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Structure and wear of Al surface layers reinforced with AlCuFe particles using ultrasonic impact peening: Effect of different particle sizes
Surface & Coatings Technology 205 (2011) 5278–5284
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Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
Structure and wear of Al surface layers reinforced with AlCuFe particles using ultrasonic impact peening: Effect of different particle sizes B.N. Mordyuk a,⁎, M.O. Iefimov b, K.E. Grinkevych b, A.V. Sameljuk b, M.I. Danylenko b a b
Kurdyumov Institute for Metal Physics, 36 Academician Vernadsky Boulevard, UA-03680, Kyiv, Ukraine Frantzevich Institute for Problems of Materials Science, 3 Krzhyzhanivsky Street, UA-03142, Kyiv, Ukraine
a r t i c l e
i n f o
Article history: Received 13 December 2010 Accepted in revised form 19 May 2011 Available online 6 June 2011 Keywords: Surface composite Quasicrystal Aluminium Microstructure Ultrasonic impact peening Wear
a b s t r a c t Near-surface layers in aluminium specimens are modified using quasicrystalline (QC) AlCuFe particles introduced into a zone of severe plastic deformation induced by ultrasonic impact peening (UIP). Two types of QC particles are used: atomized with average size of approx. 25 μm (coarse QC–c-QC) and milled — 0.3–0.5 μm (fine QC–f-QC). The effect of QC particles of different sizes on microstructure and wear resistance of subsurface composite layers in aluminium is studied in this paper. XRD, SEM and TEM studies of reinforced aluminium layers allow establishing the links between microstructural features of the layers and their sliding wear. The formed layers of composites reinforced with both types of QC particles demonstrate almost double increment in wear resistance when compared to that of annealed aluminium. It is due to the combination of several factors: (i) high hardness and high wear resistance of QC reinforcement (more efficient for c-QC); (ii) relatively strong interfacial bonding of homogeneously dispersed reinforcing QC particles; (iii) fine grain structure of the Al matrix (f-QC) or increased density of dislocations arranged in fine dislocation-cell structure (c-QC) — i.e. increased volume fraction of grain boundaries/dense dislocation walls. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Metal-matrix composites (MMCs) exhibit wide engineering applications due to their enhanced hardness, wear, creep and fatigue resistance in comparison to the matrix metal or alloy [1–3]. Composite layer in a component surface is known to be sufficient or desirable in many applications — mainly to improve the surface dependent properties of material and the operation life of component or construction [2–4]. A number of recent studies show wide possibilities in fabrication of composite layers by dispersing of hard ceramic or intermetallic particles in a surface of Al (or Al based alloys) using either melting related techniques [5,6] or methods of severe plastic deformation (SPD) [7–9]. Use of quasicrystalline (QC) particles to reinforce AlMMCs seems to be prospective because QC materials have unique combination of high hardness and modulus of elasticity, high wear resistance and relatively low density. Unlike metastable QC reinforcements usually formed as precipitates to strengthening Albased alloys by rapid solidification techniques [10–12], stable QC particles (as AlCuFe) can be dispersed in modified surface layer being introduced into a zone of SPD. As a result, more stable materials can be obtained. They would not deteriorate above some critical temperature (temperature of phase transformation ‘quasicrystalline–crystalline’)
⁎ Corresponding author. Tel./fax: + 380 44 424 0521. E-mail address: [email protected] (B.N. Mordyuk). 0257-8972/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2011.05.046
as it usually occurs in Al based alloys contained metastable QC precipitates. Our previous studies show beneficial effects of QC particles dispersed in a surface layer of Al specimens on its microhardness and damping characteristics, which are increased significantly [13]. However, a size of reinforcing particles is known to affect the matrix structure and mechanical properties of the formed composite [1–5]. The aim of this study is two-fold: to examine the effect of AlCuFe particles of different sizes on microstructure of the Al matrix in surface layers formed at the UIP process, and to establish the links between microstructures formed in the QC particle-reinforced layers and their sliding wear performance.
2. Experimental details 2.1. Preparation of composite layers Plane specimens (60 ×15× 2.5 mm) of cp aluminium (purity of 99.75%) were initially annealed at 400 °C for 1 h in vacuum of 1.3 ×10− 3 Pa. The resulting grain size was about 60 μm. Atomized powder of icosahedral AL63Cu25Fe12 quasicrystal was used for reinforcement. An average particle size of initial atomized powder was approx. 25 μm (Fig. 1a); it is denoted “coarse” below (c-QC). Other powder (denoted “fine” in this paper — f-QC) was obtained by grinding of the initial powder in an ultrasonic mill [14] down to average particle size of 0.3–0.5 μm (Fig. 1b). Suspensions of AlCuFe powders in liquid paraffin
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The surface morphology of composite layers was examined in Superprobe-733 scanning electron microscope (SEM) using characteristic X-ray radiation. Transmission electron microscopy (TEM) analysis was carried out using a JEM 100 CX-II microscope. The plane-view TEM foils prepared from the top sub-surface layer (~ 5–10 μm) were mechanically polished from the un-treated side to thickness about 30 μm followed by ionic polishing (also from the un-treated side). 2.3. Wear examinations
Fig. 1. Appearance (a, b) and XRD spectra (c) of reinforcing powders: coarse AlCuFe powder after atomization and annealing (a, spectrum 1 in c) and fine one after milling (b, spectrum 2 in c).
were used to cover the surface of Al specimen in the UIP process to form composite layers. Equipment and appropriate regimes for the UIP process are described in detail in Refs. [13,15–17]. Basically, the UIP process consists in multiple sliding impacts imposed to a surface of treated specimen; impacts are induced by the pins positioned between specimen and an ultrasonic horn in a special impact head, which rotates forcedly at given velocity. The frequency of impacts in this loading scheme is about 3 kHz. Progressive shifts of the acoustic system along the treated surface for several times allow obtaining rather low surface roughness, relatively uniform distribution of reinforcements in composite layer, and sufficient thickness of this particle-reinforced surface layer (ten UIP passes performed for each specimen in this study produce the resulting layer with thickness of approx. 50–60 μm — similarly to our previous works [13,17,18]).
2.2. Microstructure examination The XRD θ–2θ analysis of reinforcing powders and treated specimens was carried out using a DART-UM1 diffractometer with Cu Kα irradiation. Assuming a biaxial stress state, analysis of residual stress in the Al matrix was carried out using a sin 2ψ based method [6,17,19]. An annealed Si powder deposited on a surface of an Al specimen was used as reference. In-plane residual stress σ|| = σ1 + σ2 was calculated using literature values of elastic modulus (E = 70.29 GPa) and Poison's ratio (ν = 0.3) for aluminium as an intercept of a linear graph εψi = f(sin2ψ) with the axis of ordinates (at sin2ψ= 0) [19]: σ‖ = −(E εψ = 0)/ν. Besides, considering different propagation depths of Xrays depending upon the diffraction angle, through depth profiles of inplane residual stress σ‖ were roughly estimated using the following expression σ1 + σ2 = −E(d− d0)/νd0[20,21], where d0 is a reference lattice spacing in a stress-free specimen and d is the lattice spacing in the UIP-treated specimen measured at different diffraction angles.
Wear resistance was examined using an automated tribocomplex (CATC) [22]. Sliding wear tests of plane specimens were performed with Si3N4 spherical indenter (8 mm in diameter) as counter-part in inactive liquid paraffin. Those tests consisted in a reciprocating sliding movement of indenter along a track of 4 mm, and lasted for 20 min. Sliding velocity was 0.013 m/s. The CATC allowed conducting tests of two types [22]: (i) usual quasi-static wear test performed at quasistatic load (20 N) applied in a normal direction to the tested surface, and (ii) so-called dynamic wear test, which was carried out using both quasi-static and alternating components of load applied in a normal direction to the tested surface. In the case of dynamic test, the obtained results are much closer to the operation conditions of studied material. Amplitude of the alternating component of load was approx. 10% regarding to that of quasi-static one. Wear magnitudes for quasi-static (Wst) and dynamic (Wd) modes were characterized with the depth of wear tracks measured using P-201 profilograph– profilometer. The friction force (F) was estimated via measured displacement of an elastic circular element rigidly connected to tested specimen. Besides, microhardness of the specimens was measured at load of 200 g on the Vickers indenter applied to their surface. 3. Results 3.1. XRD and SEM observations Results of X-ray structural analysis of reinforcing powders presented in Fig. 1c demonstrate that after atomization and subsequent annealing an AlCuFe powder is fully quasicrystalline (spectrum 1). The milling process does not change the phase composition of the powder, it just leads to marked changes in the intensity and width (β) of diffraction peaks for QC powder (spectrum 2), and either significant diminution of the particles' size or high microstrains formed seem to be responsible for much larger β. XRD spectra for reinforcing QC powders correlate well with SEM images of their appearance (Fig. 1a, b). It is of importance that the surface of particles (Fig. 1a, b) differs significantly — it appears relatively smooth after the water atomization process (Fig. 1a), and becomes extremely rough after milling (Fig. 1b). This difference is known to be a key feature in formation of strong interfacial bonding in composites [1,3]— the higher the roughness of the particles' surface, the stronger the interfacial bonding. Sometimes, f-QC particles appear as agglomerates with a greater size, and seem to be disintegrated easily at SPD induced by the UIP process. XRD spectra for composite layers formed at the UIP process are shown in Fig. 2. Spectrum 1 corresponds to the Al specimen covered by the composite layer reinforced with f-QC powder. The c-QC particle-reinforced layer (spectrum 2) also contains appropriate reflections of AlCuFe I-phase of slightly higher intensity. No additional peaks were observed. As compared to the spectrum for the annealed specimen, diffraction peaks of aluminium in spectra for UIP-treated specimens are slightly broadened and noticeably shifted to lower diffraction angles. The SEM images of surface morphology shown in back-scattered electron images in Fig. 3a, e demonstrate uniform distribution of QC particles in both composite layers formed at the UIP process. Some
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Fig. 2. XRD spectra of surface layers of composite reinforced with fine (1) or coarse (2) QC particles.
reinforcing particles are not visible, and one can see it only in copper and iron X-ray maps (Fig. 3b, c, f, g) — it seems to be a result of their possible coating with ductile aluminium at the slice preparation. Good coincidence of both elemental maps allows being aware of the absence of probable contamination of mechanically mixed composite layer with iron from steel pins at the UIP process. Such contamination was reported to occur at dry sliding wear of Al–SiC composites against steel counter-part [23]. Moreover, presented images provide evidences that f-QC particles gather sometimes in larger conglomerates (see arrows in Fig. 3f, g). On the other hand, the particles size of c-QC powder appears much smaller than that of the initial one — i.e. refinement of some c-QC particles occurs at their implantation into Al matrix. Histograms presented in Fig. 3d, h compare typical size distribution of reinforcing particles (or conglomerates) dispersed in composite layers.
Key feature like breaking up of c-QC particle is shown in Fig. 4a. Appearance of reinforcing particles of both types and their interfacial boundaries with surrounding matrix can be also compared using the images presented in Fig. 4. It appears that the reinforcements of both types are bonded to the matrix tightly enough. However, cracking of c-QC particles could cause weakening of their interfacial bonding to matrix. Similar situation is observed in the SEM images of specimens' cross-sections (Fig. 5a, b) — just a few broken c-QC particles are visible in Fig. 4a while f-QC particles are gathered in conglomerates sometimes. One can assess a thickness of both composite layers formed — it appears to be about 50–60 μm, which correlate well to the rough estimation performed for the through-depth profiles of inplane residual stresses (Fig. 5d, e). These rough estimates (Fig. 5d, e) testify that compressive stresses are mainly concentrated within the composite layers formed, and some fluctuations are observed on the boundary between reinforced layer and the matrix beneath, which is under tension. Compressive residual stresses measured using sin 2ψ method in both composite layers are slightly different in magnitude: approx. −95 and −115 MPa for f-QC and c-QC reinforcements, respectively (Fig. 5c, Table 1). 3.2. TEM investigations Dislocation structures of Al matrix in the top subsurface (about 5– 15 μm from the treated surface) of both composite layers are shown in Fig. 6a, b in comparison to that of the UIP-treated Al (Fig. 6c). When compared to the initial state, the dislocation/grain structure of the Al specimen underwent apparent changes at the UIP process (Fig. 6c). As usual, it consists in dislocation-cell structures with low angle boundaries [13,24]. Size of subgrains is larger than 1–2 μm. Generally, density of dislocations is rather low in the cell interiors. However, dislocation bundles can be observed sometimes.
Fig. 3. SEM images of morphological features of composite surfaces reinforced with fine (a–d) and coarse (e–h) AlCuFe particles observed in back scattering electrons (a, e), in Cu Kα (b, f) and Fe Kα (c, g) X-ray irradiations along with histograms of size distributions of reinforcements (d, h). Arrows in (f), (g) indicate conglomerates of fine particles.
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Fig. 4. SEM images demonstrated appearances of coarse (a) and fine (b) QC particles.
Addition of QC particles modifies the dislocation/grain structure of the Al matrix substantially (Fig. 6a, b). Well-formed dislocation-cell structure dominates in the composite layer contained c-QC particles (Fig. 6a). Frequently, the dislocation bundles within the cell interiors turn into tangles, and some of dislocation tangles patch along the cell boundaries thickening them. Size of dislocation cells is slightly smaller (approx. 0.5–1 μm), and density of dislocations in the cell interiors is rather high. An electron diffraction pattern is usual to a coarse-grained FCC metal. On the contrary, a lot of azimuthally diffused diffraction spots which corresponded to two first interplanar distances of Al (2.33 and 2.02) is visible in electron diffraction pattern registered for the composite layer reinforced with f-QC particles (Fig. 6b), indicating the presence of many small highly misoriented Al grains within the
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diffraction area. Their average size is of 200–400 nm. The dislocation density is even lower than that in the UIP-treated Al specimen. Peculiarities of interfacial boundaries for two studied cases can also be considered using TEM images shown in Fig. 6b, d, e. The dislocation concentration is increased significantly in the Al matrix surrounded c-QC particle (Fig. 6d). In the case of reinforcement with fQC particles, interfacial regions are multiform. Some of f-QC, particles appeared as intergranular particles in the grains’ corners, seem to be strongly bonded to the Al matrix (arrows in Fig. 6b). Such an observation is consistent to the opinion of authors [12,25,26] reported an existence of several crystallographic orientation relationships between the fcc Al matrix and the intergranular QC I-phase. At the same time, along with strongly adhered sphere-shaped f-QC particles surrounded with dislocation tangles, some of them appear encircled with bright borders that indicate supposedly their low bonding to the Al matrix (Fig. 6e). It is therefore believed that the interface between AlCuFe particles and the Al matrix is semi-coherent. 3.3. Wear behaviour Fig. 7 demonstrates marked decrease of a wear loss of Al/QC composite layers as compared to the annealed Al specimen. Almost double augmentation of the wear resistance of Al/QC composite layers is observed when measured in quasi-static conditions. Even in the dynamic tests, such enhancement is observed in wear behaviour but to a lesser extent — about 20% and 30% for reinforcement with f-QC and c-QC, respectively. Magnitudes of friction force (Fig. 8) attained from the studied specimens are most interesting when viewed in tandem with the wear data. Considering results of quasi-static tests, the highest friction force of the annealed Al can evidently be attributed to a localized
Fig. 5. SEM images of cross sections of the surface layers reinforced with coarse (a) and fine (b) QC particles; and dependencies of average lattice strains on sin2ψ (c) and estimates for changes of lattice spacing (d) and through-depth profiles of residual stresses (e) for composite layers reinforced with fine (1) or coarse (2) QC particles.
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Table 1 Parameters of surface layers of Al specimens. Material
Initial Al UIP + fine QC UIP + coarse QC
Hμ, MPa
288.3 734.8 783.5
σres, MPa (sin2ψ)
Quasi-static wear test Wst, μm
Wst/Wst(Al)
Fst, N
Fst/Fst(Al)
Wd, μm
Wd/Wd(Al)
Fd, N
Fd/Fd(Al)
− − 95 − 115
47.9 31.5 25.3
1.0 0.66 0.53
18.2 16.2 11.7
1.0 0.89 0.66
198.7 165.3 134.7
1.0 0.83 0.68
3.91 3.89 3
1.0 0.99 0.77
welding of the worn debris to the Al surface, and it leads to the lowest wear resistance [3,27]. The friction force of surface composite layers is lower in 10% and 30% for reinforcement with f-QC and c-QC particles, respectively. In the dynamic test, diminution of the friction force is observed only for c-QC reinforced layer, and it achieves approx. 20%. 4. Discussion Most authors conclude that the addition of reinforcing particles to the matrix alloy greatly decreases the wear rate, and this effect is enhanced as the volume fraction of particles increases [3,28–30]. High wear resistance of hard particle reinforced MMCs is reported to be primarily attributed to decrease in real area of contact due to high hardness of composites reinforced with hard particles [31]. Besides, mechanical properties of the matrix of composite are also responsible for overall performance [32]. Thus, the wear rate of a material is known to decrease with increasing hardness [33]. However, this decrease is frequently not linear as required by Archard's law, and cannot be described using simple mixture rule [23,33]. It is due to the
Dynamic wear test
fact that several factors, viz. microstructural changes in the composite matrix [1–4,23] and interfacial reinforcement/matrix strength [12,25,28,34] should be accounted for. Reinforcement interfaces play an important role on the mechanical properties of the MMCs, and strong bonding of reinforcement with the Al matrix has been verified to be the control factor that affects the wear improvement of the composite [34]. Increased microhardness of coatings or surface layers is known to reduce the ploughing component of friction which decreases the friction force [3,35]. The low coefficient of friction (that correlates to the friction force) and enhanced wear resistance were also reported in Al/nano Al2O3 particles composite produced by FSP [27] or in Al/SiC MMCs [36]. In this study, microhardness of composite layers reinforced with both types of QC powders is nearly twice higher than that of the UIP-treated Al specimen (Table 1), and this increase is evidently caused by change of microstructure of the Al matrix and by QC particles dispersed in surface layers. Results obtained in this study regarding the wear behaviour (Fig. 7) and the friction force (Fig. 8) can be analyzed in accordance with the above-provided brief description of the literature data, and
Fig. 6. TEM observations of the matrix microstructure and appropriate diffraction patterns in top subsurface layer in Al specimens after reinforcement with coarse (a) or fine (b) QC particles, after the UIP only (c) and interfacial regions adjacent to coarse QC particle (d) and fine ones (e). Fine QC particles are indicated by arrows in (b, e).
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Fig. 7. Wear losses for annealed Al specimen (0) and for composite layers reinforced with fine (1) or coarse (2) QC particles measured in quasi-static (S) or dynamic (D) regimes.
Fig. 8. Magnitudes of friction force for annealed Al specimen (0) and for composite layers reinforced with fine (1) or coarse (2) QC particles measured in quasi-static (S) or dynamic (D) regimes.
this analysis can be based on the observed microstructural features affecting the sliding wear. These features are summed up as follows: (i) average size of reinforcing particles (in the case of c-QC particles the “load bearing effect” could be expected [30,37,38]); (ii) their volume fraction and homogeneity of distribution (similar for both reinforcements); (iii) strength of interfacial bonding (seems to be slightly higher in the case of f-QC particles); (iv) increased density of dislocations (much higher in the case of reinforcement with c-QC particles); (v) differences in the grain/cell size of the Al matrix, i.e. in the volume fraction of grain boundaries or dense dislocation walls (the former is higher in the case of f-QC reinforcement, and the latter – in the case of c-QC). Thus, analyzing the size of reinforcement one should be aware of the fact that significant difference in size of initial reinforcing particles is partially diminished at the implantation process. Frequently, c-QC particles appear broken (Figs. 3a, 4a, 5a), and their average size becomes twice lower (Fig. 3d). Nevertheless, these formed particles is still effective regarding the operation of so-called “load bearing effect”, which is frequently used to describe the enhanced hardness and wear resistance of composite materials [30,37,38]. It is important to emphasize that only sufficiently large reinforcing particles can support the contact stresses, as if diminishing the load applied at measurements of microhardness or at wear tests. Fine particles are inoperative in this sense [37–39], in spite of the fact that sometimes they were observed agglomerated in larger aggregations (Fig. 3f–h). Distribution of reinforcing particles can though be qualified as homogeneous. Moreover, SEM analysis shows almost identical volume fraction of both reinforcements in surface layers of composite formed — analyzed both along the treated surface (Fig. 3a, e) and along the layers' thickness (Fig. 5a, b). Strength of interfacial bonding can be analyzed on the bases of SEM and TEM observations. It is believed that matrix/reinforcement interface is sufficiently strong in both cases studied (Figs. 4, and 6b, d, e), although some weakening of interfacial strength seems to occur at breaking up of c-QC particles (Figs. 4a, and 5a) or when f-QC particles appear encircled with light borders (Fig. 6e). A semi-coherent interfacial bonding between the Al matrix and QC phase was also observed in Al alloys contained QC precipitates [12,25,26], and in our previous study regarding SPD of QC particle-reinforced surface layers of Al specimens [13]. Although, dislocation/grain structure of aluminium matrix (Fig. 6a– c) was analyzed in the top subsurface (~5–15 μm) in this study, one could doubtfully extrapolate these data to whole composite layers (~50–60 μm), if the results of almost constant microhardness (Hμ) within the layers of similar thickness [13] would be taken into account.
Magnitudes of Hμ were measured to be twice higher when compared to un-reinforced Al [13]. We observe similar increase in Hμ in this study (Table 1). Authors of [40] reported the large increase in yield stress of composites produced by severe plastic deformation (ECAP), it is compared to that obtained by an increase in volume fraction or decrease in the reinforcement size. Considering that the reinforcing particles deform only elastically it was pointed out that the increase in yield stress is due to the two hardening parameters — decrease in matrix grain size and increase in internal dislocation density. In this study, fine grain structure contributes greater to wear and microhardness of the layer reinforced with f-QC particles, when the latter arrest grain boundaries [41], facilitating annihilation of moving dislocations that approach these boundaries [42]. As a result, large lattice rotation occurs in neighbour grains (i.e. the misorientation of boundaries enlarges), and highly misorientated grain structure with thin grain boundaries is formed in the composite layer reinforced with f-QC particles (Fig. 6b, e). Mean grain size is essentially low (~ 200– 400 nm). This observation correlate with other works [7,12,13,41– 44], analyzing formation of fine grain structure in Al-MMCs due to the introduction of fine second-phase particles into a zone of SPD — i.e. as a result of the inhibition of grain boundary migration. On the contrary, dislocation walls are effective in the case of reinforcement with c-QC particles. These particles act mainly as additional dislocation sources (Fig. 6d) and provoke increasing density of dislocations, thickening of dislocation walls and formation of dislocation tangles in the cell interiors (Fig. 6a, d). Similar dislocation activity near interfacial regions between the Al matrix and reinforcements was also reported in Al–SiO2[45], Al–SiC [2,46] and Al-QC [13] composites. It is known to be a result of large strain misfit caused by significant difference between elastic properties of matrix and ceramic or QC reinforcements. Resulting increased dislocation density is reported to cause the growth of damping capacity, yield strength and hardness of Al-MMCs [13,45] and Al alloys [47,48]. At the same time, slower dislocation accumulation in the aluminium matrix at fine reinforcement was reported to lead to reduced stress concentrations at the matrix/reinforcement interface and to produce less weakening at those interfaces [41]. A final aspect worth discussing is a merit of compressive residual stresses formed at the UIP process in the improvement of the wear resistance. As we report in Subsection 3.1, relatively high magnitudes of compressive residual stresses (Fig. 5c–e, Table 1) were registered for both composite layers. Naturally, they should contribute to enhanced wear performance (Fig. 7). Indeed, it is well known [49] that when a load is applied to a particle-reinforced MMC, stresses
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concentrate around the coarse particles. In the wear test, voids formed around reinforcing particles tend to coalesce and form micro-cracks. Material removal is linked to the growth of such cracks. Compressive stresses formed at the UIP process essentially suppress the initiation of cracks and their subsequent coalescence. Lower increment in wear resistance of composite layers measured at dynamic tests is naturally related to the high strain rate of load used at this test scheme and its variable character (Fig. 7, Table 1). Accelerated relaxation of compressive residual stresses under the influence of alternating component of load causes the situation when the initiation of cracks around reinforcing particles proceeds more easily and their subsequent coalescence could be simplified. Presence of the alternating component of load leads also to significant lowering of the magnitude of the friction force in comparison to those observed at quasi-static tests (Fig. 8, Table 1). 5. Summarizing remarks It was shown in this study that composite layers can be formed in aluminium by addition of paraffin suspensions contained AlCuFe fine powders in a zone of SPD in the UIP process. It was also demonstrated that the size of reinforcing particles affects significantly the microstructure and wear of the Al matrix. With almost similar volume fraction of uniformly dispersed reinforcing particles of various sizes, microstructures in top subsurface layers (5–10 μm) of composite formed differ greatly. The average grain size of the Al matrix is much lower in the case of f-QC reinforcements in comparison to c-QC ones — 200–400 nm and 1–2 μm, respectively. It is due to efficient inhibition of movement of grain/subgrain boundaries by f-QC particles that promotes relatively rapid increase of subgrains' misorientation and grain refinement. On the contrary, well-formed dislocation-cell structure dominates in the composite layer contained c-QC particles, and density of dislocations in the cell interiors is rather high. Both factors, the increased volume fractions of grain boundaries and dislocation walls, play a key role in the wear behaviour of the composite layers obtained. The microstructural features affecting the sliding wear can be summed up as follows: (i) average size of reinforcing particles; (ii) strength of interfacial bonding (seems to be slightly higher in the case of f-QC particles); (iii) increased density of dislocations (much higher in the case of reinforcement with c-QC particles); (iv) differences in the grain/cell size of the Al matrix, i.e. in the volume fraction of grain boundaries or dense dislocation walls (the former is higher in the case of f-QC reinforcement, and the latter — in the case of c-QC). Homogeneous dispersion of QC reinforcing particles of different size is quite effective in enhancement of wear resistance of Al surface layers. Taking into account increased density of dislocations, developed dislocation cell structures and/or a reduced grain size of Al matrix formed at SPD induced by the UIP process one can expect that these factors would together provide a substantial increase in wear and hardness of Al-MMC layers.
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Contents lists available at ScienceDirect
Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
Structure and wear of Al surface layers reinforced with AlCuFe particles using ultrasonic impact peening: Effect of different particle sizes B.N. Mordyuk a,⁎, M.O. Iefimov b, K.E. Grinkevych b, A.V. Sameljuk b, M.I. Danylenko b a b
Kurdyumov Institute for Metal Physics, 36 Academician Vernadsky Boulevard, UA-03680, Kyiv, Ukraine Frantzevich Institute for Problems of Materials Science, 3 Krzhyzhanivsky Street, UA-03142, Kyiv, Ukraine
a r t i c l e
i n f o
Article history: Received 13 December 2010 Accepted in revised form 19 May 2011 Available online 6 June 2011 Keywords: Surface composite Quasicrystal Aluminium Microstructure Ultrasonic impact peening Wear
a b s t r a c t Near-surface layers in aluminium specimens are modified using quasicrystalline (QC) AlCuFe particles introduced into a zone of severe plastic deformation induced by ultrasonic impact peening (UIP). Two types of QC particles are used: atomized with average size of approx. 25 μm (coarse QC–c-QC) and milled — 0.3–0.5 μm (fine QC–f-QC). The effect of QC particles of different sizes on microstructure and wear resistance of subsurface composite layers in aluminium is studied in this paper. XRD, SEM and TEM studies of reinforced aluminium layers allow establishing the links between microstructural features of the layers and their sliding wear. The formed layers of composites reinforced with both types of QC particles demonstrate almost double increment in wear resistance when compared to that of annealed aluminium. It is due to the combination of several factors: (i) high hardness and high wear resistance of QC reinforcement (more efficient for c-QC); (ii) relatively strong interfacial bonding of homogeneously dispersed reinforcing QC particles; (iii) fine grain structure of the Al matrix (f-QC) or increased density of dislocations arranged in fine dislocation-cell structure (c-QC) — i.e. increased volume fraction of grain boundaries/dense dislocation walls. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Metal-matrix composites (MMCs) exhibit wide engineering applications due to their enhanced hardness, wear, creep and fatigue resistance in comparison to the matrix metal or alloy [1–3]. Composite layer in a component surface is known to be sufficient or desirable in many applications — mainly to improve the surface dependent properties of material and the operation life of component or construction [2–4]. A number of recent studies show wide possibilities in fabrication of composite layers by dispersing of hard ceramic or intermetallic particles in a surface of Al (or Al based alloys) using either melting related techniques [5,6] or methods of severe plastic deformation (SPD) [7–9]. Use of quasicrystalline (QC) particles to reinforce AlMMCs seems to be prospective because QC materials have unique combination of high hardness and modulus of elasticity, high wear resistance and relatively low density. Unlike metastable QC reinforcements usually formed as precipitates to strengthening Albased alloys by rapid solidification techniques [10–12], stable QC particles (as AlCuFe) can be dispersed in modified surface layer being introduced into a zone of SPD. As a result, more stable materials can be obtained. They would not deteriorate above some critical temperature (temperature of phase transformation ‘quasicrystalline–crystalline’)
⁎ Corresponding author. Tel./fax: + 380 44 424 0521. E-mail address: [email protected] (B.N. Mordyuk). 0257-8972/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2011.05.046
as it usually occurs in Al based alloys contained metastable QC precipitates. Our previous studies show beneficial effects of QC particles dispersed in a surface layer of Al specimens on its microhardness and damping characteristics, which are increased significantly [13]. However, a size of reinforcing particles is known to affect the matrix structure and mechanical properties of the formed composite [1–5]. The aim of this study is two-fold: to examine the effect of AlCuFe particles of different sizes on microstructure of the Al matrix in surface layers formed at the UIP process, and to establish the links between microstructures formed in the QC particle-reinforced layers and their sliding wear performance.
2. Experimental details 2.1. Preparation of composite layers Plane specimens (60 ×15× 2.5 mm) of cp aluminium (purity of 99.75%) were initially annealed at 400 °C for 1 h in vacuum of 1.3 ×10− 3 Pa. The resulting grain size was about 60 μm. Atomized powder of icosahedral AL63Cu25Fe12 quasicrystal was used for reinforcement. An average particle size of initial atomized powder was approx. 25 μm (Fig. 1a); it is denoted “coarse” below (c-QC). Other powder (denoted “fine” in this paper — f-QC) was obtained by grinding of the initial powder in an ultrasonic mill [14] down to average particle size of 0.3–0.5 μm (Fig. 1b). Suspensions of AlCuFe powders in liquid paraffin
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The surface morphology of composite layers was examined in Superprobe-733 scanning electron microscope (SEM) using characteristic X-ray radiation. Transmission electron microscopy (TEM) analysis was carried out using a JEM 100 CX-II microscope. The plane-view TEM foils prepared from the top sub-surface layer (~ 5–10 μm) were mechanically polished from the un-treated side to thickness about 30 μm followed by ionic polishing (also from the un-treated side). 2.3. Wear examinations
Fig. 1. Appearance (a, b) and XRD spectra (c) of reinforcing powders: coarse AlCuFe powder after atomization and annealing (a, spectrum 1 in c) and fine one after milling (b, spectrum 2 in c).
were used to cover the surface of Al specimen in the UIP process to form composite layers. Equipment and appropriate regimes for the UIP process are described in detail in Refs. [13,15–17]. Basically, the UIP process consists in multiple sliding impacts imposed to a surface of treated specimen; impacts are induced by the pins positioned between specimen and an ultrasonic horn in a special impact head, which rotates forcedly at given velocity. The frequency of impacts in this loading scheme is about 3 kHz. Progressive shifts of the acoustic system along the treated surface for several times allow obtaining rather low surface roughness, relatively uniform distribution of reinforcements in composite layer, and sufficient thickness of this particle-reinforced surface layer (ten UIP passes performed for each specimen in this study produce the resulting layer with thickness of approx. 50–60 μm — similarly to our previous works [13,17,18]).
2.2. Microstructure examination The XRD θ–2θ analysis of reinforcing powders and treated specimens was carried out using a DART-UM1 diffractometer with Cu Kα irradiation. Assuming a biaxial stress state, analysis of residual stress in the Al matrix was carried out using a sin 2ψ based method [6,17,19]. An annealed Si powder deposited on a surface of an Al specimen was used as reference. In-plane residual stress σ|| = σ1 + σ2 was calculated using literature values of elastic modulus (E = 70.29 GPa) and Poison's ratio (ν = 0.3) for aluminium as an intercept of a linear graph εψi = f(sin2ψ) with the axis of ordinates (at sin2ψ= 0) [19]: σ‖ = −(E εψ = 0)/ν. Besides, considering different propagation depths of Xrays depending upon the diffraction angle, through depth profiles of inplane residual stress σ‖ were roughly estimated using the following expression σ1 + σ2 = −E(d− d0)/νd0[20,21], where d0 is a reference lattice spacing in a stress-free specimen and d is the lattice spacing in the UIP-treated specimen measured at different diffraction angles.
Wear resistance was examined using an automated tribocomplex (CATC) [22]. Sliding wear tests of plane specimens were performed with Si3N4 spherical indenter (8 mm in diameter) as counter-part in inactive liquid paraffin. Those tests consisted in a reciprocating sliding movement of indenter along a track of 4 mm, and lasted for 20 min. Sliding velocity was 0.013 m/s. The CATC allowed conducting tests of two types [22]: (i) usual quasi-static wear test performed at quasistatic load (20 N) applied in a normal direction to the tested surface, and (ii) so-called dynamic wear test, which was carried out using both quasi-static and alternating components of load applied in a normal direction to the tested surface. In the case of dynamic test, the obtained results are much closer to the operation conditions of studied material. Amplitude of the alternating component of load was approx. 10% regarding to that of quasi-static one. Wear magnitudes for quasi-static (Wst) and dynamic (Wd) modes were characterized with the depth of wear tracks measured using P-201 profilograph– profilometer. The friction force (F) was estimated via measured displacement of an elastic circular element rigidly connected to tested specimen. Besides, microhardness of the specimens was measured at load of 200 g on the Vickers indenter applied to their surface. 3. Results 3.1. XRD and SEM observations Results of X-ray structural analysis of reinforcing powders presented in Fig. 1c demonstrate that after atomization and subsequent annealing an AlCuFe powder is fully quasicrystalline (spectrum 1). The milling process does not change the phase composition of the powder, it just leads to marked changes in the intensity and width (β) of diffraction peaks for QC powder (spectrum 2), and either significant diminution of the particles' size or high microstrains formed seem to be responsible for much larger β. XRD spectra for reinforcing QC powders correlate well with SEM images of their appearance (Fig. 1a, b). It is of importance that the surface of particles (Fig. 1a, b) differs significantly — it appears relatively smooth after the water atomization process (Fig. 1a), and becomes extremely rough after milling (Fig. 1b). This difference is known to be a key feature in formation of strong interfacial bonding in composites [1,3]— the higher the roughness of the particles' surface, the stronger the interfacial bonding. Sometimes, f-QC particles appear as agglomerates with a greater size, and seem to be disintegrated easily at SPD induced by the UIP process. XRD spectra for composite layers formed at the UIP process are shown in Fig. 2. Spectrum 1 corresponds to the Al specimen covered by the composite layer reinforced with f-QC powder. The c-QC particle-reinforced layer (spectrum 2) also contains appropriate reflections of AlCuFe I-phase of slightly higher intensity. No additional peaks were observed. As compared to the spectrum for the annealed specimen, diffraction peaks of aluminium in spectra for UIP-treated specimens are slightly broadened and noticeably shifted to lower diffraction angles. The SEM images of surface morphology shown in back-scattered electron images in Fig. 3a, e demonstrate uniform distribution of QC particles in both composite layers formed at the UIP process. Some
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Fig. 2. XRD spectra of surface layers of composite reinforced with fine (1) or coarse (2) QC particles.
reinforcing particles are not visible, and one can see it only in copper and iron X-ray maps (Fig. 3b, c, f, g) — it seems to be a result of their possible coating with ductile aluminium at the slice preparation. Good coincidence of both elemental maps allows being aware of the absence of probable contamination of mechanically mixed composite layer with iron from steel pins at the UIP process. Such contamination was reported to occur at dry sliding wear of Al–SiC composites against steel counter-part [23]. Moreover, presented images provide evidences that f-QC particles gather sometimes in larger conglomerates (see arrows in Fig. 3f, g). On the other hand, the particles size of c-QC powder appears much smaller than that of the initial one — i.e. refinement of some c-QC particles occurs at their implantation into Al matrix. Histograms presented in Fig. 3d, h compare typical size distribution of reinforcing particles (or conglomerates) dispersed in composite layers.
Key feature like breaking up of c-QC particle is shown in Fig. 4a. Appearance of reinforcing particles of both types and their interfacial boundaries with surrounding matrix can be also compared using the images presented in Fig. 4. It appears that the reinforcements of both types are bonded to the matrix tightly enough. However, cracking of c-QC particles could cause weakening of their interfacial bonding to matrix. Similar situation is observed in the SEM images of specimens' cross-sections (Fig. 5a, b) — just a few broken c-QC particles are visible in Fig. 4a while f-QC particles are gathered in conglomerates sometimes. One can assess a thickness of both composite layers formed — it appears to be about 50–60 μm, which correlate well to the rough estimation performed for the through-depth profiles of inplane residual stresses (Fig. 5d, e). These rough estimates (Fig. 5d, e) testify that compressive stresses are mainly concentrated within the composite layers formed, and some fluctuations are observed on the boundary between reinforced layer and the matrix beneath, which is under tension. Compressive residual stresses measured using sin 2ψ method in both composite layers are slightly different in magnitude: approx. −95 and −115 MPa for f-QC and c-QC reinforcements, respectively (Fig. 5c, Table 1). 3.2. TEM investigations Dislocation structures of Al matrix in the top subsurface (about 5– 15 μm from the treated surface) of both composite layers are shown in Fig. 6a, b in comparison to that of the UIP-treated Al (Fig. 6c). When compared to the initial state, the dislocation/grain structure of the Al specimen underwent apparent changes at the UIP process (Fig. 6c). As usual, it consists in dislocation-cell structures with low angle boundaries [13,24]. Size of subgrains is larger than 1–2 μm. Generally, density of dislocations is rather low in the cell interiors. However, dislocation bundles can be observed sometimes.
Fig. 3. SEM images of morphological features of composite surfaces reinforced with fine (a–d) and coarse (e–h) AlCuFe particles observed in back scattering electrons (a, e), in Cu Kα (b, f) and Fe Kα (c, g) X-ray irradiations along with histograms of size distributions of reinforcements (d, h). Arrows in (f), (g) indicate conglomerates of fine particles.
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Fig. 4. SEM images demonstrated appearances of coarse (a) and fine (b) QC particles.
Addition of QC particles modifies the dislocation/grain structure of the Al matrix substantially (Fig. 6a, b). Well-formed dislocation-cell structure dominates in the composite layer contained c-QC particles (Fig. 6a). Frequently, the dislocation bundles within the cell interiors turn into tangles, and some of dislocation tangles patch along the cell boundaries thickening them. Size of dislocation cells is slightly smaller (approx. 0.5–1 μm), and density of dislocations in the cell interiors is rather high. An electron diffraction pattern is usual to a coarse-grained FCC metal. On the contrary, a lot of azimuthally diffused diffraction spots which corresponded to two first interplanar distances of Al (2.33 and 2.02) is visible in electron diffraction pattern registered for the composite layer reinforced with f-QC particles (Fig. 6b), indicating the presence of many small highly misoriented Al grains within the
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diffraction area. Their average size is of 200–400 nm. The dislocation density is even lower than that in the UIP-treated Al specimen. Peculiarities of interfacial boundaries for two studied cases can also be considered using TEM images shown in Fig. 6b, d, e. The dislocation concentration is increased significantly in the Al matrix surrounded c-QC particle (Fig. 6d). In the case of reinforcement with fQC particles, interfacial regions are multiform. Some of f-QC, particles appeared as intergranular particles in the grains’ corners, seem to be strongly bonded to the Al matrix (arrows in Fig. 6b). Such an observation is consistent to the opinion of authors [12,25,26] reported an existence of several crystallographic orientation relationships between the fcc Al matrix and the intergranular QC I-phase. At the same time, along with strongly adhered sphere-shaped f-QC particles surrounded with dislocation tangles, some of them appear encircled with bright borders that indicate supposedly their low bonding to the Al matrix (Fig. 6e). It is therefore believed that the interface between AlCuFe particles and the Al matrix is semi-coherent. 3.3. Wear behaviour Fig. 7 demonstrates marked decrease of a wear loss of Al/QC composite layers as compared to the annealed Al specimen. Almost double augmentation of the wear resistance of Al/QC composite layers is observed when measured in quasi-static conditions. Even in the dynamic tests, such enhancement is observed in wear behaviour but to a lesser extent — about 20% and 30% for reinforcement with f-QC and c-QC, respectively. Magnitudes of friction force (Fig. 8) attained from the studied specimens are most interesting when viewed in tandem with the wear data. Considering results of quasi-static tests, the highest friction force of the annealed Al can evidently be attributed to a localized
Fig. 5. SEM images of cross sections of the surface layers reinforced with coarse (a) and fine (b) QC particles; and dependencies of average lattice strains on sin2ψ (c) and estimates for changes of lattice spacing (d) and through-depth profiles of residual stresses (e) for composite layers reinforced with fine (1) or coarse (2) QC particles.
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Table 1 Parameters of surface layers of Al specimens. Material
Initial Al UIP + fine QC UIP + coarse QC
Hμ, MPa
288.3 734.8 783.5
σres, MPa (sin2ψ)
Quasi-static wear test Wst, μm
Wst/Wst(Al)
Fst, N
Fst/Fst(Al)
Wd, μm
Wd/Wd(Al)
Fd, N
Fd/Fd(Al)
− − 95 − 115
47.9 31.5 25.3
1.0 0.66 0.53
18.2 16.2 11.7
1.0 0.89 0.66
198.7 165.3 134.7
1.0 0.83 0.68
3.91 3.89 3
1.0 0.99 0.77
welding of the worn debris to the Al surface, and it leads to the lowest wear resistance [3,27]. The friction force of surface composite layers is lower in 10% and 30% for reinforcement with f-QC and c-QC particles, respectively. In the dynamic test, diminution of the friction force is observed only for c-QC reinforced layer, and it achieves approx. 20%. 4. Discussion Most authors conclude that the addition of reinforcing particles to the matrix alloy greatly decreases the wear rate, and this effect is enhanced as the volume fraction of particles increases [3,28–30]. High wear resistance of hard particle reinforced MMCs is reported to be primarily attributed to decrease in real area of contact due to high hardness of composites reinforced with hard particles [31]. Besides, mechanical properties of the matrix of composite are also responsible for overall performance [32]. Thus, the wear rate of a material is known to decrease with increasing hardness [33]. However, this decrease is frequently not linear as required by Archard's law, and cannot be described using simple mixture rule [23,33]. It is due to the
Dynamic wear test
fact that several factors, viz. microstructural changes in the composite matrix [1–4,23] and interfacial reinforcement/matrix strength [12,25,28,34] should be accounted for. Reinforcement interfaces play an important role on the mechanical properties of the MMCs, and strong bonding of reinforcement with the Al matrix has been verified to be the control factor that affects the wear improvement of the composite [34]. Increased microhardness of coatings or surface layers is known to reduce the ploughing component of friction which decreases the friction force [3,35]. The low coefficient of friction (that correlates to the friction force) and enhanced wear resistance were also reported in Al/nano Al2O3 particles composite produced by FSP [27] or in Al/SiC MMCs [36]. In this study, microhardness of composite layers reinforced with both types of QC powders is nearly twice higher than that of the UIP-treated Al specimen (Table 1), and this increase is evidently caused by change of microstructure of the Al matrix and by QC particles dispersed in surface layers. Results obtained in this study regarding the wear behaviour (Fig. 7) and the friction force (Fig. 8) can be analyzed in accordance with the above-provided brief description of the literature data, and
Fig. 6. TEM observations of the matrix microstructure and appropriate diffraction patterns in top subsurface layer in Al specimens after reinforcement with coarse (a) or fine (b) QC particles, after the UIP only (c) and interfacial regions adjacent to coarse QC particle (d) and fine ones (e). Fine QC particles are indicated by arrows in (b, e).
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Fig. 7. Wear losses for annealed Al specimen (0) and for composite layers reinforced with fine (1) or coarse (2) QC particles measured in quasi-static (S) or dynamic (D) regimes.
Fig. 8. Magnitudes of friction force for annealed Al specimen (0) and for composite layers reinforced with fine (1) or coarse (2) QC particles measured in quasi-static (S) or dynamic (D) regimes.
this analysis can be based on the observed microstructural features affecting the sliding wear. These features are summed up as follows: (i) average size of reinforcing particles (in the case of c-QC particles the “load bearing effect” could be expected [30,37,38]); (ii) their volume fraction and homogeneity of distribution (similar for both reinforcements); (iii) strength of interfacial bonding (seems to be slightly higher in the case of f-QC particles); (iv) increased density of dislocations (much higher in the case of reinforcement with c-QC particles); (v) differences in the grain/cell size of the Al matrix, i.e. in the volume fraction of grain boundaries or dense dislocation walls (the former is higher in the case of f-QC reinforcement, and the latter – in the case of c-QC). Thus, analyzing the size of reinforcement one should be aware of the fact that significant difference in size of initial reinforcing particles is partially diminished at the implantation process. Frequently, c-QC particles appear broken (Figs. 3a, 4a, 5a), and their average size becomes twice lower (Fig. 3d). Nevertheless, these formed particles is still effective regarding the operation of so-called “load bearing effect”, which is frequently used to describe the enhanced hardness and wear resistance of composite materials [30,37,38]. It is important to emphasize that only sufficiently large reinforcing particles can support the contact stresses, as if diminishing the load applied at measurements of microhardness or at wear tests. Fine particles are inoperative in this sense [37–39], in spite of the fact that sometimes they were observed agglomerated in larger aggregations (Fig. 3f–h). Distribution of reinforcing particles can though be qualified as homogeneous. Moreover, SEM analysis shows almost identical volume fraction of both reinforcements in surface layers of composite formed — analyzed both along the treated surface (Fig. 3a, e) and along the layers' thickness (Fig. 5a, b). Strength of interfacial bonding can be analyzed on the bases of SEM and TEM observations. It is believed that matrix/reinforcement interface is sufficiently strong in both cases studied (Figs. 4, and 6b, d, e), although some weakening of interfacial strength seems to occur at breaking up of c-QC particles (Figs. 4a, and 5a) or when f-QC particles appear encircled with light borders (Fig. 6e). A semi-coherent interfacial bonding between the Al matrix and QC phase was also observed in Al alloys contained QC precipitates [12,25,26], and in our previous study regarding SPD of QC particle-reinforced surface layers of Al specimens [13]. Although, dislocation/grain structure of aluminium matrix (Fig. 6a– c) was analyzed in the top subsurface (~5–15 μm) in this study, one could doubtfully extrapolate these data to whole composite layers (~50–60 μm), if the results of almost constant microhardness (Hμ) within the layers of similar thickness [13] would be taken into account.
Magnitudes of Hμ were measured to be twice higher when compared to un-reinforced Al [13]. We observe similar increase in Hμ in this study (Table 1). Authors of [40] reported the large increase in yield stress of composites produced by severe plastic deformation (ECAP), it is compared to that obtained by an increase in volume fraction or decrease in the reinforcement size. Considering that the reinforcing particles deform only elastically it was pointed out that the increase in yield stress is due to the two hardening parameters — decrease in matrix grain size and increase in internal dislocation density. In this study, fine grain structure contributes greater to wear and microhardness of the layer reinforced with f-QC particles, when the latter arrest grain boundaries [41], facilitating annihilation of moving dislocations that approach these boundaries [42]. As a result, large lattice rotation occurs in neighbour grains (i.e. the misorientation of boundaries enlarges), and highly misorientated grain structure with thin grain boundaries is formed in the composite layer reinforced with f-QC particles (Fig. 6b, e). Mean grain size is essentially low (~ 200– 400 nm). This observation correlate with other works [7,12,13,41– 44], analyzing formation of fine grain structure in Al-MMCs due to the introduction of fine second-phase particles into a zone of SPD — i.e. as a result of the inhibition of grain boundary migration. On the contrary, dislocation walls are effective in the case of reinforcement with c-QC particles. These particles act mainly as additional dislocation sources (Fig. 6d) and provoke increasing density of dislocations, thickening of dislocation walls and formation of dislocation tangles in the cell interiors (Fig. 6a, d). Similar dislocation activity near interfacial regions between the Al matrix and reinforcements was also reported in Al–SiO2[45], Al–SiC [2,46] and Al-QC [13] composites. It is known to be a result of large strain misfit caused by significant difference between elastic properties of matrix and ceramic or QC reinforcements. Resulting increased dislocation density is reported to cause the growth of damping capacity, yield strength and hardness of Al-MMCs [13,45] and Al alloys [47,48]. At the same time, slower dislocation accumulation in the aluminium matrix at fine reinforcement was reported to lead to reduced stress concentrations at the matrix/reinforcement interface and to produce less weakening at those interfaces [41]. A final aspect worth discussing is a merit of compressive residual stresses formed at the UIP process in the improvement of the wear resistance. As we report in Subsection 3.1, relatively high magnitudes of compressive residual stresses (Fig. 5c–e, Table 1) were registered for both composite layers. Naturally, they should contribute to enhanced wear performance (Fig. 7). Indeed, it is well known [49] that when a load is applied to a particle-reinforced MMC, stresses
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concentrate around the coarse particles. In the wear test, voids formed around reinforcing particles tend to coalesce and form micro-cracks. Material removal is linked to the growth of such cracks. Compressive stresses formed at the UIP process essentially suppress the initiation of cracks and their subsequent coalescence. Lower increment in wear resistance of composite layers measured at dynamic tests is naturally related to the high strain rate of load used at this test scheme and its variable character (Fig. 7, Table 1). Accelerated relaxation of compressive residual stresses under the influence of alternating component of load causes the situation when the initiation of cracks around reinforcing particles proceeds more easily and their subsequent coalescence could be simplified. Presence of the alternating component of load leads also to significant lowering of the magnitude of the friction force in comparison to those observed at quasi-static tests (Fig. 8, Table 1). 5. Summarizing remarks It was shown in this study that composite layers can be formed in aluminium by addition of paraffin suspensions contained AlCuFe fine powders in a zone of SPD in the UIP process. It was also demonstrated that the size of reinforcing particles affects significantly the microstructure and wear of the Al matrix. With almost similar volume fraction of uniformly dispersed reinforcing particles of various sizes, microstructures in top subsurface layers (5–10 μm) of composite formed differ greatly. The average grain size of the Al matrix is much lower in the case of f-QC reinforcements in comparison to c-QC ones — 200–400 nm and 1–2 μm, respectively. It is due to efficient inhibition of movement of grain/subgrain boundaries by f-QC particles that promotes relatively rapid increase of subgrains' misorientation and grain refinement. On the contrary, well-formed dislocation-cell structure dominates in the composite layer contained c-QC particles, and density of dislocations in the cell interiors is rather high. Both factors, the increased volume fractions of grain boundaries and dislocation walls, play a key role in the wear behaviour of the composite layers obtained. The microstructural features affecting the sliding wear can be summed up as follows: (i) average size of reinforcing particles; (ii) strength of interfacial bonding (seems to be slightly higher in the case of f-QC particles); (iii) increased density of dislocations (much higher in the case of reinforcement with c-QC particles); (iv) differences in the grain/cell size of the Al matrix, i.e. in the volume fraction of grain boundaries or dense dislocation walls (the former is higher in the case of f-QC reinforcement, and the latter — in the case of c-QC). Homogeneous dispersion of QC reinforcing particles of different size is quite effective in enhancement of wear resistance of Al surface layers. Taking into account increased density of dislocations, developed dislocation cell structures and/or a reduced grain size of Al matrix formed at SPD induced by the UIP process one can expect that these factors would together provide a substantial increase in wear and hardness of Al-MMC layers.
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