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Three body abrasion of laser surface alloyed aluminium AA1200
Wear 290–291 (2012) 1–9
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Three body abrasion of laser surface alloyed aluminium AA1200 L.A.B Mabhali a,b, N. Sacks a,n, S. Pityana b a b
School of Chemical and Metallurgical Engineering, University of the Witwatersrand, Private Bag 3, Wits 2050, South Africa Council for Scientific and Industrial Research, National Laser Centre, PO Box 375, Pretoria 0001, South Africa
a r t i c l e i n f o
abstract
Article history: Received 17 June 2011 Received in revised form 30 May 2012 Accepted 31 May 2012 Available online 9 June 2012
Laser surface alloying of aluminium AA1200 was performed with a 4 kW Nd:YAG laser to improve the abrasion wear resistance. Aluminium surfaces reinforced with metal matrix composites and intermetallic phases were achieved. The phases present depended on the composition of the alloying powder mixture. The wear performance of the alloyed surfaces was characterised using an ASTM G65 three body dry abrasion apparatus. A maximum 82% improvement in the wear resistance of the pure aluminium was achieved with a 40 wt% Ni þ 20 wt% Ti þ40 wt% SiC composition. The three alloys which had the best wear resistance were all produced with a composition of 40 wt% SiC and Ti and Ni powders ranging from 20 to 40 wt%. No direct correlation was observed between hardness and wear resistance. Microstructural examination showed that the main wear mechanisms were intense plastic deformation with micro-fracture of the SiC particles and intermetallic phases. The wear behaviour is mainly determined by the response of the different alloy phases, either independently or in combination, to the action of the abrasive particles and the precise nature of this response is complex and requires further study. & 2012 Elsevier B.V. All rights reserved.
Keywords: Laser surface alloying Aluminium Metal-matrix composites Intermetallics Three body-body abrasion
1. Introduction Aluminium is widely used in industry due to its attractive attributes such as low density, high strength to weight ratio, high thermal conductivity and good formability [1]. Its drawback is the poor surface properties such as hardness and wear resistance. These surface properties can be improved by depositing hard reinforcement materials on its surface in an attempt to provide a composite material which has a tough matrix reinforced with hard particles which should lead to improved wear resistance. This can be achieved by a laser alloying process. Alloying materials are deposited as powders into a melt pool generated on the material surface by a focused laser beam. The beam is scanned over the part and the deposited material resolidifies resulting in good bonding between the substrate and the alloyed layer. The surface of the material is modified by changing the composition and microstructure without affecting the bulk properties of the material. Process parameters such as laser power, laser beam spot size, laser scanning speed and powder feed rate have to be controlled to achieve the desired metallurgical bonding and alloyed surface properties. Research on the wear properties of aluminium alloys reinforced with either ceramic or metallic materials has been done by
n
Corresponding author. E-mail address: [email protected] (N. Sacks).
0043-1648/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.wear.2012.05.034
several authors [1–22]. Staia et al. [8] alloyed aluminium A356 with 96 wt% WC, 2 wt% Ti and 2 wt% Mg. The wear behaviour of the alloyed material against AISI52100 steel balls under a load of 5 N was investigated. For high laser scanning speeds, large WC particles were formed which served as the load carrying particles and severely abraded the steel. For low laser scanning speeds, the WC particle size decreased and the wear mechanisms changed as the aluminium matrix participated in the transference process. The wear mechanisms of the aluminium A356 was adhesive with high quantities of aluminium transferred to the steel counterface. Almeida et al. [9] studied the dry sliding wear mechanisms of Al–Mo deposited on an aluminium substrate by laser surface alloying. The wear mechanisms were predominantly adhesion followed by material detachment and transfer, oxidation and some abrasion, mainly by hard intermetallic compound particles on the steel counterbody. Elleuch et al. [10] studied the abrasive wear of aluminium alloys rubbed against sand. The author reported that the wear rate increased by three times as the incident angle of sand was increased from 01 to 451. S- ahin [11] studied the abrasive wear of aluminium AA2014 reinforced with SiC particles of different sizes (9 mm, 14 mm and 33 mm). The authors observed that the wear rate decreased with increasing hardness of the MMC while it increased with increasing abrasive particle size of the counter surface, applied load and sliding distance (up to a maximum value, then either decreased or remained constant). Miyajima and Iwai [12] reported that SiC and Al2O3 reinforcement caused severe wear of a steel counter surface
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during sliding wear of the aluminium matrix composite on a pinon-disc wear tester. Metal matrix composites with low volume fractions of embedded ceramics produced severe wear characterised by plastic deformation and large grooves. MMCs with high volume fractions of embedded ceramics did not wear severely and the surfaces were smooth and flat without large grooves. The embedded ceramic acted as inhibitors against plastic flow and adhesion of the matrix metal to the steel counter surface. Majumdar et al. [13] laser alloyed Ti with Si, Ti with Al, and Ti with SiþAl and studied the wear behaviour on a ball-on-disc wear testing machine. Surfaces alloyed with Al suffered extensive amounts of abrasive and adhesive wear while surfaces alloyed with Si þAl suffered abrasive wear. No significant amount of adhesive or abrasive wear was observed in samples alloyed only with Si. This decrease in the wear rates was associated with the formation of the Ti5Si3 phase which was found in higher volume fractions in the surfaces alloyed only with Si. Anandkumar et al. [15] conducted three body abrasion wear tests on Al–Si/SiC composite coatings using SiC as the abrasive. The coatings were produced using two different specific energies, namely 26 MJ/m2 and 58 MJ/m2 which yielded hardness values of 120 and 250 VH respectively. The higher hardness was attributed to the larger proportions of Al4SiC4 and Si precipitates. However the harder coating had the highest wear rate. The authors concluded that this was due to the significant difference between the SiC abrasive and the composite hardness. The lower wear rate achieved by the softer coating was attributed to the presence of SiC particles, which although do not provide an improvement in hardness, do resist scratching by the abrasive. The friction and wear behaviour of aluminium laser surface alloyed with SiC and SiCþAl was also studied by Majumdar et al. [21,22]. It was found that some the SiC dissociated to form Al4C3 which led to an increase in the aluminium hardness from 25 VHN to 200–250 VHN. The resulting wear resistance was improved up to three times and was attributed to the dispersion strengthening effect provided by the Al4C3 phase. Shipway et al. [14] studied the sliding wear mechanisms of aluminium reinforced with TiC particles against a carbon-manganese steel (BS 080A15) counterface and found that the wear rates increased as the applied load increased and delamination cracks were observed in the MMC layer. The addition of the TiC particles resulted in counter surface wear as hard particles acted as abrasives in the sliding process. The wear rates of the steel counter surfaces increased as the volume fraction of the TiC particles increased. These TiC particles resulted in ploughing and cutting of the steel counter surfaces in the sliding direction. In dry sliding wear tests of Al–12Si/TiB2 laser clads with an AISI 440C tool steel counterbody by Anandkumar et al. [17] the laser clads displayed ultra mild wear. The addition of TiB2 led to an improvement in the hardness and wear resistance of the Al–12Si alloy. The authors attributed this to the role played by the reinforcement phase in protecting the Al-based alloy against severe plastic deformation. Similar results were found by Majumdar et al. [20] who studied the abrasion wear resistance of aluminium reinforced with TiB and TiB2 particles where the volume fractions were varied between 7 and 18%. The addition of the TiB improved the hardness threefold using a laser power of 1.2 kW and a scan speed of 500 mm/min. The wear resistance was improved and decreased with an increase in applied load. This was attributed to the increased hardness provided by the TiB particles. Molybdenum has been shown to improve the mechanical properties of aluminium. In a study done by Almeida et al. [9] on a range of Al–Mo alloys (14.8–19.1 wt% Mo) the hardness was shown to increase from 85 to 160 VH while Young’s Modulus increased from 84 to 92 GPa. These increases were attributed to
increasing volume fractions of the intermetallic compounds. The wear resistance of the alloys were tested under dry sliding wear conditions using applied loads of 0.15 and 1 N respectively. The wear resistance was found to increase with an increase in the volume fraction of the intermetallic compounds. However higher wear rates were observed at the lower load. This result was attributed to the resulting wear mechanisms which included adhesion, oxidation as well as detachment and transfer of the surface material. Almeida et al. [18] reported that the addition of niobium to aluminium provided a good combination of a very high hardness and moderate toughness which was attributed to the volume fraction and dendritic spacing of the Al3Nb phase. Several authors have attempted to model the wear behaviour of metal matrix composites and general multiphase materials as some of the prevailing wear theories do not provide adequate insight into multiphase material behaviour under abrasive wear conditions [16,23–28]. Recently Colaco and Vilar [16] developed a mathematical model to provide understanding of the functional dependence of wear resistance on the properties of the reinforcement phases in metal matrix composites. The model was based on the Rabinowicz approach [29] on defining abrasive wear mechanisms of multiphase materials and also considered the mechanisms proposed by Hutchings [28] and Zum-Ghar [24] which reported that material loss occurs by wear of the matrix, wear of the reinforcement particles and extraction of the reinforcement particles from the matrix. The model was tested against experimental data based on three body abrasion wear of Fe-0.25%C-15%Cr coatings reinforced with different volume fractions of Nb. The authors reported a good correlation between the model and experimental data and provided detailed understanding of the functional dependence of the wear resistance on the proportion of reinforcement particles. The main conclusions showed that while the average hardness increases as the volume fraction of reinforcement increases, increased wear rates are possible when the volume fraction of reinforcement is increased where cracking and extraction of the particles may increase due to interaction with the abrasive. Published work is generally limited to higher aluminium alloys (i.e. 2xxx to 7xxx series). To the authors knowledge no published work is available on the abrasive wear of Aluminium AA1200 which has been laser alloyed with ceramic and metallic materials simultaneously. Therefore the aim of this work was to investigate the abrasive wear resistance of Aluminium AA1200 by laser surface alloying with mixed Ni, Ti and SiC powders. Intermetallic phases are formed when aluminium is laser alloyed with metallic materials (e.g. Ni and Ti) while metal matrix composites are formed when alloying with ceramics (e.g. SiC).
2. Material and methods Laser alloying of the aluminium AA1200 surface was performed with a Nd:YAG laser. Aluminium AA1200 plates were initially sand blasted to enhance the absorption of laser energy by the aluminium substrate. The chemical composition of the aluminium plates was 0.12 wt% Cu, 0.13 wt% Si, 0.59 wt% Fe and the balance was Al. The laser parameters used were 4 kW of power, a beam diameter of 4 mm and a scanning speed of 10 mm/s. These parameters ensured that sufficient laser energy was supplied for the dissolution of the powders. The alloying powder mixture consisted of different compositions of Ni, Ti and SiC powder mixtures as listed in Table 1. The powder feed rate was 2.5–3 g/ min (depending on the powder composition) which ensured sufficient supply of powder during the alloying experiments. Argon was used as the carrier and shielding gas to prevent oxidation during the alloying process. Ten overlapping tracks with a track overlap of 50% were made.
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Table 1 Starting powder compositions.
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Table 2 Vickers hardness of the alloyed surfaces.
Sample no.
Powder composition
Powder composition
Hardness (HV1)
Al 1 2 3 4 5 6 7 8 9 10 11 12 13 14
Untreated aluminium AA1200 10 wt% Ni þ70 wt% Tiþ 20 wt% SiC 20 wt% Ni þ30 wt% Tiþ 50 wt% SiC 20 wt% Ni þ40 wt% Tiþ 40 wt% SiC 30 wt% Ni þ30 wt% Tiþ 40 wt% SiC 30 wt% Ni þ40 wt% Tiþ 30 wt% SiC 33.3 wt% Ni þ 33.3 wt% Tiþ33.3 wt% SiC 40 wt% Ni þ20 wt% Tiþ 40 wt% SiC 40 wt% Ni þ30 wt% Tiþ 30 wt% SiC 40 wt% Ni þ40 wt% Tiþ 20 wt% SiC 50 wt% Ni þ20 wt% Tiþ 30 wt% SiC 60 wt% Ni þ30 wt% Tiþ 10 wt% SiC 70 wt% Ni þ10 wt% Tiþ 20 wt% SiC 70 wt% Ni þ20 wt% Tiþ 10 wt% SiC 80 wt% Ni þ15 wt% Tiþ 5 wt% SiC
Untreated aluminium AA1200 10 wt% Ni þ 70 wt% Ti þ20 wt% SiC 20 wt% Ni þ 30 wt% Ti þ50 wt% SiC 20 wt% Ni þ 40 wt% Ti þ40 wt% SiC 30 wt% Ni þ 30 wt% Ti þ40 wt% SiC 30 wt% Ni þ 40 wt% Ti þ30 wt% SiC 33.3 wt% Ni þ33.3 wt% Ti þ33.3 wt% SiC 40 wt% Ni þ 20 wt% Ti þ40 wt% SiC 40 wt% Ni þ 30 wt% Ti þ30 wt% SiC 40 wt% Ni þ 40 wt% Ti þ20 wt% SiC 50 wt% Ni þ 20 wt% Ti þ30 wt% SiC 60 wt% Ni þ 30 wt% Ti þ10 wt% SiC 70 wt% Ni þ 10 wt% Ti þ20 wt% SiC 70 wt% Ni þ 20 wt% Ti þ10 wt% SiC 80 wt% Ni þ 15 wt% Tiþ 5 wt% SiC
24.03 7 0.27 143.697 9.32 149.537 10.10 172.387 13.63 184.487 6.92 192.217 8.19 197.437 7.53 198.197 13.49 218.90 7 6.41 239.687 10.28 250.10 7 7.75 264.277 9.39 290.42 7 11.86 295.817 11.18 318.647 9.13
Fig. 1. Wear response of untreated aluminium and laser alloyed surfaces.
The hardness was measured using a 1 kgf load and the average of eleven indentations is reported as the surface hardness. The microstructures of the alloyed surfaces were studied using scanning electron microscopy (SEM) and x-ray diffraction (XRD). Samples were ground and polished to a 0.04 mm (colloidal silica OP-S suspension) surface finish and etched with Keller’s reagent (3 ml HClþ2 ml HFþ5 ml HNO3 þ190 ml HO2) prior to microstructural characterisation. Three body abrasion wear tests were conducted using an ASTM G65-04 [30] dry sand rubber wheel apparatus shown in Fig. 1. Silica sand obtained from Rolfes (Pty) Ltd. with a particle size range of 0.3–0.65 mm was used as the abrasive. Prior to wear testing, the silica sand was sieved with a MACSALAB electronic sieve shaker to determine the particle size and distribution. Majority of the sand particles were in the 500–600 mm range with a D50 of 525 mm and were generally angular in shape. The sand feed rate was 2.3 g/s. This feed rate ensured that sufficient sand was introduced into the rubber wheel/specimen contact area in the test chamber. The samples, 70 20 6 mm in size, were polished to a 1 mm surface finish prior to the wear tests. The applied load was kept constant at 9.8 N. The duration of each wear test was 60 min and the mass of the samples was recorded at 10 min intervals. Three tests were conducted to determine the average for each alloy. The wear scars were examined using a JEOL JSM-840 scanning electron microscope.
3. Results Laser alloying was successfully performed with the laser processing parameters used. Tables 2 and 3 list the hardness
and phases of the alloyed surfaces. Metal matrix composites reinforced with SiC particles and intermetallic phases were formed in the alloyed layers. The predominant intermetallic phases were Al3Ni, Al3Ni2, Al3Ti, TiC and Ti5Si3. The Al3Ni, Al3Ni2 and Al3Ti phases were formed due to the reaction of Al with Ni and Ti. The dissociation of SiC particles and reactions with Ti resulted in the formation of TiC and Ti5Si3 phases. In alloys with a high SiC content, Al4C3, Al4SiC4, Ti3SiC2 and Al3TiC2 intermetallic phases were also formed from the reaction of the dissociated SiC with either Al or Ti. A comprehensive study of the microstructures produced by each powder composition and the methods used for phase identification (EDS, XRD, phase diagrams, thermodynamic calculations) can be found in Ref. [31]. The hardness results show that laser alloying improved the surface hardness for all the compositions used. Grain refinement, due to the rapid heating and cooling rates associated with laser alloying plays a role in increasing the hardness of the laser alloyed surfaces [32]. The highest hardness was obtained when alloying with high Ni contents. Fig. 1 shows the typical wear behaviour of the materials during the wear test. The wear response of all the alloys was initially high during the first 10 min of testing followed by a levelling off of the wear rate for the remainder of the test. Fig. 2 illustrates the wear rates of all the alloyed surfaces after 60 min of wear. A reduction in wear rate (or improved wear resistance) was found in all the alloyed surfaces with the maximum reduction in wear of 82% observed when laser alloying with 40 wt% Niþ20 wt% Tiþ40 wt% SiC. No correlation was observed between wear response and hardness. The 80 wt% Niþ15 wt% Tiþ5 wt% SiC alloy which had the highest hardness did not show the best wear resistance. This alloy has the ninth highest wear rate of the fifteen materials. The best wear resistance was shown by the alloy which only had the eighth highest hardness. From the SEM images of the worn surfaces similar features were found on the wear scars of all the alloys. The degree to which the different types of mechanisms occurred depended on the alloying composition and phases present. Figs. 3–6 illustrate the typical wear scars observed. Fig. 3 shows the wear scars on the pure Al. Due to the high ductility of aluminium, some of the fragmented SiO2 particles were embedded in the deformed microstructure. During testing the rubber wheel pushes the sand against the specimen leading to point indentation of the SiO2 particles on the aluminium surface under the applied load. Fig. 3(a) and (b) shows plastic deformation of the surface which is characterised by non-uniform ridges and plastic flow. The extent of plastic deformation on the surface is also evident on the cross-sections of the aluminium shown in
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Table 3 Phases identified in the alloyed surfaces. Powder composition
Phases observed
10 wt% Ni þ 70 wt% Tiþ 20 wt% SiC 20 wt% Ni þ 30 wt% Tiþ 50 wt% SiC 20 wt% Ni þ 40 wt% Tiþ 40 wt% SiC 30 wt% Ni þ 30 wt% Tiþ 40 wt% SiC 30 wt% Ni þ 40 wt% Tiþ 30 wt% SiC 33.3 wt% Ni þ 33.3 wt% Tiþ 33.3 wt% SiC 40 wt% Ni þ 20 wt% Tiþ 40 wt% SiC 40 wt% Ni þ 30 wt% Tiþ 30 wt% SiC 40 wt% Ni þ 40 wt% Tiþ 20 wt% SiC 50 wt% Ni þ 20 wt% Tiþ 30 wt% SiC 60 wt% Ni þ 30 wt% Tiþ 10 wt% SiC 70 wt% Ni þ 10 wt% Tiþ 20 wt% SiC 70 wt% Ni þ 20 wt% Tiþ 10 wt% SiC 80 wt% Ni þ 15 wt% Ti þ5 wt% SiC
a-Al, SiC, TiC, Al3Ni, Al3Ti and Ti5Si3 a-Al, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti, Ti5Si3, Al3TiC2 and Al4SiC4 a-Al, Si, SiC, TiC, Al3Ni, Al3Ti, Ti5Si3, Ti3SiC2 and Al4SiC4 a-Al, Si, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti, Ti5Si3, Ti3SiC2 and Al4SiC4 a-Al, Si, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti, Ti5Si3, Ti3SiC2 and Al4SiC4 a-Al, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti and Ti5Si3 a-Al, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti, Ti5Si3, Ti3SiC2 and Al4SiC4 a-Al, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti, Ti5Si3 and Ti3SiC2 a-Al, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti, Ti5Si3 and Ti3SiC2 a-Al, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti, Ti5Si3 and Al4C3 a-Al, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti and Ti5Si3 a-Al, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti, Ti5Si3 and Al4C3 a-Al, SiC, Al3Ni, Al3Ni2 and Al3Ti a-Al, SiC, Al3Ni, Al3Ni2 and Al3Ti
Fig. 2. Wear rate of untreated and laser alloyed aluminium surfaces. Numbers correspond to alloy compositions listed in Table 1.
Fig. 3(c) and (d) which also highlights the extent and depth to which the fragmented SiO2 abrasive particles were embedded in the deformed aluminium. A few micro-cracks close to the surface indicate that material removal occurred by progressive detachment of the deformed material. The worn surfaces of the alloyed layers was characterised by plastic deformation mechanisms, fracture of the embedded SiC particles and micro-fracture of the intermetallic phases as presented in Figs. 4–6. The plastic deformation observed in Figs. 4 and 5 is characterised by a non-uniform ridge structure where partial delamination (Fig. 5(c)) of the deformed surface layers can also be seen. Random paths of smearing are also present; these paths were likely caused by the movement of wear debris during the wear test. Sub-surface plastic deformation occurred in the laser alloyed materials. This was evident in the cross-sections of the worn surfaces where distortions in the microstructures were observed, for example uniform changes in grain orientation in the direction of motion of the abrasive. This deformation is caused by the cyclic action of the abrasive over the surface under the influence of the applied load. Lateral cracking seen in the cross-sectional views was transgranular and contributed to delamination of the worn surface layer as seen in Fig. 6(a) and (b). Fracturing of the SiC particles (Fig. 6(c–e)) and transgranular cracking of the intermetallic matrix (Fig. 6(f)) was also observed.
4. Discussion A reduction in wear rate was observed for the laser alloyed surfaces compared to the pure aluminium metal under three body
abrasion wear conditions. A maximum reduction in wear rate of 82% (compared to pure Al) was observed when laser alloying with a 40 wt% Niþ20 wt% Tiþ40 wt% SiC powder composition. This enhanced surface abrasive wear performance may be attributed to a good combination of the tough aluminium matrix reinforced with hard metal matrix composites and intermetallic and ceramic phases. One of the main results of the study was that no direct correlation was observed between hardness and wear resistance. Similar trends were found by others [16]. The alloy with the highest hardness namely, the 80 wt% Niþ 15 wt% Ti þ5 wt% SiC composition had the ninth highest wear rate of the fifteen materials tested, while the best wear resistance was shown by the alloy which only had the eighth highest hardness. The general increase in hardness of the alloys is attributed to the presence of the carbide (TiC, Al4SiC4,) and aluminide (Al3Ni, Al3N2) phases. TiC has a hardness range of 2000–3000 HV [33] and was present in alloys having both high hardness and high wear resistance, except for the two hardest alloys which showed no traces of TiC. The aluminides, Al3Ni and Al3Ni2, have hardnesses of 732 HV and 1013 HV respectively [34] and were present in most alloys. One of the main phase differences in the five alloys which had the best wear resistance and the five alloys which had the highest hardness was the presence of the Ti3SiC2 and Al4SiC4 phases in the alloys with the best wear resistance; these phases were not present in alloys having high hardness values. Ti3SiC2 is known to have a low hardness, but has a high Young’s modulus and a high damage tolerance [35] while Al4SiC4 has a hardness of 1200 HV [15]. The presence of Al3Ti in all the alloys also led to a strength increase in the alloys [36]. Despite the high hardness values of some of the individual phases the overall average hardness of the alloyed layers was in the range of 144–319 HV1. Detailed analysis and discussion of the phases produced in the alloyed surfaces and their contribution to hardness can be found in references [31,37]. For the current discussion on the response of the materials to abrasive wear it is sufficient to state that the test results indicate that there is not a linear relationship between alloy hardness and wear rate and therefore cannot be used as a primary parameter to control the wear resistance of these alloys. The wear behaviour is mainly determined by the response of the different alloy phases, either independently or in combination, to the action of the abrasive particles. It has been shown that there is a direct correlation between the microstructural properties, bonding and volume fraction of the reinforcement particles in metal matrix composites and intermetallic phases and their relationship to abrasive wear [16,24,33]. The influence of volume fraction of reinforcement phase on the abrasive wear has been studied by several authors;
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Fig. 3. Worn surfaces of the untreated aluminium in (a) and (b) plan view; (c) and (d) cross-sections showing gross plastic deformation as well as embedded SiO2 abrasive particles.
Fig. 4. SEM plan view images of wear scars of the laser alloyed surfaces showing typical features observed in all the alloyed layers. Images taken at low magnifications.
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Fig. 5. SEM plan view images of wear scars of the laser alloyed surfaces showing typical features observed in all the alloyed layers. Images taken at high magnifications.
however contradictory results have been obtained. Some authors reported that as the volume fraction increased the wear resistance increased [23], while others observed the opposite trend [24]. In an attempt to address these contradictions Colaco and Vilar [16] developed a model which showed that while the increase in volume fraction of reinforcement phase led to increased hardness, the wear rates were influenced by the action of the abrasive on the reinforcements which can cause cracking and detachment of the reinforcements which may subsequently lead to increased wear rates. To a certain extent these factors determine if the alloyed aluminium surfaces would respond homogenously or heterogeneously under the abrasive wear test conditions. The type of response expected has been related to a comparison between the size of the reinforcement particles and the dimension of damage caused by the abrasive [24,28]. If the sizes are similar then a heterogeneous response is expected and the hardness properties of abrasive and reinforcement will determine the wear rates. If the size of the reinforcement particles is much smaller than the dimensions of the damage then the response is homogenous. In
the latter case the wear rates are influenced by the brittle behaviour of the particles and matrix, the hardness of the reinforcements and the abrasive as well as the volume fraction of reinforcements present. These factors influence the wear mechanisms which will prevail and determine the resulting wear rates. In the current study the worn surfaces did not display parallel grooving which is typically observed under abrasion wear testing. Rather gross plastic deformation as well as brittle fracture of the SiC and intermetallic phases was observed. Thus the dimension of damage done by the SiO2 abrasive could not be calculated. To accurately assess this aspect it should be possible to conduct scratch tests on the surfaces and to analyse the resulting wear behaviour. This was not attempted in the current study. A partial understanding of the material response can be gained by studying the worn surfaces and comparing properties between the phases present and the abrasive used. The largest reinforcement particle size present was that of the retained SiC particles which had a maximum size of at least 50 mm. The average SiO2 abrasive particle size was 525 mm. On this basis it is likely that the alloyed
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Fig. 6. SEM images of cross-sections of the worn surfaces of the laser alloyed surfaces showing fractured SiC grains and micro-fracture of the intermetallic phases.
layers responded homogenously during wear. The precise nature of the response in terms of linking increasing wear rates to particle hardness, reinforcement volume fraction and phase brittleness is more complex due to the wide range of reinforcement phases present. If the alloyed surfaces were treated as having a ‘pure’ homogeneous response and the hardness of the SiO2 abrasive, 800 HV [33], is compared to the composite hardness, which ranged from 144 to 319 HV1, the ratio of the hardness of the abrasive particles to the hardness of the alloyed surfaces (Ha/Hm) is in the range of 2.5–5.6. These calculated ratios are greater than 1.2 which signifies the boundary condition between soft and hard abrasion as defined by Hutchings [33] and would therefore place the abrasive wear of the alloyed surfaces in the hard abrasion category. This classification would confer with the extent of hard abrasion damage mechanisms observed on the wear scars of the alloyed surfaces. If the individual reinforcement hardnesses are used, then the Ha/Hm ratios span a range of values on either side of the
1.2 boundary condition and the wear rates are then influenced by the ability of the SiO2 abrasive to plastically deform, fracture or cause detachment of the particles from the matrix and would be strongly influenced by the volume fraction of each reinforcement phase present, the matrix properties and the applied stress during testing. The applied stress was calculated to be 0.02 MPa but the actual pressure is expected to be higher than this as the real contact area between individual abrasive particles and the alloyed surface is lower than the surface area of 420 mm2 used to calculate the nominal stress. For each alloyed layer produced the strength properties of the specific phases (matrix and reinforcement) present, would determine that phase’s response to the applied pressure. This would lead to different wear mechanisms being experienced by the different phases and ultimately result in a variety of measured wear rates. The influence of increasing reinforcement volume fraction showed a variable wear response and cannot be easily quantified as the alloying compositions were not kept constant for a wide
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range of compositions. However a few deductions can be made. When the Ni volume fraction was kept constant at 40 wt%, the wear rate increased as the Ti volume fraction increased along with a simultaneous decrease in the SiC volume fraction. When the Ti volume fraction was kept constant at 40 wt%, the wear rate increased with a decrease in SiC content, but no linear effect was observed as the Ni volume fraction was simultaneously decreased. When the SiC volume fraction was kept constant at 40 wt%, no linear relationships were observed with increasing the Ti and decreasing the Ni volume fractions simultaneously. It is noted that the three alloys which had the best wear resistance were all produced with a composition of 40 wt% SiC and Ti and Ni powders ranging from 20 to 40 wt%. Due to the influence of laser processing conditions on the amount of dissolution and reaction rates, the final proportions of each phase in the alloyed layers are likely to change. Thus in order to state which composition would ultimately produce the best wear resistance the final proportions of each phase in the alloyed layers should be accurately determined. Due to the complex nature of the alloy system investigated and the variety in phase morphologies present this proved to be difficult; even qualitative results by EDS is considered misleading. Abrasive wear generally occurs by plastic flow, plastic deformation, brittle fracture or combinations of the different modes [18] and in this work all modes were observed. The presence of intermetallic phases such as Al3Ni and Al3Ti together with the soft a-Al in the inter-dendritic regions constituted a favourable combination with these hard phases resisting abrasion and the soft aluminium phase suppressing macro-crack growth in these regions. The same behaviour was observed by Man et al. [36] who studied the abrasive wear of aluminium AA6061 laser surface alloyed with NiTi and reported that wear generally occurred by micro-cracking and flaking of the surface. The brittleness of the material has been shown to increase with increasing proportion of reinforcement particles [23]. In a study on the relationship between hardness and brittleness of intermetallic phases Guo et al. [38] reported that hardness and brittleness of the intermetallic phases play a role in determining which phase would have more wear resistance. Phases that have a high hardness and are less brittle will have better wear resistance but when ductility is low (for brittle phases) fracture occurs leading to high wear rates. This may be the reason for the high wear rates observed by the high-Ni alloys which showed the highest hardness, but did not have the best wear resistance mainly due to the brittle fracture of the aluminides. The role of the mean free path between reinforcements has also been reported to play a role in the wear resistance of multiphase materials. Axe´n and Zum Gahr [39] studied the sliding wear resistance of a tool steel 90MnCrV8 reinforced with TiC particles and concluded that the mean free path (defined as the average spacing between reinforcement particles) also plays an important role during wear. Matrix protection is improved by a reduced mean free path between the reinforcing particles. Large reinforcement particles result in an increased mean free path which ultimately results in reduced wear resistance. The effect of the size of the reinforcements was also found to depend on the matrix structure. Small reinforcement particles are of greater advantage within a hard matrix than large particles within a soft matrix [39]. It is well known that a high volume fraction of small particles will result in low mean free paths which would generally lead to a high wear resistance; this is based on the theory of dispersion strengthening [33]. In the current study the alloyed layers with small Al mean free paths had high wear rates as brittle fracture was promoted by the formation of high densities of the hard and brittle intermetallic and ceramic phases in which cracking of the intermetallic phases and fracturing of the SiC
particles were observed. The Ti containing and the Al–SiC intermetallic phases resulted in large Al mean free paths. This promoted wear by plastic deformation. In a few instances the hard phases, e.g. the equiaxed dendritic Al3Ni2 grains and the SiC particles, were also seen to interrupt plastic deformation. Fracture of the SiC particles and intermetallic phases can be reduced if narrower distributions of starting powders were used to ensure either complete dissolution of the SiC particles or to limit retention of larger particles. It has been shown that if the extent of cracking or extraction of reinforcement particles is negligible, then the abrasive wear resistance may be improved [16]. The hard intermetallic phases probably cracked due to repeated loading and the stresses applied during wear and these micro-cracks facilitated material removal as wear progressed. Fracturing of the embedded SiC particles also occurred due to cyclic loading but the strong bonding between the particles and the matrix ensured that decohesion (interfacial debonding and pull-out) of the SiC particles was prevented. A strong interfacial bond between the reinforcements and the matrix is important. When the bond is weak the reinforcements are easily pulled out during wear and can act as additional abrasives thereby increasing the wear rate, as was observed by reference [39]. Due to the strong interfacial bonding and prior to fracture, the embedded SiC particles in the MMCs probably led to lower wear rates in alloys where these particles were present in high volume fractions. This strong bond between the matrix and the reinforcements is formed by a reaction between the ceramic particles and the metal matrix at the metal/ceramic interface and therefore it is not easy to detach these particles from the matrix even though they have fractured. In wear tests conducted on Al–7 wt% Si laser alloyed with SiC, Anandkumar et al. [15] observed discontinuous wear grooves confined to the matrix on the worn surfaces and reported that the hard SiC reinforcement particles interrupted the abrasive grooving action. Deeper analysis is required to fully understand the response of the alloyed layers to abrasive wear and to quantify the role of each individual phase in contributing to the overall wear performance. Existing models based on the abrasive wear of multiphase materials are generally limited to the addition of one reinforcement phase in varying proportions to a constant matrix. In the current work three reinforcement materials were added to a constant matrix which produced a wide variety of reinforcement phases. Due to the complex nature of the investigated systems a model to understand the abrasive wear has not been attempted in the present work. Existing models may be used to explain the behaviour of individual phases, e.g. fracture of the SiC and microcracking of intermetallic particles. However the application of these models is limited as they do not provide a holistic view of both the individual and the combined phase response of the investigated system under abrasive wear conditions.
5. Conclusions In general laser alloying improved the wear resistance of the aluminium AA1200 surface with an improvement of 19–82% was achieved. The three alloys which had the best wear resistance were all produced with a composition of 40 wt% SiC and Ti and Ni powders ranging from 20 to 40 wt%. No direct correlation was observed between hardness and wear resistance. The alloy with the highest hardness had the ninth highest wear rate while the best wear resistance was shown by the alloy which only had the eighth highest hardness. The predominant wear mechanisms were intense plastic deformation with micro-fracture of the SiC particles and intermetallic phases. The wear behaviour is mainly determined by the response of the different alloy phases, either
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independently or in combination, to the action of the abrasive particles and the precise nature of this response in terms of linking increasing wear rates to specific phases, particle hardness, reinforcement volume fraction and phase brittleness is complex due to the wide range of reinforcement phases present and requires further study.
Acknowledgements The authors gratefully acknowledge the financial support of the Department of Science and Technology and the Council for Scientific and Industrial Research in South Africa. The University of the Witwatersrand in South Africa is also acknowledged for technical support. References [1] S. Tomida, K. Nakata, Surface and Coatings Technology 174–175 (2003) 559–563. [2] L.R. Katipelli, A. Agarwal, N.B. Dahotre, Applied Surface Science 153 (2000) 65–78. [3] H.C. Man, S. Zhang, F.T. Cheng, Materials Letters 61 (2007) 4058–4061. [4] A. Almeida, et al., Surface and Coatings Technology 70 (1995) 221–229. [5] N. Ravi, et al., Materials and Manufacturing Processes 15 (3) (2000) 395–404. [6] M.H. Staia, M. Cruz, N.B. Dahotre, Wear 251 (2001) 1459–1468. [7] M. Kok, Journal of Materials Processing Technology 161 (2005) 381–387. [8] M.H. Staia, M. Cruz, N.B. Dahotre, Thin Solid Films 377–378 (2000) 665–674. [9] A. Almeida, et al., Surface and Coatings Technology 200 (2006) 4782–4790. [10] K. Elleuch, S. Mezlini, N. Guermazi, Ph. Kapsa, Wear 261 (2006) 1316–1321. [11] Y. S- ahin, Tribology International 43 (2010) 939–943. [12] T. Miyajima, Y. Iwai, Wear 255 (2003) 606–616.
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[13] J.D. Majumdar, B.L. Mordike, I. Manna, Wear 242 (2000) 18–27. [14] P.H. Shipway, A.R. Kennedy, A.J. Wilkes, Wear 216 (1998) 160–171. [15] R. Anandkumar, et al., Surface and Coatings Technology 201 (2007) 9497–9505. [16] R. Colaco, R. Vilar, Wear 254 (2003) 625–634. [17] R. Anandkumar, et al., Surface and Coatings Technology 205 (2011) 3824–3832. [18] A. Almeida, et al., Materials Science and Engineering A 303 (2001) 273–280. [19] A. Almeida, R. Vilar, Scripta Materialia 63 (2010) 811–814. [20] J.D. Majumdar, et al., Surface and Coatings Technology 201 (2006) 1236–1242. [21] J.D. Majumdar, et al., Wear 262 (2007) 641–648. [22] J.D. Majumdar, et al., Materials Science Engineering A 433 (2006) 241–250. [23] J.W. Liou, L.H. Chen, T.S. Lui, Journal of Materials Science 30 (1995) 258–262. [24] K.-H. Zum-Ghar, Tribology International 31 (10) (1998) 587–596. [25] M.M. Krushkov, Wear 28 (1974) 69–88. [26] W.M. Garrison, Wear 82 (1982) 213–220. [27] N. Axen, I.M. hutchings, S. Jacobson, Tribology International 29 (6) (1996) 467–475. [28] I.M. Hutchings, Materials Science and Technology 10 (1994) 513–517. [29] E. Rabinowicz, Friction and Wear of Materials, Wiley, New York, 1965. [30] American Standard Testing Methods, Standard Test Method for Measuring Abrasion Using the Dry Sand/Rubber Wheel Apparatus, ASTM G65-04, 2004. [31] L.A.B. Mabhali, Laser Surface Alloying and In-Situ Formation of Aluminium Metal Composites Reinforced With Ceramics and Intermetallics, Ph.D. Thesis, University of the Witwatersrand, South Africa, 2011. [32] Z. Brytan, et al., Journal of Achievements in Materials and Manufacturing Engineering 40 (1) (2010) 70–78. [33] I.M Hutchings, Tribology: Friction and Wear of Engineering Materials, Edward Arnold, 2002. [34] M.A. Zamzam, et al., Materials and Design 23 (2002) 161-138. [35] J. Lis, et al., Ceramics International 19 (1993) 219–222. [36] H.C. Man, S. Zhang, F.T. Cheng, Materials Letters 61 (2007) 4058–4061. [37] L.A.B. Mabhali, et al., Molecular Crystals and Liquid Crystals 555 (1) (2012) 138–148. [38] B. Guo, et al., Surface and Coatings Technology 202 (2008) 4121–4129. [39] N. Axe´n, K.H. Zum Gahr, Wear 157 (1992) 189–201.
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Three body abrasion of laser surface alloyed aluminium AA1200 L.A.B Mabhali a,b, N. Sacks a,n, S. Pityana b a b
School of Chemical and Metallurgical Engineering, University of the Witwatersrand, Private Bag 3, Wits 2050, South Africa Council for Scientific and Industrial Research, National Laser Centre, PO Box 375, Pretoria 0001, South Africa
a r t i c l e i n f o
abstract
Article history: Received 17 June 2011 Received in revised form 30 May 2012 Accepted 31 May 2012 Available online 9 June 2012
Laser surface alloying of aluminium AA1200 was performed with a 4 kW Nd:YAG laser to improve the abrasion wear resistance. Aluminium surfaces reinforced with metal matrix composites and intermetallic phases were achieved. The phases present depended on the composition of the alloying powder mixture. The wear performance of the alloyed surfaces was characterised using an ASTM G65 three body dry abrasion apparatus. A maximum 82% improvement in the wear resistance of the pure aluminium was achieved with a 40 wt% Ni þ 20 wt% Ti þ40 wt% SiC composition. The three alloys which had the best wear resistance were all produced with a composition of 40 wt% SiC and Ti and Ni powders ranging from 20 to 40 wt%. No direct correlation was observed between hardness and wear resistance. Microstructural examination showed that the main wear mechanisms were intense plastic deformation with micro-fracture of the SiC particles and intermetallic phases. The wear behaviour is mainly determined by the response of the different alloy phases, either independently or in combination, to the action of the abrasive particles and the precise nature of this response is complex and requires further study. & 2012 Elsevier B.V. All rights reserved.
Keywords: Laser surface alloying Aluminium Metal-matrix composites Intermetallics Three body-body abrasion
1. Introduction Aluminium is widely used in industry due to its attractive attributes such as low density, high strength to weight ratio, high thermal conductivity and good formability [1]. Its drawback is the poor surface properties such as hardness and wear resistance. These surface properties can be improved by depositing hard reinforcement materials on its surface in an attempt to provide a composite material which has a tough matrix reinforced with hard particles which should lead to improved wear resistance. This can be achieved by a laser alloying process. Alloying materials are deposited as powders into a melt pool generated on the material surface by a focused laser beam. The beam is scanned over the part and the deposited material resolidifies resulting in good bonding between the substrate and the alloyed layer. The surface of the material is modified by changing the composition and microstructure without affecting the bulk properties of the material. Process parameters such as laser power, laser beam spot size, laser scanning speed and powder feed rate have to be controlled to achieve the desired metallurgical bonding and alloyed surface properties. Research on the wear properties of aluminium alloys reinforced with either ceramic or metallic materials has been done by
n
Corresponding author. E-mail address: [email protected] (N. Sacks).
0043-1648/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.wear.2012.05.034
several authors [1–22]. Staia et al. [8] alloyed aluminium A356 with 96 wt% WC, 2 wt% Ti and 2 wt% Mg. The wear behaviour of the alloyed material against AISI52100 steel balls under a load of 5 N was investigated. For high laser scanning speeds, large WC particles were formed which served as the load carrying particles and severely abraded the steel. For low laser scanning speeds, the WC particle size decreased and the wear mechanisms changed as the aluminium matrix participated in the transference process. The wear mechanisms of the aluminium A356 was adhesive with high quantities of aluminium transferred to the steel counterface. Almeida et al. [9] studied the dry sliding wear mechanisms of Al–Mo deposited on an aluminium substrate by laser surface alloying. The wear mechanisms were predominantly adhesion followed by material detachment and transfer, oxidation and some abrasion, mainly by hard intermetallic compound particles on the steel counterbody. Elleuch et al. [10] studied the abrasive wear of aluminium alloys rubbed against sand. The author reported that the wear rate increased by three times as the incident angle of sand was increased from 01 to 451. S- ahin [11] studied the abrasive wear of aluminium AA2014 reinforced with SiC particles of different sizes (9 mm, 14 mm and 33 mm). The authors observed that the wear rate decreased with increasing hardness of the MMC while it increased with increasing abrasive particle size of the counter surface, applied load and sliding distance (up to a maximum value, then either decreased or remained constant). Miyajima and Iwai [12] reported that SiC and Al2O3 reinforcement caused severe wear of a steel counter surface
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during sliding wear of the aluminium matrix composite on a pinon-disc wear tester. Metal matrix composites with low volume fractions of embedded ceramics produced severe wear characterised by plastic deformation and large grooves. MMCs with high volume fractions of embedded ceramics did not wear severely and the surfaces were smooth and flat without large grooves. The embedded ceramic acted as inhibitors against plastic flow and adhesion of the matrix metal to the steel counter surface. Majumdar et al. [13] laser alloyed Ti with Si, Ti with Al, and Ti with SiþAl and studied the wear behaviour on a ball-on-disc wear testing machine. Surfaces alloyed with Al suffered extensive amounts of abrasive and adhesive wear while surfaces alloyed with Si þAl suffered abrasive wear. No significant amount of adhesive or abrasive wear was observed in samples alloyed only with Si. This decrease in the wear rates was associated with the formation of the Ti5Si3 phase which was found in higher volume fractions in the surfaces alloyed only with Si. Anandkumar et al. [15] conducted three body abrasion wear tests on Al–Si/SiC composite coatings using SiC as the abrasive. The coatings were produced using two different specific energies, namely 26 MJ/m2 and 58 MJ/m2 which yielded hardness values of 120 and 250 VH respectively. The higher hardness was attributed to the larger proportions of Al4SiC4 and Si precipitates. However the harder coating had the highest wear rate. The authors concluded that this was due to the significant difference between the SiC abrasive and the composite hardness. The lower wear rate achieved by the softer coating was attributed to the presence of SiC particles, which although do not provide an improvement in hardness, do resist scratching by the abrasive. The friction and wear behaviour of aluminium laser surface alloyed with SiC and SiCþAl was also studied by Majumdar et al. [21,22]. It was found that some the SiC dissociated to form Al4C3 which led to an increase in the aluminium hardness from 25 VHN to 200–250 VHN. The resulting wear resistance was improved up to three times and was attributed to the dispersion strengthening effect provided by the Al4C3 phase. Shipway et al. [14] studied the sliding wear mechanisms of aluminium reinforced with TiC particles against a carbon-manganese steel (BS 080A15) counterface and found that the wear rates increased as the applied load increased and delamination cracks were observed in the MMC layer. The addition of the TiC particles resulted in counter surface wear as hard particles acted as abrasives in the sliding process. The wear rates of the steel counter surfaces increased as the volume fraction of the TiC particles increased. These TiC particles resulted in ploughing and cutting of the steel counter surfaces in the sliding direction. In dry sliding wear tests of Al–12Si/TiB2 laser clads with an AISI 440C tool steel counterbody by Anandkumar et al. [17] the laser clads displayed ultra mild wear. The addition of TiB2 led to an improvement in the hardness and wear resistance of the Al–12Si alloy. The authors attributed this to the role played by the reinforcement phase in protecting the Al-based alloy against severe plastic deformation. Similar results were found by Majumdar et al. [20] who studied the abrasion wear resistance of aluminium reinforced with TiB and TiB2 particles where the volume fractions were varied between 7 and 18%. The addition of the TiB improved the hardness threefold using a laser power of 1.2 kW and a scan speed of 500 mm/min. The wear resistance was improved and decreased with an increase in applied load. This was attributed to the increased hardness provided by the TiB particles. Molybdenum has been shown to improve the mechanical properties of aluminium. In a study done by Almeida et al. [9] on a range of Al–Mo alloys (14.8–19.1 wt% Mo) the hardness was shown to increase from 85 to 160 VH while Young’s Modulus increased from 84 to 92 GPa. These increases were attributed to
increasing volume fractions of the intermetallic compounds. The wear resistance of the alloys were tested under dry sliding wear conditions using applied loads of 0.15 and 1 N respectively. The wear resistance was found to increase with an increase in the volume fraction of the intermetallic compounds. However higher wear rates were observed at the lower load. This result was attributed to the resulting wear mechanisms which included adhesion, oxidation as well as detachment and transfer of the surface material. Almeida et al. [18] reported that the addition of niobium to aluminium provided a good combination of a very high hardness and moderate toughness which was attributed to the volume fraction and dendritic spacing of the Al3Nb phase. Several authors have attempted to model the wear behaviour of metal matrix composites and general multiphase materials as some of the prevailing wear theories do not provide adequate insight into multiphase material behaviour under abrasive wear conditions [16,23–28]. Recently Colaco and Vilar [16] developed a mathematical model to provide understanding of the functional dependence of wear resistance on the properties of the reinforcement phases in metal matrix composites. The model was based on the Rabinowicz approach [29] on defining abrasive wear mechanisms of multiphase materials and also considered the mechanisms proposed by Hutchings [28] and Zum-Ghar [24] which reported that material loss occurs by wear of the matrix, wear of the reinforcement particles and extraction of the reinforcement particles from the matrix. The model was tested against experimental data based on three body abrasion wear of Fe-0.25%C-15%Cr coatings reinforced with different volume fractions of Nb. The authors reported a good correlation between the model and experimental data and provided detailed understanding of the functional dependence of the wear resistance on the proportion of reinforcement particles. The main conclusions showed that while the average hardness increases as the volume fraction of reinforcement increases, increased wear rates are possible when the volume fraction of reinforcement is increased where cracking and extraction of the particles may increase due to interaction with the abrasive. Published work is generally limited to higher aluminium alloys (i.e. 2xxx to 7xxx series). To the authors knowledge no published work is available on the abrasive wear of Aluminium AA1200 which has been laser alloyed with ceramic and metallic materials simultaneously. Therefore the aim of this work was to investigate the abrasive wear resistance of Aluminium AA1200 by laser surface alloying with mixed Ni, Ti and SiC powders. Intermetallic phases are formed when aluminium is laser alloyed with metallic materials (e.g. Ni and Ti) while metal matrix composites are formed when alloying with ceramics (e.g. SiC).
2. Material and methods Laser alloying of the aluminium AA1200 surface was performed with a Nd:YAG laser. Aluminium AA1200 plates were initially sand blasted to enhance the absorption of laser energy by the aluminium substrate. The chemical composition of the aluminium plates was 0.12 wt% Cu, 0.13 wt% Si, 0.59 wt% Fe and the balance was Al. The laser parameters used were 4 kW of power, a beam diameter of 4 mm and a scanning speed of 10 mm/s. These parameters ensured that sufficient laser energy was supplied for the dissolution of the powders. The alloying powder mixture consisted of different compositions of Ni, Ti and SiC powder mixtures as listed in Table 1. The powder feed rate was 2.5–3 g/ min (depending on the powder composition) which ensured sufficient supply of powder during the alloying experiments. Argon was used as the carrier and shielding gas to prevent oxidation during the alloying process. Ten overlapping tracks with a track overlap of 50% were made.
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Table 1 Starting powder compositions.
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Table 2 Vickers hardness of the alloyed surfaces.
Sample no.
Powder composition
Powder composition
Hardness (HV1)
Al 1 2 3 4 5 6 7 8 9 10 11 12 13 14
Untreated aluminium AA1200 10 wt% Ni þ70 wt% Tiþ 20 wt% SiC 20 wt% Ni þ30 wt% Tiþ 50 wt% SiC 20 wt% Ni þ40 wt% Tiþ 40 wt% SiC 30 wt% Ni þ30 wt% Tiþ 40 wt% SiC 30 wt% Ni þ40 wt% Tiþ 30 wt% SiC 33.3 wt% Ni þ 33.3 wt% Tiþ33.3 wt% SiC 40 wt% Ni þ20 wt% Tiþ 40 wt% SiC 40 wt% Ni þ30 wt% Tiþ 30 wt% SiC 40 wt% Ni þ40 wt% Tiþ 20 wt% SiC 50 wt% Ni þ20 wt% Tiþ 30 wt% SiC 60 wt% Ni þ30 wt% Tiþ 10 wt% SiC 70 wt% Ni þ10 wt% Tiþ 20 wt% SiC 70 wt% Ni þ20 wt% Tiþ 10 wt% SiC 80 wt% Ni þ15 wt% Tiþ 5 wt% SiC
Untreated aluminium AA1200 10 wt% Ni þ 70 wt% Ti þ20 wt% SiC 20 wt% Ni þ 30 wt% Ti þ50 wt% SiC 20 wt% Ni þ 40 wt% Ti þ40 wt% SiC 30 wt% Ni þ 30 wt% Ti þ40 wt% SiC 30 wt% Ni þ 40 wt% Ti þ30 wt% SiC 33.3 wt% Ni þ33.3 wt% Ti þ33.3 wt% SiC 40 wt% Ni þ 20 wt% Ti þ40 wt% SiC 40 wt% Ni þ 30 wt% Ti þ30 wt% SiC 40 wt% Ni þ 40 wt% Ti þ20 wt% SiC 50 wt% Ni þ 20 wt% Ti þ30 wt% SiC 60 wt% Ni þ 30 wt% Ti þ10 wt% SiC 70 wt% Ni þ 10 wt% Ti þ20 wt% SiC 70 wt% Ni þ 20 wt% Ti þ10 wt% SiC 80 wt% Ni þ 15 wt% Tiþ 5 wt% SiC
24.03 7 0.27 143.697 9.32 149.537 10.10 172.387 13.63 184.487 6.92 192.217 8.19 197.437 7.53 198.197 13.49 218.90 7 6.41 239.687 10.28 250.10 7 7.75 264.277 9.39 290.42 7 11.86 295.817 11.18 318.647 9.13
Fig. 1. Wear response of untreated aluminium and laser alloyed surfaces.
The hardness was measured using a 1 kgf load and the average of eleven indentations is reported as the surface hardness. The microstructures of the alloyed surfaces were studied using scanning electron microscopy (SEM) and x-ray diffraction (XRD). Samples were ground and polished to a 0.04 mm (colloidal silica OP-S suspension) surface finish and etched with Keller’s reagent (3 ml HClþ2 ml HFþ5 ml HNO3 þ190 ml HO2) prior to microstructural characterisation. Three body abrasion wear tests were conducted using an ASTM G65-04 [30] dry sand rubber wheel apparatus shown in Fig. 1. Silica sand obtained from Rolfes (Pty) Ltd. with a particle size range of 0.3–0.65 mm was used as the abrasive. Prior to wear testing, the silica sand was sieved with a MACSALAB electronic sieve shaker to determine the particle size and distribution. Majority of the sand particles were in the 500–600 mm range with a D50 of 525 mm and were generally angular in shape. The sand feed rate was 2.3 g/s. This feed rate ensured that sufficient sand was introduced into the rubber wheel/specimen contact area in the test chamber. The samples, 70 20 6 mm in size, were polished to a 1 mm surface finish prior to the wear tests. The applied load was kept constant at 9.8 N. The duration of each wear test was 60 min and the mass of the samples was recorded at 10 min intervals. Three tests were conducted to determine the average for each alloy. The wear scars were examined using a JEOL JSM-840 scanning electron microscope.
3. Results Laser alloying was successfully performed with the laser processing parameters used. Tables 2 and 3 list the hardness
and phases of the alloyed surfaces. Metal matrix composites reinforced with SiC particles and intermetallic phases were formed in the alloyed layers. The predominant intermetallic phases were Al3Ni, Al3Ni2, Al3Ti, TiC and Ti5Si3. The Al3Ni, Al3Ni2 and Al3Ti phases were formed due to the reaction of Al with Ni and Ti. The dissociation of SiC particles and reactions with Ti resulted in the formation of TiC and Ti5Si3 phases. In alloys with a high SiC content, Al4C3, Al4SiC4, Ti3SiC2 and Al3TiC2 intermetallic phases were also formed from the reaction of the dissociated SiC with either Al or Ti. A comprehensive study of the microstructures produced by each powder composition and the methods used for phase identification (EDS, XRD, phase diagrams, thermodynamic calculations) can be found in Ref. [31]. The hardness results show that laser alloying improved the surface hardness for all the compositions used. Grain refinement, due to the rapid heating and cooling rates associated with laser alloying plays a role in increasing the hardness of the laser alloyed surfaces [32]. The highest hardness was obtained when alloying with high Ni contents. Fig. 1 shows the typical wear behaviour of the materials during the wear test. The wear response of all the alloys was initially high during the first 10 min of testing followed by a levelling off of the wear rate for the remainder of the test. Fig. 2 illustrates the wear rates of all the alloyed surfaces after 60 min of wear. A reduction in wear rate (or improved wear resistance) was found in all the alloyed surfaces with the maximum reduction in wear of 82% observed when laser alloying with 40 wt% Niþ20 wt% Tiþ40 wt% SiC. No correlation was observed between wear response and hardness. The 80 wt% Niþ15 wt% Tiþ5 wt% SiC alloy which had the highest hardness did not show the best wear resistance. This alloy has the ninth highest wear rate of the fifteen materials. The best wear resistance was shown by the alloy which only had the eighth highest hardness. From the SEM images of the worn surfaces similar features were found on the wear scars of all the alloys. The degree to which the different types of mechanisms occurred depended on the alloying composition and phases present. Figs. 3–6 illustrate the typical wear scars observed. Fig. 3 shows the wear scars on the pure Al. Due to the high ductility of aluminium, some of the fragmented SiO2 particles were embedded in the deformed microstructure. During testing the rubber wheel pushes the sand against the specimen leading to point indentation of the SiO2 particles on the aluminium surface under the applied load. Fig. 3(a) and (b) shows plastic deformation of the surface which is characterised by non-uniform ridges and plastic flow. The extent of plastic deformation on the surface is also evident on the cross-sections of the aluminium shown in
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Table 3 Phases identified in the alloyed surfaces. Powder composition
Phases observed
10 wt% Ni þ 70 wt% Tiþ 20 wt% SiC 20 wt% Ni þ 30 wt% Tiþ 50 wt% SiC 20 wt% Ni þ 40 wt% Tiþ 40 wt% SiC 30 wt% Ni þ 30 wt% Tiþ 40 wt% SiC 30 wt% Ni þ 40 wt% Tiþ 30 wt% SiC 33.3 wt% Ni þ 33.3 wt% Tiþ 33.3 wt% SiC 40 wt% Ni þ 20 wt% Tiþ 40 wt% SiC 40 wt% Ni þ 30 wt% Tiþ 30 wt% SiC 40 wt% Ni þ 40 wt% Tiþ 20 wt% SiC 50 wt% Ni þ 20 wt% Tiþ 30 wt% SiC 60 wt% Ni þ 30 wt% Tiþ 10 wt% SiC 70 wt% Ni þ 10 wt% Tiþ 20 wt% SiC 70 wt% Ni þ 20 wt% Tiþ 10 wt% SiC 80 wt% Ni þ 15 wt% Ti þ5 wt% SiC
a-Al, SiC, TiC, Al3Ni, Al3Ti and Ti5Si3 a-Al, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti, Ti5Si3, Al3TiC2 and Al4SiC4 a-Al, Si, SiC, TiC, Al3Ni, Al3Ti, Ti5Si3, Ti3SiC2 and Al4SiC4 a-Al, Si, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti, Ti5Si3, Ti3SiC2 and Al4SiC4 a-Al, Si, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti, Ti5Si3, Ti3SiC2 and Al4SiC4 a-Al, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti and Ti5Si3 a-Al, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti, Ti5Si3, Ti3SiC2 and Al4SiC4 a-Al, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti, Ti5Si3 and Ti3SiC2 a-Al, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti, Ti5Si3 and Ti3SiC2 a-Al, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti, Ti5Si3 and Al4C3 a-Al, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti and Ti5Si3 a-Al, SiC, TiC, Al3Ni, Al3Ni2, Al3Ti, Ti5Si3 and Al4C3 a-Al, SiC, Al3Ni, Al3Ni2 and Al3Ti a-Al, SiC, Al3Ni, Al3Ni2 and Al3Ti
Fig. 2. Wear rate of untreated and laser alloyed aluminium surfaces. Numbers correspond to alloy compositions listed in Table 1.
Fig. 3(c) and (d) which also highlights the extent and depth to which the fragmented SiO2 abrasive particles were embedded in the deformed aluminium. A few micro-cracks close to the surface indicate that material removal occurred by progressive detachment of the deformed material. The worn surfaces of the alloyed layers was characterised by plastic deformation mechanisms, fracture of the embedded SiC particles and micro-fracture of the intermetallic phases as presented in Figs. 4–6. The plastic deformation observed in Figs. 4 and 5 is characterised by a non-uniform ridge structure where partial delamination (Fig. 5(c)) of the deformed surface layers can also be seen. Random paths of smearing are also present; these paths were likely caused by the movement of wear debris during the wear test. Sub-surface plastic deformation occurred in the laser alloyed materials. This was evident in the cross-sections of the worn surfaces where distortions in the microstructures were observed, for example uniform changes in grain orientation in the direction of motion of the abrasive. This deformation is caused by the cyclic action of the abrasive over the surface under the influence of the applied load. Lateral cracking seen in the cross-sectional views was transgranular and contributed to delamination of the worn surface layer as seen in Fig. 6(a) and (b). Fracturing of the SiC particles (Fig. 6(c–e)) and transgranular cracking of the intermetallic matrix (Fig. 6(f)) was also observed.
4. Discussion A reduction in wear rate was observed for the laser alloyed surfaces compared to the pure aluminium metal under three body
abrasion wear conditions. A maximum reduction in wear rate of 82% (compared to pure Al) was observed when laser alloying with a 40 wt% Niþ20 wt% Tiþ40 wt% SiC powder composition. This enhanced surface abrasive wear performance may be attributed to a good combination of the tough aluminium matrix reinforced with hard metal matrix composites and intermetallic and ceramic phases. One of the main results of the study was that no direct correlation was observed between hardness and wear resistance. Similar trends were found by others [16]. The alloy with the highest hardness namely, the 80 wt% Niþ 15 wt% Ti þ5 wt% SiC composition had the ninth highest wear rate of the fifteen materials tested, while the best wear resistance was shown by the alloy which only had the eighth highest hardness. The general increase in hardness of the alloys is attributed to the presence of the carbide (TiC, Al4SiC4,) and aluminide (Al3Ni, Al3N2) phases. TiC has a hardness range of 2000–3000 HV [33] and was present in alloys having both high hardness and high wear resistance, except for the two hardest alloys which showed no traces of TiC. The aluminides, Al3Ni and Al3Ni2, have hardnesses of 732 HV and 1013 HV respectively [34] and were present in most alloys. One of the main phase differences in the five alloys which had the best wear resistance and the five alloys which had the highest hardness was the presence of the Ti3SiC2 and Al4SiC4 phases in the alloys with the best wear resistance; these phases were not present in alloys having high hardness values. Ti3SiC2 is known to have a low hardness, but has a high Young’s modulus and a high damage tolerance [35] while Al4SiC4 has a hardness of 1200 HV [15]. The presence of Al3Ti in all the alloys also led to a strength increase in the alloys [36]. Despite the high hardness values of some of the individual phases the overall average hardness of the alloyed layers was in the range of 144–319 HV1. Detailed analysis and discussion of the phases produced in the alloyed surfaces and their contribution to hardness can be found in references [31,37]. For the current discussion on the response of the materials to abrasive wear it is sufficient to state that the test results indicate that there is not a linear relationship between alloy hardness and wear rate and therefore cannot be used as a primary parameter to control the wear resistance of these alloys. The wear behaviour is mainly determined by the response of the different alloy phases, either independently or in combination, to the action of the abrasive particles. It has been shown that there is a direct correlation between the microstructural properties, bonding and volume fraction of the reinforcement particles in metal matrix composites and intermetallic phases and their relationship to abrasive wear [16,24,33]. The influence of volume fraction of reinforcement phase on the abrasive wear has been studied by several authors;
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Fig. 3. Worn surfaces of the untreated aluminium in (a) and (b) plan view; (c) and (d) cross-sections showing gross plastic deformation as well as embedded SiO2 abrasive particles.
Fig. 4. SEM plan view images of wear scars of the laser alloyed surfaces showing typical features observed in all the alloyed layers. Images taken at low magnifications.
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Fig. 5. SEM plan view images of wear scars of the laser alloyed surfaces showing typical features observed in all the alloyed layers. Images taken at high magnifications.
however contradictory results have been obtained. Some authors reported that as the volume fraction increased the wear resistance increased [23], while others observed the opposite trend [24]. In an attempt to address these contradictions Colaco and Vilar [16] developed a model which showed that while the increase in volume fraction of reinforcement phase led to increased hardness, the wear rates were influenced by the action of the abrasive on the reinforcements which can cause cracking and detachment of the reinforcements which may subsequently lead to increased wear rates. To a certain extent these factors determine if the alloyed aluminium surfaces would respond homogenously or heterogeneously under the abrasive wear test conditions. The type of response expected has been related to a comparison between the size of the reinforcement particles and the dimension of damage caused by the abrasive [24,28]. If the sizes are similar then a heterogeneous response is expected and the hardness properties of abrasive and reinforcement will determine the wear rates. If the size of the reinforcement particles is much smaller than the dimensions of the damage then the response is homogenous. In
the latter case the wear rates are influenced by the brittle behaviour of the particles and matrix, the hardness of the reinforcements and the abrasive as well as the volume fraction of reinforcements present. These factors influence the wear mechanisms which will prevail and determine the resulting wear rates. In the current study the worn surfaces did not display parallel grooving which is typically observed under abrasion wear testing. Rather gross plastic deformation as well as brittle fracture of the SiC and intermetallic phases was observed. Thus the dimension of damage done by the SiO2 abrasive could not be calculated. To accurately assess this aspect it should be possible to conduct scratch tests on the surfaces and to analyse the resulting wear behaviour. This was not attempted in the current study. A partial understanding of the material response can be gained by studying the worn surfaces and comparing properties between the phases present and the abrasive used. The largest reinforcement particle size present was that of the retained SiC particles which had a maximum size of at least 50 mm. The average SiO2 abrasive particle size was 525 mm. On this basis it is likely that the alloyed
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Fig. 6. SEM images of cross-sections of the worn surfaces of the laser alloyed surfaces showing fractured SiC grains and micro-fracture of the intermetallic phases.
layers responded homogenously during wear. The precise nature of the response in terms of linking increasing wear rates to particle hardness, reinforcement volume fraction and phase brittleness is more complex due to the wide range of reinforcement phases present. If the alloyed surfaces were treated as having a ‘pure’ homogeneous response and the hardness of the SiO2 abrasive, 800 HV [33], is compared to the composite hardness, which ranged from 144 to 319 HV1, the ratio of the hardness of the abrasive particles to the hardness of the alloyed surfaces (Ha/Hm) is in the range of 2.5–5.6. These calculated ratios are greater than 1.2 which signifies the boundary condition between soft and hard abrasion as defined by Hutchings [33] and would therefore place the abrasive wear of the alloyed surfaces in the hard abrasion category. This classification would confer with the extent of hard abrasion damage mechanisms observed on the wear scars of the alloyed surfaces. If the individual reinforcement hardnesses are used, then the Ha/Hm ratios span a range of values on either side of the
1.2 boundary condition and the wear rates are then influenced by the ability of the SiO2 abrasive to plastically deform, fracture or cause detachment of the particles from the matrix and would be strongly influenced by the volume fraction of each reinforcement phase present, the matrix properties and the applied stress during testing. The applied stress was calculated to be 0.02 MPa but the actual pressure is expected to be higher than this as the real contact area between individual abrasive particles and the alloyed surface is lower than the surface area of 420 mm2 used to calculate the nominal stress. For each alloyed layer produced the strength properties of the specific phases (matrix and reinforcement) present, would determine that phase’s response to the applied pressure. This would lead to different wear mechanisms being experienced by the different phases and ultimately result in a variety of measured wear rates. The influence of increasing reinforcement volume fraction showed a variable wear response and cannot be easily quantified as the alloying compositions were not kept constant for a wide
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range of compositions. However a few deductions can be made. When the Ni volume fraction was kept constant at 40 wt%, the wear rate increased as the Ti volume fraction increased along with a simultaneous decrease in the SiC volume fraction. When the Ti volume fraction was kept constant at 40 wt%, the wear rate increased with a decrease in SiC content, but no linear effect was observed as the Ni volume fraction was simultaneously decreased. When the SiC volume fraction was kept constant at 40 wt%, no linear relationships were observed with increasing the Ti and decreasing the Ni volume fractions simultaneously. It is noted that the three alloys which had the best wear resistance were all produced with a composition of 40 wt% SiC and Ti and Ni powders ranging from 20 to 40 wt%. Due to the influence of laser processing conditions on the amount of dissolution and reaction rates, the final proportions of each phase in the alloyed layers are likely to change. Thus in order to state which composition would ultimately produce the best wear resistance the final proportions of each phase in the alloyed layers should be accurately determined. Due to the complex nature of the alloy system investigated and the variety in phase morphologies present this proved to be difficult; even qualitative results by EDS is considered misleading. Abrasive wear generally occurs by plastic flow, plastic deformation, brittle fracture or combinations of the different modes [18] and in this work all modes were observed. The presence of intermetallic phases such as Al3Ni and Al3Ti together with the soft a-Al in the inter-dendritic regions constituted a favourable combination with these hard phases resisting abrasion and the soft aluminium phase suppressing macro-crack growth in these regions. The same behaviour was observed by Man et al. [36] who studied the abrasive wear of aluminium AA6061 laser surface alloyed with NiTi and reported that wear generally occurred by micro-cracking and flaking of the surface. The brittleness of the material has been shown to increase with increasing proportion of reinforcement particles [23]. In a study on the relationship between hardness and brittleness of intermetallic phases Guo et al. [38] reported that hardness and brittleness of the intermetallic phases play a role in determining which phase would have more wear resistance. Phases that have a high hardness and are less brittle will have better wear resistance but when ductility is low (for brittle phases) fracture occurs leading to high wear rates. This may be the reason for the high wear rates observed by the high-Ni alloys which showed the highest hardness, but did not have the best wear resistance mainly due to the brittle fracture of the aluminides. The role of the mean free path between reinforcements has also been reported to play a role in the wear resistance of multiphase materials. Axe´n and Zum Gahr [39] studied the sliding wear resistance of a tool steel 90MnCrV8 reinforced with TiC particles and concluded that the mean free path (defined as the average spacing between reinforcement particles) also plays an important role during wear. Matrix protection is improved by a reduced mean free path between the reinforcing particles. Large reinforcement particles result in an increased mean free path which ultimately results in reduced wear resistance. The effect of the size of the reinforcements was also found to depend on the matrix structure. Small reinforcement particles are of greater advantage within a hard matrix than large particles within a soft matrix [39]. It is well known that a high volume fraction of small particles will result in low mean free paths which would generally lead to a high wear resistance; this is based on the theory of dispersion strengthening [33]. In the current study the alloyed layers with small Al mean free paths had high wear rates as brittle fracture was promoted by the formation of high densities of the hard and brittle intermetallic and ceramic phases in which cracking of the intermetallic phases and fracturing of the SiC
particles were observed. The Ti containing and the Al–SiC intermetallic phases resulted in large Al mean free paths. This promoted wear by plastic deformation. In a few instances the hard phases, e.g. the equiaxed dendritic Al3Ni2 grains and the SiC particles, were also seen to interrupt plastic deformation. Fracture of the SiC particles and intermetallic phases can be reduced if narrower distributions of starting powders were used to ensure either complete dissolution of the SiC particles or to limit retention of larger particles. It has been shown that if the extent of cracking or extraction of reinforcement particles is negligible, then the abrasive wear resistance may be improved [16]. The hard intermetallic phases probably cracked due to repeated loading and the stresses applied during wear and these micro-cracks facilitated material removal as wear progressed. Fracturing of the embedded SiC particles also occurred due to cyclic loading but the strong bonding between the particles and the matrix ensured that decohesion (interfacial debonding and pull-out) of the SiC particles was prevented. A strong interfacial bond between the reinforcements and the matrix is important. When the bond is weak the reinforcements are easily pulled out during wear and can act as additional abrasives thereby increasing the wear rate, as was observed by reference [39]. Due to the strong interfacial bonding and prior to fracture, the embedded SiC particles in the MMCs probably led to lower wear rates in alloys where these particles were present in high volume fractions. This strong bond between the matrix and the reinforcements is formed by a reaction between the ceramic particles and the metal matrix at the metal/ceramic interface and therefore it is not easy to detach these particles from the matrix even though they have fractured. In wear tests conducted on Al–7 wt% Si laser alloyed with SiC, Anandkumar et al. [15] observed discontinuous wear grooves confined to the matrix on the worn surfaces and reported that the hard SiC reinforcement particles interrupted the abrasive grooving action. Deeper analysis is required to fully understand the response of the alloyed layers to abrasive wear and to quantify the role of each individual phase in contributing to the overall wear performance. Existing models based on the abrasive wear of multiphase materials are generally limited to the addition of one reinforcement phase in varying proportions to a constant matrix. In the current work three reinforcement materials were added to a constant matrix which produced a wide variety of reinforcement phases. Due to the complex nature of the investigated systems a model to understand the abrasive wear has not been attempted in the present work. Existing models may be used to explain the behaviour of individual phases, e.g. fracture of the SiC and microcracking of intermetallic particles. However the application of these models is limited as they do not provide a holistic view of both the individual and the combined phase response of the investigated system under abrasive wear conditions.
5. Conclusions In general laser alloying improved the wear resistance of the aluminium AA1200 surface with an improvement of 19–82% was achieved. The three alloys which had the best wear resistance were all produced with a composition of 40 wt% SiC and Ti and Ni powders ranging from 20 to 40 wt%. No direct correlation was observed between hardness and wear resistance. The alloy with the highest hardness had the ninth highest wear rate while the best wear resistance was shown by the alloy which only had the eighth highest hardness. The predominant wear mechanisms were intense plastic deformation with micro-fracture of the SiC particles and intermetallic phases. The wear behaviour is mainly determined by the response of the different alloy phases, either
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independently or in combination, to the action of the abrasive particles and the precise nature of this response in terms of linking increasing wear rates to specific phases, particle hardness, reinforcement volume fraction and phase brittleness is complex due to the wide range of reinforcement phases present and requires further study.
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