Effect of grain size on friction and wear of nanocrystalline aluminum

MATERIALS SCIENCE & EIIGINEI~ING ELSEVIER

Materials Science and Engineering A206 (1996) 302-313

A

Effect of grain size on friction and wear of nanocrystalline aluminum Z.N. Farhat, Y. Ding, D.O. Northwood, A.T. Alpas* Department of Mechanical and Materials Engineering, University of Windsor, Windsor, Ont., Canada N9B 3P4

Received 2 June 1995; in revised form 24 July 1995

Abstract

The friction and wear characteristics of nanocrystalline aluminum were investigated as a function of grain size. Nanocrystalline aluminum samples with an average diameter of 16.4 nm were produced using an r.f. magnetron sputtering technique. The grain size was increased (up to 98.0 nm) by an isothermal annealing treatment at 573 K. Hardness measurements were performed using an ultra-microhardness indentation system and it was observed that within the grain size range of 15-100 nm the hardness-grain size data could be well represented by the Hall-Petch relationship. Friction and wear measurements were made using a miniature pin-on-disk type tribometer under unlubricated conditions both in air and in vacuum. The coefficient of friction of aluminum tested against a stainless steel pin varied with the sliding distance. At the early stages of sliding the coefficient of friction rose to a peak value, and this was followed by a decrease to a steady-state value. The transition on the friction curve corresponded to a similar transition from a severe wear regime to a mild wear above a characteristic sliding distance on the cumulative volume loss versus sliding distance curve. The value of the peak coefficient of friction decreased from/~p = 1.4 for aluminum with a coarse grain size (10 6 nm) to/lp = 0.6 for the nanocrystalline aluminum with a grain size of 16.4 nm. The coefficient of friction of nanocrystalline aluminum showed a 30% increase when tested in vacuum. In the nanocrystalline grain range, the wear rates were found to be linearly dependent on the square root of the grain size. An empirical equation based on the Archard's Law is proposed to describe the effect of grain refinement on the wear resistance under unlubricated sliding conditions. A qualitative understanding of wear processes is developed in terms of the variation of the surface morphology and subsurface strength with sliding distance. Keywords: Grain size; Friction; Wear; Hardness; Nanocrystals; Aluminium

I. Introduction

The influence o f grain size on the yield and flow behaviors o f metals and ceramics have been studied extensively since the early w o r k o f Hall [1] and Petch [2]. The majority o f experimental data on the yield strength, ay, o f polycrystalline materials within the conventional grain size range, i.e. d > 1 /tm, have been shown to obey the H a l l - P e t c h equation O'y = O"o +

kd -°5

(1)

where d is the grain diameter, a o is a friction stress and k is a constant which is usually interpreted as a measure o f the contribution o f the grain boundaries to the strength [3]. The applicability o f the H a l l - P e t c h equation to materials with grain sizes smaller than 1 p m is * Corresponding author.

a question o f significant technological importance in view o f the recent advances in material processing techniques, such as rapid solidification, v a p o r deposition and sputtering which have m a d e it possible to produce materials with ultrafine grain sizes (i.e. nanocrystalline materials) which could ultimately be utilized in structural or wear resistant components. A study o f mechanical properties o f these materials is important to explore the upper limit o f strengthening that can be achieved by microstructural refinement and delineate the range o f applicability o f k n o w n strengthening mechanisms. Experimental studies have indicated that considerable increases in the strengths and hardnesses o f metals can be obtained by refining the grain size o f metals to the n a n o m e t e r scale. F o r example, an increase from 0.9 to 6.9 G P a in the hardness o f nickel was observed when the grain size was decreased f r o m 12.5 # m to 12 n m [4]. 0921-5093/96/$15.00 © 1996 - - Elsevier Science S.A. All rights reserved SSDI 0921-5093(95)10016-4

Z.N. Farhat et al. / Materials Science and Engineering A206 (1996) 302-313

Similarly, an increase from 0.5 to 2.5 GPa was achieved in the nanocrystalline copper by reducing the grain size from 50 p m to 6 nm [5]. These and other studies (e.g. Refs. [6] and [7]) indicate that within the nanoscale grain range the observed variation of hardness with grain size can most conveniently be described by the H a l l - P e t c h equation. An important aspect of strengthening by microstructural refinement is that the strength increase need not to be at the expense of the ductility. Experimental evidence for increased ductility in nanoscale ceramics has been reviewed by Siegel [8]. Fan [9] recently proposed that the plane strain fracture toughnesses, Klc, of metals and alloys are inversely dependent on the grain size. However, it is also well documented in the literature that below a certain grain size the H a l l - P e t c h slope may start to decrease and may even become negative [10-12]. The type of processing method used for varying the grain size, such as heat treatment [10] as well as presence of imperfections such as triple junctions [13,14] or porosity [15] have been cited among the possible causes for the observed inverse Hall-Perch behavior. Despite these complexities in, and lack of a full understanding of, the mechanical behavior, it is important to initiate studies on the friction and wear characteristics of nanocrystalline structures because of the perceived potential engineering applications of these materials as tribological surface coatings. It can be expected that microstructural refinement should lead to an improvement in the wear resistance according to Archard's Law of Wear [16], i.e. L W = K-H

(2)

where W is volume worn per unit sliding distance; L is the applied load; H is the hardness of the softer of the materials in contact; and K is the wear coefficient. Korshunov et al. [17] reported an increase of 2 - 5 times in the wear resistance of a quenched and tempered steel when the grain size was reduced from 100 to 10 pm. They have also observed a slight (15%) reduction in the average value of the coefficient of friction within this grain size range. Otherwise, experimental data relating wear resistance to grain size is almost non-existent in the open literature. Investigation of wear of nanocrystalline materials may also shed light on the role of tribological layers, which were found to have nanocrystalline microstructures [18,19], on the wear behavior of tribological materials. In this study nanocrystalline aluminum films with a grain size range of 15-100 nm were produced by an r.f. sputtering technique and their deformation, friction and wear characteristics under unlubricated sliding conditions were determined.

303

2. Experimental 2. I. Fabrication of nanocrystalline aluminum films Aluminum films were fabricated from high-purity cast aluminum (99.99%) targets using an r.f. magnetron sputtering technique. The sputter chamber was evacuated to < 10 - 7 Torr prior to admitting the sputtering gas (argon). During the sputtering operation, the gas pressure was maintained at 1.5 x 10 -3 Torr. The system was equipped with a thermionic emission source (a Ta cathode) to assist the glow discharge process. More details about the fabrication procedures can be found in Ref. [20]. The r.f. magnetron sputtering method provided a higher degree of control over the film thickness and uniformity compared to thermal or e-beam deposition, but the deposition rates were rather slow (10 nm min - ~). The aluminum films were deposited on substrates made of a commercial 1100 A1 alloy, The films were in the form of 20 x 20 mm 2 coupons suitable for the friction and wear tests. The total thickness of the films was about 10 /~m. In order to vary the grain size of the as-sputtered aluminum films, the films were annealed at 573 K for the time intervals of 5, 10 and 50 h. in vacuum sealed quartz tubes. The mechanical and tribological properties of samples taken from the high purity aluminum used as the sputtering target were also characterized. These samples had a very large grain size of 1.0 _+ 0.2 mm and will be referred to as coarse-grained aluminum.

2.2. Material characterization Two different methods, namely X-ray diffraction (XRD) line broadening and transmission electron microscopy (TEM) imaging techniques were used to determine the grain size of aluminum samples before and after the isothermal annealing treatment. X R D analyses were performed using a Rigaku X-ray diffractometer and a Cu-K~ radiation. X-Ray diffraction patterns of the (111) peaks were recorded by scanning aluminum films at small steps of the Bragg's angle, i.e. 0 = 0.02 °. In order to eliminate broadening effect caused by the Ko~ 2 radiation, the Rachinger [21] correction method was employed and K~ 2 profile of the (111) peak was filtered out from the X-ray spectrum. Patterns of the (111) peaks of a fully annealed polycrystalline aluminum powder with an average diameter of 2/~m were obtained in order to determine the extent of instrumental broadening which is subsequently separated from the broadening due to the sample using an empirical relation proposed by Wagner and Aqua [22]. The broadening associated with the profiles of K ~ peaks was determined by measuring the integral peak breadth, fl which is the ratio of the total area under the diffraction peak to the maximum peak height. Then the grain size, d, was calculated from the Scherrer formula

Z.N. Farhat et al. / Materials Science and Engineering A206 (1996) 302-313

304 2 d= - -

p cos0

(2)

where 2 is the X-ray wavelength. T E M observations were carried out using a JEOL 100 TEM. TEM samples were prepared using the films sputter deposited on NaC1 substrates which were subsequently dissolved in water. The average grain size was measured using the conventional linear intercept method. Selected area electron diffraction (SAED) patterns of the samples were also determined. 2.3. Nanoindentation tests The hardness of nanocrystalline aluminum films was measured using an ultra-microhardness indentation system (UMIS 2000). With this instrument, the applied load and the displacement of the indenter could be measured with a resolution of 0.1 m N and 1.0 nm respectively [23]. The instrument uses a Berkovich type (three-sided pyramid) indenter with an angle of 65.3 ° between the tip axis and the faces of the pyramid. During the test the applied force was increased in square root increments to a maximum of 20 raN. An average of 20 measurements were performed on the surface of each film to determine the load vs. displacement curves at each grain size. The nanohardness (H) at the maximum load was computed using the relationship between the plastic depth of penetration (hp) and the applied load (P), i.e. 14 = r / A

(4)

where A is given as 24.5 for the Berkovich indenter [23]. The total depth of penetration was always less than 10% of the total thickness of the films. 2.4. Friction and wear tests Friction and wear tests were performed using a miniature pin-on-disk type tribometer designed and constructed at the University of Windsor to study the tribological properties of thin films. Unlubricated sliding tests were made under a constant load of 1.0 N which was applied to the rotating sample through a stationary pin made of AISI 304 stainless steel (19 wt% Cr, 11 wt% Ni, 1.5 wt% Mn, 0.07 wt% C, 0.02 wt% N the balance) with a tip radius of r = 2.15 mm. The hardness of the pin was measured as 4.5 GPa. It is because of this high hardness and also high oxidation resistance that the AISI 304 was selected as the counterface material. The material also has a high stiffness which makes it possible to transmit the applied load to the sample with minimum deflection. AISI 304 steel which has a tendency for galling at high loads and velocities did not show galling under the current testing conditions. A circular wear track of about 6 mm in diameter was formed on the contact surface of the

aluminum. During the wear tests, a constant speed of 1.5 x 10 -2 m s - ' was maintained at the center of the wear track. The width o f the wear track was measured at regular intervals during the test using a low power optical microscope (magnification × 20) and the volume loss (V) from the surface of material was calculated according to the ASTM Standard G99 as follows g=2~RIr2sin-l(~r)-(4)(4r2-w2)°51

(5)

where V is the wear volume (ram3); R and r are the radii of the wear track and the pin, respectively; and w is the width of the wear track. The instantaneous values of the normal (P) and tangential (T) forces were measured by means of the strain gauges mounted to the arms of the tribometer and recorded with a digital data acquisition system. The data was processed to determine the variation of the coefficient of friction (/l = T / P ) as a function of sliding distance. Most tests were performed under the ambient laboratory conditions (21°C, 35% relative humidity). However, to study the effect of vacuum on the friction and wear of nanocrystalline aluminum, the pin-on-disk machine was mounted in the specimen chamber of a SEMCO Nanolab 7 scanning electron microscope whose stage was rotated by an external d.c. motor at a speed of 1.5 x 10 -3 m s '. The pressure inside the scope during the wear tests w a s 1 0 - 6 Torr. Topographical and microstructural changes in the wear tracks were characterized using a scanning electron microscope (SEM) equipped with an energy dispersive spectroscope (EDS) for semi-quantitative chemical analysis. Microhardness measurements of the material below the worn surfaces of the coarse grained aluminum were made using a Knoop indenter at a load of 0.10 N.

3. Results 3.1. Microstructure and mechanical properties The grain size of the as-sputtered aluminum was measured as 16.4 + 1.5nm by the X-ray line broadening technique. The grain size was increased to 31.7 + 3.0 nm and 98.0 + 5.0 nm after annealing for 5 and 50 h, respectively at 573 K. Fig. l(a) shows the dependence of grain size on the annealing time. Measurements made by the X-ray line broadening technique are in g o o d agreement with the TEM measurements. TEM bright field micrographs of the as-sputtered and annealed (for 18 h) samples are shown in Fig. l(b). It is seen that the microstructures before and after annealing are effectively free of porosity and consist of a uniform distribution of equiaxed grains.

305

Z.N. Farhat et al. / Materials Scienee and Engineering A206 (1996) 302-313

Fig. 2 shows the force-displacement curves o f alum i n u m with different grain sizes. Each curve represents an average o f twenty indentations. The effective plastic depth o f penetration, hp, was c o m p u t e d by plotting a regression line t h r o u g h the unloading curve over the

,

.

.

.

.

.

.

.

,

.

,

,

,

.

.

.

.

.

.

.

.

.

.

200

,

,

400

600

,

800

,

1000 1200

,

1 2

,

,

3

4

/ 25 ] - 20 /

g ,

0 25

.

.

15

100

15

a: l o a d i n g I0

I0

. - , 8o

5

60

5

N 0

40

~

,

0

i

t

200

400

/~..~

,

600

PENETRATION 20

(a)

10 20 ANNEALING

30 TIME

40 50 (HOURS)

,

I

(,

I /

0

1000 1200

DEPTH

(rim)

Fig. 2. Force displacement (penetration depth of the indenter) curves for aluminum with different grain sizes: (1) 15.4_+ 1.5 nm; (2) 31.7 + 3.0 nm: (3) 98.0 ± 5.0 nm; (4) [.0 + 0.2 mm.

, , , , l , , , * l , , , , l , , , , J , , , , l , , , ,

0

I

800

60

upper one-third o f the data and then the hardness at the m a x i m u m load (20 m N ) was calculated from the value o f hp using Eq. (4). The hardness data is plotted in Fig. 3 as a function o f the reciprocal square root o f grain size. The highest level o f hardness obtained was 1.70_+0.06 G P a (in the specimen with 16.4 n m grain size). This corresponds to an increase o f a b o u t 130% by c o m p a r i s o n to a grain size o f 98 nm. According to Fig. 3, within the 1 5 - 1 0 0 n m grain size range, the hardness data obeys the H a l l - P e t c h equation H= 34(MPa)+0.21(MPam°5)d

1800

.

.

.

.

I

'

'

'

'

°5(m o 5 )

I

'

'

'

.

.

(6)

.

.

.

1500

1200

900

Z 600

300 s "

(b) Fig. I. (a) Grain size versus annealing time at T= 573 K: (o) by XRD; (V) by TEM. (b) TEM bright field images of aluminum: (i) as-sputtered; (ii) annealed at 573 K for 18 h. Inserts show SAED patterns.

O

2000

4000

6000

d-O.5 (m-O.5)

Fig. 3. Nanohardness versus grain size (d o5).

8000

Z.N. Farhat et al. / Materials Science and Engineering A206 (1996) 302-313

306 1

.

0

. .. . . .

... . . ~ .i. . . ,

.1. . . 1 . ., . . 1 . ., . .

....

....

1.0

I ....

0.9

0.9

0.8

0.8

~p

0.7 E

0.6

8

0.5

o

0.4

u0.Ti

~

0.6

8

0.5

~

0.4

Nss

0.3

~

0.3

0.2

0.2

0.1

0.1

0.0''"

0.0

0

30 ~ 50 60 70 80 SMDINGDIST~C£(m) I , = l l l l l l l l l l l . l .... I .... 0 1000 2000 3000 ~00

(a)

10

20

90

......... 10 20

100

I. 5000

i i , l l l , , , l l

1000

0

(b)

~£ROFCYC~S 1.0

' . . .'. '. . . . ' . . . . . . . . . . . . . ' ......... : 30 ~ 50 60 70 8 0 9 0 100 SLIDINGDIST~CE(m) .... r ....

2000

3000

NUMB~

OF~C~S

t ....

t,

5000

~00

1.5 1.4

0.9

1.3 0.8

1.2 1.1

0.7

1.0 E

0,6

~

0.5

o

0.4

~

0.9 o U

0.8 0.7 0.6 0.5

0.3

0.4 0.2

0.3 0.2

0.1

0.1 ...,

0.0

....

10

J ....

20

, ....

30

i ....

40

, ....

i ....

50

60

SLIDING DISTANCE i

. . . .

0

(C)

i

. . . .

] 000

i

. . . .

2000

i ....

~

i ....

70

80

i

....

90

...,

0.0

....

10

100

, ....

20

(m)

. . . .

3000

i

. . . .

4000

i

t ....

.

5000

W O U ~ E R OF CYCLES

0

(d)

i ....

1O00

J ....

, ....

, ....

, ....

, ....

J ....

30 40 50 60 70 SLIDING DISTANCE ( r n ) i ....

i ....

2000

i

3000

N U M B E R OF

J

80 .

4000

.

.... 100

90 .

.

J ,

5000

C Y C L E S

Fig. 4. Coefficient of friction versus sliding distance curves for aluminum with different grain sizes. Load = 1 N, sliding speed = 1.5 × 1 0 - 2 m s - i . (a) 16.4+ 1.5 nm; (b) 3 1 . 7 + 3 . 0 nm; (c) 9 8 . 0 + 5 . 0 nm; (d) 1 . 0 + 0 . 2 mm.

which suggests that the Hall-Petch strengthening in polycrystalline aluminum extends to grain sizes as small as 16 nm. Data in Fig. 3 suggests also a positive deviation from the hardness of coarse grained aluminum. However, caution must be exerted in comparing nanohardness results over a broad range of grain sizes because when a load of 20 m N is applied the indentation impression covers approximately 5 - 5 0 grains for the nanocrystalline material, which is clearly not the case for the material with a grain size of 106 nm. The elastic modulus, E, of the aluminum measured from the slope of unloading parts of curves in Fig. 3 following a method proposed by Doerner and Nix [24] was 62 + 3 GPa and did not vary significantly with the grain size. 3.2. Friction and wear

The coefficient of friction versus sliding distance curves for aluminum with different grain sizes tested in air against a stainless steel pin are given in Fig. 4. The

basic shape of friction curve remains similar for all grain sizes. Each curve is characterized by two friction regimes. Initially, the coefficient of friction increases until it reaches a maximum value/tp (peak coefficient of friction). This is followed by a gradual decrease to a steady-state value/t .... where the coefficient of friction is independent of sliding distance. The principal effect of reducing the grain size is to decrease the value of/Zp. For a grain size of 98 nm,/~p = 0.94. This is reduced to 0.62 in a sample with a grain size of 16.4 nm. Compared to the coarse grained aluminum which exhibits the highest value for 1.34) this represents a total decrease about 55%. However, only a slight variation in the values of the steady-state coefficient of friction was observed. Within the broad grain size range studied, /1.... varied between 0.1-0.3. A third attribute of the friction curves is the characteristic distance, St, above which a steady-state friction condition is attained. According to Fig. 4, nanocrystalline aluminum shows a transition to steady-state in a shorter time (e.g. st ~ 50 m for d = 16.4 nm) compared jt/p

(~/p

=

Z.N. Farhat et al. / Materials Science and Engineering A206 (1996) 302 313

307

Table 1 Coefficients o f friction at different grain sizes Coefficient of friction ~

G r a i n size (nm) Nanocrystalline

l~p /L...... /l//p •~t (m)

Conventional

16.4

43.1

0.62 _+ 0.06 b 0.25 _+ 0.05 0.37 ± 0.11 ~ 50 (2600)c

0.58 O. 14 0.44 ~70

98.0 _+ 0.03 ± 0.03 ± 0.06 (3600)

0.92 O. 14 0.78 ~75

l0 s' -- 0.02 ± 0.02 ± 0.4 (4200)

1.34 O. 12 1.22 ~82

_+ 0.02 ± 0.06 ± 0.08 (4500)

= m e a n value of peak coefficient of friction;/Ls.s. = m e a n value of steady-state coefficient of friction; A/Lp - (/Lp - t~ • ); s, = t r a n s i t i o n distance to steady-state. b D e n o t e s fluctuation a r o u n d the mean. " N u m b e r o f cycles. " #p

with coarse grained aluminum for which St ~ 80 m. The values of parameters which characterize the friction curves at different grain sizes are given in Table 1. In addition to the experiments performed in air, some additional friction experiments were conducted in vacuum ( 1 0 - 6 Torr) inside the specimen chamber of a scanning electron microscope using the same pin-ondisk tribometer. Fig. 5 shows the initial stages (up to 4 m) of the friction curves for nanocrystalline aluminum tested in air and vacuum using a load of 1.0 N and a speed of 1.5 x 10-3 m s 1. In vacuum, the coefficient of friction increases from an initial value towards a maximum, but the rate of increase is higher compared to the tests in air. There is almost 40% difference between the coefficients of friction measured in the air and vacuum for all the grain sizes investigated (Table 2). The cumulative volume loss versus sliding distance curves of aluminum are given in Fig. 6. Each curve is identified by two distinct wear regions. Initially the wear rates (slope of the curves) are high (severe wear) but after a certain sliding distance, the slope decreases to a lower value (mild wear). The values of severe and mild wear rates measured at different grain sizes are given in Table 3, which shows that grain refinement has a significant influence on the wear resistance. Nanocrystalline aluminum shows an approximately two orders of magnitude decrease in the mild wear rates and about a 300% decrease in severe wear rates. This large improvement in the wear resistance indicates that nanocrystalline structures may be of value in certain tribological applications. The dependence of the wear rates on the grain size in the nanocrystalline range will be discussed in detail in Section 4.

3.3. Examination of worn surfaces In this section the changes in the coefficients of friction and wear rates are related to the changes in the surface topography and microstructural changes in the

material layers adjacent to the contact surfaces. Worn surfaces were examined at two characteristic sliding distances, namely at 12 m which corresponds to a sliding distance where coefficients of friction and wear rates are high and at 90 m where wear becomes mild and proceeds under steady-state conditions. Worn surfaces of coarse grained aluminum tested to 12 m were subjected to severe deformation and they were heavily damaged as shown in Fig. 7(a). The surface consisted of a series of parallel grooves, of about 100 /tm wide each, extending longitudinally towards the sliding direction, i.e. it was subjected to microplowing. The material within the grooves was displaced to the edges forming a topography which can be described as extruded ridges. It is not clear whether one or more wear mechanisms were responsible for the formation of the observed grooves and ridges but the origin of microplowing can be attributed to a process of abrasive wear by the hard steel counterface. The presence of continuous scratch marks along and inside the grooves can be regarded as compelling evidence for the abrasive wear. The large scale deformation processes observed on the worn surfaces appears to have led to the removal of the wear particles by the fracture of the extruded ridges. The loose debris particles detached from the worn surfaces were in the form of plate-like particles of approximately 5 - 2 0 /zm diameter. Examples of debris particles adhered on the wear track can be seen in Fig. 7(a). In contrast, nanocrystalline aluminum samples tested under the same conditions displayed smoother wear tracks. Microgrooves and ridges can also be seen on the worn surfaces of 16.4 nm aluminum (7(b)), but these are now much narrower (and shallower) indicating that nanocrystalline aluminum resisted damage caused by plastic deformation and abrasion more effectively. Furthermore, the debris particles are finer in size ( < lctm) than those shown in Fig. 7(a). Figs. 7(c) and (d) show, respectively, the worn surface morphologies for coarse and fine grain sizes at a

Z.N. Farhat et al. / Materials Science and Engineering A206 (1996) 302-313

308

1.0L

. . . .

,

. . . .

,

. . . . . . . .

,

. . . .

,

. . . .

,

. . . .

Table 2 Effect of vacuum on the friction coefficient # and wear rate, W, at the early stage of sliding

, . .

0.9 0.8

Coefficient o f friction

Grain size (nm)

0.7

0.6

Nanocrystalline

Conventional

0.5

16.4

98

106

1.06 ___0.04 0.69 + 0.06 1.77+0.11

1.30 + 0.05 0.88 + 0.10 8.10+0.46

1.60 + 0.36 1.12 + 0.15 33.33+2.50

0.45-t-0.2

3.63_+0.27

18.67_+ 1.63

0.4

J

~ 0.3 0.2 0.1 0.0 i 0.0

. . . i. . . . . i . . ii . . . . i . . . .i .

0.5

~.0

1.5

f

2.O

= . . . . . .i

3.0

2.5

1

3.5

¢.0

SLZDING DLST~CE (m) i ....

i ....

0

(a)

I ....

20

i ....

40

i ....

t ....

i ....

i ....

i ....

t ....

i..1

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

i

1.1

1.0 0.9 0.8 0,7 0.6 0.5 0.4 0,3 0,2 0.1

0.0 0.0

'

'

I

. . . .

0.5

i

. . . .

1.0

*

. . . .

1.5

I

. . . .

2.0

I

I

l

l

l

2.5

l

l

l

l

3.0

l

l

l

,

3.5

4,0

8L,IDING DIRT~CE (m) i

.

,

,

i

,

0

. . . .

i

.

.

i

r

100

50

(b)

I

i

i

,

i

150

,

200

Wu'M~ER OF CYCLES 1.6

. . . .

,

,

. . . .

. . . .

,

,

. . . .

,

. . . .

,

. . . .

. . . .

, . . .

1.4

1.2

~ 1.0 i .~)

"/1 and W measured at a sliding distance of 3.5 m (number of cycles = 190).

60 80 100 120 140 160 180 200 NU'~ER OF CYCLES

1,2 "

p a (vacuum, 10 - 6 Torr) p (ambient, 1 atm) W(×10 -4 )mm 3m -I (vacuum, 10 -6 Torr) W(xl0 4) m m 3 m - 1 (ambient, 1 atm)

o.81

~ 0.6

sliding distance of 90 m, i.e. within the mild wear regime. The most noticeable feature of the wear track of coarse grained aluminum is its smoothness (Fig. 7(c)). The surface is almost featureless and the wear grooves are hardly visible which suggests that during sliding the surface of aluminum became polished by the stainless steel pin. Debris particles rejected from both sides of worn area consist o f a mixture of both platelike and submicroscopic particles. Thus the surface of the coarse grained aluminum gradually assumes a similar morphology to that of its nanocrystalline counterpart (Fig. 7(d)). EDS (energy dispersive X-ray analyses) of the worn surfaces did not reveal the presence of any elements from the stainless steel on the worn surfaces of aluminum. Aluminum transfer to the steel pin was negligible even in the severe wear regime. In order to study the change in the flow strength of the material adjacent to the worn surfaces, as a function of sliding distance, the hardness of the material below the center of the wear track in a coarse grained aluminum sample was measured. The results are shown in Fig. 8(a). There is considerable hardening of the material layers near the contact surface. For example, the hardness measured as 350 MPa at a depth of 35 p m increases to almost 450 MPa at 5 p m beneath the wear track after a sliding distance of 1 m. The variation of hardness with the depth can be described by the following equation

0.4

H = Hs exp ( - AZ)

(7)

0.2 0.0

....

0.0

i . . . .

0.5

i ....

I ....

1.0

1.5

i . . . .

2.0

i . . . .

2.5

i , , , , i , , ,

3.0

3.5

4.0

SlZDING DISTJ~ICE (m) i

o

(c)

....

i ....

20

i,..,i

....

i

....

i

....

i

....

i

....

i

....

i ....

i,..

40 so 80 ~00 120 a 4 0 1 e 0 1 8 0 2 0 0 m a m ~ OF c y c r . ~

Fig. 5. Comparison of coefficient of friction versus sliding distance curves for aluminum tested in vacuum (10 6 Torr) and in air. L o a d = 1 N, sliding s p e e d = 1.5 x l0 - 3 ms - I . (a) 16.4 nm, air; (b) 16.4 nm, vacuum; (c) 98 nm, vacuum.

where Hs is the extrapolated surface hardness and Z is the depth measured from the contact surface. The thickness of the plastically deformed and hardened subsurface layer, i.e. the depth of material where subsurface hardening is first detected, is about 30 p m after sliding to 1 m. The depth of the plastic zone extends to 60 ~ m at 110 m. The effect of the sliding distance on the subsurface hardness (measured at 10 pm) is given in Fig. 8(b). The rate of hardening is fast at the initial

Z.N. Farhat et al. / Materials Science and Engineering A206 (1996) 302-313 0.12

I

'

'

'

'

l

'

'

'

'

l

'

'

'

'

I

'

'

'

'

l

'

'

'

'

l

'

'

'

'

l

'

'

'

309

Table 3 Mild (Win) and severe (Ws) wear rates at different grain sizes

'

Wear rate

g'- 0 . 1 0

Grainsize (nm)

(mm 3 m -1) Nanocrystalline

Conventional

0.08

16.4 0

Wm ( x 1 0 -5 ) 2.77+0.39 Ws ( x 1 0 -3 ) 1.82 +0.37

0.06

43.1

98.0

11.36+0.69 32.36+_1.57 2.37 +_0.38 3.24_+0.50

106 342.46_+6.94 10.08_+1.78

0.04 O ~>

0.02

0.00

0

20

(a)

60

80

100

120

140

SLIDING DISTANCE ( i n ) 0.7

g--

40

'

'

'

'

I

'

'

'

'

I

'

'

'

'

I

. . . .

I

'

'

'

'

I

'

'

'

'

I

'

'

'

'

0.6

~

0.5

r~

0.4

0 ~

0.3

~

0.2

o 0.1 0.0

, , , I

0

(b)

. . . .

20

I . . . .

40

I . . . .

60

I

eO

. . . .

I...,I,,

100

,,

120

140

SLIDING DISTANCE ( m )

Fig. 6. Cumulative volume loss versus sliding distance curves for (a) nanocrystalline aluminum ((O) 16.4 nm; (O) 43.1 nm; ( , ) 98 nm); (b) coarse grain size aluminum.

stages of wear (up to 20 m) but then starts to increase monotonously at a smaller rate with the sliding distance and reaches 550 MPa after a sliding distance of 120 m.

4. D i s c u s s i o n

The principal objective of this section is to discuss the friction and wear behavior of the nanocrystalline aluminum in terms of observed microstructural and mechanical parameters. However, it is appropriate to first comment briefly on the applicability of the HallPetch equation to the nanocrystalline grain size range.

The hardness versus grain size data in Fig. 3 covers a grain size range 15-100 nm which has not been explored previously. As a first approximation, assuming that the hardness H is related to the yield strength, 0- by H = 30-, the Hall-Petch parameters for nanocrystalline aluminum deduced from Eq. 6 (o-o=ll.3 MPa, k = 0.07 MPa) can be compared with the published results on aluminum with coarser grain sizes (Table 4). Both the intercept strength and the Hall-Petch slope (Eq. 6) of nanocrystalline aluminum are in agreement with those reported for a grain size range of 1-200/~m [25-29] which suggests that the grain boundaries may be continuing to provide strengthening to aluminum down to grain sizes as small as 15 nm. However, this comparison is rather simplistic and does not take into account for example the possibility of smaller work hardening rates expected in the specimens with ultrafine grain size. There is currently a controversy on the applicability of the Hall-Petch equation to the nanocrystalline materials. Based on a survey of existing literature on grain size and hardness data of nanocrystalline materials, Fougere et al. [12] concluded that an inverse Hall-Petch behavior (i.e. softening with decreasing grain size) should be expected in nanoscale materials for which the grain size is increased by annealing. The present results disagree with this suggestion. One important difference between the materials that have reported to exhibit inverse Hall-Petch behavior and nanocrystalline aluminum studied here is that the present materials produced by rf sputtering were effectively porosity free and were less susceptible to complications such as shrinkage and internal necking that occurs during subsequent heat-treatment in material produced in powder form. Experimental results indicate that the basic form of friction curve for polycrystalline aluminum (worn against a stainless steel slider) remains relatively unchanged within the broad range of grain sizes. The peak value of the coefficient of friction and the rate of wear (severe wear) at the initial stages of sliding decrease with decreasing the grain size. However, the wear rates at the later stages of sliding (mild wear) and the steadystate value of coefficient of friction are less sensitive to the initial grain size (and hence to the bulk hardness of samples).

310

Z.N. Farhat et al. / Materials Science and Engineering A206 (1996) 302-313

(a)

(b)

(c)

(d)

Fig. 7. Worn surface morphologies.(a) Coarse grain size aluminum, sliding distance 22 m; (b) 16.4 nm, sliding distance 22 m; (c) coarse grain size aluminum, sliding distance 90 m; (d) 16.4 nm, sliding distance 90 m. The shape of friction curve shown in Fig. 4 is not peculiar to aluminum, but also observed in other tribological systems in which the bulk hardness of the stationary slider is higher than that of the rotating sample [30]. There have been numerous studies [31-35] to rationalize the time and sliding distance dependent processes that control the shape of the friction curves and friction transitions since the early work of Bowden and Tabor [34]. For example, Suh [31] has suggested that the magnitude of the coefficient of friction and time to reach various characteristic curve features are controlled by the combined effects of surface deformation (/~d) plowing by wear particles and hard surface asperities (Ppl) and adhesion between flat surfaces (p,). Based on the metallographic observations, the time-dependent features of the friction curves in Fig. 4 can be qualitatively discussed as follows. As indicated before, surface deformation of aluminum starts as soon as the sliding process commences. The surface of coarse grained aluminum exhibits signs of heavy damage resulting from severe plastic deformation and plowing (Fig. 7(a)). These processes would increase the magnitude of /~d and ~t/pl and hence the overall value of the coefficient of friction. In the nanocrystalline samples with surface hardness about six times larger than the coarse grained aluminum deformation and microplowing cannot occur as readily (Fig. 7(b)) so that the sliding of the pin can progress with a

reduced amount of energy expenditure. Attention should also be given to the role of surface films on the coefficient of friction. Experiments performed in vacuum clearly indicate that the coefficient of friction is higher (about 30%) when the surfaces are clean and free of oxides, such as aluminum and iron oxides, which may provide a solid lubrication action (Table 3). The stress-strain distribution in the material adjacent to the contact surfaces changes during the wear process. The hardness of the subsurface region continuously increases (Fig. 8) first at a high rate then steadily from 370 to 550 MPa. Thus, plastic deformation and associated surface damage processes become progressively more difficult to operate which in turn translate to a drop in the coefficient of friction and increase in the wear resistance. Therefore, at the later stages of sliding the damage processes depend largely on the properties of the material within the deformed subsurface zone rather than the initial material properties before the test. This trend, i.e. the stronger dependence of the tip t o the (bulk) hardness, is shown in Fig. 9. Previous work has shown that subsurface hardening can improve both abrasive [35] and dry sliding wear [36] resistance of aluminum. Kato et al. [37] observed that the dislocation density in the subsurface region increases during sliding of aluminum from about l08 to 1 0 z° c m - 2 (at a depth of 0.2 pm). The authors argued

Z.N. Farhat et al./ Materials Science and Engineering A206 (1996) 302-313 Table 4 Comparison of Hall-Petch parameters for aluminum

570

X7 4 540

m Material

so =1":1 [] .o =1 t

•

'510

•

90

m

% (MPa)

k (MPa m °-5)

Grain size

Ref.

(~m)

-~

99.97% A1 99.999% A1 99.99% AI t100 AI AI 6'<0 Ni 99.99% AI

~-~480

~

311

450

420

15.0 15.5 22.4 14.3 11.0 11.3

0.07 0.04 0.07 0.07 0.14 0.07

[25] [261 [271 [281 [291 this work

20 150 150 250 30-100 20 200 0.5-20 0.015 ~0.1

390 1.6

360 = ,

330 (0)

,

,I,

,

20 30 40 DEPTH BELOW THE WORN 0

10

50 60 70 SURFACE (/~m)

Z; 0

. . . .

I ' ' ' ' 1 ' ' ' ' 1

. . . .

I ' ' ' ' 1

. . . .

I . . . .

1.4

I,=-I

iiiii

1.2

¸

1.0

600

' ' ' ' 1 ' ' ' ' 1 ' ' ' ' 1 ' ' ' ' 1 ' ' ' ' 1 ' ' ' '

C)

"]DEPTH BELOW WORN SURFACE=I0 ~ m I

E.-,

0.8

Z ._.550

U

0.6

0.4 "-'500 0

0.2 ....

0.0

Z 450

0.0

~ ....

0.5

i ....

1.0

i

.',-..-i'.-,'.','r:-r-~-~4--

1.5

2.0

2.5

......

3.0

:].5

H -I (GPa -I) 400

Fig. 9. Dependence of coefficientof frictionon hardness (O)/~p; (()) /As.s/

350

i1,1,,,,I,,,,1,,,,I,,,,I,,,,

20

(b)

40

60

OO

100

120

SLIDING DISTANCE ( m )

Fig. 8. (a) Hardness gradients below the worn surface in aluminum with a coarse grain size.; (b) variation of subsurface hardness with sliding distance.

that steady-state conditions are established when a cell structure begins to form. Similar to the coefficient of friction in the early stages of sliding, the bulk hardness of nanocrystalline aluminum appears to be a good index o f the wear resistance. Fig. 10(a) shows the variation of the wear rates within the 15-100 nm grain size range with the hardness. A simple linear regression analysis of the data in Fig. 10(a) indicates that the rate of wear in the severe wear regime (Ws) obeys the following expression Ws = 0.85 x 10 3 + (1.73 x 10 3 ) H

1

x 10-6+(1.74x

10-4)H

1

Ws = 0.85 x 10 3 + (2.44 x 10-4)d °5

(9a)

and for the mild wear Win=0.69 x 1 0 - 6 + ( 0 . 2 5 x 10 4 ) d ° 5

(9b)

The linear relationship between Ws, Wm and the grain size is shown in Fig. 10(b). Eq. (8), relating the volume of the material worn per unit sliding distance, W, to the hardness H of the nanocrystalline material, is an Archard-type wear (Eq. (2)). Combining Eqs. (8) and (9), Archard's Wear Law can be re-expressed in the formalism of the H a l l - P e t c h relationship, i.e. W = Wo + K ( H ° + kd_0.5)

(10)

(8a)

The mild wear rates are less sensitive to the initial bulk hardness. The mild wear rate (Wm) versus H plot can be described as follows Wm=0.69

We can also establish a relationship between the grain size and wear rates, i.e. for the severe wear

(8b)

Eq. 10 describes the effect of microstructural refinement on the wear rates. The constant K is termed wear coefficient and is related to the severity of the sliding wear process. According to Eq. (10), under the ambient conditions the wear coefficients of the nanocrystalline

312

Z.N. Farhat et al. / Materials Science and Engineering A206 (1996) 302-313 400

,,,,i,,,,i,,,,i,,,,i,,,,i,,,,lO,,,i,,,

obtained through a study of the microstructure of the plastically deformed zone beneath the contact surfaces and by determining stress-strain states in the deformed material near the contact surfaces.

350 t~

300 250 tt~ I

5. Conclusions 200

The main conclusions from this study were as follows:

150 E-

100 50 0 o.i

0.2

0.4

0.6

0.8

1.o

1.2

1.4

1.6

H -1 ( G P a - 1 )

(a)

400

' ' ' ' I ' ' ' ' I

....

| ' ' ' ' I ' ' ' ' I ' ' ' ' I

....

I "°'

350 ¢'2

300

•- - ,

250

v~

200

tt~ I

150

1oo 50 o

. . . . . . . . . . .

4 (b)

5

6

. . . . . . . . . . .

7

8

, ....

9

,

I0

11

d ° ' 5 ( n m °'5)

Fig. 10. (a) Wear rate versus hardness (I/H) in the nanocrystalline range; (b) Wear rate versus grain size (d°5) in the nanocrystalline range (O) severe wear; (V) mild wear. aluminum are 1.73 x 10 -3 and 1.74 x 10 - 4 for severe and mild wear regimes respectively (note that here the applied load L = 1.0 N). The term Wo represents the grain size independent component o f the wear rate and reflects the fact that since the largest hardness achievable by grain reduction in nanocrystalline aluminum is likely to be less than that o f the steel counterface (4.5 GPa) the surface of aluminum will .always be subjected to wear. Although Eq. (10) establishes a simple empirical relationship between the wear rates and the grain size in t e r m s of H a l l - P e t c h parameters, it is important to emphasize that bulk hardness is only one of the many material parameters that influence the wear resistance. A more complete understanding of the friction wear mechanisms in nanocrystalline materials could be

(1) Nanocrystalline aluminum films were produced by an r.f. magnetron sputtering technique. Grain size was varied between 15-100 nm by isothermal annealing. Samples showed a uniform distribution of equiaxed grains and were free of porosity. (2) Within the grain size range of 15-100 nm, the hardness is inversely dependent on the square root of grain size, i.e. it follows a H a l l - P e t c h type relationship. (3) The coefficient of friction rises to a peak value (/iv) after a short sliding distance then settles down to a steady-state (/~s.s.). /~p decreases about 55% with decreasing the grain size from 106 to 16.5 nm. Coefficients of friction become higher when the sliding test is performed in vacuum. (4) Similar to the friction curves, wear rate versus sliding distance curves revealed a transitional behavior from severe wear to mild wear above a critical sliding distance. Severe wear (and /Zp) was sensitive to both the grain size and the bulk hardness of the material. (5) Based on the experimental evidence, a modified Archard-type equation which incorporates the effect of grain size is proposed in order to describe the wear rates of nanocrystalline aluminum. (6) The severity of wear damage, i.e. plastic deformation, microplowing and microfracture, decreases during sliding as the subsurface material undergoes hardening. Consequently, mild wear rates and values of steady-state coefficient of friction become less sensitive to the initial bulk hardness of aluminum.

Acknowledgments

This work is supported by The Natural Sciences and Engineering Research Council of Canada (NSERC) by a Strategic grant.

References

[1] E.O. Hall, Proc. Phys. Soc., London, B64 (1951) 747. [2] N.J. Petch, J. Iron Steel Inst., 174 (1953) 25. [3] R.W. Armstrong, in T.N. Baker (ed.), Yield, Flow and Fracture of Polycrystals, Applied Science, London 1983.

Z.N. Farhat et al. / Materials Science and Engineering A206 (1996) 302-313

[4] G.D. Hughes, S.D. Smith, C.S. Pande, H.R. Johnson and R.W. Armstrong, Scripta Metall., 20 (1986) 93. [5] G.W. Nieman, J.R. Weertman and R.W. Siegel, Seripta Metall., 23 (1989) 2013. [6] J.S.C. Jang and C.C. Koch, Scripta Metall. Mater., 24 (1990) 1599. [7] R.Z. Valiev, F. Chmelik, F. Bordeaux, G. Kapelski and B. Bandelet, Scripta Metall. Mater., 24 (1992) 855. [8] R.W. Siegel, Mater. Sci. Eng., BI9 (1993) 37. [9] Z. Fan, Mater. Sci. Eng., A191 (1995) 73. [10] A.H. Chokshi, A. Rosen, J. Karch and H. Gleiter, Scripta Metall. Mater., 23 (1989) 1679. [11] K. Lu, W.D. Wei and J.T. Wang, Scripta Metall. Mater., 24 (1990) 2319. [12] G.E. Fougere, J.R. Weertman and R.W. Siegel, Scripta Metall. Mater., 26 (1992) 1879. [13] G. Palumbo, U. Erb and K.T. Aust, Scripta Metall. Mater., 24 (1990) 2347. [14] N. Wang, Z. Wang, K.T. Aust and U. Erb, Acta Metall. Mater., 43 (1995) 519. [15] R.Z. Valiev, N.A. Krasilnikov and N.K. Tsenev, Mater. Sei. Eng., A 137 (1991) 35. [16] J.F. Archard, J. Appl. Phys., 24 (1953) 981. [17] L.G. Korshunov, A.M. Polyakova, N.L. Chernenko and V.M. Umova, Phys. Met. Metall., 61 (1986) 160. [18] D.A. Rigney, L.H. Chen, M.G.S. Naylor and A.R. Rosenfield, Wear, 100 (1984) 195.

313

[19] S.K. Ganapathi and D.A. Rigney, Seripta Metall., Mater., 24 (1990) 1675. [20] Y. Ding, Z. Farhat, D.O. Northwood, A.T. Alpas, Sur[i Coat. Technol., 68/69 (1994) 459. [21] W. Rachinger, J. Sci. Instrum., 25 (1942) 254. [22] C.N. Wagner and E.N. Aqua, Adv. X-Ray Anal., 7 (1963) 46. [23] T.J. Bell, J.S. Field and M.V. Swain, Mater. Forum. 17(1993) 127. [24] M.F. Doerner and W.D. Nix, J. Mater. Res., l (1986) 601. [25] P.R. Carreker and W.R. Hibbard, Trans. Met. Soc., A1ME, 209 (1957) 863. [26] N. Hansen, Acta Metall., 25 (1977) 863. [27] H. Fujita and T. Tabata, Acta Metall., 21 (1973) 335. [28] A.W. Thompson and M.I. Baskes, Phil. Mag., 28 (1973) 301. [29] D.J. Lloyd, Metal Science, May (1980) 193. [30] N.P. Suh and H.C. Sin, Wear, 69 (1981) 91. [31] N.P. Suh, Tribophysies, Prentice-Hall, New Jersey, 1986, p. 63. [32] P.J. Blau, J. Tribology, 109 (1987) 537. [33] J.P. Hirth and D.A. Rigney, in F.R.N. Nabarro (ed.), Dislocations in Solids, North-Holland, Amsterdam 1983, p. 3. [34] F.P. Bowden and D. Tabor, Friction and Lubrication of Solids~ Part II, Clarendon, Oxford, 1964, p. 52. [35] M.A. Moore, in D. Scott (ed.), Wear (Treatise on Materials Science and Technology, Vol. 13, Academic Press, New York, 1979, p. 217. [36] J. Zhang and A.T. Alpas, Mater. Sei. Eng., (1993) 25. [37] K. Kato, T. Kayaba and Y. Ono, in K.C. Ludema (ed.), Wear ~/' Materials, ASME, New York, 1985, p. 463.