TiN composites

Surface and Coatings Technology, 68/69 (1994) 459—467

459

Mechanical properties and tribological behaviour of nanolayered Al/A1203 and Ti/TiN composites Y. Ding, Z. Farhat, D.O. Northwood, A.T. Alpas Engineering Materials Group, Mechanical Engineering Department, University of Windsor, Windsor, Ont. N9B 3P4, Canada

Abstract Mechanical and tribological properties of nanolayered composites of Al/A1203 and Ti/TiN were investigated. Alternating layers of metals and ceramics were deposited using an r.f. magnetron sputtering technique. Nanoindentation tests wereperformed to determine force—displacement curves which were used to calculate elastic moduli and nanohardness of composites as a function of distance between layers. It was observed that both elastic modulus and hardness of composites increased with decreasing layer thickness. A good agreement was found between experimentally determined values for elastic modulus and predictions based on the rule of mixtures for isostress conditions. The hardness of Al/Al203 and Ti/TiN could be described in the formalism of the Hall—Petch-type equation indicating that ceramic layers inhibited slip transfer across metallic layers. A deviation from the Hall—Petch type of strengthening was observed in Al/Al203 at small interlayer spacings. Friction and wear behaviour of laminates was studied using a pin-on-disc type of tribometer. A systematic increase in wear resistance with decreasing layer thickness was observed. The peak friction coefficient decreased about 70% in Al/A1203 (with 200 nm Al layer thickness) while a significant 60% improvement in steady state friction coefficient was measured in Ti/TiN (with 150 nm Ti layer thickness) in comparison with as-sputtered monolithic metallic films. These observations indicated that the nanolaminated films are suitable for applications where a combination of low coefficient of friction, high wear resistance and hardness is required.

1. Introduction Mechanical properties of thin films containing nanolayers of metallic and/or ceramic (oxide, nitride) phases are becoming a focus of interest because of the high strength levels attainable [1]. Conventional composite theories such as rule of mixtures have been used to predict the strength and hardness of nanolayered materials [2,3]. However, it has been shown that when layer thickness is the smallest dimension in the structure the strength obeys a Hall—Petch-type relationship [4,5]. Lehoczky [6] (following the predictions based on Koehier’s analysis [7]) found that the tensile strength should increase inversely with layer thickness until a critical thickness value, which corresponds to a thickness where a Frank—Read dislocation source ceases to operate, is reached. In this paper, the mechanical properties, namely strength and the modulus, of two types of nanoscale laminates Al/Al203 and Ti/TiN produced by r.f. sputtering were studied. Since microlaminated materials of this type may be deposited on the surface of bulk materials and may thus have significant technological potential as wear resistant surface coatings, the study was extended to the characterization of the tribological properties of these Al/A1203 and Ti/TiN microlaminates. 2. Experimental details The Al/A1203 and Ti/TiN microlaminated composites were fabricated by an r.f. magnetron sputtering tech-

SSDI 0257-8972(94)08105-8

nique. The vacuum system was pumped to a pressure of iO~Torr prior to sputtering, and a pressure of (1.5—5.0) x iO~Torr was maintained during the deposition process. Alternating Al and A1203 layers were deposited by positioning the substrates mounted on a rotating substrate table over each target source for a predetermined time, which depended on the required layer thickness. An Al target was prepared from pure Al (99.999%) and an A1203 target was prepared from sintered A1203 (99.99%). Since each target source (Al or A1203) is sputtered by a separate cathode, the bombarding ion current densities could be varied independently. Al films were deposited using only high purity Ar (99.999%) as the sputter gas, whereas the Al203 films were deposited using a sputter gas containing both Ar (99.999%) and 02 (99.999%). Without the compensation of the 02 partial pressure, the resulting A1203 films were oxygen deficient. Full details of the fabrication of the laminated Al/Al203 composites can be found in Ref. [8]. The titanium layers in the Ti/TiN composites were directly sputter deposited from a Ti (99.999%) target. The TiN layers were deposited using the same Ti target, but this time nitrogen was used as a reactive sputtering gas. Details of sputtering conditions are given in Table 1. Al and Al/A1203 films were deposited on a 1100 Al alloy whereas Ti and Ti/TiN films were deposited on AISI 310 type stainless steel plates and silica glass slides. Ti/TiN films deposited on the silica glass substrate were used for the friction and wear tests. For these tests the

Elsevier Science S.A.

460

Y Ding

et al.

/

Al/Al

203 and Ti/TiN composites

Table 1 Details of the sputtering conditions

Sputtering gas Reactive gas Working pressure Target Target diameter Target-to-substrate distance Substrate temperature Deposition Power density rate

Al/A1203 laminated composites

Ti/TiN laminated composites

Al layer

—

A1203 layer

Ti layer

TiN layer

Ar (99.999%)

Ar (99.999%)

(1.5—5.0) x iO~Torr Al (99.999%) 25.0mm 50 mm

Ar (99.999%) 02 (99.999%) (1.5—5.0) x i0~Torr Sintered A1203 (99.99%) 25.0mm 50mm

(2—4.5) x i0~Torr Ti (99.999%) 25.0 mm 50 mm

Ar (99.999%) N2 (99.999%) (2—4.5) x iO~Torr Ti (99.999%) 25.0 mm 50 mm

320 K 1 15—20Wcm~2 10 nm min

350 ~0.6Knm min’ 10—l5Wcm2

32010Knm min’ 10—20Wcm2

350 ~0.5Knm min1 10—2OWcm2

films were deposited in the form of 20 x 20 mm2 square coupons

Table 2 Mechanical properties of monolithic Al and Ti (as-sputtered) films, and laminated Al/A1 203 and Ti/TiN composites

Cross-sectional specimens were prepared for scanning electron microscopy (SEM) examination to confirm the thickness of each layer in Al/A1203 and Ti/TiN laminated composites (Fig. 1). The periodicity and continuity of the layers are clearly seen. The layer thicknesses of the films are given in Table 2. X-ray diffraction (XRD) patterns of Al/Al203 and Ti/TiN composites were obtained using a Rigaku X-ray powder diffractometer

_______________________________________________________

Hardness (GPa) E (GPa)

Monolithic Al film Al(500 nm)/A1203 (20 nm) Al(200 nm)/Al203 (20 nm) Al(70 nm)/A1203 (10 nm) Monolithic Ti film Ti(450 nm)/TiN (40 nm) Ti(lSOnm)/TiN (2Onm)

a _____________

_______

-~

____________

_________

______

—.

—I b Fig. 1. Scanning electron micrographs of the cross-section of (a) Al/Al203 and (b) Ti/TiN.

H/E l0~)

(x

1.70 3.29 4.52 4.83

±0.03 ±0.06 + 0.08 + 0.08

—

11.96 ±0.25 12.04 ±0.20 12.80±0.25

60.92 ±1.50 61.69 ±1.50 66.99 + 1.50 84.75 ±2.00

27.9 53.3 67.5 57.0

113.10±2.5 119.55 ±2.7 117.42±2.5

105.7 100.7 109.0

—

and Cu Kit radiation. XRD patterns of the films are shown in Fig. 2. It is seen that Ti/TiN films exhibit clearly defined TiN peaks albeit with some broadening. Thus Ti and TiN layers in these films have crystalline structures. On the contrary, the XRD pattern of Al/Al203 composites is similar to that of the as-sputtered monolithic Al and does not reveal the presence of a crystalline oxide phase. The non-crystalline nature of the aluminium oxide layers has also been shown by selected area diffraction studies by transmission electron microscopy [8]. The total thicknesses of the films used in the wear tests were 10—12 j.tm. Therefore, depending on the thickness of the individual layers, the films were made up of a minimum of 20 pairs of layers (e.g. in Ti(450 nm)TiN(40 nm)) to 150 pairs (in Al(70 nm)/A1203(10 nm)). Investigations of the mechanical properties of the composites were performed using a commercially able nanohardness measurement instrument, availthe UMIS-2000 developed by CSIRO, Australia [9]. With this instrument, displacement of the indenter and the load could bethe measured independently with a resolution of 1 nm and 10_i mN respectively. The indenter was a Berkovich diamond pyramid with an angle of

Y Ding et al.

/

Al/A1

203 and Ti/TiN composites

~ Al/A1203 composites

S. ~ ~

~I

i~i

~

~.

~

~.

Aluminum film fl

ri.

.~ .

~,-r, ~

20

________

40

____________________

140 60

80

a

100

~ ..~

II

~.

~.JL~._

—

~

~

~.

_________

~.

—

As-sputtered Ti Film

—

—

.

_______________________

~< 30.

b

40.

50.

60.

70.

60.

~

100.

110.

120.

130.

140.

2 0 (Deg.)

Fig. 2. XRD patterns of (a) as-sputtered aluminium and Al/A1203 and (b) as-sputtered Ti and Ti/TiN.

65.3°between the tip axis and the faces of the triangular pyramid. The force steps were applied in a square root sequence to produce approximately equal increments of penetration depth. For each test indentations were produced in line at a constant interval of 30 p.m. The indenter tip was then moved to a new location and the process was repeated until an average 17 measurements were made on a given up specimen. Continuous loading—unloading indentations to a load of 20 mN for A1/Al 203 and 50 mN for Ti/TiN were made. The largest depth of indentation was always less than 10% of the total thickness of specimens. Hardness H was determined using the relationship between plastic depth h~of penetration and indenter load P as follows [9]: H P‘A 1 — —

1/E* = (1 v2)/E + (1 v?)/E~ (~) where E and v are elastic modulus and Poisson’s ratio for the specimens, while E 1 and v~are the same parameters for the Berkovich diamond indenter. In the current analysis E~and v~ were taken as 1050 GPa and 0.3 respectively [12]. Friction and wear tests were carried out using a miniature pin-on-disc type wear rig designed and built at the University of Windsor. The coefficient of friction, defined as the ratio of the tangential to the normal load, was measured as a function of sliding distance. The friction action was provided by a spring loaded stainless steel ball (radius, 2.15 mm) under a constant normal load of 1 N and a sliding speed of 1.3 x 10-2 ms* Tests were performed under ambient conditions. The nominal diameter of the wear track was about 6 mm. Wear track widths were measured by a low power optical microscope (x 10) at different time intervals to calculate the volume of material lost during wear.

Ti/TiN (l5Onm/2Onm) Composites

~

elastic recovery which occurs during the unloading. The unloading curve was used to determine the elastic modulus E of composites using equations developed by Doerner and Nix [10] and Sneddon [11]. The slope of the force P and displacement h curve on unloading is givenby 5 (2) dP/dh= 1.67 E*AO where E* is defined as effective modulus of the system:

120

20 (Deg.)

2

461

/

where A is the area of contact and A = kh~(the constant k for a Berkovich-type indenter is 24.5). The total penetration depth consists of a plastic component and

3. Results and discussion

3.1. Mechanical properties Force—displacement curves for Al/A1203 composites with different aluminium layer thicknesses are shown in Fig. 3(a) and similar curves for Ti/TiN composites are given in Fig. 3(b). Each curve in these figures represents an average of 17 measurements taken from each specimen. The nanoindentation measurements show a consistent trend: as the spacing between metallic layers becomes smaller, the depth of penetration of the indenter decreases (at a constant load). The hardnesses of the 13max were determined composites at the maximum load using Eq. (1) and are summarized in Table 2. Table 2 also lists the elastic moduli of the specimens calculated from the slope of the upper one-third portion of unloading part of the force—displacement curves using Eqs. (2) and (3). It is clear that, compared with as-sputtered monolithic metallic films, microlaminated composites incorporating Al203 or TiN layers exhibit higher hardnesses and elastic moduli. For example, a 68% increase in hardness and 39% increase in elastic modulus are observed in laminates consisting of 70 nm thick Al layers with 10 nm A1203 relative to as-sputtered aluminium

462

/

Y Ding et al.

25

0

200

400

600

I

I

-

4

5

20

800

2

3

1000

Al/A1

203 and Ti/TiN composites

1200

25

____________________________

1

z

E

Table3 Elastic moduli and hardnesses of composites calculated using the rule of mixtures

Al/A1203

15

15

~~10 0

5 0

0

200

a:b: loading unloading 600 800 1000

400

PENETRATION

b.

20 510

0 1200

DEPTH (nm)

0

200

400

600

800

1000 60

I

432 50

a

0.04 0.09

62.96 65.73

2.63 3.89

70

10

0.88

0.12

67.48

4.49

Ti/TiN )~ 450 (nm) )~ 40 0.92 f~ 0.08 E 120.31 (GPa)fractions H (GPa) _____________________________ 150 20 (nm) respectively. 0.88 0.12 124.16 metallic 2 and and 2~are ceramic layer components thicknesses, fmfm and f,~ are volume of

nated composites consistently increase with decreasing metal layer thickness. Since indentations were made perpendicular to the plane of layers, it can be argued

Eiaminate=

1/(fm/Em +f~/E~)

(4a)

50

40

30 ~

10

‘~m/(’~m +

fc

=

i~c/(2m + )Le)

30 20

the moduli fractions of the of the layers. monolithic To predict metal the films elasticand moduli A1203 of and TiN films were used (i.e. EAI = 60.92 GPa, ET~=

10

113.1 GPa, EAI2O3 = 320 GPa, ET1N = 450 GPa). The elastic moduli of A1/Al203 and Ti/TiN composite films, calculated using Eqs. (4), are given in Table 3. There is

0

0

=

the laminates the experimentally determined values of

b

0

2~) and

(4b) and ‘tm’ ~ are the thicknesses of metallic and nonmetallic layers, and fm, ,f~are the corresponding volume

frn

0

0.96 0.91

where

1

loading

Z40

20

20 20

that isostress conditions should prevail in the layers and thus the modulus of the compressed laminates would be

a 60

500 200

200 400 600 800 1000 PENETRATION DEPTH (nm)

b Fig. 3. (a) Force—penetration curves: curve 1, pure aluminium (target); curve 2, pure aluminium film (as sputtered); curve 3, Al(500 nm)/ A1203 (20 nm); curve 4, Al(200 nm)/A1203 (20 nm); curve 5, Al(70 nm)/Al203(10 nm). (b) Force—penetration curves: curve 1, pure titanium (target); curve 2, pure Ti (as sputtered); curve 3, Ti(450 nm)/TiN(40 nm); curve 4, Ti(150 nm)/TiN(20 nm).

films (with an average grain size of 100 nm). It is also noted that the Ti/TiN composites have hardness values as high as 13 GPa which are about 2.5 times higher than the maximum hardness reached in Al/Al203 laminates. As-sputtered monolithic and composite films show a smaller degree of hysteresis during unloading compared with conventional grain size metals (indicated by curves 1 in Fig. 3). This is consistent with a higher degree of elastic contribution to the total displacement in the as-sputtered metals. Table 2 indicates that the elastic moduli of the lami-

athereasonable agreement between the of elastic moduli of laminates (seemeasured Table 2) values and the predictions based on the rule of mixtures. Thus the elastic modulus, which is a microstructure-insensitive property, is controlled by the volume fractions of the phases and, since isostress conditions prevail, then the contributions of the oxide and nitride layers at relatively low volume fractions to the elastic moduli of the composites are relatively small. The dependence of hardness, and thus flow strength, on the layer thickness is more complex and cannot be successfully predicted using equations similar to Eqs. (4). This is shown in Table 3 where the rule of mixtures underestimates the hardnesses of composites compared with the measured values (given in Table 2). Thus the role of oxide and nitride layers is not limited to supporting the applied load. In these fine structures the layers could play additional roles in restricting the plastic deformation by either inhibiting dislocation sources from becoming operative [7] or restricting slip transfer across the adjacent metallic layers [13, 14]. The hardness data for Ti/TiN and A1/A1203 composites are plotted in Fig. 4

Y Ding et al.

14

z

Al/A1

203 and Ti/TiN composites

....~

13

~

/

I,

14

expected to occur, can be calculated using the relationship

13

t<32ItbAQ2B

12

12

ii

11

6’’

8

5 —

4

:

5

T —

-

V

TI/TIN

•

A1/A1203

463

~

(6)

/1A)/(/2B +~tA)

confirmed These theoretical by Lehoczky predictions [6]. have According been experimentally to Koehier’s model [7], for a dislocation moving on a glide plane in a material A with a low modulus into a material B with a higher modulus, the minimum stress required for yield is given by =

(VA + VBEB/EA)(am + a~) 83t(/,LB +

(7)

VA and VB are the where O~m = PBCUB EA — /IA)/ volume fractions, and EB are the elastic moduli, and U,~and 11n are the shear moduli of A and B respectively; cr~is the stress caused by frictional forces in B. Using the values in Table 4, the yield strengths of the Al/A1 203 laminates are calculated as 1.57 GPa, 1.90 GPa and /~A).

3

2

~Z

__________________________________

1600

x

2400

—0.5

3200

GPa for A1(500 nm)/A1203(20 nm), A1(200 nm)/ which arenm) A1203(20 of and the order A1(70 nm)/A1203(10 of measured values nm) respectively, (assuming 2.10

4000

_0.5)

(m

H/3) in Tablemechanism 2. However,should according to Eq. (7) this given strengthening only beay =

come operative with aluminium layers thinner than 41 nm (or thinner than 50 nm for Ti) which is smaller than the smallest A tested for the Al/A1203 laminates.

as a function of reciprocal square root of 4metallic layer 5)).According to Fig. the hardness thickness m° data can be described by a vs. layer(2_05 thickness Hall—Petch-type equation. For Ti/TiN

In the laminated composites, different strengthening mechanism(s) might be operative in different microstructure(s) and several mechanisms might operate simultaneously. Further work is required to rationalize the differences in strengthening mechanisms in terms of differences in the structure of the hard layers, e.g. the aluminium have an and amorphous whereas theoxide TiN layers is crystalline thus thestructure, TiN—Ti interfaces may have an epitaxial structure which would influence the dislocation movement across the layers as

H(Ti/TiN)=11 080 MPa+(0.70 MPa m°5)2°5 (5a)

has been demonstrated by Kung et al. [17].

Fig. 4. Hardness vs. A°~ plots for Al/A1203 and Ti/TiN.

For A1/A1 203 (with the exception of laminates with the smallest aluminium layer thickness) 5)2°~5 (5b) H(A1/A1203)= 1700 MPa + (1.09 MPa m° These results are in agreement with previous observations for Al/A12O3 laminates [4, 5]. Fig. 4 shows that there is no significant further increase in hardness of the A1/A1203 laminates for very thin aluminium layers (2<200 nm) once the strength reaches a high proportion of the theoretical strength (e.g. E/H ~ 20). For the situation where there is no further increase in hardness with decreasing A it is appropriate to consider Koehier [7] strengthening. Koehier [7] predicted that dislocations will require a large externally applied force which causes them to move from the material of lower elastic constant to the material of higher elastic constant. The critical thickness, below which Koehier strengthening is

3.2. Coefficient offriction and wear behaviour of friction vs. sliding distance curves forThe Al, coefficient Al/Al 203, Ti and Ti/TiN all have the same overall shape, but with different values for peak friction coefficient, steady state friction coefficient and transition time to steady-state. Figs. 5(a) to 5(d) show friction

Table 4 Elastic moduli, Poisson’s ratio, shear moduli and Burger’s vector of Al, A1203, Ti and TiN _____________________________________________________

Al E (GPa) ji (GPa) v b (nm)

60.92 25.4 0.345 2.86

Al203

[16] [16]

320 155.0 0.3

Ti

[16]

113.1 44.0 0.345 2.90

TiN

[16] [16]

____________________________________________________

450 [15] 173.0 0.3 —

464

Y. Ding et al. / Al/A1

203 and Ti/TIN composites 1.0

15 1.4

0.9 0.8

1.3

1.2 i_1.1

_______________________________

0.0

o:

0.0

~

o

_______________________________

10

0

20

30 40 50 80 70 80 SUDINGDISTANCE(rn?

1000

2000

3000

4000

90

100

10

20

30

1000

40

~

1.0

50

2000

b

NUI~ERoF CYCLES

a

0

5000

,,..i..,,

~

0

60

70

3000

80

4000

90

100

5000

or cycr~s

1.0

0.9

.

.

0.9

0.8

.

,

0.8 ~O.7. ~O.8

_________________

0

10

20

30

I..,.

4.0 50

60

70 80

90

100

0

10 20

SLIDING DISTANCE (m) I..

.1

0

C

1000

2000

3000

..I.

4000

30

40 50

60

SLIDING DISTANCE .1.

I...

•

5000

0

1000

NUIIBER OF CYCLES

d

2000

70 80

90

100

(ni)

3000

4000

5000

NU~EROF CYCLES

Fig. 5. Friction coefficient vs. sliding distance for (a) Al (target), (b) as-sputtered Al, (c) Al( 500 nm)/A1203(20 nm) and (d) Al(200 nm)/Al203(20 nm).

curves for the four Al-based specimens, namely Al target material, as-sputtered Al and A1/Al203 (with 500 nm and 200 nm Al layer thickness) respectively. The target aluminium (grain size = 5 mm) exhibits a high peak friction coefficient of 1.34 ±0.02 ( ±0.02 represents the variation in the coefficient of friction around a mean value). However, the as-sputtered aluminium shows a decrease in the peak friction coefficient to a value of 0.62 ±0.06, a reduction of about 54%. As for the microlaminated A1/A1203 films, the peak friction

coefficient decreases from 0.52 ±0.02 to 0.36 ±0.03 as the aluminium layer thickness decreases from 500 to 200 nm with a net reduction of about 73% in the peak friction coefficient between the target aluminium (grain size = 5 mm) and the 200 nm microlaminated A1/A1203 composites. The steady state friction coefficient does not vary significantly between the target Al and the microlaminated Al/A1203 composites (0.08 and 0.12). However, the as-sputtered aluminium shows a higher value of 0.25 ±0.05. Data for as-sputtered Ti and the

Y Ding et al.

/

Al/A1

Ti/TiN microlaminated composites, having 450 and 150 nm Ti layer thickness (Fig. 6), show that there is no significant decrease in the peak friction coefficient. However, the steady state friction coefficient is reduced from 0.39 ±0.03 to 0.23 ±0.03 (about 40% reduction) as the Ti layer thickness decreases from 450 to 150 nm. Moreover, a 57% reduction in the steady state friction coefficient between the 200 nm Ti/TiN microlaminated composite and the as-sputtered Ti is achieved. A system1.0

atic decrease in the sliding distance (time) to steady state with decreasing layer thickness is observed, for both A1/A1203 and Ti/TiN microlaminated composites. The value of the friction coefficient j.t is determined by the sum of the individual contributions of the principal friction components [18, 19], namely adhesion (12a), ploughing (j~) and deformation (ud). Hence, /1

(8)

_____________________________________

1.0

.

0.9

~: :; 0.5

z

8

20.4

004

~0.3

~0.3

t

0.2

o.~

0.2

0.1

0.0

+f~i.t~ +fd/.Ld

fapa

—J•.—.—.,..

0.9

465

203 and Ti/TIN composites

2

0.1

o

~

10

20

30

40

50

60 70

SLIDING DISTANCE

80

90 100

0.0

(na)

~

0

~.

0

~O0O

a

2000

3000

4000

10 20

30

40

50

60

SLIDING DISTANCE

.1.1...

.1...,

I

70

(m)

.1...

80

I

90 100

I..

5000

b

NU~ERor c~c~

0

1000

2000 OF3000 NUMBER CYCLES 4000

5000

1.0 0.9 0.8 0.7

0.6

8

o.~

~

0.4

0.3 0.2 0.1 010 I....

0

C

20 I....

1000

30 40 50 60 70 60 SLIDING DISTANCE Cm) ~

90

100

•

2000 3000 4000 NUMBER OF CYCLE3

6000

Fig. 6. Friction coefficient versus sliding distance for (a) as-sputtered Ti, (b) Ti(450 nm)/TiN(40 nm) and (c) Ti( 150 nm)/TiN(20 nm).

466

Y. Ding et a!.

/

Al/A1

203 and Ti/TiN composites

where coefficients fa,p,d indicate the relative contribution of each mechanism. The ac.tual values of the contributions are determined by the characteristics of the wear mechanism(s) taking place. As the pin slides on the surface of the specimen, its asperities plough and deform the specimen surface generating wear debris. At the peak friction value, the surface roughness and the amount of wear debris between the two surfaces reach maxima. A low peak coefficient of friction means a higher resistance to plastic deformation and ploughing. The coefficient of friction then slowly decreases as the asperities of both surfaces are polished away. At the steady state both surfaces are polished to a smooth finish, This behaviour is also reflected in the volume loss vs. sliding distance curves which show two distinct regions. The first region is associated with a high wear rate and it is commonly classified as “severe mode of wear” (this is the initial wear rate ~ in Table 5). This region is also

S LI Eli Nil DIRE CTI0N,j,

_________

S1l111N6

DIRECTIUN1~

associated with a large friction coefficient and a rough wear track (as shown in Fig. 7(a) for aluminium). The second region can be classified as “mild mode of wear” (this is the steady state wear rate W~in Table 5) since it is associated with low wear rate, low friction coefficient and smooth wear track (Fig. 7(b)). The transition occurs at sliding distances of about 10—20 m, which correspond to the sliding distances at which there is a reduction from peak friction coefficient to steady state friction coefficient. The wear tracks in Al/A1203 with 200 nm Al layer thickness (Fig. 7(c)) are smoother than for aluminium at the peak friction coefficient because of its higher yield strength and resistance to ploughing. The general wear rate vs. sliding distance behaviour of Ti and Ti/TiN is similar to that of Al and A1/A1203 (Table 5). The initial and the steady state wear rates ~ and W~for both types of laminates are reduced as the metal layer thickness is reduced. Furthermore, there is

~SLl0 INO Dl

~

SLIDING DIREcTI~I’~

Fig. 7. SEM images of wear tracks of lal aluminium target at a sliding distance of 22 m, I b) aluminium target at a sliding distance of 9Dm. (c) Al(200 nm)/Al203(20 nm) at a sliding distance of 15 m and (d) Al(200 nm)/Al203(20 nm) at a sliding distance of 90 m.

Y Ding et a!.

/

Al/A1

203 and Ti/TIN composites

ances than monolithic metallic films. The peak coefficient of A1/Al203 is 70% lower than that of monolithic

Table 5 Initial and steady state wear rates l4~and 1’~’~ 3mm3m1)

467

(x 10

(x105mm3 m~)

Al (target)a Al (as sputtered)b Al(500 nm)/A1 203(20 nm) Al(200nm)/Al203(20 nm)

7.69 ±0.66 1.82 ±0.37

342.46 ±6.94 2.77 ±0.39

1.46 ±0.24 1.31 ±0.29

2.35 ±0.33 1.57 ±0.19

Ti (as sputtered)° Ti(450 nm)/TiN(40 nm) Ti(lSOnm)/TiN(2Onm)

0.09 ±0.02 0.08 ±0.02 0.07±0.02

0.32 ±0.01 0.18 ±0.01 0.10±0.02

Grain size, 5.0 ±2.0 mm. 50 ±10 nm.

b

aluminium. A 60% decrease in steady state friction coefficient observed inTi. Ti/TiN with distance respect to as-sputteredis unreinforced The sliding to steady state is reduced in both types of composites.

Acknowledgements

This work is supported by National Research Council

Grain size, 100 ±20 nm.

Grain size,

of Canada (NSERC) through a Strategic Grant. The authors would like to thank Mr. J. Robinson for his help in fabrication of composites and mechanical property measurements. References

a reduction in the initial and the steady state wear rates due to the incorporation of the ceramic layers. The wear rates for Ti and Ti/TiN microlaminated composites are lower than observed for Al and Al/A1203 microlaminated composites because Ti and Ti/TiN have higher hardness and can resist plastic deformation more efficiently.

4. Summary and conclusions Al/A1203 and Ti/TiN-type nanolaminated composites are fabricated using an r.f. magnetron sputtering method. Nanoindentation experiments are performed to study elastic and plastic deformation characteristics as a function of layer spacing. Tribological properties of laminates are investigated using a pin-on-disc type of tribometer. The main results can be summarized as follows. (1) The elastic moduli of both Al/Al203 and Ti/TiN increase as the volume fraction of ceramic layers increase. The elastic moduli of laminates could be predicted using a simple rule of mixtures model for isostress condition. (2) Nanohardnesses of laminates increase linearly with reciprocal root square of thickness of metallic layers suggesting that the flow strength of Ti/TiN and Al/Al203 can be described by a Hall—Petch-type equation. When the interlayer spacing of ceramic phases becomes small (less than 200 nm in Al/A1203), deviations from the Hall—Petch slope occur. (3) Laminates consistently exhibit higher wear resist-

[1] P.H. Townsend, T.P. Weihs, J.E. Sanchez, Jr., and P. (eds.), Thin Films: Stress and Mechanical Properties IV, Res. Soc. Symp. Proc., 308 (1993). [2] J. Grilhé, in Mater. Sc!. Forum, 59—60 (1990) 481. [3] T.P. Weihs, T.W. Barbee, Jr., and M.A. Wall, Mater. Symp. Proc., 308 (1993) 753. [4] A.T. Alpas, J.D. Embury, D.A. Hardwick and R.W. J. Mater. Sc!., 25 (1990) 1603. [5] Y. Ding, DO. Northwood and A.T. Alpas, Proc. Interfaces 2, Ballarat, November 1993, in press.

Borgesen in Mater.

Res. Soc. Springer, Conf. on

[6] S.L. Lehoczky, J. App!. Phys., 49 (1978) 5479. [7] J.S. Koehler, Phys. Rev. B, 2 (1970) 547. [8] Y. Ding, DO. Northwood and A.T. Alpas, Surf. Coat. Technol., 62 (1993) 448. [9] T.J. Bell, J.S. Field and M.V. Swain, Mater. Forum, 17(1993)127. [10] M.F. Doerner and W.D. Nix, J. Mater. Res., 1(4) (1986) 601. [11] IN. Sneddon, Int. J. Eng. Sc!., 3 (1965) 47. [12] [13] [14] [15]

[16]

[17]

[18] [19]

N. and Phys. T.J. Bell, App!. Phys., (7) (1992) E.O.Savvides Hall, Proc. Soc.J.London B, 64 72 (1951) 747. 2791. N.J. Petch, J. Iron Steel Inst., 174 (1953) 25. L. Chollet and C. Biselli, in D. Dowson, C.M. Taylor and M. Godet (eds.), Mechanics of Coatings, Tribology Series 7, Elsevier, Amsterdam, 1990, p. 217. H.J. Frost and M.F. Ashby, Deformation-Mechanism Maps: The Plasticity and Creep of Metals and Ceramics, Pergamon, Oxford, 1982, ~ 20, 43, 98. H. Kung, M. Nastasi, T.R. Jervis, K.M. Hubbard, R.M. Messner, T.E. Mitchell and J.D. Embury, Mater. Res. Soc. Symp. Proc., 308 (1993) 747. E. Rabinowicz, Friction and Wear of Materials, Wiley, New York, 1965, ~. 68. N. Suh, Tribophysics, Prentice-Hall, Englewood Cliffs, NJ, 1986, p. 69.