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Surface strengthening in aluminium single crystals coated with evaporated films
SURFACE STRENGTHENING IN ALUMINIUM SINGLE CRYSTALS COATED WITH EVAPORATED FILMS TAKAYUKI
and OSXMU
TAKASC’GI
IZUMI
The Research Institute for Iron. Steel and Other Metals. Tohoku University, Sendai. Japan (Receirrti
16 frmr
1975: in
rerised_forrn 1 April 1976)
.&tract-Using evaporated silver. copper and nickel films covering a broad spectrum of shear moduli and film strength, the effect of surface coatings on the deformation of aluminium single crystals was studied. From the anaiyses of stress-strain behaviour, surface marking observation, slip length measurement, disiocation density measurement and X-ray diffraction, a distinct difference was clearly observed in the surface strengthening in accordance with the coating materials. The surface strengthening by coating with silver and copper. which revealed a slight effect. was explained by the short range intetaction, while the strengthening by nickel coating was interpreted by adding more of the long-range interaction to the short-range one. The short-range interaction component is thought to be due to an image force and pile-up force. On the other hand. the long-range interaction component is thought to be due to the plastic constraint which is correlated with the differences in shear modulus and strength between the substrate and the Film. and to the deformation mode of the film itself. RPsum&--On a Ctudii l’effet d’un revetement superficiel de monocristaux d’aluminium sur leur dtformation, en utilisant des films evapores d’argent, de cuivre et de nickel qui presentaient une grande gamme de modules de cisaillement et de resistances. L’analyse des courbes contrainte-deformation, des observations de marquages superficiels. des mesures de la longueur glissement et de la densite de dislocations. et des diffractions de rayons X ont permis de mettre en evidence une nette difference dans Ie durcissement selon le materiau de revttement. On a expliqui le petit effet de durcissement supetficiel dd a un revetement d’argent et de cuivre par l’interaction Q courte distance, alors qu’on a interpret6 le durcissement par un revetement de nickel en ajoutant une interaction a longue distance a l’interaction a courte distance. On pense que la composante &interaction a courte distance provient dune force-image et dune force d’empilement. On pense par contre que la composante a longue distance de l’interaction provient de la contrainte plastique, lice 6 des differences de module de cisaillement et de resistance entre le substrat et le film, et au mode de deformation du film lui-meme. Zu~mmenfa~~g-per EinfluB van O~rfl~chenschichten auf das Ve~ormungsverhalten von Aluminium Einkristallen wurde mittels aufgedampften Silber-. Kupfer- oder Nickelfilmen untersucht. Es wurden in Abhlngigkeit vom Uberzugsmaterial ausgeprlgte Unterschiede in der Oberflachenhlrtung beobachtet, angezeigt durch das Spannungs-Dehnungsverhalten, die Oberflachenbeobachtungen. die L2ngen der Gleitlinien, die Versetzungsdichte und die Riintgenbeugung. Der kleine EinfluB der Beschichtung mit Silber oder Gold wurde durch eine kunreichende Wechselwirkung erklart. der EinRuB der Nickelbeschichtung wurde mit einer Zunahme der langreichweitigen Wechselwirkung gegeniiber der kurzreichenden interpretiert. Die kurzreichende Wechselwirkung wird einer Bildkraft und einer Kraft herriihrend van Aufstauungen zugeschrieben. Andererseits wird angenommen. da13 die langreichweitige Wechselw~rkungskomponente von einer plastischen Einschrankung herriihrr, die mit dem Unterschied in Schermodul und Festigkeit zwischen Substrat und Film und der Ve~o~ungsa~ der Schicht selbst zusammenhIngt.
1. IXTRODUCTION
It is well known that surface conditions can influence plastic deformation of metallic single crystals. The strengthening produced by the presence of a surface film is much greater than what could be expected simply by the strength contribution due to the film. The reader is referred to a number of review articles [l-6]. The mechanisms which have been proposed to explain the surface strengthening effects fall into five
categories: (1) the surface dislocation source blocking mechanism [73, (2) the image force mechanism [S], (3) the surface damage zone mechanism [9], (4) the surface barrier mechanism against the egression of dislocations [lo] and (5) the plastic constraint mechanism (compatibility mechanism) [l 11. However none have achieved universal acceptance. It has been suggested[ll, 121 that strengthening could not be attributed
the
surface
to a single mechanism, but to several mechanisms. Takasugi and
I II37
110s
TAKASUGI
AND
IZUMI:
SURFACE STRENGTHESING
Izumi [12] proposed that the surface strengthening in aluminium single crystals coated with electro-deposited nickel film consists of long- and the short-range interaction stress components. The former could be explained by the plastic constraint theory [13]. The latter could be explained by the image force mechanism. It is expected that both interaction components will be altered by various types of films. The purpose of the present work is to explain these interaction components more clearly using different materials as evaporated films, with a range of shear moduli and film strength. 2. EXPERIMENTAL
WITH EVAPORATED
FILMS
lated with hardening mechanism. The procedure of the measurement was as follows. The plated and unplated specimens were strained in tension to a desired strain then unloaded. Both specimens were then electro-polished (here, surface films of plated specimen were removed). Slip line lengths were measured at the middle portion of the wide surface of specimens. Etch-pitting and X-ray observations of dislocation were carried out before and after the deformation both on the unplated and plated specimens. In the case of plated specimens, dislocation density was measured after removing the films by electro-polishing [12].
PROCEDURES
3. RESULTS
2.1 Preparation of substrate ana’ surface film
3.1 Snucture of substrates and surface jilms
Single crystal specimens were produced by a technique described previously [12]. The specimens were placed on a specimen holder in a vacuum evaporator and the residual surface contamination was removed by argon ionic bombardment at 2 kV with a current density of approx. lOpA/cm’. As the evaporating materials, 99.9% Ag, 99.99% Cu and 99.9% Ni as shown in Table 1 were chosen. The heating boats of molybdenum sheet for silver and copper films, and of tungsten wire coated by alumina for nickel film were used. The surface thin films were evaporated onto both surface of the annealed aluminium crystals with an evaporation rate of 300 A/mm _ 600 AImin in vacuum of 10m6Torr. The evaporation point source was covered during the initial heating period in order to drive-off impurities and to achieve a steady evaporation rate. The film thickness evaporated on the aluminium substrate was determined from weighing method. Most film thickness used in the present work was approx. 1.6pm.
The average values of dislocation density and of subgrain size obtained by etch-pitting technique in aluminium single crystal substrate were 1.0 x 10’ cm-’ and 0.3 mm dia., respectively. X-ray diffraction of the surface of specimens coated with 1.6pm Ag, Cu and Ni indicated that each film consisted of tine polycrystalline grains. The grain size and the lattice strain were measured, and calibrated against annealed pure metals [ 141. The values of grain size and the lattice strain were obtained using Barret’s equations [lS], and are given in Table 1. The grain sizes of silver, copper and nickel film were obtained as >0.2, 0.17 and 0.09 ym, the lattice strains of those being 0.21, 0.18 and 0.82 x lo-“, respectively. The grain sizes and the lattice strains of silver and copper films are nearly equal, while those of nickel film are smaller and higher than others.
2.2 Tensile test The tensile deformation was carried out at room and at - 196°C by using an Instron-type testing of 8.3 x machine, and an initial strain-rate 10m6 set-i. 2.3 Obsewations Optical and a scanning electron microscopy were carried out in order to find the relationship between the flow modes in films (crack-like or glide-like) and the slip lines of substrate. Furthermore, the mean slip length was measured both on the unplated and the plated specimens after deformation, and was corre-
3.2 Stress-strain behaviour Typical relationships between resolved shear stress and shear strain (on the primary slip system of aluminium single crystal substrate) at room and liquid nitrogen temperatures are shown in Fig 1 both for the specimens with and without films. The initial yield stress at room temperature of the unplated specimen is 210 g/mm’, while those coated with silver, copper and nickel film are, in sequence, 275, 305 and 445 g,mm’. Also at liquid nitrogen temperature, they are, in sequence, 285, 350, 490 and 570g;mm’. The specimens coated with silver and copper strainharden linearly as the unplated specimen does, though the coated specimens show somewhat higher level of flow stress than the unplated one. On the other hand, the specimen coated with nickel strain-
Table 1. Evaporation conditions and film properties Purity
Rate
Material
(x4
(A/min)
Boat
Ni cu Ag
99.9 99.99 99.9
-300 -300 -600
(coated A1203) MO sheet MO sheet
Grain size
Lattice strain
(pm)
(x10-y
0.09 0.17 > 0.2
0.81 0.1s 0.11
W wire
T.K_ASCGI
.=,YDIZL3fI:
SURF,ACE STREXGTHEYIXG
WITH EV.\POR;\TED
FILMS
f It9
i
Fig. I. The stres-strain
beha\iour of the specimens with and without films at room and liquid mtrogtn temperatures.
hardens parabolically. Such a behaviour has also been observed for the specimen coated with electro-deposited nickel previously Cl?]. The reason why the Yi curve becomes more nearly linear at 77 K is considered to be caused by the embrittlement of Ni film at low temperature. The tendency of crack initiation by the embrittlement of the film and the diferential thermal contraction at low temperature is thought to mask the effective longrange strengthening. After the plated specimens were strained, the films were removed, then the bared specimens were strained further. The stress-strain behaviours during those processes are shown in Fig. 2. The reioaded stress-strain curves of the bared specimens initially coated with silver and copper coincide with those of the originally unplated specimen. while the flow stress of the bared specimen initially coated with nickel does not, indicating much higher strain-hardening state. The increased stresses of the specimens coated with silver and copper consists only of the stress component recovered after removing films, while that of the specimen coated with nickel involves the com-
/
/ 3
2
3
&oed
5 rhex
-
?
s&in, %
a
(
9
3
/I
Fig. 2. The retained and the recovered stress after removing films.
ponent retained recovered.
after removing
film. as u-e11 as that
3.3 Sltrfilce obserurion The optical microscopy photographs of the plated and unplated specimens deformed to approx. 8”, shear strain at room temperature are shown in Fig. 3. In this observation, the difference was also recognized between the surface marking of the specimens coated with silver and copper and the specimen coated with nickel. The film surfaces of the former two specimens show the “glide-like” markings corresponding to the slip length and spacing of the unplated specimen. On the other hand, the film surface of the latter specimen shows the -‘crack-like” markings. The mean slip length was measured both on the plated and the unplated specimens which were strained at room temperature. The results are given in Table 2. For two crystals, the mean s!ip lengths of the unplated and the silver- and copper-plated specimens are nearly equal. But the mean slip lengths of the specimens coated with nickel are shorter than others. According to Seeger et al. [ 161 and Kelly [ 173. the reciprocal of mean slip length is proportional to strain for aluminium single crystals. Therefore, the shorter slip length in the present case means the more strained state (more hardened state). Thus. she difference in mean slip length can be related to the retained stress components.
For the crystal C, dislocation densities of the plated and the unpiated specimens a’ere measured by the etch-pitting technique. The results are given in Table 2. The initial dislocation density of this specimen was approx. 1.0 x lo- cm-‘. The dislocation distributions in coated specimens were relative!y homogeneous and did not show the existence of ‘-damage layers” which
Ag iitm
Nfi film
Fig. 3. The optical microscopy photographs of the plated and unplated specimens deformed to approx. So, shear strain at room temperature. tai indicates photographs for unplated specimen, (b) Appfated specimen, (c) Cu-plated specimen and {dd)Wi-piared specimen.
of the specimens which were coated with silver, copper and nickel, then removed the EIms and further strained to 8.0”, shear strain were 6.7 x IO’,
silver- and copper-coated specimens show the similar densities to the unplated specimen, whiie the nickeicoated specimen shows higher value than others. Thus, both from this observation and from the slip
7.0 x lo- and 2.0 x 10’ cmw3: respectively. (31 the other hand, the density of the specimen \vhich was Thus, the o&inally unpiated was t5.8 x 10: cm-‘.
length measurement, an essential difkrcnce has been recognized between the specimens coared with silver and copper, and tvith nickel. Furthermore> the diger-
was reported
by Bilello [9]. The dislocation
densities
Table 2. Dislocation densities and mean slip lengths in the Ag-, Cu- and X-plated.
axx4
the unplatzd specimens Slip length !,=-4 unplated Ag plated Cu ptated Ni plated unplated Ag plated Cu plated Ni plated
SKI522 O.OSi3L 0.0534 0.0321 0.0476 0.043 1 0,@4!32 O”O2il
Disfocation density (1 km’) 6.8 6.7 7.0 2.0
x x x x
lo10’ IOIO”
f,,ln.>*:~Resolved shear strain uhtre film wvas removed. Q,_.: Rtsolwd shear strain u!xre mean slip length and dislocation density were measured.
TAK.YSUGI
AND
IZUMI:
WITH EVAPORATED FILMS
SURFACE STRENGTHENISG
1111
(b) Ni film
(a)
Ag.Cufilm
I
I Plated
Unploted
tt
L
,
1
NI Ni I
CRTK-196-3
c
c
Fig. 1. The schematic illustration of the role of the longrange interaction stress component, rL and the short-range component. ?p (a) indicates the case which the specimens coated with Ag and Cu were deformed, (b) indicates the case in which the specimens coated with Ni were deformed. ence in dislocation density, Ap, between unplated and nickel-plated specimen can be correlated with the retained stress component, Aa, as discussed later. After the slip line length was measured, X-ray backreflection Laue spots were observed at the same area of the specimens, i.e. at the center portion of the surface. No marked difference was recognized between the spots from the silver- and copper-coated specimens and from the unplated specimen; that is. all spots were diffused in the same way. However, the deformed specimen coated with nickel showed asterisms. Such a result means that. in the nickel-coated specimen, lattice bending (lattice curvature) was introduced during deformation.
4. DISCUSSION 4.1 General interprerarion of surface strengthening due to the evaporated films
From the present experimental results described above, it is clear that different coating materials show the different effects on the surface strengthenings. In the previous work [12] the authors have emphasized that the origin of the film strengthening should be divided into two components; that is, a retained stress component, arising from a long-range film-substrate interaction, and a recovered stress component, from a short-range film-substrate interaction. The former could be explained the plastic constraint due to the film [13], while the latter could be explained by the image force mechanism [S]. At yield point the surface strengthening is controlled only by the shortrange interaction component. As the plastic strain increases, the long-range interaction component increases. From results represented above, the surface strengthening in the specimens coated with silver, copper and nickel is schematically shown in Fig. 4.
Fig. 5. Relation between the increment of yield stress and the image force term. The former is expressed as the ratio of r p,;l,edrunplrled calculated from experimental values. and the latter as (G2 - G,)/(G, + G,) given by the values in Table 3.
The surface strengthening in specimens coated with silver and copper as shown in Fig. 4(a) seems to be due to the short-range interaction component not only at yield point range but also at plastic strain one. The reason why the long-range interaction component was lacking in these specimens may be that the film-substrate plastic constraint was accommodated by glide-like deformation in film. Such a lack of the long-range interaction component is also supported by the experimental results, such as (1) a linear strain hardening, (2) the lack of retained stress, (3) the glide-like deformation in film, (4) lower dislocation density and (5) a lack of asterisms. On the other hand, the surface strengthening in specimen coated with nickel as shown in Fig. 4(b) seems to be controlled only by the short-range interaction component at yield point, but with increasing plastic strain the long-range interaction component increases parabolically. Such a consideration is rationally explained by the facts: (1) a parabolic strain hardening, (2) the occurrence of the retained stress component, (3) the higher dislocation density and (4) the existence of asterisms. 4.2 Controlling factor of yield stress As obviously shown in Fig. 4, the factor which controls the yield stress is the short-range interaction component (the image force mechanism). Image force due to an abrupt change of shear modulus at the boundary of different materials depends on their modulus’ difference, and increases as the difference increases. The values of the shear modulus for aluminium, silver, copper and nickel at room temperature are given in Table 3. As shown in Fig. 1, the yield stresses in specimens coated with silver. copper
Table 3. Shear modulus at room temperature
(dyn&‘)
Al
Ag
cu
Ni
0.260 x10-t*
0.313 X lo-‘2
0.419 X lo-‘2
0.77 X lo-”
II 12
TAKASUGI
ASD
IZUMI:
SURFACE
STRENGTHENING
and nickel take in sequence higher values. The amounts of increased yield stress plotted against the magnitude of image force, (G, - G1)/(G2 + G,) given by Table 3 are shown in Fig. 5. It is clear that the yield stress increases with increasing the modulus’ difference. Therefore, the yield stress seems to be connected with the image force acting at the boundary of different materials. However, in this figure, the specimen coated with silver shows a higher value of yield stress than that of the unplated specimen, although no difference of shear modulus can be seen between aluminum and silver. Barnett [18], using the model of continuously distributed screw dislocations near the boundary of a bi-metallic medium with different shear moduli, calculates the stresses generated in the softer phase (containing the pile-up), and in the harder phase, due to image forces. The image term increases with increasing difference in shear modulus. When the medium is perfectly homogeneous, the image term is zero, but the real pile-up term takes the following tinite value Tij
=
SjLJp,
(1)
where T indicates an applied shear stress, L is a pile up length which is taken equal to half the specimen thickness, and p is the distance from interface. Thus the increased yield stress in specimen coated with silver seems to originate in the real pile-up term at the interface. It is concluded that the higher yield stresses in specimens coated with copper and nickel were caused by adding further the image force term to the pile-up one. 4.3 Controlling factor of jlow stress As represented in Fig. 4, the flow stresses in specimens coated with silver and copper were controlled only by the short-range component, while the flow stress in specimen coated with nickel was controlled both by the parabolically increasing long-range component* and by the short-range component. That is, the long-range component (plastic constraint) was not effectively introduced on the specimens coated with silver and copper, while it was effective on the specimen coated with nickel. The conditions for introducing the plastic constraint may in general be characterized by the following three points; (1) the extent of the constraint increases as the difference in modulus and strength between two materials increases, (2) the voids or strip pings do not occur on the interface during deformation, (3) the strain accommodation does not occur in the substrate and film or in the one side material * Whatever the physical hardening mechanism of this component may be, the hardening cannot be established without an accumulation process of dislocations by a stress accommodation. This accumulation is induced by the existence of a rigid film. Therefore, although the specimen shows the stress level corresponding to this dislocation structure immediately after stripping Ni film, the hardening curve thereafter does not follow parabolically.
WITH
EVAPORATED
FILMS
during deformation. From condition (I ,!_we can easily recognize that the plastic constraint on the specimen coated with nickel is most remarkable. Furthermore, because any strippings after deformation could not be observed in the present specimens. the different behaviours of deformation are impossible to be interpreted from condition (2). Finally, in condition (3), as shown in Fig. 3, the films in the specimens coated with silver and copper showed the “glide-like” markings corresponding to the slip lines of unplated specimen. Therefore the behaviours of the films in these specimens seem to be accommodated with the strain during deformation. On the other hand the film surface of the specimen coated with nickel showed the “crack-like” markings which do not correspond to the slip lines of unplated specimen. Such a deformation in the nickel film indicates that the strain-accommodation was not established. Thus, it is concluded that the long-range interaction component (plastic constraint) in the specimen coated with nickel was introduced by the conditions (1 and 3). Finally, let’s consider the film deformation modes which correlate with the activities of the long-range interaction component. It should be considered that the “glide-like” markings in the specimens coated with silver and copper were originated from the high stress field (stress concentration) due to the pile-up dislocations in aluminium substrate. Moreoever, such a high stress field would inject into the grains in the film. However, the “crack-like” markings in the specimen coated with nickel may be originated not by the high stress field due to piled-up dislocations in aluminium substrate, but by an applied tensile axis stress. Such a consideration is supported by the observations that the strain at which the crack appears did not depend on the primary shear strains of aluminium single crystal substrate, but it depended on the tensile axis strains. That is, they were approximately the same, = 1%.
5. SUMMARY The surface strengthening mechanism of aluminium single crystal coated with evaporated silver, copper and nickel films was studied. From the stress-strain behaviour, the surface observation, dislocation density measurement, X-ray Laue spots observation and the mean slip length mexurement, the following results were obtained. (1) The yield and flow stresses in specimens coated with silver, copper and nickel showed in sequence, higher values than those in the unplated specimen. (2) The specimens coated with silver and copper showed linear strain hardenings similar to the condition of the unplated specimen, while the specimen coated with nickel showed the parabolic strain hardening. (3) After removing the fihns. the specimens coated with silver and copper showed the “recovered stress”,
TXKASUGI
.&ND
IZUMI:
SURF;\CE
STRENGTHENING
while the specimen coated with nickel showed both the “recovered” and the “retained stress”. (4) Flow mode in silver and copper film was “glidelike”, while that in nickel film was “crack-like”. (5) From the dislocation density measurement, the mean slip length measurement and the X-ray spots observation. a greater strained condition was observed in the deformed specimen coated with nickel. (6) The surface strengthening in specimens coated with silver and copper could be explained by the short-range interaction component not only at yield point but also at plastic strain range. On the other hand, that with nickel could be interpreted by adding more of the parabolically increasing long-range component to the short-range one. (7) The short-range interaction component is thought to be due to the image force and pile-up force, while long-range interaction component is due to the plastic constraint. The latter component is correlated with the shear modulus’ and the strength’s differences between the substrate and the film. and with the deformation mode of the film. REFERENCES 1. I. R. Kramer 131 (1961).
and L. J. Demer, Prog. Marer. Sci. 9,
WITH
EVAPORATED
FILSIS
1113
2. E. S. htachlin, Srrrngrhening Mechanisms in Solids. p. 375. ASM. Cleveland. OH (1962). A. R. C. Westwood. Encironmmr-Sensiticr 2fechanicnl Brhaciour. p. 1. Gordon & Breach, New York (1966). 4. F. R. N. Nabarro. Theory of Ctysral Dislocnrions, p. 275. Oxford University Press. London (1967). I. G. Greenfield, Snrfaces and Inrrrfuces II: Physical and Mechanical Properries. p. 61. Syracuse L-niv. Press. New York (1968). A. J. Sedricks and A. R. C. West6. R. M. Latanisoon. wood. Hondu Memorial Series on .Lfarerinl Science No. 1: Structure and Properties of Jferal Surjaces. p. 500. Maruzen. Tokyo (1973). 1. J. C. Fisher. Trans. .Lfer. Sot. ,AfJfE 19-l. 531 (1952). s. A. K. Head. Phil. Mao. 44. 92 (1953). 9. J. C. Bilello, Scripra ,cfet. 4, 49j (1973). 10. C. S. Barrett, ilcra .\fet. 1, 2 (1953). 11. J. Pridans and J. C. Bilello, ,Icm .C/rr. 20. 1339 (1972). 12. T. Takasugi and 0. Izumi. Acra .bfer. 23. 1111 (1975). 13. IM. F. Ashby. Phil. 4fng. 21, 399 (1970). 1-t. B. E. Warren and J. Biscoe, J. .Am. Cergni. Sot. 49, 49 (19381, p. 1li. McGraw15. C. S. Barrett et al.. Srrucrure oj.lfrr&, Hill. New York (1966). Phil. 16. A. Seeger, J. Diehl. S. Mader and H. Rebstock. &fag. 2, 323 (1957). 17. A. Kelly, Phil. Jfrrg. 1, 835 (1956). 1% D. Xl. Barnett. rlcra 4fet. 15, %39 (1967).
and OSXMU
TAKASC’GI
IZUMI
The Research Institute for Iron. Steel and Other Metals. Tohoku University, Sendai. Japan (Receirrti
16 frmr
1975: in
rerised_forrn 1 April 1976)
.&tract-Using evaporated silver. copper and nickel films covering a broad spectrum of shear moduli and film strength, the effect of surface coatings on the deformation of aluminium single crystals was studied. From the anaiyses of stress-strain behaviour, surface marking observation, slip length measurement, disiocation density measurement and X-ray diffraction, a distinct difference was clearly observed in the surface strengthening in accordance with the coating materials. The surface strengthening by coating with silver and copper. which revealed a slight effect. was explained by the short range intetaction, while the strengthening by nickel coating was interpreted by adding more of the long-range interaction to the short-range one. The short-range interaction component is thought to be due to an image force and pile-up force. On the other hand. the long-range interaction component is thought to be due to the plastic constraint which is correlated with the differences in shear modulus and strength between the substrate and the Film. and to the deformation mode of the film itself. RPsum&--On a Ctudii l’effet d’un revetement superficiel de monocristaux d’aluminium sur leur dtformation, en utilisant des films evapores d’argent, de cuivre et de nickel qui presentaient une grande gamme de modules de cisaillement et de resistances. L’analyse des courbes contrainte-deformation, des observations de marquages superficiels. des mesures de la longueur glissement et de la densite de dislocations. et des diffractions de rayons X ont permis de mettre en evidence une nette difference dans Ie durcissement selon le materiau de revttement. On a expliqui le petit effet de durcissement supetficiel dd a un revetement d’argent et de cuivre par l’interaction Q courte distance, alors qu’on a interpret6 le durcissement par un revetement de nickel en ajoutant une interaction a longue distance a l’interaction a courte distance. On pense que la composante &interaction a courte distance provient dune force-image et dune force d’empilement. On pense par contre que la composante a longue distance de l’interaction provient de la contrainte plastique, lice 6 des differences de module de cisaillement et de resistance entre le substrat et le film, et au mode de deformation du film lui-meme. Zu~mmenfa~~g-per EinfluB van O~rfl~chenschichten auf das Ve~ormungsverhalten von Aluminium Einkristallen wurde mittels aufgedampften Silber-. Kupfer- oder Nickelfilmen untersucht. Es wurden in Abhlngigkeit vom Uberzugsmaterial ausgeprlgte Unterschiede in der Oberflachenhlrtung beobachtet, angezeigt durch das Spannungs-Dehnungsverhalten, die Oberflachenbeobachtungen. die L2ngen der Gleitlinien, die Versetzungsdichte und die Riintgenbeugung. Der kleine EinfluB der Beschichtung mit Silber oder Gold wurde durch eine kunreichende Wechselwirkung erklart. der EinRuB der Nickelbeschichtung wurde mit einer Zunahme der langreichweitigen Wechselwirkung gegeniiber der kurzreichenden interpretiert. Die kurzreichende Wechselwirkung wird einer Bildkraft und einer Kraft herriihrend van Aufstauungen zugeschrieben. Andererseits wird angenommen. da13 die langreichweitige Wechselw~rkungskomponente von einer plastischen Einschrankung herriihrr, die mit dem Unterschied in Schermodul und Festigkeit zwischen Substrat und Film und der Ve~o~ungsa~ der Schicht selbst zusammenhIngt.
1. IXTRODUCTION
It is well known that surface conditions can influence plastic deformation of metallic single crystals. The strengthening produced by the presence of a surface film is much greater than what could be expected simply by the strength contribution due to the film. The reader is referred to a number of review articles [l-6]. The mechanisms which have been proposed to explain the surface strengthening effects fall into five
categories: (1) the surface dislocation source blocking mechanism [73, (2) the image force mechanism [S], (3) the surface damage zone mechanism [9], (4) the surface barrier mechanism against the egression of dislocations [lo] and (5) the plastic constraint mechanism (compatibility mechanism) [l 11. However none have achieved universal acceptance. It has been suggested[ll, 121 that strengthening could not be attributed
the
surface
to a single mechanism, but to several mechanisms. Takasugi and
I II37
110s
TAKASUGI
AND
IZUMI:
SURFACE STRENGTHESING
Izumi [12] proposed that the surface strengthening in aluminium single crystals coated with electro-deposited nickel film consists of long- and the short-range interaction stress components. The former could be explained by the plastic constraint theory [13]. The latter could be explained by the image force mechanism. It is expected that both interaction components will be altered by various types of films. The purpose of the present work is to explain these interaction components more clearly using different materials as evaporated films, with a range of shear moduli and film strength. 2. EXPERIMENTAL
WITH EVAPORATED
FILMS
lated with hardening mechanism. The procedure of the measurement was as follows. The plated and unplated specimens were strained in tension to a desired strain then unloaded. Both specimens were then electro-polished (here, surface films of plated specimen were removed). Slip line lengths were measured at the middle portion of the wide surface of specimens. Etch-pitting and X-ray observations of dislocation were carried out before and after the deformation both on the unplated and plated specimens. In the case of plated specimens, dislocation density was measured after removing the films by electro-polishing [12].
PROCEDURES
3. RESULTS
2.1 Preparation of substrate ana’ surface film
3.1 Snucture of substrates and surface jilms
Single crystal specimens were produced by a technique described previously [12]. The specimens were placed on a specimen holder in a vacuum evaporator and the residual surface contamination was removed by argon ionic bombardment at 2 kV with a current density of approx. lOpA/cm’. As the evaporating materials, 99.9% Ag, 99.99% Cu and 99.9% Ni as shown in Table 1 were chosen. The heating boats of molybdenum sheet for silver and copper films, and of tungsten wire coated by alumina for nickel film were used. The surface thin films were evaporated onto both surface of the annealed aluminium crystals with an evaporation rate of 300 A/mm _ 600 AImin in vacuum of 10m6Torr. The evaporation point source was covered during the initial heating period in order to drive-off impurities and to achieve a steady evaporation rate. The film thickness evaporated on the aluminium substrate was determined from weighing method. Most film thickness used in the present work was approx. 1.6pm.
The average values of dislocation density and of subgrain size obtained by etch-pitting technique in aluminium single crystal substrate were 1.0 x 10’ cm-’ and 0.3 mm dia., respectively. X-ray diffraction of the surface of specimens coated with 1.6pm Ag, Cu and Ni indicated that each film consisted of tine polycrystalline grains. The grain size and the lattice strain were measured, and calibrated against annealed pure metals [ 141. The values of grain size and the lattice strain were obtained using Barret’s equations [lS], and are given in Table 1. The grain sizes of silver, copper and nickel film were obtained as >0.2, 0.17 and 0.09 ym, the lattice strains of those being 0.21, 0.18 and 0.82 x lo-“, respectively. The grain sizes and the lattice strains of silver and copper films are nearly equal, while those of nickel film are smaller and higher than others.
2.2 Tensile test The tensile deformation was carried out at room and at - 196°C by using an Instron-type testing of 8.3 x machine, and an initial strain-rate 10m6 set-i. 2.3 Obsewations Optical and a scanning electron microscopy were carried out in order to find the relationship between the flow modes in films (crack-like or glide-like) and the slip lines of substrate. Furthermore, the mean slip length was measured both on the unplated and the plated specimens after deformation, and was corre-
3.2 Stress-strain behaviour Typical relationships between resolved shear stress and shear strain (on the primary slip system of aluminium single crystal substrate) at room and liquid nitrogen temperatures are shown in Fig 1 both for the specimens with and without films. The initial yield stress at room temperature of the unplated specimen is 210 g/mm’, while those coated with silver, copper and nickel film are, in sequence, 275, 305 and 445 g,mm’. Also at liquid nitrogen temperature, they are, in sequence, 285, 350, 490 and 570g;mm’. The specimens coated with silver and copper strainharden linearly as the unplated specimen does, though the coated specimens show somewhat higher level of flow stress than the unplated one. On the other hand, the specimen coated with nickel strain-
Table 1. Evaporation conditions and film properties Purity
Rate
Material
(x4
(A/min)
Boat
Ni cu Ag
99.9 99.99 99.9
-300 -300 -600
(coated A1203) MO sheet MO sheet
Grain size
Lattice strain
(pm)
(x10-y
0.09 0.17 > 0.2
0.81 0.1s 0.11
W wire
T.K_ASCGI
.=,YDIZL3fI:
SURF,ACE STREXGTHEYIXG
WITH EV.\POR;\TED
FILMS
f It9
i
Fig. I. The stres-strain
beha\iour of the specimens with and without films at room and liquid mtrogtn temperatures.
hardens parabolically. Such a behaviour has also been observed for the specimen coated with electro-deposited nickel previously Cl?]. The reason why the Yi curve becomes more nearly linear at 77 K is considered to be caused by the embrittlement of Ni film at low temperature. The tendency of crack initiation by the embrittlement of the film and the diferential thermal contraction at low temperature is thought to mask the effective longrange strengthening. After the plated specimens were strained, the films were removed, then the bared specimens were strained further. The stress-strain behaviours during those processes are shown in Fig. 2. The reioaded stress-strain curves of the bared specimens initially coated with silver and copper coincide with those of the originally unplated specimen. while the flow stress of the bared specimen initially coated with nickel does not, indicating much higher strain-hardening state. The increased stresses of the specimens coated with silver and copper consists only of the stress component recovered after removing films, while that of the specimen coated with nickel involves the com-
/
/ 3
2
3
&oed
5 rhex
-
?
s&in, %
a
(
9
3
/I
Fig. 2. The retained and the recovered stress after removing films.
ponent retained recovered.
after removing
film. as u-e11 as that
3.3 Sltrfilce obserurion The optical microscopy photographs of the plated and unplated specimens deformed to approx. 8”, shear strain at room temperature are shown in Fig. 3. In this observation, the difference was also recognized between the surface marking of the specimens coated with silver and copper and the specimen coated with nickel. The film surfaces of the former two specimens show the “glide-like” markings corresponding to the slip length and spacing of the unplated specimen. On the other hand, the film surface of the latter specimen shows the -‘crack-like” markings. The mean slip length was measured both on the plated and the unplated specimens which were strained at room temperature. The results are given in Table 2. For two crystals, the mean s!ip lengths of the unplated and the silver- and copper-plated specimens are nearly equal. But the mean slip lengths of the specimens coated with nickel are shorter than others. According to Seeger et al. [ 161 and Kelly [ 173. the reciprocal of mean slip length is proportional to strain for aluminium single crystals. Therefore, the shorter slip length in the present case means the more strained state (more hardened state). Thus. she difference in mean slip length can be related to the retained stress components.
For the crystal C, dislocation densities of the plated and the unpiated specimens a’ere measured by the etch-pitting technique. The results are given in Table 2. The initial dislocation density of this specimen was approx. 1.0 x lo- cm-‘. The dislocation distributions in coated specimens were relative!y homogeneous and did not show the existence of ‘-damage layers” which
Ag iitm
Nfi film
Fig. 3. The optical microscopy photographs of the plated and unplated specimens deformed to approx. So, shear strain at room temperature. tai indicates photographs for unplated specimen, (b) Appfated specimen, (c) Cu-plated specimen and {dd)Wi-piared specimen.
of the specimens which were coated with silver, copper and nickel, then removed the EIms and further strained to 8.0”, shear strain were 6.7 x IO’,
silver- and copper-coated specimens show the similar densities to the unplated specimen, whiie the nickeicoated specimen shows higher value than others. Thus, both from this observation and from the slip
7.0 x lo- and 2.0 x 10’ cmw3: respectively. (31 the other hand, the density of the specimen \vhich was Thus, the o&inally unpiated was t5.8 x 10: cm-‘.
length measurement, an essential difkrcnce has been recognized between the specimens coared with silver and copper, and tvith nickel. Furthermore> the diger-
was reported
by Bilello [9]. The dislocation
densities
Table 2. Dislocation densities and mean slip lengths in the Ag-, Cu- and X-plated.
axx4
the unplatzd specimens Slip length !,=-4 unplated Ag plated Cu ptated Ni plated unplated Ag plated Cu plated Ni plated
SKI522 O.OSi3L 0.0534 0.0321 0.0476 0.043 1 0,@4!32 O”O2il
Disfocation density (1 km’) 6.8 6.7 7.0 2.0
x x x x
lo10’ IOIO”
f,,ln.>*:~Resolved shear strain uhtre film wvas removed. Q,_.: Rtsolwd shear strain u!xre mean slip length and dislocation density were measured.
TAK.YSUGI
AND
IZUMI:
WITH EVAPORATED FILMS
SURFACE STRENGTHENISG
1111
(b) Ni film
(a)
Ag.Cufilm
I
I Plated
Unploted
tt
L
,
1
NI Ni I
CRTK-196-3
c
c
Fig. 1. The schematic illustration of the role of the longrange interaction stress component, rL and the short-range component. ?p (a) indicates the case which the specimens coated with Ag and Cu were deformed, (b) indicates the case in which the specimens coated with Ni were deformed. ence in dislocation density, Ap, between unplated and nickel-plated specimen can be correlated with the retained stress component, Aa, as discussed later. After the slip line length was measured, X-ray backreflection Laue spots were observed at the same area of the specimens, i.e. at the center portion of the surface. No marked difference was recognized between the spots from the silver- and copper-coated specimens and from the unplated specimen; that is. all spots were diffused in the same way. However, the deformed specimen coated with nickel showed asterisms. Such a result means that. in the nickel-coated specimen, lattice bending (lattice curvature) was introduced during deformation.
4. DISCUSSION 4.1 General interprerarion of surface strengthening due to the evaporated films
From the present experimental results described above, it is clear that different coating materials show the different effects on the surface strengthenings. In the previous work [12] the authors have emphasized that the origin of the film strengthening should be divided into two components; that is, a retained stress component, arising from a long-range film-substrate interaction, and a recovered stress component, from a short-range film-substrate interaction. The former could be explained the plastic constraint due to the film [13], while the latter could be explained by the image force mechanism [S]. At yield point the surface strengthening is controlled only by the shortrange interaction component. As the plastic strain increases, the long-range interaction component increases. From results represented above, the surface strengthening in the specimens coated with silver, copper and nickel is schematically shown in Fig. 4.
Fig. 5. Relation between the increment of yield stress and the image force term. The former is expressed as the ratio of r p,;l,edrunplrled calculated from experimental values. and the latter as (G2 - G,)/(G, + G,) given by the values in Table 3.
The surface strengthening in specimens coated with silver and copper as shown in Fig. 4(a) seems to be due to the short-range interaction component not only at yield point range but also at plastic strain one. The reason why the long-range interaction component was lacking in these specimens may be that the film-substrate plastic constraint was accommodated by glide-like deformation in film. Such a lack of the long-range interaction component is also supported by the experimental results, such as (1) a linear strain hardening, (2) the lack of retained stress, (3) the glide-like deformation in film, (4) lower dislocation density and (5) a lack of asterisms. On the other hand, the surface strengthening in specimen coated with nickel as shown in Fig. 4(b) seems to be controlled only by the short-range interaction component at yield point, but with increasing plastic strain the long-range interaction component increases parabolically. Such a consideration is rationally explained by the facts: (1) a parabolic strain hardening, (2) the occurrence of the retained stress component, (3) the higher dislocation density and (4) the existence of asterisms. 4.2 Controlling factor of yield stress As obviously shown in Fig. 4, the factor which controls the yield stress is the short-range interaction component (the image force mechanism). Image force due to an abrupt change of shear modulus at the boundary of different materials depends on their modulus’ difference, and increases as the difference increases. The values of the shear modulus for aluminium, silver, copper and nickel at room temperature are given in Table 3. As shown in Fig. 1, the yield stresses in specimens coated with silver. copper
Table 3. Shear modulus at room temperature
(dyn&‘)
Al
Ag
cu
Ni
0.260 x10-t*
0.313 X lo-‘2
0.419 X lo-‘2
0.77 X lo-”
II 12
TAKASUGI
ASD
IZUMI:
SURFACE
STRENGTHENING
and nickel take in sequence higher values. The amounts of increased yield stress plotted against the magnitude of image force, (G, - G1)/(G2 + G,) given by Table 3 are shown in Fig. 5. It is clear that the yield stress increases with increasing the modulus’ difference. Therefore, the yield stress seems to be connected with the image force acting at the boundary of different materials. However, in this figure, the specimen coated with silver shows a higher value of yield stress than that of the unplated specimen, although no difference of shear modulus can be seen between aluminum and silver. Barnett [18], using the model of continuously distributed screw dislocations near the boundary of a bi-metallic medium with different shear moduli, calculates the stresses generated in the softer phase (containing the pile-up), and in the harder phase, due to image forces. The image term increases with increasing difference in shear modulus. When the medium is perfectly homogeneous, the image term is zero, but the real pile-up term takes the following tinite value Tij
=
SjLJp,
(1)
where T indicates an applied shear stress, L is a pile up length which is taken equal to half the specimen thickness, and p is the distance from interface. Thus the increased yield stress in specimen coated with silver seems to originate in the real pile-up term at the interface. It is concluded that the higher yield stresses in specimens coated with copper and nickel were caused by adding further the image force term to the pile-up one. 4.3 Controlling factor of jlow stress As represented in Fig. 4, the flow stresses in specimens coated with silver and copper were controlled only by the short-range component, while the flow stress in specimen coated with nickel was controlled both by the parabolically increasing long-range component* and by the short-range component. That is, the long-range component (plastic constraint) was not effectively introduced on the specimens coated with silver and copper, while it was effective on the specimen coated with nickel. The conditions for introducing the plastic constraint may in general be characterized by the following three points; (1) the extent of the constraint increases as the difference in modulus and strength between two materials increases, (2) the voids or strip pings do not occur on the interface during deformation, (3) the strain accommodation does not occur in the substrate and film or in the one side material * Whatever the physical hardening mechanism of this component may be, the hardening cannot be established without an accumulation process of dislocations by a stress accommodation. This accumulation is induced by the existence of a rigid film. Therefore, although the specimen shows the stress level corresponding to this dislocation structure immediately after stripping Ni film, the hardening curve thereafter does not follow parabolically.
WITH
EVAPORATED
FILMS
during deformation. From condition (I ,!_we can easily recognize that the plastic constraint on the specimen coated with nickel is most remarkable. Furthermore, because any strippings after deformation could not be observed in the present specimens. the different behaviours of deformation are impossible to be interpreted from condition (2). Finally, in condition (3), as shown in Fig. 3, the films in the specimens coated with silver and copper showed the “glide-like” markings corresponding to the slip lines of unplated specimen. Therefore the behaviours of the films in these specimens seem to be accommodated with the strain during deformation. On the other hand the film surface of the specimen coated with nickel showed the “crack-like” markings which do not correspond to the slip lines of unplated specimen. Such a deformation in the nickel film indicates that the strain-accommodation was not established. Thus, it is concluded that the long-range interaction component (plastic constraint) in the specimen coated with nickel was introduced by the conditions (1 and 3). Finally, let’s consider the film deformation modes which correlate with the activities of the long-range interaction component. It should be considered that the “glide-like” markings in the specimens coated with silver and copper were originated from the high stress field (stress concentration) due to the pile-up dislocations in aluminium substrate. Moreoever, such a high stress field would inject into the grains in the film. However, the “crack-like” markings in the specimen coated with nickel may be originated not by the high stress field due to piled-up dislocations in aluminium substrate, but by an applied tensile axis stress. Such a consideration is supported by the observations that the strain at which the crack appears did not depend on the primary shear strains of aluminium single crystal substrate, but it depended on the tensile axis strains. That is, they were approximately the same, = 1%.
5. SUMMARY The surface strengthening mechanism of aluminium single crystal coated with evaporated silver, copper and nickel films was studied. From the stress-strain behaviour, the surface observation, dislocation density measurement, X-ray Laue spots observation and the mean slip length mexurement, the following results were obtained. (1) The yield and flow stresses in specimens coated with silver, copper and nickel showed in sequence, higher values than those in the unplated specimen. (2) The specimens coated with silver and copper showed linear strain hardenings similar to the condition of the unplated specimen, while the specimen coated with nickel showed the parabolic strain hardening. (3) After removing the fihns. the specimens coated with silver and copper showed the “recovered stress”,
TXKASUGI
.&ND
IZUMI:
SURF;\CE
STRENGTHENING
while the specimen coated with nickel showed both the “recovered” and the “retained stress”. (4) Flow mode in silver and copper film was “glidelike”, while that in nickel film was “crack-like”. (5) From the dislocation density measurement, the mean slip length measurement and the X-ray spots observation. a greater strained condition was observed in the deformed specimen coated with nickel. (6) The surface strengthening in specimens coated with silver and copper could be explained by the short-range interaction component not only at yield point but also at plastic strain range. On the other hand, that with nickel could be interpreted by adding more of the parabolically increasing long-range component to the short-range one. (7) The short-range interaction component is thought to be due to the image force and pile-up force, while long-range interaction component is due to the plastic constraint. The latter component is correlated with the shear modulus’ and the strength’s differences between the substrate and the film. and with the deformation mode of the film. REFERENCES 1. I. R. Kramer 131 (1961).
and L. J. Demer, Prog. Marer. Sci. 9,
WITH
EVAPORATED
FILSIS
1113
2. E. S. htachlin, Srrrngrhening Mechanisms in Solids. p. 375. ASM. Cleveland. OH (1962). A. R. C. Westwood. Encironmmr-Sensiticr 2fechanicnl Brhaciour. p. 1. Gordon & Breach, New York (1966). 4. F. R. N. Nabarro. Theory of Ctysral Dislocnrions, p. 275. Oxford University Press. London (1967). I. G. Greenfield, Snrfaces and Inrrrfuces II: Physical and Mechanical Properries. p. 61. Syracuse L-niv. Press. New York (1968). A. J. Sedricks and A. R. C. West6. R. M. Latanisoon. wood. Hondu Memorial Series on .Lfarerinl Science No. 1: Structure and Properties of Jferal Surjaces. p. 500. Maruzen. Tokyo (1973). 1. J. C. Fisher. Trans. .Lfer. Sot. ,AfJfE 19-l. 531 (1952). s. A. K. Head. Phil. Mao. 44. 92 (1953). 9. J. C. Bilello, Scripra ,cfet. 4, 49j (1973). 10. C. S. Barrett, ilcra .\fet. 1, 2 (1953). 11. J. Pridans and J. C. Bilello, ,Icm .C/rr. 20. 1339 (1972). 12. T. Takasugi and 0. Izumi. Acra .bfer. 23. 1111 (1975). 13. IM. F. Ashby. Phil. 4fng. 21, 399 (1970). 1-t. B. E. Warren and J. Biscoe, J. .Am. Cergni. Sot. 49, 49 (19381, p. 1li. McGraw15. C. S. Barrett et al.. Srrucrure oj.lfrr&, Hill. New York (1966). Phil. 16. A. Seeger, J. Diehl. S. Mader and H. Rebstock. &fag. 2, 323 (1957). 17. A. Kelly, Phil. Jfrrg. 1, 835 (1956). 1% D. Xl. Barnett. rlcra 4fet. 15, %39 (1967).