摘要:

Epitaxy of Nickel Electrodeposits on a Copper (1 10) Face, from a Sulphamate Bath, in Relation to Rate of Deposition, Deposit Thickness, Degree of Stirring, and Bath Temperature BY S. IS. VERMA AND H. WILMAN" Applied Physics and Chemistry of Surfaces Laboratory, Department of Chemical Engineering and Chemical Technology, Imperial College of Science and Technology, London, S. W.7 Received 5th June 1972 Grazing-incidence electron diffraction at 50-60 kV shows the surface structure of smooth electro- polished Cu (110) faces and of the epitaxial Ni deposits on them from the bath : The Ni deposits were all f.c.c. The effects of c.d. (up to 600 mA/cm2), thickness (100-60 0008,), temperature (50 and 20"C), and vigorous stirring, are shown in detail. The effects of codeposition of Ni(OH)2 at high c.d.are also shown. Special points of interest are : (1) at 50°C although (1 11 1 twinning of the " parallel " epitaxial Ni was strong near the interface (100-200A Ni deposits) at up to -150 mA/cm2, it was weaker at 200 and absent at 300-600 mA/cm2 (at least up to thousands of 8, thickness), and this is similar to our results for Ni on (100) and (111) Cu; (2) at up to -150mA/cm2 at 50°C with unstirred bath the strongly twinned Ni at up to lo00 8, thickness showed no " directed disorientation " though accom- panied by an extremely small trace of widely disoriented (random?) Ni, possibly associated with the epitaxial misfit and contacting crystal nuclei ; (3) at 10 OOO 8, or more, at up to -150 mA/cm2 at 50°C with unstirred bath, additional weak arcs showed a smaIl amount of (100) Ni, and at 30 000- 60 OOO8, these arcs became strong and showed a (100) Ni orientation but limited to azimuths associa- ted with those of the (111) twins of the " parallel " Ni, arising from the twins by a " directed dis- orientation " (a range of rotation round the horizontal [lTO] Ni direction, presumably due to contacts between the twins, as they grow, with misfit stresses and strains at their interface), followed by cube face development leading to preferred growth of crystals in { 100) orientation.Ni(S03*NH2)2 -4Hz0 350 gll., NiClz -6H20 5 gll., H3B03 35 g/l. In epitaxial crystal growth on an atomically highly smooth substrate, twinning in the initial very thin deposits is usually attributed to the " pseudomorphic " tendency and the associated stresses arising from the misfit between the substrate and the deposit at their interface.In addition to twinning, disorientation effects are also sometimes observed, evidently arising from the same causes, and presumably involving dislocation movements and slip processes. A " directed disorientation ", consisting of a range of rotational displacements about a particular densely-populated lattice row of the interface plane, away from the ideal epitaxial orientation, was observed by Evans and Wilman first in the case of ZnO formed as a reaction product on a zinc- blende (110) cleavage face heated in air at 520°C. No reaction occurred at 500°C; thus at 520°C the reaction must have been less rapid than in several previous studies at about 650"C, where also recrystallization must have been occurring and promoting development of the single orientation then observed.Such " directed disorienta- tions" have also been observed in this laboratory in epitaxial electrodeposits on highly smooth single-crystal faces, for Zn and Cd on Cu Cu on Ag(100) and (1 and in Fe and Cr on Cu(lll), (100) and (110).4 Kumar's observation that 185186 EPITAXY OF NICKEL ON A COPPER (110) FACE such a disorientation occurred in Cu on Ag, but not in Ag on Cu, appears to corres- pond to the Cu deposit with axial length 3.61 A being in tension and thus less stable when growing on Ag (a = 4.08 A), whereas for Ag on Cu the Ag is under compres- sion and thus more stable when constrained to fit on to the smaller Cu lattice. In the (21 1)-oriented Fe and Cr deposits on Cu (1 lo), Reddy and Wilman observed a " one-sided directed disorientation ", and various angles of rotational displacement from the ideal (211) orientation, as well as extensive continuous ranges of rotation.In the present system of Ni on Cu (1 lo), no directed disorientation was found in the initial thin epitaxial (parallel) Ni, but one occurs in the upper regions of thick layers ( N 10 000 A and higher), which is associated with the octahedral twins of the " parallel " Ni and leads to the development of a moderately high proportion of an azimuthally limited { 1001 Ni orientation. The present study provides an interesting comparison with our results for Ni on Cu (100) and Cu (111) (to be published), and on mechanically polished polycrystalline copper,6 from the same sulphamate bath.This bath is known to give relatively stress-free ductile nickel deposits, and our electron diffraction observations showed that the nickel is uniformly f.c.c. in structure, free from any close-packed hexagonal nickel such as is partly obtained when the usual sulphate bath with a higher propor- tion of NiC12 is used. This makes the consideration of the epitaxy correspondingly simpler. The results are obtained up to the unusually high current density of 600 mA/cm2, corresponding to N 1500 A/s rate of deposition. EXPERIMENTAL The experimental details were as described by Verma and Wilman.5* The electrolyte consisted of nickel sulphamate, Ni(S03NH& -4Hz0 350 g/l. ; NiClz -6H20 5 g/l. ; H3B30 35 g/l.; it was supplied by British Insulated Callender's Cables Ltd., who obtained it from Messrs. Albright & Wilson Ltd. The electrodeposition was carried out in 60ml of this electrolyte in a Pyrex beaker of 5 cm diam., surrounded by a water jacket at the desired bath temperature. The freshly electropolished copper crystal was mounted with its (1 10) face 3 mm below the surface of the electrolyte and 3 cm above the horizontal nickel-sheet anode. Before deposition from the bath at 50"C, the crystal was immersed for 5 min in some of the solution at 50°C so as to attain this temperature. After nickel was deposited at a desired c.d. for the desired time and thus thickness, the crystal was removed without switching off the current ; it was washed immediately with distilled water followed by acetone to remove the water, and then covered with isopropyl alcohol to minimize oxidation during transfer to the Finch electron diffraction camera, which was then evacuated.Electron diffraction photographs were obtained with the -55 kV electron beam at low grazing incidence and with a specimen-to-plate distance of about 48 cm. The deposit thickness was estimated from Faraday's law, allowing for the current effi- ciency of the deposition, which was determined by a series of experiments on a larger sheet- copper cathode, by weighing the deposit on an Oertling semi-microbalance. The current efficiency at 50°C was virtually 100 % up to about 300 mA/cm2 when there was no stirring of the bath (and up to 360 mA/cmZ when vigorous stirring was used), falling to about 75 % at 600 mA/cm2.6 Nickel deposits were removed by boiling in orthophosphoric acid, and the copper (1 10) face was then re-electropolished and used again as substrate.RESULTS The electron diffraction patterns from the electropolished copper (1 10) face were similar in type to fig. 8 and 13, showing the surface to be highly smooth and flat, on the atomic scale of dimensions (cf. Verma and Wilman 5 ) . The nickel electro- deposits were in all cases f.c.c. Above a certain c.d. and thickness range, some nearly amorphous greenish nickel hydroxide was co-deposited (see fig. 1, 14, 15).FIG. 2.-Electron diffraction from 200 8, Ni at 100 mA/cm2 at 50°C unstirred ; beam along [lYO] of Cu ; note twinning spots. FIG. 3.-As fig. 2 but 10008, Ni.FIG. 4.-As fig. 2 but 10 000 8, Ni ; note faint 200 arc in plane of incidence. FIG. 5.-(a) (6) (4 extending from spots, FIG. 5.-As fig. 2 but 61 000 8, Ni : (a) beam IICu[lTO] ; (b), (c) beam11 Cu(211). Note strong arcs [To face page 1 86FIG. S.-5008, Ni at 300 mA/cm2, 50°C, unstirred bath ; note absence of twinning. FIG. 9.-As fig. 8 but 51 000 8, Ni ; strong twin- ning but smooth horizontal surface. FIG. 10.-10 200 8, Ni at 600 mA/cm2 ; 50°C, FIG. 11.-30 600 8, Ni at 100 mA/cm2 ; 5OoC, bath stirred vigorously ; surface region mainly in azimuthally limited (100) orientn. unstirred ; after Ni(OH)2 etched away. FIG. 12.-Optical micrograph ( x 800) of Ni de- FIG. 13.-34 000 8, Ni at 400 mA/cm2 ; 50°C, posit of fig. 11. bath stirred vigorously.S . K .VERMA AND H . WILMAN 187 NICKEL ELECTRODEPOSITS AT 50°C WITH THE BATH UNSTIRRED Fig. 1 shows the nickel orientations indicated by the electron diffraction photo- graphs, in the surface region (to < about 100 A depth) of nickel deposits prepared at various c.d. and thicknesses. d T T d T Region I(A) Region D(B) %t I o n I3 rl 0 100 2 0 0 3 0 0 400 5 0 0 600 current density/(m A /cm2) FIG. 1 .-Electron diffraction observations on the surface structure of nickel electrodeposits from the sulphamate bath on Cu (110) at 50°C with bath unstirred. 0 = parallel epitaxial single-crystal Ni (region I) ; t , T, a trace of, or much, octahedral twinning of the " parallel " Ni ; r, R, a small amount of, or much, random polycrystalline Ni ; / = one-degree {loo} orientation but azimuthally limited, with a cube edge near to a <211> direction of the Cu (110) substrate ; 0 = Ni(OH)2 surface layer obscuring the Ni (region 11)-the symbol inside the circle shows the Ni orientation after etching away the hydroxide, and region II(A) is where this Ni is still entirely epitaxial (parallel), while in region II(B) twinning and/or random Ni is found.In region I, parallel epitaxial crystal growth of the nickel occurs, and in regions I(A) and I(B), i.e., up to about 250 mA/cm2 at thicknesses up to 25 OOOA, and up to higher c.d. at larger thicknesses, the nickel was mainly in epitaxial orientation with its cube axes parallel to those of the copper, but with also strong { 11 11 twinning even in deposits only about 200 A thick, at up to about 150 mA/cm2 (cf.fig. 2-5). At 200 mA/cm2 the twinning in the very thin initial deposits was much less strong, and at -250-600 mA/cm2 it was absent up to a certain thickness (region I(C) in fig. l), as is seen from fig. 8 ; though twinning developed eventually at large thicknesses ; e.g., it was present at 51 000 A at 300 mA/cm2, as is shown by fig. 9-or else co- deposition of nickel hydroxide occurred (region II in fig. 1). The codeposition of nickel hydroxide (region I1 in fig. 1) was also soon accom- panied by strong twinning of the nickel and a progressively increasing proportion of randomly-oriented polycrystalline nickel, until at 60 000 A deposit thickness the nickel grew entirely as randomly oriented crystals. In region I(A) also, a barely visible trace of the nickel diffraction-ring pattern was present in photographs such as fig.3, from deposits up to a few 1000 A thick at up to188 EPITAXY OF NICKEL ON A COPPER (1 10) FACE 100 mA/cm2 or slightly more. This indicates that a very small trace of randomly oriented nickel was present, denoted by r in fig. 1. Patterns such as fig. 4 from the surface of deposits 10 000 A thick, showed a faint 200 diffraction arc centred on the plane of incidence. This arc became strong as the thickness was increased, as in fig. 5 from a 61 000 deposit at 100 mA/cm2. In fig. 5(a), obtained with the electron beam along the [liO] direction of the copper substrate, the spots from the " parallel " single-crystal nickel are strong, but are limited to a narrow [lTO] zero-order circular Laue zone which passes through the undeflected-beam spot and is centred about 24- 3 cm above the undeflected-beam spot (this copper face was then about 2" from a (1 10) plane, and the 220 Ni Bragg reflection is absent).Most of the strong spots are due to the (111) twins of this nickel lattice, forming component patterns of centred- J2-rectangle type (see fig. 6). \ I A \ FIG. 6.-Diagram showing (not to scale) the component centred-d2-rectangle spot patterns in fig. 5(a) (those of the two twins ABCO, AB'C'O, and the vertical-rectangle pattern from the " parallel " Ni), and the 200 arc E between A and A'. The arc pattern also present in fig. 5(a) appears to correspond to a preferred (100) nickel orientation but not an azimuthally random one which might have been expected to arise from the early small trace of random nickel if these crystals developed cube faces (nickel deposits on mechanically polished copper polycrystalline substrates were epitaxially random, like the copper substrate, but a preferred ( 100) orientation developed as the thickness increased).For example, fig. 5(a) does not show a strong arc on the 220 ring position, at the same level as the strong 200 arc (E in fig. 6), as would be expected from an azimuthally random (100) orientation. Fig. 5(b) and (c) were obtained at a (211) type of azimuth of the copper (and the parallel Ni), from neighbouring parts of the surface. They show that at this azimuth the electron beam was not far from along a cube axis of some of these (100)-oriented nickel crystals, thus giving rise to the square pattern of hkO arcs seen prominently in fig.5(c) and less strongly in fig. 5(b), where the vertically elongated nickel spots in J(8/3) vertical-rectangle array are strongest and confirm that this azimuth is indeed of (21 1) type, relative to the copper. The (100)-oriented nickel crystals present in these discrete azimuths must evidently be associated with the (Ill)-twin lattices of the " parallel " nickel lattice. Two of these twins, on (1 11) planes perpendicular to the Cu (110) substrate, have still a (110) plane parallel to the Cu (110) substrate, but the other two have a cube face at 19" 28' to the substrate and have a cube edge in thatS . K . VERMA A N D H . W I L M A N 189 cube face in a vertical plane through an azimuth of the Cu (1 10) substrate only a few degrees from a (211) type of direction, as is seen from fig.7. At this azimuth these two twins thus evidently give rise to the two spots of 200 type in fig. 5(b) and (c), one on each side of the plane of incidence on which the central 200 arc lies, with which they form a group similar to AEA' of fig. 6 (cf. fig. 5(a)). PLAN FIG. 7.-Side-view of cubic unit cell seen along [lYO] in the (horizontal) (1 10) plane, together with the { 11 1)-twin unit cell ; and plan view of these showing azimuthal relation of (21 1> and the cube edges of the twin. Since the arcs are longest in fig. 5(a), it seems probable that this arcing corresponds to a directed disorientation consisting of a range of rotation from the above two twin lattices, about the [lTO] lattice row which is parallel to the beam in fig.5(a). The strong spots from the twins in fig. 5(a) (cf. fig. 6 ) indeed show a noticeable tailing-off along the arc, on the side corresponding to a rotation in the sense which would bring the cube face nearer to the specimen surface (Cu (110) substrate) than their initial position at 19" 28' to it. The two component arc patterns from the two mirror- symmetrical twins, in fig. 5(a), are rotated in opposite sense, and the 200 arcs overlap across the plane of incidence. The strong (100) orientation (in these azimuths) must then be due to a development of cube faces on these twin crystals and thus a prefer- ential lateral growth and predominance of those twin crystals which (within the rotational range present) have the cube face most nearly normal to the incoming ion stream (cf.Verma and Wilman 6). The refractive drawing-out (downwards) of the strong 200 arc on the plane of incidence in fig. 5(a), (b), (c) is indeed evidence of such cube-face development. This one-sided directed disorientation from the twin orientations seems likely to190 EPITAXY OF NICKEL ON A COPPER (1 l o ) FACE be due to contacts between these two twins during their growth, with consequent generation of dislocations at their interface, and stresses in the interface region. At 300 mA/cm2, in spite of the very high rate of deposition (1000 A/s) and the general tendency for cube faces to be formed on the deposit crystals in the nickel deposits from this bath on polycrystalline copper substrates,6 the photographs such as fig.8 and 9 show by the long vertical elongation of the diffraction spots that the atomically highly smooth surface of the deposit (parallel to the Cu (1 10) substrate) at the early stages of growth is maintained up to large thicknesses. The { 1 1 1) twins in the thick deposits (51 OOOA for fig. 9) are also evidently bounded by a smooth face parallel to the Cu (1 10) substrate face, which is a (1 10) face of two of the twin lattices but a (1 14) face of the other two twin lattices giving the spots present in fig. 9 addi- tional to those from the " parallel " nickel in fig. 8 (cf. fig. 6 and 7). At 600 mA/cm2 (about 1500 A/s rate of nickel deposition), 1000 A deposits gave patterns similar to fig. 8, showing only the parallel epitaxial nickel with smooth (1 10) surface.Fig. 10, from a lO2OOA deposit, was obtained after etching away the amorphous codeposited Ni(OH)2, and shows that the epitaxial nickel and its (111) twins still tend to be bounded by a surface parallel to the copper substrate, though the vertical spot elongation is now shorter and the spots broader, indicating that these elements of surface are of smaller extent, i.e., the surface is less continuously flat. The rings also present in fig. 10 show the presence of much randomly oriented poly- crystalline nickel. At 61 000 A thickness, only a ring pattern from random nickel was observed when the obscuring Ni(OH), was etched away. NICKEL ELECTRODEPOSITS AT 50°C WITH THE BATH STIRRED VIGOROUSLY Fig. 14 shows the electron diffraction results.The locus to the right of which the codeposition of nickel hydroxide occurs is now displaced to much higher current densities, above about 500 mA/cm2. In region I to the left of this locus, the epitaxial " parallel " nickel continues during growth up to 60 000 A or more, accompanied by some (111) twinning, and also, as fig. 11 shows, by much (100) orientation, azimuth- ally limited, at 100 mA/cm2 and about 30 000 A thickness (similarly at 400 mA/cm2 and 50 000 A or more). Fig. 12 is an optical micrograph ( x 800) of the deposit which gave fig. 11, and it shows the rough topography of the deposit on the microscopic scale, as distinct from the submicroscopic form of the tips of these projections shown by fig. 11. Fig. 13 shows the highly smooth (1 10) surface of the " parallel " nickel, and the very small trace of twinning, in the surface region of a 34 000 A deposit at 400 mA/ cm2.A 25 500 A deposit at 500 mA/cm2 showed a similarly smooth surface but a larger proportion of twinning, both the parallel nickel and the twins being bounded by smooth faces parallel to the Cu (110) substrate surface. NICKEL ELECTRODEPOSITS AT 20°C WITH THE BATH UNSTIRRED Fig. 15 shows the electron diffraction results, which were made for comparison with the 50°C results of fig. 1. As in fig. 1 the region I where epitaxial nickel is obtained is bounded by a steeply rising locus on its right-hand limit, which now is at about 120mA/cm2. To the right, in region 11, there is codeposition of Ni(OH)2, accompanied by increasing twinning and random polycrystalline nickel.In region I there is again strong twinning of the " parallel " nickel, even at a few 100 A deposit thickness, at the lower current densities, and it is still considerable at 100 mA/cm2.S . K . VERMA A N D H . WILMAN 6 0 - 5 0 - 40- 2 2 -5 2 2 0 - 0" i? 3 0 - * 8 v1 44 .C( rn 10- 191 dt Region I t t t 1 I 0 100 2 0 0 3 0 0 400 5 0 0 6 0 0 current density/(mA/cm2) FIG. 14.-Electron diffraction observations on the surface structure of nickel electrodeposits from the sulphamate bath on Cu (110) at 50°C with the bath stirred vigorously ; symbols as in fig. 1. Region I r I 0 100 2 0 0 current density/(mA/cm2) FIG. 15.-Electron diffraction observations on the surface structure of nickel electrodeposits from the sulphamate bath on Cu (110) at -20°C ; bath unstirred.192 EPITAXY OF NICKEL ON A COPPER (110) FACE DISCUSSION There appear to have been no previous electron-diffraction results on the structure and growth of nickel electrodeposits from the sulphamate bath, apart froin our recent observations on nickel deposits on Cu (100) and on mechanically polished poly- crystalline copper.6 We find that the deposits from this bath are uniformly f.c.c., and this simplifies consideration of the conditions of growth at the interface with the highly smooth electropolished Cu (1 10) face. The strong twinning of the initial nickel deposits only a few lOOA thick at c.d.up to 100, or possibly 150 mA/cm2, seems to be attributable to the misfit of -2.5 % between the nickel and the copper lattices. This twinning, and possibly the barely detectable trace of random polycrystalline nickel present (cf. fig.l), may perhaps originate in the Ni-Cu interface regions where neighbouring initial nickel crystal nuclei meet together as they grow larger. The reduction in twinning at higher c.d. -200 mA/cm2, and its eventual absence in the initial deposits at 300-600 mA/ cm2, is an observation of a type which has not apparently been observed before, and indeed most studies have not extended up to such high c.d., i.e., high rates of deposition. It seemed likely hitherto that the probability of twinning would increase with increase in rate of deposition, there being then more chance of some atoms being trapped in the alternative set of potential troughs on the growing surface to nucleate a twin lattice, and being overlaid by other atoms and thereby to become built in.However, at 300-600mA/cm2 or more (1000-1500 A/s), the rate of arrival of energy at the cathode must be so high that there will be appreciable rise in temperature in the solution near the cathode, and in the growing cathode surface itself, so that there will be an increase in atomic (or adion) mobility on the growing cathode. In a sense, this can be said to be of the nature of an annealing effect. This decrease in twinning at these high c.d. is matched in our results for epitaxial nickel growth on electropolished Cu (1 11) faces (to be published), and also on the Cu (100) face where, in the latter case, we observed only an occasional faint trace of twin- ning at the lower c.d., and none at the higher c.d. range.On the Cu (100) face, we observed a greatly reduced tendency of the deposit to develop undulating or obliquely facetted surface in the higher c.d. range, at least up to large thicknesses; and this is also attributable to a rise in temperature at and near the cathode surface, with a corresponding diminution of the " outward-growth " tendency.' These results therefore reinforce our earlier conclusions as to the appreciable temperature rise in the surface region of growing deposits.s-lo The absence of any " directed disorientation " of the initial epitaxial nickel in the interface region is interesting, since the nickel is in tension; however, the misfit is only - 2.5 %, thus much smaller than for Cu on Ag, where it is - 11.5 % and a directed disorientation is obser~ed.~ In our results on the epitaxy of Fe and Cr on C U , ~ the misfit was only -2.7 % along the most densely populated atom rows, but about + 12 % in the perpendicular direction ; and extensive directed disorientations were observed. The observed axis of rotational disorientation is not always easily explained in terms of normal slip systems and processes, and it is not clear whether the rotations are of multiple translational slip type, or flexural translational slip (about an axis parallel to the slip plane and normal to the slip direction), or of lamellar rotational- slip type 11-15 with the slip lamellae perpendicular to the axis of rotation.' The directed disorientation in the thicker epitaxial twinned deposits seems to be similarly related to interface regions where two of the twins meet as they grow bigger.Menzer 16* l7 has discussed the fitting together of two such differently oriented twinsS . K . VERMA AND H . WILMAN 193 of a f.c.c. lattice such as nickel, and he concluded that certain superlattices should be possible at the interface region. However, there seems scope for an appreciable misregister of the lattices to occur, with consequent introduction of dislocations and stresses, which might cause such disorientations to develop and to increase during further growth of the deposit. The effects of stirring and of increase in bath temperature in extending region I to higher c.d. (cf. fig. 1, 14, 15) are similar in type to those discussed for Ni on Cu We thank Mr. E. H. Reynolds, Dr. R. M. Hinde and Mr. R. E. Davies of the Central Research and Engineering Division, British Insulated Callender’s Cables Ltd., for initiating the research and providing the materials. We also thank BICC Ltd for the bursary which enabled one of us (S. K, V.) to carry out the experiments. D. M. Evans and H. Wilman, Proc. Phys. SOC. A., 1950,63,298. A. P. Goswami, Ph.D. Thesis (University of London, 1950). D. N. Kumar, 1955, D.I.C. Thesis (Imperial College, London) ; see also H. Wilman, Acta Cryst., 1957, 10, 842. A. K. N. Reddy and H. Wilman, Trans. Inst. Metal Finishing, 1958-9, 36, 97. S. K. Verma and H. Wilman, J. Phys. D.: Appl. Phys., 1971, 4, 1167. S. K. Verma and H. Wilman, J. Phys. D.: Appi. Phys., 1971, 4, 2051. G. I. Finch and D. N. Layton, J. Eiectrodepositors’ Tech. Suc., 1951, 27, 215. H. P. Murbach and H. Wilman, Pruc. Phys. Soc. B, 1953, 66,905. * D. M. Evans and H. Wilman, Acta Cryst., 1952, 5, 731. lo H. Wilman, Pruc. Phys. SOC. B, 1955, 68,474. l 1 H. Wilman, Nature, 1950, 165, 321. l 2 H. Wilman, Proc. Phys. SOC. A, 1951, 64, 329. l 3 A. D. Whapham, J. Inst. Metals, 1956, 84, 109. l4 A. D. Whapham and H. Wilman, Nature, Lond., 1955, 176,460. l5 A. D. Whapham and H. Wilman, Proc. Roy. SOC. A, 1956,237, 513. l6 G. Menzer, Naturwiss., 1938. 26, 385. l7 G. Menzer, 2. Krist., 1938, 99, 378, 410.

ISSN:0370-9302    

DOI:10.1039/S19720200185

出版商:RSC    

年代:1972

数据来源: RSC

摘要:

Interface Crystallography of Iron-Group Metals and Alloys Electrodeposited on Copper BY J. M. O'SULLIVAN" AND D. P. OXLEY~ Received 12 June, 1972 A wide range of binary Co, Fe, Ni alloys has been prepared by electrodeposition on to cube- textured copper substrates. The alloys, after removal from the substrate were examined by trans- mission electron microscopy. All alloys were found to be of either b.c.c., f.c.c. or c.p.h. structure according to their composition. The interface crystallography of these alloys is discussed in terms of the relationships between prominent directions in the iron-group alloys, and the direction of the copper face diagonals in the cube-textured substrate. The metals and binary alloys of the iron, nickel, cobalt group are of interest because of their ferromagnetic properties,l.which have important applications in the thin foil state in magnetic recording and computer memory devices. Although these metals and alloys have been used for such applications, the crystallography at the interface between substrate and deposit foil has not been surveyed for all possible alloys of all the possible phases in this group. It is thought that a detailed knowledge of the crystallographic structure is important in interpretation of the ferromagnetic properties of such materials especially in the thin foil condition. Previously, most attention has been given tc alloys of the nickel +iron system, in which both f.c.c., b.c.c. interface crystallography with respect to copper have been investigated. In the other binary systems containing cobalt, measurements of magnetic parameters have been made but structural factors have not had such attention.The purpose of this paper is to extract the general crystallographic relations from the interface situations available in the iron, nickel, cobalt group electrodeposited on to electropolished cube-oriented copper foil. The group provides examples of f.c.c., b.c.c. and c.p.h. foils. Since lattice constants of the metals and alloys of a given phase are close, the crystallographic relations are described conveniently by means of three simple two-dimensional models of the interface structure for f.c.c., b.c.c. and c.p.h. phases. These are based on the transmission electron diffraction pattern of a typical metal or alloy in each phase. EXPERIMENTAL All metals and alloys were deposited from static aqueous salt baths at a current density of 10 mA cm-2.The foils were chemically removed from their substrates and examined by transmission electron diffraction. Analyses of foil compositions were carried out by atomic absorption spectrophotometry. * Department of Physics, University of Warwick, Coventry. t School of Physics, Leicester Polytechnic, Leicester. 194J . M. O'SULLIVAN AND D. P. OXLEY 195 RESULTS F . C . C . FOIL DEPOSITION O N COPPER F.c.c. foils of nickel, cobalt, nickel-iron,2 iron-cobalt and nickel-cobalt may be deposited on to copper so that (100) [Ol 11 alloy 11 (100) [Oll] Cu. Some evidence of (1 11) microtwinning was also apparent in all systems. Grain size was comparable to that of the substrate, -0.01 mm.Plate l(a), a diffraction FIG. 1. pattern of a (100) 94 nickel, 6 % cobalt alloy, shows extra diffraction features at A and B indicating twinning. The geometry and associated diffraction effects have been thoroughly discussed previ~usly,~ and these are not included in the model (fig. 1) of the arrangement of planes in copper and alloy in this situation. B . C . C . FOIL DEPOSITION O N COPPER The iron+cobalt and ironfnickel systems give b.c.c. alloys in the iron-rich regions. Some previous work has been carried out on iron+nickel alloys of this phase,2 but this has not been extended to the iron+cobalt system. It was found in the present work that the interface structure of both systems in the b.c.c. phase is identical and can best be stated as where the last term expresses the misfit between two sets of orthogonal (110) deposit orientations in the foil plane as in fig.2. When @ = 19" 30', the structure consists of two orthogonal ( 1 lo] deposit orientations. Grain size in these alloys was small, - 500 A in extent, with no marked growth habit. A typical diffraction pattern of a 43 % cobalt, 57 % iron alloy foil is given in plate l(b). 2-G*196 IRON-GROUP ALLOYS ELECTRODEPOSITED ON COPPER - ~ l l o y - ----- Copper FIG. 2. FIG. 3.PLATE 1 [To face page 1 96J . M. O'SULLIVAN AND D. P. OXLEY 197 C.P.H. FOIL DEPOSITION ON COPPER As representative of the c.p.h. phase, only cobalt was stripped from the substrate successfully, even then only with great difficulty, in small pieces. The interface crystallography is described by (001)[210]aCo~~{ 100)(011)Cu, so that there are six equivalent positions of the cobalt basal plane.One such position is shown in fig. 3, which is based on plate l(c). The crystallographic relations found here are in agreement with those determined by a reflection electron diffraction method for cobalt foils on copper single crystal (001) faces.' Due to difficulties in stripping cobalt from the substrate, it was not possible to measure the grain size accurately. DISCUSSION Matching in the interface planes may be discussed in a unified fashion with refer- ence to a particular dimension in the copper substrate (100) surface. Taking the lattice parameter for copper as 3.61 A, typical f.c.c. as 3.55A, b.c.c. as 2.86A and c.p.h.(ao) as 2.50A, a slightly more quantitative model may be used to support that found from the diffraction experiments reported above. From the diffraction experi- ments and approximate lattice parameters, the copper cube-face diagonal may be used to describe the interface crystallography. For f.c.c. deposition the misfit along cube- face diagonals is small, i.e., 5.02 A in the alloy against 5.12 A in copper. For b.c.c. deposits, misfit is small between the body diagonal, 4.95 A, and the copper cube-face diagonal. In the c.p.h. deposit there is closer agreement between a basal-plane diagonal and the copper cube-face diagonal. Thus, the variety of structures available in thin foil metals and alloys of the iron group may be correlated with a tendency to minimize misfit along prominent directions in substrate and electrodeposit. Hence, the main structures obtainable in thin foil form in the iron group electrodeposited on to copper foil have been described with respect to the interface crystallography between substrate and deposit.The importance of copper cube-face diagonals in this description has been noted. It was found that close correlation occurred in the length of the copper cube-face diagonal and prominent directions in the various iron group structures. The bulk of the above work was carried out in the Applied Sciences Department, Lanchester Polytechnic, Rugby. The authors are especially grateful to Dr. R. F. Y. Randall, Head of Department, for provision of laboratory facilities and encouragement during the project, and to Mr. N. F. Hall who carried out analyses of foils. H. W. Fuller and M. E. Hale, J. Appl. Phys., 1960, 31, 238. G. A. Jones, D. P. Oxley and R. S. Tebble, Phil. Mag., 1966, 14, 881. I. M. Wolf, J. Appl. Phys., 1962, 33, 1152s. J. M. O'Sullivan and D. P. Oxley, J. Elect. Micr., to be published. ' R. D.. Burbank, R. D. Heidenreich, Phil. Mag., 1960, 5, 373. ti G. A. Jones, Phys. Stat. Sol., 1967, 19, 811. J. Goddard and J. G. Wright, Brit. J. Appl. Phys., 1964, 15, 807.

ISSN:0370-9302    

DOI:10.1039/S19720200194

出版商:RSC    

年代:1972

数据来源: RSC

摘要:

Mechanism of Formation and Some Surface Characteristics Electron Bombardment of Thin Polymer Films Formed on Metal Surfaces by BY S. FROST, W. J. MURPHY,? M. W. ROBERTS," J. R. H. Ross AND J. H. WOOD School of Chemistry, The University of Bradford, Bradford Received 23rd June, 1972 The formation of thin films by electron bombardment has been studied and some of the important factors controlling the nature and surface characteristics of the films have been identified. A number of experimental techniques have been developed in order to obtain information on both the species likely to be involved in film growth and also the nature of the film surfaces. The surface character- istics have been monitored by contact angle studies, The possible importance of negative ions in film formation is illustrated with Si(CH&, where over 20 different negative ions were detected, the major ones being H-, C2H- and Sic;.Such results also have an important bearing on electron affinity data obtained by the magnetron method. Thin films formed from perffuorobut-Zene have surface characteristics which vary with the monomer pressure, the substrate temperature, and electron and ion energies. Information on the relative importance of the role of negative ions compared with electron and adsorbed monomer reaction is obtained. Some suggestions are discussed regarding possible surface structures inferred from wetting characteristics. Thin polymer-like films have been of interest for some years and methods of preparation have included glow discharge, u.-v.radiation and electron bombard- ment.', The object of this investigation is an understanding of the mechanism by which thin films (e g., < 1000 thick) are formed by electron-monomer interaction, the factors that control both the adhesion of the film to the substrate (in this work a metal) and the surface characteristics of the film. Although the monomer tetramethyl tetraphenyl siloxane is technologically important, we have chosen to start with struc- turally simpler, but related molecules, e.g., tetramethyl silane Si(CH& Such thin films are of interest 3* as insulators, dielectrics, computer storage elements and " chemically inert " coatings. We have adopted a number of distinct experimental approaches. Adsorption studies of the monomer (in the presence and absence of an electron beam) on a number of atomically clean metal substrates give direct information on the bonding at the metal/polymer interface and therefore are relevant to the question of adhesion.Secondly, we considered it important to investigate whether any reactions occurring at the electron emitting filament were significant in determining the nature of the polymer film. Thirdly, the surface characteristics of the thin films were monitored by determining the critical surface tension yc which Zisman 5 9 considered to reflect the molecular naure of solid surfaces. The possible variation of these surface char- acteristics with the conditions of film preparation (monomer pressure, substrate temperature, electron energy) was investigated. j. present address : Department of Chemistry, New University of Ulster, Coleraine.198FROST, MURPHY, ROBERTS, ROSS AND WOOD 199 These objectives led to the pursuit of a number of more fundamental questions, namely the nature of species resulting from molecular interaction with an electron emitting filament and the role of negative surface ionization, the significance of electron affinity values determined by the magnetron method and the significance of the critical surface energy yc of low energy polymer surfaces. EXPERIMENTAL There are a number of distinct experimental problems to be resolved, each with its special requirements. Adsorption studies were performed in a UHV (- lo-' Torr) all glass apparatus with mass-spectrometric facilities for gas analysis and for following H2/D2 exchange.' Identification of the current carriers participating in film formation required us to dis- tinguish between electrons and negative ions.A magnetron similar to that used by Page and co-workers * was constructed for this and was used for pressures in the range loF6 to N Torr. At a later stage of the work, a smaller version of the magnetron was coupled to an EAI Quadrupole mass-spectrometer which had been modified to monitor negative ions. This enabled the molecular nature of the negative ions to be determined; pressures in the range 5x to -5x Torr were used. Electrons emitted by the tungsten filament were diverted by the magnetron in order that they did not interfere with the recording of the negative-ion mass-spectrum. FIG. 1.-Details of the reaction vessel used for " polymer " film formation.Insert : Schematic representation. The thin polymer-like films formed specifically for surface studies were prepared in the reaction vessel shown in fig. 1. It consists of an optical glass plate on to which a substrate film of tungsten is evaporated from a filament ; electrical contact to the substrate (anode) is made via platinum wires. The temperature of the substrate can be varied by circulating liquid nitrogen, hot water, etc. The spherical reaction vessel had appendages for an electron gun, a Pirani gauge, facilities for introducing the monomer through a capillary leak and for200 SURFACE CHARACTERISTICS OF POLYMER FILMS fast pumping. The ultimate pressure in this system was -lo-’ Torr. The surface char- acteristics of films formed on the anode were obtained as described elsewhere for bulk solids.The “ standard conditions ” of film preparation were : electron gun (W) filament temperature 2300 K ; anode potential (V*)+ 500 V ; grid potential ( VG)+ 35 V; substrate temperature 298 K; monomer pressure in reaction vessel 5 x Torr ; time of film preparation 30 min. RESULTS AND DISCUSSION CHEMISORPTION STUDIES We have reported recently on the adsorption of tetramethyl silane Si(CH3)4 on tungsten and iron surfaces. Fragmentation of the molecule occurs on W at 293 K with the desorption of H,(g) and CH,(g), while more extensive breakdown occurs at higher temperature. With Fe, Si(CH3), is adsorbed temperatures than with W. We have good evidence dissociated surface species on W include CH3 CH3 CH3 CH3 I / H,C- i \ / Si but is less fragmented at all to suggest that at 293 K the and j The bonds formed to the surface atoms * are strong and are likely to lead to good adhesion. How these surface groupings are perturbed under polymerization con- ditions is not known, but some observations on the influence of electron interaction are relevant.When a W film surface, which had chemisorbed Si(CH3)4 at 295 K, was bombarded with electrons (100 pA at 150 V) for - 16 min, H,(g) and CH,(g) were desorbed. The number of product molecules corresponded to one per impinging electron. On the other hand, when physically adsorbed Si(CH3)4 present in muItilayers on tungsten at 77 K was bombarded with electrons, the extent of desorption of CH4(g) and H,(g) was small but extensive and efficient conversion of the monomer to “ poly- mer ” occurred.The latter is concluded from the fact that not all the Si(CH,),(g) was recovered on warming the “ bombarded layer ” to 293 K. We might therefore anticipate quite distinct surface properties with films formed under different condi- tions and we will return to this point later. Further Si(CH3)4 was adsorbed reversibly. ROLE OF NEGATIVE IONS IN FILM FORMATION That ions are involved in film growth was first suspected from current measure- ment during the interaction of Si(CH3)4 (P- to 5 x Torr) with an electron- emitting W filament. The rate of disappearance of Si(CH3)4 was close to the rate of molecular impingement with the W filament. This could clearly have been fortuitous, but its significance was emphasized when the current increased with increasing Si(CH3)4 pressure.By making use of the magnetron to distinguish between electron and ion current, we concluded that with the majority of the molecules studied, including Si(CH3)4, a high proportion of the carriers are negative ions in the pressure range Torr. These ions are accelerated to the anode (fig. 1) and presumably participate in film formation. The next question to consider was the nature of these ions ; we will not concern ourselves here with the detailed mechanism by which they form. The molecular nature of the ions leaving the W filament was investigated using the modified quadrupole mass-spectrometer coupled to the magnetron ; table 1 The nature of these current carriers is clearly important.toFROST, MURPHY, ROBERTS, ROSS AND WOOD 20 1 summarizes the major negative ions (with their relative abundances) observed with .- c .$-a r C 'C H E 0 0 -5 E 6 H tetramethyl silane. - With-perfluorobut-2-ene the F; in the respective proportions 1500 : 28 : 5 : in this molecule will become clear later. u u B H U C H 3- C H r -CH2 I r.CH- -CH-,;r.CH ;z: 2 B 2 -CHOH -CHg -COsCHg r..CHm TABLE 1 .-NEGATIVE IONS OBSERVED WITH Si(CH3)4 -2000 K main ions were F', C ; , CF- and 80. The reasons for our interest AND A w FILAMENT MAINTAINED AT ion H C CH CH2 C2H SiH SiH2 relative intensity (arb units) 760 8 23 6 690 103 48 ion SiCH SiCH2 SiCH3 c4 C4H Sicz SiC2H SiH3 1 SiC2H2 f4 c3 45 SiC3H 69 C3H 38 SiC3H2 34 C3H2 12 Si(CH3)4 140 relative intensity (arb units) 83 8 4 24 120 345 103 c..THE POLYMER SURFACE. WETTING STUDIES BASIS OF THE APPROACH Zisman et aL5* 6* lo have shown that for a homologous series of pure liquids, a plot of cos 8 against yLv (8 = contact angle, yLv the liquid surface tension) was linear with an intercept on the cos 0 = 1.0 line which Zisman designated yc and termed the critical surface tension of the solid. yc has been shown to be a useful empirical parameter related to the surface energy of the solid ys. Of particular significance to the present work was Shafrin and Zisman's correlation of yc with the molecular constitution of the solid surface (fig. 2). This " wettability spectrum " shows a distinct empirical correlation between yc and the chemical constitution of low-energy surfaces. The lowest values of yc (i.e., less wettable and higher contact angle) are202 SURFACE CHARACTERISTICS OF POLYMER FILMS obtained with surfaces comprising C and F atoms only (6 to 19 dyn cm-l) ; hydro- genation increases yc to the range 15 to 28 dyn cm-l.The next higher range of yc values is observed with surfaces containing only C and H atoms (22 to 35 dyn cm-') while higher values are obtained with C, C1, H, F; C, 0, H and C, N, 0, H surfaces. We note in particular the following : CF3 (yc = 6 dyn cm-l), CF2 (yc N 18 dyn cm-'), CH3 (yc N 24 dyn cm-l) and CH2 (yc N 30 dyn cm-l). the wetting behaviour of a number of polymer surfaces using both butanol solutions and a series of pure liquids, the yc values observed with the solutions were much lower than with the pure liquids and what is more significant, were virtually independent of the solid.This led naturally to the conclusion that yc, at least with the solutions, reflected the perturbation of the molecular nature of the surface by adsorption. This was shown convincingly to be the case when yc for any one (low energy) solid varied in a predictable manner when solutions of BuOH, MeOH and 2,2,2, trifluorethanol were used (table 2). It followed that the adsorbed molecules were likely to be oriented with the hydrocarbon groupings pointing away from the solid surface, viz : polymer- - - - -HO-CH2-CH3. Although such an orientation, with the hydrocarbon end pointing away from the surface, is not generally accepted, Ottewill and Vincent have suggested recently that both orientations may occur in the polystyrene + BuOH system.Clearly, yc obtained with solutions is not necessarily diagnostic of the molecular nature of the film surface. When we investigated TABLE 2.-yc VALUES (dyn crn-') FOR A NUMBER OF SOL^ SURFACES DETERMINED BY DIFFERENT LIQUIDS AND SOLUTIONS pure liquids BuOH EtOH MeOH 2,2,2,-trifluoroethano polystyrene 41-44 25.3 - 23.0 - polymethylmethacr ylat e 39-41 26.0 26.5 23.0 24.7 polyethylene [- CH2-CH2-], 31-32 27.5 - - 27.5 Further support for adsorption being important was obtained by combining the Gibbs adsorption isotherm with the conditions determining the equilibrium contact angle 8 to give eqn (1) where r2 is the surface excess of solute (component 2) at the interface. At low concentrations r2 is equivalent to the amount of solute adsorbed.Thus, gradients of plots of yLv cos 8 against In a (a = activity of solute) will be related to the difference in the surface excess of solute molecules at the solid-liquid and solid-vapour interfaces. For example, a variation in the contribution of adsorption at the solid/liquid and solid/ vapour interfaces with concentration would be reflected in the form of the plot. Such a variation is observed with different surfaces, and is evident in fig. 3, 4 and 5. For the present paper, however, we base our discussion mainly on the variation of Ow, the the contact angle observed with pure H20 (yLv N 72 dyn cm-l). STUDIES OF THIN FILMS BASED ON PERFLUOROBUT-2-ENE The possibility of obtaining from perfluorobut-2-ene, CF3-CF=CF-CF3, thin films whose surface properties were analogous to polytetrafluoroethylene (PTFE), yc E 18 dyn cm-l, and apparently involving CF2 and possibly CF3 surface groupings (table 3) was attractive. The physical properties of perfluorobut-Zene enabled itFROST, MURPHY, ROBERTS, ROSS AND WOOD 203 to be handled quantitatively in our mass-spectrometric and magnetron studies.Thin films were prepared using the " standard procedure " already outlined and their surface characteristics investigated by contact angle studies using various alcohol solutions and pure water. These characteristics were compared with films prepared from other related fluorocarbon " monomers " and also with those reported for fluorocarbon surfaces thought to have CF3 and CF2 groupings (table 3).The influence of such preparation conditions as monomer flux, substrate temperature and accelerating potentials were also investigated. Fig. 3(i) shows " Zisman " plots using BuOH solutions for two samples each of films formed from (a) octafluorobutane, (b) perfluorobut-Zene and (c) tetrafluoro- methane. These plots can be compared with those for films prepared from (a) methane, (b) ethane, (c) ethylene, (d) butane, (e) butadiene and (f) neopentane (fig. 3(ii)). We conclude that although the apparent yc is not particularly sensitive to the monomer, whether hydrocarbon or fluorocarbon, (not surprising in view of our adsorption model ') there are distinct differences in the manner in which cos 8 varies with yLv within the fluorocarbon series (fig. 3(i)). In all cases when Ow decreased, the apparent yc also decreased. 1.c 0.E 0, t 8 0 . 4 0.2 Table 4 summarizes Ow and yc values, the latter obtained with pure liquids, for hydrocarbon surfaces with characteristic surface groupings and also for thin films formed from various hydrocarbons. It is clear that as the degree of unsaturation of204 SURFACE CHARACTERISTICS OF POLYMER FILMS I .o 0. s 0.7 m y, 0.5 8 0.: 0. FIG. 3.-(ii) Zisman plots (as in (i)) for thin films formed from (a) methane (0), (b) ethane (01, (c) ethylene m), (d) butane (+), (e) butadiene (- - -), (f) neopentane (- - -). 0.8 0.6 y, 0.4 8 0.2 0 Q 2 0 4 0 6 0 8 0 YLvldYn cm-' c \ x 10 IV xz x 103 FIG. 4.--Influence of the variation of the monomer flux on the wetting behaviour of perfluorobutene-2- films.(a) Zisman plot, (b) a plot of [(F2)SL- (r2)sv] against the mol fraction XZ.FROST, MURPHY, ROBERTS, ROSS AND WOOD 205 8 LOOOV . \"I -6.0 1 FIG. 5.-Influence of the anode potential on (a) the Zisman plot and (b) a plot of [(r2)s~- (I',)sv] against the mol fraction X,. the surface grouping increases (i.e., CH3 to CH2 to CH) & decreases from - 110" to - 55", and yc increases from -22 dyn cm-' to - 34 dyn cm-I, reflecting an increase in surface energy. TABLE 3.-cOMPARISON OF Bw(deg) AND ')Jc (dyncrn-') FOR (U) SOME THIN FLUOROCARBON POLYMER FILMS WITH (6) " STANDARD " FLUOROCARBON SURFACES OW Yo * (a) thin films (monomer) CF4 69.0 CF,CF=CF CF3 59.0 - 35 F2 F2 I I I I 49.5 - (b) perfluoroacid monolayer (--CF,) 102.0 -6 (-CF2-) 108.0 - 18 poly tetrafluoroet h ylene * values obtained using pure liquids206 SURFACE CHARACTERISTICS OF POLYMER FILMS On this basis we generalize and suggest that a surface which exposes " unsatur- ated " groupings will be of higher energy than that exposing " saturated " groupings.We therefore accept that Ow will reflect both the molecular nature of the surface and yc, TABLE 4.-cOMPARISON OF Ow(deg) AND 'yc(dyn Cm-') FOR (a) SOME " STANDARD " HYDRO- CARBON SURFACES AND (6) THIN FILMS FORMED FROM A SERIES OF HYDROCARBON MONOMERS ow Y C * (a) standard hydrocarbon surfaces CH3 (hexatricontane) CH3 (paraffin) CH2 (polyethylene) CH2, CH (polystyrene) CH (anilene monolayer) (b) thin films (monomer) CH4 C2H6 C4HIO 110 22 108 22 94 32 87 34 53 35 74 35 71 69 - - * obtained using pure liquids and consider what influence film preparation conditions have on the magnitude of Ow and the adsorption characteristics as assessed from the application of the Gibbs adsorption isotherm.In the present paper the latter is used merely as a qualitative pointer of the adsorption behaviour. VARIATION OF MONOMER FLUX A marked variation in the wetting behaviour was observed (fig. 4) as the flux (pressure) of perfluorobut-2-ene was increased by factors of 10 and 20. 8, increased, which is in keeping with the view that increasing the flux rate encouraged the formation of a more saturated surface (see later). Analysis of the data in terms of the Gibbs- Adsorption Isotherm provides further evidence regarding variation in the surface characteristics (fig. 4). VARIATION OF ACCELERATING POTENTIALS Fig.5 shows that both the Zisman plots and the Gibbs adsorption isotherm data are influenced by the potential V, applied to the anode. As the anode potential increases from 300 to 1000 V, the contact angle Ow for pure water (yLv N 72 dyn cm-l) decreases. In other words, as V, increases the film surface becomes more wettable. These changes in Ow are also reflected in the Gibbs adsorption isotherm plotted according to eqn (1) (fig. 5). When the grid potential V, was varied, keeping the anode potential fixed at the " standard value " of 300 V, the surface characteristics of the thin films formed changed. The contact angle of water, for example, decreased as the voltage VG was decreased from 300 to 100 to 35V. This change in surface behaviour was also reflected in the Gibbs adsorption plot.SUBSTRATE TEMPERATURE Preliminary data obtained with the substrate maintained at 77 K indicate that the film exhibits (fig. 6) a very low energy surface (y,<20 dyn cm-l, see fig. 2). In thisFROST, MURPHY, ROBERTS, ROSS AND WOOD 207 case it would appear that films formed from perfluorobut-2-ene are more similar to PTFE (table 3) and could indicate that we are concerned here with the interaction of electrons with multilayers " adsorbed " on the substrate. Extensive adsorption of the monomer will not occur at 298 K and so electron/adsorbed monomer interactions will be less significant under these conditions. 298K 77K 3 FIG. 6.-The influence of substrate temperature on the surface characteristics as assessed from contact angles (0) observed with aqueous solutions of BuOH.Apparent y, values are also indicated. GENERAL CONCLUSIONS As to the mode of formation of the thin polymer films, there is clear evidence for the participation of gaseous negative ions in the growth process. Previous work has emphasized the formation of active species by electron interaction with adsorbed monomer, Haller and White arguing that growth involved a surface reaction between adsorbed monomer and active sites (positive ions) produced by electron bombard- ment. With condensed multilayers of perfluorobut-2-ene at 77 K we believe that the electron/adsorbed monomer reaction is important, but at 298 K ion-ion reactions appear to be more significant. We shall return to this point. Ionized species have of course been freely invoked in studies of polymer formation induced by electrical discharges but neither the nature nor the abundance of the ions involved has been clarified. The surface characteristics of the polymer films are clearly dependent on the nature of the monomer and ions, the monomer flux and substrate temperature. Infra-red studies did not give any evidence on their structure.We therefore have relied heavily on contact angle studies and the way in which various parameters control the wetting behaviour and presumably their molecular nature. For example, the anode potential (fig. 1) will determine the energy of the ions and electrons so that an increase in V, is likely to lead to more extensive cross-linking, giving a surface of higher energy with a smaller contact angle 8,; this is observed (fig.5). Variations in the grid potential will determine the ratio ze/& reaching the substrate as well as the total flux Ie+IN. Thus when V, decreased from 300 to 35 V the electron and negative ion currents will also decrease. We do not, however, know how the ratio varies; it will depend on such factors as the electron affinities S of the desorbed ion(s), work function q5 of the filament, etc., terms which are involved in the relationship (eqn (2)) which we have considered dN' ew-4o-l- JZ) - = AB,Eexp dt kT208 SURFACE CHARACTERISTICS OF POLYMER FILMS elsewhere.12 Although further studies to elucidate the detailed mechanism are envisaged, and also an investigation of the fate of the ions (e.g., those listed in table 1) impinging on the substrate surface, tentative mechanisms for the formation of two polymer films formed from perfluorobut-2-ene but with distinctly different surface characteristics, are given below.They illustrate our general approach rather than provide firm conclusions. In the case of film formation using perfluorobut-2-ene as monomer, an electron emitting W filament and a substrate maintained at 298 K, we suggest that the main film growth process involves negative ions (e.g., CF-) and extensive cross-linking. The surface groupings would be C4F8(g) * CF-@ 1 3 . 1 I I l l 1 I I I I I l l 11111111111111111 substrate -CF CF-CF CF-CF CF-CF -CF CF-CF unsaturated, which, according to our suggestions based on Zisman's wettability spectrum, exhibit a high surface energy.The value of Ow in this case was -59" and yo using pure liquids, -35 dyn em-l, both values in keeping with a high energy surface. On the other hand, when multilayers of C4F8 are bombarded with electrons at 77 K, no significant concentrations of gaseous negative ions are involved and we suggest that film growth occurs by electron initiation in the adsorbed monomer. This process is very efficient for polymer formation using multilayers of Si(CH3)4. I C,F,(adS) + e-, CF3-CF-CF-CF3 I F3 C-CF-CF-CF3 I llllllllllllllllllllllt t substrate Such films have characteristics approaching those of PTFE, yc using aqueous BuOH solutions was N 18 dyn cm-1 and Ow N 75". The proposed structure exhibits " satur- ated " surface groupings and according to Zisman's wettability spectrum the surface would be expected to be of " low energy ". Finally, our results indicate the need for extreme caution in determining electron affinity data by the magnetron method.* We have explored a number of molecules in addition to those reported here ; in all cases extensive fragmentation is observed. Furthermore, some of the ions are formed by molecule+electron interaction in the gas phase rather than by surface ionization. This is observed. We acknowledge generous support of this work by T.R.W. Inc. (California) through a grant (N.A.S.A. contract 7-717) to one of us (M. W. R.).FROST, MURPHY, ROBERTS, ROSS AND WOOD 209 I. Haller and P. White, J. Phys. Chem., 1963, 67, 1784. D. S. Allam and C. T. H. Stoddart, Chem. in Britain, 1965, 1,410. H. T. Mann, J. Appl. Phys., 1964,35, 2173. R. W. Christy, J. Appl. Phys., 1960,31, 1680. W. A. Zisman, Adv. Chem. Ser. (Amer. Chem. SOC.), 1964,43, 1. E. G. Shafrin and W. A. Zisman, J. Phys. Chem., 1960,64,519. M. W. Roberts and J. R. H. Ross, J.C.S. Faraday I, 1972,68,221. Interscience, N.Y. 1969). W. J. Murphy, M. W. Roberts and J. R. H. ROSS, J.C.S. Faraday I, 1972,68, 1190. * See, for example, Negative Ions and the Magnetron, F. M . Page and G. C. Goode (Wiley- lo M. K. Bernett and W. A. Zisman, J. Phys. Chem., 1959, 63, 1241. l 1 R. H. Ottewill and B. Vincent, J.C.S. Faraday I, 1972, 68, 1533. l2 M. W. Roberts and J. R. H. Ross, Chem. Comm., 1970, 1170.

ISSN:0370-9302    

DOI:10.1039/S19720200198

出版商:RSC    

年代:1972

数据来源: RSC

摘要:

Mechanical Degradation of Thin Polystyrene Films BY ROBERT J. NASH AND DARRYL M. JACOBS Xerox Corporation, Rochester Research Center, Webster, New York 14580, USA Received 1 st June, 1972 Thin films of polystyrene drastically degrade, on a molecular scale, when subjected to mechanical forces. An examination, by gel permeation chromatography, of the molecular weight distributions of the degraded polymers shows that the polymer molecules do not undergo the random scissions or bisections that have been postulated in the many prior theoretical treatments of polymer degradation, but rather degrade via non-random multiple scissions. The mode of degradation is such that while medium molecular weight polystyrene molecules degrade directly to a limiting low molecular weight species, high molecular weight polystyrene degrades both indirectly through an intermediate species and directly to the limiting species.An analysis of the degradation data from a wide molecular weight distribution polystyrene sample yields a rate equation which is a decreasing function of molecular weight. A model is proposed which predicts the form of the degradation rate equation on the basis of molecular size : the mode of degradation is explained by the concept of polymer chain entanglement. Possible degradation pathways are differentiated using a computer simulation. The phenomenon of mechanical degradation of polymer molecules is well known. It is manifest in such diverse examples as the bond scission which can occur when polymer solutions are merely stirred or shaken,2 and in the more practical problem of molecular breakdown generated during the spinning or extrusion of polymer melts.Polymer degradation has been the subject of several and many experi- mental 6-25 and theoretical Unfortunately, a majority of the experi- mental studies have followed the degradation process through viscometric measure- ments which, by their nature, yield only limited information about the molecular weight distribution of the degraded species. This lack of distribution data has hindered theoretical developments, both in the formulation of models and in the testing of hypotheses. However, the recently developed experimental technique of gel permeation chr~matography,~~ which yields rapid measurements of molecular weight distributions, will alleviate this problem.To date, only a few studies of shear 22 and ultrasonic induced 21 degradation of polymer solutions have utilized this technique. The present study is the first systematic application of gel permeation chromatography to the process of mechanical degradation of thin, solid polymer films, and indeed appears to be the first mechanical degradation study on such films. EXPERIMENTAL MATERIALS Narrow distribution polystyrene samples having weight average molecular weights, M,, of 2~ lo5, 6 . 7 ~ lo5, and 8 . 6 ~ lo5, and molecular weight distributions, MWD, less than 1.2 were obtained from the Pressure Chemical Co. The National Bureau of Standards supplied a wide distribution polystyrene sample, NBS-706, having an Mw of N 3.6 x lo5 and a n MWD of -2.(For brevity, the samples will be further referred to as PSZOO, PS670, PS860 and NBS-706 respectively). 210R . J . NASH AND D . M. JACOBS 21 1 COATING A N D DEGRADATION PROCEDURES In a preliminary series of experiments, 2 pm films of polystyrene were solvent-cast onto the inner surface of glass jars. The addition of 2 mm diameter steel beads to the jars fol- lowed by rolling on a roll-mill served to mechanically degrade the films. However, this procedure gave poor results, partly because of non-uniformities in the films but chiefly because of the tendency of the films to flake from the glass substrate. While the flaking was reduced by etching the glass prior to coating, the method was abandoned because uniforrn films could not be reproducibly obtained.This problem was eliminated by casting the polystyrene films on the surface of spherical stainless steel shot (250 pm diameter, 9 x m2 g-' surface area; manufactured by the Nuclear Metals Division of the Whittaker Corporation). The coating vessel, a 0.25 1. stainless steel beaker, was screwed to the driving element of a vortex generator. Excellent coatings were obtained by the following steps : (a) warm the beaker plus shot, (b) add a toluene solution of polystyrene, (c) remove the solvent by evaporation combined with vigorous agitation, ( d ) remove residual solvent by heating the coated shot at 95°C in uacuo. The weights of shot and polymer were chosen so as to give a nominally 2 pm coating thickness. A coating experiment typically yielded 150 g of coated shot, and this was a sufficient supply for all of the degradation experiments.For these experiments, 15 g of coated shot and 45 g of 2 mm diameter uncoated steel shot were placed in glass bottles (length = 9 cm, diam. = 4 cm) and rolled for various timed intervals at 275 r.p.m. on a roll mill at room temperature. A fresh jar and polymer sample was used for each degradation experiment. The large steel balls were added merely to accelerate the degradation process since excessive times were required when the coated shot was rolled alone. GEL PERMEATION CHROMATOGRAPHY After the required milling time, the polymer was recovered by the addition of tetrahydro- furan to the contents of the milling bottle. The resulting polymer solution was then analyzed using a Waters Associates Model 200 gel permeation chromatograph which had been pre- viously calibrated using standard polystyrene samples of known molecular weights.In brief, a gel permeation chromatograph separates polymer molecules on the basis of molecular size. The molecules elute from the column in order of decreasing size, and if (as in the present study) the effluent is monitored by a differential refractometer, the area of the chromatogram is proportional to the concentration of the injected polymer solution. The effluent is also measured by a syphon which places a tick mark-" elution count "-on the chromatogram for every passage of a fixed volume of carrier liquid. The elution volume (i.e., elution count number multiplied by syphon volume) can be related to the molecular weights of the eluting polymer molecules by the equation In (molecular weight) = a+ b (elution volume) where a and b are constants. RESULTS GEL PERMEATION CHROMATOGRAMS The various polystyrene samples were milled for a series of times up to 23 h.Since polymer degradation can be thermally induced,50 it was pertinent to check that the results obtained at extended milling times were not caused by possible milling- generated heating. As a test, therefore, a sample was milled for eight hours, the time being divided into alternate intervals of 30 min of milling and 30 min of resting. The result thus obtained matched that from a continuous eight hour milling experiment indicating an absence of any long-term thermal contributions to the degradation process. (The possibility still exists, however, that each violent impact of a milling ball causes local heating in the polymer film.) Fig.1,2 and 3 show the gel permeation chromatograms, for selected milling times,212 DEGRADATION OF POLYSTYRENE FILMS 0 9 2 12 5 FIG. 1.-NBS-706 gel permeation chromatograms for the degradation times noted (in hours). The ticks on the abscissa of each chromatogram mark, from left to right, the elution points for polystyrene molecules of molecular weights lo3, lo4, lo5 and lo6. 2 12 16 FIG. 2.-PS200 gel permeation chromatograms for the degradation times noted (in hours). The ticks on the abscissa of each chromatogram mark, from left to right, the elution points for polystyrene molecules of molecular weights lo3, lo4, lo5 and lo6.R . J . NASH A N D D.M. JACOBS 213 of the degradation experiments using the NBS-706, the PS200, and the PS860 poly- styrene samples : the PS670 sample gave results similar to those of the PS860 sample. In each case, the 23 h chromatogram closely resembled the 16 h chromatogram, the only apparent difference being a narrowing of the low molecular weight peak at the longer time. The chromatograms show that for each molecular weight sample examined, prolonged mechanical degradation produces the same limiting, low molecular weight species (M, z 8.4 x lo3). The generation of this latter species presumably accounts for the observed change in the polymer coating from a smooth to a powdery texture. 7 1 (d) 2 (b) 12 5 FIG. 3.-PS860 gel permeation chromatograms for the degradation times noted (in hours).The ticks on the abscissa of each chromatogram mark, from left to right, the elution points for polystyrene molecules of molecular weights lo3, lo4, lo5 and lo6. DISCUSSION PRIOR HYPOTHESES Even a superficial examination of fig. 1 , 2 and 3 reveals that many of the concepts of previous theoretical treatments of polymer degradation cannot apply to the present system. Thus, a commonly held hypothesis has been that polymer molecules, under mechanical stress, degrade by single scissions. The simplest view has the breakage point located at random along the polymer An alternative view has the break at the centre of the molecule,27 while a more refined treatment 30 holds that there is a range of possible positions, the probability for scission rising from zero near the ends of the molecule to a maximum at the centre.None of these views is in accord with the present data since, regardless of the location of the point of scission, single breaks will cause only a continuous shifting of the maximum in the molecular weight214 DEGRADATION OF POLYSTYRENE FILMS distribution to lower values. By contrast, fig. 1, 2 and 3 suggest that in fact mech- anical degradation occurs via concurrent multiple breakages. For example, in fig. 3, the molecular weights assignable to the three peak maxima in the chromatogram are in the approximate ratio of 70 : 7 : 1, and a traverse from fig. 3a to fig. 3fmust involve at least 10l8 bond breaks per gram of initial polymer. Considering the limited number of impacts possible during that time span, each impact must cause multiple scissions.That the chromatogram in fig. 3d is trimodal further implies a non- random distribution of scission sites. DEGRADATION MODEL BASED O N THE PRESENT DATA RATE OF DEGRADATION Possible parameters controlling the rate of degradation of thin polymer films include the time of degradation, the number and efficiency of impacts, the thickness of the film and the molecular weight of the polymer. In theory, a search for a possible molecular weight dependence would best be made using the data from the narrow distribution samples, PS860, PS670 and PS200. However, this approach is precluded by the variability in the coating thickness from sample to sample. The problem can be eliminated by use of the NBS-706 data since the molecular weight distribution of this sample must contain contributions from a wide spectrum of molecular species within a single film.As an initial step, consider fig. 4, which is a plot against degrad- ation time of the relative curve height (Nt/No where Nt and No are the curve heights at time t and time zero) at various elution counts (i.e., at various molecular weights). 0 : 0 4 degradation time/h FIG. 4.-NBS-706 degradation data for the elution counts noted. Since the heights are those taken directly from the chromatograms, they may contain contributions from instrumental effects and from overlapping neighbouring peaks : there is no ready method to correct for these contributions which, however, should be minor for the highest elution counts. WhiIe the plots for counts above 27 (equivalent to a molecular weight of M 1.4 x lo5) show a net increase with time caused by the arrival of fragments from the degradation of higher molecular weight species, the plots for counts below 27 apparently represent pure degradation.The relative peak heights at these latter counts decay exponentially with time, the rate decreasing withR. J . NASH AND D . M . JACOBS 215 decreasing molecular weight. At this point it will be helpful to develop a model to aid in the rate analysis. The model must explain the following observations : (a) The rate of degradation varies exponentially with time and is a decreasing function of molecular weight. (b) The number of bond scissions per impact is large. (c) The bond scissions per molecule are multiple and non-random. Since the same fixed coating weight was used for each of the various polystyrene films, the films of the largest molecular weight species will contain the fewest number of molecules.This fact alone is sufficient to cause the rate of degradation to be a decreasing function of molecular weight. Thus, consider a film of side lo and depth d. If each successful milling event is assumed to cause degradation within a zone having orthogonal dimensions b,, by, b,, then where P is the probability that a polymer molecule of size I will be at least partially within the degradation zone. The rate of degradation will be given by and hence dN/dt = -MPNt, Nt = No exp (-MPt), where M is the number of impacts per unit time having sufficient energy to cause degradation, No is the initial number of molecules and Nt is the number of molecules remaining at time t.Under the experimental conditions of the present study, the degradation zone appears to extend throughout the entire film depth, and thus the last term in eqn (1) can be taken as unity. If the zone dimension is assumed to be b in both the x and y directions then (2) Fig. 5 , the experimental data plotted according to eqn (2), yields (3) Nt = No exp [ - (Mt/Z~)(Z+b)2]. Nt = No exp [ -9.242 x 10-12t[(0.67M3) + 1920l2], 400 8 0 0 1200 (molecular weight)+ F I ~ . 5.-NBS-706 degradation data.216 DEGRADATION OF POLYSTYRENE FILMS where M is the polymer molecular weight, t is the degradation time in seconds, and the factor 0.67 converts M3 to molecular dimensions in A.Thus, with the above assumptions, the degradation zone appears symmetrical since the film thickness is 2000A and the deduced zone size b is 1920A. Eqn (3) correctly predicts the rate of degradation not only for the selected components within the NBS-706 chromato- gram, but also for the PS860, PS670 and PS200 chromatograms. Fig. 6 shows the NBS-706 data and eqn (3) plotted so as to indicate the dependence of degradation on time and molecular size. -2.0 - -2.5 , 1 I 0 5 0 0 1000 1 5 0 0 (molecular weight)J FIG. 6.-The points are NBS-706 data ; the lines are those calculated using eqn (3), (see text). The numbers are the degradation times in hours. MODE OF DEGRADATION The chromatograms in fig. 1,2 and 3 show that selective bond scission must occur.Since bond breakage will occur at points least able to relieve the applied stresses, the data suggests that these “ weak ” points must be regularly distributed along the poly- mer molecule. A possible identification of these points can be made by noting that in an amorphous polymer film the molecules, being highly intertwined, will be most constrained in the regions of entanglement.” These regions have indeed been postulated as the sites for bond breakage in the shear-induced degradation of polymer While the mean distance between entanglements will not be a function of polymer molecular weight, the number of entanglements per molecule will increase with molecular weight. Thus, while large molecules will be able to relieve applied stresses by bond breakage at only a portion of the many regions of entanglement, smaller molecules under the same conditions will be forced to break at the majority of the regions.This view could account for the bimodal splitting of the degradation products from samples PS860 and PS670 and the single mode from PS200. An alternate, though complementary explanation can also be given from a probability standpoint. Thus, the probability for a molecule to be completely hit will be a function of the size of the molecule and of the impact zone. For the latter, the two extreme cases will be (a) the impact occurs at a point, (b) the impact zone greatly exceeds the dimensions of any of the impacted molecules. Case (a) will generateR . J . NASH AND D . M. JACOBS 217 large degradation fragments regardless of molecular size, while case (b) will cause all molecules to degrade directly to small fragments, provided that sufficient energy is available.Between these two cases there must be a zone size large enough com- pletely to degrade PS200, but small enough only partially to degrade PS860. Addi- tionally, since in any actual impact zone the stresses will vary from point to point, this variation will produce a range of modes of bond breakages. As molecular size decreases, the entanglements per molecule will diminish until, at some characteristic size, the molecules no longer drastically hinder one another. Molecules below this size will hence be expected to resist mechanical degradation provided that the applied stresses are not severe enough to produce bond breakage by some mechanism other than entanglement.The limiting molecular weight species observed in the present study may, therefore, be viewed as being in this sufficiently disentangled state, with the molecular weight distribution reflecting the distribution of inter-tangle spacings. It should be noted that many properties of polymers reflect the existence of molecular entanglement and the concept of a " critical entanglement molecular weight " is well established. COMPUTER SIMULATION By combining the results of the rate and mode examinations, an overall picture of the degradation process can be obtained. Fig. 7, a plot against time of the areas of the three component peaks in the chromatograms of sample PS860, suggests that degradation time/h weight peak 0 ; intermediate molecular weight A ; low molecular weight peak U.FIG. 7.-PS860 degradation data : variation of component peak areas with time. High molecular the degradation process can be represented as a series of consecutive, irreversible reactions. If the high, intermediate and low molecular weight peaks are labelled A, B, and C, then two possible pathways appear possible : (a) A+B+C (b) A+B C w For each pathway, expressions for the concentration of each species as a function of218 DEGRADATION OF POLYSTYRENE FILMS time can be derived by an application of Laplace-Carlson transforms to the compon- ent rate equations.52 Unfortunately, by an appropriate choice of rate constants, the experimental results can be correctly described by both pathways. Therefore, a kinetic analysis based on the degradation data of one molecular species alone cannot reveal which is the correct pathway.However, if the rates calculated using eqn (3) are inserted into the kinetic analysis, then (b) emerges as the likely mode. To illus- trate this point, the following computer simulation of the gel permeation chromato- grams was made for both pathways. For the simulation of the PS860 data, peaks A, B and C are approximated by three Gaussian curves of fixed base widths in the proportion 5 : 8 : 8. Peak C is assumed to be comprised of non-degradable molecules : the small amount of degradation observed experimentally is simulated by a gradual shift in the peak maxima to lower values as time progresses. To calculate the amount of degradation occurring for peaks A and B, eqn (3) was rewritten in the following form : f= 1 -exp [-9.242 x 10-12[(0.67M9)+ 1920l2] (4) wherefis the fraction of molecules of molecular weight M which degrade in unit time.Rather than calculate the extent of degradation for selected values of M within peaks A and B, only the value of M at the peak maximum was considered, the 16 FIG. 8.-Computer simulation, using pathway (a), (see text), of PS860 degradation chromatograms for the times noted (in hours). The ticks on the abscissa of each chromatogram mark, from left to right, the elution points for polystyrene molecules of molecular weights lo3, lo4, lo5 and lo6. entire distribution being calculated on the basis of the behaviour of this single value. This tacitly assumes that the molecules within a distribution all degrade similarly : the experimental results suggest that this is a reasonable assumption.219 At each pass through the program, (simulating one hour of degradation), the previous height values for peaks A and B are multiplied by the required values off.For peak B, the resulting " fragment " is added to peak C ; for peak A, the " frag- ment " is divided, after scaling to allow for the differences in peak base widths (the sum of the peak areas, being proportional to polymer concentration, must be main- tained constant to simulate mass balance), between peaks B and C according to a ratio Y. Experimentally, the area under the total chromatogram decreases slightly with time ( z 5 % loss after 23 h of degradation) prcsumably because of production of fragments below the resolution of the chromatograph (i.e., < 1000 molecular weight). To simulate this loss, a small amount of the " fragment " from A is removed from the computed chromatogram.Following each degradation step, the program calculates the curve envelope for the sum of the three component peaks. The program was run in a time-shared interactive mode on a Xerox Sigma 7 computer, the results being displayed on a Hewlett-Packard 7200A graphic plotter R . J . NASH AND D. M. JACOBS 2 FIG. 9.-Computer simulation, using pathway (b), (see text), of PS860 degradation chromatograms for the times noted (in hours). The ticks on the abscissa of each chromatogram mark, from left to right, the elution points for polystyrene molecules of molecular weights lo3, lo4, lo5 and lo6.driven by an ASR-33 Teletype. In this way, the effect of various values of Y could be readily visualized. To simulate pathway (a), Y was taken as zero. The computed chromatograms, fig. 8, using this value do not match the experimental chromatograms, in particular the computed peak B grows too rapidly. This occurs because eqn (4) predicts that the rate of degradation of peak A will exceed that of peak B. The closest simulation, fig. 9, was achieved using pathway (b) with the degradation products being distributed 65 % to peak B, 31.5 % to peak C and 3.5 % as non-detectable fragments. The simulation thus suggests that fragmentation of the PS860 polystyrene mole- cules occur via two parallel, consecutive paths, one involving on the average w7220 DEGRADATION OF POLYSTYRENE FILMS breaks per molecule, the other ~ 7 0 breaks.Presumably, the ratio of medium to small fragments is related to the dimensions of the impacted molecule and the impact zone, and the stress distribution within the zone. The present results are not compre- hensive enough to reveal the form of such a relationship : experiments on polymer films of various thicknesses or different chemical constitution might offer a way to explore this point. Throughout this work, the polymer samples have been referred to as “thin films ”, however, a better description would be “ thin coatings ”. As such they are doubtless influenced by the physical properties of the substrate : pre- sumably a change of substrate could cause a change in the degradation process.From a practical point of view it should be noted that solids are often given protective coatings of thin polymer films. The present study shows that the deterioration of such coatings by mild mechanical forces can occur not only by macroscopic but also by molecular degradation. R. S. Porter, M. J. R. Cantow and J. F. Johnson, J. Polymer Sci. C, 1967, 16, 1. R. M. Thomas, J. C. Zimmer, L. B. Turner, R. Rosen and P. K. Frolich, Znd. Eng. Chem., 1940, 32, 299. N. K. Baramboim, Mechanochemistry of Polymers (Maclaren and Sons Ltd., London, 1964). H. H. G. Jellinek, Degradation of Vinyl Polymers (Academic Press, New York, 1955). N. Grassie, Chemistry of High Polymer Degradation Processes (Interscience Publishers, New York, 1956). N. K. Baramboim, Zhur.Fiz. Khim., 1958,32, 806. H. W. W. Brett and H. H. G. Jellinek, J. Polymer Sci., 1956, 21, 535. R. J. Ceresa and W. F. Watson, J. Appl. Polymer Sci., 1959, 1, 101. G. Gooberman and J. Lamb, J. Polymer Sci., 1960, 42, 35. lo H. Grohn, K. Bischof, M. Loesche and K. Moeckel, PIaste Kautschuk, 1961, 8, 593. l 1 H. Grohn and K. Bischof, Wiss. Z. Tech. Hochsch. Chem. Leuna-Merseburg, 196112, 4, 153. l2 H. H. G. Jellinek and G. White, J. Polymer Sci., 1951, 6, 757. l3 H. H. G. Jellinek and G. White, J. Polymer Sci., 1951, 7, 21. l4 H. H. G. Jellinek and G. White, J. Polymer Sci., 1951, 7, 33. l 5 W. R. Johnson and C. C. Price, J. Polymer Sci., 1960, 45, 217. l 6 K. Mackenzie and A. E. Jemmett, Wear, 1971, 17, 389. M. A. K. Mostafa, J. Polymer Sci., 1958, 28, 499.A. Nakano and Y. Minoura, J. Appl. Polymer Sci., 1971, 15,927. l9 R. S. Porter, M. J. R. Cantow and J. F. Johnson, J. Appl. Phys., 1964, 35, 15. 2o R. S. Porter and J. F. Johnson, J. Appl. Phys., 1964, 35, 3149. 21 R. S. Porter, M. J. R. Cantow and J. F. Johnson, J. Appl. Polymer Sci., 1967, 11, 335. 22 R. S. Porter, M. J. R. Cantow and J. F. Johnson, J. Polymer Sci. C, 1967, 16, 1. 23 R. S. Porter, M. J. R. Cantow and J. F. Johnson, Polymer, 1967, 8, 87. 24 H. Staudinger and W. Heuer, Ber., 1934, 67, 1159. 25 H. Staudinger and E. Dreher, Ber., 1936, 69, 1091. 26 P. E. M. Allen, G. M. Burnett, G. W. Hastings, H. W. Melville and D. W. OvenaI1, J. Polymer 27 F. Bueche, J. Appl. Polymer Sci., 1960, 4, 101. 2 8 G. Gooberman, J. Polymer Sci., 1960, 42,25. G. Gooberman, J. Polymer Sci., 1960, 47, 229. 30 G. J. Heymach and D. E. Jost, .I. Polymer Sci. C, 1958, 25, 145. 31 H. H. G. Jellinek and G. White, J. Polymer Sci., 1951, 6, 745. 32 H. H. G. Jellinek, J. Polymer Sci., 1956, 22, 149. 33 H. H. G. Jellinek, J. Polymer Sci., 1959, 37, 485. 34 H. H. G. Jellinek, J. Polymer Sci., 1962, 62, 281. 35 J. Malac, J. Polymer Sci. C, 1971, 33, 223. 36 M. A. K. Mostafa, J. Polymer Sci., 1956, 22, 535. 37 M. A. K. Mostafa, J. Polymer Sci., 1958, 27, 473. 38 M. A. K. Mostafa, J. Polymer Sci., 1958, 28, 499. 39 M. A. K. Mostafa, J. Polymer Sci., 1958, 28, 519. 40 M. A. K. Mostafa, J. Polymer Sci., 1958, 33, 295. 41 M. A. K. Mostafa, J. Polymer Sci., 1958, 33, 311. 42 M. A. K. Mostafa, J. Polymer Sci., 1958, 33, 323. Sci., 1958, 33, 213.R . J . N A S H A N D D . M . JACOBS 22 1 43 M. Okkuama and T. Hirose, J. Appl. Polymer Sci., 1963, 7, 591. 44 D. W. Ovenhall, G. W. Hastings and P. E. M. Allen, J. Polymer Sci., 1958, 33, 207. 45 D. W. Ovenhall, J. Polymer Sci., 1960, 42, 455. 46 N. Sata and M. Okuyama, 2. Elektrochem., 1954,58, 196. 47 G. Schmid and 0. Rommel, 2. phys. Chem., l939,185A, 97. 48 G. Schmid and 0. Rommel, 2. Elektrochem., 1939,45,659. 49 H. Determann, Gel Chromatography (Springer-Verlag, New York, 1969). 5 0 H. H. G. Jellinek, J. Polymer Sci., 1948, 3, 850. 51 R. S. Porter and J. F. Johnson, Chem. Rev., 1966, 66, 1. s2 N. M. Rodiguin and E. N. Rodiguina, Consecutive Chemical Reactions (D. Van Nostrand Company, Inc., Princeton, 1964).

ISSN:0370-9302    

DOI:10.1039/S19720200210

出版商:RSC    

年代:1972

数据来源: RSC

摘要:

GENERAL DISCUSSION Dr. H. Wilman (Imperial College, London) said: There is presumably much uncertainty in the estimation of the current density in the electrodeposition on to very small convex tips of metal wires used in ion-microscope experiments. I wonder if Farr can state what was the order of magnitude of the “ high ” and “ low ” current densities used by Rendulic and Muller (his ref. (21)), to whom he refers, regarding epitaxy in cathodic processes? It is also likely that the tendency to cracking of epitaxial layers strained by some misfit would be increased when the deposits are made on such substrates of very small radius of curvature, and also subjected to the high field stresses in the ion microscope, as Farr indicates. I find it difficult to reconcile fig. 2(a) with the statement in the legend that “ a (100) region lies centrally, .. .”. Although there is some degree of symmetry about a horizontal line through the centre of the photograph, and also about a vertical line through it, there does not seem to be 4-fold symmetry of the pattern features about the central region of the photograph, to give square dispositions of similar 4-fold repeated features. Can the authors clarify this? Nanis and Javet (their ref. (20)) show a clear pattern of 4-fold symmetry from an iridium tip in a (100) orientation, and it differs clearly from this. I would also emphasize that the “ liquid-like ” coalescence observed by Pashley, Matthews and others in the growth of island crystal nuclei during condensation of vapour in vacuum, was under conditions of high atomic mobility on substrates in vacuum at considerably raised temperatures-about 400°C.It is unlikely that such high mobilities and coalescence effects will occur in electrodeposition except at com- parable temperatures in non-aqueous electrolytes (molten salts) or at very high current densities. With respect to the use by Farr and Rowe of the term “pseudomorphism”, I would stress that this term is not normally confined to cases of strict conformity of lattice spacings of the deposit crystal with those of the substrate crystal at the interface -if indeed any such cases of exact correspondence are known in practice. “ Pseudo- morphism” denotes that the deposit atoms, under the forces exerted by the 2- dimensional array of substrate surface atoms (with some contribution from lower layers), must take up positions tending towards the consecutive potential troughs of the substrate surface, but not in general exactly at the minima, since the deposit atoms are also acted on by the forces of cohesion with other laterally adjacent deposit atoms.It therefore seems inappropriate to say that “ pseudomorphism is, however, no longer regarded as necessary for epitaxy ”. I have also a query on the results under the heading Electrochemical Experiments in their paper, where evidence was quoted for deposition of copper at electrode potentials anodic to the equilibrium potential on some substrates. Were the metal substrates in these cases polycrystalline? If so, their surfaces would in general be uneven and consist of elements of different crystallographic forms on different crystal grains, on which presumably the equilibrium potentials would be also different, so that deposition would occur on some surface elements while still not on others.Dr. 6. P. G. Farr (Birmingham University) said: Most potential sweep experiments on electrodeposition so far reported have involved polycrystalline cathodes. There is general agreement (see our ref. (40)) with Wilman’s suggestion of initially localized 222GENERAL DISCUSSION 223 deposits on particular surface elements. Our own work on single crystals suggests that even here a variety of initial growth sites is normally available (fig. 1). One may question the use of the term “equilibrium potential” at the early stages of formation of the solid-solid interface. In our review we refer to the equilibrium potential as, for example, that of the Cubulk/Cu2+ electrode and we accept that an “ under-voltage ” is an indication of the adsorption energy between depositing atoms and a foreign substrate. One might expect a spread in under-voltage across a sub- strate surface, a dependence on surface preparation, and some variation as growth proceeds beyond a monolayer until the deposit is indistinguishable from bulk metal.Dickson, Jacobs and Pashley (ref. (27) of our paper) interpret the similarity in the early stages of electrolytic growth of gold on (111) silver substrates to growth by evaporation in terms of a surface migration of gold over silver and over gold islands. Farr and Loong have not observed liquid-like coalescence of copper or nickel nuclei on Ag.These metals grow rather differently from each other, copper particularly at discrete nuclei and nickel in continuous layers on the silver substrate with typical misfit dislocations. (This work will be published in detail elsewhere.) In the first part of our paper we mention some deficiencies in the characterization of adspecies from the electrode kinetics of aqueous systems. A further discussion of the role of adspecies may be obtained from ref. (2), (3) in our paper. In some early work, Gerischer found it necessary to ascribe 30-40 % ionic character to silver adspecies on silver and suggested that this was due to a strong interaction between adsorbed atoms with polar water molecules.1 suggest that at very high current densities the growth process is governed by the rate of charge transfer rather than the rate of nucleation or the surface diffusivity of an adspecies.The uncertainty in estimating a current density in electrodepositing on to a field ion microscope tip was recognized by Rendulic and Muller (ref. (21) of our paper) who estimated a mean current density i, between 0.1 A/cm2 (low) and 10 A/cm2 (high) at 50°C. They mention that the epitaxial layer of Pt on Ir ruptured easily at a thickness of 100-200 A ; the corresponding non-epitaxial layer on W was much more stable. However, the cracks within the epitaxial Pt were not thought to be the result of field stress. The apparent difference between our field-ion image (our fig. 2a) of a (100) oriented iridium tip and that shown by Nanis and Javet (ref. (20), fig.4) results from a rotation of about 45” in the plates as printed and, more important, a disparity in the dimensions of the tips. Our tip was smaller than that of Nanis and Javet and so there are fewer rings, or atomic ledges, between the central (002) plane and the next significant four (113) poles. In both cases the (111) planes referred to are at the peripheries of the images and in ours only one (1 11) plane is imaged, i.e., at the right- hand side of fig. 2a. Both illustrations show the “ zone decoration ” characteristic of the region around the (002) plane, i.e., bright spots on the images running from (002) towards { 113). It will be appreciated that in the field ion microscope an image of a 3-dimensional surface is projected on to a 2-dimensional screen and this projec- tion may not be simple; furthermore, our image is truncated compared witn that of Nanis and Javet, because our tip was asymmetric.Some authors have adopted a more restricted use of the term “ pseudomorphism ” than does Wilman. In our paper we have followed Pashley (ref. (25) of our paper), who accepts that pseudomorphism prevails where lattice misfits can be accommodated by elastic strain alone. However, this allows the “ basal plane pseudomorphism ” Various authors H. Gerischer and H. Fischer, 2. Elektrochem., 1957, 61, 1159. see, e.g., K. Vetter in Electrochemical Kinetics, Academic Press, 1967, p. 325.224 GENERAL DISCUSSION of Finch and Quarrell to be a rather frequent phenomenon.Basal plane pseudo- morphism would be expected to be lost more rapidly as a deposit thickened if the lattice misfit were considerable or if there were differences in crystal symmetry. For Ni (100) on Ag (100) where there is a misfit of some 15 %, Loong and Farr have observed increasing numbers of typical misfit dislocations in thickening films and suggest that partial coherence may occur in the initial stages of growth. For Ni on Ag it would not be expected (see ref. (25) of our paper) that pseudomorphism in Pashley's sense would be retained beyond the basal plane. It is thought that most solid precipitates within bulk metal have some coherence with the matrix.2 Unusual crystal forms and particular precipitate morphologies may result. Clearly, epitaxy may often be a result of basal plane pseudomorphism.Occasionally it may be that where deposition conditions or crystallographic factors do not permit this, a close-packed island of depositing atoms may form a 2-dimen- sional nucleus, so that a highly textured deposit may grow without there having been initial coherence. This may explain the [ l l l ] oriented Pt crystals obtained by Rendulic and Muller (ref. (21) of our paper) on Ir. We thank Wilman for reminding us of the review by K. Lawless in Physics of Thin Films, 1967, 4 ; that by K. L. Chopra (in Thin Film Phenomena, McGraw-Hill (1969) p. 193) is also relevant. Prof. M . W . Roberts (Bradford University) said : With reference to Wilman's remark concerning the growth model, there is now evidence from Bassett's field-ion work that dimers, trimers and " long chain " monomers are involved in growth of one metal on another.The predominance of any one type of growth nucleus (if I recall correctly) depended on the nature of the adatom, and also possibly on the im- pingement rate (i.e. , Wilman's current density). Dr. H . Wilman (Imperial College, London) said : Can O'Sullivan and Oxley state the mean thickness of their electrodeposited films under discussion? The form of expression used to define the epitaxial orientation of the b.c.c. alloy deposits seems unnecessarily complicated and confusing. This expression is indeed the same as is stated by Jones who, however, made it clear that $ is the azimuthal angle between the two sets of orthogonal b.c.c. (110) orientations on the Cu (001) substrate face (e.g., the angle between the radii to the innermost strong pairs of arcs in patterns such as fig.(b) of plate 1 of the paper of O'Sullivan and Oxley). Jones also stated that this angle was observed to lie between 15 and 18", so that the value of f (19'30'- $)/Z was only between 0 and about 2.5". It thus would be much clearer to define these epitaxial orientations simply as, in the mean : (1 10)[1T1] alloy 11 (100)[011] or [Oil] Cu in some deposits, but in others differing from this by up to about 2.5" in azimuthal orientation, in both cases there being a few degrees spread of azimuthal orientation about the mean, as shown by arcing in the pattern. In such electron diffraction patterns obtained by transmission through isolated thin films there is often some ambiguity as to the interpretation.An arcing of the diffractions can arise as a result of cylindrical curvatures (or bending) of the film. An appreciable inclination of the mean film plane away from the setting normal to the electron beam could cause a small change in the observed angle $ between the radii to the arc centres concerned, like the above small deviations, even if the ideal G. I. Finch and A. G . Quarrel], Proc. Roy. SOC. A, 1933, 141, 398. see, e.g., A. Kelly and R. B. Nicholson, Prog. Materials Sci., 1963, 10, 151 ; also the A. S. M. Seminar on Phase Transformations, ed. H. I. Aaronsen (1968)).GENERAL DISCUSSION 225 epitaxial b.c.c. (1 10) orientation stated above were strictly followed. The four-fold symmetry of the arc pattern would, however, be slightly impaired in this case and this should be an indication that the film had become tilted away from the normal setting.Similarly, such an inclination from the normal setting could suggest azimuthal disorientations or displacements of the kind concluded here, even if the disorientation were really about an axis in the film plane, as we have observed in some cases by grazing-incidence electron diffraction of films in situ on their substrates (see my paper with Verma, this Discussion). I am also not clear as to the reference to the case of " when 11/ = 90". . .". Did the authors in fact observe such (1 10) orientations in their electrodeposited b.c.c. alloys? As a further comment, ref. (2) is evidently incorrect and should be presum- ably Phil.Mag., 1965, 11, 993. Dr. J, M. O'Sullivan (Stainless Equipment Co.) said: The mean thickness of alloy foils used in our experiments was lOOOA. This was measured by application of Faraday's laws of electrolysis and direct weighing of cathodes. We are grateful to Wilman for his suggestion of an alternative and simpler method of expressing the interface crystallography between b.c.c. deposits and the copper substrate. His remarks about arcing and the effects of inclination of the mean film plane from a setting normal to the electron beam are of general application to transmission diffraction work, and it is salutary to be reminded of them. " When $ = 90" . . ." should have read ". . . when 11/ = 19'30'. This has been corrected in the present paper. Ref. (2) is indeed incorrect.Prof. W. C. Wake (City University) said: I have a few comments on the methods used by Roberts et al. for characterizing the deposited and bombarded surfaces. The use of Zisman's critical surface tension can be pushed too far by way of inter- pretation. For example, it is usually believed that yc will decrease with increasing density of packing but recent work at Brunel University has shown this is not so.' Using highly branched compounds of the type CH3 I CH3-C-CH3 CH3 1 I CH3-C-CH and rather simpler fluorinated derivatives, it was found that yc decreased with de- creasing density. The point is that under electron bombardment highly branched structures are likely and yc will be accordingly influenced. M. A. Parrish, unpublished work for M. Tech (1970).226 GENERAL DISCUSSION Prof.M. W. Roberts (Bradford University) said: Thank you for the information regarding the work at Brunel University. I agree that the correlation between yc and molecular environment (fig. 2) should be used cautiously. Nevertheless, if by controlling the chemistry we can predict surface characteristics, this is a step in the right direction. The present paper illustrates how control of some parameters (e.g. , the nature of the monomer, the substrate temperature, the nature of negative ions) leads to distinct surface characteristics as reflected by yc. It is also worth mentioning that the ll term (the spreading pressure) can not, as is often assumed, be equated to zero. The next stage is to obtain direct information on the molecular nature of the sur- face and we have had recently some success in this direction.Films formed from C,F,(g) and a substrate temperature of - 300 K have been shown to contain little, if any, fluorine. The films are therefore essentially composed of a carbon lattice and not built around CF as suggested in the paper. The electron spectroscopy informa- tion l is compatible with our mass-spectrometric studies in that CF ions were the pre- dominant gas phase ions observed. Prof. W. C. Wake (City Uiziversity) said: If the net result is carbon deposition then it is known that the normal alkanes used for determining yc show a mobile and not a fixed site (Langmuir) type adsorption. This I demonstrated some years ago by calculation of the appropriate entropies of adsorption.In this case the spreading pressure term in the Young equation would be appreciable and must influence the contact angle. Prof. R. Sh. Mikhail (Ain Shams University, Egypt and University of Salford) said: There is an unfortunate use in nomenclature when the authors used the terms “ high energy ” and “ low energy ” surfaces. The surface may have a weak interaction with one adsorbate and a strong interaction with another adsorbate, and it seems quite inappropriate to describe the surface itself to be low energy or high energy when this energy is a characteristic of the bond and not the surface. Prof. M. W. Roberts (Bradford University) said: Zisman drew a distinction between different solid surfaces on the basis of the observed yc (the critical surface tension, i.e., the value of y l v at cos 8 3 1 .O) using a series of non-polar liquids.Values below 100dyncm-l he referred to as “low energy ” surfaces. First, it is clearly an arbitrary definition. Secondly, as we have mentioned in our paper and else- where, the measured yc reflects not only the molecular nature of the substrate but how that substrate is perturbed by adsorption from the liquid phase. Furthermore, on this basis we put forward some tentative suggestions as to the orientation of the adsorbed molecules on the polymer film surfaces. It should be emphasized that yc should not be equated with ys the surface energy of the solid. Prof. W. C. Wake (City University) said: With regard to the paper by Nash et al., degradation by mechanical or ultrasonic means from high molecular weight to a limiting low molecular weight has been identified in morphological terms by Schoon and his fellow workers at Wurzburg. Using electron microscopy and a variety of ’ M. W. Roberts and C. R. Brundle, unpublished work. W. J. Murphy, M. W. Roberts and J. R. H. Ross, J.C.S., Faraday Trans. I, 1972, 68, 1190. Th. G. F. Schoon and R. Kretschmer, KoIIoidZ. 2. Polymer, 1964,197,51 ; Th. G. F. Schoon, British Polymer J,, 1970, 2, 86 ; Th. G. F. Schoon and G . Rieber, Angew. Makromol. Chem., 1971, 15, 263.GENERAL DISCUSSION 227 sample preparative techniques, they have shown that bulk polymer contains structures not apparent in solution studies, that there are several stages of ordered structure, and that a fundamental structural unit exists the size of which corresponds roughly to the limiting low molecular weight identified in the work of Nash et al. The views of Schoon et al. differ from those of Nash et al. in that they postulate a structure which is broken into pre-existing units rather than a mechanism determining a zone size. A recent paper by P. H. Geil of the Case Institute of Cleveland comes to similar conclusions with respect to polyvinyl chloride.

ISSN:0370-9302    

DOI:10.1039/S19720200222

出版商:RSC    

年代:1972

数据来源: RSC

摘要:

AUTHOR INDEX * Allen, G., 127. Allen, K. W., 38. Alsalim, K. S., 38. Bader, M. G., 165. Barnett, S. V., 144. Bessell, T., 137. Biggs, W. D., 117. Bishop, P. T., 124 Blundell, D. J., 127. Bowden, M. J., 127. Bowyer, W. H., 165. Briscoe, B. J., 7, 56, 57, 58, 59. Cameron, A,, 26. Eley, D. D., 58, 60, 124, 126. Farr, J. P. G., 177, 222. Found, M. S., 77. Frost, S., 198. Harris, S. J., 144, 175. Hull, D., 137. Hutchinson, F. G., 127. Jacobs, D. M., 210. Jeffs, G. M., 127. Leveson, R., 175. Lin, D. S., 46. Linford, R. G., 57, 61, 174. McGarry, J?. J., 90. Mandell, J. F., 90. Mikhail, R. Sh., 123, 175, 226. Morley, J. G., 109. Murphy, W. J., 198. Nash, R. J., 210. O’Sullivan, J. M., 194, 225. Owen, M. 3., 77, 124, 126. Oxley, D. P., 194. Proctor, B. A., 63. 123. Roberts, M. MI., 57, 174, 198, 224, 226. ROSS, J. R. H., 198. Rowe, G. W., 177. Shortall, J. B., 137, 159. Sutcliffe, M. J., 26. Tabor, D., 7, 56, 57, 58, 59. Verma, S. K., 185. Wake, W. C., 38, 60, 123, 126, 174, 225, 226. Weaver, C., 18. Weaver, J. V., 44. Wilman, H., 46, 56, 59, 60, 62, 185, 222, 224. Wood, J. H., 198. Yip, H. W. C., 159. Young, J. R., 58. * References in heavy type indicate papers submitted for discussion. 228

ISSN:0370-9302    

DOI:10.1039/S19720200228

出版商:RSC    

年代:1972

数据来源: RSC