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Wear behaviour and electrical conductivity of β-Sialon-ZrN composites fabricated by reaction bonding and gas pressure sintering process
Ceramics International 45 (2019) 2266–2274
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Ceramics International journal homepage: www.elsevier.com/locate/ceramint
Wear behaviour and electrical conductivity of β-Sialon-ZrN composites fabricated by reaction bonding and gas pressure sintering process Li Yin, Wei Gao, Mark Ian Jones
T
⁎
Department of Chemical and Materials Engineering, Faculty of Engineering, The University of Auckland, Auckland 1023, New Zealand
ARTICLE INFO
ABSTRACT
Keywords: β-Sialon-ZrN composites Wear resistance Electrical resistivity
β-Sialon-ZrN composites with different levels of substitution and ZrN content have been formed by a process of reaction bonding and post gas pressure sintering, and the wear properties and electrically conductive properties have been investigated. The results showed that an appropriate higher sintering temperature was beneficial for increasing wear resistance, and the composites with β-Sialon (Z = 1) had better wear properties than those with β-Sialon (Z = 4). The incorporation of the ZrN particles was observed to have an effect on the wear properties of the composites. The best wear properties was observed for the composites sintered at 1700 °C with lowest wear rate of 2.4 × 10−5 mm3N−1m−1 for β-Sialon (Z = 1)-20 wt% ZrN and 5.0 × 10−4 mm3N−1m−1 for β-Sialon (Z = 4)-30 wt% ZrN, respectively. The wear resistance was influenced by numerous factors, including phase composition, microstructure, hardness and fracture toughness, and the material was mainly removed by delamination, micro-fracture and micro-cracks. At a given ZrN content, a continuous electrically conductive network was formed and had an effect on the electrical resistivity of the composites, where the electrical resistivity decreased from around 1012 Ω∙m for monolithic β-Sialon to around 10 Ω∙m for 23 vol% ZrN.
1. Introduction Sialon ceramics receive significant attention in engineering applications due to their excellent physical and chemical properties, such as high strength, high hardness, good oxidation resistance, excellent wear resistance and chemical corrosion resistance [1–4]. Many efforts have been made to fabricate Sialon materials with high density and high mechanical properties to broaden their applications. As one of the most common researched Sialon materials, β-Sialon is derived from β-Si3N4 and can be described by the formula Si6-zAlzOzN8-z (where Z ranges between 0 and 4.2) due to the substitution of z Si-N bonds by z Al-O bonds [5]. Compared to silicon nitride, β-Sialon can be completely densified by conventional ceramic techniques because of the intermediate glass phase, created at the grain boundaries, which has a great effect on the mechanical properties of β-Sialon. At present, there are numerous researches focused on the fabrication and mechanical properties of Sialons using commercial Si3N4, Al2O3 and AlN powders as raw materials and sintered by hot pressing and spark plasma sintering (SPS) [6–9]. However, these raw materials are expensive and the sintering methods are not suitable for samples with complex shapes and large sizes, and require complicated equipment. Generally, Sialon and Sialon based ceramics are widely used as cutting tools and wear components, and also used for seals and ⁎
bearings, due to their excellent physical and chemical properties, including high hardness, and wear resistance [9–11]. However, application of these materials is limited due to the high cost of production and their poor machinability by traditional machining methods [4,12]. In order to overcome these drawbacks, commercial Si3N4 powder can be replaced by a cheaper Si powder [13–16] and a suitable electrically conductive phase (i.e. nitrides and carbides, including TiN, ZrN, SiC, TiC and TiCN [17–20]) can be employed to form conductive composites, which should allow for the possibility of being machined into complex shape components by electrical discharge machining (EDM). β-Sialon is described by the formula Si6-zAlzOzN8-z, where Z represents the substitution of Z Si-N bonds by Z Al-O bonds [5]. The degree of substitution in β-Sialon materials, as indicated by the Z-value, plays an important role on their physical and chemical properties, and also tailors their microstructures [21]. Jiang reported that an increase in z-value could increase crystallite size and weaken bonding strength, resulting in reduced mechanical properties [22]. Kudyba-Jansen et al. proposed a β-Sialon with z-value equal to 0.5 with a strength of 850 MPa, a fracture toughness of 4.4 MPa∙m1/2 and a hardness about 15.8 GPa [23]. β-Sialon (z = 2), synthesized by combustion synthesis and post sintered by SPS at 1600 °C under 30 MPa pressure, had a 14.8 GPa hardness and 4.4 MPa∙m1/2 [24]. Ekström et.al fabricated dense single-phase β-Sialon with various z-values by glass-encapsulated
Correspondence to: 2-6 Park Avenue - Bldg 529, Level 1, Room 131, 2-6 Park Avenue Grafton, Auckland 1023, New Zealand. E-mail address: [email protected] (M.I. Jones).
https://doi.org/10.1016/j.ceramint.2018.10.140 Received 10 July 2018; Received in revised form 17 October 2018; Accepted 17 October 2018 Available online 18 October 2018 0272-8842/ © 2018 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
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substitution on the phase, microstructure and mechanical properties, and how they influence the wear and electrically conductive properties. 2. Experiment work 2.1. Sample preparation The starting powders of the composites were β-Sialons, pre-synthesized by reaction bonding of silicon as described in reactions 1 and 2 (details of the preparation process for β-Sialons have been described in previous work [26]). 15Si + Al2O3 + AlN + 10N2(g) → 3β-Si5AlON7 (Z=1)
(1)
6Si + 4Al2O3 + 4AlN + 4N2(g) → 3β-Si2Al4O4N4 (Z=4)
(2)
The ZrN content was introduced using commercially available ZrN powders (Sigma-Aldrich Ltd., D50 ≈ 3 µm) and was varied from 0% to 50 wt% in increments of 10 wt%. Sintering additives were composed of Y2O3 and Al2O3 with a molar ratio of 3:5 (i.e. YAG composition), and a total amount of 8 wt%. The powder mixtures were ball-milled in isopropanol for 4 h with Si3N4 balls, then dried in a rotary evaporator and passed through a 250 µm sieve. The sieved powders were uniaxially pressed into pellets with sizes of ø12 mm × 7 mm, and then pressed by cold isostatic pressing at 200 MPa. The green compacts were placed in a BN crucible and sintered in a graphite pressure furnace at 1600 and 1700 °C under a nitrogen pressure of 0.7 MPa for 6 h. Following sintering, the samples were ground and subsequently polished using a 1 µm diamond suspension to remove the outer layer.
Fig. 1. The diagram for derivation of wear volume loss in the wear testing.
hot isostatic pressing method, and stated that grain size increased with increase in z-value, whereas an increase in z above 1 caused both the hardness and fracture toughness to decrease significantly [25]. As a result, although there are numerous reports focusing on the fabrication and properties of β-Sialon materials, there has been little focus on the investigation of fabrication, mechanical and wear properties of β-Sialon or β-Sialon-based composites with extremely high z-value (where z equals to 4, around the highest z = 4.2) In previously reported work [26], the authors investigated the effect of post-sintering temperature and ZrN content on the fabrication, phase assemblage and mechanical properties of β-Sialon (Z = 1)-ZrN composites produced using reaction bonding to form Sialon from inexpensive Si powders. In this work, we focus on the influence of ZrN contents and sintering temperature on the wear and electrical conductivity of the composites since these properties are essential for both the application and the machinability of the materials. Using the same techniques as previously reported, this work also investigates the formation of SiAlON-ZrN composites with different degree of substitution in the Sialon phase through modification of the z-value of β-Sialon. This allows for a comprehensive study of the effect of different levels of
2.2. Characterisation Wear tests were conducted using a computer controlled a Tribometer (T50, NANOVEA, USA) in pin-on-disk configuration under a reciprocating sliding motion. An alumina ball with a diameter of 6 mm was chosen as the friction counterpart. The wear testing was conducted under a load of 10 N for 15 min, with a sliding speed of 100 rpm under non-lubricated conditions (based on ASTM G99-17 testing standard [27]). Following the tests, wear tracks were observed by optical microscopy and the wear track width was measured using ImageJ software. Based on the Fig. 1, the wear volume and wear rate were
Fig. 2. XRD patterns of the β-Sialon (Z = 4)-ZrN composites sintered at 1600 and 1700 °C under a nitrogen pressure of 0.7 MPa; (a): 1600 °C and (b) 1700 °C.
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Table 1 Phase assemblage of the composites sintered at 1600 and 1700°C under a nitrogen pressure of 0.7 MPa, with the data of β-Sialon (Z=1)-ZrN from previous work [26]. β-Sialon
Z=1 Z=4
ZrN (wt.%)
T = 1600 °C
T=1700°C
Major phases
Intermediate phases
Major phases
Intermediate phases
0 10–50 0
β-Sialon (Z = 1) β-Sialon(Z = 1) ZrN β-Sialon (Z = 4)
Y3Al5O12
Y5Si4Al2O17N
10–50
β-Sialon (Z = 4)
15R-Sialon 21R-Sialon ZrAl3O3N Zr0.82Y0.18O1.91 Y3Al5O12
ZrN
Al2O3
β-Sialon (Z = 1) β-Sialon(Z = 1) ZrN β-Sialon (Z = 4) 15R-Sialon β-Sialon (Z = 4) ZrN 15R-Sialon
ZrAl3O3N Y4SiAlO8N Zr0.82Y0.18O1.91 Al2O3 12H-Sialon
Fig. 3. SEM morphologies of the selected samples: (a) β-Sialon (Z = 4)− 0 wt% ZrN sintered at 1600 °C, (b) and (c) β-Sialon (Z = 4)− 50 wt% ZrN sintered at 1600 and 1700 °C, and (d) β-Sialon (Z = 1)− 50 wt% ZrN sintered at 1700 °C.
calculated by the following Eqs. 1 [27] and 2 [28]:
V = Aw × L =
K=
r2 w arcsin 180 2r
V V = PD P×L×t×s
()
w 2 2
w r2 2
measuring the electrical resistivity, using the two-probe method at room temperature on disk shape samples. A precision source/measure unit (B2901A, Agilent Technologies) was used to record the I-V curve to calculate the electrical resistance. Silver paste electrodes were deposited on both sides of the samples. The phase assemblage of the polished samples were analysed by XRD diffraction with a Bruker 2D Phaser with CuKα radiation. The data were collected over the range 2θ = 10–80° with a step count of 0.2 s and a step size of 0.02°, and analysed using DIFFRAC EVA v1.4 software. A field-emission gun scanning electron microscope (ESEM, FEI QUANTA 200) was used to investigate the microstructures of the samples. The Archimedes’ method was employed to determine the bulk densities and open porosities of the polished samples. Relative density was determined by comparison to the theoretical density of the
L (1) (2)
where V is the wear volume in mm3, K is the specific wear rate in mm3 N−1 m−1, Aw is the cross-section area of the wear track, r is the radius of the ball in mm, w is the wear scar width (mm), L is the length of the track (mm), P is the load in N, t is sliding time in min and s is sliding speed in rpm. The electrical conductivity of the composites was evaluated by 2268
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Table 2 summary of the hardness and fracture toughness of the β-Sialons composites with different z-values and ZrN contents sintered at 1600 and 1700 °C under a nitrogen pressure of 0.7 MP, where the hardness and fracture toughness of the β-Sialon (Z = 1)-ZrN composites referred from previous work [26]. Properties
Hardness (GPa)
Fracture toughness (MPa·m1/2)
ZrN content (wt%)
0 10 20 30 40 50 0 10 20 30 40 50
β-Sialon(Z = 1)
β-Sialon(Z = 4)
T = 1600 °C
T = 1700 °C
T = 1600 °C
T = 1700 °C
14.84 15.23 14.84 14.52 14.16 13.22 3.47 3.76 3.96 4.02 4.35 4.64
15.25 15.69 16.05 15.38 14.96 14.41 3.74 4.35 4.55 5.00 5.23 5.35
6.27 11.06 10.49 10.92 10.12 9.95 3.03 3.50 3.73 3.65 3.73 3.78
– 11.96 12.87 12.39 10.94 11.11 – 3.33 3.72 4.01 4.06 4.12
15R-Sialon (SiAl 4 O2 N4 )+Al2 O3+N2 21R-Sialon (SiAl 6O2 N6 )+
3 *O2 2
3 3 ZrN+ *Al O3 ZrAl3O3 N+ *O2 2 4 2 ZrN+
1 9 *O2+ *Y2O3 2 82
(4) (5)
50 1 *Zr0.82Y0.18O1.91+ *N2 41 2
(6)
When the sintering temperature was increased to 1700 °C, new intermediate phases were detected. For the sample without ZrN, a new phase (Y4SiAlO8N) was observed, and can be described by reaction 7.
-Si2Al 4 O4 N4 +liquid phase 15R-Sialon (SiAl 4 O2 N4 )+Y4 SiAlO8 (7)
N+gas phase
With increasing ZrN content, the peak intensities of the β-Sialon (Z = 4) and 15R-Sialon (SiAl4O2N4) decreased, while a new phase (12H-Sialon (SiAl5O2N5)) appeared in the samples, which were formed by the following possible reactions 8 and 9:
Fig. 4. The specific wear rate of β-Sialons composites with different z-values and ZrN contents sintered at 1600 and 1700 °C under a nitrogen pressure of 0.7 MPa. The inset image shows enlarged view of the wear rate of the β-Sialon (Z = 1)-ZrN composites.
5* -Si2 Al 4 O4 N4
2*15R-Sialon (SiAl 4 O2 N4 )+Al2 O3+N2 2*12H-Sialon (SiAl5O2 N5 )+ composites determined from a rule of mixtures for each composition. Vickers hardness tests were carried out under a load of 98 N for 15 s (according to ASTM C1327-15 testing standard [29]). Fracture toughness (KIC) was determined by the indentation-fracture (IF) method under the same load, where the details of the calculation of fracture toughness were based on the work of G. R. Anstis et.al [30].
3 2
*O2 (9) Fig. 3 illustrates the morphologies of several selected composites sintered at 1600 and 1700 °C. As shown in Fig. 3d, the composites with β-Sialon (Z = 1) were composed of elongated matrix grains and granular ZrN particles, with both phases having good crystalline morphologies (discussed in detail in previous work [26]). Compared with the composites with β-Sialon (Z = 1), it could be concluded that the z-value of β-Sialon had remarkable influence on the microstructures of the composites as shown in Fig. 3a-c. At 1600 °C, the composites with βSialon (Z = 4) had poor sintering behaviour (a large amount of largesize pores with an average size ranging from 10 µm to 17 µm, depending on the ZrN contents); and it was difficult to distinguish the shape or size of Sialon matrix and reinforcing ZrN grains. At a higher sintering temperature (1700 °C), a smaller amount of pores with smaller sizes were observed because of grain boundary migration or formation of gaseous phases by thermal decomposition of β-Sialon (Z = 4) and pore filling with intermediate phases. The composites with β-Sialon (Z = 4) were composed of elongated Sialon matrix grains with average lengths between 30 and 45 µm and aspect ratios ranging from 3 to 6, and granular reinforcing particles with sizes between 6 and 10 µm. The grain sizes were larger than those of the composites with β-Sialon (Z = 1), and also had more intermediate phase (shown as the grey contrast regions in the images). The existence of a large amount of
3. Results and discussion 3.1. Phase assemblage and morphology Fig. 2 shows X-ray diffraction patterns of the composites with βSialon (Z = 4) sintered at 1600 and 1700 °C under a nitrogen pressure of 0.7 MPa. A summary of the phases identified is given in Table 1. As shown by these results, the phases present in these samples were more complex than those reported previously for Z = 1 samples [26]. This may be a result of the different starting compositions and thermal decomposition of the β-Sialon (Z = 4). At 1600 °C, the samples were composed of β-Sialon (Z = 4), ZrN, AlN-polytype Sialons and other intermediate phases (i.e. ZrAl3O3N, Zr0.82Y0.18O1.91, Y3Al5O12 and Al2O3) which were dependent on the ZrN content. The reactions that could lead to the presence of these phases are shown below:
6* -Si2 Al 4 O4 N4 4*21R-Sialon (SiAl 6O2 N6 )+8*SiO+4*O2
(8)
4*12H-Sialon (SiAl5O2 N5 )+6*SiO+3 *O2
(3) 2269
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Fig. 5. SEM images of worn surfaces of the sintered composite samples following pin-on-disk experiment at 10 N.
residual pores, large amount of intermediate phases and large grain size of the samples had a negative effect on the mechanical properties and led to decreased wear resistance (discussed below).
composites with β-Sialon (Z = 4) sintered at 1600 °C presented the poorest fracture toughness of around 3 MPa∙m1/2 with no ZrN, then increased to 3.7 MPa∙m1/2 with increase in ZrN content. The composites with β-Sialon (Z = 4) sintered at higher sintering temperature (1700 °C) showed better fracture toughness and reached the highest value around 4 MPa∙m1/2. Adding ZrN produced a positive influence on the fracture toughness; however, the increased toughening effect was only slight when the ZrN content was higher than 30 wt%. These results are at least 3 GPa in hardness and 1 MPa∙m1/2 in fracture toughness, lower than those of the composites with β-Sialon (Z = 1). Several reasons are attributed to the degradation in mechanical properties of the composites with β-Sialon (Z = 4), including more complex phase assemblage (shown in Fig. 2 and Table 1), existence of larger amount of
3.2. Hardness and fracture toughness Table 2 indicates the differences in the hardness and fracture toughness of the composites with different Z values. At 1600 °C, the monolithic β-Sialon (Z = 4) had a hardness of 6.27 GPa, then fluctuated between 10 and 11 GPa with increase in ZrN content, which was lower than the hardness of the composites sintered at 1700 °C (10–13 GPa). This illustrated that increase in sintering temperature (1700 °C) was advantageous for enhancing hardness of the composites. Similarly, the 2270
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experiment, all the wear tests were conducted under the same experimental conditions. The wear behaviour would be primarily affected by the composite's characteristics, such as, phase composition, microstructure, relative density, open porosity, hardness and fracture toughness. In previous work [26], it was described how the β-Sialon (Z = 1)ZrN composites sintered at 1600 and 1700 °C had similar phase assemblage, microstructures, relative density and open porosities, such that the difference in the wear rate of these materials was not related to those factors, and is thought to be due to the difference in mechanical properties, where the samples sintered at 1700 °C had better hardness and fracture toughness [26]. There are numerous reports focusing on the influence of elastic modulus, hardness and fracture toughness on the wear behaviour of advanced ceramic materials [6,31–34]. Several models [35] are proposed to estimate the sliding wear behaviour in the tribo-contact of brittle materials in the absence of any transfer layer formation, such as the sharp indenter model and blunt indenter models. For these two models, the wear volume of the brittle material can be estimated by following equations [35]:
Fig. 6. TGA Data for ZrN powders.
P 9/8
Vs =
1/2 5/8 KIC H
E H
4/5
S
or
Vs =
2s
2 3
3PR 4E
2/3
4 2EP 2 3 2KIC R
1
where P is the given load, R is the radius of the ball, S is the total sliding distance, H is the hardness, S is the stroke length of the sliding indenter, KIC is the fracture toughness, E is the elastic modulus, is a material property depending upon Poisson's ration of the material, and β are constants. It can be seen from these equations that a higher hardness and fracture toughness would be expected to result in a higher wear volume, or lower wear rate. As shown in Table 2, for the composites with β-Sialon (Z = 1) sintered at 1700 °C, the hardness of the samples increased with ZrN content up to 30 wt% and reached the highest hardness for 20 wt% ZrN, and then decreased for higher ZrN content. This is in agreement with the effect of ZrN content on wear rate of the β-Sialon (Z = 1)-ZrN composites, due to higher ZrN content made a degradation in hardness leading to decrease in wear rate. As seen in Fig. 4, the composites with β-Sialon (Z = 4) had far higher wear rates than those of the composites with low z-value. This could be caused by their lower hardness and fracture toughness, but may also be influenced by the more complex phase assemblage, larger amount of intermediate phases, larger amount of pores and poorer crystallization morphologies. Fig. 5 presents SEM images of the worn surface of the composites against the alumina ball under a load of 10 N, showing that tribolayers covered the worn surfaces of specimens and some scattered wear debris could be observed on the worn surfaces. Abraded grooves, micro-fracture and micro-cracks were observed on all specimens. These worn morphologies can be used to discuss the wear characteristics of the composites. During unlubricated sliding wear testing, the material removal mechanisms are predominately controlled by mechanical fracture and tribochemical reactions, which are believed to occur simultaneously and interact with each other. However, Q.T Sun et al. [36] reported the dry sliding wear behaviour of β-Sialon from room temperature to 800 °C, and stated that no new phases were detected at room temperature, indicating that no tribochemical reactions took place. This illustrated that the adhesive, delamination and abrasive wear were considered as the primary wear mechanism for the β-Sialon material. These similar wear mechanisms were reported in other composites [37]. Acikbas et al. [7] reported that during unlubricated wear tests of α/β-SiAlON-TiN composites, a fine grain size was positive for the wear
Fig. 7. Electrical resistivity of Sialons composites with different z-value and ZrN contents sintered at1600 and 1700 °C under a nitrogen pressure of 0.7 MPa.
intermediate phases and the larger amount of residual pores (shown in Fig. 3a-c). As shown in Figs. 2 and 3, a large amount of intermediate phases were formed and distributed in the composites with β-Sialon (Z = 4) to form weak interfaces between various phases, resulting in decrease in hardness and fracture toughness. 3.3. Wear properties The specific wear rates of β-Sialon composites with different z-value and ZrN contents sintered at 1600 and 1700 °C under 0.7 MPa are shown in Fig. 4. As can be seen, the β-Sialon-ZrN composites had better wear properties when sintered at a higher temperature (1700 °C) for both Z = 1 and Z = 4; and the wear rate of the composites with βSialon (Z = 4) was between 16 and 25 times larger than those of the composites with β-Sialon (Z = 1). At 1700 °C, compared to the monolithic β-Sialon, the addition of ZrN resulted in an increased wear resistance up to 30 wt% ZrN, with the lowest wear rate of 2.4 × 10−5 mm3 N−1 m−1 for the Z = 1 sample 20 wt% ZrN and 5.0 × 10−4 mm3 N−1 m−1 for Z = 4 with 30 wt% ZrN. The wear rate increased for higher ZrN content (≥ 40 wt%), which could be related to the degradation in mechanical properties. In general, wear behaviour is dependent on the material's physical and chemical characteristics, as well as the wear test parameters. In this 2271
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Fig. 8. The distribution of ZrN particles in the composites sintered at 1700 °C.
resistance due to lower formation of reaction layers (which would be broken in the form of fine grains). In the present work, the similar results were also presented, where the composites with β-Sialon (Z = 1) had greater wear resistance than those of the corresponding β-Sialon (Z = 4), due to their smaller grains shown in Fig. 3c and d. During the wear testing, the produced wear debris could accelerate the wear loss of the composites due to the three-body abrasive wear. During the sliding wear testing, temperature would increase at the local contacts between the contacting surfaces and alumina ball, due to the frictional heat. According to Sun et al. [36], at room temperature and normal load of 5 N, the calculated local sliding contact temperature for β-Sialon (Z = 1) can be 360 °C. It is generally accepted that β-Sialon material has excellent oxidation properties, implying that the frictional heat is not high enough to cause tribochemical reactions. However, the TGA data of ZrN powders in Fig. 6 shows that there was around 18 wt% mass increase during the heating process of ZrN in air atmosphere, close to the theoretical change in mass (17.1 wt%) due to oxidation as shown in reaction 10; and that this oxidation began to occur at around 400 °C.
2*ZrN+4 *O2
2*ZrO2 +N2
Such oxidation mechanisms, also observed during wear test for other materials [38], should be considered as one of the dominant wear mechanisms for the high wear rates of high ZrN composites. The oxidation of ZrN also causes the worse wear properties of the composites with a higher amount of ZrN. 3.4. Electrical conductivity Fig. 7 illustrates the electrical resistivity of the β-Sialon-ZrN composites sintered at 1600 and 1700 °C under a nitrogen pressure of 0.7 MPa, showing the β-Sialon (Z = 1)-ZrN composites sintered at 1600 and 1700 °C had the similar electrical conductivity, where the electrical resistivity of the samples without ZrN ranged between 1010 and 1012 Ω ∙m. The resistivity decreased to around 107 Ω∙m with the introduction of a small amount of ZrN (less than 30 wt%), but there was a dramatic change for higher ZrN contents, with the resistivity decreasing to around 10 Ω∙m at 40 wt% ZrN (around 23 vol% ZrN), indicating that the percolation threshold is reached at this concentration [39]. Similar behaviour was observed for the composites with β-Sialon (Z = 4) sintered at 1600 °C, where the electrical resistivity decreased from 1010 Ω∙m to 25 Ω∙m at a ZrN content of 40 wt%. However, for the β-Sialon (Z = 4)-ZrN composites sintered at higher temperature (1700 °C), the electrical resistivity was greater than 108 Ω∙m, even for the samples with 50 wt% ZrN. The distribution of ZrN particles shown in Fig. 8 demonstrates that a continuous network of particles was formed for the composites with a higher ZrN content. Combining the results in Figs. 2–4 and Tables 1 and 2, it can be considered that the electrical conductivity of the composites were predominately determined by the formation of a continuous network of ZrN particles (room-temperature electrical resistivity of ZrN is 12.0
(10)
Micro-cracks and craters could be observed on the worn surfaces and the reinforcing ZrN particles tended to be pulled out from the matrix. This resulted in the subsequent formation of wear debris on the worn surfaces of the composites and then exposed more ZrN particles, leading to three-body abrasion to accelerate wear. In this condition, it was also possible to have sufficient time to generate enough high temperature to oxidize ZrN particles to form tribochemical reaction layers and easily detachable reaction products. The greater propensity to oxidation for high ZrN composites may lead to increased wear by the removal of the oxide film by adhesive wear, whilst exposing the surface to further oxidation, resulting in further decrease in wear rate [34]. 2272
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µΩ·cm [40]). For β-Sialon (Z = 1)-ZrN composites with simple phase assemblage it was possible to form this continuous conducting network when the ZrN content is above 40 wt% (23 vol% ZrN). For the β-Sialon (Z = 4)-ZrN composites after sintering at 1600 °C, some ZrN reacted with other phases to form intermediate phases (ZrAl3O3N and Zr0.82Y0.18O1.91), but it was still possible to form the conducting network. These samples had similar conductivity as β-Sialon (Z = 1)-ZrN composites. However, as shown in Fig. 2b and Table 1, the β-Sialon (Z = 4)-ZrN composites at 1700 °C were composed of more complex phases where most of these intermediate phases have high electrical resistivity (1011 Ω∙m for Y3Al5O12 [41] and 1015 Ω∙m for Al2O3 [42]), and the continuous network may be destroyed, resulting in poor electrical conductivity. For the composites with same amount of ZrN shown in Fig. 8d and f, β-Sialon (Z = 4)-ZrN composites were comprised of larger matrix and reinforcing grains than those of β-Sialon (Z = 1)-ZrN. This might also be negative for the formation of the continuous conducting network leading to poor electrical conductivity.
[4] [5] [6] [7] [8] [9] [10] [11]
4. Conclusions [12]
We have investigated the wear and electrical properties of β-SialonZrN composites with different levels of substitution and ZrN contents, sintered at 1600 and 1700 °C under a nitrogen atmosphere. All composites showed better hardness and wear property when sintered at a higher temperature (1700 °C). At 1600 and 1700 °C, the composites with β-Sialon (Z = 1) had better wear properties than the composites with β-Sialon (Z = 4), with specific wear rate 16–25 times lower. The composites with β-Sialon (Z = 4) were composed of more complex phase assemblage (including a large amount of intermediate phases) and had poorer mechanical property, leading to the decrease of wear resistance. The addition of ZrN increases the wear properties up to a certain content, with the lowest wear rate of 5.0 × 10−4 mm3 N−1 m−1 for β-Sialon (Z = 4)-30 wt% ZrN and 2.4 × 10−5 mm3 N−1 m−1 for βSialon (Z = 1)-10 wt% ZrN. The wear rate increased for higher ZrN content, attributed to the decrease in hardness due to weaker interfaces between the matrix and reinforcing particles, and also by the oxidation of reinforcing particles. The electrical properties of the composites were mainly influenced by the ZrN content, phase assemblage and microstructures of the composites. The electrical resistivity of the composites (excluding the βSialon(Z = 4)-ZrN sintered at 1700 °C) decreased from around 1012 Ω∙m for monolithic β-Sialon to around 10 Ω∙m for the composites with more than 40 wt% ZrN (23 vol% ZrN), due to the formation of a continuous electrically conductive network. This network was not formed for βSialon(Z = 4)-ZrN composites sintered at 1700 °C, and these samples showed no step-like change in conductivity. This work shows that it was possible to fabricate β-Sialon-ZrN composites with satisfactory wear resistance and electrical properties, and which should be possible to machine by EDM for complex component shapes.
[13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26]
Acknowledgements
[27]
Li Yin would like to acknowledge the China Scholarship Council (CSC) (2013060400019) for providing a doctoral scholarship. The authors are grateful to technical staff in the Chemical & Materials Engineering Department at the University of Auckland for their assistance with preparation, property testing and data analysis.
[28] [29] [30]
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Ceramics International journal homepage: www.elsevier.com/locate/ceramint
Wear behaviour and electrical conductivity of β-Sialon-ZrN composites fabricated by reaction bonding and gas pressure sintering process Li Yin, Wei Gao, Mark Ian Jones
T
⁎
Department of Chemical and Materials Engineering, Faculty of Engineering, The University of Auckland, Auckland 1023, New Zealand
ARTICLE INFO
ABSTRACT
Keywords: β-Sialon-ZrN composites Wear resistance Electrical resistivity
β-Sialon-ZrN composites with different levels of substitution and ZrN content have been formed by a process of reaction bonding and post gas pressure sintering, and the wear properties and electrically conductive properties have been investigated. The results showed that an appropriate higher sintering temperature was beneficial for increasing wear resistance, and the composites with β-Sialon (Z = 1) had better wear properties than those with β-Sialon (Z = 4). The incorporation of the ZrN particles was observed to have an effect on the wear properties of the composites. The best wear properties was observed for the composites sintered at 1700 °C with lowest wear rate of 2.4 × 10−5 mm3N−1m−1 for β-Sialon (Z = 1)-20 wt% ZrN and 5.0 × 10−4 mm3N−1m−1 for β-Sialon (Z = 4)-30 wt% ZrN, respectively. The wear resistance was influenced by numerous factors, including phase composition, microstructure, hardness and fracture toughness, and the material was mainly removed by delamination, micro-fracture and micro-cracks. At a given ZrN content, a continuous electrically conductive network was formed and had an effect on the electrical resistivity of the composites, where the electrical resistivity decreased from around 1012 Ω∙m for monolithic β-Sialon to around 10 Ω∙m for 23 vol% ZrN.
1. Introduction Sialon ceramics receive significant attention in engineering applications due to their excellent physical and chemical properties, such as high strength, high hardness, good oxidation resistance, excellent wear resistance and chemical corrosion resistance [1–4]. Many efforts have been made to fabricate Sialon materials with high density and high mechanical properties to broaden their applications. As one of the most common researched Sialon materials, β-Sialon is derived from β-Si3N4 and can be described by the formula Si6-zAlzOzN8-z (where Z ranges between 0 and 4.2) due to the substitution of z Si-N bonds by z Al-O bonds [5]. Compared to silicon nitride, β-Sialon can be completely densified by conventional ceramic techniques because of the intermediate glass phase, created at the grain boundaries, which has a great effect on the mechanical properties of β-Sialon. At present, there are numerous researches focused on the fabrication and mechanical properties of Sialons using commercial Si3N4, Al2O3 and AlN powders as raw materials and sintered by hot pressing and spark plasma sintering (SPS) [6–9]. However, these raw materials are expensive and the sintering methods are not suitable for samples with complex shapes and large sizes, and require complicated equipment. Generally, Sialon and Sialon based ceramics are widely used as cutting tools and wear components, and also used for seals and ⁎
bearings, due to their excellent physical and chemical properties, including high hardness, and wear resistance [9–11]. However, application of these materials is limited due to the high cost of production and their poor machinability by traditional machining methods [4,12]. In order to overcome these drawbacks, commercial Si3N4 powder can be replaced by a cheaper Si powder [13–16] and a suitable electrically conductive phase (i.e. nitrides and carbides, including TiN, ZrN, SiC, TiC and TiCN [17–20]) can be employed to form conductive composites, which should allow for the possibility of being machined into complex shape components by electrical discharge machining (EDM). β-Sialon is described by the formula Si6-zAlzOzN8-z, where Z represents the substitution of Z Si-N bonds by Z Al-O bonds [5]. The degree of substitution in β-Sialon materials, as indicated by the Z-value, plays an important role on their physical and chemical properties, and also tailors their microstructures [21]. Jiang reported that an increase in z-value could increase crystallite size and weaken bonding strength, resulting in reduced mechanical properties [22]. Kudyba-Jansen et al. proposed a β-Sialon with z-value equal to 0.5 with a strength of 850 MPa, a fracture toughness of 4.4 MPa∙m1/2 and a hardness about 15.8 GPa [23]. β-Sialon (z = 2), synthesized by combustion synthesis and post sintered by SPS at 1600 °C under 30 MPa pressure, had a 14.8 GPa hardness and 4.4 MPa∙m1/2 [24]. Ekström et.al fabricated dense single-phase β-Sialon with various z-values by glass-encapsulated
Correspondence to: 2-6 Park Avenue - Bldg 529, Level 1, Room 131, 2-6 Park Avenue Grafton, Auckland 1023, New Zealand. E-mail address: [email protected] (M.I. Jones).
https://doi.org/10.1016/j.ceramint.2018.10.140 Received 10 July 2018; Received in revised form 17 October 2018; Accepted 17 October 2018 Available online 18 October 2018 0272-8842/ © 2018 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
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substitution on the phase, microstructure and mechanical properties, and how they influence the wear and electrically conductive properties. 2. Experiment work 2.1. Sample preparation The starting powders of the composites were β-Sialons, pre-synthesized by reaction bonding of silicon as described in reactions 1 and 2 (details of the preparation process for β-Sialons have been described in previous work [26]). 15Si + Al2O3 + AlN + 10N2(g) → 3β-Si5AlON7 (Z=1)
(1)
6Si + 4Al2O3 + 4AlN + 4N2(g) → 3β-Si2Al4O4N4 (Z=4)
(2)
The ZrN content was introduced using commercially available ZrN powders (Sigma-Aldrich Ltd., D50 ≈ 3 µm) and was varied from 0% to 50 wt% in increments of 10 wt%. Sintering additives were composed of Y2O3 and Al2O3 with a molar ratio of 3:5 (i.e. YAG composition), and a total amount of 8 wt%. The powder mixtures were ball-milled in isopropanol for 4 h with Si3N4 balls, then dried in a rotary evaporator and passed through a 250 µm sieve. The sieved powders were uniaxially pressed into pellets with sizes of ø12 mm × 7 mm, and then pressed by cold isostatic pressing at 200 MPa. The green compacts were placed in a BN crucible and sintered in a graphite pressure furnace at 1600 and 1700 °C under a nitrogen pressure of 0.7 MPa for 6 h. Following sintering, the samples were ground and subsequently polished using a 1 µm diamond suspension to remove the outer layer.
Fig. 1. The diagram for derivation of wear volume loss in the wear testing.
hot isostatic pressing method, and stated that grain size increased with increase in z-value, whereas an increase in z above 1 caused both the hardness and fracture toughness to decrease significantly [25]. As a result, although there are numerous reports focusing on the fabrication and properties of β-Sialon materials, there has been little focus on the investigation of fabrication, mechanical and wear properties of β-Sialon or β-Sialon-based composites with extremely high z-value (where z equals to 4, around the highest z = 4.2) In previously reported work [26], the authors investigated the effect of post-sintering temperature and ZrN content on the fabrication, phase assemblage and mechanical properties of β-Sialon (Z = 1)-ZrN composites produced using reaction bonding to form Sialon from inexpensive Si powders. In this work, we focus on the influence of ZrN contents and sintering temperature on the wear and electrical conductivity of the composites since these properties are essential for both the application and the machinability of the materials. Using the same techniques as previously reported, this work also investigates the formation of SiAlON-ZrN composites with different degree of substitution in the Sialon phase through modification of the z-value of β-Sialon. This allows for a comprehensive study of the effect of different levels of
2.2. Characterisation Wear tests were conducted using a computer controlled a Tribometer (T50, NANOVEA, USA) in pin-on-disk configuration under a reciprocating sliding motion. An alumina ball with a diameter of 6 mm was chosen as the friction counterpart. The wear testing was conducted under a load of 10 N for 15 min, with a sliding speed of 100 rpm under non-lubricated conditions (based on ASTM G99-17 testing standard [27]). Following the tests, wear tracks were observed by optical microscopy and the wear track width was measured using ImageJ software. Based on the Fig. 1, the wear volume and wear rate were
Fig. 2. XRD patterns of the β-Sialon (Z = 4)-ZrN composites sintered at 1600 and 1700 °C under a nitrogen pressure of 0.7 MPa; (a): 1600 °C and (b) 1700 °C.
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Table 1 Phase assemblage of the composites sintered at 1600 and 1700°C under a nitrogen pressure of 0.7 MPa, with the data of β-Sialon (Z=1)-ZrN from previous work [26]. β-Sialon
Z=1 Z=4
ZrN (wt.%)
T = 1600 °C
T=1700°C
Major phases
Intermediate phases
Major phases
Intermediate phases
0 10–50 0
β-Sialon (Z = 1) β-Sialon(Z = 1) ZrN β-Sialon (Z = 4)
Y3Al5O12
Y5Si4Al2O17N
10–50
β-Sialon (Z = 4)
15R-Sialon 21R-Sialon ZrAl3O3N Zr0.82Y0.18O1.91 Y3Al5O12
ZrN
Al2O3
β-Sialon (Z = 1) β-Sialon(Z = 1) ZrN β-Sialon (Z = 4) 15R-Sialon β-Sialon (Z = 4) ZrN 15R-Sialon
ZrAl3O3N Y4SiAlO8N Zr0.82Y0.18O1.91 Al2O3 12H-Sialon
Fig. 3. SEM morphologies of the selected samples: (a) β-Sialon (Z = 4)− 0 wt% ZrN sintered at 1600 °C, (b) and (c) β-Sialon (Z = 4)− 50 wt% ZrN sintered at 1600 and 1700 °C, and (d) β-Sialon (Z = 1)− 50 wt% ZrN sintered at 1700 °C.
calculated by the following Eqs. 1 [27] and 2 [28]:
V = Aw × L =
K=
r2 w arcsin 180 2r
V V = PD P×L×t×s
()
w 2 2
w r2 2
measuring the electrical resistivity, using the two-probe method at room temperature on disk shape samples. A precision source/measure unit (B2901A, Agilent Technologies) was used to record the I-V curve to calculate the electrical resistance. Silver paste electrodes were deposited on both sides of the samples. The phase assemblage of the polished samples were analysed by XRD diffraction with a Bruker 2D Phaser with CuKα radiation. The data were collected over the range 2θ = 10–80° with a step count of 0.2 s and a step size of 0.02°, and analysed using DIFFRAC EVA v1.4 software. A field-emission gun scanning electron microscope (ESEM, FEI QUANTA 200) was used to investigate the microstructures of the samples. The Archimedes’ method was employed to determine the bulk densities and open porosities of the polished samples. Relative density was determined by comparison to the theoretical density of the
L (1) (2)
where V is the wear volume in mm3, K is the specific wear rate in mm3 N−1 m−1, Aw is the cross-section area of the wear track, r is the radius of the ball in mm, w is the wear scar width (mm), L is the length of the track (mm), P is the load in N, t is sliding time in min and s is sliding speed in rpm. The electrical conductivity of the composites was evaluated by 2268
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Table 2 summary of the hardness and fracture toughness of the β-Sialons composites with different z-values and ZrN contents sintered at 1600 and 1700 °C under a nitrogen pressure of 0.7 MP, where the hardness and fracture toughness of the β-Sialon (Z = 1)-ZrN composites referred from previous work [26]. Properties
Hardness (GPa)
Fracture toughness (MPa·m1/2)
ZrN content (wt%)
0 10 20 30 40 50 0 10 20 30 40 50
β-Sialon(Z = 1)
β-Sialon(Z = 4)
T = 1600 °C
T = 1700 °C
T = 1600 °C
T = 1700 °C
14.84 15.23 14.84 14.52 14.16 13.22 3.47 3.76 3.96 4.02 4.35 4.64
15.25 15.69 16.05 15.38 14.96 14.41 3.74 4.35 4.55 5.00 5.23 5.35
6.27 11.06 10.49 10.92 10.12 9.95 3.03 3.50 3.73 3.65 3.73 3.78
– 11.96 12.87 12.39 10.94 11.11 – 3.33 3.72 4.01 4.06 4.12
15R-Sialon (SiAl 4 O2 N4 )+Al2 O3+N2 21R-Sialon (SiAl 6O2 N6 )+
3 *O2 2
3 3 ZrN+ *Al O3 ZrAl3O3 N+ *O2 2 4 2 ZrN+
1 9 *O2+ *Y2O3 2 82
(4) (5)
50 1 *Zr0.82Y0.18O1.91+ *N2 41 2
(6)
When the sintering temperature was increased to 1700 °C, new intermediate phases were detected. For the sample without ZrN, a new phase (Y4SiAlO8N) was observed, and can be described by reaction 7.
-Si2Al 4 O4 N4 +liquid phase 15R-Sialon (SiAl 4 O2 N4 )+Y4 SiAlO8 (7)
N+gas phase
With increasing ZrN content, the peak intensities of the β-Sialon (Z = 4) and 15R-Sialon (SiAl4O2N4) decreased, while a new phase (12H-Sialon (SiAl5O2N5)) appeared in the samples, which were formed by the following possible reactions 8 and 9:
Fig. 4. The specific wear rate of β-Sialons composites with different z-values and ZrN contents sintered at 1600 and 1700 °C under a nitrogen pressure of 0.7 MPa. The inset image shows enlarged view of the wear rate of the β-Sialon (Z = 1)-ZrN composites.
5* -Si2 Al 4 O4 N4
2*15R-Sialon (SiAl 4 O2 N4 )+Al2 O3+N2 2*12H-Sialon (SiAl5O2 N5 )+ composites determined from a rule of mixtures for each composition. Vickers hardness tests were carried out under a load of 98 N for 15 s (according to ASTM C1327-15 testing standard [29]). Fracture toughness (KIC) was determined by the indentation-fracture (IF) method under the same load, where the details of the calculation of fracture toughness were based on the work of G. R. Anstis et.al [30].
3 2
*O2 (9) Fig. 3 illustrates the morphologies of several selected composites sintered at 1600 and 1700 °C. As shown in Fig. 3d, the composites with β-Sialon (Z = 1) were composed of elongated matrix grains and granular ZrN particles, with both phases having good crystalline morphologies (discussed in detail in previous work [26]). Compared with the composites with β-Sialon (Z = 1), it could be concluded that the z-value of β-Sialon had remarkable influence on the microstructures of the composites as shown in Fig. 3a-c. At 1600 °C, the composites with βSialon (Z = 4) had poor sintering behaviour (a large amount of largesize pores with an average size ranging from 10 µm to 17 µm, depending on the ZrN contents); and it was difficult to distinguish the shape or size of Sialon matrix and reinforcing ZrN grains. At a higher sintering temperature (1700 °C), a smaller amount of pores with smaller sizes were observed because of grain boundary migration or formation of gaseous phases by thermal decomposition of β-Sialon (Z = 4) and pore filling with intermediate phases. The composites with β-Sialon (Z = 4) were composed of elongated Sialon matrix grains with average lengths between 30 and 45 µm and aspect ratios ranging from 3 to 6, and granular reinforcing particles with sizes between 6 and 10 µm. The grain sizes were larger than those of the composites with β-Sialon (Z = 1), and also had more intermediate phase (shown as the grey contrast regions in the images). The existence of a large amount of
3. Results and discussion 3.1. Phase assemblage and morphology Fig. 2 shows X-ray diffraction patterns of the composites with βSialon (Z = 4) sintered at 1600 and 1700 °C under a nitrogen pressure of 0.7 MPa. A summary of the phases identified is given in Table 1. As shown by these results, the phases present in these samples were more complex than those reported previously for Z = 1 samples [26]. This may be a result of the different starting compositions and thermal decomposition of the β-Sialon (Z = 4). At 1600 °C, the samples were composed of β-Sialon (Z = 4), ZrN, AlN-polytype Sialons and other intermediate phases (i.e. ZrAl3O3N, Zr0.82Y0.18O1.91, Y3Al5O12 and Al2O3) which were dependent on the ZrN content. The reactions that could lead to the presence of these phases are shown below:
6* -Si2 Al 4 O4 N4 4*21R-Sialon (SiAl 6O2 N6 )+8*SiO+4*O2
(8)
4*12H-Sialon (SiAl5O2 N5 )+6*SiO+3 *O2
(3) 2269
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Fig. 5. SEM images of worn surfaces of the sintered composite samples following pin-on-disk experiment at 10 N.
residual pores, large amount of intermediate phases and large grain size of the samples had a negative effect on the mechanical properties and led to decreased wear resistance (discussed below).
composites with β-Sialon (Z = 4) sintered at 1600 °C presented the poorest fracture toughness of around 3 MPa∙m1/2 with no ZrN, then increased to 3.7 MPa∙m1/2 with increase in ZrN content. The composites with β-Sialon (Z = 4) sintered at higher sintering temperature (1700 °C) showed better fracture toughness and reached the highest value around 4 MPa∙m1/2. Adding ZrN produced a positive influence on the fracture toughness; however, the increased toughening effect was only slight when the ZrN content was higher than 30 wt%. These results are at least 3 GPa in hardness and 1 MPa∙m1/2 in fracture toughness, lower than those of the composites with β-Sialon (Z = 1). Several reasons are attributed to the degradation in mechanical properties of the composites with β-Sialon (Z = 4), including more complex phase assemblage (shown in Fig. 2 and Table 1), existence of larger amount of
3.2. Hardness and fracture toughness Table 2 indicates the differences in the hardness and fracture toughness of the composites with different Z values. At 1600 °C, the monolithic β-Sialon (Z = 4) had a hardness of 6.27 GPa, then fluctuated between 10 and 11 GPa with increase in ZrN content, which was lower than the hardness of the composites sintered at 1700 °C (10–13 GPa). This illustrated that increase in sintering temperature (1700 °C) was advantageous for enhancing hardness of the composites. Similarly, the 2270
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experiment, all the wear tests were conducted under the same experimental conditions. The wear behaviour would be primarily affected by the composite's characteristics, such as, phase composition, microstructure, relative density, open porosity, hardness and fracture toughness. In previous work [26], it was described how the β-Sialon (Z = 1)ZrN composites sintered at 1600 and 1700 °C had similar phase assemblage, microstructures, relative density and open porosities, such that the difference in the wear rate of these materials was not related to those factors, and is thought to be due to the difference in mechanical properties, where the samples sintered at 1700 °C had better hardness and fracture toughness [26]. There are numerous reports focusing on the influence of elastic modulus, hardness and fracture toughness on the wear behaviour of advanced ceramic materials [6,31–34]. Several models [35] are proposed to estimate the sliding wear behaviour in the tribo-contact of brittle materials in the absence of any transfer layer formation, such as the sharp indenter model and blunt indenter models. For these two models, the wear volume of the brittle material can be estimated by following equations [35]:
Fig. 6. TGA Data for ZrN powders.
P 9/8
Vs =
1/2 5/8 KIC H
E H
4/5
S
or
Vs =
2s
2 3
3PR 4E
2/3
4 2EP 2 3 2KIC R
1
where P is the given load, R is the radius of the ball, S is the total sliding distance, H is the hardness, S is the stroke length of the sliding indenter, KIC is the fracture toughness, E is the elastic modulus, is a material property depending upon Poisson's ration of the material, and β are constants. It can be seen from these equations that a higher hardness and fracture toughness would be expected to result in a higher wear volume, or lower wear rate. As shown in Table 2, for the composites with β-Sialon (Z = 1) sintered at 1700 °C, the hardness of the samples increased with ZrN content up to 30 wt% and reached the highest hardness for 20 wt% ZrN, and then decreased for higher ZrN content. This is in agreement with the effect of ZrN content on wear rate of the β-Sialon (Z = 1)-ZrN composites, due to higher ZrN content made a degradation in hardness leading to decrease in wear rate. As seen in Fig. 4, the composites with β-Sialon (Z = 4) had far higher wear rates than those of the composites with low z-value. This could be caused by their lower hardness and fracture toughness, but may also be influenced by the more complex phase assemblage, larger amount of intermediate phases, larger amount of pores and poorer crystallization morphologies. Fig. 5 presents SEM images of the worn surface of the composites against the alumina ball under a load of 10 N, showing that tribolayers covered the worn surfaces of specimens and some scattered wear debris could be observed on the worn surfaces. Abraded grooves, micro-fracture and micro-cracks were observed on all specimens. These worn morphologies can be used to discuss the wear characteristics of the composites. During unlubricated sliding wear testing, the material removal mechanisms are predominately controlled by mechanical fracture and tribochemical reactions, which are believed to occur simultaneously and interact with each other. However, Q.T Sun et al. [36] reported the dry sliding wear behaviour of β-Sialon from room temperature to 800 °C, and stated that no new phases were detected at room temperature, indicating that no tribochemical reactions took place. This illustrated that the adhesive, delamination and abrasive wear were considered as the primary wear mechanism for the β-Sialon material. These similar wear mechanisms were reported in other composites [37]. Acikbas et al. [7] reported that during unlubricated wear tests of α/β-SiAlON-TiN composites, a fine grain size was positive for the wear
Fig. 7. Electrical resistivity of Sialons composites with different z-value and ZrN contents sintered at1600 and 1700 °C under a nitrogen pressure of 0.7 MPa.
intermediate phases and the larger amount of residual pores (shown in Fig. 3a-c). As shown in Figs. 2 and 3, a large amount of intermediate phases were formed and distributed in the composites with β-Sialon (Z = 4) to form weak interfaces between various phases, resulting in decrease in hardness and fracture toughness. 3.3. Wear properties The specific wear rates of β-Sialon composites with different z-value and ZrN contents sintered at 1600 and 1700 °C under 0.7 MPa are shown in Fig. 4. As can be seen, the β-Sialon-ZrN composites had better wear properties when sintered at a higher temperature (1700 °C) for both Z = 1 and Z = 4; and the wear rate of the composites with βSialon (Z = 4) was between 16 and 25 times larger than those of the composites with β-Sialon (Z = 1). At 1700 °C, compared to the monolithic β-Sialon, the addition of ZrN resulted in an increased wear resistance up to 30 wt% ZrN, with the lowest wear rate of 2.4 × 10−5 mm3 N−1 m−1 for the Z = 1 sample 20 wt% ZrN and 5.0 × 10−4 mm3 N−1 m−1 for Z = 4 with 30 wt% ZrN. The wear rate increased for higher ZrN content (≥ 40 wt%), which could be related to the degradation in mechanical properties. In general, wear behaviour is dependent on the material's physical and chemical characteristics, as well as the wear test parameters. In this 2271
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Fig. 8. The distribution of ZrN particles in the composites sintered at 1700 °C.
resistance due to lower formation of reaction layers (which would be broken in the form of fine grains). In the present work, the similar results were also presented, where the composites with β-Sialon (Z = 1) had greater wear resistance than those of the corresponding β-Sialon (Z = 4), due to their smaller grains shown in Fig. 3c and d. During the wear testing, the produced wear debris could accelerate the wear loss of the composites due to the three-body abrasive wear. During the sliding wear testing, temperature would increase at the local contacts between the contacting surfaces and alumina ball, due to the frictional heat. According to Sun et al. [36], at room temperature and normal load of 5 N, the calculated local sliding contact temperature for β-Sialon (Z = 1) can be 360 °C. It is generally accepted that β-Sialon material has excellent oxidation properties, implying that the frictional heat is not high enough to cause tribochemical reactions. However, the TGA data of ZrN powders in Fig. 6 shows that there was around 18 wt% mass increase during the heating process of ZrN in air atmosphere, close to the theoretical change in mass (17.1 wt%) due to oxidation as shown in reaction 10; and that this oxidation began to occur at around 400 °C.
2*ZrN+4 *O2
2*ZrO2 +N2
Such oxidation mechanisms, also observed during wear test for other materials [38], should be considered as one of the dominant wear mechanisms for the high wear rates of high ZrN composites. The oxidation of ZrN also causes the worse wear properties of the composites with a higher amount of ZrN. 3.4. Electrical conductivity Fig. 7 illustrates the electrical resistivity of the β-Sialon-ZrN composites sintered at 1600 and 1700 °C under a nitrogen pressure of 0.7 MPa, showing the β-Sialon (Z = 1)-ZrN composites sintered at 1600 and 1700 °C had the similar electrical conductivity, where the electrical resistivity of the samples without ZrN ranged between 1010 and 1012 Ω ∙m. The resistivity decreased to around 107 Ω∙m with the introduction of a small amount of ZrN (less than 30 wt%), but there was a dramatic change for higher ZrN contents, with the resistivity decreasing to around 10 Ω∙m at 40 wt% ZrN (around 23 vol% ZrN), indicating that the percolation threshold is reached at this concentration [39]. Similar behaviour was observed for the composites with β-Sialon (Z = 4) sintered at 1600 °C, where the electrical resistivity decreased from 1010 Ω∙m to 25 Ω∙m at a ZrN content of 40 wt%. However, for the β-Sialon (Z = 4)-ZrN composites sintered at higher temperature (1700 °C), the electrical resistivity was greater than 108 Ω∙m, even for the samples with 50 wt% ZrN. The distribution of ZrN particles shown in Fig. 8 demonstrates that a continuous network of particles was formed for the composites with a higher ZrN content. Combining the results in Figs. 2–4 and Tables 1 and 2, it can be considered that the electrical conductivity of the composites were predominately determined by the formation of a continuous network of ZrN particles (room-temperature electrical resistivity of ZrN is 12.0
(10)
Micro-cracks and craters could be observed on the worn surfaces and the reinforcing ZrN particles tended to be pulled out from the matrix. This resulted in the subsequent formation of wear debris on the worn surfaces of the composites and then exposed more ZrN particles, leading to three-body abrasion to accelerate wear. In this condition, it was also possible to have sufficient time to generate enough high temperature to oxidize ZrN particles to form tribochemical reaction layers and easily detachable reaction products. The greater propensity to oxidation for high ZrN composites may lead to increased wear by the removal of the oxide film by adhesive wear, whilst exposing the surface to further oxidation, resulting in further decrease in wear rate [34]. 2272
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µΩ·cm [40]). For β-Sialon (Z = 1)-ZrN composites with simple phase assemblage it was possible to form this continuous conducting network when the ZrN content is above 40 wt% (23 vol% ZrN). For the β-Sialon (Z = 4)-ZrN composites after sintering at 1600 °C, some ZrN reacted with other phases to form intermediate phases (ZrAl3O3N and Zr0.82Y0.18O1.91), but it was still possible to form the conducting network. These samples had similar conductivity as β-Sialon (Z = 1)-ZrN composites. However, as shown in Fig. 2b and Table 1, the β-Sialon (Z = 4)-ZrN composites at 1700 °C were composed of more complex phases where most of these intermediate phases have high electrical resistivity (1011 Ω∙m for Y3Al5O12 [41] and 1015 Ω∙m for Al2O3 [42]), and the continuous network may be destroyed, resulting in poor electrical conductivity. For the composites with same amount of ZrN shown in Fig. 8d and f, β-Sialon (Z = 4)-ZrN composites were comprised of larger matrix and reinforcing grains than those of β-Sialon (Z = 1)-ZrN. This might also be negative for the formation of the continuous conducting network leading to poor electrical conductivity.
[4] [5] [6] [7] [8] [9] [10] [11]
4. Conclusions [12]
We have investigated the wear and electrical properties of β-SialonZrN composites with different levels of substitution and ZrN contents, sintered at 1600 and 1700 °C under a nitrogen atmosphere. All composites showed better hardness and wear property when sintered at a higher temperature (1700 °C). At 1600 and 1700 °C, the composites with β-Sialon (Z = 1) had better wear properties than the composites with β-Sialon (Z = 4), with specific wear rate 16–25 times lower. The composites with β-Sialon (Z = 4) were composed of more complex phase assemblage (including a large amount of intermediate phases) and had poorer mechanical property, leading to the decrease of wear resistance. The addition of ZrN increases the wear properties up to a certain content, with the lowest wear rate of 5.0 × 10−4 mm3 N−1 m−1 for β-Sialon (Z = 4)-30 wt% ZrN and 2.4 × 10−5 mm3 N−1 m−1 for βSialon (Z = 1)-10 wt% ZrN. The wear rate increased for higher ZrN content, attributed to the decrease in hardness due to weaker interfaces between the matrix and reinforcing particles, and also by the oxidation of reinforcing particles. The electrical properties of the composites were mainly influenced by the ZrN content, phase assemblage and microstructures of the composites. The electrical resistivity of the composites (excluding the βSialon(Z = 4)-ZrN sintered at 1700 °C) decreased from around 1012 Ω∙m for monolithic β-Sialon to around 10 Ω∙m for the composites with more than 40 wt% ZrN (23 vol% ZrN), due to the formation of a continuous electrically conductive network. This network was not formed for βSialon(Z = 4)-ZrN composites sintered at 1700 °C, and these samples showed no step-like change in conductivity. This work shows that it was possible to fabricate β-Sialon-ZrN composites with satisfactory wear resistance and electrical properties, and which should be possible to machine by EDM for complex component shapes.
[13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26]
Acknowledgements
[27]
Li Yin would like to acknowledge the China Scholarship Council (CSC) (2013060400019) for providing a doctoral scholarship. The authors are grateful to technical staff in the Chemical & Materials Engineering Department at the University of Auckland for their assistance with preparation, property testing and data analysis.
[28] [29] [30]
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