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Mechanics of Deformable Solids / Механика деформируемого твердого тела
УДК 620.18
Moshkovich A., Perfilyev V., Lapsker I., Meshi L., Rapoport L.
ALEXEY MOSHKOVICH, Ph.D., Holon Institute of Technology, Holon. 5810201, Israel, e-mail: alexeym@hit.ac.il VLADYSLAV PERFILYEV, Ph.D., Holon. Institute of Technology, Holon, 5810201, Israel, e-mail: vladper@hit.ac.il IGOR LAPSKER, Ph.D., Holon Institute of Technology, Holon 5810201, Israel, e-mail: lapsker@hit.ac.il LOUISA MESHI, Ph.D., Department of Materials Engineering, Ben Gurion University of the Negev, Beer Sheva 84105, POB 653, Israel, e-mail: louisa@bgu.ac.il
LEV RAPOPORT, Prof., Ph.D., Holon Institute of Technology, Holon, 5810201, Israel (corresponding author), e-mail: rapoport@hit.ac.il
Superplastic deformation of a/p brass under friction in lubrication conditions
Copper (Cu) and its alloys are widely used for various moving machine parts. Friction and wear of machine parts is mainly determined by the microstructure and plastic deformation of surface layers. The dislocation structure and plastic deformation of a/p brass after friction in boundary lubrication (BL) conditions were studied. The friction of brass was attributed to superplastic deformation (according to our estimation close to 500%) in thin surface layers of the a-phase grains. Intragranular slip, as the accommodating mechanism, occurred in a-phase, whereas a little deformation was observed in the P-phase. The accommodation of sliding was accompanied by growth and coalescence of voids, formation and propagation of cracks, leading finally to delamination of the wear particles. The equivalent strain vs. the depth was evaluated. Large deformation of thin surface layers (s = 10-12) led to formation of the nanocrystalline structure (d = 35 nm). The SPD of surface layers is accompanied by formation of thin shear bands practically parallel to the direction of friction. The thickness of wear particles was found to be close to the thickness of shear bands. The hardness of surface layers of Cu slightly decreased, while the hardness of brass increased with an increasing the contact temperature. Deformation hardening of brass was significantly larger than for Cu.
Key words: brass; microstructure; plastic deformation; friction.
Introduction
Copper (Cu) and its alloys are widely used as engineering materials for various rubbed machine parts. Friction and wear of Cu and Cu alloys are accompanied by severe plastic deformation (SPD) and fracture in surface layers [e.g. 12, 13, 15-18, 32]. The dislocation structure and work hardening of Cu during friction with lubricant were studied in [21, 27]. It has been shown that a laminated structure with thin and elongated grains is formed in the elasto-hydrodynamic lubrication (EHL) region. Significant grain refinement (d = 20-100 nm) was observed in the surface layers after friction in the boundary lubrication (BL) region. The SPD of surface layers was accompanied by formation of shear bands in the
© Moshkovich A., Perfilyev V., Lapsker I., Meshi L., Rapoport L., 2015
region of contact spots. Recently, it was found that the deformation hardening of nanocrystalline surface layers of Cu under friction in the BL region resembles the behavior of the refined structure of Cu obtained by different methods of SPD [25]. Grain refinement following SPD in face-centered cubic (FCC) bulk metals and alloys is usually achieved through the accumulation of dislocations and subsequent rearrangement by dynamic recovery or recrystallization [e.g. 30, 33, 34, 36, 38]. For these materials, the ultra-fine grain size and saturated crystalline defects can lead to superplastic deformation at high temperatures [28]. Using Equal Channel Angular Pressing (ECAP) for grain refinement, a superplastic elongation of 640% was recently achieved for the 60%-40% CuZn alloy in tension and at a temperature of 673 K [28]. It is 200° lower in comparison to previously reported results [3, 8, 29], where superplastic ductility was exhibited at testing temperatures of 800-900 K. The temperature of superplasticity during tension is close to that observed at the points of contact during friction. It is known that the flash temperature in the interface of rubbed surfaces can achieve 750-800 °C [2]. For instance, Wang et al. [35] demonstrated the temperature rise to hundreds of degree Celsius close to the surface. Thus, it is expected that the SPD of a/p brass due to friction can also lead to superplastic deformation of surface layers at relatively high flash contact temperature. In spite of the large number of papers, dedicated to superplasticity in tension, little effort has been directed toward elucidating the mechanism of plastic deformation and microstructure of brass in lubrication friction conditions. The current research was undertaken with three main objectives: first, to briefly demonstrate the friction behavior of a/p brass in the BL region; second, to evaluate the microstructure and plastic deformation of surface layers after friction, to compare the structure of these layers with FCC materials undergone to SPD; third, to study the effect of friction on hardness and deformation hardening of Cu and brass.
Experimental Procedure
For friction pin-on-disk tests, the disks made from steel (AISI 1040) and hardened up to HRc = 40, slide against the flat pin of commercial a/p brass. This alloy was chosen as an example of two-phase engineering material with a possible superplastic deformation at elevated temperatures. The composition of brass is listed in Table 1. Hardness and composition of different phases of brass is shown in Table 2. In order to obtain the same virgin grain size (d = 30-40 pm), the Cu samples were annealed at 400 °C during one hour and the brass samples were annealed at 300 °C during two hours. The grain virgin structure of brass is shown in Fig. 1. The relationship between a- and P-phases in virgin state of brass is 70:30.
Table 1
Chemical composition of brass, in % wt
Cu Pb Zn Al As Cd Fe Mn Ni Sn As+Sb Other
58.5 2.86 37.97 0.005 0.01 0.01 0.256 0.01 0.093 0.235 0.027 0.024
Table 2
The composition and hardness of a-and ß-phases
Phase Zn, wt% Microhardness, MPa
a-phase 38 ± 0.5 1500±50
ß-phase 44 ± 0.5 1900±50
The sliding velocity was constant, v = 0.35 m/s and the load varied in wide range (150-3200 N). Some drops of the synthetic oil, PAO-4 have been supplied to contact region. The bulk temperature was measured by fixing a thermocouple to the block in close proximity (1.5^10-3 m) to the tested surface. The virgin roughness of disks and pins, (Ra), was estimated as 0.1 p,m. The variation in Vickers hardness with depth was measured using cross-sections normal to friction surface. In order to measure the hardness of thin surface layers, a load of 0.05 N was applied (a depth of the indenter penetration was about ~ 1^m). Thirty measurements were performed for decreasing the spread of the results. The equivalent plastic strain, a, at a given depth of subsurface layers was calculated from the shear angle of the interface 0 [33, 24].
^ a a —— tand
3
Fig. 1. General view of the microstructure of a/p brass sample at virgin state. Black spots are the Pb inclusions
(1)
The stress-strain, of = f(a), relationship for deformed layers was analyzed. The flow stress, of, was assumed to be related to the hardness value, H, as of =H/3 [20], where H is the Vickers hardness.
The microstructural evolution of Cu and brass samples before and after friction was characterized by optical (light) microscopy (OM), scanning electron microscopy (SEM) and transmission electron microscopy (TEM) techniques. Cu galvanic coatings were used in order to protect the rubbed surfaces from damage during preparation. The Cu-Zn blocks were cut carefully parallel and normal to sliding direction. SEM images were obtained on a Stereoscan-430i SEM (20kV acceleration voltage and standard Everhart-Thornley detector). The arrows in all SEM images indicate the direction of sliding motion. TEM samples were prepared in cross-sectional orientation, by mechanical polishing followed by Ar-ion thinning in a Gatan PIPS ion milling machine, and examined in a JEOL TEM-2010 electron microscope operating at 200kV.
Results
Virgin microstructure of Cu and Brass
TEM study has revealed that virgin brass's microstructure was complicated, see Fig. 2
(a) (b) (c)
Fig. 2. Investigation of the microstructure of virgin brass revealed (a) large equiaxial grains with high dislocation density (b) dislocation networks and (c) dense twins with varying width. Brass exhibited a relatively high dislocation density (1014-1015 m-2, see Fig. 2a for illustration) and varying grain size: 2^15 ^m equiaxial grains and 20^800 nm wide twins (see image in Fig. 2c and matrix/twin diffraction
pattern in Fig. 3a)
Fig. 3. PED patterns taken from a brass (a) matrix and twin at [011] orientation and (b) one
grain orientated along [001] zone axis
As shown in Fig. 2a and b - the dislocations form the complex nets and frameworks. In order to verify the crystallographic structure of a-brass and matrix-twin orientation relationship - precession electron diffraction (PED) was used. All PED patterns taken from different grains of brass were successfully indexed in terms of FCC cubic brass structure (see indexed [001] pattern, as an example in Fig. 3b) and, using superimposed matrix-twin PED pattern (shown in Fig. 3a), it was found that (111) is the twin plane, which is consistent with known matrix twin relationship for FCC crystals.
The morphology of subsurface layers after friction
The morphology of rubbed surface of brass after friction in the BL region is shown in Fig. 4. The SPD of surface layers is associated with a strong shearing in relatively thick sheets. It can be easily seen in Fig. 5 that this shearing is accompanied by severe plastic deformation of subsurface layers up to the depth of about 20 pm.
Fig. 4. SEM image of the surface of brass Fig. 5. Cross-section of the worn Cu-Zn surface after friction in the BL region. Arrow shows the layers after friction in the BL region. Contact direction of friction spots of contra body (marked by 1) are shown
schematically. White spots in the brass (designated as 2) are the Pb inclusions
The magnified image of surface layers of brass is depicted in Fig.6. The plastic deformation in the contact regions is characterized by formation of macro-scale shear bands due to high contact pressure. A width of the bands is decreased from 10 pm to 1-2 pm close to the surface. The formation of macro-scale bands in brass is accompanied by enlargement of voids and microcracks both inside the bands and in their boundaries.
Fig. 6. Magnified image of cross-section perpendicular to the wear track revealed a coalescence of voids and microcracks in direction of flow lines around the contact spots. Insert demonstrates clearly the direction in the coalescence of voids and growth of cracks
It seems that intensive void nucleation and growth precede the macrocracks formation. A size of the region of SPD, void's formation and their coalescence is about 20 pm. A void development is observed just in the a-phase. The cracks between a and P phases are practically absent. The amount of the P phase in surface layers is less than 10%.
Subsurface deformed layers sheared in the direction of friction are shown in Fig. 7. Strong gradient of plastic deformation up to the depth of about 95 pm was observed (see line marked by 1 in Fig. 7).
Fig. 7. SEM images of subsurface layers in the direction of friction. Direction of
friction
is shown by an arrow. 2 and Q mark flow line and shear angle, respectively
The flow line (marked by 2 in Fig. 7) and shear angle (marked by 0 in Fig. 7), are added in order to demonstrate the method of s calculation (see eq.1 in experimental procedure section). The brass grains were severely deformed by bending in the direction of sliding. A dependence of the equivalent strain, s on a given depth below the surface is shown in Fig. 8.
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Fig. 8. A variation of the equivalent strain vs. the depth bellow the wear surface
Three regions of the strain variation can be revealed: first one is a region up to ~20 |im (s<2); region two up to ~ 5 |im with slow strain increasing (s ~ 6); and third region from 5 |im to the surface where the strain increased slowly up to maximal value (s ~ 12). At this region, the SPD of the a-phase is accompanied by strong shearing and cracks propagation in the direction of friction, Fig. 9. The coalesced voids and microcracks propagate in the direction of shear bending (flow lines) up to the depth about 20 ^m. During shearing, the virgin a-grains (d~30 |im) decreased to the thickness of bands (1-2 |im). The length of these bands elongates up to more than 150 |im. Since the grains elongated from d = 30 |im to d = 150 |im, it can be inferred that the superplastic deformation of approximately 500% occurred. Similar results, obtained in tensile tests of brass, presented in [8]. If to suggest that the virgin grains transformed into thin shear bands (~2 |im), the superplastic deformation is probably even larger. The shear bands in the depth of 1-5 |im are practically parallel to the friction surface, see Fig. 10. It can be easily seen that the cracks propagate mainly inside the thin shear bands in the a-phase. A cracking between the shear bands is responsible for formation the wear particles and their delamination. The wear particles and their cross section are shown in Fig. 11. Comparing Fig. 10 and 11, it can be concluded that a width of the bands corresponds to the thickness of the wear particles. The size of wear particles was varied in the range of 10 |im to 250 |im. It is interesting to note that these wear particles consisted of thin micro-scale sheets (see Fig. 11b). X-ray map in the insert in Fig. 11a demonstrates a presence of the Pb inclusions, uniformly distributed on the surface of wear particles.
0
2
Fig. 9. SEM images of surface layers of brass. Direction of friction is shown by an arrow
Fig. 10. SEM image of shear bands close to the rubbed surface
Based on the analysis of the microstructure of surface layers after friction it can be concluded that the superplastic deformation occurs mainly in the a-phase. The deformation of P-phase is limited. Some cracks and cavities appeared just in the boundaries between the a and P phases, and inside the grains of P-phase.
Fig. 11. SEM image of the wear particles of brass (a), cross section of the wear particle (b). In the insert, is the distribution of Pb on the surface of wear particles
The microstructure of brass after friction in the BL region
In order to understand the evolution of the microstructure of brass after friction, the detailed TEM investigation was performed, Fig. 12. The nanosized grains are clearly seen at the outer surface. Ring electron diffraction pattern shown in the insert points the complete misorientation of these nanograins. The depth of the damaged surface layers is deeper than 10 |im. It can be seen that even 10 |im below the surface the morphology is still different from the virgin state.
It can be seen that they appear near the cracks. In the insert -ring electron diffraction pattern taken from the surface layers of brass is shown. Damaged layers of rubbed brass in the BL region can be visually divided into several layers. First outer layer has a thickness of ~1.5 |im. It contains the nanosized equiaxial grains with d = 35 nm. Then, the grains become elongated in the direction of friction (see Fig. 13 for higher magnification).
These grains have varying width of 100^200 nm (wider at the bottom layers) and length up to 1 micron. Elongated grains are the substructure of shear bands, as can be seen in Fig. 13. The size of the shear bands (one band is marked by an ellipse in Fig. 13, for example) is about 1^2 |im in width and this correlates well with SEM observations listed earlier in this paper, see Fig. 10. It is clear that these bands are those micro-scale bands which constitute the macro-scale shear bands observed in SEM in Figs. 6 and 9. Interestingly, we could identify the twins only in the lower layers, were the morphology resembled the virgin brass state.
Mechanical properties of the surface layers In order to evaluate the deformation hardening of Cu and brass after friction, a variation of the hardness vs. a depth was measured, Fig. 14. The hardness (stress) of thin surface layers for brass was found to be significantly higher than for Cu samples (2750 MPa and 1450 MPa, respectively). The deformation hardening (a difference between the hardness of deformed and virgin samples) for brass was 1400 MPa, while it was about 500 MPa for Cu samples.
Fig. 12. TEM montage showing the microstructure of brass after friction in the BL region. Black particles marked by arrows are the Pb inclusions
It is well know that an increase of the load during friction tests leads to a raise of the temperature. The effect of temperature on the hardness of surface layers of Cu and brass is shown in Fig. 15. Principally different effects are revealed for Cu and brass samples. Slight softening is observed in Cu, whereas the deformation hardening is revealed in brass.
Friction of brass is attributed to superplastic deformation of thin surface layers of the a-phase grains due to compression and shearing. Intragranular slip as the accommodating mechanism occurs in the aphase, whereas little deformation is observed in the P-phase. The accommodation of superplastic deformation in sliding is accompanied by the growth and coalescence of voids, the formation and propagation of cracks, leading finally to delamination of wear particles. Our results indicate that intragranular slip in the a-phase grains is the main mechanism of superplastic deformation under friction of brass in the BL region. As a result of this superplastic deformation, the shear bands appeared practically parallel to the direction of friction, see Figs. 9a and 10. Two view-points were proposed in the literature in order to explain the superplasticity appearing during tension of a/p brass at temperature about 600 °C [3, 8, 29]: (1) the intragranular slip in the P phase; and, (2) the boundary sliding on the interface between the a and P phases. In the current work, a superplasticity of the a-phase's grains during friction contradicts to superplastic ductility of the P-phase observed at high temperature in tension. It is known that P phase at the temperature lower than of 454 °C is hard and brittle [6]. In fact, the deformation of P-phase during friction was limited. Based on these results, it can be concluded that the flash temperature of brass-steel contact was lower than 454 °C. The measured temperature near the contact was not more than 100 °C. Therefore, it is plausible to suggest that superplastic deformation of brass is attributed to the SPD of surface layers and ductility of the a-phase at relatively high flash temperatures of contact, lower than 454 °C.
The observed microstructure of surface layers after friction correlates well with the equivalent strain. Large deformation (s = 10-12) led to formation of the nanocrystalline structure (d = 35 nm) at thin surface layers (1-5 pm). The microcracks observed at these layers. The twins were not formed at large strain where a development of the shear bands was dominant. These results are in-line with the results received in plane strain compression of copper-based alloy [9]. With decreasing the strain (s ~ 6), the elongated grains are formed (d = 100/200 nm-1000nm). The grain size increased as a function of the depth.
Fig.13. Montage of several TEM images showing the subsurface layer where the shear bands are appeared. White arrow shows a direction of friction
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Fig. 15. The effect of temperature on the hardness of Cu and brass rubbed in
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Likewise, our results agree well with the analysis of the microstructure of the low carbon 316 stainless steel subjected to dry friction [31]. The structure and plastic deformation was explained there by relatively low SFE of low carbon 316 stainless steel. A similarity in the damage development during friction of FCC a-phase of brass and stainless steel hints that the SFE plays a noticeable role in the plastic deformation and strain hardening of brass It was shown that the hardness and strength of Cu-Al, Cu-Zn and Cu-Ge alloys subjected to the SPD can be simultaneously enhanced by lowering the SFE [1, 30, 34, 38]. However, the plastic deformation of these materials even in the early stages of subsequent deformation was accompanied by plastic instability due to formation of the shear bands having a direct effect on the failure initiation and crack development [4, 10, 22]. Plastic instability-shear banding during tension/compression leads to necking and, finally, to fracture. Superplastic deformation of brass during friction is also accompanied by formation of shear bands, crack propagation and delamination of the wear sheets. Principal difference between the fracture in the static, or dynamic (fatigue) and wear tests, is a continuation of plastic deformation after delamination of wear particles. A delamination and, thus, formation of "fresh" surface can also stimulate the plastic deformation during friction.
Present research showed significantly higher deformation hardening and saturation hardness (stress) for brass in comparison to pure Cu. The recovery processes in Cu lead to decrease in hardness as a function of the temperature of contact. An increase of the deformation hardening of brass during friction corresponds well with the analysis of the work hardening and saturation stress during plastic deformation of FCC materials with low SFEs [1, 30, 33, 34, 36]. It is known that strain hardening of FCC metals involves the intersections of dislocations and cross slip. An increase of the saturation stress in the low SFE's materials was explained previously by a difficulty in cross slip and, thus, the deceleration of the recovery process even under relatively high temperatures [e.g. 14, 37]. Generally, the deformation hardening also depends on the concentration of alloying elements. However, it was shown that the alloying elements play a negligible role in strengthening (< 15% [19]) or even smaller [11]. Recently, the negligible effect of the alloying element on the strengthening of FCC alloys has been confirmed in Ref. [9]. Thus it is suggested that the SFE decreasing plays a main role in the deformation hardening of a/P brass.
Many voids and cracks were observed both in the boundary and inside of the sheared a-grains after friction in the BL region. It is naturally to anticipate that the shear bands, formed during friction, initiate the voids, their coalescence, crack formation and their development. High contact pressure in the friction spots led to large strain in the surface layers of brass and Cu [27], similar to that observed for FCC materials subjected to the SPD processing [23]. The voids expand, grow and coalesce in the direction of flow lines. It can be proposed that the voids nucleate mainly in the boundaries of the shear bands. Compression and shearing on the contact spots activate grain-boundary sliding. In accordance with the Cocks-Ashby model [5], the voids grow due to the flow of vacancies along boundaries and surfaces or dislocations. The density and the average size of voids around the contact spots increased with a repeating sliding. The voids and their coalescence initiated crack-like early growth due to the local plastic constraints. The cracks in the boundaries of shear bands, shown in Fig. 10, led to the delamination of wear particles.
One of the surprises of this research's outcomes was the role of the Pb inclusions in the brass. Usually, the Pb inclusions play a dual role in friction and wear. From one hand, the Pb is supplied to rubbed surfaces and thus can lead to decrease of the friction. From other hand, an appearance of Pb inclusions near the microcracks can accelerate their propagation. Here, the Pb inclusions were found both on the rubbed surface and on the surface of wear particles. Although Pb was revealed on rubbed surfaces, the friction coefficient remained high, as usually observed in the BL region. It can be suggested that the plastic deformation in much thicker layers limited the effect of Pb as lubricated film, while it effected severely the propagation of cracks.
Finally, we would like to note that the superplastic ductility of two-phase's materials with ultrafine grain structure are especially attractive in metal forming processes. However, the superplastic deformation in friction leads to formation of thick and large wear particles and, therefore, to high wear rate. Materials, such as Cu, with a localized deformation at thin surface and ultra-fine grain layers can provide low friction and wear.
Conclusions
1. The dislocation structure and plastic deformation of a/p brass rubbed in the BL regime were studied.
2. Friction of a/p brass is characterized by superplastic deformation of surface layers. The superplastic deformation of brass is attributed to intragranular sliding inside of the a-phase grains and accompanied by the growth and coalescence of pores, formation and propagation of cracks, and development of wear particles.
3. Strong shearing in surface layers of brass is accompanied by formation of shear bands parallel to the direction of friction. A thickness of wear particles is close to the width of shear bands.
4. The saturation hardness of Cu slightly decreases with a temperature, while the hardness of brass increased. The deformation hardening of brass is significantly larger in comparison to Cu.
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Механика деформируемого твердого тела
А. Мошкович, В. Перфильев, И. Лапскер, Л. Меши, Л. Рапопорт
АЛЕКСЕЙ МОШКОВИЧ - Ph.D. (Технологический институт, г. Холон, Израиль). Holon 5810201, Israel, e-mail: alexeym@hit.ac.il
ВЛАДИСЛАВ ПЕРФИЛЬЕВ - Ph.D (Технологический институт, г. Холон, Израиль). Holon 5810201, Israel, e-mail: vladper@hit.ac.il
ИГОРЬ ЛАПСКЕР - Ph.D. (Технологический институт, г. Холон, Израиль). Holon 5810201, Israel, e-mail: lapsker@hit.ac.il
ЛУИЗА МЕШИ - Ph.D., Департамент материаловедения (Университет им. Бен Гуриона, г. Бер Шева, Израиль). Beer Sheva 84105, POB 653, Israel, e-mail: louisa@bgu.ac.il
ЛЕВ РАПОПОРТ - Professor, Ph.D. (Технологический институт, г. Холон, Израиль). Holon 5810201, Israel, e-mail: rapoport@hit.ac.il
Сверхпластическая деформация a/ß латуни при трении в условиях смазки
Медь (Cu) и ее сплавы широко используются для различных подвижных частей машин. Трение и износ деталей машин в основном определяется микроструктурой и пластической деформацией поверхностных слоев. В работе исследовалась структура дислокаций и пластическая деформация a/ß латуни после трения в режиме граничной смазки (BL). Трение латуни было связано с сверхпластической деформацией (по нашим оценкам, примерно 500%) в тонких поверхностных слоях зерен a-фазы. Внутризеренное скольжение, как механизм подстройки микроструктуры, происходило в a-фазе, в то время как незначительная деформации наблюдалась в ß-фазе. Внутризеренное скольжение сопровождалось ростом и слиянием пор, формированием и распространением трещин, что привело, в конце концов, к расслоению частиц износа. В работе оценена относительная деформация как функция глубины от поверхности. Большая деформация тонких поверхностных слоев (в = 10-12) привела к формированию нанокристаллической структуры (размер зерен d = 35 нм). Жесткая пластическая деформация поверхностных слоев сопровождается образованием тонких полос сдвига практически параллельно направлению трения. Толщина частиц износа близка к толщине полос сдвига. Микротвердость поверхностных слоев Cu несколько снизилась, в то время как микротвердость латуни увеличивается с повышением температуры контакта. Деформационное упрочнение латуни было значительно больше, чем меди.
Ключевые слова: латунь, микроструктура, пластическая деформация, трение.
