www.ptcer.pl/mccm
A
GATAG
OLONKA1*, N
ORBERTM
OSKAŁA1, P
IOTRP
UTYRA2, W
ALDEMARP
YDA11AGH University of Science and Technology, Faculty of Material Science and Ceramics, al. Mickiewicza 30,
30-059 Kraków, Poland
2Institute of Advanced Manufacturing Technology, Materials Engineering Department, ul. Wrocławska 37A,
30-011 Kraków, Poland *e-mail: ahg@agh.edu.pl
1.
Introduction
Modern materials used for manufacturing cutting tools should mainly be of the highest hardness, strength, and wear resistance. Some oxide ceramics, including alumina matrix composites conform to those requirements, and its usage is economically benefi cial.
Incorporation of a small addition of metallic nano-particles improves the properties of polycrystalline oxide ceramics due to a mechanism of strengthening which is different when compared to the mechanism attributed to metal particles of micrometric size [1-3]. When soft metal in a form of microme-tric particles, e.g. nickel, is being introduced into a ceramic matrix, e.g. alumina one, hardness of the resultant material decreases, following the law of mixtures, and the bridging mechanism of fracture toughness improvement appears.
Effects of nano-nickel and MgO addition to
alumina based matrices on properties of
the resultant composites
Abstract
Ni/Al2O3 composites were produced by using alumina nano- and micropowders, and 1.7 vol.% Ni nano-particles prepared by an oxalate
method. Uniform mixtures of alumina and nano-nickel particles were consolidated by the SPS method for 7 min at 1400 °C in argon. Effects of second-phase nano-particles, MgO-dopant and particle size of alumina powder on the microstructure, and in turn on the resultant density, hardness, fracture toughness and wear resistance of the composites were studied. XRD, SEM, EDS, image analysis, Vickers indenting and wear resistance testing were used to determine the microstructural and mechanical properties of the studied materials. The MgO-dopant controlled the microstructure of Al2O3 polycrystals limiting abnormal grain growth, but in case of the Ni/Al2O3 composites a pinning effect,
coming from nickel nano-particles, was more pronounced in reducing alumina grain sizes than a solute drag mechanism incorporated by the MgO additive. The alumina matrix composite hardened with nickel nano-particles was produced which showed a hardness of 18.3 GPa and very high wear resistance, as a result of the Hall-Petch effect.
Keywords: Al2O3, Nickel nano-particle, Nanocomposite, Microstructure - fi nal, Cutting tool
WPŁYW DODATKU NANO-NIKLU I MGO DO OSNOWY Al2O3 NA WŁAŚCIWOŚCI WYNIKOWYCH KOMPOZYTÓW
Wytworzono kompozyty Ni/Al2O3, wykorzystując nano- i mikroproszki tlenku glinu oraz nanocząstki niklu w ilości 1,7% obj., które
otrzy-mano metodą szczawianową. Jednorodne mieszaniny cząstek tlenku glinu nano-niklu konsolidowano metodą SPS przez 7 min w 1400 °C w argonie. Zbadano wpływ nanocząstek drugiej fazy, dodatku MgO i rozmiaru wyjściowego cząstek tlenku glinu na mikrostrukturę kompo-zytów i, jako skutek, na ich gęstość, twardość, odporność na pękanie i odporność na ścieranie. Wykorzystano metody XRD, SEM, EDS, analizę obrazu, metodę Vickersa i badania odporności na ścieranie, aby scharakteryzować mikrostrukturę i właściwości mechaniczne badanych materiałów. Dodatek MgO kontrolował mikrostrukturę polikryształów Al2O3, ograniczając nieciągły rozrost ziaren tlenku glinu,
ale w przypadku kompozytów okazał sie mniej skuteczny niż dodatek nanocząstek niklu, które działały poprzez mechanizm kotwiczenia granic. Wytworzono kompozyt z osnową tlenku glinu utwardzony nanocząstkami niklu o twardości 18,3 GPa i bardzo dobrej odporności na ścieranie, jako wynik działania efektu Hall-Petcha.
Słowa kluczowe: Al2O3, nanocząstka niklu, nanokompozyt, mikrostruktura fi nalna, nóż do obróbki skrawaniem
However, if a size of metal particles ranges between 20 nm and 100 nm, reinforcement is observed due to the Hall-Pech effect, when concentration of metal particles is below thre-shold of percolation of ~15%.
Such specifi c impact of nanometric particles of metallic phase in the composite caused a lot of interest for these kind of materials [1-9]. It has been shown that the effect of hardness can only be achieved by strict control of the con-centration and size distribution of the nano-particles [1-5]. In case of the Ni/Al2O3 system the biggest increase of hardness
is obtained in nanocomposites with a metallic phase content of about 3 vol.% [3]; over this threshold, the coalescence among metallic nano-particles into larger aggregates occurs, and the hardening effect is lost.
A sintering atmosphere affects the microstructure and properties of the Ni/Al2O3 composites [6-8]. Formation of
secondary phases does not occur, when the sintering is car-ried out in a reducing atmosphere, preferably hydrogen, and the highest density and hardness are obtained. Oxidation of nickel particles proceeds during sintering in oxidizing atmo-spheres [7, 8]. The NiO and NiAl2O4 phases were detected
after sintering in argon and air, respectively. The formation of both NiO and NiAl2O4 spinel makes diffi cult receiving sintered
materials of high-density, and thus hardness of the compo-sites decreases [7]. However, the appearance of the spinel phase can be advantageous in terms of fracture toughness which is improved at low spinel contents.
Other factors besides the sintering atmosphere can be listed that infl uence the microstructure, and therefore the engineering properties of the nickel-alumina composites. It is well known that the microstructure of the sintered material is infl uenced strongly by the microstructure of the green body, which in turn is dependent generally on the powder characteristics and the forming method [10]. In the case of composites, the former factor comprises physical and chemi-cal characteristics of composite powders, including particle size distributions both of matrix (primary) and reinforcement (secondary) phases, a state of agglomeration, chemical pu-rity and dopants incorporated as sintering additives.
The use of fi ne second-phase particles, and in particular, the use of dopants form two of the most effective approaches for controlling grain size during solid phase sintering due to the pinning and solute drag mechanisms, respectively [10]. In the nickel-alumina composites, the presence of nano-nic-kel particles generate strong effect on the grain boundary mobility via the pinning mechanism [e.g. 1, 9], and in turn on the microstructure with respect to the alumina grain size and pore size distributions. On the other hand, small additions of MgO (0.25 wt%) to Al2O3 allowed the production of alumina
with theoretical density and controlled grain size, preven-ting the occurrence of abnormal grain growth [10, 11]. It is interesting that the nickel/MgO-doped Al2O3 system allows
studying effects of simultaneous acting of the two mentioned approaches of the grain size control on the microstructure and related application properties.
Therefore, the aim of the work was to produce nano-nic-kel-alumina composites for studying effects of second-phase nano-particles and MgO-dopant on alumina grain sizes in the microstructure of the composites, and in turn on the resultant density, hardness, fracture toughness and wear resistance. The nickel oxalate method was used for preparation of na-no-nickel particles uniformly distributed within the alumina matrix [9] which originated from alumina powders composed of micro- or nano-particles, and was SPS consolidated.
2. Experimental procedure
Ni/Al2O3 composites containing 1.7% by volume of
me-tallic phase were studied. The composites were obtained by using α-Al2O3 powders that had characteristics shown in
Table 1 and Fig. 1. The powders differed in particle size of alumina particles, and MgO concentration added to modify the sintering process. Alcoa A16 powder is classifi ed as the submicron one, and its particle size distribution is shifted towards 7 μm when the powder is agglomerated as indicated by comparison with the BET equivalent particle size (Table 1). The agglomeration process due to the incorporation of MgO into Al16 via milling in a colloid mill in ethanol is responsible for a decrease in the specifi c surface area and an increase in the BET equivalent particles size measured for the Al16Mg powder. TM-DAR powder had the nanometric particle size (d50 = 0.15 μm), and the narrow symmetric particle size
dis-tribution with a maximum particle size of 0,35 μm.
Powders of nickel precursors were obtained by precipi-tation of nickel oxalate from nickel(II) nitrite (PA) and oxalic acid (PA), and the deposits were dried at 105 °C to the constant weight. Ref. [9] delivers details of the method. The nickel oxalate powders were incorporated to the alumina powders in the appropriate amount to obtain 1.7 vol.% of nano-nickel particles in the composite, and ground together for 1 h in a ball grinding mill of attrition type for uniformity of the mixtures. The grinding was carried out with 3Y-TZP grinding media (TOSOH) in anhydrous ethanol. The mixtures were heated for 1 h at a temperature of 360 °C in an atmo-sphere of Ar + 10% H2 to decompose nickel oxalate, and to
obtain metallic nickel nano-particles. The low decomposition temperature allows producing nickel particles of ~30 nm in size as reported in Ref. [9]. The reduction was carried out in a vacuum-pipe furnace with graphite heating elements. In this study, the SPS method (Spark Plasma Sintering) was applied to consolidate both the composite and pure alumina powders. The SPS consolidation was carried out in a ma-chine HP D5 made by FCT-System GmbH. In this method, an impulse current is used to heat up the powder being sintered. During pulsating the fl ow of current, spark discharges are generated in pores that remove from the surface of particles adsorbed gases and oxides, and facilitate the formation of active contacts between sintered particles of powder [10, 12]. It leads to the reduction of time and the decrease of sintering temperature, and is especially useful in case of the studied composite mixtures when dense composites are intended to be produced. The sintering was carried out for 7 min at 1400 °C in the argon atmosphere of purity 5.0 (99.999%).
Table 1. Chemical composition and size reduction state of matrix powders. Tabela 1. Skład chemiczny i stan rozdrobnienia proszków osnowy.
Sample Chemical composition Type/Producer/purity
Specifi c surface area [m2/g]* BET equivalent particle size [μm] Size of particles d50** [μm]
A16 100% α-Al2O3 A16-SG /Alcoa/99.8% 11.0 ± 0.2 0.14 0.34
A16Mg 99.5% α-Al2O3 + 0.5% MgO
A16-SG/Alcoa/99.8% +
MgO/Fluka/99.99% 8.3 ± 0.2 0.18 0.34
TM 100% α-Al2O3
nanometric size TM-DAR/
Taimei/99.99% 13.1 ± 0.1 0.12 0.15
The sintering was carried out under a pressure of 18 MPa. A quantitative analysis of the phase composition of sintered composites was carried out by means of an X-ray diffractometer produced by PANalytical. Monochromatic radiation with a 0.1540598 nm wavelength corresponded to the emission line Kα1 of copper; the analysis step and the range in the 2θ scale amounted to 0.008° and 20-90°, respectively. A diffractometer software based on the Rietveld refi nement was used.
Examination of the microstructure of sintered materials was performed by means of a scanning electron microscope Nova NanoSEM230 produced by FEI. Energy dispersive spectroscopy (EDS) delivered information on the chemical composition of areas that represent particular phases of the composite microstructure. The particle sizes of nickel secondary phase and zirconia contamination one were determined by using images of polished surfaces of the studied composites and a software program MATLAB for image binarization. A Saltykov method has been applied in calculations of particles size distributions [13]. The same method of mathematical calculations has been applied for semi-quantitative grain size description of fractures of the studied materials with respect to the alumina grains. This methodology introduces some systematic error connected to the assumption of equivalency between a grain cross section and a grain fragment which is visible in the fracture, but should remain observed tendencies unchanged.
Before XRD and a part SEM measurements, the surface of the sintered samples was ground with diamond abrasive materials, and fi nally polished by using a polishing cloth damped with colloidal silica: SPU – Suspension produced by Struers with grains of 1 μm in size.
Apparent densities of sintered materials were determined by means of the Archimedes’ method. Selected mechanical properties of the composites were assessed. Measurements of hardness and fracture toughness were carried out by means of a FV-700 apparatus produced by Futur-Tech from Japan. The Vickers pyramid was applied for indenting the samples under 9.807 N and 98.7 N for the measurement of hardness and fracture toughness, respectively. 10 punctures applied for 10 s each were made for each sample during both
the hardness and toughness tests. In order to determine a critical stress intensity factor KIc, the values of Young’s
modulus were estimated for the composites by using the Ravichandran’s equation [14].
Wear resistance of the sintered materials was examined by the Dry Sand test. The test refers to an ASTM G 6585 standard. The volume of material worn a sample off during 2000 rotates of a wheel lined with rubber is a measure of wear resistance in this test. A volume decrement was calculated from the difference between the sample mass before and after the test, and its apparent density.
3. Results and discussion
XRD analysis indicated the presence of α-Al2O3 and
nickel, and some amount of the tetragonal polymorph of zirconia in the studied composites (Fig. 2). NiO and NiAl2O4
were not detected. The nickel amount determined by the Rietveld method ranges from 3.1 wt% to 3.7 wt%. It agrees well with the target value of 3.6 wt% (1.7 vol.%) in the range of both batching and measurement errors, and suggests very limited oxidation of the Ni particles during the 7 min of SPS consolidation in argon if at all existed. The tetragonal ZrO2 comes from the grinding media, that wear down
dur-ing attrition milldur-ing of the composite powders. This attrition wear incorporates 1.5-2.9 wt% of zirconia particles into the studied composites.
Both fracture and polished surfaces of sintered bodies were observed to study the evolution of sizes of Al2O3 grains
and nickel particles (Figs. 3 and 4) in relation to the amount of MgO, the secondary nickel particles and the particle size of starting alumina powder. Propagation of cracks through the composite microstructure was observed on the polished surfaces of the samples (Figs. 4b, 4d, and 4f).
A microstructure of the A16 polycrystal involves the lar-gest grains (Fig. 3a). Symptoms of abnormal grain growth are also visible as a number of areas of transgranular, fl at fracture, having features of single grains in majority of cases, and a size signifi cantly larger (even 6.5 times) than the grains revealed by intergranular fracturing. The 0.5 wt% additive of MgO to Al2O3 limited grain growth during sintering in the
A16Mg bodies, eliminating almost completely the areas of transgranular fracture referred to as abnormal grains (Fig. 3b), but the reduction of alumina starting particle size down to the nanometric range decreased even stronger the alumina grain size in the sintered material (Fig. 3c). Moreover, the symptoms of abnormal grain growth remained as a feature of the microstructure originated from the alumina nanopow-der with no MgO additive. A number of grains were even 12 times larger than the average size of fi ne grains in the TM polycrystal (Fig. 3c).
Quite different behaviour is observed in the case of alumina composites containing ~1.7 vol% nickel and ~1.5 vol% tetragonal zirconia particles (Fig. 4). Very large grains and transgranular fracture are not observed and alumina grain sizes are extremely low which indicates that the total additive of ~3.2 vol% of submicrometric in size and inert Ni and zirconia particles completely eliminates abnormal grain growth during sintering, and limits Al2O3 grain sizes below
the values observed in the alumina systems modifi ed only with MgO (Fig. 3).
Fig. 1. Particle size distributions of TM-DAR Taimei and A16 Alcoa powders.
Rys. 1. Rozkłady wielkości cząstek w proszkach TM-DAR i A16 Alcoa.
The results of semi-quantitative determinations of alumi-na particle size distributions shown in Fig. 5 indicate median values of ~6.4 μm, ~3.1 μm, and ~2.1 μm for the A16, A16Mg and TM samples, respectively. The presence of secondary populations of larger transgranular fracture areas in both the alumina unmodifi ed (A16) and modifi ed with MgO (A16Mg), having the median sizes of ~10 μm and ~8,5 μm, respec-tively, is also shown. In case of the TM microstructure, the secondary population is residual.
Alumina composites originated from the A16 powder and containing submicrometric particles of Ni and ZrO2 (1.7Ni/
A16 and 1.7Ni/A16Mg) show alumina grains of ~0,6 μm in size despite of the MgO additive. But the decrease in size of the starting alumina particles brought by the application of 0.15 μm alumina powder shifted the distribution down to ~0.45 μm in case of the 1.7Ni/TM composite.
Direct SEM observations of the nano-Ni/Al2O3 composites
in the LVD mode (Fig. 4) were not capable of distinguishing nickel particles from zirconia ones, basing only on the con-trast coming from the difference in atomic numbers of Ni (28) and Zr (40). The linear EDS analysis (Fig. 6) delivered information that nickel particles were generally lower than zirconia ones. This statement allows interpreting the inclu-sion size distributions determined for the studied Ni/Al2O3
composites (Fig. 7), which were bimodal. As a consequence, the populations having modal sizes of ~70 nm and ~139 nm have been ascribed to nickel and zirconia particles, respecti-vely. The particles of nickel compared to zirconia were more spherical as indicated by the data of Fig. 6.
The nickel nano-particles have sizes that contain in a ran-ge of ~30-130 nm, and the larran-gest ones are greater in size only ~4 times than the nickel nano-particles in the composite powder (~30 nm). This can indicates that the coalescence among nickel nano-particles into larger particles (aggrega-tes) was very limited during the SPS consolidation process.
According to Fig. 7, the modal sizes of both nickel and zirconia particles are not infl uenced by the MgO-dopant as indicated by comparison of the 1.7Ni/A16 and 1.7Ni/A16Mg composites. The decrease of the alumina particle size in the starting powder reduces only a width of the zirconia particle size distribution (1.7Ni/TM).
The MgO-dopant and the nanometric size of starting alu-mina powder increased densities of the single-phase alualu-mina polycrystals up to 98.4% and 98.1%, respectively, as shown in Table 2 for the A16Mg and TM samples, when compared to the undoped A16 material which showed a density of 94.5% after 7 min SPS sintering at 1400 °C. The incorporation of nano-nickel particles permitted to increase, slightly reduce and remain unchanged densifi cation of the polycrystals de-rived from the A16, A16Mg and TM powders, respectively (Table 2). The relative density of the SPS sintered nano-Ni/ Al2O3 composites was not greater than 98.0%.
The same tendencies were observed in the case of hardness (Table 2). The measured values of hardness followed changes in density of the studied materials, with the exception of the 17Ni/TM nanocomposite, which has been manufactured from the nanometric alumina powder, and showed an extra increase in hardness up to the highest measured value of 18.3 GPa.
The 17Ni/TM nanocomposite showed also the highest increase in fracture toughness when compared to the undo-ped A16 alumina single-phase polycrystal (Table 2). A critical stress intensity factor of 6.0 MPa·m1/2 has been measured
for this material. The TM alumina single-phase polycrystal and other studied nano-Ni/Al2O3 composites showed the
KIc average values ranging from 5.6-5.8 MPa·m1/2, that
exceeded KIc values measured for the alumina single-phase
polycrystals made from the undoped (A16) and MgO-doped (A16Mg) powders.
a)
b)
c)
Fig. 2. X-ray diffraction patterns of bodies after being SPS sintered for 7 min at 1400 °C: a) 1.7Ni/A16, b) 1.7Ni/A16Mg and c) 1.7Ni/TM. Rys. 2. Dyfraktogramy rentgenowskie materiałów spiekanych me-todą SPS w 1400 °C przez 7 min: a) 1.7Ni/A16, b) 1.7Ni/A16Mg i c) 1.7Ni/TM.
a) b) c) d) e) f)
Fig. 3. SEM images of fractures of alumina polycrystals SPS sintered for 7 min at 1400 °C before being broken: a) A16 - general view, b) A16 – details of exaggerated grain growth and microcracking, c) A16Mg - general view, d) A16Mg - details of homogeneous microstructure, e) TM – general view, f) TM – details of inhomogeneities in submicrometric scale.
Rys. 3. Obrazy SEM przełomów polikrystalicznego tlenku glinu spiekanego metodą SPS przez 7 min w 1400 °C: a) A16 – widok ogólny, b) A16 – szczegóły nieciągłego rozrostu ziaren i mikropękania, c) A16Mg - widok ogólny, d) A16Mg – szczegóły jednorodnej mikrostruktury, e) TM – widok ogólny, f) TM – szczegóły niejednorodności w skali submikronowej.
a) b) c) d) e) f)
Fig. 4. SEM images of the microstructure of nano-nickel/alumina composites SPS sintered for 7 min at 1400 °C: a) fracture surface of 1.7Ni/A16, b) crack propagation in 1.7Ni/A16, c) fracture surface of 1.7Ni/A16Mg, d) crack propagation in 1.7Ni/A16Mg, e) fracture of surface 1.7Ni/TM, f) crack propagation in 1.7Ni/TM; TM corresponds to the alumina powder with the smallest initial particle size, Mg indicates the addition of 0.5 wt% MgO as a sintering additive, Ni indicates the addition of 1.7 vol.% nano-Ni particles.
Rys. 4. Mikrofotografi e SEM mikrostruktury kompozytów nano-Ni/Al2O3 spiekanych przez 7 min w 1400 °C metodą SPS: a) przełom
kom-pozytu 1.7Ni/A16, b) bieg pęknięcia w kompozycie 1.7Ni/A16, c) przełom komkom-pozytu of 1.7Ni/A16Mg, d) bieg pęknięcia w kompozycie 1.7Ni/A16Mg, e) przełom kompozytu 1.7Ni/TM, f) bieg pęknięcia w kompozycie 1.7Ni/TM; TM koresponduje z proszkiem tlenku glinu o najmniejszym rozmiarze cząstki, Mg wskazuje na MgO wprowadzone w ilości 0,5% wag. jako dodatek do spiekania, Ni wskazuje dodatek nanocząstek Ni wynoszący 1,7% obj.
Wear resistance increased signifi cantly only after the incorporation of 1.7 vol.% nano-nickel particles into the alumina matrix. A wear of 1.8 mm3 has been measured for
the 1.7Ni/TM nanocomposite when compared to 13.9 mm3
for the A16 polycrystal. The 1.7Ni/A16 and 1.7Ni/A16Mg
Fig. 5. Alumina grain size distribution in relation to initial particle size of alumina powder, MgO and nano-Ni additions in composites SPS sintered for 7 min at 1400 °C; TM corresponds to materials originated from alumina powder with the smallest initial particle size, Mg indicates the addition of 0.5 wt% MgO as a sintering additive, Ni indicates the addition of 1.7 vol.% nano-Ni particles.
Rys. 5. Rozkład wielkości ziaren tlenku glinu w zależności od wyj-ściowego rozmiaru cząstki proszku tlenku glinu i dodatków MgO i nano-Ni w kompozytach spiekanych metodą SPS przez 7 min w 1400 °C; TM wskazuje na materiały otrzymane z proszku tlenku glinu o najmniejszym rozmiarze cząstki, Mg wskazuje na MgO wpro-wadzone w ilości 0,5% wag. jako dodatek do spiekania, Ni wskazuje dodatek cząstek nano-Ni wynoszący 1,7% obj.
Fig. 6. EDS linear scan for O, Al, Zr and Ni across microstructure of 1.7Ni/A16 composite, evidencing the largest particles as being of ZrO2 s.s.
Rys. 6. Wynik liniowej analizy EDS zawartości O, Al, Zr i Ni w przy-padku kompozytu 1.7Ni/A16 wskazujący, że największe cząstki wtrąceń są roztworem stałym dwutlenku cyrkonu.
Fig. 7. Volume fractions of inclusion particles in composites as a function of size. Populations of nano-Ni and zirconia s.s. particles are indicated.
Rys. 7. Rozkłady wielkości cząstek wtrąceń w badanych kompo-zytach. Wskazano populacje cząstek nano-Ni i roztworu stałego dwutlenku cyrkonu.
nanocomposites showed larger values of wear of 3.7 mm3
and 4.2 mm3, respectively.
4. Discussion
Nano-nickel-alumina composites were successfully produced by the nickel oxalate method and SPS sintering to study the effectiveness of the solid solution and second--phase nano-particles approaches applied simultaneously for control the microstructure and the application properties. A particle size of alumina starting powder was the additional variable, infl uencing the microstructure.
4.1. Microstructure control in single phase
alumina polycrystals
Let us start the treatment of microstructural control with an analysis of grain growth in single-phase alumina polycrystals with no dopants that comprise the A16 and TM samples prepared during the experiments. The A16 and TM polycrystals originated from micro- and nano-powders of undoped alumina, respectively, and contained a number of very large grains in a fi ne-grained microstructure after SPS heat treatment. This was a result of abnormal grain growth in which the large grains had a much faster growth rate relative to the surrounding fi ne-grained matrix.
Abnormal grain growth has often been explained in terms of the grain size distribution in starting material. It was infer-red basing on the Hiller theory [10] that a grain larger than twice would grow abnormally. But it was proved for isotropic systems that size difference alone is not a suffi cient criterion for abnormal grain growth. Theoretical analysis and computer simulations show that true abnormal grain growth can occur as a result of variable (anisotropic) grain boundary energy or mobility [15-17]. Abnormal grain growth is favored for grains with boundaries that have a higher mobility or a lower energy than the surrounding matrix grains.
The specifi c ways by which variable grain boundary properties can arise in practice and initiate abnormal grain growth are as follows [10]:
Table 2. Density, hardness, fracture toughness and wear resistance of alumina matrix and Ni/Al2O3 composites SPS sintered for 7 min at
1400 °C.
Tabela 2. Gęstość, twardość, odporność na pękanie i podatność na ścieranie osnowy Al2O3 i kompozytów Ni/Al2O3 konsolidowanych
metodą SPS przez 7 min w 1400 °C.
Sample Density [%] Hardness [GPa] Fracture toughness KIc [MPa·m1/2] Wear [mm 3] A16 94.45 ± 0.01 12.2 ± 0.8 4.7 ± 0.3 13.9 A16Mg 98.43 ± 0.01 16.3 ± 0.8 5.1 ± 0.3 14.9 TM 98.11 ± 0.01 16.3 ± 1.5 5.6 ± 0.1 20.6 1.7Ni/A16 96.10 ± 0.01 15.1 ± 0.8 5.8 ± 0.2 3.7 1.7Ni/A16Mg 98.01 ± 0.01 15.4 ± 0.9 5.6 ± 0.2 4.2 1.7Ni/TM 97.99 ± 0.01 18.3 ± 1.0 6.0 ± 0.6 1.8
1. The structure and misorientation of the grain boundary. Special and low-angle grain boundaries possess a lower energy than general boundaries that may have high miso-rientation angle. Transfer of matter from the surrounding grains to the low energy boundaries results in abnormal grain growth, producing often anisotropic and faceted grains.
2. Preferential segregation of dopants and impurities to different types of boundaries can also change the relative grain boundary mobility or energy.
3. The release of solutes, second-phase particles, and pores from moving grain boundaries leads to a sudden increase in the boundary mobility.
4. The presence of liquid phase.
5. Physical and chemical inhomogeneities, such as inhomo-geneous packing and non-uniform distribution of dopants and second phase particles, commonly lead to local variation in the microstructure. The local microstructural variations produce inequalities in the boundary mobility or energy.
From the listed mechanisms, the inhomogeneous pac-king was most probably responsible for relatively stronger abnormal grain growth effects observed in the TM polycrystal when compared to the A16 one. The alumina nano-powder was used for the production of the TM polycrystal, and ge-nerally nano-powders present the smaller ability for dense and uniform particle packing than micro-powders. The SEM observations (Figs. 3a and 3e) reveal that although the grain size in the sintered polycrystal decreases with decrease in the particles size of the alumina starting powder, the size ratio of very large and fi ne grains increases from 6.5 to even 12 for the A16 and TM polycrystals, respectively.
Abnormal grain growth has been nearly completely sup-pressed in the A16Mg polycrystal by the 0.5 wt% additive of MgO (Fig. 3c). Some remains of abnormal grain growth should be attributed to Mechanism 5.
The specifi c mechanisms by which MgO acts in Al2O3 to
suppress abnormal grain growth are described in detail in Ref. [10, 18]. The single most important role of MgO is the reduction in the grain boundary mobility by a solute drag mechanism [19]. However, it is also recognized that MgO plays an important role in reducing anisotropies in the surface and grain boundary energies and mobilities [20, 21] and in stabilizing the microstructure against the consequences of inhomogeneous densifi cation and the resultant differential densifi cation during sintering.
It is worth to recall that the solute drag effect is produced by the segregation of dopants, (e.g., MgO referred to as the solute) to the grain boundaries. Symmetrical distribution of the dopant will be observed in the region of a hypothetical stationary boundary. If the boundary starts to move, the solute concentration profi le becomes asymmetric if the diffu-sion coeffi cient of the dopant atoms across the boundary is different from that of the host atoms. The asymmetric distri-bution produces a retarding force or drag on the boundary that reduces the driving force for migration. The boundary can break away from the region of increased solute con-centration (a solute cloud) when the driving force boundary migration is high enough, and then its mobility approaches the intrinsic value, Mb.
According to the Cahn’s model for grain boundary mi-gration controlled by solute drag, the boundary mobility
M
b/comprises the intrinsic component Mb and the solute drag
component Ms, following an expression in the form [22]
1 /
1
1
S b bM
M
M
(1) where
C
Q
T
k
N
D
M
b g v b S
4
(2)where Db is the diffusion coeffi cient for the solute atoms
across the boundary, Nv is the number of host atoms per
unit volume, δgb is the grain boundary width, Q is a partition
coeffi cient (>1) for the dopant distribution between the boundary region and the bulk of the grain, such that the solute concentration in the boundary region is Q∞, k is the
Boltzmann constant and T is the absolute temperature. Equ-ation (2) predicts that dopants are most effective for reducing the boundary mobility when the diffusion coeffi cient of the rate-limiting species Db is low and the segregated solute
concentration Q∞ is high.
4.2. Microstructure control in
nano-nickel-alumina composites
Because of its infl uence on the boundary mobility, the use of nano-nickel particles formed a very effective approach for controlling alumina grain size during solid-state sintering in the SPS regime. If a grain boundary moving under the driving force of its curvature encounters a suffi cient number of nickel
nanoparticles, the boundary will be pinned, and boundary migration will, therefore, cease.
The Zener model and computer simulations are the basic approaches that have been used to describe particle inhibi-ted grain growth in a system of fi ne second-phase particles dispersed randomly in a polycrystalline in which they are insoluble and immobile [10, 23].
The Zener model indicates that a limiting grain size, GL,
will be reached, the magnitude of which is proportional to the inclusion size and inversely proportional to the inclusion volume fraction, according to the following equation
f
r
G
L3
2
(3)where α is a geometrical shape factor (e.g., α = 2 for spherical grain), f - the volume fraction of the inclusions in the solid, r is the radius of the inclusion. Boundary migration will cease, when the limiting grain size GL is reached, because net
driving force per unit area of the boundary, coming from its curvature, is equal to zero. Changes in temperature would not affect this equilibrium relationship but would affect only the rate at which the system approaches the equilibrium condition. If a limiting grain size is reached, further grain growth can occur only if (1) the inclusions coarsen by Ost-wald ripening, (2) the inclusions go into solid solution in the matrix, or (3) abnormal grain growth occurs.
Computer simulations employing the Monte Carlo procedure were used by Srolovitz et al. to analyze particle inhibited grain growth [24-26]. The technique delivers another expression for the limiting grain size GL in the form
1/3f
r
K
G
b S L
(4)where KS is a constant, r is the radius of the inclusions, f is
the volume (or areal) fraction of the inclusions, and
b is the fraction of inclusions that are located at the boundaries. The simulations also reveal that the time taken to reach the limiting grain size decreases with increasing values of f, and the value of
b decreases with time. This effect becomes more signifi cant for smaller values of f.The results presented earlier indicate that the nickel nano-particles incorporated to the studied composites were in majority spherical, contained in the rather narrow size range, immobile and randomly distributed in the polycry-stalline alumina matrix, so the nickel nano-particles fulfi l the requirements of the Zener model. However, the precise calculation of limiting grain sizes for the studied Ni/Al2O3
nanocomposites is not possible due to the presence of the second-phase particles of zirconia. ZrO2 particles are quite
mobile at Al2O3 grain boundaries, and they migrate with the
boundaries [10]. The particles also tend to coalesce, so the particles size does not remain constant. It is also questiona-ble that the system approaches the equilibrium conditions during 7 min of the SPS sintering.
A rough calculation of limiting grain size give a value of 2.9 μm for the Ni/Al2O3 composite. This value is charged with
the assumption that the zirconia inclusions had the same characteristics as the nickel nanoparticles which allows summarizing volume fractions of the nickel and zirconia particles (f = 0.032) and taking r = 70 nm. Due to this and the reasons mentioned above the value of limiting grain size
is overestimated when compared to the measured alumina grain size in the studied composites (~0.6 μm for the 1.7Ni/ A16 and 1.7Ni/A16Mg composites, and ~0.45 μm for the 1.7Ni/TM composite).
It is worth also noticing that the Zener analysis overesti-mates the driving force for grain growth because it neglects the effect of the inclusions on the curvature of the boundary and considers only the work involved in dragging the inclusion [10]. As a result, for a given value of f, the Zener relationship predicts more grain growth and a larger limiting grain size.
There was no infl uence of the MgO additive on the alumi-na grain size in the studied Ni/Al2O3 composites in the range
of measurement error. This suggests much stronger infl uence of the nickel nanoparticles on grain boundary migration than the MgO dopant. The pinning effect from the nickel nanopar-ticles more effectively inhibits grain growth than MgO solute drag in the studied system, and at the applied consolidation conditions. However, the MgO dopant plays the important role as a microstructural homogenizer, reducing anisotropies in the surface and grain boundary energies, and decreasing effects of inhomogeneous densifi cation, so it should be ap-plied in the Al2O3 based systems.
The infl uence of particles size of the starting alumina powder on alumina grain size in the microstructure of sintered polycrystals has been detected in the studied Ni/Al2O3
com-posites. It proves that the non-equilibrium microstructures were obtained during the applied SPS consolidation, and the rate at which the system approaches the equilibrium microstructure depends also on the starting particle size.
4.3. Mechanical properties
The mechanical properties of the studied alumina po-lycrystals and the Ni nano-particle contained composites are strongly dependent on the microstructure. The most important feature of the microstructure controlling hardness, fracture toughness and wear resistance of the studied materials was the amount of porosity. The MgO-dopant signifi cantly controlled the evolution of microstructure in the single-phase polycrystal, but was a secondary factor in the case of composites.
The size of alumina grains appeared very important only in the case of the nano-Ni/Al2O3 composite produced using
the alumina nano-powder. The 17Ni/TM nanocomposite had the hardness of 18.3 GPa, showing an extra increase when compared to a value of 17.6 GPa calculated according the rule of mixture (the hardnesses of monolithic alumina, zirconia and nickel used for the calculation were 18.0 GPa, 10.8 GPa and 0.6 GPa [27], respectively; the average volume fraction of zirconia particles was 0.15 vol.%). This notable increase in the hardness comes from the extremely low nickel content of 1.7 vol.% in the form of nanoparticles randomly dispersed among the alumina matrix grains of ~0.45 μm in size. The same situation has been reported in Refs. [1-5] and attributed to hardening known as the Hall-Petch effect.
It is worth noticing that in order to obtain the extra incre-ase in hardness of 0.5 GPa by using only 1.7 vol.% of nickel nanoparticles, they should present a hardness of 42 GPa, which is similar to a value of 40 GPa predicted by a hardening model proposed by Pecharroman et al. [5].
5. Conclusions
Alumina matrix composites hardened with nickel na-noparticles were successfully produced. The composites contained 1.7 vol.% nickel nano-particles of ~30-130 nm uniformly distributed in the composite microstructure.
Both second-phase nano-particles and MgO-dopant affected alumina grain sizes in the microstructure of the composites. The size of starting alumina particles was also the important factor controlling the microstructure.
Abnormal grain growth was the feature of microstructures originated from the alumina nano- and micropowders with no MgO additive.
The 0.5 wt% MgO-dopant controlled the microstructure of Al2O3 polycrystals limiting abnormal grain growth, but in
case of the Ni/Al2O3 composites a pinning effect, coming from
nickel nano-particles, was more pronounced in reducing alu-mina grain sizes than a solute drag mechanism incorporated by the MgO additive.
Mechanical properties of the studied alumina polycrystals and composites depended on the microstructure, and espe-cially on the amount of porosity and the alumina grain size.
The 1.7 vol.% nano-nickel/Al2O3 composite showed
a hardness of 18.3 GPa when the alumina grain size and density were 0.45 μm and 98.0%, respectively. The hardening of the composite is here attributed to the Hall-Petch effect. The composite showed also the highest wear resistance, and it should be considered as a good candidate for cutting tool applications.
Acknowledgements
The work in part was financially supported by the European Union within a framework of European Regional Development Found under Innovative Economy Operating Programme; grant no. POIG.01.03.01-00-024/08. Some part of the work was also fi nancially supported by the statutory means of AGH WIMiC nr 11.11.160.615 in 2014. The authors would like to thank Mrs. B. Trybalska and dr. Ł. Zych for the assistance in SEM and BET measurements, respectively, and Ewa Drożdż-Cieśla for helpful discussion, concerning the preparation of nano-nickel.
References
[1] Moya, J. S., Lopez-Esteban, S., Pecharromán, C.: The chal-lenge of ceramic/metal microcomposites and nanocomposites,
Prog. Mater. Sci., 52, 7, (2007), 1017–1090.
[2] Lopez-Esteban, S., Rodriguez-Suarez, T., Esteban-Betegón, F., Pecharromán C., Moya J. S.: Mechanical properties and interfaces of zirconia/nickel in micro- and nanocomposites, J.
Mater. Sci., 41, 16, (2006), 5194-5199.
[3] Esteban-Betegón, F., Lopez-Esteban, S., Requena, J., Pechar-román, C., Moya, J. S.: Obtaining Ni Nanoparticles on 3Y-TZP Powder from Nickel Salts, J. Am. Ceram. Soc., 89, 1, (2006), 144–150.
[4] Rodriguez-Suarez, T., Díaz, L. A., Torrecillas, R., Lopez- Este-ban, S., Tuan, W. H., Nygren, M., Moya, J. S.: Alumina/tungsten nanocomposites obtained by Spark Plasma Sintering, Comp.
Sci. Tech., 69, (2009), 2467–2473.
[5] Pecharromán, C., Esteban-Betegón, F., Bartolomé, J. F., Richter, G., Moya, J. S.: Theoretical Model of Hardening in Zirconia-Nickel Nanoparticle Composites, Nano Letters, 4, (2004), 747-751.
[6] Konopka, K., Miazga, A., Właszczuk, J.: Fracture toughness of Al2O3-Ni composites with nickel aluminate spinel phase NiAl2O4,
Composites, 11:3, (2011), 197-201.
[7] Konopka, K., Kuczyński, P., Miazga, A., Riegert, D.: Effect of Sintering Atmosphere on Microstructure and Properties Al2O3-Ni
Composites, Materiały Ceramiczne/Ceramic Materials, 64, 1, (2012), 89-93.
[8] Wang, T. C., Chen, R. Z., Tuan, W. H.: Oxidation resistance of Ni-toughened Al2O3, J. Eur. Ceram. Soc., 23, (2003), 927-934.
[9] Golonka, A., Drożdż-Cieśla, E., Moskała, N., Pyda, W.: Effects of the nickel oxalate precursor and SPS consolidation process on properties of Ni/Al2O3 composites, Materiały Ceramiczne/
Ceramic Materials, 65, 2, (2013), 227-234.
[10] Rahaman M. N.: Sintering of Ceramics, CRC Press, Taylor & Francis Group, Boca Raton, FL, 2008.
[11] Coble, R. L.: Sintering crystalline solids. II Experimental test of diffusion models in powder compacts, J. Appl. Phys., 32, (1961), 793.
[12] Yao, X., Huang, Z., Chen, L., Jiang, D., Tan, S., Michel, D., Wang, G., Mazerolles, L., Pastol, J.-L.: Alumina–nickel com-posites densifi ed by spark plasma sintering, Materials Lett., 59, 18, (2005), 2314-2318.
[13] Saltykov, S. A.: Stereometric Metallography, 3rd edn., Metallurgy,
Moscow, 1970 (in Russian)
[14] Ravichandran, K. S.: Elastic Properties of Two – Phase Com-posites, J. Am. Ceram. Soc., 77, (1994), 1178–1184.
[15] Rollet. A. D., Srolovitz, D. J., Anderson, M. P.: Simulation and theory of abnormal grain growth – anisotropic grain boundary energies and mobilities, Acta Metall., 37, (1989), 1227. [16] Yang, W., Chen, L., Messing, G. L., Computer simulation of
anisotropic grain growth, Mater. Sci. Eng. A, 195, (1995), 179. [17] Kunaver, U., Kolar. D.: Three-dimensional computer simulation
of anisotropic grain growth in ceramics, Acta Mater., 46, (1998), 4629.
[18] Harmer, M.P.: A history of the role of MgO in the sintering of α-Al2O3, Ceram. Trans., 7, (1990), 13.
[19] Berry, K. A., Harmer, M. P.: Effect of MgO solute on microstruc-ture development in Al2O3, J. Am. Ceram. Soc., 69, (1986), 143.
[20] Handwerker, C. A., Dynys, J. M., Cannon, P. M., Coble, R. L., Dihedral angles in magnesia and alumina: distributions from surface thermal grooves, J. Am. Ceram. Soc., 73, (1990), 1371. [21] Bennison, S. J., Harmer, M. P., Effect of MgO solute on the
kinetics of grain growth in Al2O3, J. Am. Ceram. Soc., 66, (1983),
C90.
[22] Cahn, J. W.: The impurity drag effect in grain boundary motion, Acta Metall., 10, (1962) 789.
[23] Smith, C. S.: Grains, phases, and interfaces: an interpretation of microstructure, Trans. AIME, 175, (1948), 15.
[24] Srolovitz, D. J., Anderson, M. P., Grest, G. S., Sahni, P. S.: Computer simulation of grain growth – III. Infl uence of particles dispersion, Acta Metall., 32, (1984), 1429.
[25] Doherty, R. D., Srolovitz, D. J., Rollett, A. D., Anderson, M. P.: On the volume fraction dependence of particle-inhibited grain growth, Scripta Metall., 21, (1987), 675.
[26] Anderson, M. P., Grest, G. S., Doherty, R. D., Li, K., Srolovitz, D. J.: Inhibition of grain growth by second-phase particles: three-dimensional Monte Carlo simulations, Scripta Metall., 23, (1989), 753.
[27] Samsonov, G. V.: Handbook of the physicochemical properties
of the elements, New York, USA, IFIPlenum. 1968.
