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Wear transition diagram for silicon carbide
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Tribology
Vol. 28, No. 8, pp. 559-512. 1995 Copyright 0 1995 Elsevier Science Ltd Printed in Great Britain. All rights reserved 0301-679X195/$10.00 +O.OO
International
N
0301-679X(95)00082-8
ar transition silicon carbide X. Dong*,
S. Jahanmir+
diagram
for
and L.K. Ives’
To obtain information on the tribological behaviour of silicon carbide at elevated temperatures, unlubricated ball-on-flat wear tests were conducted on sintered silicon carbide in self-mated sliding in air. The contact load was varied from 3.2 to 98.0 N, and a temperature range of 23°C to 1000°C was used. Scanning electron microscopy, Fourier transform infrared spectroscopy and energy-dispersive spectroscopy were used to elucidate the wear mechanisms. The results of the tests and observations were employed to construct a wear transition diagram, which provides a sumn-ary of tribological information including friction coefficient, wear coefficient and wear mechanisms as a function of temperature and load. The wear transition diagram for the sintered silicon carbide studied is divided into four regions plus one transiTion zone. At room temperature, under high loads and high environmental humidity, the tribological behaviour is controlled by tribochemical reactions between the silicon carbide surface and water vapour in the environment. Under low loads and at temperatures below 25O”C, wear occurs by ploughing and polishing. At temperatures about 250°C and under low loads, tribooxidation and formation of cylindrical wear particles control the triboltsgical behaviour. Wear occurs by microfracture when the load is increased above a critical value; and both the friction coefficient and the wear coefficient increase. Keywords:
sintered
silicon
carbide,
tribo-oxidation,
Introduction Silicon carbide materials, due to their attractive mechanical properties and chemical stability, are being i-lcreasingly used in high-temperature tribological systems. However, effective implementation of these materials in new applications requires detailed information on their tribological behaviour. In particular, information on friction coefficient and wear rate under * Graduate Research Assistant, Deuartment of Mechanical Engineering, University of Maryland at dollege Park, College Park, MD 20742, USA T Ceramics Division, National Institute of Standards and Technology, Gaithersburg, MD 20899, USA
Tribology
microfracture
a broad range of conditions is needed to ensure efficient and reliable performance. Furthermore, an understanding of the wear mechanisms and the relationship between these mechanisms and the microstructure of the material is needed to develop improved materials with enhanced performance. Several studies of the tribological behaviour of silicon carbide materials at high temperatures in self-mated slidinglP4 and in sliding against other ceramics or metalP have been published. The influence of temperaturel, normal loadlo, sliding speedll, environment12, and chemical composition of the contacting materials13 were investigated. Several wear mechanisms, for example, microfracture14, plastic deformation,
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ploughing, and the formation and removal of oxidation products l5 , have been identified. In addition to this information, sudden changes in the tribological behaviour, such as an increase in the friction coefficient and the wear rate, have been reported2z9. The mechanisms for such transitions, however, have not been identified. There is considerable variation in the published data, especially regarding the influence of temperature, on the tribological behaviour of silicon carbide materials, For example, the friction coefficient during self-mated sliding at temperatures below 400°C has been reported to range from 0.16 to 0.45’z9, from 0.3 to 0.62, or from 0.7 to 0.84. The wide variation in the published data is apparently due to differences in the composition and type of silicon carbide studied and in test parameters such as contact geometry, sliding speed, and applied load. These disparate results suggest that a systematic approach is needed to elucidate the effects of test parameters and microstructure on the tribological performance and wear mechanisms. In two recent studies on alumina16 and silicon nitridei’, contact load and temperature were systematically varied to obtain data and mechanistic information on the effects of these parameters on the tribological behaviour. The results were compiled into ‘wear transition diagrams’ to show the effects of load and temperature on the wear mechanisms and the values of friction coefficient and wear rate. A similar approach has been adopted for the present study on silicon carbide.
Experimental
procedure
Wear tests were conducted with a linear-reciprocating ball-on-flat high-temperature tribometer, described in a previous paperr6. The material used for both the balls and the flats was sintered alpha silicon carbide. Selected properties of this material are summarized in Table 1. The chemical composition and microstructure of this material have been reported in several references18~22. Boron and carbon are used as sintering aids, resulting in a final composition of 0.5 wt% B and 0.42 wt% free carbon in the sintered material19. The microstructure consists of silicon carbide grains, Table 1 Selected properties carbide (Hexoloy SAP
of sintered
Bulk density (g/cm) Hardness ( GPajb Young’s modulus (GPa) Shear modulus (GPa) Poisson’s ratio Flexure strength (MPa) ‘Fracture toughness (MPa ml’*) Coefficient of thermal expansion Thermal conductivity (W/m/K)
(lO-6/oC)
silicon
3.15 36 406 26 0.14 400 3.7 4.02 120
“The information on the manufacturer or supplier is included here to identify the material. This does not imply endorsement of the product by the National Institute of Standards and Technology “The hardness was measured as a part of the present study. The other values were obtained from reference 18
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with an average grain size of 4 km, and about 8 ~01% of what appear to be ‘dark’ features in the SEMi9. These regions have been identified as amorphous carbon, B4C, and void?. The average diameter of the boron carbide regions were determined to be around 10 Frn; and the silicon carbide grains were found to be saturated with boron in solid solution”‘. Impurities in this material have been reported as 0.2% Fe and 0.06% A115; and the density has been reported to be 98% of the theoretical valueis. The flat specimens used for this investigation were 12.7 mm squares with a thickness of 3 mm. After final polishing with 1 pm diamond paste, the flat test surface had a roughness of 0.1 pm rms. The ball specimens were 12.7 mm in diameter, and were used in the asreceived condition with a roughness of 0.2 p,m rms. Prior to testing, the polished flats and the balls were cleaned to remove surface contamination. The cleaning procedure consisted of the following steps: rinsing twice with hexane, then ultrasonic cleaning in fresh hexane, followed by ultrasonic cleaning with acetone, and final rinsing with hexane. The experiments were conducted at a relatively slow sliding speed of 1.4 x low3 m/s to minimize frictional heating. A stroke length of 10 mm was used in all tests. The normal load was varied from 3.2 to 98.0 N; and a temperature range of 23°C (room temperature) to 1000°C was used. Each test was conducted for a total sliding distance of 10 m (500 reciprocating cycles of 1000 passes) in laboratory air. The relative humidity, measured at room temperature, varied from 50% to 70%. Since the test samples were located in a partially open system, both the relative humidity and the absolute humidity decrease rapidly as the test temperature is increased. Prior to testing, several thermocouples were used to monitor the temperature of the ball, the flat surface, the heated chamber and the specimen holder. These data and the output of a thermocouple positioned under the flat specimen were used to calculate the surface temperature of the flat. Before initiation of sliding, the specimens were allowed to stabilize at the test temperature. The friction force was monitored during the tests with load cells, which measured the tangential force exerted on the stationary ball. The friction coefficients reported in this paper are the average steady state values for the final 2.5 m of sliding. The wear volume on each flat specimen was calculated from surface profile traces (usually three to five) across the wear track and perpendicular to the sliding direction. These data were then used to calculate the wear coefficient, which is a dimensionless quantity defined as wear volume per unit sliding distance per unit load normalized by the indentation hardness. Wear coefficient is used in this paper for convenience; its use does not imply a relationship between wear and hardness. The indentation hardness values were measured as a function of temperature following the procedure described previously 16. These results, shown in Fig 1, indicate that hardness decreases as the temperature is raised. The hardness values at each test temperature were used for the calculation of the wear coefficients. After wear testing, the wear tracks were examined by
Volume 28 Number 8 December 1995
Wear
40
transition
diagram
0 e
I
7
1.0
for silicon
3.2 4.9
.
N N
carbide:
19.6 39.2
N N
* -
X. Dong
70.4 98.0
N N
,
1
et al.
.
I
.
1
800
1000
0.0I 04
I C
200
/
/
400
600
TEMPEHATURE
0
I 000
200
400
600
TEMPERATURE
1000
("C)
('C) o 3.2 o 49N
Fig 1 .Effect of temperature on Vickers hardness of sinterea silicon carbide determined with a 9.8 N load in vacuum of 10p3 Pa
N
.
19.6 39.2
N N
. .
78.4 98.0
N N
1o-2
scanning electron microscopy (SEM) to elucidate the wear process. Energy-dispersive X-ray spectroscopy (EDXS ) and micro-Fourier transform infrared (FTIR) spectroscopy (in reflectance mode) were used to analyse the wear tracks and the wear debris.
Results
(b)
t
t’
and discussion
1o-5-f
The friction coefficients and the wear coefficients are plotted as a function of temperature and load in Fig 2. The data in this figure indicate that the effect of load or. the results can be separated into two groups, where I he friction coefficients and the wear coefficients are load-independent within each group. In Fig 2, a solid curve is drawn for the results obtained at loads lower tian 10 N, and a dashed curve is used for the values obtained at higher loads. At low loads, the average friction coefficient increases slightly as the temperature is raised. This is followed by a sharp drop in the friction coefficient from 0.75 to 0.3a at about 200-300°C. This drop in friction is an indication of a sudden change in the controlling mechanism. As the temperature is increased further, the friction coefficient slowly increases. The wear coefficient, however, is not as sensitive as the friction coefficient to the changes in temperature. The f&ion coefficient at the higher loads ranges from 0.20 to 0.26 at room temperature and quickly increases to about 0.80 as the temperature is raised. The friction coefficient attains a maximum value at approximately 7OO’C, and then drops as the temperature is increased further. At high loads, the friction coefficient and the wear coefficient follow a similar trend as a function of temperature, with the exception of the data at temperatures higher than 7OO”C, where the friction coefficient falls and the wear coefficient rises with increasing temperature. These results show that the tribological behaviour of sinterel? silicon carbide is complex and should be explored in greater detail. The results of the wear Tribology
/ 0
I
I 200
/
1
)
400
TEMPERATURE
/
/
I
600
800
1000
("C)
Fig 2 The effect of temperature and load on (a) friction coefficient and (b) wear coefficient. The solid curves are drawn through the data points at loads less than 10 N; the dashed curves are drawn through the data points obtained at loads larger than 10 N
tests and analysis of the wear mechanisms by SEM, EDXS, and FTIR were used to construct the wear transition diagram shown in Fig 3. The diagram shows the effect of load and temperature on the tribological behaviour of the sintered silicon carbide used in this investigation. Based on the results, which will be presented shortly, the wear transition diagram in Fig 3 is divided into four regions (identified as I to IV) and a transition zone (the cross-hatched area). The controlling mechanism in each region is distinct: tribochemical reaction between the silicon carbide and water vapour in the environment in region I, ploughing and polishing in region II, tribo-oxidation and formation of cylindrical-shaped wear particles in region III, and microfracture in region IV. The last is separated into two parts by the transition zone. The average values of the friction coefficient, f, and the wear coefficient, K, for the flats in each region are indicated in the figure. The standard deviation for the friction coefficient values listed in each region was approximately equal to 10% of the mean value. The standard deviation for the wear coefficient, however, was larger; it ranged from 40% to 60% of the mean value in each
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f=0.23,
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et a/.
. .o
K=lO-'
'0°4[
I1
z
,I.::; 0
200
400
504 P 2 B 0.2
t 0
800
600
TEMPERATURE
i
20
40 60 LOAD (N)
80
100
1000
("C)
Fig 3 Wear transition diagram for sintered silicon carbide, showing four distinct regions plus one transition zone. The friction coefficient, f, and wear coefjicient, K, for each region are indicated in the figure
Y9 6 N, 23Oc t
region. In the following sections, the test results and the observations on the wear surfaces are analysed to eIucidate the underlying mechanisms in each region. Region
I: tribochemical
reactions
0.0
At room temperature and loads larger than 10 N, the friction coefficient ranges from 0.20 to 0.26 with an average value of 0.23. The wear coefficient ranges from 2 x 10Vs to 2 x 10P3 with an average value of 4 x 10P4. This region is observed only at relative humidities greater than 60%. The friction coefficient in this region is independent of load, as shown in Fig 4(a). A typical plot of the friction coefficient versus sliding distance in region I is shown in Fig 5(a). The friction coefficient starts at about 0.75 and rapidly decreases to a steady-state value of 0.22. The wear tracks on the flat specimens tested in this region, shown in Fig 6(a), are smooth with some indications of damage. Observation of the damaged areas at higher magnifications, Figs 6(b) and 6(c), indicates that the surface is covered with a thick film and that the damage is associated with fracture of this film. However, the debris found near the edge of the wear track and those at the end of the track were very small and quite different in appearance. The wear tracks generated in region I were examined with FTIR to determine the composition of the surface films. For reference, a spectrum was first obtained from an area outside the wear track. This spectrum, designated as (a) in Fig 7, shows the SIC peak (due to an IR reflectivity phenomenon). The spectrum (b) for an area inside the wear track of a sample tested in region I is also shown in Fig 7. The major peaks in this spectrum are identified as Si-OH and Si-0 according to reference 23. These peaks suggest the existence of hydrated silicates. The absence of the SIC reflectance peak in this spectrum indicates that the surface film on the wear track is sufficiently uniform 562
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~
40
50 RELATIVE
60 IiUMIDlTY
7:: (X)
80
Fig 4 (a) Effect of load on friction coefficient at room temperature and at 70% relative humidity; (b) effect of relative humidity on friction coefjkient for samples tested at room temperature under a load of 19.6 N and thick to mask the underlying Sic. Since hydrated silicates are usually unstable, the ‘mud cracks’ that were seen on the wear track were probably formed as a result of water loss from the film either after the test or in the SEM. The details for the formation and the structure of the tribochemical reaction products are not known. However, oxidation studies on silicon carbide at elevated temperatures 24.25 have revealed that the oxidation rate increases by addition of water vapour to the oxygen environment. This increase was attributed to the diffusion and solution of OH- ions through the surface oxide film. In our study, the FTIR spectra confirmed the formation of oxidation products by tribochemical reactions. Others have also found evidence for such reactions on the wear tracks in silicon carbide and silicon nitride”6P29. When the relative humidity or the load was decreased at room temperature, the friction coefficient increased from 0.23 to 0.60 (Fig 4). At low relative humidities, the number of OH- ions needed to increase the oxidation rate was probably too small. Therefore, the production rate of hydrated silicate was insufficient to provide a uniform coverage and a low friction coefficient. Examination by scanning electron microscopy of flat specimens tested at lower humidities indicated 8 December
1995
Wear
transition
diagram
for silicon
carbide:
X. Dong
et a/.
8
10
1 .o 19.6
0.0
I 0
I 4
2
(a)
4
SLIDING
N.
23°C
I
6
8
DISTANCE
10
(m)
0.0i 2
4 SLIDING
6
8
DISTANCE
10
0
(m)
2
-4 SLIDING
CC)
6 DISTANCE
(m)
1.o 98.0
N.
19.6
3OO’C
0.c
Cd)
00
-c
0
2
4 SLIDING
6 DISTANCE
8
10
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--0
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4 SLIDING
(m)
Fig 5 iypical plots of friction coeficient (d) region IVa, and (e) region IVb
N,
-I n
t
6 DISTANCE
8
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(m)
(e)
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Fig 6 Scanning electron micrographs of the wear track on the flat specimen tested in region I at 23°C. (a) Sample tested under a load of 19.6 N, showing a smooth wear track with fractured regions; (b) sample tested under a Eoad of 39.2 N, showing microcrack networks on the wear track at a high magnification; (c) higher magnification of an area in (a) that the wear tracks were mostly smooth, but contained areas where material had been removed by microfracture. It will be discussed later that this dual nature of the wear track appearance is characteristic of the transition zone between regions II and IV. Therefore, at low relative humidities, region I becomes part of the transition zone. Accordingly, the boundary between region I and the transition zone in Fig 3 is drawn by a dashed line to indicate this phenomenon. The load effect can be explained by considering the effect of load on frictional heating. Although in our tests a low speed was used to minimize frictional heating, the contact temperature depends on the applied load during sliding. It is postulated that at low loads the contact temperature is too low to allow for the formation of sufficient oxide film for friction reduction. Therefore, the conditions needed for formation of an oxide film for friction reduction consist of high relative humidities and a high load. Since the relative humidity in our tests decreases rapidly as the temperature is raised, any possible benefit from formation of hydrated silicates is lost as the temperature is increased by frictional heating. 564
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Region
II: ploughing
and
polishing
This region extends from room temperature to approximately 250°C at loads lower than 10 N. The friction coefficient ranges from 0.50 to 0.78 with an average value of 0.67, and the wear coefficient varies between 6 X 10W5 and 2 X 10e3 with an average value of 2 x 1OW. Both the friction coefficient and the wear coefficient were found to be independent of load in this region. A typical plot of friction coefficient vs. sliding distance for region II is shown in Fig 5(b). At the initiation of sliding, the friction coefficient starts at a value between 0.80 and 0.90 and decreases slightly during the test. The friction coefficient remains much larger than that observed in region I. The appearance of the wear track on a flat tested in this region is shown in Fig S(a). The wear surface is smooth, but contains several ploughing or abrasion marks. The microstructure, i.e. the silicon carbide grains and ‘dark’ regions, which are associated with boron carbide, carbon and porosityi are also observed. The wear debris collected at one end of the wear track is shown in Fig S(b),
Volume 28 Number 8 December 1995
Wear
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X. Dong
et al.
1 Si C:
I
(e)
,L
1-
(d) (cl
I i c 00 0
I I
I
I
3200
1600
2400
8130
WAVE NUMBER
Fig 7 P’i’TR spectra obtained on: (a) unworn surface, (b) werv track at room temperature and 39.2 N load, region f; (c) wear track at 100°C and 4.9 N load, region i”r; (d) wear debris accumulated at one end of wear trrrck at 500°C and 4.9 N load, region III; and (e) wecr track at 1000°C and 4.9 N load, region III indicating agglomeration of very fine wear particles. The FTIR spectrum (c) in Fig 7 shows no evidence for for:na.tion of silicates or other reaction products. Therefore, based on the ploughing marks observed on the we,ar track, it is hypothesized that wear in this region occurs by ploughing and polishing of the silicon carbide surface. There is, however, insufficient evidence to completely rule out a possible role for oxidation, since a very thin oxide film cannot be detected by FTIR. Regior 111:tribo-oxidation
Fig 8 Scanning electron micrographs of the wear track on the flat specimen tested in region II at 200°C under a load of 3.2 N showing (a) the ploughing marks, and (b) wear debris accumulated at one end of wear track
reactions
This region is observed at temperatures ranging from 250°C :o 1000°C (and possibly higher) and at loads lower than 10 N. The friction coefficient is much lower in this region than in region II. As shown in Fig 2, the friction coefficient in this region slowly increases from 0.30 to 0.50 as the temperature is increased. The wear coefficient, however, remains relatively constant; it ranges from 2 x IO-’ to 5 x lop4 with an average value of 3 X lo-“. A typical plot of friction coefficient versus sliding distance for region III is shown in Fig 5(c). In this particular test, the friction coefficient started at approximately 0.85 and decreased with sliding distance, ending at a steady-state value of 0.32. This reduction in the friction coefficient is an indication for formation of chemical reaction products. To check this possibility and to define the wear mechanism, SEM was used to examin,? the wear surfaces. A typical wear track is shown in Fig 9(a). The surface appears polished, similar to the micrograph in Fig Z(a) obtained for the wear track in region II. However, the number of ploughing marks are fewer and numerous cylindricalshaped wear particles are observed. These wear Tribology
particles have an average diameter around I km and a length of about 10 km. An SEM micrograph showing the debris accumulated at one end of the wear track (Fig 9(b)) suggeststhat the debris were pushed towards the end of the track and that some of the debris were fractured to produce finer and more equiaxed particles. The wear scar on one of the balls tested in this region is shown in Fig 10(a). The wear scar contains the same features as those on the flat; namely, a smooth surface with cylindrical debris oriented perpendicular to the sliding direction. Figure 10(b) shows the wear debris accumulated at the edge of the wear scar. The morphology of these debris is also similar to that observed on the flat in Fig 9(b). Examination of balls tested in other regions of the wear transition diagram also revealed that wear on both the balls and the flats occurred with the same mechanism. To investigate the origin of the cylindrical wear debris, EDXS and FTIR were used. The EDXS spectrum in Fig 11(a) was obtained with the electron beam located on a single cylindrical particle, and in Fig 11(b) with the electron beam located on an adjacent area on the
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Fig 9 Scanning electron micrographs of the wear track on the fiat specimen tested in region III at 500°C under a load of 4.9 N, showing (a) the smooth surface with cylindrical wear particles and (b) wear debris accumulated at one end of the wear scar wear track. The relatively high concentration of oxygen in the particle suggests that tribo-oxidation was responsible for the formation of the cylindrical particles. However, FTIR analysis, for example spectrum (d) in Fig 7, shows no evidence of silicates or other reaction products in the wear debris accumulated at the end of the wear track. Only the SIC reflection peak can be identified in the spectrum. This could be due to the overwhelming abundance of silicon carbide wear particles at the end of the wear track. However, the spectrum (e) in Fig 7 obtained from the wear track tested at 1000°C shows a small Si-0 peak, confirming increased oxidation rate at higher temperatures. Since the sintered silicon carbide used in our study contained boron and born carbide, we expect oxidation of both boron and boron carbide to have occurred and to contribute to the composition of the oxidation products. However, boron could not be detected by FTIR and the sensitivity of the EDXS system to boron is extremely low. Nevertheless, boron oxide can provide lubricatior?‘*“‘, especially in humid environments, by formation of boric acid32. However, the lubrication 566
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Fig 10 Scanning electron micrographs of the wear scar on the ball tested in region III at 600°C under a load of 9.8 N, showing (a) the smooth surface with cylindrical wear particles and (b) wear debris accumulated at the edge of wear scar
property of boron oxide is strongly temperature dependent. The friction coefficient of boron oxide film rapidly increases to a value higher than 0.9 when the temperature is increased from 200°C to 400°C and then quickly reduces to 0.1 at 600”C30.31. The temperature dependence of the friction coefficient in our tests is different from the expected behaviour of boron oxide. This indicates that formation of boron oxide alone does not explain the results at high temperatures in region III. The oxidation behaviour is complex because of the sintering aids and impurities in the material. It has been found that although oxide films grown at elevated temperatures consist primarily of cristobalite, they contain a glassy phase containing the sintering aides and impurities such Al, Fe, Na, K, etc.25. These elements were found to increase the oxidation rate of silicon carbide2’. It is therefore proposed that formation of a mixed oxide film, containing boron oxide and silicon oxide with other sintering aids and impurities, such as carbon, controls the tribological behaviour in region III. 8 December
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X. Dong
et al.
dependence on load and temperature. The average friction coefficient in this region is 0.78 and the average wear coefficient is 2 x 10-3. At higher temperatures, from 600°C to 1000°C in region IVb, the friction coefficient decreasesfrom 0.90 to 0.70 as the temperature is raised; but the wear coefficient increases from 8 x lop4 to 3 x 10-3.
400
0.0
0.5
1.0 ENERGY
1.5
2.0
Figure 5(d) shows a typical plot of friction coefficient versus sliding distance for region IVa. The friction coefficient quickly reaches a steady-state value of 0.78 without much variation in the average value. In contrast with this plot, the friction coefficient for region IVb, shown in Fig 5(e), starts at a low value and quickly rises to 0.90, then drops to a lower value and stabilizes at around 0.70. Although the average steady-state value of friction coefficient is lower in region IVb, the friction trace is much more erratic and variable than the friction trace observed for region IVa.
2.5
(keV)
(b)
0 ;> 0
0.5
1.0 ENERGY
1.5
2.0
2.5
(keV)
Fig 11 EDXS spectra qf wear track for a ball tested in region 111 at 400°C under 58.8 N: (a) on a cylindrical wear particle and (b) on an adjacent area on the wear track Formation of cylindrical wear particles or ‘rolls’ has been observed by others for different ceramics tested under various test conditions9.16,17J6.Z9J3x34, where the compcsition of the rolls were found to depend on the material: test condition and environment. The process for the formation of these rolls is not entirely understood. Microscopic examination of rolls has shown that tl-ey consist of an aggregation of small crystallites bonded together with an amorphous oxide phase29.j3.34. Irrespective of the process for the formation of rolls, the existence of such cylindrical particles is accompanied by a relatively low friction coefficient. Althorngh the process for the reduction of friction by the roj.ls is not clear, it has been suggested that the cylindrical particles act as miniature rolling element bearings, providing a low friction9.‘9%34. Whether this explacation can be used for rolls generated with different ceramics tested under all sliding conditions is not known. Region
IV: microfracture
This r:gion is observed at loads larger than 10 N and is separated by the transition zone into two parts, designated regions IVa and IVb. At lower temperatures, from 50°C to 400°C in region IVa, both the frictio-n coefficient and the wear coefficient show little Tribology
A typical SEM micrograph for a flat specimen tested in region IVa is shown in Fig 12(a), which indicates that the wear track is somewhat rough and it is aPmost completely covered by wear debris. The structure of the compacted debris is shown at a higher magnification in Fig 12(b). This figure shows that what appears as plastic flow at the lower magnification is actually agglomeration of the fine wear debris. The debris layer was found to be strongly attached to the wear track since it could not be removed by ultrasonic cleaning. The micrograph in Fig 12(c) shows an area of the wear track where the debris layer has been removed by wear prior to ultrasonic cleaning. This micrograph suggests a microfracture dominated wear mechanism for the generation of wear particles in tkis region of the wear transition diagram. The wear tracks for the flat samples tested in region IVb were also covered with a debris layer, shown in Fig 13(a). In contrast to Fig 12(a), this micrograph shows that the debris forms long ridges along the sliding direction. The higher magnification micrograph irn Fig 13(b) shows that the debris ridge is composed of compacted wear particles. Most of the fine debris was removed by ultrasonic cleaning (Figure 13(c)); but the debris ridges were not removed, suggesting that they were strongly attached to the wear track. The structure of a region in the middle of the wear track shown in Fig 13(d) suggests that wear has occurred by microfracture. Therefore, the primary difference between regions IVa and IVb is in the behaviour of the wear debris. At lower temperatures, the wear debris forms an adherent layer almost completely covering the wear track, whereas at the higher temperatures, most of the debris is loosely held with the exception of the adherent ridges along the sliding direction on the wear track. In both cases, however, the wear particles are generated by microfracture. Milling, oxidation and adhesion of the fractured particles contribute to the formation of the compacted debris layer in region IVa and debris ridges in region IVb. In the latter, formation and removal of such debris ridges could produce the observed erratic friction trace.
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Fig 12 Scanning electron micrographs of the wear track on the flat specimen tested in region Na at 300°C under a load of 98.0 N: (a) the general appearance of the wear track, (b) higher magnification of an area in (a) showing the structure of compacted debris, and (c) higher magnification of an area in {a) showing microfracture
Transition
zone
The transition zone is shown in Fig 3 as the crosshatched area. In addition to this area, region I becomes part of the transition zone at low relative humidities. In the transition zone, wear occurs by a combination of the mechanisms observed in the two neighbouring regions; one mechanism is enhanced and the other is diminished as the test variables (load and temperature) are varied. Because of this combination of wear mechanisms, it was found that the values of friction coefficients and wear coefficients in the transition zone were in between those expected for the two neighbouring regions. In the transition zone between regions II and IV, as the applied load is increased the dominant mechanism changes from plowing to microfracture. The micrographs in Fig 14 illustrate this dual mechanism. Figure 14(a) shows that the wear track contains smooth regions as well as fractured ones. These regions are shown at a higher magnification in Figs 15(b) and 14(c) confirming wear by ploughing and by microfracture, respectively. 568
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In the transition zone between regions III and IV the dominant mechanism changes from tribo-oxidation reactions to microfracture. Figure 15(a) shows the features associated with the two mechanisms. These areas are shown at a higher magnification in Figs 15(b) and 15(c). The smooth area with a few cylindrical particles on the wear track, in Fig 15(b), is typical for region III where tribo-oxidation was observed. Figure 15(c), however, shows a fractured surface structure in a rough area, which is typical of wear by microfracture in region IV. Summary Unlubricated wear tests were conducted in air to investigate the effect of temperature and load on the tribological behaviour of sintered silicon carbide in self-mated sliding tests. The contact load was varied from 3.2 to 98.0 N, at a constant but low sliding speed of 1.4 x 1O-3 m/s to minimize frictional heating. The specimen temperature was varied from 23°C to 1000°C. The wear surfaces and wear debris were subsequently analysed by SEM, FTIR and EDXS to elucidate 8 December
1995
Wear
transition
diagram
for silicon
carbide:
X. Dong
et a/.
Fig I3 Scanning electron micrographs of the wear track on the fiat specimen tested in region IVb at iOOo”C under a load of 19.6 N: (a) The general appearance of the wear track showing wear debris compaction in ridges along the sliding direction, (b) higher magnification of an area in (a) showing the structure of the compacted debris, (c) the general appearance of the wear track after ultrasonic cleaning, and (d) higher magnification of an area in (c) showing microfracture
the w:ar mechanisms. The results of the tests and observations were used to construct a wear transition diagram: which provides a summary of tribological inforn ation including friction coefficient, wear coefficient and wear mechanisms as a function of temperature a.ld load. The wear transition diagram for the sintered silicon carbide studied is divided into four regions plus one transition zone. At room temperature and al loads above 10 N, formation and removal of hydrated silicate films were found to be the dominating wear processes, resulting in an average friction coefficient of 0.23 and a wear coefficient of lo-‘. However, this region was only observed at high relative humidities. At loads below 10 N and temperatures up to 250°C. the dominant wear mechanism was found to be plonghing and polishing. In this region, the average friction coefficient was 0.67 and the wear coefficient was approximately lo-“. In the temperature range from 250°C to 1000°C at loads below 10 N, the tribolagical behaviour was controlled by tribo-oxidation and formation of cylindrical wear particles. The friction coefficient in this region was 0.40 and the wear Tribology
coefficient was approximately 10PJ. When the load was increased above 10 N, microfracture dominated the wear process, resulting in a high friction coefficient of 0.78 and a higher wear coefficient of approximately 10-3. The wear transition diagram proposed for silicon carbide in this paper and for alumina and silicon nitride in the two previous papers’6.‘7 can be used to obtain tribological information for each material. These diagrams contain information on the friction coefficient, wear coefficient and wear mechanisms as a function of load and temperature. Since in most engineering applications a low-friction coefficient and a low wear rate are required, the wear transition diagrams can be useful in determining the range of contact conditions where an acceptable level of performance can be expected. For example, for the sintered silicon carbide used in this study in self-mated tests, high loads and high temperatures should be avoided since the friction coefficient and the wear rate are excessive under these conditions.
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Fig 14 Scanning electron micrographs of the wear track on the flat specimen tested in the transition zone at 100°C under a load of 19.6 N: (a) Overall view of the wear track showing smooth areas and fractured regions, (b) higher magnification of the smooth area in (a) showing ploughing and (c) higher magnification of the fractured area in (a) showing microfracture
It should be pointed out that the wear transition diagram presented in this paper has been obtained for one specific material and is based on tests conducted at a low sliding speed. The effect of such factors as chemical composition, microstructure, sliding speed and contact geometry was not investigated. Any significant change in these factors could alter the shape of the diagram. For example, the boundary between regions II and III could be sensitive to the composition of the material, and compositional changes may shift this transition boundary. The effect of speed, however, is complex. First, an increased speed would increase the rate of tribochemical reactions, but at the same time, it could promote microfracture through increased thermoelastic stresses. The effect of these factors on the tribological behaviour of silicon carbide requires further study.
The authors would like to thank Dr Ramesh Divakar of Carborundum Company for providing the test samples and Dr Hong Liang of the University of Maryland for obtaining the FTIR spectra.
References 1. Gee M.G., measurement temperature.
Matharu C-S., Almond E.A. of sliding friction and wear Wear 1990, 138, 169-187
and Eyre T.S. The of ceramics at high
2. Dufrane K.F. and Glaeser W.A. Wear of ceramics in advanced heat engine applications. In: Wear of Materials, Ludema K.C. (Ed.), American Society of Mechanical Engineers, New York, 1957, pp. 285-291 3. Mivoshi K. and Bucklev D.H. Considerations in friction and wear, In: New Directions in Lubrication, Materials. Wear, and Surface Inreracrions. Loomis W.R. (Ed.), Noyes Publications, Park Ridge. NJ, 1985, pp. 252-319
Acknowledgements
4. Yust C.S. and Carignan F.J. Observations on the sliding of ceramics. ASLE Trans. 1984. 28, 245-252
Financial support received from the National Science Foundation under Grant No. MSS-9005868 to the University of Maryland is gratefully acknowledged.
5. Gray S. Friction and wear of ceramics pairs under high temperature conditions representative of advanced engine components. Ceramic Engineering and Science Proceeding 1985, 6. No. 7-8, 965-974
570
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carbide:
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et al.
Fig 1.5 Scanning electron micrographs of the wear track on a fiat tested in the transition zone at 500” under a load of 19.6 N: (a) Overall view of the wear track showing smooth areas and fractured regions, (b) higher magni$cation of the smooth area in (a) showing cylindrical wear particles, and (c) higher magnification of the fractured area in (a)
showing microfracture
6. Miyoshi K., Buckley properties of sintered carbide. Am. Ceram.
D.H. and Srinivasan M. Tribological polycrystalline and single-crystal silicon Sot. Bull., 1983. 62, 494-500
7. Slirey H.E., Jacobson T.P., Deadmore D. and Miyosbi K. Tribology of selected ceramics at temperatures to 900°C. Ceramic Et@zeering and Science Proceedings 1986,7, No. 7-8,1039-1051 8. Mi!poshi K. and Buckley D.H. Tribological properties and surface chemistry of silicon carbide at temperatures to 1500°C. ASLE Trcns., 198:. 26. 53-63 9. Both P., Platen F. and Kapelski G. Effect of temperature on tribological properties of Al,O, and Sic ceramics. In: Euro-Ceramics, Vol. 3. deWith G., Terpstra R.A. and Metselaar R. :Eds), Elsevier Applied Science, New York, 1989 10. Derby J., MacBetb J. and Seshadri S. Tribological of alpha silicon carbide engine components, IMechE 133-137 11. Denap J. and Lamon J. Sliding Sci 1990, 25, 3592-3604
friction
behaviour 1985. C65,
of ceramics.
J. Mat.
12. Ajayi 0.0. and Ludema K.C. Formation of transfer film during ceramics/ceramics repeat pass sliding. In: Wear of Materials, Lul;iema K.C. (Ed.), American Society of Mechanical Engineers, New York, 1989, pp. 349-359 13. Adewoye 0.0. and Page T.F. Frictional deformation and fracture in solycrystalline Sic and Si,N,. Wear 1981, 70, 37-51
Tribology
14. Miyoshi sliding
K. and Buckley D.H. contact. ASLE Trans.
Ceramic wear in indentation 1984. 28, 296-302
and
15. Cranmer D.C. Friction and wear properties of monolithic silicon-based ceramics. J. Mater. Sci. 1985, 20. 2029-2037 16. Dong X. and Jahanmir S. Tribological characteristics alumina at elevated temperatures. J. Am. Ceram. Sot. 74, 1036-1044 17. Dong X. and Jahanmir S. Wear nitride. Wear, 1993, 165, 169-180
transition
diagram
of (Y1991.
for silicon
18. Strom R.S., Boecker W.D.G., McMurtry C.H. and Srinivasan M. Sintered alpha silicon carbide ceramics for high temperature structural application. ASME Paper No. 85%IGT-127, 1985 19. Browning R., Smialek J.L. and Jacobson N.S. mapping of o-Sic by scanning auger microscopy. Mat. 1987, 2, 773-779
Multi-element Adv. Ceramic
20. More K.L., Carter C.H. Jr, Bentley J., Wadlin W.H., LaVanier L. and Davis R.F. The occurrence and distribution of Bcontaining phases in sintered alpha silicon carbide. J. Am C’er. sot. 1987 Lane J.E., Carter C.H. Jr, Bentley J., Wadlin 21. Davis R.F.. W.H., Griffis D.P., Linton R.W. and More K.L. Microanalytical and microstructural analysis of boron and aluminum regions in sintered alpha silicon carbide. Scanning Electron Microscopy 1984, III, 1161-1167
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Wear 22.
transition
diagram
for silicon
carbide:
X. Dong
wu c..
Rice R-W., Platt B.A. and Carrie S. Wear and microstructure of Sic ceramics. Ceramic Engineering ans Science Proceeding 1985, July-August
23. Tsuge A., Uwamino Y. and Ishizuka T. Determination of silicon oxide in silicon carbide by diffuse reflectance infrared Fourier transform spectroscopy. Appl. Spectroscopy 1986, 40, 311-313 24. Singhal S.C. Effect of water vapor on the oxidation of hotpressed silicon nitride and silicon carbide. J. Am. Cer. Sot. 1976, 59. 81-82 25. Singhal J. Mat.
S.C. Oxidation kinetics Sci. 1976, 11, 1246-1253
of hot-pressed
silicon
carbide.
26. Tomizawa H. and Fischer T.E. Friction and wear of sihcon nitride and silicon carbide in water: hydrodynamic lubrication at low sliding speed obtained by tribochemicai wear. ASLE Trans. 1986, 30, 41-46 27. Hegemann B.E., Jahanmir S. and Hsu SM. Microspectroscopy application in tribology. In: Microbeam Analysis. Newbury D.E. (Ed.). San Francisco Press. San Francisco, CA, 1988, pp. 191-19; 28. Cranmer analysis Ceramic
572
D.C. and Galuska A.A. Quantitative wear surface of SIC using a multiple-technique approach. Advanced Materials 1986. 1. 247-251
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et a/. 29. Danyluk silicon Friction Dekker, 30. Peterson contacts General
S., McNallan M., and Park D.S. Friction and wear of nitride exposed to moisture at high temperatures. In: and Wear of Ceramics, Jahanmir S. (Ed.), Marcel New York. 1994, pp. 6’1-78 M.B. Investigation at temperatures Electric, 1958
31. Rahinowicz temperatures. 1961
E.
of solid film lubricants and sliding above 1000°F. Report No. SSGL51,
and Imai M. Friction WADC Technical Report
and wear at elevated 59-603, Part II, May
32. Erdemir A., Fenske G.R. and Erck R.A. and self-lubrication mechanisms of boric Coatings Tech. 1990, 43144, 588-596
Study of the formation acid films. Surface and
33. Arbabi H., Watts J.F. and Briggs J. A chemical study of wear particles and other features associated with the wear of ceramics. Tribology Internatiorzal 1990, 23, 235-243 34. Le Mope Th., Martin J.M., Platon F. and Berthier Y. Rol! formation with triboceramics: which parameters and what mechanisms? In: Wear Particles: From the Cradle to the Grave, Dowson D., Taylor C.M.. Childs T.H.C. Godet M. and Daimaz G. (Eds), Eisevier. Amsterdam, Tribology Series 21, 1992. pp. 285-292
8 December
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Vol. 28, No. 8, pp. 559-512. 1995 Copyright 0 1995 Elsevier Science Ltd Printed in Great Britain. All rights reserved 0301-679X195/$10.00 +O.OO
International
N
0301-679X(95)00082-8
ar transition silicon carbide X. Dong*,
S. Jahanmir+
diagram
for
and L.K. Ives’
To obtain information on the tribological behaviour of silicon carbide at elevated temperatures, unlubricated ball-on-flat wear tests were conducted on sintered silicon carbide in self-mated sliding in air. The contact load was varied from 3.2 to 98.0 N, and a temperature range of 23°C to 1000°C was used. Scanning electron microscopy, Fourier transform infrared spectroscopy and energy-dispersive spectroscopy were used to elucidate the wear mechanisms. The results of the tests and observations were employed to construct a wear transition diagram, which provides a sumn-ary of tribological information including friction coefficient, wear coefficient and wear mechanisms as a function of temperature and load. The wear transition diagram for the sintered silicon carbide studied is divided into four regions plus one transiTion zone. At room temperature, under high loads and high environmental humidity, the tribological behaviour is controlled by tribochemical reactions between the silicon carbide surface and water vapour in the environment. Under low loads and at temperatures below 25O”C, wear occurs by ploughing and polishing. At temperatures about 250°C and under low loads, tribooxidation and formation of cylindrical wear particles control the triboltsgical behaviour. Wear occurs by microfracture when the load is increased above a critical value; and both the friction coefficient and the wear coefficient increase. Keywords:
sintered
silicon
carbide,
tribo-oxidation,
Introduction Silicon carbide materials, due to their attractive mechanical properties and chemical stability, are being i-lcreasingly used in high-temperature tribological systems. However, effective implementation of these materials in new applications requires detailed information on their tribological behaviour. In particular, information on friction coefficient and wear rate under * Graduate Research Assistant, Deuartment of Mechanical Engineering, University of Maryland at dollege Park, College Park, MD 20742, USA T Ceramics Division, National Institute of Standards and Technology, Gaithersburg, MD 20899, USA
Tribology
microfracture
a broad range of conditions is needed to ensure efficient and reliable performance. Furthermore, an understanding of the wear mechanisms and the relationship between these mechanisms and the microstructure of the material is needed to develop improved materials with enhanced performance. Several studies of the tribological behaviour of silicon carbide materials at high temperatures in self-mated slidinglP4 and in sliding against other ceramics or metalP have been published. The influence of temperaturel, normal loadlo, sliding speedll, environment12, and chemical composition of the contacting materials13 were investigated. Several wear mechanisms, for example, microfracture14, plastic deformation,
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ploughing, and the formation and removal of oxidation products l5 , have been identified. In addition to this information, sudden changes in the tribological behaviour, such as an increase in the friction coefficient and the wear rate, have been reported2z9. The mechanisms for such transitions, however, have not been identified. There is considerable variation in the published data, especially regarding the influence of temperature, on the tribological behaviour of silicon carbide materials, For example, the friction coefficient during self-mated sliding at temperatures below 400°C has been reported to range from 0.16 to 0.45’z9, from 0.3 to 0.62, or from 0.7 to 0.84. The wide variation in the published data is apparently due to differences in the composition and type of silicon carbide studied and in test parameters such as contact geometry, sliding speed, and applied load. These disparate results suggest that a systematic approach is needed to elucidate the effects of test parameters and microstructure on the tribological performance and wear mechanisms. In two recent studies on alumina16 and silicon nitridei’, contact load and temperature were systematically varied to obtain data and mechanistic information on the effects of these parameters on the tribological behaviour. The results were compiled into ‘wear transition diagrams’ to show the effects of load and temperature on the wear mechanisms and the values of friction coefficient and wear rate. A similar approach has been adopted for the present study on silicon carbide.
Experimental
procedure
Wear tests were conducted with a linear-reciprocating ball-on-flat high-temperature tribometer, described in a previous paperr6. The material used for both the balls and the flats was sintered alpha silicon carbide. Selected properties of this material are summarized in Table 1. The chemical composition and microstructure of this material have been reported in several references18~22. Boron and carbon are used as sintering aids, resulting in a final composition of 0.5 wt% B and 0.42 wt% free carbon in the sintered material19. The microstructure consists of silicon carbide grains, Table 1 Selected properties carbide (Hexoloy SAP
of sintered
Bulk density (g/cm) Hardness ( GPajb Young’s modulus (GPa) Shear modulus (GPa) Poisson’s ratio Flexure strength (MPa) ‘Fracture toughness (MPa ml’*) Coefficient of thermal expansion Thermal conductivity (W/m/K)
(lO-6/oC)
silicon
3.15 36 406 26 0.14 400 3.7 4.02 120
“The information on the manufacturer or supplier is included here to identify the material. This does not imply endorsement of the product by the National Institute of Standards and Technology “The hardness was measured as a part of the present study. The other values were obtained from reference 18
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with an average grain size of 4 km, and about 8 ~01% of what appear to be ‘dark’ features in the SEMi9. These regions have been identified as amorphous carbon, B4C, and void?. The average diameter of the boron carbide regions were determined to be around 10 Frn; and the silicon carbide grains were found to be saturated with boron in solid solution”‘. Impurities in this material have been reported as 0.2% Fe and 0.06% A115; and the density has been reported to be 98% of the theoretical valueis. The flat specimens used for this investigation were 12.7 mm squares with a thickness of 3 mm. After final polishing with 1 pm diamond paste, the flat test surface had a roughness of 0.1 pm rms. The ball specimens were 12.7 mm in diameter, and were used in the asreceived condition with a roughness of 0.2 p,m rms. Prior to testing, the polished flats and the balls were cleaned to remove surface contamination. The cleaning procedure consisted of the following steps: rinsing twice with hexane, then ultrasonic cleaning in fresh hexane, followed by ultrasonic cleaning with acetone, and final rinsing with hexane. The experiments were conducted at a relatively slow sliding speed of 1.4 x low3 m/s to minimize frictional heating. A stroke length of 10 mm was used in all tests. The normal load was varied from 3.2 to 98.0 N; and a temperature range of 23°C (room temperature) to 1000°C was used. Each test was conducted for a total sliding distance of 10 m (500 reciprocating cycles of 1000 passes) in laboratory air. The relative humidity, measured at room temperature, varied from 50% to 70%. Since the test samples were located in a partially open system, both the relative humidity and the absolute humidity decrease rapidly as the test temperature is increased. Prior to testing, several thermocouples were used to monitor the temperature of the ball, the flat surface, the heated chamber and the specimen holder. These data and the output of a thermocouple positioned under the flat specimen were used to calculate the surface temperature of the flat. Before initiation of sliding, the specimens were allowed to stabilize at the test temperature. The friction force was monitored during the tests with load cells, which measured the tangential force exerted on the stationary ball. The friction coefficients reported in this paper are the average steady state values for the final 2.5 m of sliding. The wear volume on each flat specimen was calculated from surface profile traces (usually three to five) across the wear track and perpendicular to the sliding direction. These data were then used to calculate the wear coefficient, which is a dimensionless quantity defined as wear volume per unit sliding distance per unit load normalized by the indentation hardness. Wear coefficient is used in this paper for convenience; its use does not imply a relationship between wear and hardness. The indentation hardness values were measured as a function of temperature following the procedure described previously 16. These results, shown in Fig 1, indicate that hardness decreases as the temperature is raised. The hardness values at each test temperature were used for the calculation of the wear coefficients. After wear testing, the wear tracks were examined by
Volume 28 Number 8 December 1995
Wear
40
transition
diagram
0 e
I
7
1.0
for silicon
3.2 4.9
.
N N
carbide:
19.6 39.2
N N
* -
X. Dong
70.4 98.0
N N
,
1
et al.
.
I
.
1
800
1000
0.0I 04
I C
200
/
/
400
600
TEMPEHATURE
0
I 000
200
400
600
TEMPERATURE
1000
("C)
('C) o 3.2 o 49N
Fig 1 .Effect of temperature on Vickers hardness of sinterea silicon carbide determined with a 9.8 N load in vacuum of 10p3 Pa
N
.
19.6 39.2
N N
. .
78.4 98.0
N N
1o-2
scanning electron microscopy (SEM) to elucidate the wear process. Energy-dispersive X-ray spectroscopy (EDXS ) and micro-Fourier transform infrared (FTIR) spectroscopy (in reflectance mode) were used to analyse the wear tracks and the wear debris.
Results
(b)
t
t’
and discussion
1o-5-f
The friction coefficients and the wear coefficients are plotted as a function of temperature and load in Fig 2. The data in this figure indicate that the effect of load or. the results can be separated into two groups, where I he friction coefficients and the wear coefficients are load-independent within each group. In Fig 2, a solid curve is drawn for the results obtained at loads lower tian 10 N, and a dashed curve is used for the values obtained at higher loads. At low loads, the average friction coefficient increases slightly as the temperature is raised. This is followed by a sharp drop in the friction coefficient from 0.75 to 0.3a at about 200-300°C. This drop in friction is an indication of a sudden change in the controlling mechanism. As the temperature is increased further, the friction coefficient slowly increases. The wear coefficient, however, is not as sensitive as the friction coefficient to the changes in temperature. The f&ion coefficient at the higher loads ranges from 0.20 to 0.26 at room temperature and quickly increases to about 0.80 as the temperature is raised. The friction coefficient attains a maximum value at approximately 7OO’C, and then drops as the temperature is increased further. At high loads, the friction coefficient and the wear coefficient follow a similar trend as a function of temperature, with the exception of the data at temperatures higher than 7OO”C, where the friction coefficient falls and the wear coefficient rises with increasing temperature. These results show that the tribological behaviour of sinterel? silicon carbide is complex and should be explored in greater detail. The results of the wear Tribology
/ 0
I
I 200
/
1
)
400
TEMPERATURE
/
/
I
600
800
1000
("C)
Fig 2 The effect of temperature and load on (a) friction coefficient and (b) wear coefficient. The solid curves are drawn through the data points at loads less than 10 N; the dashed curves are drawn through the data points obtained at loads larger than 10 N
tests and analysis of the wear mechanisms by SEM, EDXS, and FTIR were used to construct the wear transition diagram shown in Fig 3. The diagram shows the effect of load and temperature on the tribological behaviour of the sintered silicon carbide used in this investigation. Based on the results, which will be presented shortly, the wear transition diagram in Fig 3 is divided into four regions (identified as I to IV) and a transition zone (the cross-hatched area). The controlling mechanism in each region is distinct: tribochemical reaction between the silicon carbide and water vapour in the environment in region I, ploughing and polishing in region II, tribo-oxidation and formation of cylindrical-shaped wear particles in region III, and microfracture in region IV. The last is separated into two parts by the transition zone. The average values of the friction coefficient, f, and the wear coefficient, K, for the flats in each region are indicated in the figure. The standard deviation for the friction coefficient values listed in each region was approximately equal to 10% of the mean value. The standard deviation for the wear coefficient, however, was larger; it ranged from 40% to 60% of the mean value in each
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1995
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Wear
transition
diagram
f=0.23,
for silicon
carbide:
X. Dong
et a/.
. .o
K=lO-'
'0°4[
I1
z
,I.::; 0
200
400
504 P 2 B 0.2
t 0
800
600
TEMPERATURE
i
20
40 60 LOAD (N)
80
100
1000
("C)
Fig 3 Wear transition diagram for sintered silicon carbide, showing four distinct regions plus one transition zone. The friction coefficient, f, and wear coefjicient, K, for each region are indicated in the figure
Y9 6 N, 23Oc t
region. In the following sections, the test results and the observations on the wear surfaces are analysed to eIucidate the underlying mechanisms in each region. Region
I: tribochemical
reactions
0.0
At room temperature and loads larger than 10 N, the friction coefficient ranges from 0.20 to 0.26 with an average value of 0.23. The wear coefficient ranges from 2 x 10Vs to 2 x 10P3 with an average value of 4 x 10P4. This region is observed only at relative humidities greater than 60%. The friction coefficient in this region is independent of load, as shown in Fig 4(a). A typical plot of the friction coefficient versus sliding distance in region I is shown in Fig 5(a). The friction coefficient starts at about 0.75 and rapidly decreases to a steady-state value of 0.22. The wear tracks on the flat specimens tested in this region, shown in Fig 6(a), are smooth with some indications of damage. Observation of the damaged areas at higher magnifications, Figs 6(b) and 6(c), indicates that the surface is covered with a thick film and that the damage is associated with fracture of this film. However, the debris found near the edge of the wear track and those at the end of the track were very small and quite different in appearance. The wear tracks generated in region I were examined with FTIR to determine the composition of the surface films. For reference, a spectrum was first obtained from an area outside the wear track. This spectrum, designated as (a) in Fig 7, shows the SIC peak (due to an IR reflectivity phenomenon). The spectrum (b) for an area inside the wear track of a sample tested in region I is also shown in Fig 7. The major peaks in this spectrum are identified as Si-OH and Si-0 according to reference 23. These peaks suggest the existence of hydrated silicates. The absence of the SIC reflectance peak in this spectrum indicates that the surface film on the wear track is sufficiently uniform 562
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~
40
50 RELATIVE
60 IiUMIDlTY
7:: (X)
80
Fig 4 (a) Effect of load on friction coefficient at room temperature and at 70% relative humidity; (b) effect of relative humidity on friction coefjkient for samples tested at room temperature under a load of 19.6 N and thick to mask the underlying Sic. Since hydrated silicates are usually unstable, the ‘mud cracks’ that were seen on the wear track were probably formed as a result of water loss from the film either after the test or in the SEM. The details for the formation and the structure of the tribochemical reaction products are not known. However, oxidation studies on silicon carbide at elevated temperatures 24.25 have revealed that the oxidation rate increases by addition of water vapour to the oxygen environment. This increase was attributed to the diffusion and solution of OH- ions through the surface oxide film. In our study, the FTIR spectra confirmed the formation of oxidation products by tribochemical reactions. Others have also found evidence for such reactions on the wear tracks in silicon carbide and silicon nitride”6P29. When the relative humidity or the load was decreased at room temperature, the friction coefficient increased from 0.23 to 0.60 (Fig 4). At low relative humidities, the number of OH- ions needed to increase the oxidation rate was probably too small. Therefore, the production rate of hydrated silicate was insufficient to provide a uniform coverage and a low friction coefficient. Examination by scanning electron microscopy of flat specimens tested at lower humidities indicated 8 December
1995
Wear
transition
diagram
for silicon
carbide:
X. Dong
et a/.
8
10
1 .o 19.6
0.0
I 0
I 4
2
(a)
4
SLIDING
N.
23°C
I
6
8
DISTANCE
10
(m)
0.0i 2
4 SLIDING
6
8
DISTANCE
10
0
(m)
2
-4 SLIDING
CC)
6 DISTANCE
(m)
1.o 98.0
N.
19.6
3OO’C
0.c
Cd)
00
-c
0
2
4 SLIDING
6 DISTANCE
8
10
1OOO’C I
--0
2
4 SLIDING
(m)
Fig 5 iypical plots of friction coeficient (d) region IVa, and (e) region IVb
N,
-I n
t
6 DISTANCE
8
10
(m)
(e)
versus sliding
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(c) region HI,
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et al.
Fig 6 Scanning electron micrographs of the wear track on the flat specimen tested in region I at 23°C. (a) Sample tested under a load of 19.6 N, showing a smooth wear track with fractured regions; (b) sample tested under a Eoad of 39.2 N, showing microcrack networks on the wear track at a high magnification; (c) higher magnification of an area in (a) that the wear tracks were mostly smooth, but contained areas where material had been removed by microfracture. It will be discussed later that this dual nature of the wear track appearance is characteristic of the transition zone between regions II and IV. Therefore, at low relative humidities, region I becomes part of the transition zone. Accordingly, the boundary between region I and the transition zone in Fig 3 is drawn by a dashed line to indicate this phenomenon. The load effect can be explained by considering the effect of load on frictional heating. Although in our tests a low speed was used to minimize frictional heating, the contact temperature depends on the applied load during sliding. It is postulated that at low loads the contact temperature is too low to allow for the formation of sufficient oxide film for friction reduction. Therefore, the conditions needed for formation of an oxide film for friction reduction consist of high relative humidities and a high load. Since the relative humidity in our tests decreases rapidly as the temperature is raised, any possible benefit from formation of hydrated silicates is lost as the temperature is increased by frictional heating. 564
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Region
II: ploughing
and
polishing
This region extends from room temperature to approximately 250°C at loads lower than 10 N. The friction coefficient ranges from 0.50 to 0.78 with an average value of 0.67, and the wear coefficient varies between 6 X 10W5 and 2 X 10e3 with an average value of 2 x 1OW. Both the friction coefficient and the wear coefficient were found to be independent of load in this region. A typical plot of friction coefficient vs. sliding distance for region II is shown in Fig 5(b). At the initiation of sliding, the friction coefficient starts at a value between 0.80 and 0.90 and decreases slightly during the test. The friction coefficient remains much larger than that observed in region I. The appearance of the wear track on a flat tested in this region is shown in Fig S(a). The wear surface is smooth, but contains several ploughing or abrasion marks. The microstructure, i.e. the silicon carbide grains and ‘dark’ regions, which are associated with boron carbide, carbon and porosityi are also observed. The wear debris collected at one end of the wear track is shown in Fig S(b),
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1 Si C:
I
(e)
,L
1-
(d) (cl
I i c 00 0
I I
I
I
3200
1600
2400
8130
WAVE NUMBER
Fig 7 P’i’TR spectra obtained on: (a) unworn surface, (b) werv track at room temperature and 39.2 N load, region f; (c) wear track at 100°C and 4.9 N load, region i”r; (d) wear debris accumulated at one end of wear trrrck at 500°C and 4.9 N load, region III; and (e) wecr track at 1000°C and 4.9 N load, region III indicating agglomeration of very fine wear particles. The FTIR spectrum (c) in Fig 7 shows no evidence for for:na.tion of silicates or other reaction products. Therefore, based on the ploughing marks observed on the we,ar track, it is hypothesized that wear in this region occurs by ploughing and polishing of the silicon carbide surface. There is, however, insufficient evidence to completely rule out a possible role for oxidation, since a very thin oxide film cannot be detected by FTIR. Regior 111:tribo-oxidation
Fig 8 Scanning electron micrographs of the wear track on the flat specimen tested in region II at 200°C under a load of 3.2 N showing (a) the ploughing marks, and (b) wear debris accumulated at one end of wear track
reactions
This region is observed at temperatures ranging from 250°C :o 1000°C (and possibly higher) and at loads lower than 10 N. The friction coefficient is much lower in this region than in region II. As shown in Fig 2, the friction coefficient in this region slowly increases from 0.30 to 0.50 as the temperature is increased. The wear coefficient, however, remains relatively constant; it ranges from 2 x IO-’ to 5 x lop4 with an average value of 3 X lo-“. A typical plot of friction coefficient versus sliding distance for region III is shown in Fig 5(c). In this particular test, the friction coefficient started at approximately 0.85 and decreased with sliding distance, ending at a steady-state value of 0.32. This reduction in the friction coefficient is an indication for formation of chemical reaction products. To check this possibility and to define the wear mechanism, SEM was used to examin,? the wear surfaces. A typical wear track is shown in Fig 9(a). The surface appears polished, similar to the micrograph in Fig Z(a) obtained for the wear track in region II. However, the number of ploughing marks are fewer and numerous cylindricalshaped wear particles are observed. These wear Tribology
particles have an average diameter around I km and a length of about 10 km. An SEM micrograph showing the debris accumulated at one end of the wear track (Fig 9(b)) suggeststhat the debris were pushed towards the end of the track and that some of the debris were fractured to produce finer and more equiaxed particles. The wear scar on one of the balls tested in this region is shown in Fig 10(a). The wear scar contains the same features as those on the flat; namely, a smooth surface with cylindrical debris oriented perpendicular to the sliding direction. Figure 10(b) shows the wear debris accumulated at the edge of the wear scar. The morphology of these debris is also similar to that observed on the flat in Fig 9(b). Examination of balls tested in other regions of the wear transition diagram also revealed that wear on both the balls and the flats occurred with the same mechanism. To investigate the origin of the cylindrical wear debris, EDXS and FTIR were used. The EDXS spectrum in Fig 11(a) was obtained with the electron beam located on a single cylindrical particle, and in Fig 11(b) with the electron beam located on an adjacent area on the
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Fig 9 Scanning electron micrographs of the wear track on the fiat specimen tested in region III at 500°C under a load of 4.9 N, showing (a) the smooth surface with cylindrical wear particles and (b) wear debris accumulated at one end of the wear scar wear track. The relatively high concentration of oxygen in the particle suggests that tribo-oxidation was responsible for the formation of the cylindrical particles. However, FTIR analysis, for example spectrum (d) in Fig 7, shows no evidence of silicates or other reaction products in the wear debris accumulated at the end of the wear track. Only the SIC reflection peak can be identified in the spectrum. This could be due to the overwhelming abundance of silicon carbide wear particles at the end of the wear track. However, the spectrum (e) in Fig 7 obtained from the wear track tested at 1000°C shows a small Si-0 peak, confirming increased oxidation rate at higher temperatures. Since the sintered silicon carbide used in our study contained boron and born carbide, we expect oxidation of both boron and boron carbide to have occurred and to contribute to the composition of the oxidation products. However, boron could not be detected by FTIR and the sensitivity of the EDXS system to boron is extremely low. Nevertheless, boron oxide can provide lubricatior?‘*“‘, especially in humid environments, by formation of boric acid32. However, the lubrication 566
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Fig 10 Scanning electron micrographs of the wear scar on the ball tested in region III at 600°C under a load of 9.8 N, showing (a) the smooth surface with cylindrical wear particles and (b) wear debris accumulated at the edge of wear scar
property of boron oxide is strongly temperature dependent. The friction coefficient of boron oxide film rapidly increases to a value higher than 0.9 when the temperature is increased from 200°C to 400°C and then quickly reduces to 0.1 at 600”C30.31. The temperature dependence of the friction coefficient in our tests is different from the expected behaviour of boron oxide. This indicates that formation of boron oxide alone does not explain the results at high temperatures in region III. The oxidation behaviour is complex because of the sintering aids and impurities in the material. It has been found that although oxide films grown at elevated temperatures consist primarily of cristobalite, they contain a glassy phase containing the sintering aides and impurities such Al, Fe, Na, K, etc.25. These elements were found to increase the oxidation rate of silicon carbide2’. It is therefore proposed that formation of a mixed oxide film, containing boron oxide and silicon oxide with other sintering aids and impurities, such as carbon, controls the tribological behaviour in region III. 8 December
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dependence on load and temperature. The average friction coefficient in this region is 0.78 and the average wear coefficient is 2 x 10-3. At higher temperatures, from 600°C to 1000°C in region IVb, the friction coefficient decreasesfrom 0.90 to 0.70 as the temperature is raised; but the wear coefficient increases from 8 x lop4 to 3 x 10-3.
400
0.0
0.5
1.0 ENERGY
1.5
2.0
Figure 5(d) shows a typical plot of friction coefficient versus sliding distance for region IVa. The friction coefficient quickly reaches a steady-state value of 0.78 without much variation in the average value. In contrast with this plot, the friction coefficient for region IVb, shown in Fig 5(e), starts at a low value and quickly rises to 0.90, then drops to a lower value and stabilizes at around 0.70. Although the average steady-state value of friction coefficient is lower in region IVb, the friction trace is much more erratic and variable than the friction trace observed for region IVa.
2.5
(keV)
(b)
0 ;> 0
0.5
1.0 ENERGY
1.5
2.0
2.5
(keV)
Fig 11 EDXS spectra qf wear track for a ball tested in region 111 at 400°C under 58.8 N: (a) on a cylindrical wear particle and (b) on an adjacent area on the wear track Formation of cylindrical wear particles or ‘rolls’ has been observed by others for different ceramics tested under various test conditions9.16,17J6.Z9J3x34, where the compcsition of the rolls were found to depend on the material: test condition and environment. The process for the formation of these rolls is not entirely understood. Microscopic examination of rolls has shown that tl-ey consist of an aggregation of small crystallites bonded together with an amorphous oxide phase29.j3.34. Irrespective of the process for the formation of rolls, the existence of such cylindrical particles is accompanied by a relatively low friction coefficient. Althorngh the process for the reduction of friction by the roj.ls is not clear, it has been suggested that the cylindrical particles act as miniature rolling element bearings, providing a low friction9.‘9%34. Whether this explacation can be used for rolls generated with different ceramics tested under all sliding conditions is not known. Region
IV: microfracture
This r:gion is observed at loads larger than 10 N and is separated by the transition zone into two parts, designated regions IVa and IVb. At lower temperatures, from 50°C to 400°C in region IVa, both the frictio-n coefficient and the wear coefficient show little Tribology
A typical SEM micrograph for a flat specimen tested in region IVa is shown in Fig 12(a), which indicates that the wear track is somewhat rough and it is aPmost completely covered by wear debris. The structure of the compacted debris is shown at a higher magnification in Fig 12(b). This figure shows that what appears as plastic flow at the lower magnification is actually agglomeration of the fine wear debris. The debris layer was found to be strongly attached to the wear track since it could not be removed by ultrasonic cleaning. The micrograph in Fig 12(c) shows an area of the wear track where the debris layer has been removed by wear prior to ultrasonic cleaning. This micrograph suggests a microfracture dominated wear mechanism for the generation of wear particles in tkis region of the wear transition diagram. The wear tracks for the flat samples tested in region IVb were also covered with a debris layer, shown in Fig 13(a). In contrast to Fig 12(a), this micrograph shows that the debris forms long ridges along the sliding direction. The higher magnification micrograph irn Fig 13(b) shows that the debris ridge is composed of compacted wear particles. Most of the fine debris was removed by ultrasonic cleaning (Figure 13(c)); but the debris ridges were not removed, suggesting that they were strongly attached to the wear track. The structure of a region in the middle of the wear track shown in Fig 13(d) suggests that wear has occurred by microfracture. Therefore, the primary difference between regions IVa and IVb is in the behaviour of the wear debris. At lower temperatures, the wear debris forms an adherent layer almost completely covering the wear track, whereas at the higher temperatures, most of the debris is loosely held with the exception of the adherent ridges along the sliding direction on the wear track. In both cases, however, the wear particles are generated by microfracture. Milling, oxidation and adhesion of the fractured particles contribute to the formation of the compacted debris layer in region IVa and debris ridges in region IVb. In the latter, formation and removal of such debris ridges could produce the observed erratic friction trace.
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Fig 12 Scanning electron micrographs of the wear track on the flat specimen tested in region Na at 300°C under a load of 98.0 N: (a) the general appearance of the wear track, (b) higher magnification of an area in (a) showing the structure of compacted debris, and (c) higher magnification of an area in {a) showing microfracture
Transition
zone
The transition zone is shown in Fig 3 as the crosshatched area. In addition to this area, region I becomes part of the transition zone at low relative humidities. In the transition zone, wear occurs by a combination of the mechanisms observed in the two neighbouring regions; one mechanism is enhanced and the other is diminished as the test variables (load and temperature) are varied. Because of this combination of wear mechanisms, it was found that the values of friction coefficients and wear coefficients in the transition zone were in between those expected for the two neighbouring regions. In the transition zone between regions II and IV, as the applied load is increased the dominant mechanism changes from plowing to microfracture. The micrographs in Fig 14 illustrate this dual mechanism. Figure 14(a) shows that the wear track contains smooth regions as well as fractured ones. These regions are shown at a higher magnification in Figs 15(b) and 14(c) confirming wear by ploughing and by microfracture, respectively. 568
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In the transition zone between regions III and IV the dominant mechanism changes from tribo-oxidation reactions to microfracture. Figure 15(a) shows the features associated with the two mechanisms. These areas are shown at a higher magnification in Figs 15(b) and 15(c). The smooth area with a few cylindrical particles on the wear track, in Fig 15(b), is typical for region III where tribo-oxidation was observed. Figure 15(c), however, shows a fractured surface structure in a rough area, which is typical of wear by microfracture in region IV. Summary Unlubricated wear tests were conducted in air to investigate the effect of temperature and load on the tribological behaviour of sintered silicon carbide in self-mated sliding tests. The contact load was varied from 3.2 to 98.0 N, at a constant but low sliding speed of 1.4 x 1O-3 m/s to minimize frictional heating. The specimen temperature was varied from 23°C to 1000°C. The wear surfaces and wear debris were subsequently analysed by SEM, FTIR and EDXS to elucidate 8 December
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Fig I3 Scanning electron micrographs of the wear track on the fiat specimen tested in region IVb at iOOo”C under a load of 19.6 N: (a) The general appearance of the wear track showing wear debris compaction in ridges along the sliding direction, (b) higher magnification of an area in (a) showing the structure of the compacted debris, (c) the general appearance of the wear track after ultrasonic cleaning, and (d) higher magnification of an area in (c) showing microfracture
the w:ar mechanisms. The results of the tests and observations were used to construct a wear transition diagram: which provides a summary of tribological inforn ation including friction coefficient, wear coefficient and wear mechanisms as a function of temperature a.ld load. The wear transition diagram for the sintered silicon carbide studied is divided into four regions plus one transition zone. At room temperature and al loads above 10 N, formation and removal of hydrated silicate films were found to be the dominating wear processes, resulting in an average friction coefficient of 0.23 and a wear coefficient of lo-‘. However, this region was only observed at high relative humidities. At loads below 10 N and temperatures up to 250°C. the dominant wear mechanism was found to be plonghing and polishing. In this region, the average friction coefficient was 0.67 and the wear coefficient was approximately lo-“. In the temperature range from 250°C to 1000°C at loads below 10 N, the tribolagical behaviour was controlled by tribo-oxidation and formation of cylindrical wear particles. The friction coefficient in this region was 0.40 and the wear Tribology
coefficient was approximately 10PJ. When the load was increased above 10 N, microfracture dominated the wear process, resulting in a high friction coefficient of 0.78 and a higher wear coefficient of approximately 10-3. The wear transition diagram proposed for silicon carbide in this paper and for alumina and silicon nitride in the two previous papers’6.‘7 can be used to obtain tribological information for each material. These diagrams contain information on the friction coefficient, wear coefficient and wear mechanisms as a function of load and temperature. Since in most engineering applications a low-friction coefficient and a low wear rate are required, the wear transition diagrams can be useful in determining the range of contact conditions where an acceptable level of performance can be expected. For example, for the sintered silicon carbide used in this study in self-mated tests, high loads and high temperatures should be avoided since the friction coefficient and the wear rate are excessive under these conditions.
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Fig 14 Scanning electron micrographs of the wear track on the flat specimen tested in the transition zone at 100°C under a load of 19.6 N: (a) Overall view of the wear track showing smooth areas and fractured regions, (b) higher magnification of the smooth area in (a) showing ploughing and (c) higher magnification of the fractured area in (a) showing microfracture
It should be pointed out that the wear transition diagram presented in this paper has been obtained for one specific material and is based on tests conducted at a low sliding speed. The effect of such factors as chemical composition, microstructure, sliding speed and contact geometry was not investigated. Any significant change in these factors could alter the shape of the diagram. For example, the boundary between regions II and III could be sensitive to the composition of the material, and compositional changes may shift this transition boundary. The effect of speed, however, is complex. First, an increased speed would increase the rate of tribochemical reactions, but at the same time, it could promote microfracture through increased thermoelastic stresses. The effect of these factors on the tribological behaviour of silicon carbide requires further study.
The authors would like to thank Dr Ramesh Divakar of Carborundum Company for providing the test samples and Dr Hong Liang of the University of Maryland for obtaining the FTIR spectra.
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Fig 1.5 Scanning electron micrographs of the wear track on a fiat tested in the transition zone at 500” under a load of 19.6 N: (a) Overall view of the wear track showing smooth areas and fractured regions, (b) higher magni$cation of the smooth area in (a) showing cylindrical wear particles, and (c) higher magnification of the fractured area in (a)
showing microfracture
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