ere prepared for microscopic examination.

Small and sharp notches were machined within the grip section of the specimens. It was found that when the specimens were broken at -10 deg. centigrade at the location of the notches a large percentage of the fracture surface is intercrystalline. From one specimen approximately 10 fractured surfaces were created. The broken specimens were examined in a scanning-transmission electron microscope (STEM) using the scanning mode. It was found that even as-received specimens contain very small microvoids having various shapes and locations, there is a large variation in 120

the early stages of creep and then increases in the tertial region. There can be possible reasons for the increasing void density: 1) new cavities are created and 2) all the cavities grow, now even the very small ones which were not seen before are large enough to be detected.

The area of the average cavity multiplied by the density gives the area fraction of cavities. This value is shown in Fig.6 as a function of creep time.

DISCUSSION As mentioned in the introduction, the fact that creep failure is caused by grain boundary cavitation during creep is well known, however the kinetics of cavitation is still not completely understood. Even from the few experimental data reported here the following trends can be summarized:

1) Microvoids exist in the as-received material. 2) The microvoids are seen mostly in grain boundaries. 3) During the first and second stages of creep neither the density nor the average size of the cavities change significantly. 4) The last stage of creep is accompanied and probably caused by very rapid cavity growth. It would be difficult to draw final conclusions from partial results and therefore the above listed four points of conclusion should be treated only as indications.

The investigation is now extended to the following directions: The role of temperature and stress on cavity growth and the effect of swaging on the phenomenon. Unfortunately, the study require? time consuming and rather expensive use of scanning electron microscopy and therefore we are seeking the critical experiment which will give us the solution to the problem. ACKNOWLEDGEMENT All specimens were manufactured, prepared and heat treated by Metallwerk Plansee, GmbH, Reutte, Tirol. This work is sponsored by Metallwerk Plansee, GmbH and their support is highly appreciated. REFERENCES 1) Hull, D., Rimmer, D.E., The growth of grain boundary voids under stress, Philosophical Mag. 1959, 4,673. 2) Muller, D., Langdon, T., Independent and sequential cavity growth mechanism. Scripta Metallurgica, 1980, 14,pp J43-148 3) Svenson, L.E., Dunlop, G.H. Growth of intergranular creep cavities. International Metals Review, 1981, 2, pp 109-131. 4) Bendersky, L., Rosen, A. The effect of Thermo-mechanica1 Processing on the creep ductility of Mo-5%W alloy. Material Science and Engi- neering, Vol. 62, no. 2, pp 211-216. 5) Konig, G., Blum, W. Comparison between cell structure produced In aluminium by cycling and monotonic creep. Acta Metallurgica, 1980, 28, pp 519-537. 121

Fig.l. As received specimen. Fig.2. Fractured specimen.

OS 10 1.5 20 OIAHETER OF CAVITIES I pm I

Fig.3. Cumulative size distribution Fig.4. Average void diameter of diameter of cavities. versus time of creep.

S-ISO-

% - 'no

20 10 60 30 00 IZO 20 40 60 80 "CO 120 TIME OF CREEP I HOURS) TIME OF CREEP (hsursl

Fig,5. Mean number of cavities Fig.6. Area fraction of versus time of creep. cavities versus time of creep. 122

ISOCHRONOUS CREEP OF COPPER-BERYLLiUM-NICKEL ALLOY (Cu-O.^Be-2-ONi) SOLUTION TREATED AND AGED

Nahum Ni r Rafael - A.D.A. Haifa

ABSTRACT

Short-time, constant-load uniaxial creep tests were performed on a high intensity electromagnet coil candidate [1] alloy Cu-0.4Be-2.0Ni in the aged [2] 1750HT condition". The test temperatures varied from room temperature to 250°C. The stresses used at each temperature were in the 0.2% yield strength range. These tests were designed to simulate periodic short-time loading conditions of the coil turns.Isochronous strength (stress-strain) curves were generated from these creep test results for time scales of 0.5sec to 8.0sec and compared to the tensile (standard) stress-strain curves of the alloy at these temperatures. It was found that at each temperature the stresses in all stress-strain curves (standard one as well as isochronous) were within

EXPERIMENT AND ANALYSIS Test Procedure Short-time, cyclic creep tests (not reaching steady-state) were carried-out using MTS tensile test machine and furnace. Most data was obtained at 150°C. However, a few trend-check tests were also caried out at room temperature, 200cC, and 250°C. Tests were load controlled, monitoring strains with an extensometer and loads with a load cell. Stress levels were reached in 0.5sec, kept constant for lOsec of transient creep, and unloaded in 0.5sec. Uniaxial specimens were used after being heat reated. Each loading cycle was repeated 10 time (until having an apparently repeating curve), with 2-3min between cycles. Typical creep curves at 150°C are shown in Fig. 1. Isochronous strength curves were generated from the creep data for creep times of 0.5 to 8.0sec, simulating the pulsed creep of the magnet coil during loading. Simulating coil working conditions, 150°C was chosen to be

solutionize and quench, k0% cold work and *tS2°C age. 123

the major testing temperature, at which several stress levels were applied. At the other temperatures (room temperature, 200°, and 250DC) only two stress levels were tested to check the trend due to change of temperature.

Results There was little difference between the isochronous strength curves in the range of 0.5sec to 8.0sec. This can be seen from the two curves, solid line and dashed line, on Fig. 2. These curves were generated from the first cycle of each test and are the boundaries of the time scale mentioned. It was also found that the same type of curves, when generated after a large number of creep cycles, resulted in an even narrower difference; and all match, within a reasonable accuracy, the lower isochronous curve (of the 0.8sec creep time of the first cycle). It can also be seen, from Fig. 2, that the scatter band of the isochronous curves is not sensitive to temperature change in the relevant range. Also, the higher the temperature, the lower the isochronous strength curve. It is to be noted that both the yield stress (ay) and the elastic modulus (E) reduce as the temperature increases. Moreover, in our case, the amount of their change with temperature and different batches is roughly the same [3] [E(T)/CTY(7)=16Q±10]. Therefore, normalizing the stress with

Analysis The unique normalized stress versus strain behavior of the Cu-0.4Be-2.0Ni alloy at 1750HT condition was also plotted on a log-log scale using the minimum vahiu of the scatterband of the curve in Fig. 3- These values, together with minimum values of the lower elastic range of our measured stress versus strain for the alloy, generated a more complete, yet conservative curve from0.\% total strain to strains of about 8% (and more!)- This curve is shown in Fig. 4. It is to be noted that the proportional limit in all the measured stress-strain curves varied from 0.72av to 0.76ciy. In our conservative curve (Fig. 4) it was taken to be 0.75tfy for best curve fit. The temperature and batch independent normalized stress versus strain curve, shown in Fig. 4, is divided into three parts: 124

1. Linear elastic range-up to 0.75oy(T) (which is the normalized proportional limit). 2. Transition range from the proportional limit up to about o/ay(T)=l (which is also of the order of about \% total strain). 3. Uniform plastic deformation range, starting at the end of the transition range (at about 1% total strain).

The uniform plastic deformation range, can be described by the following equation:

n —°—=KE , (1) ay(T) where a/oy(T) is the applied stress normalized by the 0.?% yield stress, which is a function of temperature. The total strain at the particular stress level is shown by e. The strain hardening exponent n is found to be 0.038 and the strength coefficient K is found to be 1.208 for Fig. *». In the transition range, the curve deviates from Eq. (l). This stress deviation is plotted on a log-log scale, against the total strain, and shows a straight line with a slope of -k.M as shown in Fig. 5- That transition deviation is described as sol lows:

(2) where A[a/ay(T)] is the deviation of normalized stress from Eq. (l), e has the same definition as in Eq. (1) and the constants are measured from Fig. 5 to be m=-k.\2 and A=7.5*10"11. In reality, the whole non-elastic range of the curve (starting at the end of the proportional limit) can be described by subtracting Eq. (2) from Eq. (1), to obtain one equation for that range, as follows:

°y(T)

The linear elastic range described here as

°y(T) a " where E(j)/ay(T) is here taken to be 150, which is the minimum measured value in the range of 150 to 170 measured for the different batches and at different temperatures. 125

CONCLUSIONS The uniaxial isochronous creep strength of the Cu-O.'iBe-Z.ONi alloy, 1750HT condition, was measured by simulating high intensity electromagnetic coil cyclic loading conditions, with hold times of up to 8sec, and temperatures of up to 250°C. There was almost no effect of the creep on the strength performance of the alloy that usually is given by a standard tensile stress-strain curve. By normalizing the applied stresses to the Q.2% yield stress at the given temperature, a unique curve was found and mathematically expressed that describes the complete stress-strain relationship of the alloy. That curve is independent of both temperature and batch of material within the working temperature range and specification of material. Its use is recommended in calculating both short-time creep as well as regular stress- strain behavior since it was generated also from available stress-strain data of that alloy at the 1750HT condition.

ACKNOWLEDGMENT The auther gratefully acknowledges the assistance of Mr. R. Akin as well as the discussions with Mr. S.N. Rosenwasser, Mr. R.D. Stevenson and Miss J.E. McGregor. The permission to publish these results granted by the Inesco Inc. management is also acknowledged.

REFERENCES [1] The Riggatorn T'M" Tokamak Design, Inesco Inc. 11077 N- Torrey Pines Rd. La Jolla, Ca 92037- [2l Heat Treatment Routin, Developed by Inesco Inc. (1981) . [31 Inesco Inc. Metals Data File, (I98I).

IJL

2-3 MINUTES. BETWEEN CYCLES

0.004 10 (Sec) - 113.3(K5l) STRAIN

Fig. I: Representative Cyclic Creep Raw Data. 126

120 •% a-- 110 40.6.1C 100

90 SOCHROKQUS 03 (sec) B.0 (Stc) SO - 22 (c-JtR.T.) S 70 a 150 (rt 0 CO TIM KUXC WAS • 150 (y tt5 (si.) _ 200 (C1) 50 NOT COvEftFD BV e a SHORT* Tine CR££P ISOCcl 8.0 (»c) 250 (rf ® 40 a FOR ISOCHRONOUS 30 STOEMTH CURVES 20

10 0 i 4 i ft IS fc 14 15 16 l> 18 19 20

Fig. 2: Isochronous Strength Curves.

1.05

100

REGULAR STRES -STRAIN SATf: — BATCH * ottom 0S601 -.55 ROOM TCMP A 150 (y i 200 fcl 0 ROOM TEMP. 02% VIELO STRESS 113J-IM I19-1IS< / T SOCHRONOUS 6.5 (5..18.0 (Sic) ROOT TEMP. 9 a B KOfclO 0 •.80 200(1)0 a 250ft) S s -7J (*) All 5p«iIm«M en from l»t(i<^ 0440)73

STRAIN

Fig. 3- Normalized Stress vs. Strain Data from Both Regular Tensile Tests and Isochronous Creep. 127

Fig. 4: Normalized Stress vs. Strain Characteristics of Cu-0.4Be-2.0Ni Alloy, 175OHT Condition, Independent of Temperature and Batch.

03-

QD3-

002

3 « 5 7 10

STRAIN

Fig. 5: The Transition Range Characteristics Shown by the Normalized Stress Deviation from the Plastic Range Equation Shown in Fig. k. 128

THE HOT TEARING OF LEAD ALLOYS

P. Ari-Gur* and F. Weinberg

Department of Metallurgical Engineering, University of British-Columbia, Vancouver B.C., Canada

The minimum temperature for which hot tearing occurs, is related to the ductile-brittle transition temperature, TQB- Measurements of TQB have been made in as cast binary Pb alloys containing Sb,Bi,Sn,Ma and Ca. The difference between TDB and the solidus temperature was found to be greatest for the PbSb alloys and least for the PbNa and PbCa alloys. Homogenization markedly changed TQB in the PbSb alloys, indicating that both non-equilibrium and equilibrium solute segregation at the grain boundaries contributed to hot tearing. For a certain maximum stress near TnBr the PbSb and PbCa alloys were much less suseptible to hot tearing than PbSb.

INTRODUCTION In one of the containment procedures being considered for the disposal of irradiated CANDU fuel, the fuel bundles would be sealed in a corrosion resistant metal container filled with lead1. It is important that the lead cast around the fuel bundles will be free of voids or cracks. During the solidification process most impurities and solutes are concentrated at the grain boundaries, where the final liquid freezes, lowering the solidification temperature of this liquid. If a tensile strain develops across the boundary with liquid present, the boundary will open, producing a hot tear. At present the concentration of solute at the grain boundary cannot be established quantitatively in a given system, nor can the final solidification temperature of the liquid at the grain boundary2. In some system it is possible to calculate the local strains due \o thermal contraction during solidification using heat transfer models3.

The present investigation was undertaken to measure the temperature at which the ductile-brittle transition occurs i,t lead and lead alloys. The ductile-brittle transition temperature is a direct indication of the temperature at which the final liquid solidifies and therefore the local temperaturt- above which hot tearing can occur if the material is strained. Because of shortage in space, only a brief description is given, full details are given elsewhere k.

PROCEDURE The materials investigated were Pb with binary alloy additions of Sb,Sn,Bi,Ca,and Na. Test samples of the alloys listed in Table I were produced, in shapes as shown in Fig. 1.

•• Now with Armament Development Authority, Haifa, ISRAEL. 129

The mechanical tests were performed with a table type Instron machine. During the tests, specimens were kept hot in a well stirred oil bath. After the failure the cross sectional area at the failure was measured. The appearance of the fractured surface was examined with SEM.

RESULTS The variation of the reduction in crossectional area with test temperature for the Pb alloys is presented in Fig. 2. Values of JQQ from Fig. 2 are listed in Table I, as well as the solidus temeprature T solidus and the difference AT between TQB and T^jj^g. The values of AT are large for the PbSb alloys, negligible for the PbNa and PbCa alloys. Increased solute concentration results in an increase in AT, with the exception of PbNa alloy. For the Pb \% Sb alloy, TQB is at the eutectic temperature. Homogenization reduces AT. Alloy additives significantly increase the strength and decrease the creep rate of pure lead. cruTS ^or ductile samples (close to TDB), are listed in Table I. Na additions have no effect on crm-_, whereas alloying with 1% Sb gives the largest increase from 1.17 to 3.45 MPa.

CONCLUSIONS Solute additions to lead decrease the ductile-brittle transition temperature TQB near the solidus. The temperature interval &T between the solidus and T[JB, during which hot tearing can occur, increases with increasing solute content. The interval varies appreciably for the different alloys examined. Homogenizing reduces AT indicating that hot tearing is associated with both equilibrium and non-equilibrium solute segragation at the grain boundaries. The maximum stress in tension of a given lead alloy is not directly related to AT. Accordingly a given maximum stress level can be attained with either little probability for hot tearing (PbNa and PbCa) or with a high probability (PbSb).

REFERENCES [1] D.J. Cameron, J.L. Crosthwaite and K. Nuttall; Can. Met. Quarterly, 1983 - in press. [2] F. Weinberg; Progress in Materials Science 1980, Vol. 30, PP. 295-328. [31 A. Grill, J.K. Brimacombe and F. Weinberg; Ironmaking and Steelmaking, 1976, No. 1. [4] P. Ari-Gur and F. Weinberg; Can. Met. Quarterly, 1983 - in press. TABLE,: DATA FOR LEAD ALLOYS. Sn Na LEAD 0.02 0.03 ! 2 0.25 0.75 0 wt* 0.5

06 „, 15 M " 1 55 1.17

Maximum 1.85 Stress MPa

homogenized sample- 131

:••'•••'•:„*!'/;.<",•'/.••/

r Grip -Pin O L • Sample

Fig. 1: Gripping system for tensile tests.

Fig. 2: The change in percent reduction in area with test temperature for the Pb alloys shown (weight % concentration). 132

-o-o-o

PbO 73N3

325 310 315 320 320 325 310 315 320 Temperature °C

Fig. 2 (b)

a <

C60

- PbO 02 Ca PbO 03 Ca PbO5SbH

I 1 330 320 325 330 31U 315 Temperature C

Fig. 2 (c) 133

THE EMBRITTLEMENT OF STEELS BY LOW MELTING POINT METALS Dr. Norman N. Breyer Illinois Institute of Technology In the late 1930's it was found that lead in small quantities, about 1/4% by weight, added during the teeming or ingots would improve the machinability of steel. It was determined, after extensive testing, that the mechanical properties of low carbon unalloyed steels were unaffected by the lead addition1. Improved production rates using automatic and semiautomatic screw machines resulted in lower unit part costs. The introduction of lead to alloy steels had been initiated in the post World War II period with the machining experience showing similar cost improvements. Unfortunately, the leaded alloy steels occasionally failed, both during processing and during subsequent service^>3. In some cases, the unexpected failures were catastrophic and occurred after extensive and expensive machining operations to produce the part. Figure lisa complex qear from which the teeth on one of the helical iears spalled during the heat treatment. The same type of violent failure can also occur in service under similar stress- temperature conditions. Several service failures which have come to light recently include jet engine failures and extensive cracking of dies in forging production operations^.

Extensive research has shown that the reason for such failures lies in the action of lead as a "liquid" metal embrittler in high strenath steels, a phenomenon which has been found to occur in many other metal- metal systems. The embrittlement, which can start, in some cases, 110°C (200°F) below the melting point, will exhibit a trough with a minimum just above or at the melting point has been called metal induced embrittlement (MIE), with that portion occurring up to the melting point labeled solid metal induced embrittlement, SMIE^. At temperatures above the melting point the liquid metal induced embrittlement (LMIE) is a maximum and eventually will decrease with a return to ductility, the so-called recovery temperature. (See Figure 2) Extensive testing has revealed that the steel can fail in a brittle manner if three conditions coexist: a) The presence of lead. b) Tensile loading. c) Temperatures between 200°C (400°F) and 480°C (900°F). The lead need not be present internally in the steel to cause the embrittlement; externally applied lead has also been found to cause brittle failures. Many other low melting metals, including Cd, Zn, Sn and In, will embrittle steel below the melting point; all of those tested revealing a rapid decrease in RA to a minimum at the melting point^. Additions of tin, antimony or zinc as dilute second elements in the lead and bismuth or antimony in the tin markedly increase the extent of embrittlement. As an example, the alloying with tin of the externally applied lead results in increasing severity of embrittlement. (See Figure 3) There was not a consistent trend in the SMIE or onset region. On the other hand, a systematic change in the brittle-to- ductile recovery temperature was found to accompany changes in the composition, with increasing tin alloyed with the lead progressively increasing the transition temperature. 134

Although the embrittlement is a kinetic phenomenon with the rate of arrival of the low melting metal atoms to the crack tip controlling the extent of loss in ductility, the effect of changing the testing strain rate is not consistent with a strict rate of arrival effect. Figure 4 shows the results of testing AISI 4145 leaded steel specimens over the temperature range at which an embrittlement trough was found. Using three orders of strain rate change it was found that the same general shape of the RA trough was produced with some inconsistent results. As would be expected, the shorter test times accompanying high strain rates shifted the onset (SMIE) region to higher temperatures, consistent with shorter exposure times. What was unexpected was the finding that high strain rates shifted the recovery temperature to higher values, a revelation inconsistent with shorter exposure times. (The complicating effects of strain rate on the properties of the base metal, steel, should play a part in the eventual answer.) The same atoms which cause embrittlement had been found to be responsible for other embrittlement phenomena including tempered martensite embrittlement and temper embrittlement3. Cold working the steel surprisingly tends to alleviate sensitivity to the metal induced embrittlement (MIE)8. (See Figure 5) This is true in spite of the fact that, cold work increases the strength level of the steel, an effect which normally increases the sensitivity to embrittlement. The fracture characteristics reflect the increased resistance to fracture. REFERENCES 1. "Properties and Machinability of a Leaded Steel", T.J. Dolan and B.R. Price, Metals and Alloys, (January 1940) p. 20. 2. "The Effect on Lead on the Mechanical Properties of 4145 Steel", S. Mostovoy and N.N. Breyer, Trans. ASM, 61 (1968), pp. 219-232. 3. "Some Effects of Certain Trace Elements on the Properties of High Strength Steels", N. Breyer, Proceedings of the 31st Electric Furnace Conference, (1973), pp. 183-189. 4. "Lead Induced Brittle Failures of High Strength Steels", N. Breyer and P. Gordon, Proceeding of the Third International Conference on the Strength of Metals and Alloys, (August 1"73), Cambridge, England. 5. "Solid Metal-Induced Embrittlement of Steel", J.C. Lynn, W.R. Warke and P. Gordon, Material Science and Engineering, 18 (1975), pp. 51-62. 6. "Liquid Metal Embrittlement of 4145 Steel of Lead-Tin and Lead- Antimony Alloys", N.N. Breyer and K.L. Johnson, Journal of Testing and Evaluation, 2, No. 6, 1974, pp. 471-477. 7. "Environmental Sensitivity of Structural Metals: Some Dynamic Aspects of Liquid Metal Embrittlement", K.L. Johnson, N.N. Breyer and J.W. Dally, Proceedings of Conference; Environmental Degrada- tion of Engineering Materials, Virginia Polytechnic Institute, 1977, pp. 91-103. 8. "Effect of Cold Work on Liquid Metal Embrittlement by Lead Alloys on 4145 Steel", M, Watkins and K.J. Johnson and N.N. Breyer, IV Interamerican Conference on Materials Technology, June 29-July 4, 1975, Caracas, Venezuela. 135

Fig. 3. Reduction of area, RA, for 200-ksi (1379-M Pa) 4145 steel surface wetted with various Pb-Sn alloys as a function of test temperatures.

O H.T. 70 - A 10% C.W. O 20% C.W. 40 O 30% C.W. O 50% C.W.

40 s O 200 400 600 800 1000 TEST TEMPERATURE, °F o 130 Fig. 5. Ductility Properties of Internally Leaded 4145 3 20 Steel Processed to 200 ksi Nominal UTS by Heat Treat- ment Alone and by Heat Treatment Plus 10, 20, 10 30, and 50% Reductions by Die Drawing.

400 500 600 700 Temperoture, °F 136

Fig. 1. Complex Leaded Alloy Steel Helical Gear with Spalled Teeth Resulting from Heat Treatment.

Fig. 2. Schematic Representation of Ductility Trough Unembrittled ».

o G

\ Ductility \ trough / Recovery . -J— — * temperature Test Temperature

TEST TEMPERATURE, "C 100 200 300 400 500 120 I Fig. 4. Effect of Strain Rate ' ' KEY € IN/IN/WIN M/M/SEC on Ductility-Temperature .0025 .0000417 Profiles for Lead Steel. o: 80 ~- .025 .000417

.23 .00383 1 f 1 ui 2.46 .041 i _ ££40

'•' 1 1 t 200 400 600 800 1000 TEST TEMPERATURE, *F 137

THE EFFECT OF SUBSTRUCTURE ON CREEP PROPERTIES OF THE T Z M ALLOY

D. Agronov, E. Freund and A. Rosen Department of Materials Engineering Technion, Israel Institute of Technology, Haifa

INTRODUCTION

Previous investigation with Mo-5%W alloy has revealed that the creep strength is strongly dependent on the thermo-mechanical history of the alloy (1,2). It was found that swagging at high temperatures introduces a subgrain structure and the higher the percentage of reduction by swagging the greater is the average dislocation density in the subgrain boundaries and the higher is the creep strength of the alloy. Mo-5%W is a solid solution alloy and therefore the only possible hardening mechanism is the one which stems from dislocation rearrangements. The study reported here involves an entirely different moly alloy, TZM, which is strengthened by dispersion (3,4,5). It is well known that TZM is much stronger than the moly-tungsten alloy and its creep resistance is much higher (6). However, little is known about the effect of high temperature pre-deformation, such as swagging on its creep resistance. The purpose of this investigation was to find out whether the creep strength of this alloy is also affected by thermo-mechanical treatment and to compare the results with that of Mo-5%W obtained during the previous investigation.

EXPERIMENTAL

The specimens were made from powder compacted, sintered and swaged TZM, with the following chemical composition in weight percentage: 0.5% Ti, 0.07% Zr and 0.01-0.05% C. Five series of specimens were prepared by Metallwerk Plansee. The difference between the series was in the final reduction by swagging after the last recrystallisation treatment. The final reductions were 18.6, 36.0, 50.2 and 75.0 percent. The fifth series was received in the recrystallised condition. The complete production schedule of the specimens as well as the various grain sizes in the longitudinal and transverse direction are given in detail elsewhere (7). It is to be noted that the grain size of TZM is considerably smpller than the grain size of Mo-5%W. All creep experiments were carried out in a high temperature creep apparatus and the procedure of testing are explained in detail in (8). Tests were performed at the temperature range of 1000-1300°C and in the stress range of 70 - 450 MPa. Specimens for TEM were prepared from the as- received as well as from the crept specimens. Thin films were obtained from both transverse and longitudinal directions and were examined in a Jeoi lOOCX scanning transmission electron microscope.

RESULTS

Fig.1,2 and 3 exhibit the variation of the steady state creep rate (SSCR) with stress for various temperatures for the recrystallised, for the least 138

(18%) and for the most severely swaged conditions (75%), respectively. Similar behavior was observed for the other batches of specimens, i.e. swagged to 36% and 50% reduction. The main features of these diagrams are summarized below. There are enough data points at 1300°C to indicate that they can not be connected by a single straight line. Instead, two straight lines can well represent the results. The line which connects the high-stress results is very steep (high n-va.lue), while the line which connects the lower stress results has a much smaller slope. This tendency is very well known and was published by several investigators (9,10,11). It is difficult to determine the exact value of the slope for the lower stresses since the number of experiments is small in this range, The variation of the n-values with temperature or the amount of swagging is not systematic and therefore we assumed that the slope n is independent of these factors. The average slope n = 13.6 which is very high. On the other hand, the n value of dispersion hardened alloys is known to be very high (12,13). The practical implementation of the high n value is that the stress range of the measurable creep rates is narrow. This point will be discussed later in detail.

The recrystallised alloy has the lowest creep resistance at every temperature. The creep strength of the swaged alloy increases with increasing amount of swagsing. At low stresses however, where the creep rate is small the degree of reduction by swagging has almost no effect on SSCR. Similar behavior was observed for the Mo-5%W alloy (14). The subgrain size for all four types of specimens in the as-received condition is listed below. The values were obtained by measuring a large number of samples.

amount of reduction by swagging (%) 18.6 36 50.2 75

average subgrain diameter (ym) 0.74 0.88 1.03 0.76

The subgrain size 6Q is practically the same for all the specimens and its average value is 6O = 0.85 pm. The subgrain size of Mo-5%W was also independent of the amount of swagging, however &o of Mo-5%W was 2.1 p about three times larger than 6O of TZM. The Disorientation between the subgrains was measured by the Kikuchi line shift method, which was developed for the Mo-5%W alloy (8). Fig.4 represents the cumulative distribution function of the subgrain misorientation, 7(0)), versus the subgrain Disorientation angle, a). It is to be noted that the distribution shown in Fig.4 is very similar to that obtained for the Mo-5%W alloy,(8). For both alloys the percentage of small misorientation angles decreases with increasing reduction by swagging and the average misorientation angle increases. The distribution curves for the 211 and 212 TZM alloys (50 and 75% reduction) are practically identical. The distribution curves of the 111, 112 and 122 Mo-5%W alloy (50,75 and 85% xeduction) were also similar. This fact indicates that swagging ov-z 30% reduction in both alloys could not be very effective in improving the creep strength. 139

We have also started to measure the subgrain size and the distribution of misorientation angles after creep, however at the time of preparation of this report we have only partial results.

DISCUSSION

Since the study is not completed, we do not want to draw final conclusions. Nevertheless, the results which were obtained until now indicate that the creep mechanism in the dispersion hardened TZM and in the solution hardened Mo-5% alloy is probably the same. Accordingly the idea of effective subgrain size can be applied also for TZM. A detailed explanation of the theory is published elsewhere (8). One of the important factors, according to the theory, is the strengthening effect of the subgrain boundaries and the importance of the subgrain size. Undoubtedly, the very high creep strength of swaggedTZM compared to that of the swaged Mo-5%W alloy is due to its finer subgrain size. We calculated the factor of strengthening only due to the finer subgrain size in TZM and obtained a two order of magnitude increase in creep rate. On the other hand experimental results show a four order of magnitude increase in creep rate (1200°C, 140 MPa) (7). The discrepancy between the calculated and measured values is due to the contribution of dispersion or precipitation. We believe that when the investigation is completer1 it will be possible to evaluate this contribution quantitatively.

Another important observation of the recent study is the very high stress dependence of the SSCR in the high stress region. This phenomenon is typical for dxspersion hardened alloys and has been reported in the literature (12,13). For TZM the average value, independent of temperature and swagging conditions, was n = 13.6, while in Mo-5%W alloy n=8.5. As mentioned before, due to this very high slope the stress range of the measurable creep rate is small. For example, at 1300°C under the stress of 160 MPa, the 212 (50.2% reduction) specimen creeps at a rate of 10~^ hr-1, while under the stress of 270 MPa the creep rate is 10"1 hr~l. A 40% reduction of the load results in a three order of magnitude drop of the creep rate. A further reduction of the load causes creep at extremely slow rates, in any case below the sensitivity of our measurements . The investigation is now conducted in the following directions: (i) the effect of creep deformation on the distribution of misorientation angles (a>), (ii) quantitative measurements of the free dislocation length in subgrain boundaries (A) and the relationship between oj and X and finally (iii) the interaction between dislocations and dispersoids or precipitates during creep.

ACKNOWLEDGEMENTS

All specimens were manufactured, prepared and heat treated by Metallwerk Plansee, GmbH, Reutte, Tirol. This work is sponsored by Metallwerk Plansee, GmbH, and their support is highly appreciated. 140

REFERENCES

1. Bendersky, L., Komem, Y., and Rosen, A., High Temperatures - High Pressures, 13, 511-520 (1981). 2. Bendersky, L., Rosen, A., and Mukherjee, A.K., Scr. Met., 16, 467-470 (1982). 3. Chang, W.H., Trans, of the ASM, 218, 254-256 (1960). 4. Chang, W.H., Trans, of the ASM, 56_, 107-124 (1963). 5. Wilcox, B.A., and Gilbert, A. Acta Met., l^, 601-606 (1967). 6. Eck, R., Planseeberichte fur pulvermetallurgie, 27^ 53-74 (1979). 7. Agronov, D., Bendersky, L., and Rosen, A. To be published in International Journal of Refractory & Hard Metals. 8. Bendersky, L., Dr. of Science Thesis, Haifa (1982). 9. Weertman, J., Trans, of the ASM, 6^, 681-694 (1968). 10. Sherby, O.D., Burke, P.M. Progress in Materials Science, Vol. 13, 340-350, (1967), Pergamon Press, Oxford. 11. Mukherjee, A.K., Bird, J.E., Dorn, J.E., Trans, of the ASM, 62^ 155-178 (1969). 12. Lagneborg, R., Bergman, B., Metal Science, January, 20-28 (1976). 13. Gibeling, J.C., Nix, W.D. Materials Science and Engineering, 45_, 123-135 (1980). 14. Bendersl.y, L., Rosen, A., and Mukherjee, A.K. Stress and Micro- structure Dependence of the Creep Resistance of Mo-5%W Alloy, in Strength of Metals and Alloys., ed. by R.C. Gifkins (ICSMA 6, Melbourne), 585-600 (1982). 141

Fig.l. Steady state creep rate Fig.2. Steady state creep rate versus stress for TZM in versus stress for TZM in swagged recrystallised condition. condition.

Fig.3. Steady state creep rate Fig.4. The cumulative distribution versus stress for TZM in swagged function F(o)) versus disorientation condition. angle u) for TZM in as received condition. 142

FRACTURE TOUGHNESS EVALUATION OF BRITTLE MATERIALS USING INDENTATION METHOD

Z. Nisenholz

A.D.A. HAIFA

A simple and economic technique for the evaluation of fracture toughness (Kjc) for brittle materials has been applied to several brittle materials.

The technique consisted of a two step procedure: (i) Introduction of a radial crack pattern into the surface by means of a Vickers indenter and (ii) measurement of the crack length under a microscope.

The materials tested by this method were: Alumina based ceramics, ZnS, WC-Co, and Zirconia-Ytria. The results were in good agreement with those obtained by other methods.

1. INTRODUCTION

During indentation with a sharp indenter, two orthogonal radial cracks evolve within the elastic/plastic field under the indenter. The field contributes 2 components to the driving force on the crack system: an elastic (reversible) and a residual (irreversible) component [1-6].

At the indentation surface the elastic stress is compresive while the residual stress is tensile and therefore the radial cracks grow to their final length due to the residual driving force as the ir.denter is unloaded. This force is suitably characterized by the residual stress intensity factor (3):

3/2 Kr=xrP/C (I)

r (2)

where

P - Indentation Load c - crack length A - material independent constant I|J - the characteristic half angle of indenter (7A° in Vickers indenter) E - Young's Modulus H - Vickers hardness 143

3- RESULTS

Table 1 summerises the results for the various materials, under different test conditions (load and environment).

Table 1: results under different test conditions.

Material Environment Indentation Load 1 ' l (N) (MPa-m*)

AD-85(A.R.) ai r 50 3.0+0.11 ii oil 50 3.2+0.40 ZnS(H.P.) air 2 0.56±0.06 II oil 2 0.62+0.02 II air 5 0.64±0.06 II oil 5 0.67+0.02 ai r 10 0.67+0.10 AD-85(H.T.) air 50 4.19±0.32 II air 70 4.22+0.22 II air 100 4.30+0.17

Average indentation toughness results for the materials which have been studied in this work, are listed in Table 2, with results obtained for these materials using other toughness measurements.

Table 2: Fracture Toughness Results —^_—_—____——~— Material E(GPa) H(GPa) K (indentation) K Tr lc (Other Method) (MPa-m*) (MPa-m*) i 1 AD-85(A.R.) 220 9 5 3.0 23 (D.C.B.)1 AD-85(H.T.) 220 11 0 4.2 -

Al20a-10^Ti02 300 17 1 3.46 - ZnS (H.P.) 95 1 8 0.6 - ZnS(C.V.D.) 102 1 9 1.04 1.0 (D.T.)2 WC-6£Co 686 17 5 12.3 12.0 (D.T.) ZrO2-4£Y2O3 250 12 6.0 5.6 (C.N.)3 1 144

At post indentation equilibrium conditions: C=C0 and K =K and we obtain the basic equation for evaluating material toughness:

2//2 (3) The average calibration factor according to Lawn et al [3] and Anstis et al [5] for Viclcers indenter is:

2/3 4 /2 A(cotip) =Kc(H/E) Co /P=0.015 The fracture toughness value can be calculated directly, knowing the indentation load (P) and the crack length (c): K=0.0l5(E/H)Vc3/2

2. EXPERIMENTAL

The fracture toughness of 5 different brittle materials was determined using the post indentation method.

1. Al203-(AD-85)

2. Ai2o3-iorno2 3- ZnS k. WC-63XO 5- 4 All specimens were shaped to a form with flat parallel surfaces. One of the surfaces was polished with diamond paste. The minimum sizes of specimen surface area and thickness were 7mm2 and 2mm respectively. 5 measurements or more were made on each specimen. Effect of atmosphere on crack growth was examined by comparison of results from specimens with and without protection. A drop of immersion oil was placed on the indentation site of the following specimen: AD-85 as recieved (A.R.), heat treated AD-85 (H.T.) and Hot Pressed (H.P.) ZnS. Part of the AD-85 specimens were studied after the following heat treatment: rapid cooling from T»50oC, stress relief at 800°C and reheating to 1150°C. This treatment was found to improve significantly their toughness.

Several different indentation loads were applied on these specimens. 145

1. Double Cantilever Beam. 2. Double Torsion. 3. Chevron Notch.

4. DISCUSSION

The oil immersed specimens produced slightly higher toughness results than those obtained from specimens which were exposed to air. This difference is insignificant considering the higher standard deviation of the results. The indentation toughness values seem to be independent of indentation load. This load independence was already confirmed by Anstis et al [5] for other materials. The similarity between the indentation toughness results and the results obtained by other toughness measurements as listed in table 2, leads to the conclusion that this relatively simple fracture toughness measurement method can and should replace the other methods for K. measuring of brittle materials. c

5. REFERENCES

[1] A.G. Evans, T.R. Wilshaw, Acta Metall, 24(10), 939 [2] M.W. Swain, J. Mater Sci, H_(12), 2345 1X976). [3l B.R. Lawn et al, J. Amer. Cer. Soc., 63(9-10), 574 (I980). [4] B.R. Lawn et al, J. Aust. Cer. Soc. j|Jl) 4, (I980). [51 G.R. Anstis et al, J. Amer. Cer. Soc. 64(9), 533 (1981). [6] D.B. Marshall, B.R. Lawn, J. Mater. Sci. 14(8), 2001 (1979), 146

FAILURE OF WELDED INCONEL-600 PIPE IN THE COOLING SYSTEMS OF A NUCLEAR REACTOR

G. Kohn, B. Herrmann, A. Stern, E. Rabinovitz, and S. Addess

Nuclear Research Centre - Negev, P.O.Box 9001 Beer-Sheva, ISRAEL

ABSTRACT

Serious leaks were detected in the inlet and outlet pipes of the heat exchanger in the primary cooling loop of the nuclear reactor at the NRCN. Non-destructive tests were conducted which included: ultrasonic testing, tests with dye penetrants and radiography. The flawed part was replaced and mechanical tests were performed on it. The crack areas of the Inconel 600 tube were examined using optical and scanning electron microscopy. Chemical analysis of both cracked and intact tubes were carried out. It is concluded that stress-corrosion cracking was the main cause of failure, while minor evidence of fatigue was encountered as well. Measures for the pre\'ention of similar failure in the future are suggested.

INTRODUCTION

Routine inspection of the cooling system of the nuclear reactor at NRCN detected serious leaks of heavy water both in the inlet and outlet pipes of the heat exchangers in the primary cooling loop. A schematic drawing of the defective area is shown in Figure 1. The upper part of the outlet pipe which contained the cracks was sectioned out and replaced by a new pipe and flange. The aim of this paper is to describe the tests carried out in order to locate the cracks, determine their cause and suggest measures to avoid their future occurrence.

EXPERIMENTAL

Upon dismantling the heat exchanger it was submitted to. a series of non destructive tests (NDT) which included ultrasonic testing, dye penetrants and radiography. Two cracks were detected on the outlet pipe: one was a circumferential crack parallel to weld No. 1 and at a distance of about 15 mm from it. The length of the crack was about 310 mm on the inside and about 5 mm on the outside of the pipe. A second crack about 20 mm long was detected near weld No. 2. Further destructive tests were conducted on the part of the pipe which was sectioned out and replaced by a new pipe. Chemical analysis of the pipe material has been carried out both on samples near the crack 147

and away from it. Mechanical testing was conducted both on specimens from areas near the cracked region and away from it. The latter ones were taken both from parts of the pipe containing a weld and parts without any weld. Tensile specimens were machined according to ASTM E-8 (1) and the tests were performed on an Instron machine. Curved bending specimens were machined from areas containing the longitudinal weld and from unwelded areas. Vickers micro-hardness tests have been carried out at a load of 10 gm on a part containing the longitudinal weld. The microstructure of the pipe material has been studied using both conventional optical microscopy and scanning electron microscopy (SEM). Fracture surfaces of the tensile specimens were compared to the surface of the cracked area and the different modes of fracture were established.

RESULTS

The results of the chemical analysis of the pipe material established that the composition of the steel was in accordance with regular specifications of Inconel 600 (Table 1) .

Table No. 1: Chemical Composition of Inconel 600

Sample Concentration in w/o Ni Cr Fe Mn Si Cu

Specification: >72.0 14-17 6-10 <1 <0.5 0.5 <0.08 At crack: 74.6 14.15 5.9 0.18 0.2 0.015 0.08 Away from crack: 76.5 14.8 6.3 0.20 0.2 0.009 0.08

Both tensile and hardness tests proved that the properties of the pipe away from the circular weld were as expected of Inconel 600 (Table 2), and the bending tests indicated that welding did not have an adverse effect on the pipe material.

Table No. 2: Tensile Properties of Inconel 600 Specimens

Specimen Yield Strength Tensile Strength Elongation 0.2% (kg/mm 2) (kg/mm 2) (%)

Unwelded 28.5 65.3 33.5 Unwelded 25.4 61.1 34.7 Welded 46.5 66.1 28.7 Welded 34.3 63.6 33.0

Typical values from tht literature are yield strength of 25.3 kg/mm 2 and Tensile Strength of 63.3 kg/mm 2. 148

Optical micrograph of the longitudinal deffect-free weld is shown in Figure 2 while Figure 3 shows a series of micrographs taken from the cracked area of the pipe. From the first picture it is clear that the weld area is surrounded by a small heat affected zone (HAZ) beyond which no grain growth can be observed. On the other hand, from the second picture it is evident that while welding did not affect much the grain size on the upper part of the weld (the heavy flange side), it had a strong effect on the lower part, containing the crack in which large grains are clearly seen. The intergranular nature of the crack is also evident in this picture. figure 4 shows a fracture surface of a deffect-free sample typical both to welded and unwelded areas. The mode of fracture which is dimple rupture is typical of ducticle material with transgranular tearing. Frac- ture surfaces of the cracked pipe showed a completely different mode of fracture. Figure 5 shows a SEM micrograph taken from the crack area. It is clearly seen that the mode of fracture is intergranular with little or no signs of tearing. Al rich layers (as proved by EDAX measurement), can be seen on the intergranular surfaces, while Cr rich particles containing up to 3 w/o sulfur are also evident on the grain boundaries. Some evidence of fatigue crack propagation was evident as seen in Figure 6.

DISCUSSION

The results of the tests carried out on the pipe material indicated that away from the circular weld the properties of the Inconel were as expected from a sound and ductile material. The micrographs presented show that the crack propagated in an intergranular mode from the inside of the pipe towards the outside. Intergranular fracture is the typical mode of failure of Inconel steels when submitted to stress corrosion (2-4). Stresses on the pipe in this cooling system originated probably from the weld and from external loads such as the weight of the pipe itself. The large difference in grain size between the two sides of the circular weld are probably due to the different thickness of the material on both sides of the weld which lead to different heat extraction capacity. The larger grains had Al rich layers on their boundaries which originated probably from alumina which is known to be dissolved in the water. It has been previously established that alumina can absorb large quantities of chlorides and increase their local concentration (5). This, plus the fact that the Cr rich particles found at the crack area were also rich with sulfur may indicate the presence of a corrosive environment leading to crack propagation.

CONCLUSIONS

It has been concluded that stress corrosion cracking was the main- cause of failure. Since Inconel 600 is not very susceptible to stress corrosion cracking and since the time to failure was very long it was recommended that the external tension loads on the pipe system be made as small as possible with maximum elimination of mechanical vibrations and distortion. Heat treating complete parts to eliminate internal stresses was not practical due to the large size of the components and was 149 therefore not recommended. It was also recommended to carry out periodical NDT to uncover cracks before they become critical.

REFERENCES ANSI/ASTM E8-77a, in 1977 Annual Book of ASTM Standards, part 10, American Society for Testing and Materials, Philadelphia, Pa,, 1977, p. 154. S. H. Bush and R. L. Dillon, in Steess Corrosion Cracking and Hydrogen Embrittlement of Iron Base Alloys, R. W. Staehle, J. Hochmann, R. D. McCright, and J. E. Slater, eds., NACE, Houston (1977), p. 61. J. R. Cels, Corrosion, 34, 198-209 (1978) J. Blanchet, H. Coriou, L. Grail, C. Mahieu, C. Otter, and G. Turluer, p. 1149 in Ref. 2. W. E. Berry, Corrosion in Nuclear Applications, John Wiley, New-York, 1971, p. 184.

flange

No. 1

weld No. 2

Fig. 1 Inconel-600 Outlet Pipe

Fig. 2 Micrograph Containing the Longitudinal Weld 150

Cb) Ca)

(c)

n Cβ)

Fig. 3 A Series of Micrographs from the Cracked Area O60) 151

Fig. 4 Fracture Surface (x625)

Ca)

(xl35) (xlllO) Fig. 5 Crack area (a] Intergranular fracture (b) Solid layers

Fig. 6 Crack Area Showing Fatigue Marks (xl670) 152

EFFECTS OF METALLURGICAL VARIABLES ON HYDROGEN EM3RITTLEMENT TN TYPES 3l6, 321 and STAINLESS STEELS

P. Rozenak and D. Ellezer

Department of Materials Engineering Ben-Gurion University of the Mepev, Beer Sheva, Israel

SUMMARY

Hydrogen embrittlement of 3l6, 321 and 3^7 types austenitic stainless steels has been studied "by charging thin tensile specimens with hydrogen through cathodic polarization.throughout this study we have compared solution annealed samples having various prior f.usteniiip /rra i • -sii.c with Rfinples given the additional sensitization treatment. The results shov that refinriu 'rains i~-r roves the reni •--tance to hydrogen cracking regardless of the failure node. The sensitized specimens were predominantly intergranular, while the annealed specimens show massive regions of microvoid coalescence producing ductile rupture. 3^7 type stainless steel is much more susceptible to hydrogen embrittlement than 321 type steel, and 316 type is the most resistant to hydrogen embrittlement. The practical implication of the experimental conclusions are discussed.

1. INTRODUCTION

The occurrence of hydrogen embrittlement in austenitic stainless steels is substantial with ductility losses, and by the appearance of nonductile fracture surfaces [1-6]. The deleterious effect on mechanical properties caused by hydrogen charging has been found to depend strongly on metallurgical factors [7-11].

The object of the study was to evaluate the effect of grain size and heat treatment on the hydrogen susceptibility of 3l6, 321 and 3^7 types austenitic stainless steels.

2. EXPERIMENTAL PROCEDURE

Commercial austenitic stainless steels of types 3l6, 321 and 3^7 having the compositions shown in Table I were used for the present study. Various austenitic grain-sizes, were obtained following various times at austenizing temperatures. The grains size as measured by ASTM E-112 method were 7, 9 and 11 ASTM in 3l6 stainless steel, 8, 9 and 13 ASTM in 321 stainless steel and 7, 9 and 13 ACTM in 3^7 stainless steel. 153

Table 1 Chemical composition of AISI type 3l6, 321 and 3^7 austenitic stainless steels

Element Amount ( wt .%) of element in the following stainless steels AISI AIEI AISI type 316 type 321 type 3^7

Cr 17.98 18.07 17.90 Hi 12.09 10.11 9.00 Mn 1.71 1.32 1.51 Si 0.879 1.033 1.025 C 0.05 0.072 0.061| Mo 2.08 _ Ti O.Ul _ Nb - - 0.6l

Some samples of each group were given a further senitization heat treatment at 65O C for 2.X hr. The samples were tensile tested at room temperature at an extension rate of 0.005 cm min" while undergoing cathodic polarization. The hydrogen charging cell contained IN H?S0, solution with 0.25 grl of NaAsO . A platinum counter electrode and a current density of 50 mA/cm were used. For comparison purposes, specimens were tensile tested in air at room temperature. After failure the fracture surfaces were examined with a scanning electron microscope (SEM). Microstructure from the various starting conditions were characterized using transmission electron microscopy (TEM).

3. RESULTS AND DISCUSSION

Sensitization for 2k hr. at 650 C produced discontinuous grain boundary carbides in 3l6 stainless steel (Fig. la) which varied in size from 100 2 to 1(00 2. Diffraction patterns from these particles were readily indexed as M__C,-. Exanina- tion of the microstructures of sensitized both 321 and 3^7 types revealed grain boundary carbides (M _C/-) and a dense distribution of (MC) carbides which precipitated on matrix dislocations (Fig. lb, c).

The results given in Fig. 2 and Table II demonstrate the influence of grains size and heat treatment on the mechanical properties of 316, 321 and 3^7 types austenitic stainless steels under cathodic polarization. A significant feature of the results is that decreasing prior-austenite grain-size increases the mechanical properties and the resistance of 3l6, 321 and 3^7 steels to hydrogen embrittlement. The dependence of the tensile properties on the grains size is 154

shown for annealed and sensitized 316, 321 and 3!'T types steels in Fig. 2. V it-Id s«.r.. r (fnt of samples with grain siz-. A5TM 9 were decreased ai i.ut 15f^, as compared to "! ii e grained size (ASTM 11, 13) in sensitized 31o, 321 and 3^7 ;./ne steels, but were unaffected whether the material wa;; ur.neai ed or sensitized in various grain-size groups, ii'.h •••_•; i. ••-•+, >o Mi- tensile prootrties, ujdvogaa h .. •- '<•"•.•;"/ IU-A-KV •! --ffect on the ultimate tensile strength. UTS or fine grained size (AGTM 11) decreased 35^ as coirmared to coarse grained (ACTM T) 3l6 type sensitized specimens and h2% reduction of elongation. As can be seen from Fig. 2 and Table Ii, the sensitized coarse grained 3^7 type steel wMch is the most susceot ible steel, resulted in 7855 reduction of elongation and 31* reduction in ultimate strength while the sensitized coarse grained 3?1 type steel resulted in 75? reduction of elongation and ?8% reduction of ultimate strength. Most resistant were sensitized coarse grained 316 type specimens that resulted in ^k% reduction of elongation and 2?% reduction of ultimate strength. However, the total reduction of elongation at fracture as compared to those of uncharged specimens is reduced about 60% in fine grained sairnles whether the material was sensitized or not.

Following testing, the fracture surface of each specimen was examined to determine the modes of failure. The fractures of the specimens tested while cathodically charged show considerable differences between heat treatments and prior austenite grain-size range. Ductile dimpled rupture wit}: raicro- void coalescence was the main feature mode in sensitized fine grained (ASTM 11) 3l6 type steel (Fig. 3(a)) with narrow inter granular zones of about 5 pra on both fracture surface sides. A completely brittle fracture was observed in sensitized coarse grained (ASTM 7) 3l6 type (Fig. 3(b), (c)). The initial hydrogen fracture is mostly intergranular with only small transgranular cleavage-like areas .

U. CONCLUSIONS

( l) Refined grain-size improves the resistance to hydrogen cracking regardless of the failure node.

(2) Examination of the mi crostructures of beth 321 and 3!t7 types using TEM revealed discontinuous grain boundary carbides (Mp_C^) and a dense distribution of ('1C; carbides with precipitates on matrix dislocation. It is suggested that the improved ductility in sensitized 3l6, 321 and 3^7 steels strongly depends on grain size and carbide morphology.

(3) Examination of the fracture surfaces of coarse grained specimens tested while cathodically charged shot.- considerable differences between the annealed and the sensitized specimens. 3l6 type sensitized specimens were nrodominantly intergranular, while the annealed specimens show massive region of microvoid coalpscencr T-,ror1uc i r.p, dunt""le runturn. 155

(It) 3I17 type stainless steel is much more susceptible to hydrogen embrittlement than 321 type steel, and 3l6 type is the most resistant to hydrogen embrittlement.

REFERENCES

1. M.L. Holzworth, Corrosion 26 (1969) 107. 2. M.B . Whiteman and A.R. Troiano, Corrosion 21 (1965) 53- 3. M.R. Louthan, Jr., Hydrogen in Metals, I.M. Bernstein and A.W. Thompson, eds. (A.S.M., Metals Park, Ohio, 197*0 53. k. A.W. Thompson, Hydrogen in Metals, I.M. Bernstein and A.W. Thompson, eds. (A.S.M., Metals Park, Ohio, 197*0 91- 5. H. Hanninen and T. Hakkarainen, Corrosion 36 (l980) hi. 6. D. Eliezer, D.G. Chakrapani, C. Altstetter and E.N. Pugh, Met. Trans. lOA (1979) 935- 7. C.L. Briant, Hydrogen Effects in Metals, I.M. Bernstein and A.W. Thompson, eds. (TMS-AIME, Pennsylvania, 1950) 527- 8. H. Hannir.en, ^. Hakkarp.ipen ;nd P. "enonen, Hydrogen Effects in Metals, Met. Soc. of AIME, I98O, p. 575. 9. I.M. Bernstein and A.'-/ . Thompson, els. v ''MS-AIME, Pennsylvania, 1980) S75. 10. C.L. Briant and A.M. Ritter, Scripta Met. 13 (1979) 177. 11. C.L. Briant, Met. Trans. PA (1978) 731.

TA3.LE II

Percentage reduction in tensile properties o** charged specimens compared with those of uncharged specimens for AISI type 3l6, 321 and 3*»7 stainless steels Type of Grain size Heat Reduction(#) in thefollowing stainless treatment- <• ensile properties steel Yield Ultimate Elongation strength tensile strength AISI 3l6 ASTM 11 Annealed k 11 5 5 AISI 316 ASTM 11 Sensitized 2 q 52 AISI 321 ASTM 13 Annealed 0 21 68 AISI 321 ASTM 13 Sensitized 5 16 57 AISI 347 ASTM 13 Annealed 0 16 67 AISI 31*7 A f-TH 13 Sensitized 5 18 63 AISI 316 ASTM Q Annealed 0 O 59 AISI 316 ASTM 9 Sensi tized 3 15 AISI 321 ASTM 9 Annealed 0 15 65 AISI 321 ASTM 9 Sensitized 10 20 57 AISI 31*7 ASTM 9 Annealed n 21 63 AISI 3l*7 ASTM 9 Sensitized 15 15 55 AISI 316 ASTM 7 Annealed 0 2 k6 AISI 316 ASTM 7 Sensitized 0 2? 5h AISI 321 ASTM 8 Annealed n 2U C>3 AISI 321 ASTM 8 Hens'tized •-) 2S 75 AISI 3't7 ASTM 7 Annealed c 23 58 AISI ASTM 7 Sens it i zed 0 31 78 156

Fig. 1. TEM micrographs showing carbides in the sensitized AISI type 316, 321 and 3^7 stainless steels: (a) grain boundary carbides in the sensitized AISI type 316 stainless steel; (l>) precipitation of carbides within the grains, on dislocations and at grain boundaries in the sensitized AISI type 321 stainless steel; (c) precipitation of carbides in the sensitized AISI type 3^7 stainless steel. (c)

ANNEALED --

1 SENSITIZED

Jib Asm lap "fiP i g 113 ASTH 11 1 ^ 9 »STi g *^^^*~- •"— • 4 — ^. V Alii 318 • ASTH T l r" — AISI 31B «— AlSt 347 A5TM AISI MT Aisi m

(a) DISPLACEMENT (nun) (b) DISPLACEMENT imml

Fig. 2. Engineering stress vs. displacement curves illustrat- ing a comparison between (a) annealed and (b) sensitized AISI type 316 (grain sizes, ASTM 7,9 and 11) (-.-), 321 (grain sizes, ASTM 8,9 and 13) (—) and 3U7 (graineizes, ASTM 7,9 and 13) (- - -) stainless steels tensile tested while undergoing cathoaic charging. 157

Fig. 3. Micrographs of fracture surface of AISI type 3l6 stainless steel, grain sizes(a) ASTM 11 and (b), (c) ASTM 7S tensile tested while undergoing cathodic charging after a sensitization heat treatment. 158

MARTENSITIC TRANSFORMATION IN 3O4L and 316L TYPES STAINLESS STEELS CATHODICALLY HYDROGEN CHARGED

E. Minkovitz and D. Eliezer

Department of Materials Engineering, Ben - Gurion University of the Negev, Beer-Sheva, Israel

ABSTRACT This paper reports a TEM study on the role of phase transitions at the crack tip in 304L and 316L Types Stainless Steels cathodically hydrogen charged in the absence of any eternally applied forces. The possible role of o-'and e martensite phases in the fracture mechanism is discussed. 1. INTRODUCTION Hydrogen induced martensite phase transformation in austenitic stainless steels have been extensively studied with respect to the relative stability of the austenite p(f.c.c) phase, which in fact, is characterized by its tendency to transform to martensite on cooling or during plastic deformation below the critical M_ temperature. The effect of hydrogen on the γ-phase stability is that hydrogen decreases the y-phase stability and muy induce transformation of the y-phase to a" and e martensite (1-4) . It has been shown that most steels that form or'-- martensite are quite susceptible to hydrogen cracking, however, this issue has become controversial due to the fact that embrittlement took place when no martensite has been formed (3). This paper, based partially on recent work (5,6,9), demonstrates TEM study on the formation and propogation of microcracks caused by hydrogen charging, in connection with (a';e ) martensite phases. The possible role of a'and e martensite phases in the fracture mechanism is discussed.

2. EXPERIMENTAL PROCEDURE The stuuies were carried out on 304L and 316L types austenitic stainless steels. The steels were of commercial grade, and were received in the form of sheets 0.2 mm thick. All of the samples used in these experiments were first solution annealed for 1 h at 1100°C and then water-quenched. Specimens suitable for electron microscopy were then prepared by electrolytic polishing at 65 V in a Tenupol polishing cell using 30 cm perchloric acid, 300 cm methanol and 520 cm br.canol solution at -18 C. In the electrolytic polishing process, an attack of the hole edge is normally unavoidable. Thus, in order to avrid un- certainties about the origin of the crack and the possibility of obser- ving cracks which were formed by the thinning process, only specimens having "perfect" hole edges (Fig. la) were selected to be suitable for the hydrogen charging process. After a TEM examination, where the hole edge was checked carefully to ensure that no deformation structure had 159

been induced during the thinning process the specimens were cathodically hydrogen charged and then were investigated again by TEM. The hydrogen charging was performed in a charging cell, at room temperature in^the absence of an external force in a IN H-SO. solution with 0.25 gl" of NaAsO- added as hydrogen recombination poison. A platinum counter electrode and a current density of 0.5 A cm were used. The charging time was 15 min. TEM analysis was carried out in a JEOL-200B electron microscope operating at 150 kV.

3. RESULTS AND DISCUSSION

A TEM micrograph in the "Mesh-Image" magnification showing the TEM specimen's hole after the thinning process is represented in Fig. la. Forming cathodic charging to the same specimen in the absence of any externally applied stresses for a few minutes, revealed mainly intergranular cracks and some transgranular cracks at the hole edge of the TEM specimen. Further charging and re-examination by TEM analysis revealed crack propagation. Fig. lb represents cracks at the hole.edge which were formed following the charging conditions of 0,5 Acm for 2 h. Dark field electron micrograph of a transgranular crack in 304L, following charging conditions of 0.5 Acm" for 15 min. is shown in Fig. 2a and 2b. The Selected Area Diffraction Pattern (SADP) at the crack tip and along the crack surfaces is represented in Fig. 2c. These findings have revealed Debye rings reflections from a fine grained bcc ft'-martensite and a fee )'-matrix. The bcc phase and the surrounding a'ly interface showed a completely fine grained polycrystalline structure, while the SADP taken at about 0.5 ^m from the a'/y interface showed single crystal fee y-phase, indicating that the formation of the polycrystalline structure for the p-phase was highly localized to that area. Further charging to the same specimen and re-examination by TEM studies have revealed crack propogation through the martensite phase in front of the crack tip. The bright field electron micrograph of a transgranular crack and the microstructure around the crack tip in 316L Type following the charging conditions of 0.5 Acm" for 15 min. is shown in Fig. 3a. SADP taken at the crack tip is represented in Fig. 4a and its schematic diagrams are given in Figs. 4b and 4c. These findings have revealed single crystal spot patterns of (fee) )>-austenite and (hep)t-martensite (Fig. 4b) where the v/e orientation relationships are

(lll)v II (0002)e and [01i]y || [1120]f (6) . The SAD at the crack tip also showed Debye rings characteristic of a very fine-grained polycrystalline structure, which were indexed as (bcc) a' reflections (Fig. 4c). Fig. 3b represents a dark field electron micrograph of the e-martensite plates, taken from a (OllO) reflection spot of e-martensite. The location of this phase is in the front of the crack tip. A dark dield image (Fig. 3c) using the (211) a' ring intensities showed that the a'-phase is located ahead of the crack tip within the c-phase. Further charging to the same specimen revealed crack propogation through the mixed area of a' and e

-martensite phases, however, mainly through the (hep) e-martensite (5). No evidence of the appearance of a'-martensite phase after hydrogen

charging was found within the graiis; in the thicker region of the specimen as previously reported (6). It has been shown that the resistance to hydrogen embrittlement is improved by increasing the μ-phase stability (7,8). Transformation of the )>-phase to the a'(bcc) phase in front of the crack tip has been shown in Type 304, fractured in 10 Pa 160 of H gas (7) or when Type 304L was cathodically hydrogen charged in the absence of any externally applied stresses (9), where the fracture occurs through the o'-phase. However, recently an HVEM observation of crack tips of Type 304 sheet specimens which were severly embrittled by hydrogen pre-charging and thr.n fractured in air have shown that the crack propagated mainly along the e-martensite and partly in the region having a mixed structure of a'and e-martensite phases (10). In the case of Type 316L, the significant observation of the present experiments is that the y to e -phase transition occurs in front of the crack tip, where the crack propagation occurs mainly through the e-phase. Furthermore, it has been reported (11) that Type 310 steel was embrittled by hydrogen and that the crack proceeded along the interface between austenite and the hydrogen induced e-phase. However, crack propogation in 304L Type, occurs through the a'-phase in front of the crack tip.

4. CONCLUSIONS

This paper reports a TEM study on the formation and propogation of microcracks caused by hydrogen charging in connection with (a';e) martensite phases. The presence of α-bcc martensite in front of the crack tip of both steels support the view that embrittlement of unstable steel can have the form of an auto-catalitic process. Crack initiation occurs through triaxial stresses caused by hydrogen penetration and evidently followed by the formation of a1bcc martensite when hydrogen egresses the specimen. The existence of the bcc phase in front of the crack tip escalates hydrogen entry and crack propogation which in turn aid martensite formation. However, in the case of 316L Type, it is now recognized that the presence of the e-phase may be quite important in the fracture mechanisms. Thus, it is suggested that increasing the f-phase stability increases the relative importance of the e-phase in the fracture mechanisms.

5. REFERENCES

1. J.M. Rigsbee and R.B. Benson, J. Mater. Sci. 12, 406 (1977). 2. D. Eliezer, D.G. Chakrapani, C.J. Alstetter and E.N. Pugh, Met. Trans. A, lOA, 935 (1979). 3. C.L. Briant, "Hydrogen Effects in Metals" I.M. Bernstein and A.W. Thompson, eds., The Metallurgical Society of AIME, New York (1981), 527. 4. N. Narita, C.J. Altstetter and H.K. Birnbaum, Met. Trans. A., 13A, 1355 (1982). 5. E. Minkovitz and D. Eliezer, Scripta Met. 16, 981 (1982). 6. E, Minkovitz, M. Talianker and D. Eliezer, J. Mater. Sci. 16, 3506 (1981) . 7. N. Narita and H.K. Birnbaum, Scripta Met. 14, 1355 (1980). 8. E. Minkovitz and D. Eliezer, J. Mater. Sci. 17, 3165 (1982). 9. E. Minkovitz and D. Eliezer, J. Mater. Sci. Letters 1, 192 (1982). 10. T. Nakayama and M. Takano, Corrosion-NACE 38, 1 (1982). 11. A. Inoue, Y. Hosoya and T. Masumoto, Trans. Iron Steel Inst. Japan 19, 170 (1979). 161

Fig. 1. TEM micrographs in the "Mesh-Image" magnification of the TEM specimen. (a) After the thinning process. (b) Cracks which were formed by cathodic charging at the hole edge of the TEM specimen.

Fig. 2. Dark field electron micrographs and Selected Area Diffraction Pattern (SADP) of 304L Stainless Steel, which was cathodically hydrogen charged. (a) Dark field, taken in (200) a' (bcc) ring. Note that the martensite phase is located along the crack surfaces. (b) Dark field, taken in (200") a1 (bcc) ring. Note that the martensite phase is located in front of the crack tip. (c) SADP taken along the crack surfaces and the crack tip, which was indexed as y(fcc) matrix and a'(bcc) rjartensite. 162

Fig. 3. TEM micrographs showing a transgranular crack and the microstructure around the crack tip of Type 316L which was cathodically hydrogen charged. (a) Bright field image. (b) Dark field image. Note that the e-martensite phase is located in front of the crack tip. (c) Dark field image. Note that the a'-martensite phase is located ahead of the crack tip within the e-plates.

200

042 0110 9

1011 110 1101 242 200 211 022 310 O [012] . . - z A v 123 [] 2.A.- hcpf bcc

Fig. 4 (a] Selected area diffraction patterns of Type 316L taksn at the crack tip shown in Fig. 3. (b) Schematic diagram of electron diffraction pattern of p-austenite and e-martensite. (c) Schematic diagram of the Debye rings of a'-martensite. 163

MECHANICAL PROPERTIES DEGRADATION OF HYDROGENATED AUSTENITIC SS

I. Gilad, Y. Katz and H. Mathias

Nuclear Research Center-Negev, POB 9001 Beer-Sheva, Israel.

INTRODUCTION Austenitic stainless steels (ASS) are often preferred as structural materials, due to their beneficial combination of improved mechanical properties and aggressive environment resistance. Clearly, metallurgical effects might be significant in ASS related to phase stability, segregation and sensitization processes. This turns to be particularly important by considering forming or thermo-mechanical processes. Plastic deformation below the MJ temperature causes martensitic phase transformation following the reaction: y •* E' + a' (y, e1 and a' are the FCC austenitic phase, HCP and BCC martensitic phases, respectively). The influence of y decomposition on ASS properties are not well established. As proposed by Jewett and others (1)» ductility might be selected as a possible criterion to evaluate the susceptibility of ASS to hydrogen effects. On the other hand, ductility variation with temperature, in air, are non-regular with minimum values above 900 K (2). In spite of this ductility reduction, necking occurs during uniaxial tensile tests resulting in a ductile fracture mode. ASS uniaxial tensile tests during or after hydrogenation show higher yield strength, lower ultimate strength and ductility loss associated with alternative brittle fracture modes (3,4). This behaviour is mainly observed in AISI 304L ASS hydrogenated by high fugacity methods. For example, cathodic charging provides hydrogen concentration up to 34 \ at a depth of about 1 urn (5). The present paper is aimed to track hydrogen effects due to gas charging methods which result in relatively low hydrogen concentration. More of a mapping concept is adopted, intended to spot possible sensitivity regions along the temperature scale. For comparative purposes, effects with and without hydrogen were investigated by means of mechanical testing, X-Ray diffraction technique and fractographical study.

EXPERIMENTAL PROCEDURE Standard tensile specimens with the axis parallel to the rolling direction, were machined from commercial 304L and 316L ASS plates, with a diameter of 7.14 mm and a gauge length of 50 mm. Tensile tests were performed in an inert atmosphere at stable temperatures from 77 K up to 1050 K, utilizing a nominal strain rate of 0.06 min"1. Precharging of the specimens was performed in dry hydrogen at 800 atm and 300 K for one month, or 500 K for one day. At these conditions, the theoretical hydrogen concentration near the surface is 0.56 a-s, while the average concentration of 0.085 a-i was actually measured. 164

In order to perform X-ray diffraction of the mechanically tested specimens, samples of the fractured specimens were cut along their axis by me^.s of spark errosion and then mechanically and electrolytically polished. Reflection X-Ray profiles were obtained from the uniformly strained region, utilizing Mo Kα radiation in conjunction with a graphite monochromator. SEM was applied for the examination of the necked region and of fracture surfaces of the strained specimens.

RESULTS AND DISCUSSION The temperature dependencies of the mechanical properties for the AISI 304L, before and after hydrogenation, are illustrated in Figs.l and 2. As shown in Fig. 1, the strength of the tested 304L drops monotonically, but with varying rates at different temperature ranges. It can be seen that hydrogenation caused no significant changes of the strength, except a slight increase of the yield strength between 200 K and 300 K. Referring to Fig. 2, the total strain to fracture, e-j-jis given as the sum of the uniform strain, eu, and the strain at the neck, en. Figure 2 demonstrates that the ductility of unhydrogenated 304L is strongly temperature dependent.

Values of eu reach a maximum of 65% at 270 K, and a minimum value of 15% near 1000 K. Figure 2 clearly indicates that the same trends of ductility changes have been obtained for hydrogenated 504L, with the exception that a loss of ductility occurred between 200 K and 300 K. Notice also that at 77 K, a higher value of EU has been obtained after hydrogenation. It is well observed that for hydrogenated and unhydrogenated materials, the reduction

of area (RA) curves are very similar to the corresponding en curves but not to E curves. This is not surprising, taking into account that both

RA and eu curves correspond to the plastic deformation at the necked region. Testing of several specimens at 3C0 K and at a strain rate of 0.006 min showed a better ductility, but no change in the strength, whether the specimens were hydrogenated or not. The mechanical behaviour of the 316L ASS was similar to that of the 304L, but with less pronounced hydrogen effects.

Valume fractions, Xa, of strain induced mariensitic BCC a' phase, were

evaluated from X-ray diffraction profiles. Curves of Xa vs. test temperature are shown in Fig. 3, for hydrogenated and unhydrogenated 304L and 316L

ASS. In order to allow a comparison between Xa and eu, the corresponding

eu vs. T curves are included. There are indications that the presence of hydrogen influences the formation of the strain induced a1 phase, especially for the 316L. In spite of this, approximately the same M^ temperature, namely 300 K, was obtained for the present testing conditions. It should be mentioned that in case of as-recieved 304L and 316 ASS, Caskey (4) obtained different M^ values, as indicated by arrows in Fig. 3. Fractography after strain to failure revealed a ductile fracture mode characterized by dimples. The size of the dimples increased with the temperature, but was usually smaller in the outer circumferential region as compared to the central region. Hydrogenation resulted in a change of the fracture mode at the outer region, only. Instead of dimples (Fig. 4a), facets formation and transgranular cracking occurred (Fig. 4b). Such a brittle fracture mode was obtained by straining at temperatures between 240 K and 300 K. The complex state of stress which developes during straining at the outer periphery of the reand tensile specimens, plays certainly 165

an important role in the occurrance of the brittle fracture mode. Generally, it is well recognized that hydrogen may concentrate at preferential sites during charging and redistribute in the course of plastic deformation. Local hydrogen concentration after cathodic charging may induce phase transformations and cracking at the surface, even without an external applied stress (5). After gas charging the hydrogen concentration is supposed to be low. But during straining at specific temperatures, the hydrogen concentration at discontinuities, could exceed a critical value needed to cause damage. As shown, 300 K was the highest temperature at which hydrogen affected the mechanical properties, and corresponds to the measured M^. No hydrogen degradation has been observed below 200 K, but straining at this temperature range resulted in extensive formation of a' martensite. In the curent study, mechanical properties degradation due to hydrogenation has been observed under the following conditions: 1) the presence of biaxial state of stress, 2) low austenitic stability which enabled induced martensitic transformation to occur and finally 3) the existence of sufficiant thermal drive at which the permeability is high in order to reach the critical hydrogen concentration at preferential sites. In fact, these condition really existed in the relative strongly hydrogen affected range between 200 K to 300 K. Referring back to the mentioned proposal of considering ductility loss as criterion for hydrogen susceptibility, an additional remark should be mentioned briefly. Based on the current findings, small differences between ductility loss values are not sufficiant for a proper evaluation regarding the extend of hydrogen damage in the broad sense of hydrogen degradation potential of 304L and 316L. Consequently, it is proposed to use hydrogen induced ductility loss only as a warning sign of possible hydrogen influences which often seems to be masked in metastable austenitic stainless steels.

ACKNOWLEDGEMENT The authors wish to express their appreciation to Mr. M. Kupiec and Mr. M. Aberman, for experimental assistence.

REFERENCES 1. R.P. Jewett, R.J. Walter, W.T. Chandler and R.P. Frohmberg, Rocketdyne, Division of North American Rockwell, Canoga Park Calif. 91304 (NASA CR-2163), March 1973, p. 67. 2. V.K. Sikka, R.W. Swindeman and C.R. Brinkman, Proc. ICF4, (Waterloo, Canada, June 19-24, 1977) D.M.R. Taplin (Ed.), University of Waterloo Press, 1977, vol. 2, pp. 561-567. 3. M.R. Louthan, G.R. Caskey, J.A. Donovan and D.E. Rawl, Mater Sci. and Eng., 1972, vol.10, pp. 357-368. 4. G.R. Caskey, Proc. Symp. on Environmental Degradation of Engineering Materials, (Blacksburg, Va., Sept. 21-23, 1981) M.R. Louthan and others (Eds.), Blacksburg Va. Polytechnic Ins., 1981, pp. 283-301. 5. H. Mathias, Y. Katz and S. Nadiv, Proc. Int. Symp. on Metal-Hydrogen Systems, (Miami-Beach, Florida, April 13-15, 1981) T.N. Veziroglu (Ed.), Pergamon Press, Oxford, 1081, pp. 225-249. 166

i r —i T —1 • 1 1 1 3041 304 L - -

\ -

- •

a unhydr ogcnated O i

hydr ogenated unhydro^enated ° - 0 • • hydrogenated . ... 5 i

\ • /

13-. \

- -a - y

*• • 1 i i II in

Fiy. 1 Stranght vs. temperature. Fig. 2 Ductility vs. temperature.

^ " •jnhydroRenated hydrogenated e__

Fig. 3 Volume fracture of a1 and Fig. 4 Fracture mode at peripherial uniform strain vs. temperature. region of 304L tensile specimen. 167

TENSILE FLOW OF AUSTENITIC STAINLESS STEEL AFTER THERMAL AGING IN HYDROGEN ATMOSPHERE

Y. Rosenthal1^*, M. Mark - Mark p,w itch1, A. Stern and D. Eliezer

1 - Nuclear Research Centre-Negev, P.O.B. 9001, Beer-Sheva, Israel. * - In partial fulfilment of the requirements for a M.Sc. degree. 2 - Ben Gurion University of the Negev, Dept. of Materials Engineering.

INTRODUCTION

Blanchard and Troiano [1], Whiteman and Troiano [2], Holzworth [3], Kolts [4] Eliezer et. al. [5], Hanninen and Hakkarainen [6], Narita et. al. [7] and others (8,9) have shown that considerable hydrogen embrittlement (HE) or hydrogen-induced degradation of properties (HIDP) may be produced in any austenitic staniless steel (ASS), however, stable with respect to deformation maternsite, provided two conditions were met: (1) supply of environmentalHat high input fugacities; (2) build-up of large concentrations of internal H across a significant part of the specimen's cross-section, rather than in some shallow sub-surface layer. These results, while essentially correct, apparently led to the still popular view that high charging fugacities of internal H contents are a prerequisite to considerable H effects in AGS. Thus, relatively little interest v.as left, the purpose of this study was to investigate the effects of low contents of internal H, of the order of a few hundreds of appm, on the following1 properties of selected austenitic stainless steel (ASS) in their thermally-aged (sensitized) condition: 1) plastic flow, strainhardening, ductility and fracture characteristics; 2) microstructural stability, i.e. extent of strain-induced martensite formation and possible changes of the Md temperature. Two ASS were chosen: 304L unstable at RT and 316L stable at RT in the hydrogen free condition.

EXPERIMENTAL

The two commercial-grade ASS were received as 0.2 mm thick sheets in the bright-annealed condition. H_ gas precharging of tensile specimens was conducted at 600°C in vacuum-tight ASS retorts, equipped with thermocouple ports and inlet/outlet vacuum valves. The H-pressure in the reports was kept constant at 0.5 MPa-gage for an exposure duration of 170 hr. Control specimens were heated in an argon atmospheis. All tensile , testing was conducted in air at RT. at a strain rate of O.5«'1O" 5" . The load-extension readings were fed to a PDP-11 minicomputer programmed to reduce such data to stress-strain and strain-hardening parameters, and complete flow curves too. Hollomon's empirical power equation was employed as a best-fit function:

a = a • £ o p 168

where a- true flow stress; a = strength coefficient; £ - true plastic strain; n - strain hardening°exponent. With accurate, high-resolution load-extension readings, the Hollomon equation yields for some ASS log-log flow curves comprising three linear stages, characterized by different n and a values.

RESULTS AND DISCUSSION

Tables 1 and 2 summarize the mean values of tensile properties of the 304L and 316L ASS, respectively, thermally aged at 600°C. The properties of hydrogenated specimens are compared to those of control specimens (Ar) in terms of a "hydrogen-induced change" parameter, AH% 100(11,,-". )/H, . The additional parameters serve to emphasize the idea that hydrogen-induced or assisted fracture is but the final event of a continuous, probably complex strain history. When this idea is overlooked and the "conventional" parameters solely are invoked us criteria of H effects, an incomplete or even misleading picture might emerge.

The ductility loss of the 316L (-18.8%) is roughly half that of the 304L (-34.8%), the more stable steel would appear to be considerably superior to the metastable one, at least in the given conditions of hydrogenation and testing. On the other hand, judging by the loss in uniform elongation, Ae , the two steels appear to be embrittled to nearly the same degree: -30.#% for 316L vs. -36.6% for 304L. Thus, the super- iority of the 316L over the 304L in the annealed condition drops considerably in the sensitized condition. The somewhat larger localized (necking) strain of the 316L, e~ - e = 8% vs. approximately 3% for 304L, seems to be associated with a different fracture mechanism, as discussed later.

Further, hydrogenation apparently had little effect on the 0.2% yield stress and UTS of either steel, if these parameters are examined out of context of the respective flow curves. The very small differences in 0.2% Y.S. and UTS between hydrogenated and control specimens could readily have been dismissed as experimental errors and data scatter. The implicit and misleading conclusion would have been that hydrogenation in our experiments had little effect on the ''overall" strength of the 304L and 316L ASS. The picture is significantly altered, however, if we examine the whole complex strain history by the events of gross yielding and plastic instability (onset of necking). To this purpose we propose a new criterion of hydrogen effects, already implicit in the definition of the "hydrogen-induced property change", AH in Tables 1 and 2: the "hydrogen-sensitivity of flow stress", which is simply a plot of the hydrogen-induced change in engineering flow stress, SH " SAr As = —£ 100 vs. the log of engineering plastic strain e ., % °Ar P1 (Fig. 1).

The plot of As - log e j in Fig. 1 is drawn up to e = 40% and thus does not include AUTS:P the UTS of hydrogenated andpcontrol (Ar) specimens are not proper flow stresses as (1) they occur at different strains, and (b) they mark the transition from a uniaxial state of stress to a trivial one. The flow stress of both 304L and 316L ASS shows 169

significant H effects over a large part of the uniform strain range. For the 316L the effect is one of general "hydrogen hardening" from macro- yielding up to e = 40% (and beyond, see Table 2). The 304L exhibits slight "hydrogenpsoftening" over the entire gross yielding range, up to about 0.8% strain; the effect then reverts to one of hydrogen hardening, quite parallel to that of 316L. Beyond e , = 40%, however, As drops stejply towards a negative value of AUTS^ giving rise again to slight softening. Thus, ther- ...-barging to low contents of internal H did not affect the overall strength of both steels, and the largest effects occurred at flow stresses (s^ - s.o5.) which are not among the mechanical parameters usually reported in the literature. Any H v3ffects on the strength of a metal ultimately must originate in some interaction between the dissolved hydrogen and the mechanisms which govern the plastic deformation of the metal. Some qualitative information relevant to such interactions may be obtained from strain- hardening displays, e.g. comparative Hollomon plots (Fig. 1 and 2). Further, Table 3 displays the values of the strain-hardening exponent n in the Hollomon equation and the true plastic strain £_ marking the onset of stages 2 and 3 in three-stage Hollomon plots.N CONCLUSIONS

1. A low content of internal hydrogen, of the order of 300 appm, was shown to induce significant degradation of tensile properties in thermally aged (sensitized) 304L and 316L type austenitic stainless steels: ductility losses, hydrogen hardening (flow stress increase), decreased strain-hardening capacity, and changes of fracture mode (304)L or morphology (316L) .

2. Hydrogen hardening and the changes in strain-hardening capacity mainly occurred over an intermediate range of strain; the yield stress and UTS of both steels were little affected by hygenation and thus shown to be doubtful or even misleading criteria of hydrogen effects when evaluated out of the context of overall deformation behaviour. Uniform elongation was shown to be an important criterion of hydrogen effects, besides the elongation to fracture. 3. Thermal pxecharging to low and uniform contents of internal hydrogen, with no surface damage, and a "scan" of the entire plastic range appear to be experimental techniques particularly useful in the study of intrinsic hydrogen effects in austenitic stainless steels.

ACKNOWLEDGEMENTS The authors thankfully appreciate the competent assistance of Mssr. R. Frenkel and A. Magen (NRCN). 170

REFERENCES 1. P.A. Blarichard and A.R. Troiano, Mem. Sci. Rev. Met., 57 (1960) 409. 2. M.B. Whiteman and A.R. Troiano, Corrosion, 21 (1965) 53. 3. M.L. Holzworth, Corrosion, 25 (1969) 107. 4. J. Kolts, in Stress Corrosion - New Approaches, STP 610, ASTM, Philadelphia, Pa. 1976, p. 366. 5. D. Eliezer, D.G. Chakrapani, C.J. Altstetter and E.N. Pugh, Metall. Trans. lOA (1979) 935. 6. H. Hanninen and T. Hakkarainen, Metall. Trans. lOA (1979) 1196. 7. N. Narita, C.J. Altstetter and H.K. Birnbaura, Metall. Trans., 13A (1982) 1355. 8. E. Kinkovitz and D. Eliezer, "Grain Size and Heat Treatment Effects in Hydrogen Assisted Cracking of Austenitic Stainless Steels", J. of Materials Science, 17 (1982) 3165. 9. P. Rozenak and D. Eliezer, "Effect of Metallurgical Variables on Hydrogen Embrittlement in Type 316, 321 and 347 Stainless Steel", Materials Science and Engineering, 61 (1983) 31. 171

(I.

71 .

II .

tt .

II .

I II I II S II'

Fig. 1: Hollomon Flow curves of 316L ASS thermally aged at 600 C.

i ii' i 10* 5 It

Fig. 2: Hollomon flow curves of 304L ASS thermally aged at 600°C. i.-ible 1 Effects of Hydrogen on Tensile Properties of Thermal ]y Aged 304 L ASS

Property Kngineering stress S, MPa Fract. Uniform Total \ UTS Stress elong. elong. at Engineering strain e , \ MPa MPa Exposure 0.05 0.1 0.2 0.5 O.b 1 5.5 10 20 SO 40

266.5 21>9.3 273.2 2S4.S 2SS.5 312.0 344.5 5"6.4 420.0 4S0.5 5S8.S 617.5 695.2 696.6 639.5 4S.23 51.17 il,,'h0i'OC 2.8 2.4 2.1 2.S 2.5 4.6 t>.3 8.4 16.2 24.6 1.5 25.4 59.8 55.9 2.24 6.48

274.4 27o.5 280.4 290.3 293.0 505.4 531.5 525.6 590.3 456. S 547.1 607.1 647.5 7-17.5 696.1 76.05 7S.5O Ar/600 C S.S 9.0 9.0 11.5 12.b 12.0 14.8 15.4 17.2 15.0 13.2 9.5 11.7 9.1 14.5 0.80 1.50

Hydrogen induced property change -2.9 -2.5 -2.6 -1.9 -1.5 5.9 5.S 7. (> 5.1 7.6 10.0 7.4 -S.I -56.6 -54. S

ii.2 O.2 (1.2 0.2 0.1 0.2 0.4 0.5 0." 0.5 0.5 i>. 3 0.4 2.1 5.1

Standard deviation irrelevant, as UTS and Sf are not flow stresses. Table 2 Effects of Hydrogen on Tensile Properties of Thermally Aged 516 L ASS

Fract. Uniform Total Propi-rty Engineering stress S, MPa UTS stress elong. elong. \ at Engineering strain e , ° \ >' MPa MPa LVOburc Li. 05 0.1 0.2 0.5 0.6 1 2 5.5 5 10 20 50 40

32 S.!- 329.5 531.0 357.0 345.5 559.3 392.8 426.2 475.0 539.0 642.7 twd.O 695.0 690.7 656.6 43.01 51.24

4.4 4.5 6.2 6.7 6.4 7.5 2.2 2.1 10.8 4.2 7.6 8.3 1.7 12.2 7.K 4.55 0.4 7

525.S 326.5 326.8 522.3 556.6 345.6 56S.S 400.5 450.5 492.2 578.S 626.7 656.0 6S4.7 657.0 62.21 65.14

Ar/(i0i'°C 5.5 5.9 5.6 5..S 3.1 A.I 5.9 7.2 7.9 5.3 3.0 5.4 4.5 20.0 11.3 3.76 4.02

inJiiccd i'!\>l'CTty i-li.mgo i> 9 1.0 1.3 4.6 5.0 4.6 l>.5 6.5 10.3 9.5 11.0 ".9 5.9 0.9 -0.1 -50.9 -JS.S an 6 ______-_------

0.(i 0.U 0.(1 il.l H.l 0.1 0.: 0.2 (1.4 0.2 0.2 (1.2 0.1 * * 5.1 2.4

* Standard deviation irrelevant, as UTS and S, are not flow stresses. 174

Tables: Strain-hardening Parameters of Thermally-aged 304L and 316L ASS-

\ Material, Stage 1 Stage 2 Stage 3 Parameter Vondition

0 07 + 0.01 0.17+0 .01 0.47 + 0.01 304L Z Ar 0.07 ± 0.02 0.24+0 .02 0.50 ± 0 .04 n H- 0.05 ± 0.01 0.14 ± 0 .01 0.38 ± 0 .01 316L I Ar 0.05 + 0.02 0.22 ± 0.05 0.43 ± 0.01

304L -2.9 -31. 3 -6.6 nH • nAr , n * 00 3161 -2.2 -35. 5 -11.8 nAr / 3 2 H9 4.5-10" 6.3-10" 304L 2

2 2 e Ar / 1.7-10" 9.6-10"

H t 5.4-10"3 5.2-1O"2 316L 2 Ar 1.6-10"2 9.o.io"

* computed from three-digit values of n, in order to resolve the minute yet real H-induced decrease of n over Stage 1. 175

CORROSION BEHAVIOR OF Al-Hu ALLOY THIN FILMS IN MICROELECTRONICS

J. Zahavi and M. Rotel

Israel Institute of Metals, Technion, Haifa, Israel.

H.C.W. Huang and P.A. Totta

General Technology Division, IBM Corporation Z/48A

Hopewell Junction, NY 12533, U.S.A.

INTRODUCTION

Uuminum and aluminum-copper film alloy have been used extensively as conductor lines for the production of VLSI circuits. It is known that \1-Cu bulk alloys are susceptible to general and localized corrosion CI,2) as was also observed with Al-Cu thin film alloys by several investigators (3,4,5).

Recently (6,7), electrochemical polarization techniques have been used in assessing resistance to localized corrosion processes in bulk (6) and thin film materials (7). However, no substantial or systematic research work has been reported on the behavior of Al-Cu thin film alloys and their susceptibility to localized pitting corrosion.

The present work aimed at studying the relationship between Al-Cu film electrochemical polarization behavior and film susceptibility to localized pitting corrosion, as well as various factors such as heat treatment, copper concentration affecting film corrosion behavior and corrison resistance.

EXPERIMENTAL Al-Cu alloy thin film specimens were E-beam evaporated on oxidized silicon substrates with Cu concentration varying from 0% to 8% wt%. Film thickness was 900 nm. Specimens were annealed for 1 hr at 400°C in inert atmosphere of flowing argon gas and cooled in air stream.

Polarization was conducted in a stirred bath of deionized water containing 10~3M of Cl-ions at 25°C and at a sweep rate of 0.2 mv/sec, while nitrogen gas was purged continuously into the solution. A microprocessor-based corrosion measurement system (8) model 350A, together with IR compensator module 365, manufactured by EGGG, NJ, USA, was used in conducting the polarization tests.

Cyclic polarization scans started at initial potential, which was about 176

250 mV below the given corrosion potential, Ecorr. The scan potential increased toward the noble direction while ';he corresponding current was recorded. After reaching the corrosion potential, the scan potential continued to increase in the noble direction reaching the Flade potential, Ef, and thereafter the breakdown potential, E^, as indicated by a rapid increase in current. Thereafter when the scan potential reached a predetermined vertex potential and vertex current, a reverse scan was started and continued toward more active potentials until repassivation occurred, characterized by the protection potential, E_. The hysteresis loop was completed when final predetermined potential was reached (Fig. 1).

Log Current Density

Fig. 1: Typical Cyclic Polarization Curve.

Samples before and after polarization were examined by optical microscopy, scanning electron microscopy (SEM) and Auger electron spectroscopy (AES). AES survey and depth profile were taken with a PHI Scanning Auger Microprobe, model 590A.

RESULTS AND DISCUSSION

Typical potentiodynamically cyclic polarization curves of as deposited and heat-treated Al-Cu thin film specimens are shown in Figs. 2A and 2B, respectively. The presence of hysteresis correlated very well with the mode of corrosion observed on film surface after the polarization tests. Localized pitting corrosion processes occurred on polarized specimen that displayed a well-defined hysteresis as shown in Fig. 2. 177

FILM ALLOY : AL - 4%Cu ; 9000^ 0.4 TREATMENT: AS DEPOSITED SOLUTION : D.I. WATER+ 10"3MCI" " 0.0 -

-0.4

-0.8

LLJ q -1-2 en

£ 0.4 FILM ALLOY : AI - 4%Cu ; 9000A" o TREATMENT: HEATING, 400°C v 1 hr SOLUTION : D.I. WATER +10"3MCL~ 0.0

-0.4

-0.8

-1.2

10° 10' 1OZ 105 104 105 1Oe nA /cm2

Fig. 2: Typical c'yclic potentiodynamic polarization curves for Al-4. %Cu thin film alloys. A. As deposited film. B. Heat-treated film. A correlation between the corrosion resistance expressed by the dimensionless ratio (TL^-E^/EyEf) to the various amounts of copper concentrations in as deposited and heat-treated films is given in Fig. 3. Each data point represented at least two measurements. Increase of copper concentration up to 4% in film specimen resulted in sharp decrease of film apparent corrosion resistance, namely the dimension!ess ratio, E_-Ef/Eb-Ef, for as deposited and heated film specimen (Fig. 3). However, further increase of copper concentration up to 8% did not significantly affect the ratio values. It should be noted that Al-Cu film containing less than 1% copper concentration 178 exhibited very high corrosion resistance. Heat treated film specimen exhibited higher values of the ratio E -Ef/E, -Ef compared to non-treated or as deposited films (Fig. 3), although the nature of the correlation given in Fig. 3 did not change with specimen heat treatment. Higher values of this ratio indicated higher corrosion resistance.

1 1 l 1 FILM ALLOY : Al-Cu: MOOX 1.0 - TREATMENT : A AS DEPOSITED _ o HEATING -tOO'C, Ihr SOLUTION : D.I. WATER + 1O"1 MCl" ae - -

o.e - • o "

ui 0.2 _-

I i i 1 2 4 6 COPPER CONCENTRATION [Wt%]

Fig. 3: Dependence of Al-Cu thin film corrison resistance on film copper concentration. Figs. 4 and 5 show typical views of localized corroded flower type zones in heat treated films. It was noted that the "flower", or the de- iendritic type corroded sites resulted from lateral corrosion processes. Copper rich second phase particles (point 1, Fig. 5B) were associated directly with localized corrosion processes as shown in Fig. 5B. These particles acted as efficient cathodic sites in comparison to the Al-Cu matrix (point 3 in Fig. SB], and resulted in localized corrosion process attack at the grain boundaries which led to the removal of complete grains (point 2, Fig. 5B).

Fig. 4: Typical view of corroded zone of polarized heat treated Al-8%Cu thin film alloy (900 mm in thickness). Optical microscopy (x200). 179

Fig. 5: Typical SEM observation of localized corroded zones of polarized heat treated Al-l-2%Cu thin film alloy (900 iron in thickness). A. General view of "dendritic" or "flower"- like localized corroded zones (xlOOO). B. High magnification of localized corroded zone shown in (A), Point 1 - copper-rich particle; point 2 - missing grain; point 3 •• grain matrix (x5000).

ocalized pits in the as deposited film are shown in Fig. 6. Exposure f the silicon substrate and the absence of aluminum copper film at 180

the localized pits were confirmed through X-ray mapping for silicon (bright spots in Fig. 6B) and aluminum (bright spots in Fig. 6C) . High magnification view of aluminum film-silicon substrate at the bottom of the pit (Fig. 6D) indicated that the corrosion processes resulted in gradual film dissolution up to the silicon substrate.

Fig. 6: Typical SEM observation of corroded zones of polarized as deposited Al-4%Cu thin film (2000 nm in thickness). A. Typical view of rounded pits reaching the silicon subsrate. B. X-ray image of Si(Ka) showing the presence of silicon at the bottom of the pits (bright spots in the picture). C. X-ray image of Al(Ka) showing the absence of aluminum at the bottom of the pits (dark areas are very low concentration of bright spots). D. X-ray image of Al(Ka) of the area shown in (E) . T.right spots indicated the presence of Al. E. Typical view of the interface between corroded film and tlie silicon substrate at the bottom of the pits.

Typical AES depth profiles of pitted zones in as deposited Al-Cu films are given in Fig, 7. These depth profiles showed that the Cu-to- Al signal intensity ratio in pit's walls (Fig, 7) was about ten times higher than that for film background while pit's center is mostly silicon . 181

PIT CENTER

PIT SLOPE < LJ Q.

1C Si UJ d a. AL

4 6 8 10 12 14 16 18 SPUTTERING TIME (min)

Fig. 7: Typical AES depth profile analysis of as deposited Al-4%Cu in the pitted area shown in Fig. 6. Sputtering rate was 2.5 nm/min. A. AES depth profile taken x>om the bottom of the pit, point 1 in Fig. 6. B. AES depth profile taken from the pit's walls, point 2 in Fig. 6. C. AES depth profile taken from non-corroded polarized film surface, point 4 in Fig. 6.

Mode of corrosion processes. Aluminum alloy usually exhibited high corrosion resistance due to formation of passive stable ronconductive aluminum-oxide layers. However, when discontinuous conductive oxide films were formed, either due to embedment or fast dissolution of intermetallic particles (10,11), or to breakdown events within the oxide coating (12), severe local corrosion processes took place (2). Aluminum-copper bulk alloys containing reactive copper-rich intermetal lies usually form nonuniform oxide film with non-regular structure (10,ll)which resulted in initiation and propagation of localized corrosion process. 182

Strehblow et al. (5) indicated the importance Oi oxide coating in corrosion protection and corrosion behavior of Al-Cu thin film alloys. Coating stability and protective properties were dependent on copper concentration in Al-Cu thin film alloy. Heat treated Al-Cu thin films obtained in this study with copper concentration up to 8% contained copper-rich intermetallics (up to 1 \im in size - Fig. 5), which significantly affected the uniform and continuous formation of oxide coating.

Copper concentration significantly affected the formation of copper- rich intermetallics in heat-treated Al-Cu thin film. With increase of film copper concentration the number of copper-rich particles was increased and therefore affected film corrosion behavior and resistance to localized corrosion. Heat-treated Al-Cu thin film containing copper concentrations of 0.5 to 1% did not produce copper-rich particles which immediately affected the corrosion resistance of the film, as shown in Fig. 3.

Copper-rich particles served as very efficient cathodic sites compared to anodic copper depletion zones at film grain boundaries or film grains (2,3,9). Anodic dissolution of Al (Al •> Al+3 + 3e~), therefore, took place preferentially from grain boundaries while cathodic reactions consuming electrons such as 2H2O + 02 + 4e~ ->- 40H took place on copper-rich particles resulting in localized intergranular corrosion followed by removal of complete grains (Fig. 5) known as the missing grain phenomenon (3). The presence and remainder of copper-rich particles within the localized corroded areas (Fig. 5B) indicated that these particles actually served as cathodic sites maintaining their initial shape and size.

As deposited Al-Cu thin film alloys with copper concentration up to 8% neither showed a distinc: '.re grain structure nor the presence of copper-rich intermetallies before and after polarization. However, severe localized corrosion processes characterized by single isolated rounded pits (Fig. 6) were observed,

AES examination of corroded and non-corroded zones (Fig. 7) indicated the presence of localized copper concentration particularly within pitted zones as well as the presence of about 40 nm oxide coatings in non-corroded zones. Aluminum oxide coatings covered the Al-Cu film surface. However, oxide thickness was reduced significantly at Al-Cu thin film zones containing local high concentrations of copper. Furthermore, these were probably copper-rich coatings with conductive properties (5,13,14) acting as local sites for cathodic reactions. The aluminum oxide coating formed adjacent to the copper-rich oxide coating consisted of flaws (15) and/or breakdown events (12). The presence of chloride ions in the D.I. water used prohibited the repair or the healing processes of the oxide coating (12) which resulted in localized corrosion processes. The initiation and propagation of localized pitting corrosion process took place while the adjacent Al-Cu thin film copper-rich zones served as efficient cathodes with surface area much larger compared to the copper depleted anodic zones. Anodic dissolution of the aluminum ions resulted in pit propagation reaching the silicon wafer substrates (Fig. 6). 183

Conclusions. The contribution of this study was to evaluate the susceptibility of Al-Cu thin film alloys to localized corrosion processes making use of electrochemical potentiodynamic cyclic polarization techniques. A dimensionless ratio between perfect passivation (E -Er) and total passivation range (E^-Ef) was defined to indicate thl susceptibility to localized corrosion. The higher the ratio, the higher the resistance to localized pitting corrosion. Heat treated Al-Cu thin film exhibited higher resistance to localized corrosion compared to as deposited Al-Cu thin films. The susceptibility to localized corrosion increased with increase of copper concentrations both in heat treated and as deposited Al-Cu thin film alloys. Heat treated Al-Cu thin films exhibited shallow branched dendritic or flower- like localized corroded areas which were directly associated with the presence of copper-rich intermetallics preferentially at grain boundaries.

In as deposited Al-Cu thin film alloys, well defined isolated rounded pits were observed. These were associated with localized copper concentrations, but not with copper-rich particles or grain boundaries.

REFERENCES

1. N.D. Tomashow, "Theory of Corrosion and Protection of Metals", Macmillan, New York (1966).

2. J.R. Galvele, S.M. de DeMichelli, I.L. Muller, R.B. de Wexler and I.L. Alanis, "Critical Potentials for Localized Corrosion of Aluminum Alloys", in "Localized Corrosion", NACE, Editors R.W. Stachle, B.F. Brown, J. Kruger and A. Agrawal, Volume published in 1974.

3. P. Totta, "The Missing Aluminum Problem" IBM Technical Report TR 22.1447 (1972). 4. W.T. Lee, J.M. Eldridge and G.C. Schwartz, J. Appl. Phys. 5_2_, 4, pp. 2994-2999 (1981).

5. II.H. Strehblow and C.J. Doherty, J. Electrochem. Soc. Vol. 125, No. 1, pp. 30-33 (1978).

6. R. Baboian and G.S. Haynes, "Cyclic Polarization Measurements - Experimental Procedure and Evaluation of Test Data", Electro- chemical Corrosion Testing ASTM STP 727, F. Mansfeld and U. Bertocci, Eds., American Society for Testing and Materials, pp. 274-282 (1981).

7. J. Zahavi, M. Rotel, H.C.W. Huang and P.A. Totti?, "Corrosion Behavior of Thin Films in Microelectronics". First Annual Report, No. 5095-32, September 1983.

8. W.M. Peterson and H. Siegerman, "A Microprocessor-Based Corrosion Measurement System", in Electrochemical Corrosion Testing, ASTM STP 727, F. Mansfeld and U. Bertocci, Editors, American Society for Testing and Materials, pp. 390-406 (1981). 184

9. K.R. van Horn, "Aluminum", Vol. I, p. 49, 115, 2091, American Society of Metals (1967).

10. J. Zahavi, A Zangvil and M. Metzger, J. Electrochem. Soc. 125, 268 (1974).

11. J. Zahavi, H. Kerbel and 0. Korotkina, J. Electrochem. Soc. 129, 7 (1982).

12. J. Zahavi and M. Metzger, J. Electrochem. Soc. 121_, 268 (1974).

13. G.C. Wood and A.J. Brock, Trans. Inst. Met. Finish. 44, 189 (1966),

14. J. Cote, E.E. Howlett and H.J. Lamb, Plating, 57_, 484 (1970).

15. J.A. Richardson and G.C. Wood, Corrosion Science, 10, 313 (1970). 185

QUANTITATIVE NONDESTRUCTIVE EVALUATION USING ULTRASONIC WAVES

Laszlo Adler

Department of Welding Engineering The Ohio State University Columbus, Ohio 43210

INTRODUCTION Recent developments of materials characterization by ultrasonic waves is summarized by presenting a systematic approach to discontinuity analysis. Figure 1 presents a systematic approach to ultrasonic evaluation of material structures which may be a weld or a joint, for example: • Locate all flaws • Characterize each flaw (determine size, shape, orientation and composition) • Characterize the material (determine the elastic properties, grain size, surface roughness, etc) • Evaluate the seriousness of the flaws' presence (using fracture mechanics techniques). In order to determine the presence (or absence) of weld discontinuities, an ultrasonic image is produced. Defects will be delineated as areas of increased ultrasonic-echo return. For flaws large compared with the beam dimensions, the ultrasonic image shows the extent of the defect. Flaws on the order of and smaller than the ultrasonic wavelength scatter the incident sound beam. Some of this energy returns to the transducer and appears on the image as a weakly scattering region in the material. Regions such as these are flagged for later investigation. Next, the characteristics of the weld metal are measured. The grain structure and roughness of the ultrasonic beam entrance surface will have an effect on all subsequent measurements, and so must be determined. In addition, the mechanical properties of the material surrounding any flaws will profoundly affect the strength of the weld. Also, regions of the weld containing dense porosity or clouds of inclusions will lower weld strength. Determinations of the concentration and size distribution of voids or inclusions should permit calculations of estimated weld mechanical properties. One now returns to the suspect regions in the welded structure, namely areas of increased echo return in the ultrasonic image. A broadband ultrasonic wave is directed toward the suspect region and the backscattered signal is then processed to obtain a magnitude spectrum (amplitude versus frequency)„ If the spectrum shows deep and periodic modulation, the suspect region contains a planar discontinuity (eug. crack). If, however, the spectrum is relatively smooth, the flaw is volumetric (e.g., pore). Characterization of the defect requires further signal processing., Discrimination of planar from volumetric discontinuities permits the appropriate processing algorithm to be selected. Ultrasonic spectroscopy (using cepstral analysis) is utilized for 186 determining the size of crack-like flaws. Upon processing, the distance (along the ultrasonic beam path) is calculated which separates the near and far edges of the flaw. Interrogation of the defect from a number of angles may be used to obtain the flaw's size and shape- Volumetric defects are characterized by "Born inversion" processing.. This algorithm returns both a line-of-sight estimate of flaw radius and the cross-sectional area of the flaw (along the ultrasonic beam path). "Observation" of the defect at a number of "look angles" discloses its shape and dimensions. Once all flaws are characterized, data tabulating • Coordinates of all flaws • Shape and size of all defects • Material elastic properties (Young's and shear moduli, Poisson's ratio) • RMS surface roughness • Concentration and size distribution for areas containing clouds, of porosity, or inclusions can be output for evaluation using fracture mechanics techniques. Weld Discontinuity Analysis System

Imaging Mapping of Large Flaws and Location of Suspect Regions

Material Characterization Determine Mircostructure, Porosity and Surface Roughness

Investigate Suspect Regions Use Broadband Backscattering Technique

Planar Discontinuity Volumetric Discontinuity Ultrasonic Spectroscopy Born Inversion !0>ka>l 2.5>ka>0.5 Determine the Distance Determine the Flaw Separating the Edges of the Flaw Size and Shape Figure 1. Systematic Procedure for Ultrasonic Nondestructive Evaluation. 187

DESCRIPTION OF SYSTEM ELEMENTS IMAGING The purpose of ultrasonic imaging is to locate regions in the weld which contain flaws. A focused ultrasonic beam is directed toward the weld. Changes in material elastic properties or density will cause reflection and scattering of the incident beam. A portion of this echo energy is intercepted by the receiving transducer. The receiver output is recorded with the spatial coordinates of the scattering region (Figure 2). If the entire volume of the weld is scanned one builds up a 3-dimensional map of echo amplitude versus position.

Figure 2. Assignment of Coordinates to a Weld. Display Modes Display of echo information is normally presented one plane at a time. The image is termed a C-scan (echo amplitude at constant depth) if the image plans is parallel to the sample surface (x-y plane). It is a B-scan (display brightness indicates echo amplitude) if any other plane (e.g. x-z, y-z or oblique) is presented. It is also possible to display the echo strength as a brightness-modulated three-dimensional isometric presentation (Figure 3). Echo amplitude is shown as display height as a function of two orthogonal spatial coordinates. Because of the ease in implementing this isometric display format, it was chosen for all our B- and C-scans. 183

Figure 3. B-Scan Ultrasonic Image of a Fatigue Crack (a). The Rayleigh surface waves used (b) to investigate the crack were generated by the wedge method. Image Optimization The scattering amplitude measured for a particular element in the weld is dependent on * the ultrasonic wave mode (longitudinal, shear, surface wave), • the frequency of the incident wave, the incident angle and - the spatial distribution of the incident beam. 189

Each of these factors may be adjusted such that flaw detectibility is optimized,, Analytical as well as experimental work is needed in this area., Image Processing It is possible to accentuate the presence of flaws in an ultrasonic image by postprocessing. A method, gradient processing, for enhancing edges in an image is introduced here. Let the echo amplitude as a function of spatialposition be represented by f(x,z) [B-scan display mode],, Then the gradient is defined as the vector [1] G[f(x,z)] = 6f 5x (1) &f_ 5Z The gradient vector points in the direction of the maximum rate of change of f(x,z) and its magnitude gives the maximum rate of change in f(x,z) per unit distance in the direction of G^. For manipulation of the digital image, the magnitude of the gradient is approximated by

G[f(x,z)] = |f(x,z) - |f(x,z) - f(x,z+l) (2) The gradient of the ultrasonic image is displayed if the gradient is above a threshold value. If the gradient is below threshold (presumably the case for pixels representing unflawed regions in the weld) the image pixel is sat to zero. Flaws in the processed image stand out more clearly (Figure 4b) than in the original ultrasonic image (Figure 4a). Research is needed for adaptive threshold setting, other methods of image processing and automated flaw recognition schemes. 190

Ultrasonic B-Scan Images of a Specimen containing Multiple Defects. Image before (a) and after (b) gradient processing. MATERIALS CHARACTERIZATION The mechanical properties of the material in which a flaw is embedded may be as important as defect size in determining strength. Material characteristics which are of importance are listed in Table 1. It is possible to nondestructively determine many of these properties from ultrasonic velocity and attenuation measurements. Both analytical and experimental studies have been made to determine how frequency-dependent velocity and attenuation measurements may be used to infer the concentration and size of pores occurring in dense clouds. Inhomogeneities (such as pores and inclusions) weaken the structural components in which they occur. It is important to nondestructively determine the size distribution and concentration of pores or inclusions. A multiple scattering theory was developed for treating wave propagation through inhomogeneous material. Matrix and second phase elastic properties inclusion concentration and size distribution are used as input.. Dispersion and frequency-dependent attenuation are calculated using theory. 191

Table 1 MATERIAL PROPERTIES

Tensile Modulus Shear Modulus Tensile Strength Shear Strength Bond Strength Hardness Surface Finish Impace Strength Fracture Toughness Anisotropy Microstructure Grain Size Porosity, Void Concentration Phase Composition Hardening Depth Residual Stress Heat Treatment Profile Fatigue Damage Ultrasonic Wave Propagation in Cast Iron-Graphite Composite The same multiple scattering treatment applied to porous media may also be used to analyze the problem of an ultrasonic wave traversing a material containing clouds of solid inclusions. Cast iron containing compact flake and nodular graphite, and also specimens of gray iron is studied. Because the graphite in nodular iron appears mostly as spheres, it was thought the multiple scattering theory could be used with only minor changes to include the elastic properties of the graphite nodules. The properties of the cast iron matrix were estimated from ultrasonic velocity measurements on a sample of 1045 steel. This type of steel was chosen because it has approximately the same relative percentages of ferrite and pearlit-2. Table 2 lists pertinent material properties for the matrix (iron) and inclusions (graphite nodules). Elastic moduli were calculated from the velocities. Table 2. Properties of the Constituents in Nodular Cast Iron. Velocities (km/sec) Material Longitudinal Transverse Density (g/cr3) iron (1045 steel) 5.8 3.1 7.2 graphite 3.2 1.8 2.2

Values for hte matrix material, iron, and the inclusions, graphite spheres were input to the multiple scattering theory to determine the phase velocity and attenuation as a function of frequency. Narrowband and broadband ultrasonic experiments were performed on carefully prepared specimens of cast iron. Figure 5 plots theoretically-determined attenuation and experimental measurements. The shape of the curves are 192 identical; however, there is an offset. We ascribe this to incorrect assumptions of the wave speed in the graphite and to the fact we assumed a uniform size of graphite nodules, when there is actually a distribution of sizes. Surface Roughness The condition of the surface through which ultrasonic beam enters the material can have a profound effect on the frequency content of the pulse, the angle of the refracted energy and on the spatial distribution of the beam in the solid. Both imaging and defect characterization techniques will be affected by the surface finish on the weld under evaluation,, Much of the unreliability of contact ultrasonic testing arises from multipath effects and trapped air due to roughness on the sample surface.. We have begun studies aimed at nondestructively inferring parameters describing surface roughness (rms roughness and correlation length for randomly-rough surfaces and periodicity and peak to valley height for periodic surfaces). Surface roughness may be measured ultrasonically from the angular or frequency dependence of wave scattering. Ultrasonic measurement of rms roughness (h) may be carried out in the following manner. An ultrasonic transducer, operated in the pulse-echo mode, is used to record the intensity of backscattered ultrasonic compressional waves,. The experiment is repeated for a smooth-surfaced sample of the same material (giving I ). If low frequencies are used, the backscattered intensity is approximated as where k is the wavenumber of the ultrasound in the liquid be.th. The roughness, h, may be estimated from

-17.37k2/ (4) dhere AdB is the relative amplitude of the waves backscattered from the rough surface compared to that of the smooth surface. The advantages of the ultrasonic measurement of surface roughness are • it is nondestructive, ° determinations may be made rapidly • whereas mechanical profilometers measure roughness only along the line of stylus traverse, the ultrasonic method averages over the entire insonifiecl area (be it small or large),, 195

« — Theoretical

«- Experimental

1.0 E e CD •a 0.8

a> '5 S 0.6 U

0.4

0.2

0. .06 .08 .10 .14 .18 .22

Kpa Figure 5. Attenuation of Ultrasonic Waves in Nodular Cast Iron as a Function of Frequency (o - multiple scattering theory and x - experiment). DEFECT CHARACTERIZATION Ultrasonic images produced in the first stage of our evaluation disclosed the extent of large flaws and identified weakly scattering regions as areas potentially containing smaller defects. For a fracture mechanics evaluation, the size and shape of these small discontinuities is required (in addition to their location, as determined from the ultrasonic image). Flaw characteristics may be determined from ultrasonic scattering data; however, the defect must first be classified as volumetric (3-dimensional, e.g., pore or inclusion) or planar (2-dimensional, e.g. crack-like). The classification of defect type is carried out by identifying discriminatory features in the backscattering frequency spectrum. The procedure below is followed. 1) A broadband ultrasonic pulse is directed toward the suspect 194

region in the weld. 2) The backscattered signal from the region is sampled digitized and stored in computer memory. 3) A system normalization signal is acquired by sending the incident pulse through a unflawed region in the material toward a perfect reflector (polished solid-air surface),, 4) Perturbing effects of the data acquisition system and material intervening between the ultrasonic transducer and the defect are removed by deconvolving the defect signal by the signal from the reference reflector. 5) The magnitude spectrum is displayed. • The defect is classified as planar if the spectrum has deep, periodic modulation (explanation given later). • The defect is classified as a volumetric flaw if the spectrum is relatively smooth. 6) If there is some ambiguity in classifying the defect type (5), then the suspect region is interrogated from a number of angles. • A planar discontinuity will return large backscattering signals for orientations in which the incident ultrasonic beam is directed normal to the plane of the flaw,, • Amplitude of signals returned from a volumetric discontinuity depend less strongly on angle. 7) If the discontinuity is planar, an inversion scheme termed ultrasonic spectroscopy is followed. 8) If a volumetric flaw is to be characterized, the Born inversion algorithm is utilized. The Born inversion procedure is summarized here. Born Inversion As an ultrasonic wave strikes a volumetric flaw (Figure 6), the cross-sectional area encountered increases as the wave propagates. When the wave is just incident on the defect (time, t-j) the area encountered is minimal. As the wave moves onward (time t^, then t^, . . „) the are encountered increases [k^, Ag, . . .)„ The cross-sectional area of the flaw reaches a maximum when the wavefront reaches the center of the flaw. The area encountered decreases thereafter. The Born inversion algorithm returns the area function A. versus t^. From this, a line-of-sight estimate of flaw radius may be calculated. 19S

t, MJW

Plane Wave

AREA Area 1 Radius = '2 ( w a v e speed < I,

T i me

Figure 6. Plane Wave Encountering A Scatterer. a) Position of wavefront at times t.. b) Area of the scatterer encountered at times ti and c) Area function. Ultrasonic Spectroscopy Ultrasonic spectroscopy applied to defect characterization was pioneered by Adler, et« al. [1,2,3].. The fundamental ideas underlying the technique are summarized here (details may be found in the book by Fitting and Adler, [4]).

TRANSDUCER

REFLECTOR

Figure 7. Ultrasonic Waves Scattered from the Edges of a Planar Reflector. An ultrasonic wave directed toward a planar discontinuity will be in part specularly reflected from it, but also scattering will occur from the near and far edges (flashpoints, [5]) of the flaw (Figure 7). 196

If the echo is transformed to the frequency domain (via the Fast Fourier Transform, FFT) the time spacing of the signals may be determined. Consider a single signal y(t) which has a spectrum Y(2irf). The magnitude spectrum of two such signals, separated in time by 2tQ, has been shown [6] to be |2 cos2irf tQ| |Y(2Trf) |. That is the spectrum is modulated, and the period is determined by the time separation (Figure 8). The spacing (Af) of the frequency minima may be used to determine the time separation (At) of the ultrasonic echoes: (5) Af = l/2t0 = I/At. If the wave speed is known, then the separation of the near and far flaw edges may be calculated and the flaw dimensions determined from d = 2 Afsine (6) The orientation of the crack (e) may be determined from a number of angular measurements. EXAMPLES OF ULTRASONIC NONDESTRUCTIVE WELD EVALUATION As a test of our systematic method for evaluating welded structures (Figure 1 and previous descriptions) we began with two welds containing well characterized defects. Weld sample number 2 was carefully prepared to contain only planar discontinuities. Three artificial flaws, nominally 1/16", 1/8" and 1/4" in diameter were incorporated into the weld, along its center line (Figure 8). The plane of the defects was parallel to the sample surface.. Both the top and bottom surfaces of the specimen were ground flat.

Figure 8. Weld Sample #2 Containing Planar Discontinuities 1/16", i/8" and 1/4" in Diameter. 197

Imaging A 5 MHz focused transducer (approximately 6" focal distance, 1" diameter) was scanned in a water bath above the weld. The water path length was adjusted to focus the ultrasonic beam approximately at the depth the discontinuities were thought to occur. Spacing between transducer positions (along the weld) was 5 mm. • the ultrasonic waveform returning from the weld is sampled, digitized and displayed (Figure 8 ), • the signal is full-wave rectified (Figure 8 ), • an estimate of the envelope of the waveform is determined (Figure 8 ), • the number of positions along the weld where the transducer will be located is input (Figure a ), ' wave speed for the mode of ultrasound used in the imaging is input (Figure 9 ), " the depth increment at which the echo amplitude is to be sampled is input (Figure 9 ), ' a file name identifying the image data to be stored is given (Figure 9 ) and • the depth (time) range to be used for image display is defined (using the graphic terminal's cursor) (indicated by the vertical lines in Figure IO).

Figure 9. B-Scan Ultrasonic Image of a Weld Containing Three Planar Discontinuities (1/16", 1/8" and 1/4" in diameter). 198

-40 —

dB ii'r\ ^ \ -60 7'7 \ A / — \ / \j I V \ L / Ni 1 A \ / \\' —'"""\ \ -80 r V \ ! i 1 10 15 20 Frequency (MHz) Figure io- Backscattering Spectra from Suspect Regions in Weld #2.. Ultrasonic Spectroscopy. Broadband ultrasonic pulses from the 15 MHz unfocused transducer were coupled to the weld through a 6" water path. Backscattering signals were acquired for each flaw at 0-, 10- and 20-degree refracted angles in the solid. The spectra computed from these signals was plotted versus frequency (Figure ). The absence of deep nodulation at O-degrees indicates the plane of each flaw is approximately parallel to the surface of the weld. The average spacing of minima in the spectra was determined (Table 3) and equation (6) used to calculate the flaw diameter. Table 3. Average Spacing of Minima in the Backscattering Spectra from Flaws in Weld #2..

11 10-degree Measurement 20-degree Measurement A 9 MHz 4 MHz B 4.7 MHz 2.3 MHz C 2.55 MHz 1.13 MHz 199

Nondestructive Evaluation Summary For Weld #2

Longitudinal Wave Speed: 5939 m/sec Shear Wave Speed: 3267 m/sec Young's Modulus: 2.566 Shear Modulus: 8.219 X 1010 Poisson's Ratio: 0.283 Suspect Regions: IS. X. 1 1 A 2.0 cm 0 7.5 B 7.5 cm 0 7.5 C 12.5 cm 0 7.5 Flaw Type: IS. Type A planar B planar C planar

Flaw Orientation: ID Angle with Respect to Surface A 0° B 0° C 0°

Flaw Dimension: (assume circularly ID 10-degree Measurement 20-degree Measurement symmetric) A 1.8 mm 2.06 mm B 3.5 mm 3.6 mm C 6.36 mm 7.3 mm Compare these ultrasonically determined results with the actual position and dimensions below

ID X Y Z Diameter (mm) A 1.8 cm 0 7 .5 cm 1.6 B 7 cm 0 8 cm 3.2 C 12.05 cm 0 8 cm 6.35 200

REFERENCES [1] Ho U Whaley and L. Adler, Flaw Characterization by Ultrasonic Frequency Analysis," Mat. Eva!., 2j3 (8), (1971). [2] H. L. Whaley and L. Adler, Model for the Determination of the Size and Orientation of Reflectors from Ultrasonic Frequency Analysis," J. Acoust. Soc. Am., 48 (1), (1970). [3] H. L. Whaley and L. Adler, "A New Technique for Ultrasonic Flaw Determination by Spectral Analysis," Technical Memo, Oak Ridge National Laboratory, ORNL-TM-3056, (1970). [4] D. W. Fitting and L. Adler, Ultrasonic Spectral Analysis for Non- destructive Evaluation, Plenum Press, New York (1981)„ [5] J. D. Achenbach, L. Adler, D. K. Lewis and H. McMaken, "Diffraction of Ultrasonic Waves by Penny-Shaped Cracks in Metals: Theory and Experiment," J. Acoust. Soc, Am., 66 (6), (1979). [6] W. A. Simpson, "A Fourier Model for Ultrasonic Frequency Analysis," Mat. Eva!., 34 02), (1976). 201

ON THE HOMOGENIZATION PROBLEM IN SINTERED ALLOYS

L. Levin Department of Materials Engineering Technion, Haifa

A. Stern Nuclear Research Center, Beer-Sheva and Department of Materials Engineering Ben-Gurion Univ. of the Negev, Beer-Sheva

ABSTRACT

The present study contains a detailed analysis of the sintering process at temperatures at which one of the components is in a liquid state. Experimental results were obtained for the Fe-Cu system containing 5 wt.% copper. Completion of the sintering process is accompanied by achieve- ment of homogeneity in composition. The different stages of homogenization were observed by methods of optical microscopy, hot—stage metallography, scanning electron microscopy plus EDAX and microprobe analysis. At small amounts of copper the process consists of the following stages: dissolution of iron in the liquid copper, covering of iron granule surfaces by liquid solution; penetration of Cu into the Fe matrix; migration of the copper atoms towards a uniform distribution. At higher amounts the process is complicated by penetration of the liquid phase at the granular interfaces and along grain boundaries. It was shown that the duration of each stage can be calculated when the amount of copper, and mode of penetration are known - by solution of the diffusion equation. The model is applicable to any system with limited solubility of the components.

INTRODUCTION

During liquid phase sintering in the iron copper system a series of processes take place: dissolution of iron in the liquid copper, spreading of the liquid copper over certain surfaces in the material, penetration of the copper into the iron matrix, rearrangement of the iron particles and changes in the total porosity of the alloy^"^ in the present study attention was directed mainly at the diffusion processes that bring to a homogeneous composition of the alloy.

In an earlier work a model for diffusion penetration in the system was developed. The model was based on the assumption that the copper atoms penetrate the iron matrix from the free surfaces on melting of the copper. The assumption referred to the cast of small amounts of copper in the alloy - 2%, and its validity increases with decrease of Cu content.

In the present study an alloy with 5% Cu was chosen. The behaviour of the material during the sintering process was investigated by optical metallography, including vacuum hot stage microscopy, and by the SEM and EDAX techniques. 202

EXPERIMENTAL

Table 1 contains data on the powders used in this investigation.

Table 1. Powder Characteristics

Material Size Amount Impurities Source Total

Electrolytic 125-150Mm 66.5% 240 ppm BDH Laboratory Iron Chemicals

Electrolytic 45-63 ym 28.5% 240 ppm BDH Laboratory Iron Chemicals

Electrolytic 45-63 ]im 5% 415 ppm Baker Laboratory Copper

The average grain size of the iron particles was about 20 ym.

After blending for one hour the powder mixture was placed in a neoprene sleeve and isostatically pressed at 10° Pa. The specimens were heated in vacuum for 10 min at 750°C to achieve mechanical strength necessary for further handling.

The sintering process was carried out at 1100°C in two variants:

(1) For short times at temperature, the specimen was heated in a Reichart vacuum hot stage with the sample temperature determined by a thermocouple embedded directly in the sample. The hot stage microscope was used in an attempt to observe the sintering process in situ.

(2) For long periods up to 24 hours at temperature, the specimens were sintered in a vacuum resistance furnace.

In order to perform most of the process at the chosen temperature, the last stage of heating was rapid: from 750°G to 1100°C in 4 min. In the first variant, cooling was effected by a stream of helium, in the second.- by water quenching. Specimens sintered in the vacuum resistance furnace were examined by the above-mentioned techniques. The change in the copper distribution after sintering times of 0, 3, 10, 30 min and after 2, 6, 12, 24 hours was checked.

RESULTS

The different stages of the copper melting as obtained with the aid of the hot-stage microscope are presented in Fig.l : la - a copper granule just before -melting; lb - onset of the melting process; lc - a pore at the spot where the granule had been located; Id - paths of liquid copper 203

penetration in the iron matrix. All the stages shown in the figure are accomplished in the interval between 1083°C and 1100°C.

Tha ••..'c.roscopic structure obtained in the vacuum resistance furnace after 10 mi-i of sintering is seen ic 7ig.2. The dark spots represent pores. A Scanning Electron Micrograph after the same time is presented in Fig.3. On the topography of the structure, the profiles of Fe and Cu obtained by EDAX (Scanning along the central line) are superimposed. The same profiles are shown in Fig.5 at lower magnificatian. The valleys in both profiles indicate pores and do not belong to the "real concentration profile" developed as a result of diffusional processes. The spacing of the "real maxima" of the Cu profile is not uniform but in most cases follows the Fe grain size (V^se of copper penetration along the grain boundaries). The distribution of copper is much more uniform after 6h sintering see Fig.6. An almost uniform distribution of copper is achieved after 12h of sinter- iuij at 1100°C - Fig.4.

ESTIMATION OF TIMES NECESSARY FOR. DIFFERENT STAGES OF H0M0GENIZATI0N

For diffusion calculations we begin with, the time count when 1100"C is achieved and the copper is already melted.

The copper melt can dissolve 4 at.% of iron. The dissolution process is practically instantaneous as shown in an earlier work^. The FeCCu) melt spreads along the boundaries of the Fe particles of different types (granule surfaces, grain boundaries). The time necessary for the melt to arrive at "starting positions" for the diffusion into the Fe matrix depends on the type of position and is of the order of several minutes^.

An estimate of the thickness of the liquid layer at the starting position is essential for further calculation; Assuming that most of the melt was spread along the grain boundaries, the thickness was found as h = 0.5 ym. During penetration of the liquid copper into the iron matrix, the copper, concentration in iron can be estimated^ as

C = CQerfc (x/2/5t) (1)

The total volume of copper penetrating per unit surface of an iron particle can be expressed as

dx = h<^ (2) where c_ = 0.96 is the concentration of copper in the liquid, t. the duration of this stage of the reaction, characterized by Co = 0.08 Cmax. solubility of Copper in iron at the sintering temperature) , D is the , diffusion coefficient of Cu in Fe. According to Smithells and Brandes D = 1.8 x 10"10 cn^sec"1 at 1100"C.

Taking into account that

/ erfc zdz = 0.5642 , 5 204

the time of penetration can be calculated from eq. (2) :

t = 1.07 x 103 sec - 18 min.

The next stage of the process is migration of the copper atoms towards a uniform distribution. The relationship between the effective distance the solute atoms have to transverse to achieve homogeneity, and the duration of this process can be presented as

KDt, C3) where K depends on the assumptions concerning the shape of the particle in which the diffusion takes place and on the degree of homogeneity^.

Assuming the simple case

I = 2Dt, C4) we find the correlation presented in Table 2.

Table 2: Effective distance vs. homogenization time

Effective distance Homogenization for Homogenization time

5.3 ym 10 min

7 ym 18 min

13 ym 1 h

32 ym 6 h

45 ym 12 h

55 ym 18 h

63 ym 24 h

Taking into account that the center of mass of the copper in the iron au the moment the liquid disappears is removed from the surface by about 3' ym , we have for the case when penetration proceeds from a granule surface I - 55 liia; if penetration proceeds from agrain boundary, £ = 7 ym In the last case the diffusion time is "VL8 rain; in the first case ^18 h. This is the time necessary for full homogenization. The result is in accordance with experimental data. 205

ACKNOWLEDGEMENTS

The authors would like to thank F. Simca and D. Shmariahu for technical assistance and helpful discussions.

REFERENCES

1. L. Levin, A. Stern, S. Dirnfeld, Z. Matallkde 21 (1980), 621.

2. W.A. Kaysser, W.J. Huppmann, G. Petzow, Powder Metallurgy, Q (1980), 86.

3. W.J. Huppmann et al., Zeitschrift fur Metallkde, 2P_ (1979), pp. 707, 792.

4. C.J. Smithells and E.A. Brandes, Metall. Ref. Book, 5th ed. Butterworths, London (1978), 874.

5. H.S. Carslaw and J.C. Jaeger, Conduction of Heat in Solids, Clarendon Press, Oxford (1967), 59.

6. L.E. Larson and B. Karlson, Materials Science and Engineering, 2£ (1975), 155. 206

Fig.l: Dissolution of copper as seen under hot-stage microscope: (a) copper particle just before melting in iron environment, (b) onset of melting, (c) pore created as result of Cu melting, (d) first stages of penetration into Fe matrix. 207

Fig.2: Microstructure of the Fig.3j SEM + EDAX picture after alloy after 10 min of 10 min of sintering at sintering at 1100°C. 1100°C; frott top to bottom: Fe profile, scanning line, Cu profilc- Cu background.

rwliiJ^r^WK

Fig.4: SEM + EDAX picture after 12 hr sintering at 1100"C; structure, and concentration profiles from i£5iiSZs~Z3jM top to bottom: Fe, Cu, Cu background. Fig.5: EDAX: Fe, Cu and Cu background profiles after 10 min of sintering at 1100°C. 208

Fig.6: (a) SEM: structure after 6 hr of sintering at 1100°C, (b) EDAX: Fe and Cu profiles for scanning line in (a), 209

HEAT TRANSFER TO WATER AND ITS IMPORTANCE FOR METAL CASTING AND HEAT TREATMENT

* ** M. Bamberger and B. Prinz * Technion - Israel Institute of Technology

** Metallgesellschaft AG. West Germany

INTRODUCTION Cast metal properties depend, i.a., on the rate of cooling during solidification or on heat treatment. A too high rate of cooling down will lead to thermal s^ressts in the product and consequently to cracks or fracture; on the ether hand, in case of too low a cooling rate, the special casting process may lose its metallurgical or economical advantages or be a complete failure. Consequently, major importance should be attached to the recognition of coefficients of heat transfer in relation to different cooling methods, in order to calculate the rate of cooling and their adaptability to the production process. The preferable method of control over the cooling process is continuous casting. Therefore, most researches deal with cooling methods of this casting, i.e., either by immersion in water or by water spraying. In this article we are going to present a method of measuring heat transfer to the cooling agent, and the use of continuous casting and heat treatment.

THE PRINCIPLE OF COEFFICIENT OF HEAT TRANSFER It is impossible to directly measure the heat flow from a hot metal body to a cooling agent. Therefore the calculation is carried out by measuring the body temperature and computation of the temperature field. From these two the heat transfer from the product to the cooling agent is deduced. The experimental set-up is designed in such a way as to establish conditions of heat transfer in one direction, i.e. the equation, which defines the conduction of heat in a body, in which the temperature is being measured, is:

where t, X, a, and T are time, coordinate, thermal diffusivity and temperature, respectively.

In lieu of a continuous description of the temperature field we may assume the final number N as temperature at discrete points. The heat balance at every point in accordance with differential equation equals:

where T1 - the previous temperature at point i, and T. , T., T. the present temperature at points i+1, i, i-1. 1 210

At the boundary there exists the relation:

T -T N *N1 X 0 K AX ~ Q With the usual solution 0 is known at the boundary, therefore we have N equations with N unknown temperatures, which will have to be found by solving a set of equations.

With the cooling experiments the temperature is measured at one point T , i.e. with the set of equations we do not know N-l temperatures ana the heat flow at the boundary.

The solution of the set of equations will give us the temperature field of the product and the heat flow. Control of solution quality is carried out by means of comparing the calculated temperature field with the control temperature measured in the body.

With this method the coefficient of heat transfer is calculated by means of convections and radiation. In fig. 1 we recognize that this is in accordance with the heat transfer coefficient, which is calculated by the Sheplan-Bolzman Law. Hence we conclude the method of computation to by correct (1, 2).

HEAT TRANSFER COEFFICIENTS FOR WATER COOLING Heat transfer coefficient for bars of nickel, aluminium and copper, when immersed in water, as a function of the surface temperature of the bar, we can see in fig. 2. A definite dependence on the properties of the bar, which is in a state of cooling off, can clearly be seen: very little with nickel, a little more with aluminium, and the most intensive cooling with copper. Nevertheless, the dependence on surface temperatures is identical in these three cases.

In the a/m cases the immersion in water is carried out at 20°C; however, it is well known from literature and heat treatment experiments that the temperature of the cooling water influences the rate of cooling. Therefore experiments with water immersion at various temperatures have been carried out. In fig. 3 we see that the heat transfer coefficient of a copper bar immersed in water decreased, when the cooling water temperature goes up. An increase from 20°C to 60°C brought about a drop of 30% in cooling capacity. Furthermore, it can be seen that the heat transfer coefficient is identical in all three cases.

Spraying water on a hot metal bar surface is accepted cooling procedure, the advantage over immersion in water being the possibility to control the cooling intensity by regulating the intensity of spray. In fig. 4 we show a heat transfer coefficient for cooling by spraying water on a copper bar. We see clearly an increase of the heat transfer coefficient together with an increase of the spraying intensity. Hence, the dependence on the surface temperature is identical with the one of the cooling process by immersion in water. 211

The heat transfer coefficient depends on the following factors: a) The property of the material to be cooled. b) Surface temperature. c) Temperature of the cooling water; d) Intensity of spray. Based on all previous experiments we arrive at the following formula:

"spray = 0"69 lo* <*5> ' f1"4^ exP ^T^T > + ^ + «rad

where 'a spray - heat transfer coefficient; V - density of spray (S./m2 min); X - heat conductivity of metal; p - metal density; c - specific heat of metal; 6 - average temperature of coolant; 6 - surface temperature of metal; 0 - evaporation temperature of cooling liquid; a - heat transfer coefficient during vapour-co a , - heat transfer coefficient during radiation, rad

RATE OF COOLING AND HEAT TREATMENT It is an established fact that we can slow down the rate of cooling by adding organic matter to the cooling bath. The heat transfer coefficient drops to 15% of its value during water cooling due to an addition of 3% of organic matter and arrives at 8% of its immersion-in- water-value, when 10% of organic matter is added.

This change is utilized to control the rate of cooling, when hardening - as shown in fig. 5 - a steel plate 30 mm thick.

Quenching in water brings about a structure, which is almost wholly martensite . Adding organic liquid slows down the rate of cooling, and we, therefore, arrive at a different structure. The more organic liquid reached, the slower the cooling rate, and we have more ferrite as shown by the two curves in the diagram. For comparison there exists another curve relating to air hardening, where ferrite above 50% is achieved. 212

CONTINUOUS CASTING - RESULTS Continuous casting of aluminium includes cooling in a water-cooled metal mould and immersion in water immediately afterwards. From measuring the change in temperature of the cooling water in the mould, we derive the complete heat flow from the casting into the mould. It is well known that the heat flow is relative to the square root of dwelling time in the mould (3, 4) and thus it is possible to calculate the heat flow from the casting into the mould. During the experiment temperatures were measured from a section of an aluminium bar, having a diameter of 190 mm and a solidification front was arrived at. The actual measured solidification front and the calculated one are in complete correlation (see fig. 6) .

On the other hand there exists a discrepancy between the temperature measured at a distance of 125 mm and 80 mm from the centre and at the centre itself (fig. 7). The discrepancy is mostly up to 10%, except at the outflow area of the mould. This error could possibly be the result of a lack of exact knowledge of the thermo-physical properties of the material at high temperatures.

SUMMARY In this article we have shown the principle of measuring the heat flow from a metal bar to a cooling agent. This method could also be used for measuring the heat flow to various other cooling agents. A formula has been established for calculating the heat transfer coefficient as a function of cooled-off metal properties, cooling water temperature and surface temperature. During experiments a correlation was established between the measured temperature of a cast metal bar and the one calculated by utilizing an established formula in research. The latter fact proves the accuracy of the formulating the coefficient of heat transfer. It provides an optional planning method for the establi- shment of a cooling arrangement for continuous casting and heat treatment.

ACKNOWLEDGEMENTS The work described was carried out in the Metal1-Laboratories of Metallgesellschaft AG and with its financial assistance. We would like to thank the company for allowing publication of this work, the director of the laboratories, Prof. Dr. Ing. P. Wincierz, for promoting the investigation, and Mr. J. Fachinger and E. Braun for taking part in the execution of the experiments, as well as Mr. Dohl for his assistance in the evaluation of experimental results.

"REFERENCES 1. M. Bamberger, R. Jeschar and B. Prinz; Zeitschrift ftlr Mettallkunde Vol 70 (1979) 9, p. 553. 2. M. Bamberger, B. Prinz - The 47th Foundry Cong. Oct. 1980,Jerusalem 3. R. Alberny et al - Revue de Metallurgie - Juillet-Aout 1976,p.545. 4. G. Vogt and K. WUnnenberg - Klepzigfachbericht 80 (1972)10,p.491. 213

« Cu * Ni o Al o vtaiues of N Larrtoert and M Economopoulos for Oxidized Steel

i00 600 200 (00 500 600 KM 300 <>O0 S00 600 «• 800 900 1000 Surface temperature in °C »- Surface Temperature

105- Material- copper

Cooling in water

600 ' 1 1 > ' 1 ' 1 r- 100 200 300 400 500 600 70C 800 100 200 300 400 500 600 700 80C D Surfcce temperature TJO m*c —- Surface temperature -&0 in C

3 Immersion heat transfer coefficient as a ^ Heat Transfer Coefficient in Spray Cooling of function of cooling water temperature Copper as a Function of Temperature

e Il l - : J"

- •••£: \ Ic?""" • / « * a ic ii - - Distance From Center of Cost'ng ((.mi

- 6 CoTiJuted [JxJ Measured Sol

5. Cooling Curves of 30mm Thick Steel Plate loa ino 500 TOO Cooled by Immersion into Wpter , Water • Temperature in "c Organic Fluid nnd Air Cooling 7. Computed and Measured Temperatures

FAST, NON-DESTRUCTIVE ELECTROCHEMICAL DETECTION OF SURFACE INCLUSIONS IN METALLIC SUBSTRATES Israel Rubinstein General Electric Research and Development Center Schenectady, NY 12301

Identification of surface defects in engineering materials is of considerable importance, since such defects constitute a major factor in the determination of the fatigue life of a component. In this contribution a fast and non-destructive electrochemical method is presented for visual detection of surface inclusions in metallic substrates. The method is applicable to non-conductive inclusions, e.g., oxides, carbides, etc., in metals or alloys. It is based on electrochemical coating of the surface with thin (~1 ym) highly-colored layers. The non-conductive inclusions remain uncoated and bright, thus becoming highly visible on the dark background. Very satisfactory results are obtained by employing anodic polymerization of pyrrole in an electrochemical bath.^ Thie results in the deposition of a black, continuous pclypyrrole film on the substrate, in which the flaws appear as bright spots. The method was successfully applied'in the detection of -30-300 ym oxide inclusions in nickel-base superalloys. The resultsof typical test experiments are presented in Figure 1. The two types*of inclusions were (a) sputtered Al2O3 surface flaws, and (b) bulk AI2O3 inclusions, pre-mixed with the alloy by powder-metallurgy techniques. It is evident that the inclusions become clearly visible to the naked eye upon coating with the polymer.3

30 40 50 60 30

Figure 1. Left: Inconel-718 piece with six sputtered AI2O3 inclusions, coated with polypyrrole film. Inclusion diameter, in ym: 280 (lower, right), 200, 120, 90, 60, 30 (upper, left). Right: Rene-95 piece with bulk AI2O3 inclusions, coated with polypyrrole film. Average inclusion diameter, 270 ym. 215

References and Notes 1. Present address: Department of Plastics Research, Weizmann Institute of Science, Rehovot 76100 (Israel). 2. A.F. Diaz., J.I. Castillo, J.A. Loaan and W.-Y. Lee, J. Electroanal. Chem. 129, 115 (1981). 3. I. Rubinstein, J. Appl. Electrochem. 21> 689 (1983). 216

TELEPHONE TOKENS PRODUCED BY POWDER METALLURGY

Andy Sharon M.S.H. Sinter Enterprises, Kibutz Mefalsim H.P. Hof Ashkelon 79160, Israel

ABSTRACT The production of telephone tokens by PA' technique is surveyed.The described process was taylored as to render a good and lasting product which finds its use in every day life.Manufacturing procedure is described with emphasis on the P/M advantages over the common procedure of punching and coining out of a rolled metal sheet. INTRODUCTION The manufacture of coins,tokens and medallions by the P/M process is quite known(l), although its usage is still not widespread. The P/M has some outstanding advantages over the conventional process mainly when producing tokens of intricated geometrical shape.Telephone tokens are still used in some countries throughout the world.Although this system is inconvenient (tokens shortage), it is efficient in countries where high inflation rates cause frequent changes in prices. This paper emphasises the advantages of the P/M process over sheet punching and coining. THE TOKEN The Israeli token has an intricate geometrical shape with close tolerances. The slot on its diameter(Fig.l) and the 20' slope,makes it difficult to manufacture it by single punching and coining from a rolled sheet.By using the P/M process,the token was produced in almost its final shape during powder compaction applying relatively low compacting pressure. MATERIAL The material chosen was the cupronickel coinage alloy(75$ Cu - 255? Ni). The powder had the following chemical composition and physical properties; 74.655 - 75.155 Copper 0.025% Mn 2A.1% - 25.1% Nickel 0.06$ Sulphur 0.0155 - Carbon 0.16555 Iron Flow rate determined by Hall flowmeter - 28.0 sec/50g. Grain size; +150 microns - 1.2$ wt. +45 microns - 4-7/5 wt. -45 microns - 51.555 wt. The powder was fully prealloyed having an irregular particle shape(Fig.2). Prior to compaction,0.7$ Lithium Stearate as lubricant was admixed to the powder. 217

MANUFACTURING PROCEDURE Compaction 2 Green compacts were pressed at pressures varying between 5 and 7 tons/cm in a rigid steel die. Special adjustments had to be made in order to ensure the uniform density throughout the various sections of the compacted shape. Sintering Out of the various possible sintering cycles for the Copper-Nickel powder compacts,the highest possible temperature of 1120"C was finaly chosen.The effective dwelling time at the sintering temperature was established at 30 minutes.The protective atmosphere was dissociated ajnmonia. Coining Tumbling of the sintered parts render a smooth and shiny surface,a great contribution to the quality and appearence of the final product.The closure of the surface pores following this process,enhances the corrosion resistance of the P/M part up to the level of the one produced by casting and rolling of metal sheets. „ The coining process carried out at 10 tons/eir. renders a final product with a density of 98% of the theoretical density.The close tolerances specified by the customer were met.

MECHANICAL PROPERTIES The mechanical properties of the sintered Cupronickel coinage alloy were tested on standard M.P.I.F. tensile specimens. Comparing the U.T.S. and the Y.P. data for the sintered specimens with the coresponding data for the wrought material(3),it can be seen that these properties are higher in the sintered material.The elongation and impact properties are louver as can be expected for a sintered material (Table I).

Material Property Rolled Metal Sheet Sintered Material 2 U.T.S. (N/mm ) 290 300

2 Y.P. (N/mm ) 100 110 Elongation {%) 35 19.5

Hardness (HB) 85 60

Table I. - Mechanical Properties of the Sintered Material vs. the Wrought One.

MICROSTRUCTURE The microstructure of the sintered and coined material was compared to the wrought and coined one(Fig.3) The basic phase and twins inside the grains can be seen in both materials. The main difference is the mean grain size.The mean grain size of the sintered material is 45-50 microns compared to 200 microns in the rolled metal sheet. 218

This difference is of cardinal importance when the token has to activate an electronic device which identifies the composition of the material by frequency response such as in some vending machines and public telephones. This change in frequency is affected mainly by the number of grain boundaries - grain size present in the same geometrical shape, provided the porosity of the sintered part is low (less than 3%). This problem can be otherwise overcome by increasing the nickel content in the alloy, thus altering its paramagnetic properties. CONCLUSIONS This paper emphasizes the possibilities of the P/M technique to produce tokens and coins in large quantities at high densityes having suitable mechanical properties. By using this process the production cost per unit is considerably lower, having the following outstanding advantages: 1. No scrap left comparing to the punching process, 2. Tools life considerably longer - over 5,000,000 units for compacting tools and over 4,000,000 units for the coining tools. 3. Lower pressures required for compacting and coining enables the producer to make use of cheaper machinery. The versatility of the P/M process and the above mentioned advantages can push forward the usage of this technique for manufacture of coins or medallions.

REFERENCES 1. "Coins, Tokens and Medallions Made by Powder Metallurgy", W.V. Knoop and J.D. Shaw, Int. J. of Powder Metallurgy 3 (1) 1967. 2. Unpublished work - A. Sharon 3. DIN Standard No. 17670 4. "Laminated Cupronickel/Copper Coin Blanks From Metal Powders" B.G. Harrison and T.R. Bergstrom, Int. J. of Powder Metallurgy 3 (4) 1967. 219

SCALE 5:

Figure 1: Production Design for the Telephone Token

Figure 2:

S.E.M. Micrographs

of the Pov.'der Grains. 220

Figure 3: Mictostructure of a] Sintered Material b) Wrought material 221

AUTOMATIC DETECTION OF FLUORESCENT INDICATIONS

K.M. Jacobsen

Brent Chemicals International PLC, Ridqeway, Iver. Bucks., England

Fluorescent Penetrant Inspection (FPI) is a well established and proven Non Destructive Testing (NDT) technique which is used extensively in many industries. Its reliability as a means of finding surface discontinuities is affected by the correct choice of penetrant process materials, the equipment, and the visual acuity of the inspector.

A choice of penetrant materials may be made from a wide range of excellent commercially available products. Process parameters appropriate to the intended application are established by experience and experiment, and FPI system sensitivity is verified using test pieces with known defects as a point of reference.

Fully automated process lines which guarantee that pre-determined FPI process parameters are strictly and consistently observed, have been in operation for more than twenty years. This ensures that parts emerge at the end of the process in optimal condition for visual inspection under black light. The equipment has demonstrated its ability to accommodate parts of varying geometry and surface condition. The former is achieved by careful jig design which ensures that parts are positioned for optimal coverage, washing and drainage. Different surface conditions are dealt with by providing alternate process sequences, offering options to use penetrants of various sensitivities, variations in hydro- philic emulsifier concentrations and contact times, as well as a range of different wash times.

The point of greatest weakness and major source of unreliability of r"PI is the actual visual inspection under black light. Inspectors sit in the dark, and handle each part individually, carefully examining all its faces for fluorescent indications. An experienced inspector with 20/20 vision can perceive minute indications on the basis of which he makes an accept/reject decision, making use of the ability of the humar eye to distinguish high spots of intense fluorescence from the duller background fluorescence due to residual penetrant. Such a high level of performance can only be maintained for short periods, and is not guaranteed from day to day. It is not possible to monitor short-term fluctuations, which go undetected and can lead to indications being missed.

Real improvement of FPI reliability is possible only if the human eye can be replaced by an instrument able to detect indications in the presence of background fluorescence. 222

TM The Automatic Electronic Optical Scanner - AEOS - developed by Ardrox, a Division of Brent Chemicals International, has achieved that objective. A pre-production prototype machine has been built capable of inspecting automatically jet-engine blades and similar parts up to 25O mm long at a rate of about three a minute.

The basic principle is that of a flying spot scan produced by a collimated laser beam reflected off oscillating mirrors. Fluorescent light generated by the beam passing over entrapped penetrant is collected by photomultipliers. The signals are amplified and digitised and processed by a microprocessor. They are used to activate mechanical handling equipment and displayed on a TV monitor.

AEOS is a go:no-go system which rejects all parts with indications, however produced. These require subsequent visual inspection under black light. Typically, accepted parts requiring no further inspection represent j-7O-8O% of the through-put; with a corresponding reduction of the inspection workload.

The AEOS system may be considered as consisting of three parts: the inspection module, the mechanical handling module and the microprocessor. The inspection module consists of the lightproof inspection cabinet and optical module containing the laser.

The light source is a He-Cd laser producing a high intensity collimated beam approximately 1.5 mm in diameter, in a single wavelength of 442 nm, without any interfering spectral contribution.

The laser beam is reflected off oscillating mirrors. The first mirror oscillates about a horizontal pivot at a selectable frequency between lOO Hz and 400 Hz, so as to generate a flying spot in the vertical mode. The second mirror oscillates at a much lower frequency about a vertical axis and thus induces a horizontal drift and fly-back into the scan.

From the mirrors, the beam is reflected into the inspection chamber, where it arrives as a spot of 1.5 mm instantaneous diameter which scans an area at the required testpiece location of approximately 250 mm height and lOO mm width.

The spot diameter and mirror oscillation frequencies are chosen so as to ensure the illumination of all positions within the active area, which can be readily increased or decreased using suitable stops and baffles. To ensure that no stray visible light from the laser excitation mechanism gets into the system, a filter is mounted in the light path.

The testpiece is held in the centre of the inspection cabinet with its axis coincident with the central vertical scan line. When the beam momentarily passes over entrapped penetrant it causes it to fluoresce, emitting visible light, which is collected by an array of 6 photomultipliers placed so that they can pick up light from any point within 223

the scanned area. They are positioned so as to avoid the existence of shadow regions as might occur with a single photomultiplier. Optical filters remove any unwanted reflected or scattered light and ensure that only fluorescent light is recorded.

The Mechanical Handling System is designed to present all faces of the test object to the scan. A small robot removes the test object from the fixture on which it is held for penetrant processing, and locates it in a vertical holder. There it is rotated through 360 in steps, pausing at each one to allow the beam co scan each face. The number of steps is determined by the geometry of the test objects, three being generally sufficient for round objects, while four steps are necessary for objects having a more or less rectangular cross section. If required the rotation can be performed in a larger or smaller number of steps.

A horizontal gripper then ad^^nces and grips the test object midway along its vertical axis. The holder releases the test object and the gripper rotates it through 9O in a vertical plane presenting one end to the scanner, and through 180 to present the other end face. If the shape of the test object requires it, the angle can be varied to overcome any masking effect due to test object contours. The gripper then releases cind deposits the test object on a conveyor belt which carries it out of the inspection cabinet.

Fluorescent light collected by the photomultipliers is converted into an electrical signal which is first fed into its associated amplifier and then into a common adding amplifier, where the individual signals are combined for any instant in time. The signal then passes through an analogue to digital converter to the microprocessor, which averages the signal response over repeated scans, collecting a complete signal in time and dividing it into the elements of an M x N matrix,where M and N are numbers of the order of 250. This effectively divides each averaged vertical scan into M sections, and the horizontal scan into N sections.

The sensitivity of the system is determined by the Threshold Value which can be varied in 64 steps. If an element records a high signal level with reference to a predetermined standard - the Threshold Value - as commanded by the program, then an "indication" is registered. At its lowest setting it will cause the system to trigger when an extremely small quantity of light, such as may be due to a small amount of blue fluorescence of small particles of dust, is detected.

The co-ordinates of the <.;.fending matrix element are recorded, and the signal used to activate the mechanical handling equipment which manipu- lates the test object. For practical purposes the Threshold is set at a value which will ensure that the system is triggered by the smallest indication which the inspection is required to detect. This generally lies within the quartile above the lowe=t of the range. 224

The surface being scanned may be divided into a maximum of four separate areas to each one of which different Threshold Values can be applied. This facility permits the application of different acceptance standards within a face thus avoiding miscalls due to a higher than necessary Threshold setting for less critical areas, while maintaining the highest level of sensitivity where that is called for.

When an indication above the Threshold is detected, the microprocessor compares the response of each matrix element with those of its neighbours, and triggers a rejection only when the averaged detected signal is greater than that of the surrounding area, thus distinguishing wanted indications from background fluorescence. In a case where an indication extends over several matrix elements, this method would cause the signal to be treated as due to background and thus wrongly accept the test object. Provision is made for such an eventuality by writing into the program an overriding absolute ceiling value, which would cause the system to trigger on a signal of that intensity being detected, irrespective of the intensity of the signals of the surrounding area.

The outline of the test object is detected by separate photomultipliers sensitive to the incident wavelength reflected by the back of the inspection chamber when the scanning beam passes over the edges of the test object. The image of the test object thus produced is displayed on a television monitor in real time, and detected indications are made to pulsate in their true locations relative to surface geometry, facilitating subsequent more detailed inspection by other techniques. The images may be stored on a disc and re-displayed on a separate monitor as an aid to conventional visual inspection under black light.

AEOS has been subjected to intensive evaluation in aero engine plants. Compressor and turbine blades from production lines were double inspected by the system and skilled inspectors, and the findings recorded and compared. The trials demonstrated that the system will detect consistently very small indications in the presence of normal levels of background fluorescence. Therefore it satisfies the two major requirements for any NDT procedure, sensitivity and repeatability, and it eliminates reliance on the human eye for the detection of indications, thus improving greatly the reliability of FPI.

AEOS may be used to detect fluorescent indications irrespective of the method by which the indications are produced. Thus ferromagnetic objects, treated by a fluorescent magnetic particle process are equally suitable subjects for AEOS. 225

CRYOFORMING OF 301 AND 302 STAINLESS STEEL

T. Livni, S. Bar-Ziv, A. Rotem. Rafael, Haifa

A. Rosen, Technion, Haifa

INTRODUCTION

Metastable austenitic stainless steels transform to martensite when being plastically deformed. The volume fraction of martensite is a function of increasing deformation and decreasing temperature. At cryogenic temperature, relatively small amount of deformation is required in order to acheive almost a complete transformation From austenite to martensite. This results in a very drastic increase in the yield strength of the material. The effect in indeed used in manufacturing high strength components with good corrosion resistance [1,2]. The process is known as "Cryoforming" and is used for producing pressure vessals and springs.

The aim of this investigation was to establish the relationship between the parametersof the process and the mechanical and some corrosion properties of 301 and 302 steels.

EXPERIMENTAL

The experiments were carried ^ut on three types of specimens: A. Flat tensile specimens, made of AISI 301 sheet, (1.6mm thick, 6.25mm wide, G.L.=35mm). Specimens were taken in the longitudinal direction, and were vacuum annealed. (1010°C, 5min, Argon cooling). B. TIG Weidments(both longitudinal and transeverse) were made on pieces taken from the same sheet tensile specimen were prepared, the same way as in. A. C. Round tensile specimens, 6mm, were prepared from AISI 302 annealed rods.

All specimens were plastically deformed,in tension,to various amounts (10 to 21 pet.). Cryogonic tension device, containing liquid Nitrogen (-ig6°C) was used (Combmed with an Instron machine). Two strain rates were employed - 0.25x10 2 and 0.25x10 3sec"1. Following this "Cryoforming", martensite content was measured by x-Ray diffraction method, and tensile properties of all specimens were tested in room temperature.

Cryoforming to 13 and 17 pet. elongation in type A and C specimens, respectively, resulted in optimal combination of tensile propertiess (cTy=l50kg/mm2, P+=10%). Fracture toughness (on type A specimens), Salt spray corrosion resistance (on type A and B) and notch senstivity - NTS/oy (on type C) tests were carried out on specimen cryoformed to acheive the above mentioned properties. 226

RESULTS

Tensile tests - uriwelded material Fig. 1 demonstrates a typical stress-strain curve of specimen while being deformed in L.N., together with the variation of martensite content with the plastic strain. Immersion in L.N. resulted in imidiate formation of 5 pet. martensite. No increase in martensite content was detected in specimen deformed bellow 2.6 pet. Further elongation causes martensltic transformation. After 20 pet. elongation, an almost fully martensitic material is accepted.

Room temperature tensile properties of specimens deformed cryogenically to various amounts of elongation are shown in Fig. 2. Since the strain rate during the cryoforming process had only a negligable effect on these properties, Fig. 2 demonstrates results obtained from cryoforming in one strain rate (0.25x10 ^"1)

The insensitivity of the martensitic transformation progress to strain rate in this steel was confirmed by Bakofen et al [5] for room temperature deformation.

180-

170-

150-

IJO-

I 1—1—r O3 S IO 15 2O T ' ' ' ' CRYOG€NIC PLASTIC DEFORMATION |%|

Fig 1. Typical Stress-Strain Curve of Fig 2. Room Temperature Tensile Cryogenic Deformation and the Properties of Sheet (Type A) Respective Martensite Content. Specimens Cryoformed to : Various Amounts of Elongation. 227

Few conclusions can be drawn from Fig. 2. Increasing the amount of cryogenic deformation, increases strength levels, and decreases the gap between 6y and ^u, e.g. strain hardening coefficient in reduced.

Cryogenic deformation bigger than 16 Pet. of sheet (type A) specimens results in very poor ductility at room temperature. For the reasons mentioned above, optimal elngation for Cryoforming process was chosen to be 13^14 pet. for strain rate of 0.2x10 2sec l.

It is to be noted that using the suggested cryoforming process, tensile properties similiar to those of Cjstom 455 (H-950 condition) steel can be obtained.

Welded specimens Room temperature tensile properties of welded specimens are shown in Fig. 3.

Oj; o»

190-

180-

170- . 6r

160-

150-

140-

130-

120-

7 8 9 10 11 12 1J 14 .K 7°

Fig,3 : Room Temperature Mechanical Properties of Welded Specimens, vs. Cryogenic Elongation.

For determination of yield strength, elongation was measured with extensiometer fixed in the close vicinity of the welded zone. The dependence of mechanical properties on the amount of cryogenic deformation was found to be similar to that obtained for unwelded specimens. The major difference was that slightly smaller cryogenic deformation in welded specimens was required in order to obtain a given yield strength. 228

It is to be noted that welding has no deterious effect on mechanical properties. For a given value of yeild strength, similiar values of elongation are obtained, both for welded and unwelded material.

Specimens were fractured randomly - inside or out of the weld zones.

Toughness and corrosion data Fracture toughness (Kc) and notch sensitivity (NTS/ay) data was obtained, for specimens crvoformed to yield strength of 150kg/mm2. NTS/ay tests were carried out on type C specimens, Kc tests were performed on type A specimens, fatigue precracked in L-T direction.

Results are summerized in Table 1.

Cryoformed Steel Custom ^55 Steel a =150kg/mm (H-950)

MTS/ay (Kt=A.7) l^7 1.53

Km [ (kg/mm2) l'mm] 100 260

Salt spray corrosion tests were carried out on annealed material, and cryoformed specimens (both welded and unwelded). No difference was observed in the corrosion behaviour of the different specimens after 3 month period.

DISCUSSION

Tensile properties of the cryogenically formed steel are clearly a function of the martensite content.

In the current literature there are two different approaches to describe the strengthening effect of martensite in partially transformed steels. One approach claims that the martensitic phase acts as dispersoid, pinning dislocations and creating unsurmountable barriers [3]. The other approach simply assumes that the anstenite-martensite mixture behaves like a composite material. In order to examine which of the above theories applies, room temperature yield strength of cryoformed specimens was plotted vs. the volume fraction of martensite, see Fig. k.

The linear relationship between yield strength and martensite content is clear, and can be described in the following equation, which applies to composite materials:

wo- Fig. k: Yield Strength vs. Volume Fraction of Martensite in Cryoformed Specimens.

*>-

—r~ 1 1— T~ —I 20 to 4o UO

V.iliiii. Ir.nri ii •"" "'.in

The martens!tic transformation was found to be strain induced, since no change in martensite content was observed in the elastic region.

The instability of austenite - i.e. the ease of cryoforming induced transformation can be derived from the deformation - transformation function, which can be expressed by the following equation [A], f =AeB 1-f

where f,,, is the volume fraction of martnesite, e is the amount of plastic strain, and A and B are material's constants.

"A" represents the ease with which on austenite structure can undergo a strain induced transformation to martensite, and "B" represents the measure of the autoc^talitic nature of the transformation.

Fig. 5 exhibits the variation of fm/'-fm vs. e in a log-log scale. The plot is indeed linear, and the values of material's constants are:

B=2.38 230

Fig. 5: log-log Curve of the Variation of Martensite Content with Cryogenic Strain.

A=45O B = slope=2.38

Col O.1 1. log STRAIN

CONCLUSIONS

The amount strengthening of metastable stainless steel is governed by the volume fraction of martensite, formed during cryogenic deformation. The dependence of room temperature yield strength on martensite content is 1inear.

The principle results of the investigation are: 1) There is a possibility to obtain high strength steels with relatively good ductility and toughness, combined with good corrosion resistance characteristics. 2) Welding has no detertous effect on mechanical properties of cryoformed materials. 3) The cryoforming process can be controlled by the amount of cryogenic deformation, in order to achieve desired mechanical properties.

REFERENCES

[I] Packner, Bernstein - Handbook of Stainless Steels, McGrow-Hill 1977, p. 47-6. [2] NASA Cr 61251. 231

[3] Mangonon P.L., Thomas J.G., Met. Trans. Vol. 1 1970, p. 1577- [k] Ludwigson D.L., Berger J.A., J.I.S.I. 1/1969, p. 63. [5] Powell, Marshall, Backofen - Trans, of ASM Vol. 50, 1958. 232

AUSFORMING OF H-ll MOD. STEEL

G. Rlkabir* and A. Rosen**

* RAFAEL, Armament Development Authority P.O.B. 2250, Haifa ** Dept. of Mat. Eng. - Technlon

INTRODUCTION Ausforming is a Thermo-Mechanical Treatment (TMT) intended to improve mechanical properties of steels. Generally, ausformed steels have a better combination of strength and fracture to toughness. In addition, ausforming allows use of relatively simple steels instead of expensive ones, such as maraging steels.

The aim of this investigation was to study the ausforming process in order to use it in the future in manufacturing real parts such as rocket motor casings. The H-ll Mod. steel was selected since its use in ausforming technology in widely known. Although there is no plan today to use this technology in producing thin walled tubes due to lack of very expensive foundations, it is a very interesting and important alternative.

EXPERIMENTAL H-ll Mod steel, having the composition of 0.38%C, 4.94%Cr, 1.20%Mo, 0.60%V, 0.17%Ni and 0.79% Si have been supplied in the form of a 10 mm thick plate. After the standard acceptance tests the plate was cut into pieces of 120imr. long and 45 mm wide. Ausforralng was performed by means of hot rolling. The details of the process are listed below:

a) Austenization treatment (20 min. at 1050°C). b) Fast cooling to 530°C. c) Hot rolling at 530°C to 40%, 60% and 80% reductions. d) Oil quenching. e) Double tempering (60 min. at 450°C).

All the platelets were grinded to a final thickness of 1.3mm in order to obtain scale and decarb free, flat and smooth surfaces.

The following tests were performed:

(i) Hardness measurements, (ii) TensLle tests, (iii) Metallography and (iv) Fractography.

For the sake of comparison two alternative treatments were also carried out:

(a) Oil quenching from 1010°C, double tempering at 510°C for 2 hours each. According to the literature this treatment results in maximum strength. 233

(b) Oil quenching from 1010°C, double tempering at 620°C for 2j hours to a hardness of 35 Re and cold rolling to 40%, 60% and 80% reductions. RESULTS

Mechanical properties.

The following table suramerises the average properties for the three different treatments. For the case of ausforming and cold rolling treatment we report only the results obtained after 80% reduction, since the lower reductions gave inferior results.

Treatment °u £ Hardness M?a MPa % Re

Ausforraing 2530 2610 3 65

Quenching 1980 2040 5.5 56 and Tempering

Pre-tempered 1580 1600 5.5 48 and Cold rolled

Metallography

The microstructure of the quenched and tempered specimens is normal tempered martensite. The former austenitic grain size was ASTM 5-7. No special features were observed. The pre-tempered and cold rolled specimens exhibited a fiberous, heavily deformed microstructure. The ausformed specimens showed a rather complex microstructure of branched deformation bands and fine dispersion of carbides. It was impossible to measure former austenitic grain size.

Fractography

The fracture surface of the quenched and tempered specimens can be characterized as quasi-cleavage, i.e., a mixture of smooth cleavage surfaces and wide, shallow dimples. The pre-tempered and cold rolled specimfiis exhibited a typical brittle fracture, characterized by severe delamination. The fracture surface of the ausformed specimens were similar to that of the cold rolled specimens, however, delamination was less severe and the fraction of the dimpled areas was larger.

DISCUSSION

This investigation has proved without doubt that ausforming can significantly improve the strength of H-ll Mod. steel compared to standard heat treatment, or to pre-tempering and cold rolling. Actually, the yield strength and UTS obtained after ausforraing of this steel Is higher than that of raaraging steels. 234

The increase of strength is accompanied by a certain reduction of ductility which is probably unavoidable at these high strength levels. On the other hand it is possible to temper the ausformed steel at a higher temperature than was done in this research (450°C) and obtain higher elongation on account of losing strength.

At this stage of the investigation we do not have unough structural evidence which can point out the reasons for the improvement of strength. One has to study the micros tructure in much finer details, e.g., by transmission electron microscopy, which is quite out oE bounds for this project.

The results reported in this paper give the best combination of strength and ductility. Much more combinations were studied, such as various degrees of reductions by rolling, different quenching as well as ausforming temperatures and finally various tempering temperatures. The important fact is that variation of the above factors did not result in marked changes of strength. This fact is encouraging since upscaling of the process will require certain tolerances of heat treatments and deformation temperatures.

CONCLUSIONS

The recent investigation of TMT of H-ll Mod. steel reviewed the following tendencies:

1. Ausforraing under certain conditions improves both yield and ultimate tensile strength by 30% but causes an almost 50% decrease in tensile elongation, compared to quenched and tempered material. 2. The microstructures and fracture surfaces of ausforraed or quenched and tempered materials are similar. 3- Small variations of ausforming conditions do not effect the end- strength of the specimens. 235

THERMODYNAMIC AND KHST3TIC PHENOMENA IN ADSORBED LAYERS

M. Grunze

Laboratory of Surface Science and Technology and Dept. of Physics and Astronomy, University of Maine Orono, Maine 04469, U.S.A. and Fritz-Haber-Institut der Max-Planck-Gesellschaft Faradayweg 4-6, 1000 Berlin 33, West Germany

ABSTRACT The relationship between thermodynamics and kinetics of adsorption will be outlined for the system N_/Ui(11O), for which an extensive study involving several modern surface science techniques were performed over the last few years. With respect to kinetics of adsorption, a brief summary of "precursor" concepts and their experimental tests will be given for molecular nitrogen adsorption on nickel and rhenium surfaces. Finally, seme recent experimental and theoretical results for the identification of the intermediate states in nitrogen dissociation on an Fe (111) -surface will be presented. INTRODUCTION The elementary steps in the interaction of gases with solid surfaces are studied in many laboratories to provide an understanding of the basic physical and chemical phenomena relevant to modern technology, e.g. heterogeneous catalysis, materials modification, corrosion resistance or microelectronics. Equilibrium measurements of the surface coverage in the respective gas ambient are carried out to determine the heat of adsorption (Q) and the entropy in the adsorbed layer (S ,), both important quantities in the description of the adsorbate phase. TRe isosteric heat of adsorption (Ogm) is derived by equilibrating the chemical potentials of the gas phase (y ) with the chemical potential of the adsorbate phase (y^) and solving for the temperature dependence of the equilibrium pressure at constant coverage 8/1/: 0 = ^§ . (1) Equation (1) applies whether or not the properties of the substrate change upon adsorption, since all changes in enthalpy of the system are included in Q. However, thermodynamic equilibrium between the solid and the gas has to be established, which also requires that gas and solid have the same temperature. This requirement is often neglected in adsorption studies on single crystals, although it was discussed in the early literature by Ehrlich /2/. As shown elsewhere, using kinetic arguments /3/, the error introduced by a temperature difference between gas and solid is typically small ($ 2 % of Q ), provided the sticking coefficient of the gas does not depend on gas temperature. Fran the isosteric heat of adsorption, the temperature of the system 236

and the equilibrium pressure the differential (and integral) entropy of the adsorbate layer can be calculated. The differential entropy is given by _E - (2) bad = Sg nea'.ecting again a possible temperature gradient between gas and solid /A/. The entropy in the adsorbed layer is a particularly useful quantity, since it allows a comparison with statistical mechanical models of the adsorbate phase /1/. In steady state equilibrium between solid and gas the rates of adsorption and desorption have to be equal. The respective activation barriers for adsorption and desorption determine the adsorption energy. Fran the knowledge of the absolute adsorption and desorption rates, the equilibrium constant K, and thus the free energy of adsorption can be calculated. Thus, the Jcmetic results can be compared to the equilibrium data. The above thermodynamic and kinetic quantities include only indirect information on the interaction between the gas species and the substrate and on lateral interactions between the adsorbed particles. For a complete description of the adsorption system and a theoretical approach in a model calculation, complementary data on the structure of the adsorbate phase and the bonding geometry of the molecules are required. This information can be obtained by surface sensitive spectroscopic techniques such as low energy electron diffraction (LEED) /5/, x-ray and uv-photoelectron spectroscopy (XPS and UPS) /6/ or vibrational spectroscopy (Infrared- Reflection Absorption Spectroscopy, IRAS /!/ or high resolution electron energy loss spectroscopy, HRELLS /!/).

MOLECULAR NITROGEN ADSORPTION ON Ni(llO) Molecular nitrogen adsorption on Ni (110) has been studied by many modern surface science techniques over the past years and probably is by now one of the most thoroughly studied adsorption systems, not considering the extensive data available for OO-adsorption on metals. Although molecular N2 adsorption on Ni(110) might seem to have no direct technical re- levance, the adsorption of N_ comprises one of the elementary reaction steps in the ammonia synthesis reaction or in the nitride formation of metals. We will discuss this adsorption system in order to outline the thermodynamics and kinetic phenomena occurring in adsorbed layers. In Fig. 1, we show the isosteric heat of adsorption determined by eq. 1 from several sets of adsorption isobars using the N1s core level photoemission band as a monitor of surface coverage /8/. In addition, the change in work function, the attenuation of the Ni-d band emission due to N2 adsorption, and the decrease in the low energy ion scattering signal from the Nickel substrate upon N^-adsorption were used as N2 coverage monitors to record adsorption isobars /9/. The isosteric heat of adsorption determined with the above techniques (except for the work function results /9/) falls well within the error bars of the N1s-core level data, and thus gives a consistent picture of the energetics of adsorption. As shown in Fig. 1, Q remains approximately constant (Q ~42 kJ/mole) up to e =0.47, and tfien drops steeply to a value of Q_~ 20 kJ/mole. The coverage calibration was obtained from LEED, the steep drop in Q occurs at completion of a (2x1) over layer. This (2x1) phase then transforms, with increasing coverage, into an incannensurate c( 1.4x2) solid phase via an intermediate fluid phase /10/. Such a commensurate-fluid-inconitiensurate 237

Fig- 1. QST for N- on Ni(11O). 60-

iO- o solid phase transition 40- on a rectangular substrate lattice was 30- postulated by recent theories of two-dimensional phase transi- 20- tions /11/, and was first observed experi- 10- mentally in the KL/ Ni(11O) system/10/. Thus, the change in 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 QST at 6 ~O.5 is associated with a two- dimensioual phase transition and we can expect this phase transition to be also reflected in the entropy of the adsorbed phase and in the kinetic data. ^ Figure 2 shows the differential entropy in the adsorbed N- layer S, (eq. 2) as a function of coverage for Ni(11O) /9, 12/ and Ni(10O) /3/. te include the KL/Ni(1OO) data /3/ to demonstrate that the entropy reflects the specific changes in degrees of freedom in the adsorbed phase and is thus sensitive to the structure cf the substrate and the change of lateral interaction between the adsorbed molecules. On Ni(1OO) , the isosteric heat is similar to bhat on Ni(11O) up to a coverage of 5.5x10 molec/cm (corresponding to 8 = 0.5 on Ni(11O)). It remains constant almost up to saturation coverage, when a drop in Q__ is observed, most likely due to the formation of antiphase domain walls in the c(2x2) over layer /3/. As mentioned in the introduction, the entropy in the adsorbed phase can be correlated with statistical mechanical models /1/. We plotted the differential entropy calculated by assuming /3/ a) a lattice gas model, considering only the configurational entropy in the adsorbed layer plus the vibrational entropy associated with the six normal modes of a N_-Ni surface complex (A) and b) a two-dimensional (mobile) Volmer gas, including the vibrational modes of the individual molecules (o). Clearly, both model calculations do not explain the experimental data, in particular the high coverage region where the commensurate-fluid phase transition occurs. Presently, there are no statistical models which explain in detail the high entropy in the fluid phase. For the coverage region below 6 ~ 0.5, we concluded that collective excitations in the adsorbed phase contribute to the high entropy, since considering only the degrees of freedom of individual molecules cannot explain the experimental results /3/. The thermal desorption data recorded for the N?/Nj|11O) system show a single desorption peak up to 6 = 0.5 /9,13/. At higher initial coverages, a second peak at lower temperatures develops indicating a lowering of the activation energy of desorption. The activation energy of desorption up to 3 = 0.5 is, v/ithin the error bars, identical to the isosteric heat of adsorption, which means that adsorption of N? is a non-activated process. For the higher coverages, present data are not accurate enough for a com- 238

Fig. 2. Differential entropy in the adsorbed layer for N- on Ni(110)(«), Ni(100)(°) and statistical mechanical calculations (see text). 250-

225-

200 parison of activation energy of 175- desorption and isosteric heat. By integration o± the desorption _ 150- traces, a plot of coverage versus exposure can be made, from which 125- the sticking coefficient, s = f(9), can be calculated; 100- 89 S(6) = 75- 3Ex where Ex is the integrated flux of 50- molecules impinging onto the surface (exposure). 25- In Fig. 3, the sticking coefficient as a function of ex- 0.8 1.6 2.4 3.2 4.0 4.8 5.6 6.4 7.2 8.0 posure is displayed. S is con- molec stant up to 9 = 0.48, then drops cm2 to a value of s ~ 0.3, increases again and diminishes at saturation coverage. This behaviour is not only observed at low temperatures, but is also evident at all temperatures where the catmensurate-f luid phase change is observed in the LEED experiments. At the higher temperatures experiments involving adsorption and subsequent desorption cannot be carried out, since the residence time of the molecules on the surface becomes much too small. Therefore, we used detailed balancing arguments to extract the sticking coefficient from equilibrium measurements. Under equilibrium conditions, the rate of adsorption equals the rate of desorption:

r = s(e) ad 1.0' * i i

r, = v-N , (3) des ad CD

Fig. 3. Sticking coefficient for } • • , N2 adsorption on Ni(HO) as a .5 1 function of coverage (T = 87 K). s 0 • 239

Fig. 4. Sticking coefficient of ^ on Ni(110) as obtained from isobars +p = 1x10~ mbar,

= 5x10~ mbar, *" -7 on = 1x10 mbar.

The adsorption rate is given by the impingement rate multiplied by the sticking coefficent; the desorption rate is expressed1 in the usual Arrhenius form. An isobar gives the correlation between P, N and T (T = substrate temperature). Equa- tiSn 3 can be solved for s(6) and, knowing the desorption rate parameters exactly, s(6) can be calculated. Fig. 4 shows the result for s(8)(note, that s(8) is not an isothermal quantity anymore!) obtained from three isobars. The agreement with the s(8) data from the desorption experiments is surprisingly 0 good, considering that the data were taken in different UHV chambers and no normalization of p, T or N , was carried out for the two ixperiments. In particular, we note that at 9~ 0.5 also under equilibrium conditions a discontinuity in s(e) occurs, which thus is clearly related to the con- mensurate-fluid phase transition. The constant sticking coefficient up to saturation coverage can be explained by assuming a mobile "precursor" state, i.e. a molecule impinging onto the surface is not reflected into the gas phase, but moves over the surface for a sufficient time to accanodate its translational and internal energy (T = 300 K, T < 200 K) and becomes chemisorbed. A distinction is made^Between "intrinsic" and "extrinsic" precursor states, the former trapped over the bare surface, the latter on top of a chemisorbed layer. The kinetic formalism relating these precursor states to the experimental sticking coefficient has been reviewed in several articles /14,15/. A spectroscopic characterization of these precursor states was only carried out at low temperatures for a few adsorption systems where the chemisorbed state is formed by dissociation of the respective gases /16,17/. For dissociative adsorption the "precursor" state is identified as the molecular state, which loses its translational and internal energy before dissociation. The results of our study on precursor states for molecular N^ adsorption on Ni(100) and Re(0001) /18/ are schematically shown in Fig. 5. These experiments were carried out at T ~ 20 K, as a spectroscopic technique, XPS, was applied. On the clean metal (5a), the molecule adsorbs and converts within the time it takes to record a spectrum (ca. 5 min.), into the linearly bonded chemisorbed state, which is the state observed also at the higher temperatures. This result implies, that a possible activa- 240

(hi N Gas N Gas (b) 29 2 1

Gas

Phys. N 2 Layer

Fig. 5. Schematic model for precursor state in molecular N« adsorption on Nickel and Rhenium surfaces (see text). tion barrier between a metastable, frozen-in state (e.g. a molecule lying flat on the surface) has to be less than E ~ 3 kJ/mole /18/. As long as the coverage in the chanisorbed state is less than 2/3 of saturation cov- e erage ( rei <2/3} i molecules impinging on occupied sites will find empty chemisorption sites during the time they need to accanodate all their kinetic and internal energy, i.e. s remains constant. For 0 > 2/3, molecules are found to condense on top of the chemisorbed layer (5c) i.e. they lose their energy in the second layer before they find a chemisorption site on the metal surface. This results in a decrease of the sticking coefficient into the chemisorbed state, but the net trapping coefficient (chemisorption and condensation) ranains unity, and multilayer formation takes place (5d)„ Completion of the chemisorbed layer is hence only observed at exposures exceeding those necessary to complete a monolayer, at T > 80 K. Thus, a molecule chemisorbed in the second layer acts as an "extrinsic precursor". The heat of adsorption of this precursor in the second layer is about twice the heat of bulk condensation /19/, which increases the residence time on the surface considerably above the value expected for a purely condensed species. Our low temperature experiments explain the observed constant and high sticking coefficient for N2 on Ni(11O) up to half monolayer coverage. The discontinuity at 9 > 0.5 cannot be explained within the framework of present precursor models, since all these models would give a continuous change in s(6). Presently, computer simulations relating the kinetics of adsorption and desorption to the geometry and energy changes in the adsorbed layer at 6 > 0.5 are being carried out to develop a microscopic model for the behaviour of s(e).

N2 DISSOCIATION ON Fe(lll) Ammonia synthesis on iron is a structure sensitive reaction with the Fe(111) surface being the most active low index single crystal plane /20/. Using various methods, it has been demonstrated for clean iron surfaces, 241 that the rate determining step for the overall reaction involves nitrogen dissociation, since the subsequent hydrogeneration of atonic nitrogen to ammonia proceeds at much higher rates. The kinetics of dissociative nitrogen adsorption on Fe(111) have been studied by Ertl et al. /21/ in detail, and it was concluded that this process proceeds through a molecularly adsorbed "precursor" state (α-state) with a rather low activation barrier. Recently, it was shown /22/ that this α-state is preceded by a more weakly held physisorbed γ-state, which desorbs already at about 85 K. More im- portantly, the XPS/HRELLS study summarized below showed that the α-state is a ir-bonded molecule, with both N atoms interacting with the surface and thus forming the immediate precursor to dissociation /24/.

Fig. 6. N1S core level - c spectrum for a: the α-state, - 397 0 '-"-" •-:-••'.•,. b: γ-state, c: atonic β-state. tn . b '"''"' 39 C D [arb . In Fig. 6 we show the J- 4012 characteristic N1s spectra a> - of the three distinguishable c 405 9 ... a . -.' / nitrogen states on Fe(111). c - Spectra 6a shows the doublet structure of two bands at 405.9 and 401.2 eV below E^,, observed for the 412 406 400 394 weakly bonded γ-state. The energy belou E-FermI [ eU) observation of a doublet structure for a single species is due to final state effects in the photoemission process; the peak at lower binding energies corresponds to emission from a "screened" final state, the one at higher binding energy to emission from the "unscreened" final state /23/. A comparison of spectrum 6a with those of N_ bonded on Ni(11O)/8/ or Re(COO1)/18/ shows that intensity is transferred to the "unscreened" final state peak, indicating a weak interaction with the substrate which is also evident from the low heat of adsorption of Q $ 24 kJ/mole /22/. This T- state then transforms, as a function of time and temperature, into the «- state producing the XPS-spectrum shown in Fig. 6b. Isotopic exchange experiments show /21/ that, in the α-state, the molecular unity is preserved. The XPS-data and the HREELS results ,/24/ reveal, however, a strong interaction of the two nitrogen atoms with the surface. A comparison of

the vN stretching frequency observed by HREELS to those in ^-bonded inorganic dinitrogen complexes, .as well as CMDO-CI calculations of the N1s spectrum show that a) the molecule bonds with its molecular axis parallel to the surface, b) that bonding to the surface involves charge donation from the 1 IT orbital of the molecule to the substrate and a1"* back- bonding from the substrate to the molecule and c) that the charfe transfer into the antiJDonding (with respect to the dinitrogen bond) T^* orbital increases the N-N bond length to ~1.25 A , compared to the N-S bond distance of 1.1 A in the free molecule. Thus, the effective activation barrier for dissociation is lowered and dissociation of N_ is observed at low temperatures (T ~ 120 K), leading to the formation of the characteristic N1s-core hole emission of atonic R nitrogen (Fig. 6c). 242

Fig. 7. Potential energy diagram for N2-dissociation on Fe(111) as derived frcm experiment. inn]

From a careful study of the ft 2N., N1S spectra as a function of gas pressure, temperature and time, a potential energy diagram cis shown in Fig. 7 can be constructed. In the upper part, we show a schematic view of Y, a and atomic g nitrogen on the Fe(111) surface, the potential energy diagram includes the respective heats of adsorption and activation barriers. The individual rate constants for the reaction are indicated in the insert. The microscopic mechanism of dissociation of N~ on Fe(111) was also investigated theoretically by calculating the potential energy surface for this reaction /25/. The theoretical treatment is described in ref. /26/. It involves calculation of the total energy E of the system as a function of the height, h, of the molecule above the surface and the intramolecular distance, d /26/. The Fe(111) surface is approximated by a four atom cluster and the molecule is lowered onto the surface with different orientations of its molecular axis with respect to the substrate. The best agreement between theory and experiment is ob- N2/FE(1 tained when the molecule be- r comes adsorbed with its axis parallel to the surface in the precursor state for dissociation (Fig. 8).

Fig. 8. Potential energy surface, E (h, d), for N dissociationon Fe(111). The 0.5 dissociation path is indicated by the arrows, the equipotential lines have a separation of 0.2 eV. 243

In this intermediate a-state, charge is transferred into the \ * orbital ofQN2, leading to a calculated increase of the N-N distance of d 1.5 A.TSy surmounting a potential energy barrier of E.^ RJ0.2 eV = 20 kJ/mole (compared to experimental E^ = 28 kJ/mole) ,^he molecule dissociates and forms the atomic β-state. From this calculation it also follows that the height of the activation barrier, E^. , is critically dependent on the charge transfer into the TIT * orbital OPN2. Charge donation into this antlbonding molecular orbital by substrate electrons weakens the intramolecular bond, and thus decreases the activation barrier for dissociation. The simple dissociation reaction discussed above is the first example in which the microscopic mechanism of the heterogeneous fission of an intramolecular bond has been identified both experimentally and theoretically. From a general point of view, it may be concluded that partial occu- pation of antibonding molecular orbitals by substrate electrons comprises one of the crucial steps in dissociation. Details of the geometry and electronic overlap in the "precursor" of dissociation will, however, depend on the particular system under study.

ACKNOWLEDGEMENTS The experiment described were performed at the Fritz-Haber-Institut der Max-Planck-Gesellschaft and at the Free University of Berlin. Finan- cial support from the Deutsche Forschungsgemeinschaft, Sonderforschungsbe- reich 6, is gratefully acknowledged. REFERENCES 1. A. Clark, in "The Theory of Msorption and Catalysis", Academic Press, 1970, and other textbooks. 2. G. Ehrlich, J. Chem. Phys. 36 (1962) 1499. 3. M. Grunze, P.A. Dowben, and R. Jones, Surf. Sci., in press. 4. M. Procop and J. Vblter, Surf. Sci. £7 (1975) 514. 5. see for example: M. Henzler, "Electron Diffraction and Surface Defect Structure", in: "Electron Spectroscopy for Surface Analysis", ed. H. Ibach, Topics in Current Physics, Springer Verlag 1977. 6. B. Feuerbacher and B. Fitton, "Photoemission Spectroscopy", ibid. 7. H. Froitzheim, "Electron Energy Loss Spectroscopy", ibid. 8. M. Golze, M. Grunze, R.K. Driscoll, and W. Hirsch, Appl. Surf. Sci. 6^ (1980) 464. 9. M. Grunze, M. Golze, and W.N. Unertl, to be published. 10. M. Grunze, P.H. Kleban, W.N. Unertl, and F. Rys, Phys. Rev. Lett. 51 (1983) 582. 11. see for example: S.N. Coppersmith, D.S. Fisher, B.I. Halperin, P.A. Lee, and W.F. Brinkmann, Phys. Rev. Lett. 46 (1981) 549. 12. M. Grunze, M. Golze, and W.N. Unertl, Proc. 3rd Symp. on Surface Science, Obertraun/Austria 1983, ed. P. Braun, HTU-Druck, Wien, p. 218. 13. M. Golze, M. Grunze, and W. Hirschwald, Vacuum 31_ (1981) 697. 14. D.A. King, in "Chemistry and Physics of Solid Surfaces", ed. R. Vanselow, Vol. II, CRC Press, (1979). 15. A. Cassuto and D.A. King, Surf. Sci. 102 (1981) 388. 16. G. Ehrlich and F.G. Hudda, J. Chem. Phys. 35 (1961) 1421. 17. R. Gcmer, in "Field Emission and Field Ionisation" Harvard University, Cambridge, MA. 1961. 18. M. Grunze, J. Fuhler, M. Neumann, J. Behm, C.R. Brundle, and D.J. Auerbach, Surf. Sci., in press. 19. J. Fuhler, Diplcmarbeit, University of Osnabruck, 1984, and to be 244

published. 20. For a review, see: M. Grunze, in:"The Chenical Physics of Solid Sur- faces and Heterogenous Catalysis", ed. D.A. King and D.P. Woodruff, Vol. 4, Elsevier (1982), p. 143. 21. G. Ertl, S.B. Lee and M. Weiss, Surf. Sci. 114 (1982) 515. 22.a)M. Golze, W. Hirschwald, M. Grunze, and M. Polak, Proc. 3rd Symp. on Surfac3 Science, Obertraun/Austria 1983, ed. P. Braun, KTU Druck, Wien, p. 254. b)M. Grunze, M. Golze, J. Fuhler, M. Neumann, and E. Schwarz, Proc. 8th Int. Congress of Catalysis, Berlin (1984), in press. 23. K. Schonhammer und 0. Gunnarson, Surf. Sci. 89 (1979) 573. 24. M. Grunze, M. Golze, H.-J. Ereund, H. Pulm, U. Seip, M.C. Tsai, G. Ertl, and J. Kiippers, submitted. 25. D. Toraanek and M. Grunze, to be published. 26. D. Totnanek and K.H. Eennemann, Surf. Sci. 127 (1983) L111. 245

SUPPORTED SILVER CATALYSTS HAVE SOME IMPORTANT PROPERTIES IN COMMON WITH ROUGH SILVER FILMS: ELLIPSOMETRY AND RAMAN DATA.

P.H. McBreen*. D. Hall**, J. Lalman and M. Moskovits

Department of Chemistry, University of Toronto, Toronto M5S 1A1, Canada.

*Department of Materials Engineering, Ben-Gurion University of the Negev, Beer Sheva, Israel.

Deparment of Chemistry, Ben-Gurion University of the Negev. Beer Sheva, Israel.

ABSTRACT Rough silver films absorb much more visible light than do smooth silver films. Ellipsometry data for this phenomenon may be interpreted by attri- buting a strong optical absorption feature, centered at 550nm, to the optical properties of a surface roughness layer. This feature is at least partially the reason why the Raman signal for adsorbates on roughened silver substrates is often up to a million times larger than expected. The so-called SERS (Surface Enhanced Raman Spectroscopy) effect is used in the present study to observe the behaviour of carbon on the surface of silver catalyst particles.

INTRODUCTION

Six years ago the surprising discovery was made that Raman spectroscopy could be used to study adsorbed species. Previously it was believed that the combination of the low cross-section for Raman scattering and the small concentration of molecules probed at the surface (10 /cm ) would preclude its use as a routine surface spectroscopic technique. However, it was found that on some surfaces the observed Raman scattering signal associated with an adsorbate was often a million times larger than expected. This enormous effect is exploited in surface enhanced Raman spectroscopy (SERS), and has been the subject of intense study in recent years. In a SERS experiment a monochromatic beam of laser light is directed at the sample. Some of the scattered light is collected and is analyzed on the low frequency side of the incident beam. Observed sharp spectral features may be directly related to the excitation of molecular vibrations through the Raman scattering process. Such spectra may be used to determine the identity and the structure of adsorbed species. Since, in general, SERS involves a beam of visible light as a probe it has the important capability of being suitable for in-situ studies of electrode or catalyst surfaces. The magnitude of the surface Raman enhancement is dependent on the substrate metal; the effect is most pronounced for adsorbates on silver. Further, there is a correlation between SERS activity and the structure of the substrate. Silver films deposited onto a cold ( ^lOOK) substrate are SERS active, whereas continuous silver 246

films prepared at room temperature are found to be SERS inactive. Data Dbtained using an ellipsometric technique show that the optical properties of the two types of silver film are markedly different. The film deposited at low temperature is strongly absorbing, compared to the room temperature film, in the 320-600nm region. Calculated spectra show (Fig.l) that the observed difference may be accounted for in terms of the difference in optical properties between a film with a smooth surface and a film with a rough surface. As illustrated in the inset of Fig. 2 the surface roughness is simulated by a layer of metal spheres resting on a flat surface. £ is the effective dielectric constant of the layer. The dashed line spectrum is very similar to measured B spectra for vapor deposition of a rough silver film on top of a smooth film. The (0) - and solid - line spectra are similar to B spectra measured for partially annealed rough silver films. Thus it appears that the SERS effect is related to the optical properties of the surface roughness layer. Such a roughness layer may be considered as a two dimensional aggregate of metal ^articles. With this in mind we wished to see if SERS could be applied to a thre^ dimensional aggregate of metal particles such as a dispersed silver catalyst. Supported silver particles are unique in their ability to catalyze the reaction of ethylene with oxygen to yield ethylene oxide. Fig. 1. Calculated spectra for the difference in optical absorbance between a silver film with a smooth surface and silver films with progressively (solid, circle and dashed line) rough surfaces. The quantitity B is related to the optical absorbance and is defined in reference 2.

EXPERIMENTAL

The experiments were performed using two types of supported silver samples which differed both in the nature of the support and in the manner of preparation. Type A. samples were prepared by a novel technique which enabled the electrodeposition of metals into the porous surface of anodized aluminium foil. Type B catalysts were prepared by filtering silver colloid solution through glass fibre paper. The silver impregnated 247

100 cm-1 3000

Fig. 2. SERS spectra of silver particles, of less than 2 urn diameter, on a glass fibre support. The solid line was obtained at 300K and displays peaks associated with hydrogenated surface carbon. At 460K [dashed line) the carbon related peaks are severely attenuated.

glass fibre paper was taken as a model Ag/SiC^ catalyst.

RESULTS AND DISCUSSION

In terms of SERS per-se the striking result of these experiments is that under appropriate condition the technique may be used to study silver samples at high temperatures. As shown in Fig. 2, SERS spectra were obtained for a sample held at 460K. In fact, type B samples which were reduced in H at 775K subsequently displayed surface Raman enhancement. Similarly, a type A sample which was heated to 770K in oxygen did not lose its SERS activity. In general SERS is not observed for silver samples at or above room temperature due to the loss of the requisite morphology through sintering. However, for type A and type B samples the support serves to minimize sintering. SEM pictures of a type B samples clearly show silver particles dispersed on glass fibres. The ability of SERS to probe such samples under conditions of high pressure and high temperature makes it an important tool for studying catalysis. For instance, the commercial production of ethylene oxide is carried out in the 520-600K range using silver particles supported on low surface area refractory materials, and should therefore be amenable to SERS studies.

We will now concentrate on the behaviour of carbon on the silver particles as revealed by SERS spectra. The observed intense carbon bands 248

at ^ 1370 cm" and ^ 1590 cm" (Fig- 2) are commonly observed in SERS studies of silver surfaces. Using SERS and EELS Tsang ct. al. have identified the features as belonging to amorphous (a) carbon. The feature at -1370 cm" correlates with the Raman spectrum of crystalline graphite, and the feature at ^ 1590 cm is observed in the IR spectrum graphite. Our SERS spectra indicate that the surface carbon is to some extent hydrogenated, as a broad CH stretching band is observed at ^ 2800 cm

All of the freshly reduced samples displayed the intense a carbon Raman bands. In the 450-550K range these bands disappeared even for a sample held under 1 atm. of C_H However, on cooling the sample to 450K the full intensity of the Raman bands was restored. It was noted that heat treatment, to 750K in the presence of l.atm. of C2H4, caused a marked sharpening of the feature at ^ 1590 cm" . This change may be due to ordering of the surface carbon as a result of increased mobility at high temperature. Such an interpretation is consistent with the dependence of the Raman spectrum of carbon on long range order .

Considering the surface science results obtained by a number cf groups it appears that the observed disappearance of the carbon Raman peaks in the 450-550K region is attributable to a change in the nature of the surface carbon. Madix has shown that the interaction of acetylene with oxygen adsorbed on Ag (110) leads to the formation of adsorbed C H at 170K. The latter species was then observed to transform to C~ ana amorphous C at 275 and 550K, respectively; a LEED pattern associated with adsorbed C? was found to vanish at 550K. In light of these results by Madix we attribute the disappearance of the a carbon Raman peaks at ^ 500K to decomposition to adsorbed atomic carbon. Thus the results for the supported silver sample are in good qualitative agreement with UHV studies of a well defined Ag (110) sample.

Before concluding we note that the observed SERS signals in the present study arise from adsorbates on the silver particles. This is confirmed by comparing the Raman spectra of blank and silver impregnated samples. In conclusion we have found that SERS may be used to study silver catalysts under realistic conditions. Further, the results that we obtained for the behaviour of carbon on supported silver particles are in good agreement with results for carbon on a silver single crystal surface.

REFERENCES

1. A recent review of SERS is given by R. Miles, Surf. Int. Anal. 5 (1983).43. 2. P.H. McBreen and M. Moskovits, J. Appl. Phys.,54 (1983) 329. 3. D. Hall and M. Moskovits, presented at the 7 Canadian Symposium on Catalysts, Edmonton, October 1980. 4. J.C. Tsand, J.E. Demuth, P.N. Sanda and J.R. Kirtley, Chem. Phys. Lett. 76 (1980) 54. 5. P.C^ Painter, 0. Mahajan and P.L. Walker, Jr., Extended Abstracts, 14 Biennal Conf. on Carbon, p. 113 The Pennsylvania State Universi- ty (1979) . 6. R.J. Madix, Appl. Surf. Sci. 14 (1982-83) 41. 249

XPS STUDIES OF Si FILMS DEPOSITED FROM SiCl. BY AN RF COLD PLASMA TECHNIQUE

E. Grossman, A. Grill and M. Polak

Ben-Gurion University of the Negev, Beer-Sheva, Israel

ABSTRACT

Microcrystalline silcon films were deposited in an inductively coupled glow discharge in SiCl.. XPS measurements show that the film has oxide layers composed of SiOx with 05x$2, 12 A wide. The chlorine concentration in the interface region is higher (5%) than that in the bulk (2%) , and it acts as a p-type dopant formiiv; Si~Cl bonds.

INTRODUCTION

In the past few years microcrystalline (me-)Si received increasing attention due to its attractive properties (1,2). The relatively high conductivity of the mc-Si can be very useful in amorphous (a-) Si solar cells as a conductive surface layer. a-Si solar ce'.ls are usually prepared from SiH4 because of the essential role of the hydrogen as a dangling bond compensator (3). However the use of SiCl- as a starting material, can lead to lower cost a-Si:H,Cl solar cells, with chlorine in the film acting as a dangling bond terminator (4-6) . The thermal stability of the chlorinated film is better than that of the hydrogenated a-Si, since Si-Cl bonds (91 kcal/mole) are stronger than Si:H bonds (77.7 kcal/mole) (6). Several works have been performed on a-Si:H,CI (5-9) with emphasis on the physical, electrical and optical properties of the films. In the present work we studied the surface and bulk chemical properties of Si films deposited from SiCl. in a rf plasma, by X-ray Photoelectron Spectroscopy (XPS). Special attention was given to the Si-SiO- interface due to its importance to the microelectronic industry (10) and to the effect of the CI in the film. All samples except one were exposed to air for several weeks.

EXPERIMENTAL

The Si coatings were deposited on a stainless steel substrates from gas mixtures of Ar.H^, and SiCl-, and B-H, or PH,, for dopant purposes. The plasma was initiated in a quartz reactor inductively coupled to a Plasma Therm rf generator of 27.12 MHz. The silicon films were obtained by the decomposition of the SiCl- molecules in the plasma and at the stainless steel substrate, and their reduction by hydrogen. XPS measurements were made with PHI model 549 system with double-pass cylindrical-mirror analyzer operating in the retarding-potential mode. The typical operating pressure in this system was 2x10" torr. Electrons were excited with a Mg k X-ray source operating at an accelarating potential of 10 kV and a power of 400 W. Arjon ion beam sputtering with 0.5 keV ions was used for 250

sputter profiling of the films. The microstructure of the film was determined by Transmission Electron Microscopy (TEM) (Jeol JEM-200B).

RESULTS AND DISCUSSION

The Transmission Electron Microscopy (TEM) studies indicated that the films were microcrystalline. Figure 1 shows a TEM micrograph of a Si Film deposited on a TEM grid. Figure la shows a bright field image of the film. A dark field image which was taken in the (111) debye ring is shown in figure lb. The bright areas have a crystalline structure while the dark areas are amorphous. Figure 2 shows a diffraction pattern and schematic diagram of the crystallographic planes of the Si film. The average size of the crystallites is 150 A. About the same crystallite size was determined by x-ray diffraction (9) . The following results were obtained for samples exposed to air for several weeks. Figure 3 shows representative XP spectra of Si 2p, before sputtering (a), and after several sputtering times. Spectrum (a)consists of two components, one with binding energy of 99.1 eV, which is due to unoxidized Si, and the other one at 102.7 eV, which is attributed to SiO,, (11). Sputtering reduces the oxide component. Figure 4 shows the difference between the spectra of the external layers (fig. 3 aid), and the spectrum from the bulk (fig. 3e). In the first difference spectrum (a-e) there is a shift of 3.6 eV between the binding energies corresponding to SiO- and the unoxidized Si. The peak is asymetric indicating the existence of other components at inteimediate binding energies. In (b-e) the main component is shifted 3.1 eV, and is attributed to Si-O, (12) Contribution of SiO can be seen in (c-e) with a shift of 1.9 eV (13). Accordingly the oxidation state of Si decreases *rith depth, the topmost layer being SiO,,. The overall thickness of the oxide layers was determined by using the expression for the ratio of the peak intensities of oxidized silicon and of the underlaying

r = 0.67 ISiC expf __J H L Uox CosS where: d - oxide layer thickness, X - electron mean free path, β -effective angle between normal to the surface and the electron take-off directions. In this case X was taken as 30 A aiiJ cos^=0.70. The experimental value gave d=12 A. Figure 5 shows the concentrations ratios of the oxygen and the oxidyzed silicon, as a function of the sputtering time, for two samples. The first was exposed to air for several weeks, while the other was exposed for 1 hour. For long exposure, PS seen above, there was a gradual change of Si oxidation state, starting with SiO . On the other hand, for short exposure to air, only a thin layer of SiO was formed. Figure 6 shows the depth profile of the component concentrations from the surface down to the substrate. At the top layer there is a sharp decrease in oxygen due to the removal of oxide layers from the surface region. The chlorine concentration is about 2% and it remains the same upto the interface with the substrate. At the interface there is an increase of the chlorine concentration to about 5%. This increase is probably due to interaction of the stainless steel substrate with the 251

chlorine ions in the plasma, at the beginning of the process. Once a thin layer of Si is formed, the substrate is no more exposed to the CI ions, and there is a reduction in the chlorine concentration to a constant value. In the interface region there is also an increase in the oxygen concentration, probably due to oxygen which is originaly bonded to the surface of the stainless steel substrate. From the chemical shift of the Si peak, measured at the interface res ion, we found that the Si is bonded to the oxygen, forming SiO- layer. The binding energy of the Si 2p in the bulk is 99.1 eV. This value is 0.5 eV less than the binding energy for a-Si:H. This shift should be due to the chlorine present in the film. Because of the relatively low concentration of CI (~2%), no chemically-induced shift is expected to be observed for the Si XPS lines. "Tierefore the apparent 0.5 eV shift should be due to displacement of the Fermi level in the gap. Thus, the chlorine presence causes a downward shift of the Fermi level, so that undoped a-Si:H.Cl is behaving like p-type semiconductor (14). Inspection of the chemical-shift observed for Cl(2p) (-2.5eV) indicates Si-Cl partially-ionic (0.5) covalent bonding.

REFERENCES

1. K.E. Spear, G. Willeke and P.G. LeComber, Physica, 117B $ 118B (1983) 908. 2. S. Hasesawa, S. Narikawa and Y. Kurata, Physica, 117B 5 118B (1983) 914. 3. M.H. Brodsky, M. Cardona and J.J. Cuomo, Phys. Rev. B16 (1476) 3556'. 4. J. Chevallier, S. Kalem, S. al Dallal and J. Bourneix, J. Non-Cryst. Solids, 51 (1982) 277. 5. V. Augelli, R. Murri, S. Galassini and A. Tepore, Thin Solid Films, 69 (1980) 315. 6. S. Kalem, J. Chevallier, S. al Dallal and J. Bourneix, J. Appl. Phys. Coll. C4 suppl. No. 10. 42 (1981) 361. 7. G. Bruno, P. Cpezzuto and F. Cramarossa, Thin Solid Films, 106 (1983) . 8. R.D. Plattner, W.W. Kruhler, B. Rauscher, W. Stetter, J.G. Grabmaier, Proc. 2nd E.C. Photovoltaic Solar Energy Conference, Berlin (1978) 860. 9. E. Grossman, A. Grill and R. Avni, Plasma Chem. and Plasma Proc, 2(4) (1982) 341. 10. W.Y. Chins, Physical Rev. B, 26(12) (1982) 6633. 11. B. Carriere, J.P. Derielle, Ch. Burggrat, Analysis, 9(5) (1981) 236. 12. CM. Garner. I. Lindau, C.Y. Su.P. Pianetta and W.E. Spicer, Physical Rev. D. 19(19) (1979) 3944. 13. T. Adachi and C.R. Helms, J. Electrochem. Soc, 127(7) (1980) 1617. 14. A.E. Delahoy and R.W. Griffith, J. Appl. Phys., 52(10) (1981) 6337. 252

Figure 1 TEM micrograph of plasma deposited Figure 2 microcrystalline Si film Ca) bright Selected area diffraction pattern field image (b) dark field image, and schematic diagram, at the taken in the (111) debye ring. microcrystalline Si film shown in figure 1.

Figure 3 Xp spectra of Si 2p, before (a), and after spattering (bfe) .

Figure 4 Difference spectra curves of the external layers (fig- 3 afd), after subtracting the spectrum at the bulk Cfig. 3e) .

• -Si •-a •-o 4 - SiO2

§70 Si2O3 long exposure I GO

SiO S4O 4 \ j short exposure §30 — 1 I— 20 10 0 •»-t--"-T"?-<" 10 20 30 O 3D 30 40 50 60 70 30 330 350 370 SPUTTER TIME.mn SPUTTER TIME,min Figure 6 Figure 5 Depth profile of the long- Oxygen/oxidyzed silicon concentrations exposed Si film as a function of sputtering time for the two atmosphereic exposures 253

QUANTITATIVE XPS OF HaS-ALUMINA: EVIDENCE FOR SODIUM ANISOTROPIC SEGREGATION

Y. Grinbaum and M. Polak

Department of Materials Engineering, Ben-Gurion University of the Negev, 3eer-Sheva, Israel

ABSTRACT The surface composition of two faces of the superionic conductor sodium β-alumina has been studied using quantitative X-ray Photoelectron-Spectroscopy (XPS). The most remarkable result concerns the occurance of sodium enrichment at one face as reflected by the relative line intensities of Na and Al from the two surfaces. This conclusion is supported by quantitative calculation of relative intensities expected for crystal faces as perfect bulk planes, and by detailed in-depth composition profiles. Possible mechanism for the !Ia anisotropic segregation process in the air-exposed crystal and its implication for the performance of Ha-B alumina as a solid electrolyte, are di scussod.

Surface and interface phenomena in ionic solids have been studied to rather limited extent, as compared to metallic materials, in spite of their implications for various properties of such materials. A fundamental question concerns the surface composition and how closely it mirrors the bulk composition. Deviations from the bulk compositions can be due to extrinsic factors, such as interactions with gases and influence of applied electric fields, or to intrinsic forces causing elemental segregation to the free surface, interface or grain boundries.

cleovoge face Fig. 1. Schematic diagram of /"O'O"O"O'O"O'cy single-crystal !Ja6-alumina; edge foce sodium (•) and oxygen (o) ions /O'O'C'G-O'O'O/ constitute the exposed conduc- 'o'c'c'o'o'o'o/ tion plane at the cleavage- fact. The spinel blocks separating the conduction planes from each other consist of close- packed aluminum and oxygen ions.

We have been conducting a comprehensive study of surface phenomena of the super ionic conductor sodium β-alumina 254

(1.2 II a O'llAl 0 ) using elcctronapec tro scopic techniques (XPS, AES). Due to its exceptionally high ionic conductivity at room-temperature, its high mechanical strength and chemical stabilityj Wa-!^ alutnitiii has been used as solid-electrolyte in the high energy dc-i.d i ly sodium-sulfur battery » arid riany investigators tried to <- J u..-i.i •. -• the Ha fast diffusion mechanism fl,2]. Previously v/e reported on the unuaual response ofNa-6 alumina, crystal to electron bombardment as revealed by AES experiments [3-5] • When the electron-beam was directed onto the cleavage face (see Fig. 1), induced high levels of negative charging at the surface region caused emergence of some sodium through induced cracks and faults in the spinel blocks. On the other hand, electron bombardment of the perpendialar edge face (Fig. l) resulted in weak negative charging which rather easily forced out large quantities of sodium. Thus, in this case electroir.igrat ion of Ha" along the conduction planes to the external edge-face took place and could be followed quantitatively via changes in the Auger intensities to give the bulk diffusion constant[3]. tfhile these experiments revealed t?ie electron-beam induced anisotropic response of the Ila-g ?i i u m: i, a 'Tvstal, M s .-._,_-; u I <• d with its two-dimensional superionic conauetivity, the present study aimed at exploring possible anisotrophy in the unpertur- bed surface composition of air-cleaved TJa-0 alumina. There- fore, rather than AES with the electron beam strongly affecting the mobile sodiuia LO.-JLS in this case comparative XPS measurements of the cleavage ;:.::,t e-.ip-j ?a.-.::-j, wcr >_• ... •._•!'.. fined . Observed differences in relative line intensities between the two surfaces, had to be compared to calculated values for the case of surfaces as terminating bulk planes.

EXPERIMENTAL

Air-cleaved slab of single crystal HSL-Q xli..mina. v•. I h ?~S rup.i wide edge-face was chosen, so as to yield reasonable XPS intensities. Yet, signal averaging was needed for the high resolution spectra of both surfaces (Pass-energy 50). The crystal was mounted on the sample holder in such a way that either the cleavage face or the edge-face could be analyzed after careful alignment in front of the CMA ( PHI 549 system). Mild sputtering was achieved with 500 V Ar+ ions (>ixlO torr) operated for short time intervals in order to obtain the near surface compositions.

RESULTS A1ID DISCUSSION

Figs. 2a-b present XP survey spectra of the cleavage and edge faces. As can be seen, in both cases there is a significant carbon contamination and different relative intensities of sodium, aluminum und ox;/een- ^n particular, '^-e .-.7- '.-tri;p of the as-inserted edge fi-cir- (Fig. 2b) roi'.lcjts rtlal:--?jj small concentration of aluminum as compared to TTa, O and C. Focuss ing on the aluminum-sodium ratio, we have chosen for quantitative analysis to compare the aluminum intensity with that of 255

Fig. 2. Survey XP spectra of as- inserted and Ar+-sputtered single crystal Na 3- alumina. (Mg Kα IJOCW X-rays; CMA pass-energy 100 Volt).

ttuttceo eogt lote

i-X S-M IOC 7ji fJO V * iTO *> fJC 0

BINDING tNtKGf. eV

the sodium Auger transition (KLL) rather than of Na (1$) since the former are much closer in electron kinetic energy, and! therefore are similarly attenuated by any overlayer contamination. Most significantly, the Na(KLL)/Al( 2P) peak area ratio measured from the high-resolution spectra of the edge- face is about six times larger than for the clearage-face. This remarkable difference may reflect excess sodium accumulation at the surface region of the edge-face,but possible intensity variation associated with the greatly anisotropic crystal structure of Na-0 alumina (Fig. l) should be considered before any firm conclusion can be made. We have derived expressions for the relative XPS intensities, ^-/^i' expected for the two faces under the assumptions: a J the cleavage plane as an overlayer contains, due to symmetry, half the number of sodium ions in a bulk conduction plane, and b) the edge-face region has homogeneous and stoichiometric distribution of Na and Al. The ratio of relative intensities, (iIJa/IAl) edge/( INa/IAl) cleavage, then reads ^ Sin9exp( - "Na Na = 2 Cal 1 -\ Na 2[l-exp( -d /> EinQ)] -1

n with n and'jja as the numbers of sodium atoms per unit volume ana per unit area in a bulk conduction-plane, respectively;d is the effective "thickness" of a conduction plane, d' is the distance between adjacent conduction planes (11.3 S 'Na and - the mean-free-paths of the 256

corresponding electrons(taken from data for alumina matrix [6]) and 6 is the "effective" emission angle of the elcetrons collected by the CMA [7]. With IK, /n' calcualted from the crystal structure, R 1 for Na( KLL ) / All 2P ) is 0 . U1* , whereas = the measured peak area pave the experimental ratio, Rejjp 6±1 This significant distinction between the measured ratio and the ratio calculated for "ideal" surfaces as terminating bulk planes seems to establish the conclusion concerning Ka segregation at the edge-face of Na-B alumina.

Also, the measured relative intensities, 0(lS)/Al(2P) and C(IS)/Al(2P), indicate excess of oxygen and even more of carbon at the edge-face as compared to the cleavage-face composition. In depth composition profiles of the edge-face obtained with short-time Ar+ sputter-etching are shown in Fig. 3. During the first 15 seconds of sputtering mainly carbon was removed thus leading to enhancement of signals from all other elements, '•fhile further sputtering gradually exposed more aluminum and oxygen, the sodium profile reached a maximum and then decrea<- sed similarly to the carbon. Disregarding possible preferential sputtering effects, this behaviour is consistent with the edge I!a segregation picture indicated by the quantitative analysis of the unsputtered surfaces. Inspecting the variations in ' relative intensities with sputtering-time (Fig. h) adds further information. Thus, during the first 30 sec there is rather sharp decrease in the Ila/Al ratio, and also a clear increase in the Al/ 0 ratio, which may indicate that the as-inserted edge surface contained some excess oxygen (adsorbed or oxidic) as compared to the bulk of the crystal. This increase can not be due to the different attenuations of the 0( IS) and Al(2P) signals coiring from the bulk by outer overlayers, since 0( IS) ha;.; shallower sampling depth so that Al(2P^)(lS) should have decreased upon removal of overlayers with no oxypen. The experimental ratio (edge/cleavage) of Na/Al relative intensities, R , decreased significantly upon 13 minutes sputtering of the ecFge-face (see Figs 2b and 2c) from a value of 6 to 0.7, which is much closer to theoretical value derived above (RCal = o.U).

Fig- 3. High-resolution O(IS) XPS peak area vs. sputtering time for ;iu-3 alumina edge- face. (500 volt Ar+).

Na(KLl)

2 3 4 t> 6 SPUTTERING TIMF

ftil2P)/01IS) • • *

* t t - Fig. h. Ratios of No(KLL)/CllSI UJ peak area from Fig. 3. .•••*• ;• s 1 NotKLL)/flll2P) UJ « (r

__i i i i 1,1,1,1,1^1 01211 567B SPUTTERING TIME (mm)

•fhile the quantitative investigation shows that the anisotropic segregation of sodium at the edge surface is accompanied by accumulation of carbon and some oxygen, no definite conclusion regarding the chemical bonding of carbon and oxygen could be drawn from their XPS chemical-shift data (the chemical shifts of the sodium lines are characteristic of Ha+). Nevertheless, we can guess that the preferential accumulation of Na+ at the edge-face originates from interaction (or chemical reaction) between the mobile Na+ and atmospheric gases (e.g. H?0). Such an interaction may occur at the cleavage face as well, but due to the particular crystal structure, only at the edge-face excess sodium diffusing from the near-surface region of conduction planes can be expected to emerge (see arrows in Fig. l).

Suci. anisotropic segregation of sodium may affect the performance of polycrystalline sintered Na-f5 alumina as solid- electrolyte in the Wa/S battery. Thus, one can expect that mass and charge transport through the solid-electrolyte/ liquid-electrode interface would be affected by the particular composition of the solid-electrolyte surface.

REFERENCES

1. G.D. Mahan and W.L. Roth Eds., "Superionic Conductors" (Plenum, Nev York, 1976). 2. A. Highe, M. Polak and R.'.{ . Vaughan, in "Proc. Inter. Conf. on Fast Ion Transport in Solids", Wisconsin, 1979 (North-Holland, Amsterdam, 1979). 3. A. Livshits and M. Polak, Surface Science,119 (1982) 311*. •» . M. Polak and A. Livshitz. Appl. Surface Science, 10 (1982) kk6. 5- A. Livshits and M. Polak, Vaccuum 33 (1983) 2Ul. 6. F.L. Battye et. al. Phys. Rev. D. 9 (l97l») 2887. 7. A. Livshits and M. Polak, to be published. 258

STRUCTURE OF PROTECTIVE DIFFUSION COATING FOR NIOBIUM ALLOYS

12 12 M. Kazinets , 0. Gafri , L. Zevin , B. Rabin

Institute for Chemistry and Chemical Technology, The Institutes for Applied Research, Ben-Gurion University, Beer-Sheva, Israel

IAI Bedek Aviation, Israel.

INTRODUCTION

Nb alloys, like other refractories alloys, have good mechanical properties at high temperatures but low oxidation resistance at these temperatures. Silicide-based, multilayered coatings are used for protection of Nb alloys against oxidation. One kind of coating (R512E) for the niobium alloy D-43 was found (1) to be composed of six discrete layers, some of which appeared to be single-phased. Five of the layers contained NbSi2 or Me5Si3 (Me = Cr, Fe) with slight amounts of NbSi2, and only the sixth was Nb5Si3. In our preliminary studies such coatings showed differences in the arrangement of the layers and in their phase compositions. The results of further work are presented here.

EXPERIMENT AND RESULTS

The specimens were samples of Nb alloy C-103, with a silicide coating (Si-20Fe-20Cr) formed by the slurry-diffusion process (2) . The metallographic structure of the coatings was studied with the scanning electron microscope (SEM). A cross section of one sample (Fig. la) shows the layered structure of the coating. A diagram of these layers is shown in Fig. lb. The total thickness of the coating was about 200 mm. The surface layer A of the coating is very thin, 10-15 pm. Layer B, which is just under the surface, is two-phased (I+II) , with the regions adjoining the surface richer in phase I. Layer C has approximately equal amounts of both phases. The thickness of layers B+C is about 120 ym. The next layer D is about 40 ym thick and seems monophasic. The boundary between layers C and D is sharp. Layer E is about 30 pm thick, is adjacent to the bulk Nb, and consists of 4-5 thin diffusion layers. The composition of each layer was determined by EDA X-ray spectrometry (Table I).

Note that phase II contains appreciable amounts of Cr, Fe and Hf, while phase I contains low concentrations of these elements. The Cr distribution in a cross-section of the coating (Fig. lc) provides more evidence that phase II contains large amounts of Cr.

The phase composition of each layer was determined by X-ray diffraction from the surfaces exposed by consecutive removal of layers. The penetration of the CuKa radiation used here is such that the thickness of the layer responsible for 90% of the diffraction intensity does not exceed 0.6 vm. This is less than the thickness of the removed layers. A total of 30 layers 5-6 urn each were mechanically removed, and 30 X-ray diffraction analyses were done. 259

a) b) c)

Fig. 1. Cross section of coating: a) SEM pattern, b) Diagram, c) Cr distribution

Table 1. Weight ratios of coirponents in the layers of the coating

Element, % Layer Phase or Nb/Si sublayer Si Nb Ti Cr Fe Hf

A 38.8 31.1 0.6 14.4 11.1 3. 9 0.80

B I 44.4 54.2 0.2 1.1 0.1 - 1.22

II 33.4 30.5 0.6 19.2 12.5 3. 8 0.91

I 43.1 52.5 0.2 2.6 1.3 0. 3 1.22 C II 31.3 29.6 0.6 20.0 14.7 3. 8 0.94

D 30.9 37.7 0.6 9.1 17.7 4. 0 1.22

4 40.1 51.1 0.4 1.4 4.2 2. 8 1.27

3 38.9 54.3 0.4 0.8 1.9 3.,7 1.40 E 2 28.6 61.3 0.7 0.9 3.0 5.,5 2.14

1 80-90 260

From the chemical data we could surmise that the layers are composed of Nb-Si compounds with some Cr (Fe)-Si compounds in regions rich in these elements. The phase identification was done on the basis of the X-ray diffraction data (3). There are six compounds in the Nb-Si system. They are: NbSi2, with a hexagonal lattice, a = 4.797A, c = 6.592A, space group P6222; Nb5Si3 - tetragonal, a = 6.570A; c = 11.884A; Nb5Si3 - tetragonal, high temperature phase, a = 10.018S, c = 5.072A, space group 142 m; Nb5Si3 - hexagonal, a = 7.5368, c = 5.248S; ttt>3Si - cubic, a = 4.211A; and Nb3Si - tetragonal, a = 10.2l8, c = 5.19?. In the Cr-Si system the compound Cr Si2 with a hexagonal lattice, a = 4.42R and c = 6.35A is known. Some diffraction lines were identified on the basis of the interplanar spacings of NbSi2- It was difficult to identify the rest of the lines because none of the known Nb-Si compounds has lines that coincide satisfactorily with the deXp. We therefore indexed these lines on the basis of the high-temperature tetragonal modification of Nb5Si3 (a = 10.018A and c = 5.072A), with the assumption that the Bravais lattice of this compound is not body-centered (4) but primitive. This permits the appearance of the lines with uneven sums of Miller indices, such as 102 (d = 2.46) and 331 ( d= 2.14). Thus, there are two phases in the surface layer, MbSi2 and the high-temperature tetragonal phase of The X-ray pattern changes at once after the removal of the surface layer A. The intensity of the l^35813 peaks increases. No further changes occurred in the next five patterns. During this time about 30 ym was removed, which represents the full thickness of layer B. In the sequence of diffraction patterns from layer C the intensities of the NbSi2 lines increased, while those of NbsSi3 decreased. The surface of this section consisted of two phases with a prevalence of the NbSi2 phase. The intensities of only two pairs of lines belonging to Nb5Si3 begin to increase again in layer D, those with d = 2O49 (400 ) +2.46 (102) and 1.24 (800) + 1.23 (204). The second pair of lines is the second-order reflections of the first pair. All other lines of Nb5Si3 were relatively weak. Therefore, this layer is textured. The NbSi2 peaks decrease and disappear. The X-ray pattern obtained in the middle of layer D showed only two peaks, 400 and 800; thus layer D is single-phased, homogeneous, and textured. The diffraction pattern from layer E contains peaks of both phases, NbSi2 and Nb5Si3, and also lines belonging to metallic Nb.

DISCUSSION

A comparison of the results obtained from metallographic investigation with those of X-ray diffraction shows that phase I is Nb5Sx3 and the phase II is NbSi2. Cr, Fe and Hf occur together with NbSi2 (II) and are absent from the areas of NbqSi3 (I) . The Cr in phase II may substitute for Nb in NbSi2, because NbSi2 and CrSi2 are isostructural. The hexagonal phase of compound Nb5Si3 does not exist at all in the coatings. This is a very important result, because the hexagonal phase NbsSi3 has a low oxidation resistance (5), in contrast to MbSi2 and the high-temperature tetragonal phase of Nb5Si3, which have a high oxidation resistance. 261

The results of the compositional analyses (Table I), show that there is an excess of Si in the coating layers, but we have not yet made an assignment for this surplus Si.

There are some differences between the structure of the layer analyzed earlier (1) and that determined here. The most important of them is that Cr and Fe were previously found to be absent from the NbSi2 layers (1) . However, they coexist in significant quantities with NbSi2 in our investigation.

REFERENCES:

1. Priceman S. and Sama L. Electrochem. Tech. 6_ 315-326, 1968.

2. Gafri, I., Rabin, B., Materials Engineering Conference 1981, Dec. 20-22 Technion, Haifa, Israel.

3. Joint Committee on powder diffraction standards. Powder diffraction file.

4. Parthe, L. Monatsch Chem. 86_, 385, 1955.

5. Lyndon B. Johnson Space Center, Houston, Texas NASA Tech. Briefs, Spring/Summer, p. 395, 1982. 262

THE EFFECT OF B AND P DOPING ON THE DEPOSITION RATE AND CHLORINE

CONCENTRATION OF Si FROM SiCl4 IN RF PLASMA

R. Manory, E. Grossman, R. Avni and A. Grill

Materials Engineering Department, Ben-Gurion University of the Negev Beer-Sheva, Israel

ABSTRACT

Silicon films were deposited by a cold RF plasma from a mixture of SiCl4 + H2 + AΓ and were doped in situ by introducing diborane or phosphine in the gas feed. The film thickness was measured by Scanning Electron Microscope and the content was determined by Energy Dispersive X-ray Analysis. The deposition rate and the chlorine content were reduced with addition of diborane and both were enhanced with addition of phosphine to the gas mixture. The described phenomena have been found to be independent of substrate material or deposition period. The morphology of the deposit has been found to be different in the two types of films. B-doped films show an uniform growth proceeding by uniform coverage of the substrate, while P-doped films show an erratic growth proceeding by agglomeration of solid clusters at preferential sites. The different deposition rates are explained by differences in the surface morphology. INTRODUCTION In a series of papers published almost simultaneously ten years ago (1-5) several workers reported the inverse effect which dopants from the 3A and 5A groups have on the deposition of silicon in CVD and plasma reactors. All these reports showed similar findings: that the deposition rates of Si are diminished in the presence of doping gases containing As or P(AsH3 or PH3) and increase in the presence of boron containing gases, such as B2H6 or BCI3 . Eversteyn and Put (1) performed CVD of silicon using silane as the silicon gas, while Ray-Choudhury and Hower (2) used both silane and silicon tetrachloride (SiCl4) in CVD and obtained similar results. Hall and Koliwad (3) showed that boron lowered the temperature required for chemical vapor deposition of silicon from silane by more than two hundred degrees C and increased the deposition rate probably due to a catalytic effect on the decomposition of silane. Similar findings were reported by Yasuda et al. (4), while Farrow obtained the same behavior of the deposition rates in silane plasma (5). Recently, Knights (6) reported on the catalytic effect of diborane during plasma deposition of a:Si-H from a gas mixture of SiH4 + Ar.Several attempts were made to explain this behavior. Chuang [7) has presented one of the more accepted theories about this behavior. He explained the inverse effects of phosphine and boron by the different electronic configuration of the depositing molecules and the surface. According to [7), both SiH4 and SiCl4 molecules have Si+-H" and Si+-Cl" ionic bonds respectively, would be attracted by a positive surface potential (p-type Si) and repelled by a negative surface potential (n-type Si). 263

The growth rate of the layer is thus affected in the adsorption stage. These effects of increasing and decreasing the rate of silicon deposition by B and P respectively have become since then a well known and accepted fact. Therefore, details about deposition rates obtained during doping in the plasma are now scarcely reported. In our deposition system we employed RF plasma to deposit silicon from a gas mixture of SiCl4 + Ar + H2, and B2H6 and PH3 were used as dopant gases. The deposition rates of silicon on graphite and glass were invetigated as a function of the plasma macrovariables, together with the CI content of the films and their moxphology. In this paper we present the effect of doping gases, diborane or phosphine on the deposition rate and composition of Si films deposited by an rf plasma of Ar + H2 + SiCl4. The difference between our results and the quoted data (1-7) will be discussed.

EXPERIMENTAL The deposition was carried out in the experimental setup shown in Figure 1. This system has been previously described elsewhere (8,9), and a few details are repeated here for reasons of clarity: The basic gaseous mixture consisted of SiCl4(99.9% pure supplied by Merck Gmbh."), hydrogen and argon (UHP grade-supplied by Matheson, USA) . High purity diborane and phosphine diluted in argon were directly introduced into the gas mixture and monitored by a mass flow controller. (Matheson-MFC). The films were grown at various conditions on graphite or glass substrates using dopant to SiC^ concentrations which varied from 2-10-3 to 2.10"2. The silicon to hydrogen ratio in the gas feed was 1:3. The pressure in the reactor was 1 or 2 mbar, the input power was between 100 and ISO Watts and the position of the substrate in the reactor was upstream, at the beginning of the RF coil (position H - Ref. 10,11). The experiments were carried out in pairs, using the same conditions for deposition of p-type or n-type layers. Therefore, the external parameters of the plasma such as pressure, substrate position, input power and period of deposition affected bath types of layers in the same way, and the observed differences are attributed only to the effects of the doping gases. Thickness measurements were made by Scanning Electron Microscope (SEM) and the chlorine content of the film was measured by Energy Dispersive X-Ray Analysis (pDAX). For samples grwon on glass, the percent of CI vs. Si in the film could not be determined, because the substrate itself contained silicon

RESULTS AND DISCUSSION A SEM micrograph of two layers deposited one upon the other is shown in Figure 2A, while Figure 2B showed the line-scan profile of the CI concentration for the same sample. The upper layer was P-doped and the lower one was B-doped, and each layer was deposited during two hours at 150 Watts input power and a pressure of 2 mbar. The dopant to SiCl4 ratio in the gas stream was 10~3 for both PH3 and B2H6 gases. The ratio between the thickness of the B and P-doped layers in this sample was measured to be 1:2.8 respectively. This factor (approximately 3) was found also for the ratio of the chlorine content in the two layers shown in Figure 2B. Similar values for the two ratios were obtained also when the order of 264

doping was changed and the deposition time was reduced to 1 hour for each layer. Table I summarizes the findings for several samples deposited on graphite and on glass substrates. The growth rate on the glass substrates was slower than on the graphite substrates. The films deposited on glass were grown separately and the comparison was made between different samples. This comparison shows that the ratio of deposition rates between B and P doped samples remains 1:3 even though the dopant concentration was changed by one order of magnitude. The difference in deposition rates between undoped and n-tyne layers deposited on glass was found to be 1.3 "" '„ value shows that phosphine enhances the deposition not only as _^mpared to the effect of diboTane, but also as compared to the growth of undoped layers. The data presented in Table I also indicates that the effect of the dopant gases on the growth rate is not affected by the nature of the substrate. These findings, first presented in (11) were later confirmed by other groups (12,13) working with halogeneous silanes (SiF4 and SiClj. respectively) in plasma. The described behavior, which is inverse to the behavior reported for CVD (1-5) or for plasma of SiH4(6) indicates that the effect of the doping gases is related to phenomena in the plasma phase and to the plasma-solid interaction for plasmas of halogenated gases. Figure 3 presents SEM micrographs of films deposited for short times on glass substrates. One can observe the uniformity of the growth of the B-doped film, as compared to the erratic surface of the P-doped and the undoped layers. The latter presents a growth pattern based on pile-ups at preferential sites. These pile-ups serve as three dimensional growth canters while the B-doped films grow uniformly on unoccupied sites on the surface. The different morphology of the growing films can account for the different deposition rates considering that pile-ups can grow into all directions above the surface, while the B-doped film grows uniformly on unoccupied sites on the surface. This morphological difference can explain the almost constant ratio between the deposition rates with the two dopants, for dopant concentrations in the gas as low as 10~3f and other deposition parameters. Figure 4A shows the behavior of the chlorine concentration vs. deposition time for undoped Si films, deposited on graphite substrates, and Figure 4B shows the behavior of the deposition rate for the same deposition conditions. A detailed discussion of these figures may be found elsewhere (14). As it can be observed the deposition rate and the CI concentration behave similarly with time, indicating a correlation between these two parameters. A similar correlation is observed also during doping, as indicated by Table I. The ratio between the chlorine concentrations in the two kinds of films deposited under similar conditions, is the same as the ratio of the growth rate of these films. We conclude therefore that the changes in the CI content in the doped films are a direct result of the different deposition rates of these films. It should be stressed that the presence of chlorine in the deposit is a significant difference between Si films obtained in plasma of SiCl4 as compared to Si films obtained by CVD from SiCl4 or by plasma of SiHg. It is therefore possible that the presence of CI in the films causes the difference between the effect of the dopants on the growth rate as observed in plasma deposition and CVD. This effect has still to be studied. 265

REFERENCES CI) F.C. Eversteyn and B.H. Putt, J. Electrochem. Soc. 12(3, 107 (1973). (2) P. Ray-Chouahury and P.L. Hower, J. Electrochem. Soc. 120, 1761 (1973) . (3) L.H. Hall and K.M. Koliwad, J. Electrochem. Soc. 120_, 1438 (1973). (4) Y. Yasuda, K. Hirabayoshi and T. Moriya, Proc. 5th Conf. on Solid State Devices, Tokyo (1973). (5) R.F.C. Farrow, J. Electrochem.Snc. 121_, 899 (1974). (6) J. Knights, J. Non-Cryst. Solids, 35_ and 36_, 159 (1980). (7) Chin-An Chuang, J. Electrochem. Soc., 123, 1245 (1976). (8) E. Grossman, M.Sc. Thesis, Ben-Gurion University, Beer-Sheva (in Hebrew), (1981). (9) R. Avni, U. Carmi, R. Manory, E. Grossman and A. Grill, Proc. 6th Int. Symp. on Plasma Chem., Montreal, M.I. Boulos and R.J. Munz Eds., Vol. 3 p. 820 (1983). (10) A. Grill, E. Grossman, R. Manory, U. Carmi and R. Avni, ibid, p. 843 (1983). (11) R. Manory, E. Grossman, R. Avni and A. Grill, Ann. Meeting Israel Phys. Soc, Bulletin of IPS, _2£, p.32 (1983). (12) B. Pratt, private communication (1983). (13) P. Capezutto, G. Bruno, F. Cramarossa, Proc. 6th Int. Symp. on Plasma Chem., Montreal, M.I. Boulos and R.J. Munz, Eds., Vol.3, p. 815(1983). (14) E. Grossman, A. Grill and R. Avni, Plasma Chem. Plasma Proc, 2, p. 341 (1982). ~

PRFPARATItW OF n- AND p-lYPE HICHOTKYSTALUNE Si BY THE COLD PLASMA TECHNIQUE 1. gases introducing system 2. mass flow controller 3. mixing chamber 4. Teactor 5. sample holder 6. matching unit 7. Z7.12 MU RF generator B. induction coil D. selection valve H. heater M. pressure gauge P. pumping system T. cold trap B Fig. 1: Experimental setup for plasma deposition. Fig. 2: A. SEM micrograph of two Si layers on graphite Sample 219) Thick layer - n type Si, Thin layer •• p-type Si. 2hr deposition time EDAX Line scan profile of CI concentration vs. Si for sample 219 (For CI concentration values see Table 1). 266

Figure 3: SEM micrographs of Si samples grown on glass: A - undoped B - n-type C - p-type. Figure 4: A - Behavior of deposition rate vs. deposition time for undoped Si deposited on graphite. B - Rate of CI incorporation ( CC1) in the coating vs. deposition time.

3 4 5 6 7 8 9 3 4 5 67 6 10 TIME, hours TIME, hours

Table I. Deposition Rates and CI Concentration for Several Si Layers.'

Sample Doping Substrate Average deposition CI vs. Si CI Dopant: No. Order Dep.Rate Rate Ratio (%) Content SiCl4 A/sec CB/P) Ratio Cone. 219 I-B graphite 5 1:2.8 3.5 1:3.1 ID"3 II-P 14.2 11

220 I-B graphite 4.9 1:2.8 4.8 1:3.1 ID"3 II-P 13.9 IS

223 I-P graphite 11.7 1:2.8 7.8 1:3.4 ID"3 II-B 4.2 2.6

504 none glass 2.2 — 505 P glass 2.8 (504/505) 2-10"2 1:1.3

S06 B glass 0.9 C506/S05) 2-KT2 i 1:3

CDExperimental conditions: for graphite: 150 Watts input power, p=2 mbar, SiCl4 : 3.7 v/o; total flow 130 SCCH. for glass: 100 Watts input power, p=1 mbar, SiCl4 : 5 v/o; total flow 60 SCCM. 26:

SILICON NITRIDE COATINGS BY THE LOW PRESSURE RF PLASMA TECHNIQUE U. Carmi, A. Raveh, A. Inspektor and R. Avni Nuclear Research Center-Negev, POB 9001, Beer-Sueva, 84190 ISRAEL

ABSTRACT Silicon nitride coatings have been deposited in a low pressure r.f. plasma on stainless martensitic steels (AISI-410), from gas mixtures of silicon tetrachloride (SiCl4) and amcnia (NH3) in an argon plasma. The substrates temperature between 230-440°C, depending mainly on the induced r.f. power. The coatings obtained were identified by x-ray diffraction and analyzed by scanning electron microscopy, optical microscopy and EDAX for morphology, thickness, and silicon and chlorine contents, respectively. The influence of parameters such as the pressure in the reactor and the r.f. power input on the properties of the coating such as the deposition rate, morphology and microhardness were investigated. The correlation between plasma parameters and properties of silicon nitride is shownSome discussion on the nature of the reactions mechanism is presented.

INTRODUCTION Silicon nitride had been chosen as a fine ceramic coating material due to its special properties. Silicon nitride is characterized by its high hardnessCll, its high resistance to thermal shock aid to chemical reactions. It has high fracture thoughness(2) and high Vickers microhardness (3) . These propertied make silicon nitride suitable for high temperature wear resistant coatings on parts such as turbine blades. Several methods are available for the deposition of silicon nitride coatings, each has its virtues as well as drawbacks. Amongst others one can name Chemical Vapour Deposition (CVD)(4), glow discharge(5), and sputtering (6) (reactive and non-reactive). R.F. plasma coatings offers a technique whereby some of the disadvantages of the above methods are avoided. In this technique the coating is carried out at low temperature (300-500°C) the substrate is coa":^1. on all sides including inside pores and holes, at relatively high deposition rate and adhesion. The deposition and coating of silicon nitride by the r.f. plasma technique will be discussed with regards to working parameters and resulting properties.

EXPERIMENTAL A full detailed description of the experimental setup is given elsewhere however for convinience of the readers some details are given below. R.F. plasma is induced inside a pyxex reactor by r.f. coil wound around it. Silicon nitride is produced on a metal substrate from silicon chloride and amonia in argon atmosphere. The coated metal was examined for its microhardness chlorine content and deposition rate. Also the morphology and crystaljographic structure were examined by SEM and XRD 268

respectively.

RESULTS AND DISCUSSIONS Primary investigations of the silicon nitride deposition indicated that the sample to be coated should be placed in position H in the plasma (upstream) and that the sample should be electricaly floating. The position- ning of the sample was fixed so that the chlorine content in the sample would be minimal (8). The electrical connection were determined from^-. the comparison of deposition rate on substrates biased and floating This was corrob rated by the results obtained by EPR (Electron Paramagne- tic Resonance) sampling of the plasma, showing that the (SiC14+NH3)/Ar plasma had a large concentration of radicals(10).XRD identification of the coatings shows the existance of hexagonal silicon nitride with cell units a=7.76^ and C=5.62A". Microhardness, deposition rate and chlorine content were determined as a function of working parameters i.e. pressure, power and NH3/SiC14 ratio. Figures 1-3 represent the changes in the above properties varying the pressure. It is noted (fig.l) that the microhardness increases with pressure up to 3 Torr, where it obtains its maximum value of 2400kg/mm2. At higher pressures the microhardness decreases. k similar behaviour is obtained for the deposition rate (fig.2) and for the chlorine content (fig.3) . It should be noted here that chlorine content is in percentage of silicon only as measured by EDAX. All the above results show that at 3 Torr the reaction reaches it optimum. This optimum may be explained considering kinetics of the reactions. At low pressure the reacting particles have larger mean free path and thus higer energies for reaction. Increasing pressure increases reactant concentration but reduces its energy. Thus ancptimum regarding the propagation of the reaction is obtained as seen in figs 1-3.

Figures 4-6 show the changes in the microhardness, deposition rate and chlorine content as a function of the input power. As expected from the previous explanation, increasing rf power increase the value of the •nicrohardness to a value of 3000kg/mm2,fig.4. The deposition rate has a somewhat unexpected behaviour, in fig.5 the round points curve rises to 100 watts and then drops moderatly with increasing power. rhis may be explained by looking at the deposition rate on an earthed substrate, (square points),which rises with power. The increase in power causes particles to be more energetic, and the rate reaction by ion molecule mechanism is increased on account of radical molecule mechanism. \s expected, from the above explanation chlorine content increases with increasing power as shown in fig.6. Figure 7 describes the microhardness as function of feed ratio. It can be seen that increasing NH3 concentration 'increases microhardness to a maximum and then drons. This is because Df the stoichiometry of the reactants, and the formation of compounds jther than Si3N4 at non adequate concentration. A fructured britlle silicon nitride, indicating its high microhardness, is shown is the SEM micrograph of fig.8. Ttcan be observed that the coating is of substantial thickness (*>»8 Urn) and it is well adherent to the metal substrate. The film is dense with high homogeneity and has no pores in it. The surface of the film shows that its build-up follows the topography of the metal surface, again indicating the adherent nature of the coating. 269

REFERENCES

1. Engineering Property Data on Selected Ceramics. Vol.1, Nitrides, Battelle Colombus Lab. NTIS AD 23773 p. 47 (1976).

2. L.J. BawenandT.G. Carruthers, J. Mater. Sci. 13, 684-687, (1978).

3. K. Niihara and T. Hirai,J. Mater. Sci. 12,1243-1252,(1977).

4. F. Galasso, U. Kuntz and W.J. Croft,J. Am. Ceram. Soc. 55,431,(1972).

5. W.A. PlisMng.J. Vac. Sci. Technol. 14, 1064,(1977).

6. C.J. Mogab and E. Lugujjo J. Electrochem. Soc. 127, 1853,(1980).

7. A. Raveh, Y. Hornik, U. Carmi, A. Inspektor and R. Avni, To be published in plasma chem. Plasm. Proc. (1984).

8. R. Manory, A. Grill, U. Carmi and R. Avni, Plasma Chem. Plasma Process. 3, 235, 0-983)

9. Y. Ron, A. Raveh, U. Carmi, A. Inspektor and R. Avni3 Thin Solid Films, 107, 181, (1983)

10. N. Mayo, U. Carmi, I. Rosenthal and R. Avni. To be published J. Appl. Phys. Apr. (1984). 270

FIG 1 FIG 2

2 3 4 5 12 3 4 5 GAS PRESSURE (TDII) GAS PRESSURE (Tori I

Fig.l. Microhardness (Vicker

FIG 4

2D0D u> FIG 3 ILJ a a. < i o a looo o i

1 2 3 4 i 6 100 150 GAS PRESSURE , loir RF POWER Watt

Fig.3. Chlorine content Ccl of SiN film vs. pressure it: the reactor for a floating substrate (430°C) (SiCl4-to-NH3 ratio, 1/3). Fig.4. Microhardness (Vickers1 hardness at a loadof 50 gf) of SiN film vs. r.f. powe* for ajfoating substrate (pressure 5 Torr; SiCl4 - to - NH3 ratio, 1/3). 271

FIG 5

FIG 6

4 a.

u> o 0.

100 150 100 150 200 RF POWER (Watt) RF POWER (Watt)

Fig.5. Growth rate (dh/dt) of SiN film vs. r.f. power (pressure; 5 Torr; SiCl4-to-Ml3 ratio, 1/3): grounded substrate; floating substrate.

Fig.6. Chlorine content Cci of SiN film vs. r.f. power for a floating substrate (pressure, 5 Torr; SiCl4-to-NH3 ratio, 1/3.

FIG7

iiooo

3 4

NH3/SiCI4 RATIO

Fig.7. Microhardness (vickers' hardness at a load of 50 gf) of SiN film vs. NH3-to-SiCl4 ratio (floating substrate, pressure, 5 Torr: r.f. power, 100 W). Fig.8. An SEM micrograph showing a top view of a fractured SiN coating on steel at an inclination of 60 . 272

BORIDATION OF STEELS IN LOW PRESSURE RF PLASMA

A. Raveh, A. Inspektor, U. Carmi, E. Rabinovitz and R. Avni

Nuclear Research Centre - Negev, P.O.Box 9001, Beer-Sheva 84190, ISRAEL

Prepared for Presentation at the Second Israel Materials Engineering Conference, Ben-Gurion University, 1984.

ABSTRACT Iron boride films were produced by the boridation of martensitic stainless steel in low temperature environment (less than 500°C). The process was carried out by the r.f. plasma of BCl3/H2/Ar at low pressure. The formation rate, microhardness and chlorine content of the films were determined and correlated to the working parameters. A maximum of formation rate [4-5 pm/h) and micxohardness (2600 kg/mm2) was found.

1. INTRODUCTION Boride films provide hard coatings(1,2) with high oxidation and chemical resistance C2~3)• Their properties make these materials attractive for us^ in corrosive and abrasive environment. Today iron boride films are produced by the electrolytic process(4), by chemical vapour deposition^^) or by the pack-cementation technique(1). In these methods boron diffuses into the ferrous material at high temperature (above 800OC) and reacts with the substrate material. At these conditions of high temperature damage to the substrate material and its pre-treatment will occur. The r.f. plasma technique is a convenient environment for low temperature surface boridation(7). Boron is supplied to the plasma from the decomposition of a boron compound such as diborane (B2Hg) or boron trichloride (BCI3) in an argon or argon-hydrogen mixtures.

This work sets to present the r.f. plasma boridation of stainless martensitic steel (AISI-410), and to check the influence of the working parameters on the surface boridation.

2. EXPERIMENTAL The experimental set-up is shown elswhere(8), however for the convenience of explanation the main features are outlined below. The substrate (AISI-410) was place inside a pyrex reactor (40 mm long and 8 mm in diameter) , in which plasma was induced by a r.f. coil wound around it and energized by 0.5 MHz generator. The reactor was pumped down to working pressure by rotary vane pump. 10 vol.% BC13 (Matheson C.P) premixed with ^rgon (Matheson, U.H.P) are further diluted by argon and mixed with hydrogen in a mixing chamber and flown into the reactor. Before boridation process the steel samples were mechanically polished to 25 ym (SiC) or to 0.05 at' (A^Oj), ultrasonicly cleaned in organic solvent and pretreated in avgon plasma for 50 min. 273

In the boridation process a mixture of BCl3+H2+Ar introduced to the plasma. The concentration of BCI3 in the gas mixture was 4 vol.% while the total gas flow rate was kept constant at 200 seem"1, during 2h processing. The following parameters were varied: i) total gas pressure in the plasma reactor, between 2-6 Torr; ii) r.f. power input expressed as current flowing in the r.f. coil, between 30-50 A. The iron boride films were analyzed by x-ray diffraction for their crystallographic identification, by scanning electron microscopy (SEM) for thickness and morphology and by energy-dispersive analysis of x-rays (EDAX) for chlorine contents in the films. Microhardness were determined on Vickers scale at 50 g constant load.

3. RESULTS AND DISCUSSION Two crystalline structure of iron boride were identified by x-ray diffraction: FeB, orthorombic (a0 = 4.053 A, b0 = 5.495 A and Co = 2.946 A) and Fe2B, tetragonal (a0 = 5.099 A and bo = 4.240 A). Figure 1 shows the iron boridefilms on a martensitic stainless steel (AISI-410) substrate. It can be seen that the film is well adherent to the substrate and has a non-porous structure. The surface of the boride film shows that it consists of small microscopic spherical shapes, and reproduce the geometric features of the substrate.

Table 1 shows the anlaysis of the iron boride films at different locations in the plasma reactor witn regards to r.f. coil: in position H (-1.0 cm away from the begining of the F 1.0, upstream), in positi^:1 G, in the middle of the r.f. coil and F 1.0 cm away from the end of the r.f. coil, downstream. The formation rate of the films do not strongly differ between the location H and G, but the chlorine content reaches a maximum in position G. The substrate temperature was the highest in position G, where the highest r.f. energy is delivered to the nlasma and substrate.

From table 1 and from previous work(8) it was evident that position H is favorable for higher formation rate and minimum chlorine content in the films. Therefore, the results which follow will be presented only for steel substrate located in position H.

A maximum rate of boride formation was obtained at gas pressure of 3 Torr and shown in Fig. 2. It should be mentioned that the boride films was formed on both sides of the steel substrate, and the film thickness was about 60% on the steel side facing the gas flow and 40% on the backe.

The maximum formation rate was obtained for the steel samples polished to 25 jjm by SiC i.e. for a rougher surface [Fig. 2).

The 25 ]im surface had a higher surface area available for the gas-metal interaction. Since in the formation of iron boride the sole source for iron is the metal substrate, a higher process rate expected.

The variation of the formation rate as a function of the current in the r.f. coil is shown in Fig. 3. Figure 3 shows a similar behaviour for pressures of 3 Torr and 5 Torr. A maximum of 4.5 ym/h was obtained at 41 A and 3 Torr and at a substrate temperature of 440±20°C (Figs. 2 and 3). This behaviour indicate that the formation rate reaches a maximum a 41 A, and is not increases with increasing the r.f. power. 274

The behaviour of the chlorine content in the boride film is shown in Figs. 4 and 5 as a function of gas pressure and current in the r.f. coil, respectively. The higher chlorine content in the boride was obtained on the steel substrate polished down to 0.05 ym. In other words, in order to reduce the chlorine content a rough surface is favorable. The chlorine concentration increases linearly with increasing the current in the r.f. coil, as shown in Fig. 5. It should be noted that increasing the current from 37 A to 46 A increases temperature of the substrate from 370°C to 500°C, and thus diffusion rate of chlorine content increases too.

Figure 6 shows the behaviour of the microhardness of the films as a function of the total gas pressure. At 2-3 Torr the microhardness reaches its maximum values, and is an order of magnitude higher compared to the non-boridized surface of the AISI-410 steel (i.e. 2600 kg/mm2 vs. 210 kg/mm2, respectively). 4. CONCLUSION Iron boride fi'ms was identified by X-ray diffraction as Fe2B and FeB at boridation of stainless martensitic steel (AISI-410) in a low pressure r.f. plasma. The starting materials were gas mixtures of boron trichloride (BC13) and hydrogen in an argon plasma. A maximum of formation rate (4.5 ym/h) and microhardness (2600 kg/mm2) obtained at current flowing in the r.f. coil of 41A and gas pressure of 3 Torr, whereas the substrates were maintained at relatively low temperature (440±20°C) compared to other techniques used for boridation.

REFERENCES 1. S.C. Singhal, Thin Solid Films, 45 (1977) 321. 2. R.H. Biddulph, Thin Solid Films, 45 (1977) 341. 3. E. Randich, Thin Solid Films, 63 (1979) 309. 4. H.C. Fiedler and R.J. Sicraski, Met. Prog. 99 (1979) 101. 5. H.O. Pierson and A.W. Mullendore, Thin Solid Films, 72 (1980) 511. 6. L. Randich, Thin Solid Films, 72 (1980) 517. 7. A. Raveh, A. Inspektor, U. Carmi and R. Avni, Thin Solid Films 108 (1983), 39. 8. A. Raveh, Y. Hornik, U. Carmi, A. Inspektor and R. Avni, to be published in Thin Solid Films, 1984.

Table 1 Properties of Iron Boride films at different locations. Boridation time 2h, current in r.f. coil 39A, feed composition BCI3 (4 v/o)+ H? (17 v/o) + Ar (79 v/o), gas flow 127 Substrate gas Film chlorine substrate microhardness position pressure thickness content temperature 2 (Torr) (ym) (w/o) C°C) (kg/mm ) H 4 7 24 440 2000 G 4 6 50 550 700 F 4 1 30 400 900 H* 1 5 10 451 1500

* H2/BC13 = 6/1, No argon + Values are in weight percentage of Fe in the iron boride films. 275

2 3 4 5

gas pressure (Torr)

Fig.l. An SEM micrograph showing a top view of an iron boride film on a martensitic stainless steel (AISI-410) substrate. Fig.2. Formation rate (dh/dt) of iron boride films vs. pressure (r.f. current of 44A and 3.9vol % BCI3); - - -samples polished to 25pm by SiC; samples polished to 0.05™ by AI2O3.

k 2-

40 44 current rf coil (A)

Formation rate (dh/dt) of iron boride films vs r.f. current at 3 Torr and 5 Torr, and at 3.9 vol % BCl,. 276

Fi'_».4. Chlorine content in boride films vs. pressure (r.f. current of 44A and 3.9 vol % BC13); —Samples polished to 25 pm by SiC. Samples polished to 0.05ym by Al203.

na- pressure (Ten-)

6 _ I 1 "' y ~

o I / - x 2 y

current rf CDII (A)

Fig-5. Chlorine content in boride films vs. r.f. current at 3 Torr and 3.9 Vol% BC13. Fig.6. Microhardness in boride films vs. pressure at r.f. current of 44A and 3.9 vol% BCI3. 277

LASER INDUCED COPPER ELECTROLESS PLATING

S. Tamir and J. Zahavi

Israel Institute of Metals, Technion, Haifa, Israel

INTRODUCTION

During the last few years, laser induced plating processes have been found to be a very attractive technology in producing high rate, highly selective deposits, mainly for the electronics and microelectronics industry. High rate metal plating from aqueous solutions on thin metallic substrate was reported by R.J. Gutfeld et al (1-5). Plating rate of more than 1000 times the rate of conventional plating was found (1) in the irradiated area when argon laser was used during copper electroplating process on copper thin fiim substrate laid on glass. It was proposed [2) that primarily thermal and heating effects associated with the laser radiation were responsible for the high plating rate.

This work aimed at studying laser induced copper electroless plating process on polymeric substrate such as polyiir.ide (Kapton). Polyimide is used as a substrate for flexible printed circuit boards.

EXPERIMENTAL

A schematic representation of the experimental set-up is illustrated in Fig. 1, Two types of laser were used in ^.nis study. The first one was a continuous argon laser operating at wavelength of 0.5145 microns, while the second one was a Nd+3/YAG laser used in the pulsed mode at wavelength of 0.53 microns, pulse duration of 150 ms and repetition rate of 1-5 KHz. The plating solution used was a commercial copper electroless plating solution, while the plating processes were carried out through various stages, consisting of specimen cleaning followed by immersion in sensitization and activation solutions, and finally immersion in the electroless plating solution. In the first part of the work, argon laser was used when the radiation of the substrate took place during its immersion in the plating solution. In the second part of the work, pulsed Nd+^/YAG laser was used to radiate specimen surfaces before the plating process, while being immersed in deionized water immediately after the activation process. Scanning electron microscopy (SEM) and Auger electron spectroscopy (AtiS) were used to examine deposit morphology structure and composition.

RESULTS AND DISCUSSION

Typical SEM observation of laser induced electroless copper deposition line is shown in Fig. 2. Deposit lines of this type were built using argon laser radiation on the specimen during its immersion in the 278

electroless plating bath. This irradiated line produced at 7.9 x 103 watt/cm2 showed cracks and microcracks probably in the polyimide substrate. Copper deposit up to 60C A0 in thickness was detected at the irradiated line zone by AES examination. AES surface analysis of the deposit line and of the polyimide between lines are shown after sputtering in Fig. 3. Copper was found in the bright irradiated line while no copper was found in the background. The sputtering was performed in order to eliminate the copper which might have remained from the plating solution. AES depth profiles of laser deposit line aud the line-free polyimide surface are shown in Fig. 4. Deposit line contained about 70% atomic concentration of copper and its thickness reached values up to 600 A0 (as shown in Fig. 4). Deposit thickness on non- irradiated surface areas was found to be around 150 A0. The calculated plating rate was four times greater at the irradiated areas compared to laser non-irradiated areas.

In the second part of the work, copper deposit areas were produced with pulsed laser system. A pulsed Nd+3/YAG laser was used at repetition rate of 5 KHz, pulse width of 150 nsee and intensity of 0.7 watt (experiment No. 31). Immersion time of the specimen in the water was around 10 minutes, while plating time was around 5 minutes in the electroless solution. SEM observation of copper deposit lines is shown in Fig. 5. Copper deposited at the irradiated lines while no deposition occured at the non-irradiated area. Deposit structure was characterized by fine and dense grains, where the grain size did not exceed half a micron, as shown clearly in Fig. 5b. These grained copper deposit lines had sharp edges with continuous structure. AES depth profile of the deposited line is shown in Fig. 6.

The analysis showed copper concentration profile with thickness ranging from as much as 60% atomic concentration up to around 20% atomic concentration (AC) at thicknesses of 200 A° and 1600 A°, respectively. However, at the non-irradiated areas (Fig. 7") , the copper concentration ranged from 30% (AC) to 10% (AC) at depths of around 100 A° to 700 A°, respectively.

CONCLUDING REMARKS

Highly selective high rate electroless copper deposition on polyimide substrate was obtained with C.W. Argon laser irradiation.

Laser irradiation took place at the final stage of the plating, i.e., after the substrate was immersed in sensitization and activation solutions.

Irradiated zones were electroless plated at very high rate of deposition in the value of about 120 pm/sec compared to about 10 pm/min in conventional electroless. Furthermore, the electroless plating was selective and took place primarily at the local zones that were irradiated with the laser beam.

The major effect of laser irradiation was probably the generation of local heating at the substrate - electroless plating solution interface resulting in high rate deposition. While the electroless solution 279

was kept at relatively low temperature of 15°C to 18°C, its plating reaction rate was very low and almost non-active. However, local selective laser radiation resulted in local temperature rise followed by high rate of electroless deposition as was observed.

Laser irradiation on the activated substrate surface resulted in plating deposition on laser irradiated zones while the noa-irradiated zones were not plated during the plating stage. Pulsed laser irradiation on specimen surface was carried out while the activated substrate was immersed in distilled water. It might be that the activator material was washed out by the water while laser irradiated zones maintained their activator material followed by electroless deposition in these zones.

REFERENCES

1. R.J. Von Gutfeld, R.E. Acosta and L.T. Romankiw, "Laser-Enhanced Plating and Etching: Mechanisms and Applications", IBM, J. Res. Develop. Vol.. 26, No. 2, March (1982).

2. I.C. Puippe, R.E. Acosta and R.J. Von Gutfeld, "Investigation of Laser Enhanced Electroplating Mechanism", J. Electrochem. Soc. Vol. 128^ No. 12, 2539-2545 (1981).

3. R.J. Von Gutfeld and L.T. Romankiw, "Laser-Enhanced Plating - Applications to Gold Patterning", Gold Bull. 15 (4),(1982).

4. R.J. Von Gutfeld, E.E. Tynan, R.L. Melchet and S.E. Blum, "Laser Enhanced Electroplating and Maskless Pattern Generation", App. Phys. Lett 35 (9), November 1979.

5. L. Kulynyck, L. Romankiw and R. Von Gutfeld, "Laser-Enhanced Exchange Plating", IBM Technical Disclosure Bulletin, Vol. 23, No. 3, August 1980. 280

CU BLACK LINE AFTER SPUT

600 840 1080

KAPTON AREA AFTER SPUT Bo .1C , ISO 360 600 840 1080 Fig. 1: Schematic illustration of KINETIC ENERGY. Ev experiments1 set-up. Fig. 3: AES surface analysis of laser induced copper electroless plating on polyimide surface after sputtering 90 nm. (A) Laser line irradiated area on polyimide surface. The presence of copper is shown. (B) Non-irradiated area of polyimide surface. No copper was detected.

Fig. 2: Lines of laser induced copper electroless plating on polyimide substrate. (A) Irradiated line showing morphological damage in form of microcracks. (B) Enlargement of (A). 281

0 9 18 27 36 45 54 63 72 81 SPUTTER TIMEtMIN.) 100 -

t°60 If** KAPTON AREA SPR:IOA/MIN

20 - B 0 w Iβ 27 36 45 54 63 72 81 SPUTTER TIME (MIN.)

Fig. 4: AES depth profile analysis of laser induced copper electroless plating on polyimide. (A) Laser line irradiated area. (B) Non-irradiated area.

Fig. 5: SEM observations of pulse laser induced copper electroless plating on polyimide substrate. Laser irradiation took place after activation process. (A) Typical deposit at various areas. (B) High magnification of the deposit area shown in (A). 282

10 15 20 25 30 35 40 SPUTTER (MIN.)

Fig. 6: AES atomic concentration depth profile of pulse laser copper electroless plating on polyiraide. Laser irradiation took place after activation process.

10 15 20 25 30 35 40 SPUTTER TIME (MIN.)

Fig. 7: AES atomic concentration depth profile of pulse laser copper electroless plating on polyiraide. Non-irradiated area. 283

SURFACE HARDENING OF STEEL BY BORIDING IN A COLD RF PLASMA

I. Finberg, R. Avni and A. Grill

Materials Engineering Department, Ben-Gurion University of the Negev Beer-Sheva, Israel

T. Spalvins and D. Buckley

Tribology Branch, NASA-Lewis Research Center, Cleveland, Ohio, U.S.A.

ABSTRACT Samples of 4340 steel have been plasma borided in cold rf plasma of diborane and argon. The plasma was initiated in a gas mixture of B2H^+Ar inductively coupled to 27.12 MHz rf generator. The steel samples were treated in the plasma at a net power input of 500W for 5 hours, the bulk temperature of the samples reaching values of 550-600°C. The influence of the position of the sample in the reactor and of the gas pressure was investigated. The plasma borided surfaces were studied by scanning electron microscopy, X-ray diffractometry and microhardness measurements. It was found that as a result of the decomposition of the diborane in the plasma, boron is deposited on the surface of the steel substrates and two crystalline phases are formed: tetragonal Fe2B and orthorhombic FeB. Due to the formation of the boride phases, the surface microhardness increased from the original value of 270 Kg-mm-2 to a maximum value of 790 Kg-mm-2.

INTRODUCTION Boriding, or boronizing, has become an effective means for wear protection in many fields of industry (1-3). In addition to the increasing of the surface hardness, the boriding improves the resistance of steel to certain types of corrosion facids and molten light metals) (2,3) and improves the oxidation resistance up to 800°C (4). The high hardness of the borided layers is attained directly through the formation of borides and the process does not require subsequent quenching (5). Boriding is performed in a gas, molten salt or pack cementation process at temperatures above 900°C (2,5,6). At these temperatures the effect of previous heat treatment will generally be lost. In order to obtain the required bulk properties,heat treatment has to be done after boriding, with the possibility of spalling of the hardened surface layer. Boriding at lower temperatures could therefore be advantageous. During boriding of steel two phases are formed, Fe2B and FeB. The simultaneous presence of two phases affects adversely the quality and wear behavior of the borided samples (1,2) and attempts are therefore made to obtain only the less brittle 284

Boriding by a cold plasma process can be advantageous from two aspects: (i) Since the availability of the active species in the plasma is determined by the electrical discharge, the temperature of tho substrate can be controlled independently and can be kept at relatively low temperatures, determined by the required diffusion rate of boron into the steel. (ii) By controlling the plasma parameters, it may be possible to produce only one boride phase, like in the ionitriding process (7), resulting in an improved wear behavior of the borided samples. The aim of the present study is to investigate the feasibility of plasma boriding of steel and to determine the effect of the plasma parameters on the boriding process and properties of the borided surface.

EXPERIMENTAL The experimental set up for plasma boriding is similar to the one described elsewhere (81• The plasma was initiated in a premixed gas mixture of 2.87% B2H6 in Ar with a 27.12 MHz rf generator, inductively coupled to the quartz reactor through an impendance matching unit. The treated samples were prepared from 4340 steel sheet and polished with 1 pm diamond paste. Before plasma boriding the samples were cleaned in acetone and ultrasonically cleaned in ethanol. The bulk temperature of the substrates was measured with a chromel alumel thermocouple in contact with the substrate. X-ray diffractometry was used to identify the crystallographic phases formed at the treated surface. Its morphology was analyzed by scanning electron microscopy, SEM. The microhardness of the treated surface was measured, normal to it, with a Vickers diamond indenter at a load of 20 g. Preliminary Auger Electron Spectroscopy, AES, analysis was also performed to determine the concentration of the elements, close to the surface of the borided samples. The decomposition of the diborane, I^Hg, in the plasma, the deposition of boron on the steel substrate, and its diffusion into the steel are strongly dependent on the plasma parameters. In the present rfork the effects of two plasma parameters were investigated: (i) The position of the sample relative to the rf coil and direction of gas flow: H - at the entrance to the rf coil (upstream). G - at the center of the rf coil. F - at the exit of the rf coil (downstream). (ii) The total pressure of the gas in the reactor which was varied between 1 and 5 mbars. The other plasma parameters were kept constant through this study, namely: Total flow = 5 seem Net Power Input = 500 W Power Density = 2 W-cnr3 Treatment Time = 5 hours The temperatures reached by the substrates during plasma boriding were 550-600°C. 2S5

RESULTS AND DISCUSSION Figure 1 presents the surface morphology, typical for the plasma borided samples. The appearance of the surface indicates that an overlay coating was deposited onto the substrate, where the growth proceeded through solidification of spherical globules. The size of these globules varied between 0.6 and 1.8 ym, however no systematic dependence of the size of the globules and either position of substrate in the plasma, or gas pressure could be determined from the available data. Further investigation is required to find whether the plasma parameters affect in a systematic way the morphology of the surface. Preliminary AES analysis indicated that the layers close to the surface of the plasma treated samples are rich in boron (more than 90%). Depth profiling indicated that the concentration of boron decreases and the concentration of iron increases into the sample as shown in Figure 2. Figure 3 presents an X-ray diffractogram obtained from the surface of a plasma borided 4340 steel sample. As indicated in the figure, X-ray lines identified as belonging to the tetragonal Fe2B and orthorhombic FeB phases are obtained from the plasma borided sample. Although, as indicated by the AES measurements (Figure 2) , boron rich layers exist on the surface of the treated samples, only X-ray diffraction lines of 4340 steel, Fe2B and FeB have been obtained. This indicates that the boron atoms not bonded in Fe2B or FeB form an amorphous boron layer. The results of SEM, AES and X-ray diffraction indicate that: (i) As a result of the excitation of the plasma in the B2H5 -i Ar mixture, the diborane decomposes and boron species deposit on the steel substrates. These boron species may be either BnHm fragments (n < 2 , m < 6) or boron atoms. Identification of the BnHm species requires mass spectrometrical study of the plasma, (ii) During the continuing deposition of the boron species on the steel substrates, a diffusion process of both boron and iron atoms occurs, resulting partly in the formation of the two crystallographic phases, and FeB. The microhardness of the treated surfaces has been measured normally to the surface. A though the measurements made normal to the surface will give values which to a certain extent are averages between the hardened layer and the bulk, these values can allow a qualitative comparison between surfaces treated under different conditions. In addition to the contribution of the bulk to the measured microhardness values, these values are also affected by the morphology of ths surface. Due to the roughness of the surface as shown in Figure 1, the measured microhardness values are lower than the microhardness of the material itself. The microhardness values of the treated surfaces should be compared to the microhardness of the untreated steel samples, of VHN = 270 Kg-mm-2. In order to check for an eventual heat treatment effect on the microhardness, the microhardness was measured also on the surfaces of the treated samples which were not exposed to the plasma. It was found that the microhardness on the unexposed surface was.the same as that of the untreated sample, thus excluding any heat treatment effects on the microhardness of the plasma borided samples. Figure 4 presents the microhardness values as a function of the position of the sample in the reactor, for different boriding pressures. The results indicate that the hardening effect decreases from position H (entrance of the rf coil) towards position F (exit from the rf coil). The values of the surface ir.icrohardness attained at a given position depend on 286

the gas pressure during the plasma boriding. As can be seen the highest value of 790 Kg-mm" 2 (as compared with 270 Kg*mm~2 for the untreated sample) was obtained at position H and 3 mbar. Taking into consideration the effect of the microhardness of the bulk and of the surface roughness, as mentioned before, the microhardness of the treated surface layers is higher than 790 Kg-mm-2. This indicates that, as a result of the exposure of the steel samples to the diborane plasma, a significant hardening of the surface is obtained. The hardening of the surface layer occured at bulk temperatures lower than 600°C, as compared to conventional boriding temperatures of 900°C. Continuing studies are in progress to find the plasma parameters for further increase of the surface hardness and for the formation of a single boride phase in the plasma borided layer.

ACKNOWLEDGEMENT The authors are grateful to the US-Israel Binational Science Foundation for supporting the work by a research grant 2813/82.

REFERENCES 1. K.H. Habig, Materials in Engineering, 2^ (2), 1980. 2. P. Goeuriot, F. Therenot and J.H. Driver, Thin Solid Films, 78, 1981, 67. 3. R.H. Biddulph, Thin Solid Films, 47_, 1977, 341. 4. L.S. Lyankhovich and S.S. Bragilevskaya, in "Protective Coatings on Metals", ed. G.V. Samsonov, Consultants Bureau, NY 2, 1970, 123. 5. T.S. Eyre, Wear, 34, 1975, 383. 6. N.N Golego, A.P. Epik, V.D. Derkach and'V.F. Labunets, Idem 4, 5, 1973, 257. 7. B. Edenhofer, Heat Treatment of Metals, 1974, 59. 8. R. Avni, U. Carmi, R. Manory, E. Grossman and A. Grill, Proc. ISPC VI, Eds. M.I. Boulos and R.J. Munz, p. 820.

Figure 1 Surface Morphology of Plasma Borided 4340 Steel p = 3mbar position - H 287

Figure 2. Composition of plasma borided surface as determined by AES with depth profiling. o

40 120 200 280 360 440 520 600 SPUTTERING TIME, mm

Figure 3. X-Ray diffractograms of 4340 steel.

Pressure, (mbar) CM _ *- 1 r 800 0 Figure 4. E • -2 E Surface uicrohardness of 700 - o-3 • plasma borided 4340 steel • -5 as a function of position GO 600 in the reactor. GO ai -z. 500 - • Q * a: A < o 400 * o cc o 300 * •

200 I -2 0 + 2 F G H POSITION, cm 288

CURRENT METALLIZATION ISSUF-S IN MICROELECTRONIC DEVICES

K.N. Tu

IBM Thomas J. Watson Research Center Yorktown Heights, New York 10598

The current trend in very-large-scale integrated Si devices is miniaturization. Since it is not accompanied by a reduction of the temperatures used in fabrication and operation, there are acute structural and compositional instability issues due to thermal stress and diffusion. Three issues on metallization will be briefly reviewed. They are the delamination in polycide gates induced by thermal stress, the failure in conducting lines due to electromigration, and the degradation of ohmi' or Schottky contacts caused by interdiffusion. These issues are generic and persistant in the miniaturization of device structures.

The application of silicide in FET gate has been dominated by the combined use of refractory metal silicide and heavily doped poly-Si as the so-called 1 2 polycide gate. - The silicides such as WSi2 and TaSi2 must be a good conductor, chemically stable with the poly-Si so that the gate oxide will not be affected, and capable of surviving a high temperature process without transforming into oxides. For these requirements, the conduction and oxidation behavior3 of refractory metal silicide have been studied. The oxidation of silicide is also interesting for the formation of a surface oxide layer of gate insulation. When a composite of refractory metal disilicide (with excess Si) and poly-Si is oxidized, a layer of SiO2 grows on top of the disilicide by consuming the excess Si and the poly-3i. Since Si is the dominant diffusing species in the refractory metal disilicides, the transport of Si will cause a reverse flux of vacancies to condense at the interface between the disilicide and poly-Si interface, weaken- ing the adhesion across the interface. Therefore, peeling occurs if the gate is highly stressed. The gate tends to experience a high thermal stress since a high temperature (~1000°C) annealing is required to improve the conductivity of the disilicide film. The disilicide film is typically amorphous in the as-deposited state. Upon annealing, the amorphous film crystallizes into a highly defective structure, and it is only after the high temperature annealing that the defects are reduced and the conductivity greatly improved.4 At present, how we can lower the annealing temperature without sacrificing the conductivity is an 289

important issue of the polycide gate. Also, the choice of disilicide is unsettled yet.

Electromigration in Al thin films can be retarded by adding a few percent of Cu.5 To be effective, the Cu addition has to go specifically to grain boundaries in the Al film but not to the interior of Al grains. This is because electromigration at low temperatures ~100°C in Al films occurs mainly along grain boundaries, and any additive of Cu into Al grains will raise line resistance. The grain boundary diffusional flux due to electromigration can be decreased by reducing the grain boundary diffusivity or the effective charge number6, or by increasing the grain size. The beneficial effect of alloying Cu to Al has been found largely due to the reduction of grain boundary diffusivity of Al. Howev- er, this effect alone will be insufficient, especially in resisting eiectromigration damage in one micron or submicron lines. For this reason, additional improvements of using bamboo-type microstructure, very fine grains, and sandwich structure have been developed. In the sandwich structure, the middle layer serves as a diffusion barrier in preventing openings from running across the entire film,7 so its lifetime is improved. Clearly, the issue of electromigration will remain and it will become more serious at places where current crowding occurs such as a stepped surface.

Interdiffusion occurs naturally in a multilayered thin film structure. The consequence of interdiffusion, for example, of Al to a PtSi/n-Si contact is to change Schottky barrier height. It can be detected readily by current-voltage measurements and analyzed by the concept of parallel contacts.8'9 To prevent the Al-pene*ration, most often a diffusion barrier is introduced to slow down interdiffusion.10'12 However, we must understand the unique behavior of interdiffusion in thin films in order to solve the problem satisfactorily. Thin fiim interdiffusion tends to form a single intermetallic compound rather than to form all of them simultaneously as in bulk diffusion coupies. The phenomenon of "single" intermetallic compound formation has been verified by using high resolution transmission electron microscopy of atomic resolution.13 A kinetic explanation has been given assuming that the growth is limited by diffusion as well as interfacial reaction.14 The consequence of "single" compound formation is that the boundary conditions of its growth kinetics are simple and we can analyze it with diffusion markers to unravel the intrinsic diffusion coefficients.15 Furthermore, the marker analysis can be repeated in each of the compounds formed in sequence. The selection of the first compound and the sequence of the followers have been found not to depend on driving force (free energy change)16 but rather on kinetics.17 This has been supported by studying interdiffusion in Al-Cu thin film bilayers; the compound CuAU which has the fastest interdiffusion coefficient is the one which forms first1 \ Knowing the intrinsic diffusion coefficients, the diffusion flux and in turn the failure rate can be estimated. However, to reduce the diffusion flux and to develop an effective diffusion barrier are still important technological issues. 290

REFERENCES

1. B.L. Crowder and S. Zirinsky, IEEE Trans. Electronic Devices, ED26, 369 (1979). 2. S.P. Murarka, J. Vac. Sci. Technol. 17, 775 (1980). 3. J.E.E. Baglin, F.M. D'Heurle and C.S. Petersson, J. Appl. Phys. 54, 18 49 (1983). 4. T. Tien, G. Ottaviani, and K.N. Tu, J. Appl. Phys. 54, 7047 (1983). 5. F.M. d'Heurle and P.S. Ho, chapter 8 in "Thin Films - Interdiffusion and Reactions".. Ed. by J.M. Poate, K.N. Tu and J.W. Mayer, Interscience - Wiley, N. Y. (1978). 6. H.B. Huntington, in "Diffusion in Solids - Recent Developments" Ed. by A.S. Nowick and J.J. Burton, Academic Press, N.Y. (1975). 7. J.K. Howard, R.F. Lever, P.I. Smith and P.S. Ho, J. Vac. Sci. Technol. 73,68 (1976). 8. I. Ohdomari and K.N. Tu, J. Appl. Phys. 57, 3735 (1980). 9. J.L. Freeouf, T.N. Jackson, S.E. Laux and J.M- Vvoodall, J. Vac. Sci. Technol. 21 (2), 570 (1982). 10. M-A. Nicolet, Thin Solid Films 52, 415 (1978). 11. M. Wittmer, J. Noser and H. Melchior, J. Appl. Phys. 52, 6659 (1981). 12. K.N. Tu, Chapter 7 in "Preparation and Properties of Thin Films", Ed. by K.N. Tu and R. Rosenberg, Academic Press, N.Y. (1982). 13. K.N. Tu, G. Ottaviani, U. Goselu and H. Foil, J. Appl. Phys. 54, 758 (1983). 14. U. Gosele and K.N. Tu, J. Appl. Phys. 53, 3252 (1982). 15. U. Gosele, K.N. Tu and R.D. Thompson, J. Appl. Phys. 53, 8759 (1982). 16. K.N. Tu, G. Ottaviani, P..D. Thompson and J.W. Mayer, J. Appl. Phys. 53, 4406 (1982). 17. H.T.G. Hentzell, R.D. Thompson and K.N. Tu, J. Appl. Phys. 54, 6923 (1983); H.T.G. Hentzell and K.N. Tu, J. Appl. Phys. 54, 6929 (1983). 291

ELECTRICAL PROPERTIES OF Ti:W SC1IOTTKY BARRIER CONTACTS TO SILICON

M.O. Aboelfotoh and K.N. Tu

IBM Thomas J. Watson Research Center Yorktown Heights, New York 10598 (U.S.A)

ABSTRACT

Ti:W alloy films sputter-deposited on chemically -leaned surfaces of n- and p-type Si(100) have been annealed at temperatures up to 750°C. Schottky- barrier height measurements were performed over a wide range of temperature using forward c .tent-voltage characteristics. The barrier height on n-type Si was found to increase and that on p-type to decrease with increasing the annealir g temperature. These changes in the barrier height can be interpreted in terms of the changes in the interfacial oxide on the Si surface upon annealing. The barrier height on n-type Si was also found to decrease with increasing measurement temperature, and that on p-type Si to be independent of the measurement temperature. The temperature dependence of the barrier height is shown to be nearly equal to that of the energy band gap in Si.

INTRODUCTION Films of Ti:W alloy have found application as a diffusion barrier between platinum silicide contacts and aluminum interconnecting lines in very-large- scale integrated devices.1 This alloy is also attracting increasing interest2 for application as low Schottky-barrier contacts to silicon with stable electrical characteristics at high temperatures. In this contribution the measurements of the Schottky-barrier height and its dependence on temperature for Ti-W alloy on n- and p-type silicon are reported. The effect of annealing at temperatures in the range from 400°C to 750°C on the current-voltage characteristics of ihe Ti:W-Si Schottky-barrier contacts is also reported.

EXPERIMENTAL PROCEDURE

The Ti:W alloy films were prepared by d.c. magnetron sputtering. Films 1000-1500A thick were deposited on O.OO5-£2cm n+ and p+ (100) oriented Si wafers with 2-jum-thick, 10-ficm epitaxial layer. Buffered HF was used to 292 clean the Si wafers immediately before they were loaded in the deposition chamber. Prior to deposition, the deposition chamber was evacuated to lxl0"7 Torr before admission of lOmTorr of argon of 99.99 percent purity. The 2 Ti3oW7O target was given a 30 min sputter clean against a shutter at 3 W/cm before deposition. The Ti:W was sputtered at 0.5 W/cm2 resulting in a deposition rate of 40A/min. The wafers were grounded during the Ti:W deposition. The resistivity of the Ti:W alloy films was about

During the same deposition, two types of samples were made: bare Si samples for x-ray diffraction, Rutherford backscattering spectrometry (RBS) and Auger electron spectroscopy (AES) analysis, and SiO2 covered and pat- terned Si samples for I-V characteristics. For the I-V measurements the deposition was made through a metal mask having 2.5-mm diameter openings but the active area of the contacts was defined by oxide windows with diameter of 130, 250, 500, and 1000 micrometers. The samples were annealed at temperatures in the range from 400°C to 750°C in a furnace with flowing helium purified by hot (950°C) Ti particles.

Composition analysis of the Ti-W alloy film and detection of compound formation upon annealing were made using MeV 4He+ Rutherford backscattering spectrometry. The structure of tho phases in the film was determined by glancing-angle x-ray diffraction in a Seeman-bohlin geometry diffractometry.

The forward I-V characteristics were analyzed according to the thermionic emission theory3 of Schottky-barrier current transport: I = aT A* exp ( - q<£BO/kT) exp (qV/nkT), where V>3kT/q is the forward applied voltage, <£BO is the effective barrier height at zero voltage, A* is the theoretical effective Richardson constant for Si, T is the measurement temperature, a is the contact area, and n is an ideality factor introduced in order to take into account deviations from ideal diode behavior. The value of $BQ was determined from an extrapolation of the linear portion of a In I-V curve to zero applied voltage. The extrapolation was made using a curve fitting program, which fits the experimental data points by exp (qV/nkT). The slope of the linear portion yielded the value of n. The I-V measurements were made with the samples held at temperatures between -90°C and +20°C. The measurements were made on all four diode areas.

Experiments were performed to investigate the importance of leakage currents at the edge of the Ti:W-nSi contacts. The edge leakage occurs as a result of the high field due to accumulation of the n-type Si surface by the 4 5 4 5 positive charge in the SiO2 surrounding the contact. - It has been shown ' that for contacts to n-type Si, this edge effect dominates the I-V characteristics for reverse-bias voltage and for low forward-bias voltage. We have plotted our measured reverse currents at -0.1V against contact diameter and found that the data lie along a straight line with a slope equal to 1.7 indicating that almost all 293

of the current is flowing across the whole area of the contact. If, on the other hand, the reverse current was dominated by the edge effect, the data would lie along a straight line with a slope equal to unity. In Ti:W-pSi contacts, the edge effect is absent because the surface of the Si at the edge of the contact is depleted in this case. The depletion does not give rise to high fields as does accumulation in the case of n-type Si. Indeed, a plot of our measured reverse current against contact diameter for Ti:W-pSi contacts showed that the data lie along a straight line with a slope equal to 2, showing that the reverse currents are proportional to the contact area. We concluded that the edge effect is not important in determining the I-V characteristics of the Ti:W-Si contacts.

EXPERIMENTAL RESULTS AND DISCUSSION

X-ray analysis on the as-deposited Ti:W alloy films showed the presence of a bcc solid solution of Ti and W2. This phase persisted upon annealing up to 700°C.

01 0.2 0.5 04 0.5 FORWARD BIAS (V)

Fig. 1. Forward I-V characteristics Tor 400°C annealed Ti:W Schottky barrier contacts to n-type Si as a function of the measurement temperature. 294

The forward I-V characteristics of Ti:W schottky-barrier contacts to n- and p-type Si annealed at 400°C for 1 hr. are presented, respectively, in Figs. 1 and 2 as a function of the measurement temperature. The results show that the ideality factor n is slightly temperature dependent. The deviation of n from unity can be caused by image force and the presence of an inlerfacial layer between the metal and the semiconductor. The contribution of field dependence of image-force lowering of the barrier height to n does not exceed one percent for barriers with the Si doping concentration (1015cirr3) under consideration.3

Auger depth profile analysis revealed that an interfacial layer containing oxygen is present between the Ti-W and Si even after annealing at high temperatures. The presence of an interfacial layer is well known to cause field dependence of the barrier height3. The existence of a field dependence of the barrier height must then represent the largest contribution to deviations of n from unity. Since the barrier height obtained by extrapolation to zero voltage is affected by the electric field in the depletion region, the zero-electric-field barrier height and not the zero-voltage barrier height should provide a better characterization of the metal-semiconductor Schottky barrier. Several authors have modeled the electric-field dependence of the barrier height.6"8 More recently, Wagner et al.9 derived a simple expression relating the zero-voltage barrier height 4>BO and ideality factor n (as determined from the forward I-V characteristics) to a zero-electric-field barrier height c/>BF defined under flat- band condition:

ln (^j n - *ype Si

(1)

p — type Si

where Nc and Nv are, respectively, the effective density of states in the conduction and valence band. The change in Nc and Nγ with temperature 3//2 variation is calculated assuming a dependence law of T '". ND and NA are, respectively, the donor and acceptor concentration. The values of the barrier heights 4>gp and <£gF based on equation (1) and deduced at various measurement temperatures from the curves of Figs. 1 and 2 for the 400°C annealed samples shown, respectively, in Figs. 3 and 4. It is seen that the barrier height on n-type Si decreases with increasing measurement temperature, and that the barrier height on p-type Si is independent, within the experimental errors, of the measurement temperature. 295

'.'

42"C ( -54°Cli 67acl -Bl DC( -a4»C(( hiii

£ 10 ;'*/ • / it i Li ;/

(Tt W) • pSi(lOO) 400 = " I hr 10-9

0 01 02 03 04 05 FORWARD BIAS (V)

Fig. 2. Forward I-V characteristics for 400°C annealed Ti-'tV Scholtky barrier contacts io p- type Si as a function of the measurement temperature.

In order to understand the temperature dependence of <£BF, a comparison is made between the experimentally determined values of <£gF and those calculated on the assumption that it is related to the dependence of the Si energy band gap Eg on temperature. The calculated values are normalized to the experimental value of <£BF determined at -82°C. The reported temperature dependence of Eg in Si by Bludau, Onfon, and Heinke11 is used. The results show that the decrease in <^>Sp with increasing temperature (-2.2xlO"4eV/°K), in the temperature range -82 C to +20°C, is very close to the change in Eg.11 Hence it can be concluded that the temperature dependence of the barrier "weight of Ti:W on n-type Si, is nearly equal, within the experimental errors, to the temperature dependence of Eq.

In Figs. 3 and 4, the values of the barrier heights 0BF and <>BF for the as-deposited samples and for samples annealed at 500°C to 700°C are also shown as a function of the measurment temperature. It is seen that the value of <£BF increases and that of <£gF decreases with annealing up to 700°C, with 296

+ ec ua to 0RF ^BF J l Eg- In the as-deposited and all annealed samples, $BF is independent of the measurement temperature, and the change in $BF is nearly equal, within the experimental errors, to the change in Eg. These results imply that the Fermi level at the Ti:W-Si interface is pinned in relation to the valence band edge. Auger depth profile analysis revealed that the oxide on the Si surface is gradually reduced and the oxygen is dissolved into the Ti:W, and that intermixing of Ti:W and Si at the interface does not occur upon annealing up to 700°C. It is likely that this change in the interfacia! layer (regarding the gradual reduction of the oxide on the Si surface) can change the pinning position of the Fermi level and hence the barrier height (Fig. 3) upon annealing.

RBS and x-ray analyses showed, after annealing at 75O°C, the formation of a ternary silicide phase (TixWj_x) Si2(x~0.2) with a hexagonal WSi2 2 12 structure. ' The values of the barrier heights 0BF and <>BF for the Ti:W silicide are also shown, respectively, in Figs. 3 and 4 as a function of the measurement temperature. As in the case of the Ti:W alloy, <|>gF is independent of the measurement temperature, and $BF decreases with increasing the measurement temperature with a coefficient of -2.7xl0'4eV/°K. This is again nearly equal to the change in Eg. These results, imply that the Fermi level at the Ti:W silicide-Si interface is also pinned in relation to the valence band edge. In addition, the results of Figs. 3 and 4 show that the values of 0gF and $BF for Ti:W after annealing at 700°C, are essentially the same as those when the silicide phase is formed. Auger depth profile analysis revealed that oxygen is expelled from the interface and is dispersed throughout the reacted layer as a result of silicide formation. The result that changing the interface from that of Ti:W on Si to that of Ti:W silicide on Si does not change the value of the barrier height indicates that the barrier height is pinned at a value which depends on the Ti:W alloy and not on the silicide phase. We conclude that it is the Ti:W alloy which is important in determining th electrical properties of the Schottky barrier in the Ti:W silicide-Si system.

ACKNOWLEDGEflENTS The authors gratefully acknowledge helpful discussions with P.S. Ho. It is also our pleasure to acknowledge R.D. Thompson for valuable assistance in making the electrical measurements, J.E. Lewis for AES data, P. Saunders for RBS data and CSS Materials Laboratory at Yorktown for the Ti:W deposition.

REFERENCES

1. P.B. Ghate, J.C. Blair, C.R. Fuller and G.E. McGuire, Thin Solid Films 53, 117 (1978). 2. S.E. Babcock and K.N. Tu, J. Appl. Phys. 53, 6898 (1982). 297

3. E.H. Rhoderick, Metal-Semiconductor Contacts (Clarendon, Oxford, 1980). 4. A.Y.C. Yu and E.H. Snow, J. Appl. Phys. 39, 3008 (1968). 5. A.M. Cowley, Solid-State Electron. 12, 403 (1970). 6. J.D. Levine, J. Appl. Phys. 42, 3991 (1971). 7. C.R. Crowell, Solid-State Electron. 20, 171 (1977). 8. J.M. Shannon, Solid-State Electron. 19, 537 (1976). 9. L.F. Wagner, R.W. Young and A. Sugerman, IEEE Electron Device Lett. EDL-4, 320 (1983). 10. S.M. Sze, Physics of Semiconductor Devices, Wiley-Interscience, New York (1981). 11. W. Bludau, A. Onton and W. heinke, J. Appl. Phys. 45, 1846 (1974). 12. F. Nava, C. Nobili, G. Ferla, G. Iannuzzi, G. Queirolo and G. Celotti, J. Appl. Phys. 54, .7.434 (1983). TEMPERATURE CO TEMPERATURE CO -80 -60 -40 "20 -100 -80 -60 -40 -20 0 65 T" i r

750-C O o

065 0 60

- 500-C

S 0.60 A— A A A V-

V V ro ID " AS-DEPOSITED —' 00 0.55 0 50 — . 0 AS-DEPOSITED O AS-DEPOSITED A 4OO°C- Ihr A 400°C-lhr - a 5OO°C-Ihr V 700°C-lhr _ O 600°C-lhr T 750°C-Jhr V 700°C-lhr - v 75O°C-3hr 1 I l 1 I 0 45' 200 250 200 250 TEMPERATURE(K) 150 TEMPERATURE (K)

Fig. 3. Dependence of barrier height on measurement temperature in Ti:W-n type Si Schottky-barrier contacts ann- Fig. 4. Dependence of barrier height on ealed sequentially from 400°C to measurement temperature in Ti:W-p type 750°C. Si Scholtky barrier contacts annealed sequentially from 400°C to 750°C. 299

INTERFACIAL REACTIONS IN LASER ANNEALED Ni/GaAs CONTACTS

A. Lahav(a\ T. Brat(a\ C. Cytermann(b) and M. Eizenberg(a'b)

(a) Dept. of Materials Engineerings Technion, Haifa (b) Solid State Institute, Technion, Haifa

ABSTRACT Thin films of Ni on (100) GaAs substrate were annealed by Nd-YAG pulsed laser (A=0.53ym). Morphology and phase composition of the reaction products as a function of laser energy density were studied by a number of analytical techniques, including Transmission Electron Microscopy, Auger Electron Spectroscopy and X-ray diffraction. The results show the lack of adhesion between melted Ni-film and GaAs, indicating that this technique is unsuitable for preparation of Ni/GaAs contacts. Nickel is usually added to the eutectic composition of Au and Ge in the fabrication of ohmic contacts to n-type GaAs by short duration melting at a temperature of ^ 450°C (1,2). It was found that Ni forms (Ni,Ge)As compounds at the interface with GaAs (2,3) and thus it is considered to be a "wetting agent" which prevents "balling" and improves surface morphology of the contacts. The curiosity about the role of Ni initiated a research on the interactions between Ni and GaAs both in solid and liquid states. Here we present the results on liquid phase reactions, and these on the solid state will be reported somewhere else. Pulsed laser annealing seems to be a good technique for investigation of metallurgical aspects of the interaction of melted Ni with GaAs, since the melting temperature of Ni (1453°O is higher than that of GaAs (1238°C). tfelting of the localized area on the surface of the sample by laser beam followed by very fast solidification is an extremely nonequilibrium process but in some cases it improves contact morphology and electrical properties (5-7). Ni films 1300 A in thickness were evaporated by an electron gun on clean (100) GaAs single crystal substrates at a vacuum of 1.10 torr. Laser annealing was performed by Q-switched Nd-YAG laser (wavelength 0.53]im) tfith pulse duration of 150 nsec. The laser beam was focussed on the sample surface with a lens of f=75 mm, resulting in a beam diameter (at 1/e intensity) of 64ym. The average laser energy density was about 6 J/cm^ and it was attenuated by a polarizer. It should be noted that the laser power was not stable enough and the energy density of individual pulses varied by *v< 10% of the average value. Area coverage was obtained by computer controlled movement of the sample on a X-Y table. Characterization of the reaction products was carried out by Nomarski nicroscopy, Scanning Transmission Electron Microscopy (STEM), X-ray diffraction, Auger Electron Spectroscopy (AES) and Energy Dispersive X-ray Spectroscopy (EDS). TEM specimens were prepared by chemical thinning from the back side of the sample. X-ray diffraction pattern of the as deposited sample (Fig.la) shows peaks of polycrystalline Ni and a very wide peak, the (400) reflection, of single crystal GaAs. TEM examination (Fig.2a) reveals the average 300

grain size in the Ni film to be of the order of 200A. AES depth profiling shows that the interface between Ni and GaAs is abrupt and free from carbon and oxygen (to the limit of detection of this technique *v> 1 at.%). Morphological examination of the annealed films by Nomarski microscope shows that at laser energy densities less than 0.4 3/cvP- the Ni film remains smooth and undisturbed. At an energy of 0.4 J/cm^ small hillocks of l-4ym in diameter appear on the sample surface. TEM examination of such a hillock reveals an increase in the average grain size up to lOOOA (Fig.2b). Electron diffraction pattern from this area yields only Ni rings, and EDS analysis in the STEM mode shows that this area contains only Ni. The above results suggest that these hillocks represent areas where Ni film was melted and recrystallized. In some places, where the laser energy was larger than the average, the continuity of Ni film is broken (Fig.3a) and it is folded to the periphery of the melted spot, leaving almost a flat surface of GaAs in tha central area. Such a behaviour of melted Hi film indicates the lack of adhesion between it and GaAs, possibly due to large surface tension, or due to undetected- interfacial contamination. At laser energy density of 0.6 J/cm^ the whole surface is covered with melted spots. The enlarged Nomarski image of a few spots is shown in Fig.3b. It can be seen that the central area of a spot is quite flat but covered with tiny ripples. The spots are separated by walls of 3-5]im in thickness. X—ray diffraction of this sample (Fig.lb) yields peaks of Ni (much weaker than in the as-deposited sample), polycrystalline GaAs5 and also additional peaks that can be indexed as belonging to a hexagonal phase with lattice parameters a=3.78 and c=5.01 A. A similar phase, the ternary compound Ni2GaAs, forms as a result of solid state reaction between Ni thin film and GaAs (4,9). AES high resolution analysis shows that the central area of the laser spot is almost free of Ni and contains only Ga and As (Fig.4a), whereas the walls around the laser spot contain about 50% Ni as well as Ga and As, and are probably responsible for the X-ray reflections of the ternary phase. TEM bright field images of the characteristic structures in the central area of laser spot are shown in Fig.5. Fig.5a shows dark ripples on a white background. Electron diffraction from the background gives patterns of GaAs in single crystal and polycrystalline form. EDS analysis of the background and the ripples confirms the AES results (absence of Ni), and shows that the ripples contain about three times less As than the background. Thus it may be assumed that the ripples are formed during the regrowth of melted GaAs accompanied with Au evaporation from the surface. At some places the ripple area has a cell structure (Fig.5b) with Ga segregation on the cell boundaries, probably due to constitutional supercooling (8). The area of the walls around the laser spot couldn't be analyzed by TEM, since it was too thick to be transparent to electrons.

In conclusion, it was demonstrated that the lack of adhesion between nelted Ni film and GaAs makes the method of pulsed laser annealing unsuitable for preparation of Ni/GaAs contacts. The situation is quite different in the case of solid state reaction, where Ni was found to be very reactive and it forms an epitaxial compound at the interface with 3aAs at 150°C (4,9). It should be noted that unlike nickel, germanium 301

being pulsed laser annealed at the same conditions does not have any adhesion problems and forms a uniform reacted layer (4,7).

ACKNOWLEDGEMENT The authors gratefully acknowledge the help of R. Brener in AES analysis

REFERENCES 1. N. Braslau, J. Vac. Sci. Technology, 19, 803 (1981). 2. G.Y. Robinson, Solid-State Electron, 18, 331 (1975). 3. M. Ogawa, J. Appl. Phys. 51, 406 (1979). 4. Authors unpublished data. 5. G. Eckhardt, in "Laser and Electron Beam Processing of Materials", C.W. White and P.S. Peercy, eds., p.467. Academic Press, New York (1980). 6. 0. Aina, J. Norton, W. Katz and K. Rose, in "Laser and Electron- Beam Interactions with Solids", B.R. Appelton and G.K. Celler, eds. (1982). 7. J.S. Williams, in "Laser Annealing of Semiconductors", J.M. Poate and J.W. Mayer, eds., p.383. Academic Press, New York (1982). 8. A.G Gullis, in "Laser Annealing of Semiconductors", J.M. Poate and J.W. Mayer, eds., p.147. Academic Press, New York (1982). 9. M. Ogawa, Thin Solid Films, 70, p.181 (1980).

Fig.l. X-ray diffraction patterns, a) as deposited, b) annealed at 0.6 J/cm2. o - Ni, A - GaAs, li- Ni-GaAs. 302

Fig.2. TEM bright field image of Ni film. a) as deposited, b) annealed at 0.4 J/cm

Fig.3. Nomarski micrographs of annealed surface by an energy density of a) 0.4 J/cm^, b) 0.6 ^ 303

450 720 990 1260 KINETIC ENERGY , (EV)

Fig.4. High resolution Auger spectra of sample annealed at 0.6 J/cm^. a) central area, b) wall area.

Fig.5. TEM bright field images from central nrea of laser spot at 0.6 J/cm2. a) Ga-rich dark ripples in GaAs single crystal, b) cell structure in the ripples. 504

ENHANCED OPTO-ELECTRONIC ACTIVITY OF SEMICONDUCTORS BY MEANS OF (PHOTOELECTROCHEMICAL ETCHING

R. Tenne and V. Marcu

Department of Plastics Research The Weizmann Institute of Science Rehovot 76100, Israel

The performance of electro-optical devices such as solar cells; light emitting diodes; radiation detectors etc. is determined, in part, by electron-hole recombinations which occur at their interfaces. Various techniques were employed for the passivation of surface states and recombination centers like dipping in metal ions solutions (1); bombardment with atomic hydrogen beams (2) and different etching techniques. We have discovered (3) that a short photoelectrochemical etching of Cd-chalcogenide electrodes leads to a considerable improvement in the performance of photoelectrochemical cells based on these semiconductors. This technique is of special importance to devices which are based on thin film layers of semiconductors where surface defects and grain bounderies have adverse effect on the performance. For yet uncomprehensible reason short photoelectrochemical etching removes these defects preferentially, whereas chemical etching fails.

Te Figure 1 shows the I-V curve of an electroplated CdSeQ g5 0 %<- electrode in polysulfide electrolyte. The beneficial role of the photoetchning is clearly exhibited. It is noteworthy that well over ten semiconductors which belong to five different families have shown positive response to photoetching. These families include: Cd-chalcogenides; Zn-chalcogenides; oxides; ternary semiconductors and laminar materials. Whereas light was necessary for the photoetching of n-type semiconductors forward bias was used in the electrochemical etching of p-type semiconductors. Various spectroscopic techniques have been used recently in order to confirm the selectivity of the photoetching towards surface states and recombination centers. They include spectral response (4,5), electron beam induced current (6); cathodoluminescence (7); electroluminescence;photovoltage transients; photocapacitance (8) and electrolyte electroreflec- tance (9) • Although each technique employed different semiconductor and different etching conditions were used, the emerging consequence was that indeed surface states and recombination centers, which are not accessible to regular chemical etchants, were removed upon short photoetching. The cathodoluminescence signal of CdTe prior to and after photoetching is an example. The two-fold increase in the signal is mainly attributed to the increased lifetime of minority carriers which is caused by the removal of surface detects. 305

Apart from the favorable effect of the (photo)electrochemical etching on the performance of semiconductor devices there seems to be some fundamental questions which are related with this phenomena. First, oxidative decomposition of the semiconductors seems to be more facile and selective than reductive decomposition by surface electrons. Both CdTe and InSe, which can be prepared as n or p-type materials by appropriate doping, were selectively etched by surface of p-type InP in comparison to its n-type counterpart, under illumination.

Second problem which is probably of even more fundamental nature is the appearance of a unique morphology of photoetched surfaces. The dense pattern of etch pits (>109 cm" ) is in general not sensitive to the circumstances under which the photoetching was carried out. Recently it was shown, however, that the etch pits density depends on the donor density (conductivity) of the crystal (n-type) (12). Figures 2 and 3 show the morphology of photoetched CdSe with two different donor densities. Figures 4 and 5 show the morphology of photoetched CdTe with two different donor densities.

We suggest (13) that the etch pit pattern is the result of a non- uniform flow of charge carriers across the space charge layer of the semiconductor during photoetching. This non-uniform flow results from non- uniformities in the electric fields which are induced by the localized positive charges of the donors (n-type) within the space charge layer. Evidence in support of this model comes from different sources. For example, if one applies a strong forward bias during the photoetching, the space charge layer diminishes and the density of the etch pits is reduced considerably. Also, the etch pit density is similar for different faces of the same crystal and it is not sensitive to the duration of the photoetching; light intensity (over a certain threshold); electrolyte used, etc.

The implications of this theory are not limited to Schottky barriers. In fact non-uniform Fiicroscopic fields are expected in any semiconductor junction and they influence the charge flow of both majority and minority carriers across the space charge layer. Recent experiments in which the wavelength of the exciting light is varied are in accord with the present model.

REFERENCES

1. B.A. Parkinson, A. Heller and B. Miller, J. Electrochem. Soc., 126, 954 (1979). 2. C.H.Seagar. and D.S. Ginley, Appl. Phys. Lett., 34_, 337 (1979). 3. R. Tenne and G. Hodes, Appl. Phys. Lett., 37_, 428 (1980). 4. R. Tenne, Y. Mirovsky, Y. Greenstein and D. Cahen, J. Electrochem. Soc, U9_, 1506 (1982). 5. Y. Mirovsky, R. Tenne, G. Hodes and D. Cahen, Thin Solid Films, 91, 349 (1982). 6. G. Hodes, D. Cahen and H. Leamy, J. Appl. Phys., 54_, 4676 (1983). 7. R. Tenne and A.K. Chin, Mater. Lett., 2, 143 (1983). 306

8. R. Haak, C. Ogden and D. Tench, in "Photoelectrochemistry: Fundamental Processes and Measurement Techniques", eds. W.L. Wallace et al., Elect-ochem. Soc. Ser., 82-3, 486 (1982). 9. M.Tomkiewicz, W. Siripala and R. Tenne, J. Electrochem. Soc, in press. 10. a) N. Miller and R. Tenne, Appl. Phys. Lett., 39_, 283 (1981); b) R. Tenne, ibid., 43, 201 (1983). 11. R. Tenne, B. Theys, J. Rioux and C. Levy-Clement, submitted to J. Appl. Phys. 12. R. Tenne, V. Marcu and N. Yellin, submitted to Phys. Rev. Lett. 13. a) R. Tenne and G. Hodes, Surf. Sci., 135^, 453 (1983); b) R. Tenne, H. Flaisher and R. Triboulet, Phys. Rev. B., in press.

V(mV).-i pt

Figure 1 (+) The electrode ?.fter preparation; (x) The electrode after mild etching; (o) The electrode after photoetching. 307

i&fi

Figure 2 Figure 3

The photoetching of CdSe COOl) from Cleveland Crystals was performed in Na2S03 1M, at +1V vs. SCE and AMI light intensity.

Fig. 2: p = 1 £2-cm Fig. 3: p = 10 n> cm

Figure 4 Figure 5

The photoetching was performed in HNO,; HC1; H90 [1:4:10 volj, at +1V vs. SCE and AMI light intensity.

Fig. 4: p < 5 Fig. 5: p ~ fl-cm 308

LASER INDUCED METAL DEPOSITION ON GaAs SUBSTRATES

J. Zahavi

Israel Institute of Metals, Technion, Haifa, Israel.

INTRODUCTION

Recently, much attention has been paid to laser-induced chemical and electrochemical plating from aqueous solutions on metallic (1.) , semiconductor (2-5) and polymeric (4) substrates for possible use in microelectronic circuits and devices. Use of C.W. laser to enhance electroplating or electroless plating of nickel,copper and gold from aqueous solutions on thin metallic layers evaporated on glass substrates was carried out (1) for maskless microcircuit repair and design changes.

Pulsed Q-switched Nd /YAG laser irradiation induced metal or alloy deposition from aqueous electroplating solution or semiconductor substrates for potential formation of conductive lines, as well as ohmic and schottky barrier contacts, have recently been reported (2-5). These investigations demonstrated that semiconductor substrates such as Si (2), InP (3) and GaAs (4,5) could be selectively plated with Pt, Au, Ni-Pt, from aqueous solutions, making use of laser beam irradiation without external electric current and without masking procedures.

Mechanisms and processes associated with metal deposition via pulsed laser irradiation on semiconductor substrates immersed in standard plating solution without external electric current could be characterized by three major stages: a. Chemical reduction of metal ions from the electroplating solution by photogenerated electrons (7,8) followed by metal atom deposition at laser irradiated zones. b. Surface heating followed by interdiffusion of deposited atoms into the substrate material (2,5). c. Termination of laser pulse followed by rapid cooling rate resulting in formation of alloys or compounds of the semiconductor with deposited diffused elements to give, for example, Pt or di silicides with silicon substrate (2) .

This study aimed at characterizing and studying factors affecting laser-induced, highly selective, high-speed gold deposition from standard electroplating solution on undoped, n-type-and p-type GaAs substrates. 309

EXPERIMENTAL AND PROCEDURE

A schematic outline of an experimental system used in conducting laser induced plating is described in Fig. 1. The system consisted of a Q-switched Nd:YAG pulsed laser system, x-y computerizsd table and an electroplating cell, all of which were controlled and monitored by a computer (Fig. 1) .

Semiconductor GaAs flat specimens 15 mm x 15 mm in size were prepared from (a) intrinsic undoped GaAs wafer with resistivity of 10 Q cm, (b) N-type Tellurium doped GaAs wafer with resistivity of 2.1 x 10 n.cm and (c) p-type zinc doped GaAs wafer with resistivity of 5.5 x 10"3 fi.cm.

A commercially available potassium gold cyanide electroplating solution produced by Lea-Ronal Company, New York (known as Aurospeed CVD electroplating gold solution has been used in this study.

Specimens immersed in the plating solution were irradiated with pulsed laser beam at wavelengths of 1.06 yin and 0.53 pm, pulse duration of 15 ns, beam diameter of 0.5 mm and energy densitizer of 0.1 joule/cm^ to 0.5 joule/cm2. Continuous deposit lines or squares were produced on the various substrates.

Gold deposit surface morphology structure and composition were examined by microscopic and surface analysis techniques, such as scanning electron microscopy (SEM), transmission electron microscopy (TEM). Auger electron spectroscopy (AES), Electron spectroscopy for chemical analysis (ESCA), Rutherford back scattering (RBS), and x-ray diffraction analysis.

RESULTS AND DISCUSSION

Laser induced high speed highly selective deposition of gold on GaAs semiconductor substrate immersed in aqueous plating solution of gold without masking and external electrical current have been achieved in this work. Lines or squares of gold deposition were formed using laser energy densities of 0.1 to 0.5 joule/cm2 at wavelength? of 1.06 ym and 0.53 ym, laser pulse duration of 15 us with pulse over- lapping of 95% with beam diameter of 0.5 mm. Gold deposit surface was characterized by the presence of craters or cell-like events of several microns in size. These events resulted probably from heating or melting processes associated with laser irradiation upon the GaAs substrate surface, as shown clearly in Fig. 2. Furthermore, deposit surface roughness increased with increase of laser energy densities as observed clearly in typical SEM micrographs (Fig. 2). Typical structures of gold deposit produced via pulsed laser of GaAs were examined by TEM as shown in Fig. 3. Typical gold deposit structure was characterized by small gold polycrystallites ranging from less than 100A° to a few hundred angstroms, as revealed by bright and dark field pictures in Fig. 3A and Fig. 3B, respectively. A typical electron diffraction of the gold deposit (Fig. 3C) revealed the presence of pure metallic gold in the GaAs substrate. The presence 310

and the amount of gold denosit were also obtained by RBS techniaues. Typical RBS snectrum indicated the presence of gold on GaAs laser irradiated zones, as shown in Fig. 4. It might be noted that the presence of elemental gold metal at the laser irradiated zones was obtained by ESCA and by x-ray diffraction examinations (data are not shown).

The amounts of gold deposits were determined by RBS examinations. The results were correlated with laser energy densities as shown, for exanrale, in Fig. 5. Gold deposit concentration increased linearly with increase of laser energy density. Furthermore, it was found that gold deposit on N-GaAs subst.ate was found to be one order of magnitude higher in comparison to gold concentration found on P-GaAs substrate under the same laser operating conditions.

Typical results of current-voltage curves characteristic of Schottky barrier contacts and of ohmic resistance are shown in Fig. 6 for laser induced gold deposits on n-type Tellurium doped GaAs. Deposit resistance values were obtained from the slopes of 1-V curves (Fig. 6B) which were measured and recorded at various deposit lines. Deposit resistance values decreased linearly with increase of laser energy density as can be seen, for exanrnle, in Fig. 7.

CONCLUDING REMARKS

A. It has been demonstrated that gold could be deposited on GaAs semiconductor substrates via pulsed laser irradiated upon the GaAs immersed in standard electroplating solution.

B. Laser induced high rate highly selective gold deposition on GaAs was conducted in direct one ster> eliminating masking and external electric current.

C. Laser induced deposition took r>lace primarily through photoelectro- "hemical deposition, followed by inter-diffusion of gold atoms into the GaAs substrate through a molten surface laver which solidified upon the termination of the laser pulses.

D. Laser gold deposit showed Schottcky behavior contacts on n-type GaAs under the operational conditions employed in this work.

REFERENCES

1. T.C. Puippe, R.E. Acosta, R.J. Van Gutfeld, J. Electrochem. Soc, Vol. 128_, No.12, pp. 2539-2545 (1981].

2. Y.C. Kiang, J.R. Moulic, J. Zahavi, IBM Technical Disclosure Bulletin, Vol. 26_, No.l, p. 327 (June 1983].

3. R.F. Karlicek, V.M. Donnelly, J. Appl. Phys. 53 (2), pp. 1084-1090 (1982] . ' ' 311

4. J. Zahavi, "Laser Induced Electrodeposition on Polyiraide and GaAs Substrate", Research Report, Israel Institute of Metals, Technion, Haifa, Israel; AFWAL-TR, 1983, Wright-Patterson AFB, OH 45433, U.S.A. (August 1983), to be published.

5. J. Zahavi, "Laser Induced Gold Deposition on GaAs Semiconductor", Research Report, Israel Institute of Metals, Technion, Haifa. Israel; AFWAL-TR-1983, Wright-Patterson, AFB, OH 45433, U.S.A. (August 1983), to be published.

6. J. Zahavi, S. Tamir, "Laser Beam Technology Promoting High Speed and Selective Plating Processes", Annual Technical Report, Israel Institute of Metals, Technion, Haifa, Israel; AFWAL-TR-83, Wright-Patterson AFB, OH 45433 (September 1983), to be published.

7. R.H. Michaels, A.D. Darrow, R.D. Rauch, Appl. Phys. Lett., Vol. 39_, No. 5, pp. 418-420 (September 1981).

8. F.W. Ostermayer, Jr., P.A. Kohl, Appl. Phys. Lett., Vol. 39, No. 1, pp. 76-78 (1981). 312

Computer Laser

Plating Solution Specimen^ W////M • X-Y Table

Fig. 1: Schematic description of experimental set-up.

TEM OBSERVATION LASER QOLD DEPOSIT ON UNDOPED QaAB LASER WAVE LENGTH =0.53 pm: ENERGY = 0.23 JOULE/cm2 1.0 PULSE/SPOT: 60x60 PULSES OF 0.5mm BEAM DIAMETER

LASER PLATED QOLD ON N-Qui: Ta (100) LASER WAVE LENGTH -1.08 pn (* Iβ)

0.11 JOUU/te1

Fig. 2: Typical SEM observation Typical TEM observa- of laser plated gold on tion of laser plated GaAs gold on GaAs 313

RBS. PROFILE LASER PLATED GOLD ON N-TYPE TELLURIUM DOPED GoAi (100) SUBSTRATE

1.200 l.GOO 2.0GQ

ENEHGY We VI

Typical RBS profile showing the presence of gold at laser irradiated zones on GaAs substrate.

LASER PLATED GOLO CN N-GoAi (100) DOPED TELLURIUM

o b 2

LASER ENERGY DENSITY (JOULE/cm2)

Fig. 5: Typical correlation between amount of gold deposit and laser energy density. 314

Typical I-V curve measurements of LASER PLATED GOLD N-GaAl: T« (100) Schottky barrier WAVE LENGTH = 0.S3 |im: ' LASER ENERGY =0.28 JOULE/cm2 contacts and ohmic DEPOSIT RESISTANCE AND CONTACT resistance of gold MEASUREMENT (# 9) deposits on GaAs substrate.

TYPICAL SCHOTTKV HARRIER CONTACT. TWO POINT PROBIIH MEASURE.

TYPICAL OHMIC RESISTANCE DETERMINED FROM I-V CURVE MEASURED ON THE SAME LINE.

LASER PLATED GOLD ON N-GoAi DOPED TELLURIUM -,20 3 q Dependence of gold a ,6 deposit resistance on laser energy density.

0.1 0.2 03 LASER ENERGY DENSITY (JOULE/cm2) 315

VAPOR PHASE SOLDERING OF SURFACE MOUNTED ELECTRONIC ASSEMBLES

E. Falkenstein and I. Fainaro

filbit Computers Ltd. Materials Engineering Dept.

INTRODUCTION The ever increasing need for more compact electronic circuitry Drought up new electronic packaging solutions, such as the leadless chip carriers and the surface mounting technology in general. One of the principal steps in the implementation of this technology is the interconnecting stage of surface mounted components to the wired substrate boards. This usually consists of a soldering process which can be realized by a variety of .tieans, a highly effective one being vapor phase reflow soldering.

VAPOR PIW.SE APARATOS

The condensation soldering process was developed and introduced into production in 1973 at the Western Electric Company. It has been proven to be an effective mass soldering method (fig. 1) • The condensation soldering process, or, as more commonly known as the Vapor Phase Soldering V.P.S. has been used successfully to reflow many varieties of electronic and other assemblies. The basic system consists of a stainless steel chamber, immersion heating elements, and vapor condensing coils. A processing liquid is placed in the chamber and heated to boiling, thus producing a zone of saturated vapor above the liquid. Parts to be reflowed are introduced into this vapor which condenses, giving up its latent heat. Systems are unpressurized thus the system is always at the boiling point of the fluid. This affords absolute control of the temperature, regardless of the length of the reflow cycle. Operating temperature is determined by the fluid selected. Fluids currently used are chemically inert, and create an oxygen free heating environment. Since vapor condenses on all surfaces of the assemblies, heating is extremely uniform, independent of the geometry of the product. Recent advances in V.P.S. equipment have greatly improved the adaptability of this technology to the production environment.

PRIMARY FLUIDS The primary fluid for V.P.S. Must have a boiling point sufficiently high to consistently produce high quality metallurgical bonds between solder and metals (about 20-30°C above the liquidus of the solder in use is generally acceptable). 316

Conversely the temperature must be sufficiently low to minimize thermal damage to any part of the product. The vapor as well as the liquid must be nonflammable, inert, chemically and thermally stable, low in toxicity in the operating temperature range, and of course to have the right heat transfer properties at an affordable price.

SECONDARY FLUID The principal purpose of the secondary vapor blanket is to act as a stable barrier between the primary vapor and the surrounding air. Therefore it must have a density between that of the primary vapor and air. Because of the continuous boiling-condensing cycle it must have a boiling point lower than the primary fluid. Complete miscibility of the primary and secondary fluids are desirable.

THE ELECTRONIC ASSEMBLY Since the whole purpose is to create a sound metallurgical joint between electronic components as the Leadless Chip Carrier LCC and the Printed Wiring Board PWB, Let us take a brief look at those parts. The ceramic LCC may be thought of as the die cavity portion of the ceramic dual in line package (DIP) having conductor patterns extending out to the exterior portion of the package. The external conductor patterns are reflow soldered to the appropriate PWB surface termination pads. The PWB that was chosed is a copper clad invar board coated with layers of polyimide which embeds the micro wire interconnection between the appropriate pads and the connector which provides the electrical connection to the rest of the system. The LCC's and connector are reflow soldered to the PWB to form a module.

APPLICATION OF SOLDER CREAMS Solder creams are used to provide the joining between the component leads and PWB termination area. The creams can be applied using mainly three basic methods: syringe application, stenciling and screen-printing, we use the two latter ones. The screen basically i~ prepared in the following manner: A Photosensitive material is placed on a taut screen in a metal frame, it is photo exposed and developed to obtain the desired pattern. In case of stencil the pattern is etched in a sheet of stainless steel or brass of suitable thickness. The PWB is placed under the screen or stencil in excact registration with the screen image. The solder cream is drawn across and through the screen by a squeegee with a sharp rubber edge.

ASSEM3LY REFLOW 5 CLEANING LCC components are placed at the appropriate pads in such a manner that there will be contact by means of the solder cream between the corresponding pads on the PWB and LCC. The whole assembly is baked to drive out the volatiles from the cream, then it is run through the reflow cycle to obtain a sound metallurgical bond. After soldering the assembly is cleaned and inspected as required. 317

SOLDER CREAMS

Solder cream consist of flux binder and solder alloy. The rheology of the flux binder is such that the dispersion of heavy solder particles (85-90w/o) can be maintained under normal room temperature and pressure conditions. The flux binder contains a minimum of 60% pure rosin and 40% other ingredients including activators, solvents, lubricants and thickeners. The alloy is dispersed in the cream in form of tiny particles. REQUIREMENTS FROM A SOLDERED JOINT

1. Good metallurgical bond. 2. Complete solder coverage ot all metalized areas. 3. Adequate space between the PWB and LCC (a minimum of 0.025 mm of solder) 4. Good alignment of the termination pads. 5. Nice appearance, no cracks.

SECONDARY VAPOR (VAPOR BLANKET)

PRIMARY SATURATED VAPOR

Fig. 1 "Batch type vapor phase reflow system 31S

EXPERIMENTAL

Various aspects of the soldering and related processes were studied and their impact on the quality of the resulting solder joint was investigated. Main issues of concern included activation and pre- tinning of components and substrates, solder paste materials and properties, printing of solder pastes, placing of components on solder printed boards, drying of solder pastes, vapor phase soldering,cleaning of soldered assemblies, and environmental testing.

INITIAL CONDITIONING Wired substrate boards are usually received with pre-tinned footprint pads. Their degree of solderability largely depends on the extent of oxidation on their surfaces. In order to assure their fitness for consequent solderinc operations a tinning action is commonly carried out. Unless greasing or serious tarnishing occured, this process yields a fresh uniform coating of a highly solderable character. Preliminary cleaning of the boards in a vapor degreaser enhances the solderability and is common practice. Tinning of boards is conveniently accomplished by dipping, by wave- soldering, or by solder-paste printing followed by vapor phase reflow. We prefer the latter method since it performs more selectively with no need for masking, and also affords higher thickness of coating. The quantity of solder applied in this stage contributes substantially to the total amount of solder in the eventual joint. If solderability problems are evident, a sequence of cleaning, masking and wave-or dip- tinning can be followed up and, if necessary, repeated again. Ceramic leadless chip carrier components, of the type we use, usually come with gold plated castellations. Carefully handled, those components maintain their good solderability properties for extended periods of time. Prior to being placed on the substrate boards they are tinned by dipping or in a soiaer wave. This action dissolves the initial gold plating. After tinning, flux residue contamination is removed- in a vapor degreaser cleaning plant.

PRINTING THE SOLDER PASTE

Solder paste printing is successfully accomplished either by screening or by stencilling. Screens are elasti-. and spring back better than stencils. On the other hand, stencils are more resistant and easier to clean while permitting higher print thickness. We tried both methods anc? -.and to adopt the latter. The viscosity of the solder paste should comply with the printing method. We found a viscosity of 500000 - 600000 cps appropriate for stencilling ,and somewhat lower values for screening. Our experiments with various thicknesses of stencils indicate that a thickness of 0.2-0.3 mm is satisfactory with regard to uniformity and line definition. After the solder paste being printed, the components are placed on the sucstrate boards using vacuum pincers. The process of placing the components is critical to the extent that such defects as slight twists, translations or excessive pressure on the components may 319

cause bridging and solder balling in the soldered assembly. From a point of view of accuracy as well as speed, automation of the component placing process is highly recommended. The substrate boards, bearing the components placed on the wet solder paste, are transferred to a drying oven, heated to 80-100°C for 15-30 minutes. The drying drives out volatiles which may impair the reflow jointing process. Care should be taken not to overheat the assemblies in the oven since this may result in excessive oxidation as well as spreading of the solder paste. VAPOR PHASE SOLDERING The dried assemblies are now ready for "soldering. We use a batch type vapor phase reflow soldering unit, manufactured by Hybride Technology Corp. (Mass. USA). It operates with Fluorinert FC-70 (3M brand)liquid, producing a primary vapor of 215°C. A secondary blanket zone is produced of Freon Ti? vapors. Assemblies to be soldered are loaded onto an elevator basket which is then lowered at a speed of about lm/min into the vapor zone. It resides about 1/2 min in the primary vapor and is raised into the secondary vapor zone.- for a cooling period, prior to exiting the tank. The controlled and effective nature of heating by this method prevents over-heating and permits perfect soldering even for very densely populated boards.

CLEANING Post-solder cleaning of LCC assemblies is considerably more difficult than that of the equivalent DIP, due to the fact that the contaminated areas are hard to reach. In addition to being difficult to clean, contamination of surface mounted assemblies is also difficult to detect. The conventional conductance test may yield erroneous results since contaminants confined by the closely overlying components may not be easily leached off. We have performed some conductance tests on LCC assemblies and found that the cleaning time to reach the allowed level of ionic contamination is as much as three times longer than with DIPs.

CONCLUSION Vapor phase soldering combined with solder paste printing prove to form an extremely effective process of interconnecting surface mounted electronic assemblies. Still, its successful implementation depends on optimizing all stages of the process.

REFERENCES 1. Carmen Capillo "The Assembly of LCC to PWB" Presented at IPC Fall meeting, October 1982,San Diego, California. 2. G.M. Wenger; L.L. Mahajan "Condensation Soldering" IPC Assembly - Joining Handbook. 3. D.J. Peck, D.J. Spigarelli "Development of Continuous Vapor Phase Systems" HTC Pond Lane, Concord Mass. 4. F.J. Dance "Mounting LCC on PCBs'",Electronic Prod. June 1982. 5. C.A. Mackay "Solder Creams and How to Use them" Electronic Packaging and Production , Feb. 1981 . 320

LATTICE, GRAIN BOUNDARY AND SHORT-CIRCUITS DIFFUSION OF PHOSPHORUS IN TaSi2 THIN FILMS

J. Pelleg*

Materials Engineering Department, Ben Gurion University of the Negev, Beer-Sheva, Israel

[NTRODUCTION

TaSi2/n polycrystalline Si(Poly-Si) was found [1] to be an attractive gate metallization in device fabrication. The reduction in the gate level sheet resistance to about 2n/d, while simultaneously maintaining other beneficial properties such as patterning, stability during device fabrication and the simultaneous retrofit to the conventional Si-gate very large scale ingegrated (VLSI) processes were the prime incentive to explore the possibility of using TaSi2/n poly-Si for application in VLSI technology.

Information on dopant migration and its distribution in composite layer gate materials is lacking. As the packing density of discrete units on 3i chips increases the diffusion of dopant inpurities acquires increasing importance. Thus, in a recent work P was found to diffuse rapidly through the WSi2 [2], causing threshold voltage changes in the MOS (metal-oxide- si liconj capacitors.

Fhe objective of this work was to identify the mechanisms contributing to the overall transport, discriminate between them and derive diffusion data for them. The possibility of using cosputtered silicides directly as the gate metal on gate oxide cannot principally be ruled out [3] and therefore the diffusion specimens used in the current experiments consisted of TaSi2/poly-Si and TaSi /SiO_ configurations.

EXPERIMENTAL

3i wafers were used as substrates. A thermal oxide of 600 S thickness was grown on them in dry oxygen at 1000°C. A 3000 ft undoped poly-Si was then deposited by low pressure chemical vapor deposition (LPCVD). On another set of Si wafers only thermal oxide was grown to a 1000 8 thickness. Ta and Si in a nominal ratio of 1:2 were cosputtered on these wafers to a nominal thickness of ^> 2500 S. The wafers were then sintered at 900°C/30 min in a flowing Ar ambient to form the crystalline TaSi_. The resulting crystalline TaSi2 had a very fine grained structure of aBout 0.05-0.1 urn. The sheet resistance of the TaSi7/Si0~ was 2.75n/D and that of TaSi?/ poly-Si was 2.2n/D-

A layer of 0.2 uC carrier free P radioisotope was deposited on each of the wafers by evaporation from a tungsten filament and it was used as an

* Work performed in Bell Laboratories, Murray Hill, NJ 07974, U.S.A. 321

instantaneous source of tracer in the diffusion measurements. In order to eliminate loss of P by evaporation, the surface of each wafer was capped by a thin layer of about 60 X Ta.

Quadrant wafers of the TaSi2/poly-Si and TaSi/SiO2 specimens werg diffusion annealed simultaneously in the temperature range of 500-900 C for short times in a flowing Ar aaibienf.. This temperature is below 0.5Tm, and therefore contribution from short circuits may be expected. The specimens were submicron sectioned [4] and the activity of P was measured by liquid scintilation (LS) and also by the residual activity (RA) method.

RESULTS

For short-circuits (SC) and grain bounday diffusion (GBD) analysis of the experimental data the solution of Whipple [4] - Suzuoka [5,6] was used in the form [7,8] given by:

Here D, ., refers to the diffusivity either in the SCs (D,) or in GBs D is the lattice diffusivity, t is the diffusion annealing time, 5 is the GB width, C is the concentration of the diffusant, y is the distance and K is the segragation factor given by K=C,/C,. C, and C., are the concentrations of the diffusant impurity in theuBs(or SCs) and/the lattice. K is usually unknown. According to this expression logC vs. Y should be linear if D and D, ,, are independent of composition and position. From the slopes of such curves and known values of D, only the product K6D, ,, can be determined. Figs. 1 and 2 summarize the results of the measurements by RA technique. Each profile of Fig. 1 can be divided into three parts: the near surface region (associated with the Ta capping and no physical significance will be attached to this at present), the linear tail and the intermediate region. There is a striking difference between the slopes of these curves and those shown in Fig. 2, where an additional contributing mechanism seems to have been observed. This is indicated by the arrows. In order to evaluate KflD, ., from the profiles, D is needed and can be evaluated by the technique given by Pelleg [5]. The temperature dependence of D, K<5D, and K'AD, (here K1 has the same meaning as K and A is the width of the SC) forTaSi2/poly-Si and TaSi2/Si02 are given by:

12 2 TaSi2/poly-Si: D = 4.21xl0" exp(-0.67/kT)cm /sec 15 K

12 2 TaSi2/SiOs: D = 5.05xlO~ exp(-0.63/kT)cm /sec 17 3 K

DISCUSSION The predominant feature of all the penetration profiles of TaSi^/SiO- is the existence of a central region having a larger slope, a shorter tail region and an ill-defined near surface region if at all. It had to be postulated that an additional SC diffusion (SCD) was operating. It is believed that this is a result of segregation effects. When a buffer poly-Si exists between TaSi and the gate oxide, virtually an infinite source of Si is available for the diffusion of Si into TaSi- during the 900°C/30 min sintering, which can segregate into the open sites in GBs and SCs. The degree of plugging the high diffusivity paths in SCs and GBs depends on temperature and time. In TaSi^/poly-Si Sc channels were plugged to such an extent that the effective paths for diffusion become practically the same as the regular paths for diffusion in the lattice. GB channels, however, were affected to such a degree only, which permitted a distinct observation of GB diffusion (GBD).

In the case of TaSi- no supply of Si was available to plug SCs and GBs which remained open for rapid transport along their paths. Both, SCs and GBs could exert their effect on the overall diffusion as indicated by the penetration profiles. Indirect evidence supporting this concept can be inferred from experimental observations [1,3,6-7]. Also, recent work indicates that a solid solubility of Si in TaSi- exists [8]. In addition it has already been observed that the most stable TaSi. is not the stoichiometric but a Si rich silic:de [6].

SUMMARY

1) Si segregation has been suggested as the cause for the difference in the penetration profiles characterizing diffusion in TaSi_/poly-Si and TSi^/SiC^ 2

2) Diffusion in TaSi2/poly-Si occurs via GBs and the lattice. The presence of the poly-Si buffer layer acts as a source of Si for segregation in TaSi- which completely plugs SCs.

3. Diffusion in TaSi /SiO_ occurs in the lattice, SCs and GBs. Channels for rapid transfer of P along them remained open due to the absence of Si segregation.

4. Complete saturation of the open channels in TaSi /poly-Si by Si or by a more effective solute should slow down diffusion rates of dopants in TaSi

REFERENCES

1. S.P. Murarka, D.B. Fraser, A.K. Sinha and H.J. Levinstein, IEEE Trans. Electron Devices ED-27, 1409 (1980) . 2. M.Y. Tsai, H.H. Chack, L.M. Ephrath, B.L. Crowder, A. Cramer, R. S. Bennett, C.J. Lucchese and M.R. Wordeman, J. Electrochem. Soc, 128, 2207 (1981) . 3. S.P. Murarka, J. Vac. Sci. Technol., 17, 775 (1980). 4. J. Pelleg and S.P. Murarka, J. Appl. Phys., 54, 1337 (1983). 5. J. Pelleg, Thin Solid Films, 110, 115 (1983). 323

6. S.P. Murarka, D.B. Fraser, W.S. Lindenberger and A.K. Sinha, J. Appl. Phys., 51, 3241 (1980). 7. S.P. Murarka, M.H. Read, C.J. Doherty and D.B. Fraser, J. Electrochem. Soc, 129, 293 (1982). 8. J. Pelleg and L. Zevin, unpublished work, 1984.

Fig- 1. P penetration profiles in TaSi2/poly- Si specimens. Data points beyond the arrows are associated with GBD.

00 10 20 PENETRATION DISTANCE 6'5 MO"6cme'5)

Fig. 2. P penetration profiles in TaSi2/Si02 specimens. Three regions are seen. The segment between the arrows is related to SCD.

0.0 10 2 0 3 0 4.0 PENETRATION DISTANCE6'5 (IO"6cm6"l 324

THE LIMM TECHNIQUE FOR DETERMINATION OF THE SPATIAL DISTRIBUTION OF POLARIZATION AND SPACE CHARGE IN POLYMER ELECTRETS

Sidney Bo Lang

Ben-Gurion University of the Negev, Beer Sheva, Israel

A recently developed technique for determining the spatial distribution of polarization in polymer electrets has been extended to the case of combined polarization and space charge.

Recently, the Laser Intensity Modulation Method (LIMM) for the accurate determination of the spatial distribution of polarization in polymer electrets was developed by Lang and Das-Gupta (1). This paper presents an extension of LIMM. to the case of combined polarization and space charge. The experimental system is illustrated in Fig. 1. Its essential features are the following,, A thin, electroded polymer electret film, which possesses combined polarization and space charge, is mounted in an evacuated sample chamber containing windows through which radiant energy can be admitted. Each surface of the film, in turn, is heated by means of a laser beam which is intensity-modulated at frequencies varying from 100 Hz to 100 kHz. The heat is absorbed on the electrodes and produces temperature waves which are propagated into the film and are attenuated more strongly as the frequency of modulation is increased as shown in

Fig. 1. Schematic diagram of experimental apparatus used in LIMM. 325

0-4 0-6 0-8 1-0 POSITION (x/L)

Fig. 2. Maximum amplitudes of thermal waves in a 25.4 y-m PVDF film.

0.5 1 POSITION (X/L)

Fig. 3. Assumed polarization distribution and calculated values.

Fig. 2. The nonuniform heating is convoluted with the polarization distribution to produce a pyroelectric current which is a function of the frequency, and the thermal diffusivity and thickness of the film. An analysis was developed for deconvoluting the current-frequency data to give the polarization distribution in the form of a Fourier cosine series (1). The method is extended by using the previously-developed electrostatic and thermal analysis to find a relationship between the pyroelectric current and an integral which is a function of the space charge distri- 326

I EXAMPLE A

1 § o \ \ I \ CO

0 0.5 POSITION (X/L)

Fig. 4. Assumed space charge distribution and calculated values. bution, the modulation frequency, and several thermal, optical, and geometrical constants. The unknown space charge is represented as a modified Fourier sine series. Finally, a linear relation is developed between the component of the current in phase with the laser beam modulation and the Fourier coefficients. The combination of this relation and the corresponding one developed for polarization are used in a least-squares analysis to give both the polarization and the space charge distributions. The correctness of the theory and the stability and the accuracy of the computational technique were verified by a simulation procedure. Figures 3 and 4 illustrate an example of assumed polarization and charge distributions, respectively, as shown by the solid lines. These values were used to simulate the current-frequency curves which would result from the combination of the assumed polarization and charge. The simulated data were used in the least-squares analysis to recalculate the polarization and charge, the results of which are ^hown as super imposed points in Figs. 3 and 4. The agreement with the assumed distributions is very good except near discontinuities, thus demonstrating the correctness of the theory* Experimental measurements were made on a 25.4-ym film of polyvinylidene fluoride (PVDF) which had been poled with a voltage of 1000 V for 15 min. at 110°C, and then cooled with the field applied. Calculated 327

50 (b)

25"

N CE < o o - K = 0.001 l = 0.00254 Q_ 0.8K OR 1.1L 1.ZK OR 0.9L -25 0 0.5 POSITION (X/L)

Fig. 5. Experimentally measured polarization distribution.

Fig. 6. Experimentally measured space charge distribution.

polarization and space charge distributions are shown in Figs. 5 and 6. The effects of errors in the thermal diffusivity and film thickness are illustrated and can be seen to be negligible. k detailed mathematical analysis of the method and an interpretation of the experimental data will be presented in a forthcoming publication.

Reference

1. S.E. Lang and D.K. Das-Gupta, Ferroelectrics (in press). 328

CHALCOGENIDE INFRARED GLASS FIBERS

A. Bornstein, N. Croitoru

Department of Electron Devices, School of Engineering Tel-Aviv university, Tel-Aviv 69978, Israel

Introduction The historical evolution of optical fiber communication systems as an alternative to radio and copper-based transmission links has recently been marked by the search for materials which exhibit low Dptical loss at progressively longer wavelengths. Since 1980 the progress in materials development has been spurred by the need both in fiber optics and in high power IR laser systems for transparent materials with absorption levels in the range of = 10~ cm" and down, four orders of magnitude below the levels required in conventional optics. These goals have now been achieved by chemical purification of Nad KC1, KB", CaF^, SiF . The losses in^IR fitter optics can be divided in to two groups: intrinsic losses and extrinsic losses. The extrinsic losses can also be divided into two groups: mechanical causes such as air bubbles, structural inhomogenities', imprefection in the core pattern and cladding surface which can be removed by precisely controlling and improving the drawing technique. The second cause for extrinsic losses is primarily associated with impurity absorption, such as oxide absorption. For example, the absorption of As 0, at 81C, 934, 980, 1100, 1165 cm" , and Si-0 band in 900 cm, or the aBsorption of H Se at 2200 cm, and at 3600 cm" by OH. The absorption due to Ee - Se bonding that usually appears in the 800 cm" area occurs because of a deviation in stoichiometry. The impurities of carbon and arsenic cause broad absorption throughout almost all of the IR region. These extrinsic ab- sorptions can bs removed by using oxide catchers (such as urea)', by working with very pure reagents and drawing fibers within a pure inert atmosphere. The most important factor is a very pure reagent and inert atmosphere. The intrinsic causes can be divided into three loss mechanisms. Two of them, absorption from electronic band edge tail and Rayleigh scattering, are responsible for a decrease in optical absorption with increasing wavelength. The third mechanism is multiphonon absorption, which causes increases in absorption with wavelength. This mechanism is the most important in the infrared region. The multiphonon absorption decreases with increased atomic weight of the elements involved in the absorption , thus justifying the usage of heavy elements. Therefore the multiphonon absorption moves to lower energy as the glass composition changes from Oxide to Sulphur, to Selenide or to Tellurite. The recent interest in developing infrared optical waveguides has been stimulated by advances in the use of C0? laser in surgery, cutting, welding, and heat treatment. The availability of such a wavegu- ide will play an important role in infrared image relay, remote measurements^ remote heating etc. Its theoretical high transparency may also serve for generating long distance communication links, although this application is more a thing of the future. 329

From a practical point of view glasses rather than crystalline materials are highly desirable for long optical fibers. Low toxic flour- ide glasses can be used in the near infrared region, while chalcogenide glasses are the most promising materials for a wider range of wavelengths. Although much research on chalcogenide glass has been,, d,oge, only limited results are available concerning fiber fabrication.c>°' The fibers we have been working with, are multimode fibers and are thick in relation to the wavelength. The main uses intended for the fibers are in power transmission, up to few watts in a single fiber, maintaining high energy density at the output port. This paper describes techniques for preparation of low toxic chalcogenide glasses and pulling fibers with precautions for eliminating oxide impurities formation. We evaluate optical perfection of these glasses and fibers, determine the maximum power that the existing fibers can withstand without being damaged by constant illumination (CW) and short pulses, and the angular scattering of the radiation at the output port. We include a search for a possible dependence of the output angular scattering on the focal length of the lens focusing the laser1, illumination at the fiber's end, and the fibers length.

Glass and Fiber Preparation; The chalcogenide glasses were prepared using 99.9992 pure As Se or As, Se, elements. The reagents were weighed and placed togetner in silica glass ampoules. The ampoules were sealed after being evacuated for several hours. The sealed ampoules were heated to 900 C for 5-24 hours in a resistance furnace. After the heating, the ampoules were withdrawn from the furnace and cooled in air. Glass rods 7mm in diameter and 30-150mm in length were obtained depending on the reagent quantity. We examined different methods for purifying samples from undesir- able oxygen. One of the methods that had been tried consisted of adding CO(NH )? (urea) to the raw material in the evacuated ampoule. The urea decomposes into CO, NH , N_ and H at 900 C. These gases removed all non-metallic oxides, sucn as Aso0o, Se0o, Sbo0, by reactions such as ASp0,+3C0 2As+3C0_; Fig. 1 compares the lnirared spectra of As Se glass samples before and after treatment with urea. It is very clear from this figure that the urea removed the absorption peaks at 80„„0„ c__m - an__dJ .^r-1050n c__.-m i . For-_.r. co iaser nne the most important improvement was the removing of the ; lOcrn" peak. Two systems were used for fibers MICROMETERS drawing; 6 IP II I2 Crucible Method; This system is intended for pulling an optical Co 80 fiber out of a melt ~ C02 LASER glass in a crucible. A ^ 60- glass ampoule contain- £ ing the chalcogenite g (70mm) WITH UREA NH4 raw material was put in 40 z the crucible in the <£ (40mm)_ Y fiber drawing machine [E 20 As 0 which includes: A fur- 4 6 nace with temperature 1800 1600 1400 1200 1000 800 control and a drun that WAVENUMBERfcnrf1) collects the fibers. 330

The drum rotates on an axis and moves in a direction perpendicular to the rotation, as its axis is threaded and is spun by a stable motor, whose speed is adjusted according to the desired fiber thickness. The fiber is drawn from a nozzle which is at the bottom of the fused silica crucible. The crucible is set in the center of a temperature controlled resistance furnace. The fiber diameter depends on the glass viscosity and drawing speed. The glass viscosity has a strong dependence on temperature and composition, so the furnace temperature •nust be very precisely controlled. The Rod Method; This system is intended for pulling an optical fiber DUt of a glass rod (preform). The system consists of a rod mount, a heating coil, collecting drum, and controller. The glass rod 6mm in iiameter and 2 cm long is mounted in a clamp that can be lowered or raised by a dc motor. The heating coil is fixed in its place and can maintain a fairly steady predetermined temperature which is measured by thermocouple, and automatically controlled. The rod is lowered so that 2mm of it is inserted into the heating coil.

Measurements and Comments; The optical losses through the glass fibers were measured by the cut back technique in which the power transmitted by a fiber is recorded before and after a length of the fiber is re- noved. A C0_ laser was used as the light source for the 10.6 m wavelength region while a pyroelectric IR detector was used for detecting the light immediately behind the fiber. Precise positioning of the fiber with respect to the focused laser spot was accomplished by means of an X-Y-Z translation stage, which allowed movement of the fiber end until a maximum power throughout was detected. The fiber was exposed to a predetermined C0_ laser illumination density. The i.iput power was calculated and the output power was measured . The laser's power was raised until the fiber was damaged. This was done with CW as well as pulsed radiation. The output scattering pattern was measured a) Near field: qualitative impression could be achieved by placing a liquid crystal coated plastic sheet near a fiber end. The liquid crystal changes its color when its temperature changes, and a pinhole thin dark spot appeared on the sheet. b) Far field scattering: The output end of the fiber was mounted at the center of a rotating table, on which an IR detector was installed, so that the detector could 'move1 in an arc centered at the fiber's output end. The detector's window was 2.5 cm away from the fiber. Since no sharp angular dependence of radiation intensity was detected a large window of 10 (4 mm width) was used for measurement's convenience sake. We have got extremely different results for different samples. Some fibers were not damaged at input power density even up to 10 kW/cm . We estimate that the great differences are mainly due to differences in raw material quality. The angular dependence of the field intensity at the output appeared to be constant (within the limits of measurements accuracy), when the focal length of the lens or the fiber length were changed. These results were expected since: a) If the laser beam was a pure Gaussian, the beam's waist was for any lens in use a few mm. long, so that the input field was always plane wave like. 331

b) For unclad fiber of high refractive index, no excessive attenuation of the higher order modes have been expected. These modes are responsible for the radiation in the wider angles. Most of the energy is ra- diated at angles 20°-i)0° measured from the fiber's optical axis. High energy density is detected right at the end of the fiber. A piece of paper was burnt by radiation transmitted through a fiber, when placed close enough to its end. The highest output power density was 0.8 kW cm . Two types of experiments were performed in order to measure the attenuation due to bend in the fiber. The experimental arrangement was set so that the only condition changed during the experiments, was the bend radius. All other conditions such as the inlet and the outlet of the fiber and the power of the laser were kept fixed. In the first experimental arrangement the fiber was bent around a cylinder to form a 90° angle. Cylinders of various diameters were used down to 2 cm. In the second experimental arrangement, the fiber had one full winding. By increasing the distance between the inlet and outlet, the winding diameter was changed down to 2 cm. No significant influence of the bending on output power was determined in the measurements resolution ability.

Conclusion: In this paper we have described the preparation of chalcogenide fiber in two different methods crucible method and rod method. Various measurements of infrared absorption power transmission output angular dependence, and bend dependence were described. Attenuation in chalcogenide glasses, and some of their implications were discussed. The results indicate that the losses in the fiber tested are 0.1 dB/cm. The attenuation stays unchanged even when the fiber is bent to circular radius of 2 cm. We did not find angular dependence of the field intensity at the output when the focal length of the lens or the fiber length were changed. We obtained fibers that were not damaged at power density up to 10 kW/cm and which transmit power density of 0.8 kW/cm . We did not detect any relation between the output power or scattering pattern and the fibers' radius of bending. Such relation was not expected since any radius of curvature obtainable with glass fibers is very large compared to the wavelength. We have examined the existing fibers quality, by the various measurements described here. Great improvements are needed for making these fibers commercial. Acknowledgement This research was supported by a grant from the Nation- al Council for Research and Developnent, Israel.

References

1. A.L. Gentile et al in "Optical Properties of Highly Transparent Solids", Edit. S. S. Mitra and E. Bendow, (Plenum Press, N.Y. 1975), p. 105. 2. J.A. Savage, P.J. Webber and A.M. Pitt, IR Phys. 2Q_, 313 (1980). 3. A. Bornstein, R. Reisfield, Non Crystalline Solids, 50, 23 (1982). 4. A. Bornstein', N. Croitorii, E. Marom, Advances in IR fibers; Los Angeles. Technical Symposium by Society of Photo-Optical Instru- mental Engineering (SPIE) January 1982. 332

PROGRAMMING OF CRYSTAL DIAMETER IN CZOCHRALSKI GROWTH BY A COOLING PLOT: APPLICATION TO InSb

M. Azoulay and Z. Burshtein

Nuclear Research Centre - Negev P.O.B. 9001, 84 190 Beer-Sheva, Israel

The growth of single crystals by the Czochralski method is very common for the growth of many types of crystals (1). Its main advantage is in preventing stresses associated with some other growth methods, such as the Stockbarger (2) or Bridgman (3) methods. In the latters, the interaction of the crystal with the ampoule may result in enhanced striae and fracture. In the Czochralski method, an oriented single crystal seed is dipped in the melt, which is contained in a crucible. The seed is then gradually pulled out, usually also being rotated around its axis. The goal of rotation is mainly to improve the radial symmetry of the thermal profile.

In the following we assume radial symmetry to the furnace and pulling system. Consider a crystal of diameter rc, being pulled out of a crucible of diameter ro>rc. The pulling distance will be denoted by z, and the melt height by hs. The values z=0 and hs=hi define the situation as the pulling starts. The crucible bottom is at hs=0. In Fig. 1 we illustrate a momentary situation during the pulling.

Let Tc be the temperature at a point rc in the crystal circumference, which is in contact with the melt. Then, Tc = Tc(hs,rc,t). A full diffe- rentiation of Tc gives

dT 3T dh 3T dr 3T c c c c c dt dt 3r dt 3t CD " 3h c When the temperature in the furnace is controlled, say by a thermocouple, at a fixed point, the temperature at that very point, To, is a function of time only. Therefore, by adding and subtracting 3TQ/"3t = dT0/dt, Eq. 1 may be written as

dT dT 3T dh 3T dr c _ o c s c c 3 , . , .. dt~ " d^ + aT dt~ 3r^ dt~ + 3t CTc " V * C2} We assume now that thermal "stiffness" conditions may be realized in the growing cell, i.e. that temperature differences in the crucible vicinity are independent of the stage of the process. In other words, "stiffness" means, that any change in To would change all the temperatures in the crucible vicinity by exactly the same amount. Else, it means that the crystal growth does not change the thermal profile. Later on we shall point out some measures that should help realize such conditions. With the above assumption, the fourth term on the right handside of Eq. 2 van- ishes. 333

From the balance of solidifying mass one easily obtains

dh* - dz dF~ • • ~i r 2 dt • l J ro pl/pc - rc

where p^ and pc are the melt and crystal densities, respectively. For the isoment we assume uhat the crystal radius rc is kept constant, i.e. drc/dt=0. Under the above conditions, and inserting Eq. 3. into Eq. 2., the latter reduces to

dT dT 3T r 2 (» , c _ _£ _c . c . dz_ ,„ dt " dt " 3h 2 . 2 dt ' v J s ro pl/pc " rc M

where rcC<») is the steady-state crystal radius.

In the course of growth of pure crystals from a stoichiometric melt the temperature Tc must remain constant. Therefore, dTc/dt=0. Then, we get

dT , 3T r 2O0 dt o " dtz 8hc 2 .c 2, • ( ^} 5 ro pl/pc - rc M

The physical meaning of Eq. 5. is clear. The temperature TQ must fall at a rate that would compensate the increase in the melt surface temperature as it drops towards the hotter regions in the furnace. Eq. 5. presents a straightforward recipe to present the crystal radius rr by a cooling plot. It requires a fair measurement of the longitudinal thermal gradient at the melt surface level 3Tc/3h. This may be done by measuring simultaneously the temperatures of two different points near the crucible, of different heights, but of the same distance from the symmetry axis. It is interesting to note, that the radial thermal gradients play no role in Eq. 5. The only pertinent material properties are the melt-to-solid densities ratio pl/pc- We have examined the foregoing idea in the growth of pure InSb single crystals. The crucible was made of silica, with three thermocouples attached to its side, to measure the longitudinal temperature gradient. The radial gradient has not been measured. The growing cell was a silica tube, with a resistivity heating element surrounding its central zone. A constant 30cc/min flow of hydrogen at normal pressure through the cell served to eliminate oxides from the melt surface, as well as to improve the heat exchange. The seed was a square shaped rod oriented along the (211) crystallographic direction. In Fig. 2 we show a photograph of a crystal grown in our system following a cooling plot derived from Eq. 5. The growing conditions were: Liquid to solid densities ratio p^p = 1.14; Longi- tudinal temperature gradient 3Tc/3hs = -12.5°C/cm; Pulling rate dz/dt = 334

2.2cm/hr; Crucible radius rQ = 2.0cm. The cooling plot set in order to get rc0°) = 1.1cm was dT0/dt = 8.3°C/hr. No program was set for the broad- ening of the radius. As seen, the crystal radius grew gradually, reaching a final constant value of 1.0cm ± 5% (actually, (area/ir)'5, since the cross section is approximately trapezoidal). This value was kept constant over a length of 5 cm.

The above result is an example of the good fit that has been obtained between the predicted radiuses and the experimental results. Par- ticularly, it proves that the temperatures, measured by the thermocouples attached to the crucible side, represent the actual longitudinal temperature profile in the melt. Still, some error might be involved by our simplifying assumptions.

FIG. 1 Illustration of momentary situation of crystal and cruci- FIG. 2 ble during pulling. The thermocouple at height h is used Photograph of InSb single Q to control the heating element crystal grown with a cooling power. Other thermocouples may plot following Eq. 5. be distributed around the crucible to measure the thermal gradients in the course of crystal growth. 335

Thus, has been shown that under conditions where the thermal profile in the crucible vicinity may be assumed to be rigid over a certain temperature range, and independent of the growing stage, the radius of a crystal grown by the Czochralski method may be programmed by a simple cooling plot. The pertinent parameters are the longitudinal and radial thermal gradients, the crucible radius, and the melt-to-solid densities ratio. A correction, to account for the change in freezing temperature, is required for the growing of highly doped crystals.

The following measures should help realize the required conditions: Reducing the thermal masses of the hot parts, increasing the thermal con- ductivities, generating high uniform temperature gradients near the crucible - negative in the longitudinal upward direction, and positive in the radial direction, and pulling sufficiently slow. We have shown, that our principle is practical in growing InSb crystals of diameters around 2cm. It should presumably be the same for many other materials. The principle put forward might save the need for complicated diameter control systems, particularly in small-scale crystal growing systems for laboratory use. We have not considered here some other important effects, such as self and forced convections in the melt, entailing the thermal gradients and seed rotation speed (4,5). These might influence the crystal perfection, and should be considered independently.

REFERENCES

1. J.C.Brice, The Growth of Crystals from Liquids, Selected Topics in Solid-State Physics, Vol. 7, E.P. Wohlfarth, editor (North-Ho11and, 1973) Ch. 7. 2. D.C.Stockbarger, Rev. Sci. Instr. 7_, 133(1936) 3. P.W.Bridgman, Proc. Am. Acad. Arts Sci. 6£, 305(1925) 4. J.R.Carruthers, J. Cryst. Growth, 36_, 212(1976) 5. Y.Miyazawa, Y.Mori, S.Homma and K.Kitamura, Mat. Res. Bull., 13, 675(1978) ~ 336

TOE DEPENDENCE OF THE HIGH TEMPERATURE, HIGH SOLAR FLUX STABILITY OF MATERIALS ON SURFACE STRUCTURE AND COMPOSITION

A. Ignatiev

Department of Physics and Chemistry University of Houston Houston, Texas 77004

INTRODUCTION The stability of materials at elevated temperatures has long been a point of great interest in materials research.(l>2) Until recently, however, such interest has centered principally on the stability of a material under high temperatures generated by the absorption of either infrared or particle radiation. Such are not the conditions encountered by materials used at elevated temperatures in solar apparatus. In this case the materials are exposed to concentrated solar radiation of wavelengths from -0.3 ]im to - 2iim. Radiation from this solar spectrum is by and large absorbed over a much shallower depth (up to a factor of 100 to 1000 shallower!-*)) than infrared radiation and hence this surface localized absorption can lead to a surface temperature much higher than the measured bulk temperature. Such an effect may prove deleterious on several fronts. In addition, the ultraviolet and near ultraviolet components of the solar radiation may indues various chemical changes at the surface which may prove to be either beneficial or deliterious in terms of affecting the long term stability of the material in a solar enviroment. It is clearly an enhancement of long term stability under elevated temperatures in solar environments that is sought after in solar materials. The work described below will show that such an enhancement can become a possibility when the basic mechanisms responsible for the noted photoeffects are identified and characterized. It is witn that knowledge as an input that efforts can be made to generate new solar stable materials.

PHOTOEFFECTS IN SOLAR MATERIAL ABSORBER COATINGS We have discussed in the past specific photoi iduced effects active in 4 5 6 the black chromev - ) and black cobalt( ) and wish to note here some of the details of the studies as background for our current investigation ut photoeffects in metals and ceramics. For both electroplated black chrome and black cobalt (formed by oxidation of plated cobalt metal(7)) it was observed that the reflectance of the samples changed less from the "as prepared" case upon solar heating than upon infrared (oven) neating. Tne differences were in fact quite significant (Figure 1) e.g., for black cobalt under oven heat at 460°C in air for 50 hours there was a resultant decrease in solar absorptance Aα » -0.11, whereas for an identical solar heated sample under equivalent (bulk) temperature conditions, the solar absorptance decreased only uy Aα « -0.S. 337

100

»S PfSEP.

OVEN HEM

03 u U IB SI WtVElEMTH (pa)

Figure 1. Spectral hemispherical reflectance for black cobalt samples that were solar simulator irradiated in air at 6UQ fcH/m? and 46QUC for 50 hours, oven heated in air at 160"C fcr 50 hours ana for and "as prepared" sample.

The basis for such a photoeffect was investigated by applying the surface sensitive techniques of Auger electron spectroscopy (AES), x-ray photoelectron spectroscopy (XPS) and mass spectrometry. It was found in Qcth cases that the effect was jjhotodesorption of oxygen bearing species (CO2) from the surfaces of the absorber coatings as a result of the near ultraviolet component of the solar radiation. This photodesorption reduced the nunber of oxygen species at a surface available for oxidation and hence reduced the rate of oxidation - oxidation being the principal degradation mode in black cnrome and black cobalt. It is seen here that the photoeffect in these two materials has been quite beneficial and should therefore be utilized to its fullest extent when designing solar absorber appa.-atus. Other materials are also expected to be used in solar technology, for reflectors, insulators and structural materials and it is of importance that active photoeffects be characterized in the systems. METALS A large variety of metals are expected to be utilized in the concentrated solar enviroment. These include iron, chromium, aluminum, copper, nickel and various steels. It is of importance, therefore, to define and characterize any photoeffects that significantly affect the stability in these systems. This nas been done to date for chromiun, iron and aluminum samples. Chromium samples exposed to ~700 kW/m2 solar simulated radiation exhibited two modes of behavior depending on cleaning history of the sample surface. Samples cleaned in ultrahigh vacuum by argon ion bom- 338

bardment such that the surface was atomically clean showed no photoeffects under solar irradiation. However , samples not cleaned, i.e. retaining the native oxide and other surface impurities., showed marked photodesorption of CO2 upon solar irradiation (Figure 2). The desorption efficiency was measured by mass spectrometry to he K3 X 10-6 molecules/photon under irradiation at 3I100A and varied linearly with photon flux. The wavelength threshold for the photodesorption was near 5000A (~2.5eV) and thus the effect did not correspond to any known photo- induced effect in metals. The threshold behavior did, however, nearly correspond to that expected for photodesorption from Cr?^. There are presently discussions underway which address the discrepancy in the measured and expected photodesorsption thresholds for l>2^3. (2.5 eV vs 3.4 ev)(a>9) with the basis for the discrepancy lying in the most appropriate description of the basic mechanism responsible for the desorption (10,11)^ j^ -jS fair to say, however, that a photo desurption roechamism(s) is (are) active in Cr2(J3 and similar to that previously noted for the black chrome absorber coatings it reduces significantly the oxidation rate of the chromium.

SOUR HEAT

I OVEN HEAT A . .Ji - 1

2 II 21 44 MASS NUMBER

Figure 2. Hass spectra of solar-simulator irradiated, air exposed chromium (660 kW/m2 at 410°C in 5x10-'° torr) and infrared heated air exposed cnromium (410°C). Note the large increase in the 44 ainu (CO2) peak under solar simulated irradiation.

Iron samples exposed to solar simulated radiation showed no photodesorption effects when the surface was atonically clean. They did, however, exhibit photodesorption of CU^ <*t an efficiency almost two orders ot magnitude lower than that of"chromium for a native oxide coatad surface. On the other hand, iron samples exposed in air to solar simulated radiation at elevated temperatures showed a new photoefect. The samples exposed at 6U0 to 700 kW/rn- fluxes at 410°C, 5O5°C all showed increased rates of oxidation under solar irradiation as compared to oven heating in air. Figure 3 shows AES depth profiles (obtained by bom- barding with argon ions) of a solar simulator irradiated iron sample and equivalent sample heated in an oven. It is clear that the solar irradiated sample has an oxide coating that is "-30^ tnicker than that of trie oven 339

Heated one, i.e. it has undergone enhanced oxidation under the influence of the solar simulated radiation. The basis for this is believed to be enhanced dissociation of molecular oxygen at the iron sur-fdce by the solar photons thus generating more highly reactive oxygen atoms for participation in the oxidation process. We have then, in the case of iron a detrimental effect due to solar irradiation of a material and must clearly include this possibility in the utilization of iron in concentrated solar environments.

OVEN HEAT

-—5=-^- SOUR MEJIT

Fl w

-. C_

DIM HUTltlK tlit (all)

rtgure 3. Auyer electron spectroscopy/ion bubardment depth profiles for oven heated and solar simulator heated (in dir) iron samples (110DC. 30 minutes, 660 kH/m2). The Sputtering Tune can be converted to depth into tne materials by the factor of -50 ft/minutes. Note the thicker oxide layer (oxygen signal extending to -23 minutes) for the solar heated sample as compared to the oven heated sample.

Irradiation of aluminum by concentrated solar simulated radiation results in different effects in two specific temperature regimes, below 40n°C and above 400°C with e flux dependence also observed above 400°C. The AES depth profile of an aluminum sample solar irradiated for 2 hours at 520°C in air at ~ 1.3 MU/m2 is shown in Figure 4a. For comparison, q depth profile of a sample solar irradiated at 520°C in air at ~ 25OkW/m2 is shown in Figure 4b and d depth profile for a sample heated in an air oven at 520°^ is shown in Figure 4c. Several points can be made in the com- parisicn. 1) The high flux solar heated sample has a thinner oxide (- 250A) as determined from the ion etch rate during profiling, than the low flux solar heated sample (-3S0A) or the oven heated sample (-350A). 2) The stoichioinetry of the high flux sample exhibits a higher oxygen to dluuinun ratio in the surface region than does the low flux sample. However, the stoichioTiutry is quite similar to that of the oven heated sample. Tiie stoicniometries of the oven heated and hign flux heated samples as obtained from the figures are approximately: oven - Mβ^; high flux - AJ^OS^. Tiiese stoichiometrics are obtained from the atomic concentrations in tne depth profiles as determined from the ALS peak-to-peak heights and AES sensitivity factors!12). The sensitivity factors have ~20% uncertainty in them and are dependent on apparatus used. Therefore, it is appropriate to denote (as is expected) the surface region of the oven heated alonintin sample as fully oxidizied 340

The surface region of the high flux heated sample is then also AA2O3, however the low flux heated sample is noticably different at - AJIO, i.e. denoting a not fully oxidized surface region. 3) The top few 10's of Angstroms of the high flux heatad sample are much reduced in oxygen concentration with most of the aluminum reduced to the metallic state. The above differences, although detailed only for samples treated at ~520°C are consistently observed for samples treated at 470°C, 4()0°C and 430°C.

— M

V HIGH FLUX- 1.3 mM 52Q"C : L /\ n

Spuitf

— Al

Law Fmx-2a) KX/M' 520"C

5yut(er Tine (

fl) /

\ / riVLN \/ 520"C /\ 7 V ^y \ fl "^ g

Figure 4. AE5 depth profiles giving the atonic concentration of oxygen and aluminum for aluminum samples heated under different conditions.

Irradiation of aluminum samples below ~4(J0°C results in very little difference between o.'en heated, high flux solar heated and low flux solar heated samples. The stoichiometrics are all near AJI2U3 and the thickness are all comparable and quite thin (<50A). The mechanisms responsible for the noted behavior of aluminum exposed to concentrated solar simulated radiation are not yet fully understood, however, several points can be made. i)The observed thinner oxide under the high flux irradiation probably indicates a surface temperature m.icii higher than that measured by the thermocouple. Less oxide growth his 341

been observed in aluminum above 600°C.U3) ii)The reduced amount of oxygen within the very top layers of the high flux sample indicates the presence of an active photo-induced effect which reduces the oxide at the surface of the sample. Such an effect may be photodesorption of oxygen bearing species, and this can be and is currently being tested for by tne irradiation of an oxidized sample in vacuum while monitoring desorbed species by mass spectrometry. The lower oxygen to aluminum ratio in the low flux heated sample as compared to the high flux or oven heated samples may have its basis in lower thermal gradients in the sample due to the low flux condition, and a reduced, but still active photodesorption effect. The lower thermal gradient generating sample temperature uniformity approaching ttiat of tne oven neated sample, result in oxidation of the sample to a depth equivalent to the observed in the oven heated sample. The reduced oxygen to aluminum ratio in the surface region.however, as well as some ennanced oxide reduction in the very top surface of the sample is probably a result of photodesorption of oxygen beaming species under the low flux irradiation conditions. The decrease in reduction of the oxide at the very top surface in the low flux sample as compared to the high flux sample directly indicates the effect of the magnitude of the solar flux on the samples. Quantitative measurements are now underway to describe the effect more fully.

CERAMICS In working with the aluminum system we nave begun to grow thick (3-5ym) thermal /U2U3 layers on aluminum for the initial evaluation of specific photoeffects in ceramics. Preliminary measurements show two effects: i) photodesorption of oxygen bearing species for aluminum oxide solar irradiated at a flux of 1.5MU/m2 anc| -d^ a temperature of approximately 400°C in vacuum; ii) discoloration of tiie coating - darkening, under ~1.5 flW/m2 in air at ~400± 1UU°C. Optical measurements to define changes in reflectance are to be undertaken, as ar^ high resolution (lym spatial resolution) AES measurements to define the basis for the discoloration. The discoloration (darkening) clearly increases the solar absorptance of the aluminum oxide sample and could result in locally high energy absorsp- tion leading to the cracking and melting previously observed at Gl1 and SANDIA.ll4).

CONCLUSION It has now been clearly shown that a nunber of solars relevant iriaterials exhibit strong photo-induced effects which may be deleterious or beneficial. It is such effects that nust be identified, characterized, and understood with respect to the basic mechanisms responsible for them. With this knowledge it is then possible to design solar system components that are highly stable under concentrated solar environments. ACKNOWLEDGEMENTS The assistance of A. Mesawri and A. Zomorrodian is greatly acknowledged. Support for this work has been provided by the Department of Energy - STAKC and by the University of Houston Energy Laboratory. 342

REFERENCES

1 High Temperature OxidatJ_on_- jtesi stajnt _Coati_ngs_, ed., National •Res"e¥rcH"Co"uncfl" Tlat." Head""of Sci., (T97tT). 2. I.E. Campbell and E.M. Sherwood, High Temperature Materials and Technology, (Wiley, New York, 19677". "" ' 3. J.P. Jackson, Classical Electrodynamics, p. 225 (Wiley, New York 1%5 4. G.B. SMith, G. Zajac and" A." IgnatieV, Solar Energy 2$_, 279 (1982). 5. A. Ignatiev, G. Zajac and G.b. Smith, Proc. SPIt 3?A, 170 (1982). 6. A. Ignatiev, Yearly Prog. Report, DOE - STARC, Task"? (Houston, 1982) 7. G.B. Smith, A. Iynatiev and G. Zajac, J. Appl. Phys. _5J., 4136 (1980). 8. L. Korenblit and A. Ignatiev, Surf. .Sci (1984). 9. 6.W. Fabel, S.M. Cox and 0. Lichtman, Surf. Sci., £0, 571 (1973). 10. P. Mark, R.C.A. Review 2^, 461 (1965). 11. 0. Lichtman and Y. Shapird, in Diemistry and_ Phys_i_cs of So_l_ids_-_H, ed., R.Vanselow p. 397 (CRC PreTsTT9T8)T 12. Handbook_^of_ Auger_^p_ectros_copy_, ed. L.E. Oavis et. al. (Physical Tnectromcsr~ffinn. ,1976)." 13. P.Doherty and R. Davis, J. Appl. Phys. 34, 619 (1963). 14. High Tenp. Materials and Ther.nal Science Coord. Meeting, (SEKI, Golden 1983). 343

TRANSPARENT CONDUCTOR FILMS AS A MATERIAL FOR PHOTOVGLTAIC JUNCTIONS WITH POLYCRYSTALLINE SILICON

Z. Harzion, M. Zafrir, J. Rishpon*, S. "ottesfeld* and N. Croitoru

Tel-Aviv University, Faculty of Engineering, Dept. of Electron Devices Tel-Aviv, 69978, Israel *Dept. of Chemistry

INTRODUCTION In recent years, a great effort has been expended to find a material which might replace single crystal Si for the manufacture of terres- trial solar cells. Polycrystalline silicon [poly-Si) is one of the candidates which might yield inexpensive solar cells. Such cells are made of p type poly-Si as the base converting semiconductor, and a transparent conductive oxide (e.g. ITO) as the other element of the junction. Tcansparent conductors are degenerated wide bandgap materials, which are used as an antireflecting coating and window as well as active layer and upper contact in surface barrier devices. The poly-Si consists of small crystallites, and therefore, grain boundaries (GB) appear to be the main reason for the low conversion efficiency and the relatively short term stability of these cells compared to the single crystal silicon solar cells. Much work has been done on the theoretical modeling of GB effects in poly-Si (1,2). It has been demonstrated that cell performance is mainly controlled by intragrain defect density and less by the grain size. GB are expected to contain high density of interfacial states which act as minority carrier recombinations sites, and produce a potential barrier against the carriers transport. Increase of the efficiency will only be possible if the limitation of GB recombination can be overcome (e.g. by passivation) . The efficiency of solar cells depends largely on the life-time of the carriers generated by the incident light. One of the methods to determine life-time is to study the response of the solar cells to light pulses. The method is usually applied to a p-n junction in a single crystal device. We have recently extended this method to transient measurements of photovoltage decay (PVD), photocurrent and photoluminescence in photoelectrochemical (PEC) solar cells (3,4). We have now applied this method also to poly-Si solar cells. The PVD as well as the impedance in a wide range of frequencies (0.1 Hz - 1 MHz) were measured and compared with those obtained for single crystal solar cells.

EXPERIMENTAL The preparation of ITO solar cells proceeded as follows: a) etching for 8 hours in KOH-methanol solution; b) deposition of Al film of 2000A and diffusion for back contact; c) sputter etching and ITO deposition by sputtering in MRC 820 system. Some solar cells were passivated with iodine saturated methanol solution before the ITO sputtering. I - V characteristic of the cells was recorded. The structure of the solar cell is shown in Fig. 1. PVD measurements were done by applying a short light pulse (10 nsec) on the cell. High impedance probe was connected across the cell in the 344

open circuit photovoltage mode. Impedance measurements of the cells were done in the range of 0.1 - 500 Hz (5). More detailed description of the PVD experimental system is given elsewhere (4) . I.-AI diffused contact RESULTS 2L-Silso polysilicon From the I-V characteristics of ITO/poly-Si solar cells 3.-2TO film (Schottky contact) under illumination, Voc was 4.-AI diffused, top contact determined as 0.35-0.38 V, Fig. 1 I *> 2mA/cm and the FF was luw (5?35). Passivation v.lth iodine ITO/poly-Si solar cell structure increased Voc by I. 50 mV and doubled I (a) The PVD measurements of poly-Si and .3 - single crystal ~,olar cells can be summar- x ized as follows: 1) the V at the o oc E .4 SINGLE CRYSTAL highest pulse, intensity (2pJ/cmA) for the poly-Si solar .2 cells was between POLYCRYSTALLINE 300 and 410 mV. For f- —i 1- 1 J i- the single crystal 0.5 1.0 solar cell, higher values were observed. Tmsec (500 mV); 2) the PVD curve has a non- (b) exponential shape 1.0 and lasts for 1-30 msec; 3) PVD of poly-Si solar cell .8 reduces to half of s. its initial value POLYCRYSTALLINE within 10 usec, .6 while for single crystal, it hardly .4 decays within the same time window (Fig. 2b); 4) At .2 the highest intensity of the incident light, 20% of the 8 10 initial photovoltage of the poly-Si solar T IO6sec cell decays at a rate of 18-60 mV/ Fig. 2 ysec, while for a. PVD curve of ITO/Si junction PVD single crystal solar b.Initial 10 psec region of the 345 cell, the rate is 1 i • 1 ' 1 ' 1 1 1 ' 6 mV/ysec. To g .^25.6 characterize the f 55,3 . shunting of the ITO/ 5 100 — V — poly-Si junction, the impedance of the solar cell was measured. A jjj 34 typical Cole-Cole % - curve at zero exter- r 1 , 1,1,1 Tig.T Theiwpo^y o" " 50 ' 100 150 " 200" 250 " 300 -Si solar cell has RE Z Kohm good shunting proper- F- 3 ties. The different- Cole/Cole plot of ITO/poly-Si solar cell; xal resistance is ,~ , , , ___ ,,_ , , some frequencies are marked 7 , about250 Kfi and does ^ not limit cell performance. The cell exhibited parallel R-C behaviour (almost ideal semi-circle in the range 0.1-500 Hz)-The series resistance was found in the range 5O-2OOS2.- DISCUSSION Some analytical solutions of PVD curves for single crystal junction have been published (6-9). For a poly-Si based junction, grain boundary effects should be considered. Even for a junction in the single crystal, the analytical expression of PVD does not take into account all the boundary conditions due to the complexity of the continuity equation. By assuming a finite thickness of the diode (6-7) , the effect of recombination at the back contact and at the emitter is important. An RC influence on the time constant of the cell was also found and should be considered (8,9). In our measur.e-.:i3Pts the slow region ('vL-lO msec) of the PVD curve is controlled by the RC time constant of the cell, in accordance with the literature (9). The effect of light intensity on the effective diffusion length (Le£f) must also be considered. In poly-Si, Leff increases with light intensity either by collapse of the space charge region caused by existence of intergrain clusters(10) or by trap saturation of deep trap levol (It). In spite of the high light intensity used in our measurements, the inixial PVD rate in the ITO/poly-Si solar cell is very fast (18-60 mV/ysec) compared to a single crystal Si/ITO junction (6 mV/ysec) and Si p-n junction (0.2-2 mV/ysec) (9). This means that the effective life-time (xeff) and Lef£ in poly-Si are shorter than in single crystal Si, thus reducing the collection efficiency of the photogenerated carriers, as seen from the photocurrent. The relatively high series resistance (50-200J2) also contributes to the low Isc. Part of the series resistance is due to the resistance of the ITO film itself, which has to be improved.

CONCLUSION The PVD measurements give information about the decay rate of the photocarriers while the Cole-Cole curve measures the impedance of the cells. These are two of the main factors which determine the solar cell efficiency. Passivation with iodine improves the efficiency. PVD and Cole-Cole measurements on the passivated solar cells will be performed and compared with those obtained for non-passivated solar cells. 346

ACKNOWLEDGEMENT This research was supported by a grant from the National Council for Research and Development, Israel and K.F.A. Julich, Germany. The authors wish also to thank M. Evenor for helping in the computer programming.

REFERENCES 1. H.C. Card, J.G. Show, G.C. McGonigl, D.J. Thomson, A.W. deGroot and K.C. Kao; Proc. of the 16th IEEE Photovolt. Spec. Conf; San Diego, Calif.(1982) 633.

2. N.C. Haider; Proc. of the 16th IEEE Photovolt. Spec. Conf; San Diego, Calif.(1982) 640.

3. 1. Harzion, D. Huppert, S. Gottesfeld and N. Croitoru; J. of Electroanal. Chem. 15£(1983) 571.

4. Z. Harzion; Ph.D. Thesis, Tel-Aviv University(1983).

5. J. Rishpon and S. Gottesfeld; J. of Electrochem. Soc; Submitted.

6. 0. Von Ross; J. Appl. Phys. 52_(1981) 5833.

7. S.C. Jain and V.C. Ray; J. Appl. Phys. 54Q983) 2079.

8. A.R. Moore; RCA Rev. 40_(1980) 549.

9. J.E. Mahan and D.L. Barnes; Solid State Electron. 24_(1981) 989.

10. S.K. Agarwal, Harsh and S.C. Jain, P. De Pauw, R. Mertens and R. Van Overstraeten; Proc. of the 16th IEEE Photovolt. Spec. Conf; San Diego, Calif;(1982) 366.

11. C.T. Ho, R.O. Bell and F.V. Wald; Appl. Phys. Lett. 31(1977) 463. 347

PHOTOELECTROCHEMICAL CHARACTERIZATION OF •J

Geula Dagan, Gary Hodes- Saburo Endo,* and David Cahen

The Weizmann Institute of Science, Rehovot 76100, Israel *Dept. of Electr. Engin., Science Univ. of Tokyo, Tokyo 162, Japan

Ternary Cu-ln-chalcogenides are of considerable current interest as semiconductor materials because of the impressive performance of photovoltaic devices containing CuInSe2 (1). As part of a program to investigate the series (Cu,X) (In2X3)j_x [X=S,Se; 0

RESULTS The material we studied was invariably n-type, both as single crystal and as slurry-painted film. (This latter type of electrode is not further considered here.) Doping: Some photovoltaic response was obtained after heating the initially highly resistive crystals in a stream of H2 at 500-550°C for ca. 7 hrs, or in low vacuum at ca. 200°C. Optimal results were found for crystals that were heated in evacuated sealed ampoules for ca. 2 days at ca. 400°C. In this way resistivities dropped by 4 orders of magnitude to ca. 0.25 ohm-cm. Li-doping, using LiCl, also gave reasonably photoactive material. Surface treatments: Acid(1:3 aqua regia:H2O) etch, followed by H20 aqueous KCN H2O rinses, improved the photovoltaic response of as-doped crystals. In photoelectrochemical cells, using aqueous polysulfide electrolyte, these crystals gave short circuit currents up to 4 mA/cm and photovoltages of ca. 500 mV. Significantly better results were obtained after photoelectro- chsmical etching in tenfold diluted aqua regia at 2 V (reverse) bias vs. the etching solution potential. In this way short circuit currents above 10 mA/cm were obtained, but the photovol^c,as were slightly lower (ca. 400 mV). 348

Photoelectrochemical Characterization: When aqueous ferro/ferricyanide, rather than polysulfide was used as electrolyte, improved fill factors (0.5, rather than 0.3) and photovoltages up to 450 mV (for a photoetched electrode) were obtained, but the photoanode was not stable in this electrolyte. In polysulfide, flat-band potentials from -440 to -560 mV vs. the polysulfide solution potential (ca -750 mV vs. SCE) were measured by capacitance voltage (Moct-Schottky) plots. Freshly etched crystals gave higher values (-1.5 V vs. SCE) a? did measurements of the onset of photocurrent (-1.35 V vs. SCE) (3). These values are not too different from those obtained for n-CuInS-p) (-1.5 V vs. SCE). This similarity may indicate similar, solution-induced, shifts of the band edges that could be due to specific adsorption of oligosulfides. The improved photo-I-V characteristics in ferro/ferricyanide point to kinetic limitations in polysulfide. It is possible that these are connected with such surface adsorption. Contrary to what is found for n-CuInS^ in polysulfide, the n-CuIncS system shows a slightly negative temperature dependence. This probaBl means that the (solid state) photovoltaic temperature dependence here is stronger than that of tha solution (including the ad- and desorption equilibria) . The poor long wavelength response of single crystal electrodes (whr-n compared with that of polycrystalline thin films), can indicate that the photovoltaic quality of the crystals is still far from optimal. From spectral response measurements (in aqueous sulfide) an indirect allowed transition of 1.35 eV is derived. No clear evidence for a direct allowed transition was found, contrary to solid state measurements on similar crystals (5). The output stability of the CuIncSg/polysulfide system (at maximum power) was found to be inferior to that of the analogous CuInS^ system. After an initial two-fold increase, until some 5 KC/cm photocnar . was passed, a gradual decrease set in. In the best case this decrease was so slow that even after 17 KC/cm photocharge passed the final output was still above the initial one. After passage of this amount of photocurrent the photo-I-V showed a decrease in open circuit voltage but an increase in short-circuit current. The (poor) fill factor hardly changed. It is therefore reasonable to assume that the initial power increase is due to decreased recombination losses, possibly due to improved transfer kinetics across the semiconductor/electrolyte interface. This can be brought about by "auto-photoetching". Surface analyses show that the near-surface becomes Cu-depleted (as is observed for the analogous CuInSe2 system). This is especially pronounced after photoelectrochemical etching. Then the surface seems to become mainly indium sulfide. After a stability test sulfur/oxygen exchange is seen very near to the surface, indicating the presence of some indium oxide. The above-mentioned results from capacitance-voltage measurements also indicate the occurence of some surface changes, in this case, very shortly after imnersion in aqueous sulfide. These changes are also evident in the large increase in spectral response, during use in this electrolyte, and in the decreased hysteresis of the capacitance-voltage plots, after some use in this electrolyte. The very high "donor-density", as calculated from all these plots, indicate a near-degenerate surface. This could bo due to a highly doped oxide or oxysulfide top layer. 349

It is quite likely that, in contrast to CuInX-, the already Cu-poor spinel cannot tolerate further decrease of its Cu-content, without deleterious consequences. The decreased open-circuit voltage indicates that the potential barrier near the interface decreases* The fact that the presence of indium oxide does not prevent the decrease in output stability, suggests that the Cu-content of the material is to blame. It is thus likely that the Cu-d-orbital participation in the valence-band of the dichalcogenides is the factor that is mainly responsible for their excellent output stability, although it should be borne in mind that the oxide film in the spinel is considerably thinner than on the disulfide. Thus, while the present results make this a plausible hypothesis, more experiments are needed to prove it. REFERENCES 1. R,A. Mickelsen and W.S. Chen, Proc. 16th IEEE Photovolt. Spec. Conf. (IEEE NY 1983) pp. 781-785. 2. J.J.M. Binsma, "Crystal Growth and Defect Chemistry of CuInS2", Ph.D. thesis, Nymegen, The Netherlands (1981). 3. Y. Mirovsky, Ph.D. thesis, Feinberg Grad. School, The Weizmann Institute of Science, Rehovot, Israel (1983); Chs. 4,5; Y. Mirovsky and D. Cahen, Appl. Phys. Lett. 40, 727 (1982). 4. J.C.W. Folmer et al. J. Electrochem. Soc. 130, 442 (1983), RNP511 and to be published. 5. A. Usujima, S. Takeuchi, S. Endo, T. Irie, Jpn. J. Appl. Phys. 20, L505 (1981).

ACKNOWLE DGEMENTS We thank the German Federal Ministry for Research and Technology (BMFT) for financial support through the Nuclear Research Centre KFA, JUlich and the Israel National Council for Research and Development, and John A. Turner (Solar Energy Research Institute, Golden, CO, USA) for capacitance-voltage measurements and their interpretation. 350

MNi Fe - HYDRIDE COMPACTS FOR HYDROGEN HEAT PUMPS 4.15 0.85

Y. Josephy, Y. Eisenberg and M. Ron Department of Materials Engineering Technion, Israel Institute of Technology Haifa, Israel 32000

Hydrogen heat pumps (h.h.p.)» an alternative energy conversion means or even energy source, are an important and developing part of the future hydrogen technology.

The main characteristics required from hydrides suitable for h.h.p. are large hydrogen flow rates and concomitant thermal effects. The thermal effects are produced by the enthalpy of the hydrogen absorption/ desorption reaction.

As a result of the hydriding reaction, the thermal conductivity of most useful hydrides decreases by orders of magnitude. Rechargeable hydrides are known to undergo comminution upon repeated cycling, ultimately turning into micron-sized particles. The fine particle size of the powder facilitates the chemical reaction rate by providing a large surface area, but it further decreases the thermal conductivity of the hydride. The resulting powder mass has a thermal conductivity of the order of 1 w/m°C or less, which, in the materials meant to be utilized in h.'i.p., does not provide a sufficient heat transfer rate. In particular, the point of lowest temperature of the h.h.p. represents a bottle-neck for the heat transfer.

Presented here are hydride compacts that have a high thermal conductivity and provide large heat transfer rates. Such compacts, called p.m.h. (abbreviation of porous metal-matrix hydrides) are being developed in our laboratory. P.m.h. compacts have a complex microstructure consisting of a porous metallic matrix in which fine hydride particles are embedded. The metallic matrix material does not react with hydrogen but serves to conduct heat to and from the fine hydride particles (1,2,3) ., The microstructure of a p.m.h. compact is seen in fig.l.

Fast hydriding/dehydriding kinetics are a necessary condition for a hydride suitable to be used in a hydrogen heat pump. The kinetics of the RNi5 compounds (where R is a rare earth), are known to be relatively fast (4). In our experiments (unpublished), we found that the compound MNi4.i5FeQ_85 is also endowed with relatively fast kinetics. The commercially available hydride MNi4.15Feo.85Hx was converted into p.m.h. compact by means of an aluminum matrix. A specially developed sintering process, that makes the p.m.h. compacts stable during repeated cycling, has been described elsewhere (5,6.7).

The thermal yield of the p.m.h. compacts was evaluated under conditions similar, as far as possible, to the ones prevailing in the heat

Research sponsored by U.S.A. Israel Binational Science Foundation, Jerusalem, Israel. 351

Fig.l. Optical micrograph of a p.m.h.-compact of MNi4#i5Feg.85 with 18 wt/o aluminum; Magnification: 1050. exchanger of a hydrogen heat pump. A tube simulating a modular element of a heat exchanger was filled with the above compacts and incorporated into a high pressure measuring system. A jacket, through which water flows as heat transfer medium, was formed by a concentric tube. The following parameters were measured.

• hydrogen flow and hydrogen flow integral • hydrogen pressure within the modular element and within the hydrogen line • temperature of inlet and outlet water as well as two temperatures at a known radial distance of the hydride compacts.

Data were collected, recorded and processed by a microcomputer-based system. Plots of the simultaneously recorded parameters in the processes of hydrogen absorption and desorption are shown in Fig.2a, b and Fig.3a, b.

Fig.2a and b show the instantaneous and the integral hydrogen flows and the hydrogen pressures. The hydrogen flow is transient in character. Several values for the integral hydrogen flow are shown in Table 1, for 60, 120 and 180 seconds, for both absorption and desorptxon. 352

90 so , me mss. IHfflESS. BO IHIEOn. F19I

40 70

60 30 50

a: 30

20 10

a 0 120 ISO 240 300 560 420 340 600 660 TIME. SECONDS

Fig.2. Hydrogen flow and pressure plotted vs. time, for a modular element filled with MNi^#isFeg_35HX p.m.h.- compact with a 18 wt/o aluminum matrix; the water flow rate was 2 lit/min. a. during the process of hydrogen absorption b. during the process of hydrogen desorption

Table 1.

Fraction of hydrogen sorbed Process 60 sec 120 sec 180 sec Remarks

Absorption 0.86 0.97 0.99 Fig. 2a

Desorption 0.57 0.86 0.96 Fig. 2b

The mean water temperature excursion upon absorption and desorption did not exceed 2°C, at a water flow rate of 2 lit/min. During absorption, the temperature of the hydride increases by 17°C within 15 seconds, while during desorption the temperature drops by 6.5°C, also within 15 seconds. 353

SO j Q] DEEP TDf* n>1 ^ MIER IK IDT nti Q wiot ajT lor nti

a Ic r

TIME. SECONDS

Fig.3. Two temperatures, designated shallow and deep at a constant radial distance of the MIH.4misFeo.85HX p.m.h.- compact with 18 wt/o aluminum matrix, plotted vs. time. Water inlet and outlet temperatures vs. time. The water flow rate was 2 lit/min.

The temperature of the hydride as a function of time during the desorption process was used for the estimation of the thermal conductivity (see Fig.3b). At the point of the lowest temperature of 10°C, where dT/dt - 0, quasi steady-state conditions were assumed. A heat transfer equation including an internal heat generation term, qv, was implied (2). The term qv was derived from the measured hydrogen flow and the temperature difference, AT, from the measured temperatures at a known radial distance. As a result, a value of k = 21 W/m°C was obtained.

The fractions of desorbed hydrogen vs. time, for three types of MNi4#i5FeQ#85^x hydrides are shown in Fig.4. The three types of H hydrides were: MNi4#i5Fe0#85 x hydride in powder form, p.m.h. with a 27 wt/o aluminum matrix and p.m.h. with a 18 wt/o aluminum matrix. The p.m.h. compacts with a 18 wt/o aluminum matrix show the greatest fraction of desorbed hydrogen after given times of desorption. For these particular p.m.h. compacts, the thermal yields per kg of hydride were calculated and marked on the upper curve of Fig.4. The thermal yields were calculated using the hydrogen flow, multiplied by the enthalpy of the reaction. The enthalpy of MNi/, p^5Feo. 85 hydrogenation is known to be AH 'v. 1050 joules/liter H2 (2) . The maximum uptake was taken as 1.2 wt/o of hydrogen. The densities, p, and the calculated fractions 354

MNJ4B F»a85 H^

PMH-COMPACT WITH ISol/o ALUMINUM

PMH-CCMPACT WITH 27.1/0 ALUMINUM

• —POWDER

TIME(MIN)

Fig.4. Fraction of desorbed hydrogen (integral flow) vs. time for MNi^j^FeQ 85^x ~ powder and p.m.h.-compacts with 18 wt/o and 27 wt/o of aluminum matrix. Water at room temperature (between 15 and 20cC) and a flow of 2 lit/min. of porosity of the two compacts were: 5 gr/cnP and 4.7 gr/cm3, and 0.20 and 0.12 respectively, for the 18 and the 27 wt/o aluminum matrices. The porosity includes open-interconnected as well as isolated pores. Structural studies for establishing the distinction between open and closed porosity are under way.

DISCUSSION

For a particular hydride, the thermal conductivity of p.m.h.-compacts increases with increasing percentage of aluminum matrix. Similar results were found in this work. However, the hydrogen yield and, consequently, the thermal power are higher for the 18 wt/o than for the 27 wt/o aluminum matrix compact. This is so in spite of the fact that the thermal conductivity of tha 27 wt/o aluminum compact is higher than that of the 18 wt/o aluminum one.

The rationale explaining the apparent contradiction, that increasing thermal conductivity does not cause an increase in the hydrogen yield is seen in the following.

The porosity, and in particular, the open-interconnected porosity, is an important factor in determining the permeability (or passability) of 355

hydrogen through the porous structure of the p.m.h.-compacts. The volume fraction, and the shape and size distribution of the open- interconnected pores, strongly influence the hydrogen flow kinetics.

The porosity was found to be higher for the 18 wt/o, than for the 27 wt/o aluminum compact and it may have the effect of facilitating the hydrogen flow, swamping the effect of the higher thermal conductivity of the 27 wt/o aluminum. In order to evaluate these effects quantitatively, the volume fraction, and the shape and size distributions of the open, interconnected porosity, are currently under intensive investigation.

In general, the higher hydrogen and thermal yields of the p.m.h. compacts, compared with powders are evident from the results discussed here. The beneficial effect on the hydrogen and thermal yields, of the increased thermal conductivity of the p.m.h. compacts, is capable of further improvement by better understanding of the influence of various factors on the hydrogen flow.

REFERENCES 1. Ron, M., 11th IECEC, p.954 (1976).

2. Ron, M., Navon, U. and Levitas, I., Proc. Int. Symp. on Metal Hydrogen Systems, Miami Beach, April 1981, Pergamon, Oxford, 1982, p.701.

3. a) Groll, A. and Nonemacher, A., 17th IECEC, 1982, p.1185. b) Rudman, P.S., Sandrock, G.D. and Goodell, P.D., J. Less-Common Metals, 89_, 437 (1983).

4. Goodell, P.D. and Rudman, P.S., J. Less Comm. Met. 89_, 117-125 (1983).

5. Ron, M., Gruen, D., Mendelsohn, M. and Sheft, I., J. Less-Common Metals, 74_» 445 (1980).

6. U.S. Patent No. 113873.

7. Israeli Patent application No. 66552. Patents applied for in six other European countries as well. 356

CROSSLINK DENSITY OF POLYMERS - CAN IT BE DETERMINED BY SOLVENT SWELLING?

Moshe Gottlieb

Chemical Engineering Department, Ben Gurion LI., Beer Sheva

ABSTRACT

The swelling of polymer networks is analyzed with emphasis on the main assumptions inherent to the classical development. Flory's new theory of rubber elasticity is included. Comparison of experimental data obtained by swelling the same network by two different solvents indicate large discrepancy between theory and experiment. The entire method is shown to be unreliable for exact determination of crosslink density.

INTRODUCTION

Determination of the degree of crosslinking of a polymer network by means of the amount of solvent uptaken at saturation has long been used as a standard method (1, 2), the experimental simplicity of this technique accounting for its widespread use. In order to obtain the concentration of crosslinks from the saturation polymer volume fraction three assumptions are made: 1. a model for rubber elasticity is assumed usually "phantom" (3] or affine (4); 2. the mixing of polymer and solvent is described by a thermodynamic model - customarily by the Flory-Huggins lattice model; 3. the chemical potential of the solvent is assumed to be described by the sum of two contributions, one due to elastic deformation of the network and the other due to the mixing of the two species (this is the so called Flory-Rehner [5) additive Gibbs free energy change assumption). Some doubts have been raised recently regarding the last assumption (6, 7). Also, it has long been known that both rubber elasticity theories mentioned above fail to describe experimentally observed stress-strain behavior (8, 9). Hence, it seems appropriate to determine the validity of the swelling method for the determination of crosslink density.

THEORY

According to the Flory-Rehner assumption the chemical potential of the solvent relative to its reference value is given by

(yl "Pl3 = (P1 " ^elasticity + ^1 " ^mixing W

If we assume that the Flory-Huggins lattice model is applicable

^1 ~ Vmixing = ^ - V + V2 + XV2 <2>

where V2 is the volume fraction of polymer in the solution and X is 357

Flory's interaction parameter which is a function of V2. The elastic contribution to eq. (1) is obtained from a rubber elasticity model by

where AAe£ is the free energy change due to elastic deformation of the network, ni the number of moles of solvent and N the Avogadro number. At equilibrium eq. (1) is set equal to zero and hence

N(8AAeJl/3n1) = -[lnCl - v2) + v2 + xv]] W

For phantom network (2, 8)

ph - v2) • v2 • xVjKvl/Vjll/SCl/Vj) (6)

where £ is the number of network cycles per unit volume, \\ is the molar volume of solvent, v§ is the polymer volume fraction upon network formation (crosslinking stage). The swelling equilibrium value of V2 should be used in 2q. 6. The cycle rank ? is related to the concentration of junctions p and junction functionality $ by y/£ = 2/<|> - 2. The relationship between ? and V2 for a phantom network is indicated in Fig. 1 for the case of a Polydimethyl siloxane (PDMS) network swollen in benzene. The value of x used here is taken from Brotzinan and Eidlinger (10). A similar expression may be obtained for an affine network model

Unlike the phantom network, the last equation depends on junction functionality. For a given swollen network [y-f) £af is always smaller than Cph approaching it as •+ °°. It should also be pointed out that although fph and £af depend on the solvent (vj) and the polymer (x) the ratio 5af/fph is independent of the polymer solvent system. The dependence of £af on V2 for PDMS/benzene is depicted in Fig. 1 for 3 and 4 functional networks. As already mentioned, the phantom and affine network theories are incapable of describing stress-strain behavior of real networks due to the omission of intermolecular interactions. A new theory (8, 9) based on the concept of suppression of junction fluctuation has been proposed and found capable of describing networks under different modes of deformation. The model has two parameters K and t, both of which depend on network topology only. According to theory polymer networks show a behavior which is intermediate to the phantom and affine models. Phantom network corresponds to K = 0 and affine network toK + » also, phantom network behavior is resumed at high deformations (or swelling). The cycle rank for the Flory-Erman model is given by

1 eFE = ?ph[i + (v/sW (8) 358

where K (v^) is a function that depends on K and ? (cf. ref. 9). Eq. (8) is plotted in fig. 1 for several typical parameter values (9,10). It should be stressed again that the ratio ?FE/?ph is independent of the polymer/solvent system for a given set of parameters,.

RESULTS AND DISCUSSION

In order to test the swelling method a model network of known degree of crosslinking has to be prepared. Despite the progress in the area of model network formation, structure determination is still questionable. But even if this problem is resolved the question of network elasticity model exists. It is possible to resolve both problems simultaneously by swelling the same network by two different solvents. Since the network in both cases is identical the ratio of Ejl and EjH the cycle ranks determined by solvent I and II respectively, should he equal to one. And since the ratio of 5af/5ph an(i ?FE/?ph are independent of the system used it is possible to assess the swelling method without a commitment to a specific network elasticity model. In Fig. 2 the ratio of EJ values obtained for model networks each swollen by two different solvents (11, 12) are plotted as function of EJ as computed from the crosslinking chemical reaction and stoichiometry. The solvents used for swelling are benzene, toluene, cyclohexane and PDMS oligomers. For the latter x=0 has been assumed. The data indicate a l-arge discrepancy well beyond experimental error, for the majority of systems implying the invalidity of the method for the determination of £.. It is also evident from Fig. 2 that at high degrees of crosslinking which will correspond to relatively low swelling, EJ1/?11 -»• 1. This may be ar indication that the non-additive contribution to the free energy becomes less important at low solvent concentrations. In Fig. 3 EjI/EjH for PDMS model networks swollen by PDMS oligomers of increasing molecular weight (13) are shown as function of the solvent molecular weight ratio. From this figure it is clear that large molecular weight species are less reliable in determination of degree of crosslinking resulting in a 60% error. The main conclusion from this work is that 20-40% error is expected in the determination of crosslink density by solvent swelling. Only at very high degree of crosslinking the method is of any reliability. The data presented strongly indicate the incompleteness of swelling theory at its present level of development.

REFERENCES

1. Collins, E.A., Bares, J., Billmeyer Jr., F.W., Experiments in Polymer Science, John Wiley, N.Y. 1973; Orwoll, R.A., Rubber Chom. Tech. 1977, 50, 451. 2. Flory, P.J., Principles of Polymer Chemistry, Cornell University Press, Ithaca 1953. 3. James, H.M. and Guth, E., J. Chem. Phys. 1947, 15, 669. 4. Flory, P.J., J. Chem. Phys.,1950, 18, 108. 5. Flory, P.J. and Rehner, Jr., J., J. Chem. Phys. 1943, 11, 521; ibid 1950, 18, 112. 6. Brotzman, R.W. and Eichinger, B.E., Macromolecules 1983, 16, 1131. 7. Gottlieb, M. and Gaylord, R.J., Macromolecules 1984, 17, 000. 359

8. Flory, P.J., J. Chem. Phys. 1977, 66, 5720. 9. Flory, P.J. and Erman, B., Macromolecules 1982, 15, 800. 10. Brotzman, R.W. and Eichinger, B.E., Macromolecules 1981, 14, 1445; ibid 1982, 15, 531. 11. Gottlieb, M. and Macosko, C.W., unpublished data. 12. Meyers, K.O., Bye, M.L., and Merrill, E.W., Macroraolecu]es 1980, 13, 1045. 13. Gent, A.N. and Tobias, R.H., J. Polym. Sci. Polym. Phys. Ed. 1982, 20, 2317.

Fig. 1. The cycle rank of PDMS networks swollen in benzene as predicted from the polymer volume fraction by the different theories. 360

I.U

1 Oβ

0.6 • BENZENE - TOLUENE / • BENZENE - POMS * BENZENE - CYCLOHEXflNE /

0.4

6 8 10 12 x I05 mole/cm3

Fig. 2. Cycle ranks of the same network computed from swelling results in two different solvents.

04-

Fig. 3. Cycle ranks obtained from swelling a network in a series of oligomers of increasing molecular weight. 361

PHASE SEPARATION IN RUBBER MODIFIED FLAME RETARDANT HIGH Tg EPOXY SYSTEMS

Hemi N. Nae

Department of Plastics Research The Weizmann Institute of Science, Rehovot, Israel 76100

INTRODUCTION Properties of epoxy systems are determined by their three dimensional network formed during polymerization. Most epoxy resins are based on difunctional epoxides of the diglycidyl type (DGEBA) or tetrafunctional such as tetraglycidyl ether of diamino diphenyl methane (TGDDM). High glass transition temperature, Tg, are a result of using highly crosslinked epoxy systems cured with aromatic diamines such as diamino diphenyl sulfone (DDS). In this work, a three functional resin, three glycidyl ether of tris hydroxy phenyl methane (TEN) is used. Due to its hit,'.! orosslinking density, TEN/DDS has a Tg of 334 C (1). Epoxy resins are brittle and have relatively low impact strength. Addition of small amounts of rubber to polyfunctional epoxy systems improves crack resistance and impact strength of the cured systems. This improvement has been attributed to the in-situ formation of rubbery domains of a definite size and shape during the cure cycle (2). Phase separation occurs when the rubber and epoxy become incompatible. Rubber separates well before the liquid-to- rubber (gelation) transition (3). The size and shape of the rubbery domains depend on the time and temperature of cure. Brominated polymeric additives (BPA) of the diglycidyl type were introduced to flame retard graphite-reinforced epoxy composites ( 4 ) . The objective of this work is to study the effect of tha introduction of reactive brominated additives and rubber on the cure behavior and phase separation of high Tg epoxy systems with implications to fiber reinforced composites.

EXPERIMENTAL

Materials; Triglycidyl ether of tris (hydroxy phenyl) methane, TEN (XD 7242.00L), Dow Chemical Company and TGDDM (MY 720), Ciba-Geigy were cured with DDS, Ciba-Geigy. A BPA containing 50% Br (F2001P), Makhteshim Chemical Works, was pre-reacted with a carboxyl terminated butadiene- acrylonitrile CTBN rubber (Hycar 1300x13), B.F. Goodrich Chemical Company. All formulations were stoichiometric compositions (1 mole epoxy with 1 mole amine hydrogen) dissolved in methyl ethyl ketone (1:1.5 w/v). The cure cycle 362

was: i|5 min . at 30 C/vacuu*. 8 min at 125 C, 2 hours at 177 C. Specimens were postcured at 192 C for 4 hours. Graphite cloth impregnated with the resin solution was cured according to the same cure cycle. Dynamic Mechanical Analysis: An automated torsional braid analyzer (TBA), Plastics Analysis Instruments Inc., was used to monitor modulus and logarithmic decrement at about 1 Hz under Helium flow. Thermogravimetric Analysis (TGA); DuPont 951 TGA with 1090 Thermoanalyzer was used at 20 C/min in Nitrogen atmosphere. Morphology: Zeiss Optical Microscope, Jeol JSM 35C Electron Microscope and Philips Scanning/Transmission Electron Microscope were used to study morphology. SEM micrographs were of gold coated fractured surfaces. STEM micrographs were of samples cured on copper grids. The cured samples were stained with osmium tetroxide. A laser light scattering analyzer was developed to identify the size of the rubbery domains. Ths average size of the phase separated rubber particles was calculated from the image of the scattered light using Bragg relation.

RESULTS AND DISCUSSION

Three events are observed in TBA spectra of isothermal cure of TEN/DDS or TGDDM/DDS with 19% bromine and 2% rubber (Figure 1). These are designated as the pre-gel , gelation and vitrification transitions (5). The system containing TEN/DDS gels and vitrifies sooner than the system containing TGDDM/DDS. This is due to differences in the crosslink density of the cured macromolecules. Addition of BPA and rubber inhibits gelation and vitrification somewhat but not significantly even when the amount of additives is 30? of the system. This indicates that the cure reaction depends mainly on the reactivity of the epoxy resin and the curing agent. Apparent activation energies, calculated from log time vs. 1/T are also similar to those of the neat systems. The difference in the epoxy resin leads also to differences in the glass transition temperature, Tg, of the systems. Tg of the modified TEN/DDS is 273 C and of TGDDM/DDS 220 C compared to 3311 and 220 of the neat systems respectively. A sub-zero relaxation corresponding to the Tg of the rubber is observed. After heating above 250 C the material starts to degrade and the Tg of the rubber becomes more distinct. TGA of all systems show a catastrophic weight loss at 310-350 C. Therefore, the limiting factor for structural purposes is the Tg of the material and not its thermal stability. However, prolonged exposure above 250 C may result in partial degradation. The char yield at 300 C is 35-455, indicating the flame retadrdancy of the modified systems. The content of bromine was chosen as 195 to impart flame retardancy to the modified systems. Epoxy/DDS with BPA but without rubber form an homogeneous matrix. Upon the addition of reactive rubber phase separation, which accompanies polymerization, occurs. Its 363

extent depends on the temperature of cure, the nature of the reacting monomers and the rubber concentration. The morphology is apparently arrested before gelation and the rubber fores spheres homogeneously dispersed in the epoxy matrix. The rubber is chemically bonded to the epoxy matrix but separates due to its incompatibility with the matrix. The rubbery domains are of an average ; lameter of 2-4 m with a core diameter of 1-2 m. The size and shape of the rubber spheres are the same in 2,4 and 6? rubber (Figure 2). Crack propagation is mainly unidirectional and stops at the rubber particles. Energy dissipated by this mechanism is apparently responsible to the increased toughness of such systems. In systems containing 8? rubber there are agglomerates in the form of cells of an average diameter of 20-180 m containing many rubber particles. The cells have an envelope of an average thickness of 10-20 m (Figure 3). The envelope is probably a result of phase inversion of the rubbery moiety. An optimum in the toughness of such systems about 6J rubber is probably due to the formation of cells which accelerate failure of these systems. Phase separation is much more distinct in TEN/DDS systems than in systems containing TGDDM/DDS due to the differences in crosslinking density and the nature of the epoxy resin. Laser diffraction patterns show a short range packing of the rubber particles with an average particle size of 2.6 m which correlates well with the size observed in the SEM. Composites prepared from BPA/rubber containing systems do not show similar phase separation since the fibers are closely packed, limiting space for rubber particles development. However, rubbery domains are formed and contribute to the increased impact strength of such systems (i<). The distribution of BPA can be mapped by EDAX line scanning of bromine which shows an homogeneous distribution of the BPA in the epoxy matrix. Acknowledgement: The help of Dr. Z. Nir, Makhteshim Chemical Works, is greatly appreciated.

REFERENCES

1. H.N. Nae and J.K. Gillham, ACS, Div. Org . Coat. Plast. Chem., Prep., 48, 566 (1983). 2. C.K. Riew, E.H. Rowe and A. R. Siebert, in "Toughness and brittlement of plastics", ACS, Advances in Chemistry Series No. 154, 1976, p. 326. 3. L.T. Manzione, J.K. Gillham and C.A. McPherson, J. Appl . Polym. Sci., 26, 089 (1931); Ibid, 26, 907 (1981). 4. Z. Nir, U.J. Gilwee, D.A. Kourtides and J.A. Parker, SAMPE Quarterly 14 (3) (1933). 5. J.K. Gillham, in "Developments in Polymer Characterization-3", J.V. Dawkins, Ed., Chapter 5, Applied Science Publishers, England (1932). 364

3 0 10 30 10 50 LOO TIMEISEC.I LOG TIME [SEC I

Figure 1: TBA spectra for isothermal cure of TGDDM ( ) and TEN ( 1 svstem. (a) relative rigidity (b) log decrement.

Figure 2: SEM micrograph of TEN/DDS, Figure 4: SEM micrograph of 19% Br and 6% rubber. graphite fiber reinforced composite containing TEN/DDS with 19% Br and 6% rubber.

Figure 3; Optical micrographs of (a) TI-N/DDS and (b) TGDDM/DDS with 19% Hr and 8%' rubber. 365

RECENT ADVANCES IN THE STRENGTH AND TIME DEPENDENT FAILURE PROCESS OF KEVLAR MONOFILAMENTS AND COMPOSITES

H. Daniel Wagner, S. Leigh Phoenix, P. Schwartz

Cornell University, Ithaca, N.Y. 14853

ABSTPACT

In this paper we discuss recent theoretical efforts to characterize the statistical strength and lifetime of Kevlar monofilaments and composites. We present some new experimental data on the variability of the mechanical strength of single aramid filaments. Significant spool-to-spool variability is discovered within a single production lot. The strength ~f single filaments is dependent upon both the flaw distribution and the filament diameter variability across the yarn. It is found that the Weibull shape parameter for strength is probably not a material constant for aramid filaments. INTRODUCTION Most brittle fibers used in modern composite materials exhibit large variability in tensile strength, with measured coefficients of variation (cv) typically ranging froir. 0.05 to 0.25. Kevlar is no exception and shows considerable variability in strength Hue to the presence of flaws on the surface and within the interior oF the filament. Simple composite materials, in which fibers are arranged parallel to each other and impregnated with a binding matrix, exhibit considerably less variability than the corresponding fibers (cv = 0.03 to 0.07). In creep-rupture applications, the variability in lifetime is generally much higher than that for the strength, for both single fibers and composites. Clearly, statistical considerations are important when studying the strength and fatigue of fibrous materials. In the present paper we discuss the significance of recent experimental data generated in our laboratory with simple Kevlar fibers in the light of the results from a lnicromechanical model of the statistical failure process developed recently by Tierney [1] and Phoenix and Tierney [2]. This model predicts that a simple relationship exists between the Weibull shape parameters of the Lifetime (b*) and the strength (b) distributions, namely, that b* = (l+r)b where r is a positive constant which can be determined independently using results from stress-rupture experiments. Phoenix [3] discusses the theoretical justification of the above fiber model in terms of thn kinetic failure c- idealized molecular crystals of the form found in stiff polymeric filaments such as Kevlar. In particular, the key constant r involves an approximation to the potential function for chain scission and is shown to follow r = If/kT, where T is absolute temperature, k is Boltzman constant and II has the units of activation energy. 566

EXPERIMENTAL RESULTS Five spools of Kevlar (49 and 29) aramid yarns, taken from different production lots (Table 1) were used, and 5 cm long specimens were tested in tension with an Instron Model TM machine at a rate of 2.54 cm/min, under standard conditions (21°C, 65% RH).

Table 1; Specifications for Kevlar aramid yarns

spool type linear density filaments/yarn lot (tex) A K-49 21.6 134 74048 B K-49 21.6 134 74048 C K-49 21.6 134 74048 U K-49 42.2 267 70438 K K-29 44.4 267 74043

1. For each spool studied, the filaments were carefully removed and their linear density (an indirect measurement of diameter) measured with a vibroscope, prior to being tested.

2. A Weibull analysis was performed for breaking loads and tenacities (breaking load/linear density), and the Weibull shape and scale parameters were estimated using the method of maximum likelihood (MLE). Table 2 presents the results of this analysis.

Table 2: NLE of Weibull shape and scale parameters

spool breaking load (tenacity) shape scale (N) shape scale(N/tex) A 8.8 0.397 10.4 2.27 B 4.1 0.429 10.2 2.45 C 8.1 0.466 10.4 2.44 U 7.8 0.393 9.0 2.30 K 7.0 0.384 10.5 2.50 note: 50 specimens per spool were used (gauge length = 5cm)

In all cases the shape parameter for tenacity is greater than that for failure load, that is, the cv for tenacities is smaller than that for breaking loads, an indication of the elimination of the fiber diameter as a variable. We are not aware of any previous theoretical study based on weakest link models which makes reference to a possible variation in the diameter of individual fibres within the links. Rather, a common assumption is that the fibers in the bundle are identical. A rather large diameter variability is found among the fibers within spool H (see Figure 1).

3. In order to test the spool-to-spool variability among Kevlar 49 spools a formal test of equality of the Weibull parameters, devised by Thonan and Bain [4] was used. The conclusion reached from this analysis (which is discussed in greater depth in [5]) is that the filaments in spools A and I! are significantly weaker than those from spools R and C; that there is no significant difference between spool1; R and C; and thaf there is no significant difference between spools A and U. 367

Fig.l: Filament diameter variability in Spool B (Kevlar 49)

4. In an earlier study, Runsell [6] claimed that the scatter observed in the failure loads of specimens selected along a simple filament was the same as that observed if the fibers had been sampled randomly from the bundle, showing that the scatter was due to the distribution of faults on the fiber rather than due to variations among the fibers themselves. We performed a replica of Bunsell's experiment with spools A and B and we found that only if fiber variability across a yarn is within limits is Bunsell's argument valid. In the case of spopl B, for instance, the high degree of size variability outweighs the effects of flaw distribution and we find significantly greater scatter taking place across the yarn (cv = 0.26) than along it (cv = 0.11).

5. Finally, the effect of filament gauge length was studied for each spool. The linear behavior predicted by the Weibull theory for log (strength) vs. log (length) is not observed. Moreover, a regression analysis yields shape parameters of~18, a value which is about twice the expected theoretical value based on the data obtained frora simple gauge length tests. Interestingly, Phoenix and Wu [7] obtained a Weibull shape parameter of 18.2 by backcalculating from lifetime data for single filaments using b* = (l+r)b from the micromechanical model alluded to praviously. This is remarkably similar to the present results which were obtained in a different manner. The reason for the discrepancy is unclear. Our data show stronger fibers at shorter gauge lengths, but not as strong as predicted by the Weibull model; weaker fibers are obtained at longer gauge lengths, but not as weak as would be predicted. This behavior has been observed in the past in carbon, glass and silicon carbide fibers. Since a progressive change in the slope of the log-log dependence of strength on length is observed, a change in the Weibull shape parameter is probable, therefore suggesting that this parameter is not a material constant- in aramid monofilaments. 368

REFERENCES 1. L.J. Tierney, Adv. Appl. Prob. U_ (1982), 95. 2. S.L. Phoenix, L.J. Tierney, Eng. Fract. Mech. (1984, to appear). 3. S.L. Phoenix, Proc. 9th U.S. Nat. Congress of Appl. Mech., Cornell Univ., Ithaca, N.Y. (1982). 4. D.R. Thoman, L.J. Bain, Technom. H_ (1969), 805. 5. H.D. Wagner, S.L. Phoenix, P. Schwartz, submitted paper. 6. A.R. Bunsell, J. Mater. Sci, _10_ (1975), 1300. 7« S.L. Phoenix, E.M. Wu, UCRL-53365, LLNL, Univ. California (1983). 369

DEFORMATION PROCESSES IN IMPACT MODIFIED PVC

A. Hadas and A. Siegmann Department of Materials Engineering Technion-Israel Institute of Technology

The impact strength of PVC is usually enhanced by the addition of non compatible rubbery polymers, which form a second phase distributed in the PVC matrix. The dependence of impact strength on blends composition is not linear but has an S shape. As can be seen in Fig.l, there exists a critical rubber content (7-15 phr) at which the impact strength abruptly increases and then levels off. The effectiveness of the rubber depends mainly on its own mechanical behavior, compatibility with PVC and phase morphology. The mechanical behavior of PVC and modifier are schematically described in Fig.2. The much lower elastic modulus and much higher elongation at break of the rubber should be noticed. The relationship between the impact behavior, the blend structure and the nature of the energy absorbing processes are the subject of the present work. Low tensile strain rate behavior of PVC blends are reported. A comparison between the effects of a rubbery modifier and a plasticizer forming two and one phase systems, respectively, is made. PVC blends with two different mofifiers (EVA and acrylic rubber) and a plasticizer (DOP) were prepared by dry mixing followed by roll milling at elevated temperature. The blankets were molded to form sheets (0.7 mm in thickness) in a preheated press at 190°C foi J min. Tensile specimens were cut and tested using an Instron machine a. strain rates of 50 and 130%/min. The microdeformation processes were followed using optical microscopy. Representative stress strain curves of PVC and several of its blends with DOP are depicted in Fig.3. The major effects are seen in the yield zone and the elongation at break. The elastic modulus and yield stress (Figs.4 and 5) gradually decrease with increasing rubber content whereas first increasing followed by a steep decrease with increasing

MODIFIER u I

CONTENT STRAIN Fig.l: Impact St. as a function of Fig.2: Schematic α-e modifier content (schematic). curves. 370

PVC

Ql 0.2 0.3 0.4 0.5 0.6 0.7 OB 0.9 ID Q

Fig.3: Stress-Strain Curves at PVC/DOP blends. (E=130%). plasticizer content. Increasing modifier content results in a more ductile behavior, a more uniform deformation changing gradually from localized deformation (necking) into a uniform one. Simultaneously, the instability intensity at the yield zone is depressed namely, the difference between yield and flow stresses declines. The microdeformation mechanism in the various systems was studied via optical microscopy. In stretched PVC, adjacent to the neck, crazes were formed, the density of which decreases with distance from the neck (Fig.6). Crazing is known to be the deformation mechanism in brittle polymers. Upon further deformation the neck continues to propagate and the crazes widen (Fig.7) and some even develop into cracks - sites of failure initiation. Upon the incorporation of DOP the craze size significantly decreases. At high DOP content shear bands are formed (Fig.8) whereas at intermediate loadings both mechanisms are active and their interaction is also observed (Fig.9). Upon the incorporation of impact modifiers, in addtion to the macroscopic changes during deformation, as mentioned above, the microdeformation mechanism is predominantly shear banding as seen in Fig.8. It should be mentioned that the deformation mechanism is affected not only by the blend composition but also by the deformation rate.

In summary, impact modifiers impart large changes also in the low rate tensile behavior of PVC. Differently from impact behavior, the tensile behavior does not exhibit any abrupt changes at any modifier loading. With increasing modifier content the systems become more ductile, the deformation is delocalized and the deformation mechanism changes from crazing to shear banding. 371

150" DOP M E

111 100

5 10 15 20 CONTENT (phr)

Fig.4. Elastic Modulus of PVC/Modifier Blends (e = 1.3 min"1),

5 10 15 CONTENT (phr) Fig.5. Yield Stress of PVC/Modifier Blends (e = 1.3 min"1).

Fig.6. Craze System developed in Fig.7. Craze widening in strained Strained Rigid PVC. Rigid PVC. 372

Fig.8. Shear Banding in deformed PVC Containing 10 Fhr EVA.

Fig.9. Shear Bands arresting crazes developed in deformed PVC containing 5 Phr DOP. 373

FATIGUE CRACK PROPAGATION MECHANISMS IN POLYMERS

A. Bussiba, Y. Katz and H. Mathias

Nuclear Research Centre Negev, P.O.Box 9001, Beer-Sheva, Israel.

INTRODUCTION In recent years more research attention has been given to fatigue crack propagation rate (FCPR) behaviour in polymeric materials (1,2). In fact, paralleling similar studies in metallic systems, fracture mechanics concepts have been adopted to polymers by examining the relationship between the FCPR and the applied cyclic amplitude controlled by the stress intensity factor range AK (3). In polymers too, internal and external variables influence the fatigue behaviour. In addition, there are specific factors which are related to the time dependent nature of polymers, their unique structural changes and fracture modes. For example, particular emphasis has been given to the proper description of the craze-plastic zone formation at the vicinity of the sharp crack tip (4) . The issue of equilibrium crack - craze - plastic zone in glassy polymers seems to be very important to the static and cyclic loading behaviour, with and without environmental effects. As known, polymers are well recognized and used as structural materials. Thus, fatigue studies aimed to explore fatigue resistance, are necessary in order to improve fatigue life evaluations and to refine the available guideness for material selection. The current study investigates two kinds of brittle polymers, with the intention to meet several objectives; (i) To determine the fatigue crack propagation curve from the near threshold values, AKtn, up to the critical fatigue upper value, Kjfc, including frequency influences, (ii) To gather experimental results in order to enable comparative analyses, including BCC and HCP metallic alloys below the ductile - brittle transition temperature, (iii) To examine the polymer fatigue macro- and micromechanisms. Here, the transparency of selected polymers is well considered, clearly beneficial for better observations of the crack front during initiation and propagation stages. Actually, the problem of fracture mechanics assessment has been addresed in earlier studies (5), but seems relevant and attractive regarding several other fundamental problems. For example, crack front profiles (6) or lead interaction phenomena in fatigue crack extension (7) have been studied by realizing the unique advantages offered by polymeric systems.

EXPERIMENTAL PROCEDURE Two typical brittle polymers, polymethyl-methacrylate (PMMA) and polysty- rene (PS) were selected. Standard mechanical testing was performed by using uniform tensile specimens. In addition, fracture toughness parameters were determined on pre-cracked single edge notched (SEN) specimens from room temperature up to Rn°C. REN specimens of IS mm in thickness were 374

also utilized for cyclic loading tests. All initially machined notches were extended by a razor-blade sharp cut. Fatigue tests were performed by an electro-hydraulic closed loop system with an amplitude-controller device. Cyclic load amplitudes were controlled by the stress intensity range of sinusoidal waveform at frequencies of 15 and 80 Hz. Tests were carried out at room temperature with a load ratio R = P . /P - 0. Fatigue crack growth extension was optically tracked, and the FCPR curves da/dN vs. AK were established from the near threshold, AK-j.n, up to the critical upper bound, Kj£c. The cyclic AK was calculated according to the following relationship:

K = Y- ^ (1) BW Where AP is the range of the cyclic applied load, a is the crack length, B is the specimen thicknes, W is the specimen width and Y is the geometric correction factor given by: Y = 1.99 - 0.41(a/w) + 18.7fa/w)2 - 38.48(a/w)~ + 53.85(a/w)4 (2) Crack front and fracture modes were observed by light and scanning electron microscopy, with particular attention to microcrack extension rates, band width and mode transitions along the different stages of the FCPR curve.

EXPERIMENTAL RESULTS AND DISCUSSION Table I summarizes static and dynamic properties for the tested materials. The fatigue properties were obtained from the relevant experimental FCPR curves for the two applied frequencies. As expected, both polymers behaved in a brittle manner and indicated significant sensitivity to the temperature and to the strain rate.

Table I : Mechanical and dynamics properties of the tested polymers.

e K of Ic AKth Klfc o 2 Polymer MPa 'o MPa»mk Mpa-m MPa-mk 296K 353K 15Hz 80Hz 15Hz 80Hz

PMMA 48 0.4 1.5 1.2 0.38 0,.58 1.4 1.2 PS 35 0.3 2.0 1.4 0.5 0,.62 1.8 1.3

Fig. 1 illustrates the FCPR curve for the PMMA. This typical behaviour was also obtained for the PS, with similar tendencies regarding the frequency effects on the FCPR, on AK^ and on Kjfc. Moreover, the sigmoidal shape of the da/dN vs. AK curve, as known for metallic systems (8), was actually preserved. In addition, the whole FCPR curve approached towards a discrete critical value as reflected by the typical narrow range between AK-j.^ and Kj£c. This fact is not surprizing, since the brittle behaviour of the PMMA and the PS resulted in FCPR curves highly dependent on the applied AK, as obtained in the case of metallic alloys, below the ductile brittle transition temperature (8). The role of the frequency confirms the results by Skibo et al. (9) and Jilken et al. (10), namely, increasing the cyclic frequency causes a significant decrease of the FCPR in both tested nolvmers. 375

Figure 2 shows the dependency of the fatigue band size and the number of cycles per band, N*, on AK for the PMMA and PS at 80 Hz. Figure 3a demonstrates the initiation stage in PMMA as obtained by fracture surface observations. As shown, this stage is associated by extremely low energy fracture surface. Figure 3b shows microcracking and microcavities at the earlier stage II. Figure 3c illustrates fatigue striations at stage II,while later, at stage III,secondary cracking occurred which enhances the fatigue crack extension rate. Probably the most significant result which differ completely from metallic fatigue crack propagation behaviour, is connected to the micromechanisms in terms of the discontinuous crack extension. This issue is emphasized by N and has been addressed earlier by Hertzberg et al (11). As shown, the band size or striations in polymers are nor associated with incremental crack extension caused by one cycle only. The latter is well recognized in metallic materials with an excellent fit between the macroscopic crack extension rate and the microscopic FCPR based on striation spacings (12). Hertzberg et al (11) have attempted to explain this unique discontinuous growth bands in terms of a more general phenomenon, namely, craze formation in polymers which affects also the fatigue crack extension processes. However, it seems that more has to be done in order to extend and to refine the mentioned idea. Firstly, the exact description of the craze profile is needed in more realistic terms of a crack-craze-plastic zone configuration (4). Secondly, the craze, as such, is associated with a threshold value which should probably be incorporated in a proposed model.

As concluded by Israel et al.(13), the Dugdale plastic zone model was not fully adequate for describing craze geometries in PMMA. Therefore, a modification of the Dugdale - Rarenblatt model has been proposed(13),describing the crack tip configuration in terms of crack-craze-plastic zone, and expressed at equilibrium by the following relationship:

•nc - 2(0 - a ) cos" b/a - 2a cos" c/a = 0 (3)

where a is the applied stress, CTC and OyS are the stress within the craze and the yield strength within the plastic zone, respectively. c,b,a are the crack-craze-plastic zone geometrical values.

Actually, this model at equilibrium describes the elastic stress intensity term by means of the fictitious crack length, which includes the cohesive plastic zone contribution and the semi cohesive term associated with the localized micro-cracked and void-filled craze region. Thus, at equilibrium the craze size is related to the applied stress intensity factor, and a monotonic increasing function is obtained for the craze size dependency on the applied AK. Furthermore, this modified description indicates the role of the crack-craze-plastic zone profile on the effective localized stress intensity factor. In fact, the effective stress intensity factor is affected by craze formation, assuming the retractive stresses exerted by the craze at the crack tip vicinity. Referring back to the current fatigue findings in PMMA and PS, clearly, the micromorphology in the tested polymers indicates the role of the craze formation. On one hand, dynamic craze-plastic zone causes a reduction in the nominal stress intensity. On the other hand, subsequent incremental crack extension might be activated only after an additional cumulative damage, mainly applied to the microfibrils, and finally resulting in a discontinuous band extension. In fact, the band size increases with higher AK. Higher values of AK result in higher values of the elastic term, that must be offset by the corresponding increase of the dynamic craze-plastic zone size. This tendency was readily observed in the fatigue band size, as shown in Fig. 2. In contrast to the band size, N* decreased with the increase of AK, which can be explained by realizing the effective subsequent damaging potential, available for further crack extension for relatively high AK values. Consequently, the increase of the cyclic range, AK, increases the dynamic cra?.e~plastic zone size, but not the total fatigue resistance capacity, which ends up with the enhancement of the fatigue crack extension rates.

ACKNOWLEDGEMENT The authors wish to express their appreciation to Mr. M. Kupiec and Mr. Y. Fachima for experimental assistance.

REFERENCES 1. J.C. Radon, Int. Journ. of Fracture, Ij3, 1980, 533-552. 2. J.A. Sauer and G.C. Richardson, Int. Journ. of Fracture, j^, 1980, 499. 3. S. Arad, J.C. Radon and L.E. Culver, Int. Cong, on Fracture, ICF3, Munich 1973, Part VII, paper VI - 323. 4. R.W. Hertzberg and J.A. Manson, in Materials Experimentation and Design in Fatigue, F. Sherratt, J.B. Sturgeon and R.A.F. Farnborough Eds., Westbury House, pp. 185-198, 1981. 5. Y. Katz, P.L. Key and E.R. Parker, Trans. ASME, £0_, Ser. D, 1968, 622. 6. Y. Katz, A. Bussiba and H. Mathias, in Mechanical Behaviour of Materials, ICM3, K.I. Miller and R.F. Smith Eds. Pergamon Press, 1979, 383-391. 7. Y.W. Mai, Int. Journ. of Fracture, 15_, 1979, 103-106. 8. A. Bussiba, H. Mathias and Y, Katz, Proc. Materials Engineering Conf. Haifa, Israel 1981, 123-125, I. Minckof Ed=, Freund Publ. House. 9. M.D. Skibo, R.W. Hertzberg and J.A. Manson, Proc. Int. Conf. on Fracture, ICF4, D.M.R. Taplin Ed., University of Waterloo Press, 1977, 1127-1133. L0. L. Jilken and G.G. Gustafson, in Fatigue Thresholds, J. Backlund, A.F. Blom and C.J. Beevers Eds., Chamelion Press V.2, 1981, 715-733. LI. CM. Rimnac, R.W. Hertzberg and J.A. Manson, in Fractography and Materials Science, L.N. Gilbertson and R.D. Zipp Eds. ASIM STP 733, 1981, 291-313. 12. Y. Katz, A. Bussiba and H. Mathias, Proc. European Conf. on Fracture and the Role of Microstructures, ECF4, K.L. Maurer and F.E. Matzer Eds., EMAS Press, 1982, 503-511. 13. S.I. Israel, C.S. Kantamneni and W.W. Gerberich, in Mechanical Behaviour of Metals, ICM3, K.J. Miller and R.F. Smith Eds., Pergamon Press, 1979, 393-402. da/dN (mm/cycle) _^

Oq P I-1 ro n o 1 a> u> n> o o 13 3= 378

AUTHORS' INDEX

Aboelfotoh M. 291 Dirnfeld S.F. 39 Addess S. 146 Edmonds D.V. 108 Adler L. 185 Eisenberg M. 299 Admon U. 82 Eisenberg Y. 350 Agronov D. 119 Elkabir G. 232 137 Eliezer D. 152 Arigur P. 128 158 Avni R. 262 167 267 Endo S. 347 272 Fainaro 1. 315 283 Falkenstein E, 315 Azoulay M. 332 Finberg I. 283 Bamberger M. 35 Fink J.L. 67 209 Fishman M. 95 Bar-Ziv S. 225 Freund A. 119 Bornstein A. 328 137 Botstein 0. 71 Gafri 0. 258 Brandon D.G. 55 Gal-Or L. 95 Brat T. 299 Gilad I. 163 Breyer N, 133 Goldstein J. 63 Buckley D. 283 77 Burstein Z. 332 Gottesfeld S. 343 Bussiba A. 373 Gottlieb M. 356 Cahen D. 347 Grill A. 249 Carmi U. 267 262 272 285 Chaim R. 51 Grinbaum Y. 253 55 Grossman E. 249 Croitori N. 328 262 343 Grunbaum E. 82 Cyterman C. 299 Grunze M. 235 Dagan G. 347 Grushko B. 71 Dariel M.P. 30 88 82 Hadas A. 369 Dayan D. 100 Hall D. 245 379

Harzion Z. 343 Minkoff I. 35 Herrmann B. 146 Minkovitz E. 158 Hershitz R. 21 Moscovits M. 245 Heuer A.H. 51 Munitz A. 46 Hodes G. 347 Nae H. 361 Huang H.C.W. 175 Nir. A. K2 Ignatiev A. 74 Nissenholz Z 142 336 Notis M.R. 63 Inspektor A. 267 77 272 Nowick A.S. 11 Iram A. 95 Pelleg J. 67 Jacobsen K.M. 221 92 Josephy Y. 350 100 Katz Y. 163 104 373 320 Kazinetz M. 258 Phoenix S.L. 365 Kimmel G. 59 Polak M. 249 82 253 100 Prinz B. 209 104 Rabin B. 258 Kohn G. 146 Rabinovitz E 146 Kornblit 74 272 Lahav A. 299 Raveh A. 267 Laks C. 92 272 Lalman J. 245 Rishpon J. 343 Landau A. 59 Ron M. 350 Lang S.B. 324 Rosen A. 119 Levin L. 201 137 Livne Z. 46 225 Livni T. 225 232 Lodder G.C. 82 Rozenak P. 152 Manory R. 262 Rosenthal Y. 167 Marcu V. 304 Rotel M. 175 Mark-Markowitch M. 167 Rotem 225 Mathias H. 163 Rubinstein I 214 373 Ruhle,M. 51 McBreen P.H. 245 380

Sariel J. 59 104 Schwartz P. 365 Seidman D.N. 21 Sharon A. 216 Shikmanter L. 30 Siegman A. 369 Spalvins T. 283 Stern A. 146 167 201 Talianker M. 30 59 Tamir S. 277 Tarby S.K. 77 Tenne R. 304 Totta P.A. 175 Tu K.N. 288 291 Turnbull D. 1 Zafrir M. 343 Zahavi J. 175 277 308 Zevin L. 92 258 Zuta Y. 39 Wagner H.D. 365 Weinberg F. 128 Weiss B.Z. 71 88 Williams D.B. 63